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Hiroki Nara
Page 1
Development of Lithium Battery Materials
Using Phase Separated Electrolytes
相分離電解質を用いたリチウム二次電池材料の開発
February, 2008
by
Hiroki Nara
Graduate School of Science and Engineering, Waseda University
Applied chemistry major

Page 2

Page 3
i
Preface
Recently, the development of rechargeable batteries indicating higher
performance are strongly desired as the power supplies of mobile equipment
evolving remarkably, such as cellular phones, lap top PC and so on. The batteries
with high energy density should be achieved when lithium metal anode is applied to
batteries; because the redox potential of lithium is lowest, indicating highest
operation voltage, and electrochemical equivalent of lithium is large, indicating
high gravimetric capacity. The primary batteries with Li metal anode has been
put into practical use, and used widely. However, when the Li metal is applied to
anode for the secondary battery and charge-discharge cycles are repeated, the
dendritic lithium is deposited on the anode. The dendritic deposition of lithium
could cause the short-circuit in the battery, resulting in ignition and explosion, and
the capacity loss, so the lithium metal secondary batteries have not put into
practical use. Thus, new electrolytes and anodes have been investigated because
of the dendritic deposition control of Li.
In the field of electrolytes, polymer electrolytes and inorganic solid
electrolytes are actively investigated as substitutes of the organic electrolytes with
the danger of the ignition and the explosion by the short-circuit. Especially, the
polymer gel electrolytes with relatively higher ionic conductivity are most
practicable in these electrolytes at present. In addition, electrolytes are required
to be enough strength mechanically for physical suppression of Li dendritic
deposition and formability. However the ionic conductivity and the mechanical
strength of polymer electrolytes are incompatible characters; it is difficult to obtain
the polymer electrolyte with both high ionic conductivity and good mechanical
strength. As the preparation technique of electrolytes with the both
characteristics, the polymer blend method to obtain the polymer blend that has
mutual co-continuous micro phase separated structures is proposed, and a PEO/PS
polymer blend gel electrolyte that consists of Polyethylene oxide (PEO) and
Polystyrene (PS) is reported as an electrolyte that meets the requirement.
However the PEO/PS polymer blend gel electrolyte has an apprehension to lose
PEO, which is the ionic conductive phase, from PS matrix by the dissolution of PEO
with the addition of a large amount of plasticizer.
In the field of anodes, the carbon anode where the dendritic deposition was
controlled by storing the lithium as not the lithium metal but lithium ion has
appeared on the market widely as a product. However its capacity at operation
has almost reached the theoretical capacity of 372 mAh / g. Thus the

Page 4
ii
developments of new anode materials for higher capacity batteries are desired.
Tin anode has attracted attention as an alloy anode system, because of its high
theoretical capacity of 994 mAh / g, about 2.7 times the carbon anode. The
expansion and constriction of tin anode at alloying and dealloying with lithium lead
to its disintegration: as the results, a remarkable decrease in discharge capacity
because of the electrical charge and discharge cycle occurs. The disintegration
with the charge and discharge cycle is a matter for its practical application.
In this thesis, phase separated electrolytes were adopted to the material
preparation technique, to improve the retention of ionic conductive phase on the
polymer gel electrolyte, composed of PEO and PS, and charge-discharge cycle
durability of a tin anode. On polymer gel electrolytes, to improve the retention of
ionic conductive phase, two approaches were attempted. One is the application of
a PEO-LiX complex to the ionic conductive phase in the polymer blend gel
electrolyte. The other is the adoption of a PEO-PS di-block copolymer, composed of
PEO chain and PS chain covalently bonded, instead of blended PEO and PS. On
tin anodes, to improve its charge-discharge cycle durability, a porous tin anode was
electrodeposited with phase separated electrolyte bath composed of a surfactant and
a Sn solution. These obtained materials were characterized and discussed.

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iii
Outline of this thesis
Chapter 1 is a general introduction to the thesis. Progress of lithium
secondary batteries, especially electrolytes and anodes in development of secondary
batteries with high power and capacity and their issues were described to clarify the
standpoint of this thesis.
In chapter 2, PEO-LiBF4 complex was applied to the ionic conductive phase
in the polymer blend gel electrolyte by the addition of LiBF4 in hot-blending process.
The PEO-LiBF4 complex formation was confirmed with Raman spectra. The
mutual co-continuous micro phase separated structures was also indicated in the
case of the PEO-LiBF4/PS polymer blend system. The electrochemical properties
required as an electrolyte for lithium batteries were characterized.
In chapter 3, the PEO-PS di-block copolymer was employed as polymer
electrolyte material. The PEO-PS diblock copolymer gel electrolyte was
demonstrated to have a micro phase separated structure by its self-assembling; and
prevented PEO phase to flow out from PS matrix when soaked in plasticizer for
PEO. The PEO-PS di-block copolymer gel electrolyte was characterized. It was
revealed that an increase of the interfacial resistance between electrolyte and
lithium originated in liquefied PEO, and the fixed PEO on PS matrix was not affect
the increase of the interfacial resistance.
In chapter 4, the potential of mesoporous tin anode to improve its cycle
durability and rate property is discussed. The mesoporous structure was
introduced to tin anode by electrodepositing with phase separated electrolyte bath
composed of a surfactant and a Sn solution. The reason of improvement on the
cycle durability and rate property of tin anode was discussed.
Chapter 5 summarizes the results from chapter 2 to chapter 4, and finally
indicates the future perspectives of this research.
After the final chapter, a work that is not the subject of screening is added
as an appendix. The downsizing possibility of a hybrid power source composed of a
direct methanol fuel cell and an electric double layer capacitor is discussed by
numerical simulation.

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iv
Contents
Preface......................................................................................................................i
Chapter 1:
General Introduction................................................................................................... 1
1.1 Background............................................................................................................ 3
1.2 Polymer electrolytes .............................................................................................. 6
1.2.1 Basic properties of polymer electrolytes.......................................................... 6
1.2.2 Polymer gel electrolytes................................................................................... 9
1.3 Anode materials................................................................................................... 16
1.3.1 Carbon anodes ............................................................................................... 16
1.3.2 Alloy anodes ................................................................................................... 19
1.4 Phase separated structure .................................................................................. 26
1.4.1 Polymer blend ................................................................................................ 26
1.4.2 Liquid crystals ............................................................................................... 30
1.5 Summary ............................................................................................................. 37
Reference ................................................................................................................... 38

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v
Chapter 2:
PEO / PS polymer blend gel electrolyte .................................................................... 47
2.1 Introduction ......................................................................................................... 49
2.2 Experimental ....................................................................................................... 52
2.3 Results and Discussion........................................................................................ 54
2.3.1 Confirmation of phase separated structure .................................................. 54
2.3.2 Confirmation of PEO-LiBF4 complex formation............................................ 55
2.3.3 Electrochemical properties of IPN gel electrolyte......................................... 57
2.4 Summary ............................................................................................................. 68
Reference ................................................................................................................... 69
Chapter 3:
PEO-PS diblock polymer gel electrolyte................................................................. 71
3.1 Introduction ......................................................................................................... 73
3.2 Experimental ....................................................................................................... 74
3.3 Results and Discussion........................................................................................ 76
3.3.1 Confirmation of phase separated structure .................................................. 76
3.3.2 Electrochemical properties of IPN gel electrolyte......................................... 77
3.4 Summary ......................................................................................................... 91
Reference ................................................................................................................... 92

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vi
Chapter 4:
Mesoporous Sn Anode electrodeposited with lyotropic liquid crystals .................... 93
4.1 Introduction ......................................................................................................... 95
4.2 Experimental ....................................................................................................... 96
4.3 Results and Discussion........................................................................................ 99
4.3.1 Confirmation of mesoporous structure.......................................................... 99
4.3.2 Electrochemical properties of mesoporous Sn anode .................................. 101
4.4 Summary ........................................................................................................... 113
Reference ................................................................................................................. 114
Chapter 5:
General Conclusion ................................................................................................. 117
Appendix ............................................................................................................. 123
Trial for Hybrid System of Fuel cell and Capacitor
List of Achievements ........................................................................................... 141
Acknowledgements.............................................................................................. 147

Page 9
Chapter 1:
General Introduction

Page 10

Page 11
Chapter 1
3
1.1. Back Ground(1, 2)
A battery is composed of several electrochemical cells that are connected in
series and/or in parallel to provide the required voltage and capacity, respectively.
Each cell consists of a positive and a negative electrode (both sources of chemical
reactions) separated by an electrolyte solution containing dissociated salts, which
enable ion transfer between the two electrodes. Once these electrodes are
connected externally, the chemical reactions proceed in tandem at both electrodes,
thereby liberating electrons and enabling the current to be tapped by the user.
The amount of electrical energy, expressed either per unit of weight (Wh/kg) or per
unit of volume (Wh/l), that a battery is able to deliver is a function of the cell
potential (V) and capacity (Ah/kg), both of which are linked directly to the chemistry
of the system. The cell potential is determined by Gibbs free energy change, -ΔG.
G-
zEF
=
where z is electron number for reaction, F is Faraday constant.
When the anode materials which have strong reducing power and the cathode
materials which have strong oxidizing power are combined, -∆G increases, and the
electromotive force increases. The reducing and oxidizing power are expressed by
the electrode potential, therefore the electromotive force is determined by the
potential difference between the anode and cathode.
The capacity of a battery is determined by electrochemical equivalent,
because the electrochemical reaction follows Faraday law. Therefore the high
electromotive force and capacity battery that is high energy density battery can be
achieved by applying the anode which has low redox potential and small
electrochemical equivalent and the cathode which has high redox potential and
small electrochemical equivalent.
As described above, the voltage and current of a battery is directly related
to the types of materials applied in the electrodes and electrolyte. The propensity
of an individual metal or metal compound to gain or lose electrons in relation to
another material is known as its electrode potential. Thus the strength of
oxidizing and reducing agents are indicated by their standard potential.
Compounds with a positive electrode potential are used for anodes and those with a
negative electrode potential for cathodes. The larger the difference between the
electrode potentials of the anode and cathode, the greater the electromotive force of
the battery and the greater the amount of energy that can be produced by the
battery.

Page 12
Chapter 1
4
High specific energies are theoretically available with most metals in the
first, second third principal groups of the Periodic Table, but only metallic lithium(3)
has found wide application. Among the various existing batteries (Fig. 1.1.1),
lithium based batteries – because of their high energy density and design flexibility
– currently outperform other systems, accounting for 63 % of worldwide sales values
in portable batteries(4).
Figure 1.1.1 Comparison of the different battery technologies in terms of volumetric
and gravimetric energy density.(1)

Page 13
Chapter 1
5
The energy properties of metallic lithium are the most advantageous because
its redox potential of -3.01 V(3) with respect to the standard hydrogen electrode is
the highest and its atomic mass is small. Of all the alkali metals, it is also the
least reactive toward organic solvents. Lithium can be electrodeposited from
solutions of its salts in organic solvents with a 100 % yield. However this does not
mean that such an electrode can be used in a storage battery offering a large
number of charge-discharge cycles, because the metal is often deposited as a loose,
inhomogeneous dendrite layer containing badly with the operating electrode (the
remaining mass of metal). Moreover, some of the deposited metal may simply fall
away from the electrode surface to the bottom of the cell. So the main problem
when looking for a reversible, multiple-use lithium cell is to find a suitable lithium
anode at which lithium can be deposited in a compact, electron-conducting form at
least several hundred times, i.e. an anode with service life, which means almost
100 % useful recovery. Thus, in order to achieve a reversible lithium battery, a
suitable electrolyte and/ or anode must be developed.

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Chapter 1
6
1.2. Polymer electrolytes(5)
1.2.1. Basic properties of polymer electrolytes
Although, tremendous research activities by Sony leading to
commercialization of lithium batteries incorporating liquid electrolytes took place
over the last decade(6), many advantages of solid polymer electrolytes over their
liquid counterparts such as organic solutions and inorganic and molten salts can be
seen. The possibility of internal shorting, leaks, and producing combustible
reaction products at the electrode surfaces, existing in the liquid electrolytes, is
eliminated by the presence of a solid polymer electrolyte. Nevertheless, the
polymer electrolytes should exhibit ionic conductivities, at least, of the order of 10-3
to 10-2 S/cm at room temperature and play the role of a separator, played by the
liquids. The polymer should also allow good cycle lives, low temperature
performances, and good thermal and mechanical strengths in order to withstand
internal temperature and pressure buildup during the battery operation. The
polymers, in general, being light-weight and non-combustible materials can be
fabricated to requirements of size and shape, thus offering a wide range of designs.
Since stable thin films of the polymers can be easily made, high specific energy
(low-mass) and high specific power (less volume) batteries can be expected for use in
electric devices and EV. Internal voltage drops may be low, about 50 mV at 10
mA/cm2, when films as thin as 40 µm are made(7, 8).
The field of polymer electrolytes has gone through three stages: “dry solid
systems”, “polymer gels”, and “polymer composites”. The “dry systems” use the
polymer host as the solid solvent and do not include any organic liquids. The
“polymer gels” contain organic liquids as plasticizers which with a lithium salt
remain encapsulated in a polymer matrix, whereas the “polymer composites”
include high surface area inorganic solids in proportion with a “dry solid polymer”
or “polymer gel” system. In this thesis, more detailed aspects of the polymer gels
system with individual examples are discussed below. The advancement of the
field can be found in literature publications and reviews(7, 9-11).

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Chapter 1
7
Ionic conductivity
One of the important properties of a polymer electrolyte leading to its
development activity is the ionic conductivity. Temperature dependence on
conductivity of amorphous polymer electrolytes generally follows the
Vogel-Tammann-Fulcher (VTF) equation
(
)0
2
1
exp TTB
AT
-
-
=
-
σ
where, T0 is the glass transition temperature of the polymer electrolyte measured
by DSC, T is the temperature of measurement, A is the pre-exponential factor, and
E is the activation energy which can be evaluated either from the configurational
entropy theory or the free-volume theory, and hence relates to the segmental motion
of polymer chains(12). The ionic conductivity is usually measured by AC impedance
techniques(13).
Another important property of the polymer electrolyte is the lithium ion
transference number (tLi+) which ideally for lithium battery applications should be
unity. A value of tLi+ lower than 1 would tend to develop concentration gradients at
electrode surfaces leading to limiting currents. Thus, both the parameters, ionic
conductivity and lithium ion transference number are important in order to choose
a polymer electrolyte for a practical lithium battery. The maximum power
obtainable in a lithium cell can be related to the conductivity of the electrolyte,
whereas, the maxi-mum limiting current that can be drawn from the cell and the
cycleability of the cell can be related to the tLi+. The lithium transference number
(and its associated diffusion coefficient) measurements are usually made using
techniques such as concentration cell method(14, 15), Tubandt method(16, 17),
LiNMR(18), and electrochemical(19, 20), and electrophoretic NMR techniques(21, 22).
The techniques have their advantages and disadvantages and differ by theoretical
models used for interpretation of data. Analysis of some of the problems and
limitations associated with some of the techniques have been described by Fritz and
Kuhn(23).
An understanding to the interactions of the various species in the polymer
electrolytes has made possible to choose appropriate polymer hosts, complexing
salts and salt concentrations so that the ionic conductivities are optimized.
Evidence of salt dissolution into a polymer host is mainly based on spectroscopic
methods such as IR and Raman(24, 25). The anion-cation and ion-polymer
associations in both crystalline and amorphous phases have been studied using both

Page 16
Chapter 1
8
these which monitor changes in the vibrational modes of the polymer host and the
anions(26-29). For example, in the case of lithium triflate (LiCF3SO3)(28) and lithium
bis(trifluoromethane sulfonyl)imide [LiTFSI] complexes(27)
with oxyethylene
containing polymer chains, SO3 symmetric stretching modes, CF3 bending modes,
and Li-O stretching modes were associated to ion-ion interactions and information
on ion pairing and ion association was obtained(30).
X-ray and neutron diffraction methods have given de-tailed information on
the structural aspects of coordination around the ions in crystalline phases of
polymer electrolytes(31-33). The EXAFS has revealed the local environment of the
ion(34) in both crystalline and amorphous phases, including the existence of ion pairs
such as [MX]0, [M2X]+, [MX2]-, etc. and the determination of bond lengths of the
ion-polymer interactions(35). The EXAFS technique is more useful when theoretical
models depicting the most probable situations are at hand so that fitting of the
EXAFS spectrum of known structures similar to the chemical environment of the
polymer electrolyte is possible.
Stability of polymer electrolytes
The electrochemical stability of an electrolyte is one of the essential
parameters when rechargeable lithium batteries are concerned. The instability in
the electrolyte is known to bring out irreversible reactions and capacity fading in
the battery(36). The stability commonly expressed as “electrochemical stability
window” of the polymer electrolyte (units of volts) should be the same as that of the
electrode potential or higher so that overpotentials during re-charge are
compensated by the higher stability window. For example, for a 4-V lithium
battery a window of at least 4.5 V vs. Li / Li+ in the polymer electrolyte is required
for compatibility with a lithium metal anode and lithium-ion intercalating electrode
materials.
Thermal and mechanical stabilities of a polymer electrolyte during
charge-discharge cycles are vital for a safe and endurable battery. During charge
and discharge, heat is known to get generated in the battery which increases the
surface area of an existing passive layer at the electrode surface(37). The heat can
also melt and degrade the polymer electrolyte within the battery and cause internal
short circuits. Increase in heat due to environmental factors during storage
increases the self discharge reactions in the battery and shorten life.

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Chapter 1
9
1.2.2. Polymer gel electrolytes
In 1973, the first measurements on conductivities of poly(ethylene oxide)
(PEO) complexes with alkali metal salts were made by Wright et al.(38, 39). It was
after Armand that the potential of these new materials were realized for future
battery applications(40). PEO with high molecular weight of about 5 x 106 and 80 %
crystallinity was usually employed as the polymer host to form complexes with
lithium salts. Apart from the ability of the sequential oxyethylene group,
-CH2-CH2-O- in PEO to complex with lithium salts, polymer hosts containing
sequential polar groups such as O-, -NH- and -C-N- in the polymer chain were also
found to dissolve lithium salts(41).
The lithium salt complexed-poly(ethylene oxide) (PEO)(42)
and
-poly(propylene oxide) (PPO)(43) are the most widely investigated “dry solid polymer”
electrolyte systems in all solid state lithium batteries. The main reason to choose
these two polymer hosts is because they form more stable complexes and possess
higher ionic conductivities than any other group of solvating polymers without the
addition of organic solvents. Complex formation in PEOn, -salt (n = number of
ether oxygens per mole of salt) is governed by competition between solvation energy
and lattice energy of the polymer and the inorganic salt(44). Low lattice energies of
both the polymer and the complexing salt have been found to increase stabilities in
the resultant polymer electrolyte(44).
On dry solid polymer systems, to enhance the ionic conductivity, the
improvement of host polymers was attempted by the reduction of its crystallinity.(7,
45-50)
The performances of electrochemical cells incorporating the “dry solid”
polymer electrolytes and lithium metal electrodes were not satisfactory, and cycle
lives were as low as 200 to 300 cycles. The poor performance of the cells was
mainly attributed to the poor conductivity of the electrolytes, along with the
reactivity of the anion of the electrolyte towards the lithium metal electrodes.
Therefore, the polymer gel systems, whose crystallinity is reduced by adding
plasticizer, have attracted attention.

Page 18
Chapter 1
10
In 1975, Feuillade and Perche demonstrated the idea of plasticizing a
polymer with an aprotic solution containing an alkali metal salt in which the
organic solution of the alkali metal salt remained trapped within the matrix of the
polymer(51). Such mixings have resulted into formation of gels with ionic
conductivities close to those of the liquid electrolytes, arising from similar
conductivity mechanisms taking place in both the systems. However, one would
expect slightly lower conductivities in the polymeric gels, if more viscous solvents
are used, as compared to those used in the liquid electrolytes.
Since then, various polymeric hosts such as, poly-(vinylidene fluoride)
(PVdF)(52), poly(vinylidene carbonate) (PVdC), poly(acrylonitrile) (PAN)(53, 54),
poly(vinyl chloride) (PVC)(55), poly(vinyl sulfone) (PVS)(56), poly(p-phenylene
terepthalamide) (PPTA)(57), and poly(vinyl pyrrolidone) (PVP), have been found to
form electrolytes with conductivities ranging between 10-4 and 10-3 S/cm at 20 oC
(Table 1.2.1)(7). These systems are presently expressed in various terms such as
“plasticized polymer electrolytes”, “polymer hybrids”, “gelionics” and “gel
electrolytes”. The electrolytes are easily prepared by heating a mixture containing
the appropriate amounts of the polymer, solvents and a lithium salt to about 120 -
150 oC, a temperature above the glass transition temperature of the polymer, to
form viscous clear liquids. Films of the gels are usually made by solution casting
in hot state allowing the solution to cool under pressure of electrodes.
Less-evaporating solvents, such as ethylene carbonate (EC), propylene carbonate
(PC), dimethyl formamide (DMF), diethyl phthalate (DEP), di-ethyl carbonate
(DEC), methylethyl carbonate (MEC), dimethyl carbonate (DMC), γ-butyrolactone
(BL), glycol sulfite (GS), and alkyl phthalates have been commonly investigated as
“plasticizer” solvents for the gel electrolytes(7). The solvents have been used
separately or as mixtures.
The phenomenon of plasticization was found to increase the amorphous
phase in the polymer system with a single glass transition temperature as low as
-40 oC, the latter varying with the amounts of solvent and polymer containing in the
gels. It was generally observed that up to 80 % of solvent could be trapped into the
polymer matrix. The high permittivity solvents allowed a greater dissociation of
the lithium salt and increased the mobility of the cation. Much research work in
the area of gel electrolytes can be found in literature reviews(7, 44, 58).

Page 19
Chapter 1
11
Table 1.2.1 Ionic conductivities of some gel polymer electrolytes and polymer
composite electrolytes.

Page 20
Chapter 1
12
PEO-based gels
Plasticization of high molecular weight P(EO)n-LiX electrolytes with PC
and/or EC was found to form soft solids with poor mechanical stabilities, although
room temperature conductivities as high as the order of 10-3 S/cm were obtained(59,
60). The poor mechanical stability was accounted to be mainly due to the solubility
of the PEO in the solvents(60).
Cross-linking of polymers by methods such as, UV(61), thermal radiation(62),
photo-polymerization(63), and electron beam radiation polymerization(64) was found
to reduce the solubility of the polymers with the organic solvents and also helped to
trap the liquid electrolyte within the polymer matrix. Low molecular weight PEO
was cross-linked and plasticized with 50 % PC by Borghini et al.(65). This material
showed good mechanical proper-ties and conductivities of the order 10-4 S/cm at 20
oC were two orders of magnitude higher than the unplasticized amorphous
cross-linked PEO-LiClO4
complex. In general, the room temperature
conductivities of gels based on polymers and copolymers prepared by crosslinking
methods were found to be in the range of 10-5 - 10-4 S/cm(65).
Plasticizers such as dioctyl-, dibutyl-, and dimethyl-phthalate were recently
investigated for PEO-LiClO4
complexes(66). In comparison with the three
plasticizer-based systems, lowest crystallinity in the gel containing the dioctyl
phthalate (DOP) was observed showing the best room temperature conductivity of
9.67 x 10-5 S/cm for the composition (PEO)8-(LiClO4 : DOP, 99.9 : 0.1 wt.%).
However, the room temperature conductivities of the complexes containing the
dibutyl phthalate and the dimethyl phthalate were slightly lower than those
containing the DOP plasticizer.

Page 21
Chapter 1
13
PAN- and PVDF-based gels
The PAN-based(67-69) and PVdF-based(70) gel electrolytes are the most widely
investigated polymer gel electrolyte systems and remarkable conductivities of the
order of 10-3 S/cm at 20 oC have been obtained. A fully amorphous gel of
PAN-LiClO4 (1 : 0.2) in EC showed room temperature conductivity of 1 x 10-3 S/cm
and an estimated activation energy of 86 kJ/mol(71). Because of the absence of any
oxygen atoms in the PAN polymer matrix, the PAN-based gels were found to have
lithium ion transference numbers greater than 0.5(72). With a greater dissociation
of salts such as, LiTFSI and LiTFSM in
Figure 1.2.1 Probable interaction of the lithium ion in gel electrolytes.
PAN-based gels, transference numbers as high as 0.7 could be obtained. The
delocalization of electron density around the large organic anions in the salts is
believed to promote greater dissociation and increase the transference number.
Strong polar groups in the polymer are undesirable because of the
complexation of lithium ion by both the polymer matrix and the solvent(73). A
probable interaction of the lithium ion with the polymer chain and the solvent is
shown in Fig. 1.2.1 for a PAN-based gel system(74). Thus, the polymer in the gel
electrolytes would mainly act as an encapsulating matrix with electrostatic
interactions with the solvated lithium salt so that the lithium ion mobility is least
hindered. Yang et al.(71) studied a PAN-LiClO4-DMF gel in order to understand the
interaction of the Li+ ions with the PAN chain. Using FTIR they showed the Li+
ions to form bonds with C=N groups of the PAN as well as C=O of the DMF which
would support the structure in Fig. 1.2.1. Using the same technique, in
PAN-LiClO4-DMF, Wang et al. reported interactions between Li+ ions and the
oxygen and/or nitrogen atoms of DMF along with interactions between the O atoms
of DMF and N atoms of the nitrile of PAN(75). Unfortunately, no interactions
between LiClO4 and PAN could be detected with the use of the FTIR technique.

Page 22
Chapter 1
14
Jiang et al. recently studied PVdF-based gels consisting of EC, PC and a LiX
salt(76). As expected they found the presence of plasticizers EC and PC to
significantly disorder the crystalline structure and reduce crystallinity in the gels
than in the parent polymer host, PVdF. They also reported the mechanical
strength of the resulting gel to be depended on the PVdF content, whereas the
conductivity to be mainly influenced by the viscosity of the medium and the
concentration of the lithium salt. Room temperature conductivities as high as 2.2
x 10-3 S/cm for the LiN(CF3SO2)2 salt containing mixture were reported(76).
Increase in gel properties were observed when mixtures of polymer hosts or
copolymers were used in the gel preparation. Increase in gel properties were
observed for PVdF co-polymerized with hexafluoropropylene (HFP). The
PVdF-HFP co-polymer in the gel showed greater solubility for organic solvents, and
lower crystallinity and glass transition temperature than the PVdF polymer alone
in the gel(77).
Others
Wieczorek and Stevens studied blends polyether, PMMA, and LiCF3SO3
(78).
The electrolytes showed a maximum room temperature conductivity of 3 x 10-5 S/cm.
Morita et al. also prepared gels consisting of PMMA grafted with PEO and
containing Li salts, and room temperature conductivities of the order 10-3 were
reported(79). Addition of crown ethers such as 12-crown-4 and 15-crown-5 to the
PMMA-PEO-Li+ gel was found to increase the conductivity to some extent, whereas
the Li ion mobility was found to increase significantly with the addition of the
15-crown-5 ether(79). Lithium transference numbers greater than 0.5 were found in
the gel electrolytes based on poly(methyl methacrylate)(80)
and
poly(tetrahydrofuran)(81). The values were found to be dependent on the amounts
of organic solvent present in the gel.
Blends of polymers, PVC and PMMA were studied using PC as the
plasticizer and LiCF3SO3 as the salt(82). Due to insolubility of PVC in the solvent
PC, phase separation was observed. The inclusion of PVC into PMMA helped to
increase the mechanical stability of the gel, but unfortunately decreased the lithium
ion conductivity. The ions were found to preferentially move towards the
plasticizer-rich phase or the PMMA-rich phase.
Dimensionally stable gels consisting of PEG-PAN-PC-EC-LiClO4
were
prepared by Munichandraiah et al.(83). Compared with gels PEO-PC-LiClO4 and

Page 23
Chapter 1
15
PAN-PC-EC-LiClO4, the PEG containing gels showed lower room temperature
conductivities, but higher mechanical stabilities.
PVC-based electrolytes consisting of LiClO4, LiTFSM, or LPF6 salt and
THF/PC mixture as plasticizer gave ionic conductivity of the order 10-4 S/cm at 20 oC
with respective lithium ion transference numbers of 0.26, 0.40, and 0.45(84). Polymer
gels consisting of PVC, LiTFSI, and solvents such as dibutyl phthalate (DBP) and
dioctyl adipate (DOP) gave room temperature conductivities of the order 10-4
S/cm(85).
Low molecular ethers such as poly(ethylene glycol) dimethyl ether
(PEGDME) was used as a solvent for salts such as LiN(CF3SO2)2, LiCF3SO3, and
LiPF6 to which poly(vinylidene fluoride)-hexafluoropropene copolymer was added as
the encapsulating polymer matrix(86). PEGDME was also added to a copolymer
formed from triethylene glycol dimethacrylate (TRGDMA) and acrylonitrile (AN)
and LiCF3SO3 salt. Room temperature conductivities in the range from 10-5 to 10-4
S/cm were reported(87). The conductivity was found to increase in line with the
molar ratios of both AN : TRGDMA and (EO) : LiCF3SO3
(87).
In the gel mixture consisting of poly(p-phenylene terephthalamide) (PPTA),
polyethylene glycol (PEG), polycarbonate in PC-EC, and a lithium salt,
conductivities as high as 2.2 x 10-3 S/cm at room temperature were observed at 0.8
M LiBF4 salt per mole of PPTA content(57). Above 1 M of the salt content, the
conductivity was found to fall rapidly suggesting the ion conductivity to be due to
LiBF4 interaction at the amide bond sites of the PPTA(57).
In the polymer gel electrolytes, the incorporation of liquid electrolytes into
homo- or co-polymer hosts has allowed room temperature conductivities as high as
10-3 S/cm. The mechanical stability of the gels is determined by the ratios of the
polymer and solvent (plasticizer) in the gel. Highly mechanically stable gels with
10 - 12 wt.% of plasticizer and 70-80 wt.% of polymer, with room temperature
conductivities between 10-4
and 10-3
S/cm have been generally observed.
Plasticizer solvents such as EC and PC have been much studied due to their high
polarity and low vapor pressure, which also allow greater plasticizing effect to the
polymer host. Most of the studied plasticizers have shown electrochemical
instabilities at lithium metal electrode surfaces, and thus a search for new
plasticizers is seen necessary. Due to greater dissociation of the lithium salt, much
higher lithium ion transference numbers than in the solvent-free electrolytes (about
0.6) have been observed.

Page 24
Chapter 1
16
1.3.
Anode materials(88, 89)
1.3.1. Carbon anodes
Since the lithium ion secondary battery with carbon anode had
commercialized by Sony Corporation in 1991(6), the carbon anode has been
investigated. At present mostly carbons are used as the negative electrode of
commercial rechargeable lithium batteries: i) because they exhibit both higher
specific charges and more negative redox potentials than most metal oxides,
chalcogenides, and polymers; and ii) due to their dimensional stability, they show
better cycling performance than Li alloys. The insertion of lithium into carbon,
habitually named “intercalation”, proceeds according to Equation (1).
(1)
Due to electrochemical reduction (charge) of the carbon host, lithium ions
from the electrolyte penetrate into the carbon and form a lithium/carbon
intercalation compound, LixCn. The reaction is reversible.
The quality of sites capable of lithium accommodation strongly depends on
the crystallinity, the microstructure, and the micromorphology of the carbonaceous
material(90-107). Thus, the kind of carbon determines the current/potential
characteristics of the electrochemical intercalation reaction, and also the risk of
solvent co-intercalation.
Carbonaceous materials suitable for lithium intercalation are commercially
available in hundreds of types and qualities(90, 91, 108-111). Many exotic carbons have
been synthesized in laboratories by pyrolysis of various precursors, some of them
with a remarkably high specific charge. Fullerenes have also been tested(112, 113).
Carbons that are capable of reversible lithium intercalation can roughly be
classified as graphitic and non-graphitic (disordered). Graphitic carbons are
carbonaceous materials with a layered structure but typically with a number of
structural defects. From a crystallographic point of view the term “graphite” is
only applicable to carbons having a layered lattice structure with a perfect stacking
order of graphene layers, either the prevalent AB (hexagonal graphite, Fig. 1.3.1) or
the less common ABC (rhombohedral graphite). Due to the small transformation
energy of AB into ABC stacking (and vice versa), perfectly stacked graphite crystals
are not readily available. Therefore, the term “graphite” is often used regardless of
stacking order. The actual structure of carbonaceous materials typically deviates
LixCn
xLi+ + xe- + Cn
discharge
charge

Page 25
Chapter 1
17
more or less from the ideal graphite structure. Materials consisting of aggregates of
graphite crystallites are named graphite as well. For instance, the terms natural
graphite, artificial or synthetic graphite, and pyrolytic graphite are commonly used,
although the materials are polycrystalline. The crystallites may vary considerably
in size. In some carbons, the aggregates are large and relatively free of defects, for
example, in highly oriented pyrolytic graphite (HOPG). In addition to graphitic
crystallites, other carbons also include crystallites containing carbon layers (or
packages of stacked carbon layers) having significant misfits and misorientation
angles of the stacked segments to each other (turbostratic orientation or
turbostratic disorder)(114). The latter disorder can be identified from an increased
average planar spacing compared to graphite(115).
Figure 1.3.1 Left: schematic drawing of the crystal structure of hexagonal graphite,
showing the AB layer stacking sequence and the unit cell. Right: view perpendicular
to the basal plane of hexagonal graphite. Prismatic surfaces can be subdivided into
arm-chair and zig-zag faces.
Non-graphitic carbonaceous materials consist of carbon atoms that are
mainly arranged in a planar hexagonal network but without far-reaching
crystallographic order in the c-direction. The structure of those carbons is
characterized by amorphous areas embedding and crosslinking more graphitic
ones(116) (Fig. 1.3.2). The number and the size of the areas vary, and depend on
both the precursor material and the manufacturing temperature. Non-graphitic
carbons are mostly prepared by pyrolysis of organic polymer or hydrocarbon
precursors at temperatures below ~1500 oC. Heat treatment of most non-graphitic
(disordered) carbons at temperatures from ~1500 to ~3000 oC allows one to

Page 26
Chapter 1
18
distinguish between two different carbon types. Graphitizing carbons develop the
graphite structure continuously during the heating process, as crosslinking between
the carbon layers is weak and, therefore, the layers are mobile enough to form
graphite-like crystallites. Non-graphitizing carbons show no true development of
the graphite structure even at high temperatures (2500-3000 oC), since the carbon
layers are immobilized by crosslinking. Since non-graphitizing carbons are
mechanically harder than graphitizing ones, it is common to divide the
non-graphitic carbons into “soft” and “hard” carbons(116).
Figure 1.3.2 Schematic drawing of a non-graphitic (disordered) carbonaceous
material.

Page 27
Chapter 1
19
1.3.2. Alloy anodes
Various insertion materials have been proposed for negative electrodes of
rechargeable lithium batteries, for example, transition-metal oxides and
chalcogenides, carbons, lithium alloys, and polymers. Table 1.3.1 shows that both
the specific charges and the charge densities of lithium insertion materials are
theoretically lower than that of metallic lithium. However, considering that the
cycling efficiency of metallic lithium is ≤ 99 %, one has to employ a large excess of
lithium(91, 117, 118) to reach sufficient cycle life. The practical charge density of a
secondary lithium electrode is therefore much lower than the theoretical one, so
that it is comparable with the charge densities of alternative lithium-containing
compounds. However, the potential of the electrode materials also has to be
considered because a higher potential versus Li/Li+ of the negative electrode means
a lower cell voltage. For instance, the potential of many Li alloys is ~0.3 to ~1.0 V vs.
Li/Li+ whereas it is only ~0.1 V vs. Li/Li+ for graphite (Fig. 1.3.3).
Table 1.3.1 Characteristics of representative negative electrode materials for
lithium batteries, calculated by using data from(117, 119-121). The values are for fully
lithiated host materials except for the values in parentheses, which are for
lithium-free host materials. Li4 denotes a four-fold lithium excess, which is
necessary to reach a sufficient cycle life.
[a] In some cases a considerably lower amount of the specific charge/charge density can be cycled
reversibly in practice.

Page 28
Chapter 1
20
The replacement of metallic lithium by lithium alloys has been under investigation
since Dey(122) demonstrated the feasibility of electrochemical formation of lithium
alloys in liquid organic electrolytes in 1971. The reaction usually proceeds
reversibly according to the general scheme shown in Equation (2).
(2)
With only a few exceptions (such as hard metals, M = Ti, Ni, Mo, Nb), Li alloys are
formed at ambient temperature by polarizing the metal M, for example, Al, Si, Sn,
Pb, In, Bi, Sb, Ag, and some multinary alloys,(93, 120, 123-134) sufficiently negatively in a
Li+- containing electrolyte. In most cases even the binary systems Li-M are very
complex. Sequences of stoichiometric intermetallic compounds and phases LixM
with considerable phase range are usually formed during lithiation of the metal M,
characterized by several steps and/or slopes in the charge diagram (Fig. 1.3.3). The
formation of Li-M phases is in many cases reversible, so that subsequent steps and
slopes can also be observed during discharge.
Figure 1.3.3 Charging curves of some matrix metals (M), compared to highly
oriented turbostratic mesophase pitch carbon fibers (P 100, FMI Composites-Union
Carbide, characterized in LiClO4 / propylene carbonate). Modified and redrawn
from(135).
LixM
xLi+ + xe- + M
discharge
charge

Page 29
Chapter 1
21
A Li+ ion transfer cell with the trademark Station, announced recently by
Fujifilm Celltech Co., Ltd.,(136) uses an “amorphous tin-based composite oxide
(abbreviated TCO or ATOC)” for the negative electrode. The TCO combines both: i) a
promising cycle life and ii) a high specific charge (>600 Ah/kg) and charge density
(>2200 Ah/L). (137)
The TCO is synthesized from SnO, B2O3, Sn2P2O7, Al2O3, and other
precursors. However, only the SnII compounds in the composite oxide are said to
form the electrochemically active centers for Li insertion. The oxides of B, P, or Al,
which are electrochemically inactive, have glass forming properties and form a
network stabilizing the dimensional integrity of the composite host during
synthesis.(137) The improvement of cycle life by using composite amorphous lithium
insertion materials, such as V2O5, together with a network former, such as P2O5, has
been reported in the literature.(138-144)
In order to explain the high specific charge a mechanism can be suggested
in which the tin oxide reacts to form Li2O and metallic Sn.(145, 146) This reaction is
associated with large charge losses due to the irreversible formation of Li2O. In a
second step the Sn then alloys with lithium reversibly. On the other hand, according
to Fujifilm Celltech(137) no Li2O was found after lithium insertion. However, the idea
that the high specific charge of the TCO is due to the alloying of metallic tin has led
to a renaissance of Li-alloy negative electrode research and development.(145)
The electrochemical reduction of a tin electrode in a Li+-containing
electrolyte leads to the subsequent formation of a number of intermetallic phases
LixSny at high temperature (415 oC)(147) as well as at room temperature(145, 148-150).
Diffusion data (Table 1.3.2) disclose that the Li+ cation mobility in lithium-rich
phases is quite acceptable both at 415 oC(148) and at room temperature(123, 132)
allowing reasonable charge/discharge current densities. However, whereas very
slow measurements at near equilibrium conditions allow to characterize several
LixSny phases (particular at 415 oC, cf. data in Table 1.3.2), at room temperature
and under practical charging conditions only the lithium-poor phases Li2Sn5 and
LiSn can be clearly distinguished in the constant current charge curve (Figure
1.3.4) as well as in X-ray diffraction patterns.(145, 151) It has been suggested that the
lithium-rich phases LixSny do not form long-range ordered structures, because the
atom mobility is too low at room temperature in this case.(145, 151)

Page 30
Chapter 1
22
Table 1.3.2 Melting points, densities from X-ray data dx), specific charges, charge
densities as well as plateau potentials and maximum chemical diffusion coefficients
((Dchem-max) of several LixSny phases at 415 oC (LiCl-KCl eutectic melt as electrolyte)
and 25 oC (LiAsF6/PC or LiClO4/PC as electrolytes).
LixSny
x in
LixSn
Melting
point
(oC)
dx
(g/cm3)
Specific
charge
(Ah/kg)
Charge
density
(Ah/L)
Plateau potential
for formation
at 415 oC
(V vs. Li/Li+)
Dchem-max
at 415 oC
(cm2/s)
Plateau potential
for formation
at 25 oC
(V vs. Li/Li+)
Dchem-max at
25 oC
(cm2/s)
Li2Sn5
0.4
319
6.11
88.3
539.5
N/A
N/A
~0.760
N/A
LiSn
1
≥485
5.10
213.3
1087.8
0.569
4.10x10-6
0.660
8.0x10-8
‘LI3Sn2
1.5
465
*
*
*
*
*
*
*
Li7Sn3
2.33
500
3.67
463.6
1701.4
0.456
4.07x10-5
0.530
N/A
Li5Sn2
2.5
≥698
3.54
492.5
1743.5
0.430
5.89x10-5
0.485
5.0x10-7
Li13Sn5
2.6
716
3.46
509.6
1763.2
0.390
7.59x10-4
0.485
N/A
Li7Sn2
3.5
783
2.96
656.0
1941.8
0.284
7.76x10-5
0.420
N/A
‘Li4Sn’
4
≥684
*
*
*
*
*
*
*
Li22Sn5
4.4
758
2.56
790.2
2022.9
0.170
1.91x10-5
0.380
5.9x10-7
*Existence doubtful
Figure 1.3.4 Characteristic charge curve of electroplated Sn in 1 M LiClO4/PC,
i =0.025 mA/cm2.
The lithium-rich phases exhibit high theoretical specific charges and charge
densities (Table 1.3.2). On the other hand, the tin pulverizes rapidly when it
experiences full lithiation due to the large density decrease (is volume increase).
Limiting the degree of lithium uptake to the stoichiometry of LiSn relieves the

Page 31
Chapter 1
23
pulverization problem to some extent, because the differences in density between
metallic Sn (7.29 g/cm3) and the phases Li2Sn5 and LiSn are relatively small (Table
1.3.2). Moreover, these phases are structurally related(152), so that the consequences
of reconstitution reactions are small as well. However, further alloying with lithium
beyond the stoichiometry LiSn leads to a rapid density decrease (Table 1.3.2) and
causes major structural rearrangements, which considerably enlarge the
mechanical stresses on the host material. In order to improve the dimensional
stability and thus the rechargeability for higher lithium uptakes, the morphology
and chemistry of the host material has to be specifically designed.
Host metal properties, such as particle size, shape, texture, porosity, etc.
strongly affect the macroscopic dimensional stability during lithium alloying and
dealloying and thus the cycling behavior.(135, 153, 154) Though the volume expansions of
the metal hosts upon alloying with lithium are in the order of several 100 %, large
absolute volume changes can be avoided, when the size of the metallic host particles
is kept small (Fig. 1.3.5).
The practical feasibility of this concept was checked by employing tin
anodes of different particle sizes. These have been prepared by electroplating on
copper substrates from aqueous solutions containing Sn2+ cations. The morphology
of the metals can be drastically changed by a control of the plating conditions, for
example, by variation (i) of the chemical composition of the solutions (e.g. by
addition of leveling and complexing agents), (ii) of the solution temperature, (iii) of
the stirring conditions and (iv) of the plating current densities. The detailed
preparation procedures can be found elsewhere.(135, 154)
The thickness of the
deposits is in the range of a few micrometers, i.e. too thin for practical application in
current cylindrical and prismatic lithium ion cells. Nevertheless, since they are
binder-free they are good model substrates for understanding the relation between
metal anode properties and electrochemical performance.

Page 32
Chapter 1
24
Figure 1.3.5 Model: 1st lithiation (is 1st alloying with lithium) of a loosely packed
small particle size metallic material. Even 100% volume expansion of the
individual particles does not crack them as their absolute changes in dimensions
are still small.
The cycling performance of lithium alloys can be significantly improved if
intermetallic and/or composite hosts are employed instead of pure metals. The
basic idea is that at a certain stage of charge/discharge, i.e. at a certain electrode
potential, one (or more) components/phases of the composite or intermetallic are
able to store lithium (`reactant'), i.e. expand/contract, whereas the other
components/phases are less active or even inactive(88), i.e. perform as a `matrix'
buffering the expansions of the reactant (Fig. 1.3.6). Additional preconditions for a
successful application of this concept are (i) that the reactant is finely dispersed in
the matrix, (ii) that the matrix allows the electrons and lithium ions to move
between the reactive domains or is a mixed (electron and lithium cation) conductor
and (iii) that the reactive domains are sufficiently small, so that the above discussed
benefits of `small particle size' are effective.
Figure 1.3.6 Model: strong expansions of the `reactant' domains due to lithiation
can be buffered by the inactive or less active `matrix' domains, thus keeping the
extent of crack formation in the overall multiphase material small.
Thus, the selection of an adequate matrix is the key to a successful Sn
compound anode. Elements that are inactive against Li are assumed to suppress
the volume change effectively without much irreversible capacity. The use of Sn

Page 33
Chapter 1
25
compounds with elements such as Ni(155-157), Fe(158-162), Cu(163-166), Mn(160, 167), and
Co(160) has been investigated based on this assumption. Recently, an amorphous
ternary Sn-Co-C anode has been introduced to a practical use.(168)

Page 34
Chapter 1
26
1.4. Phase separated structure
1.4.1. Polymer blend
A kind and quantity of the polymer materials were rapidly developed in
these dozens of years. We can not live without the polymer materials over a wide
area. However, the polymer materials are said to be reached the age of puberty, the
characteristic of polymer are demanded to meet a broad applications. However,
the performance as the single material has already come to the limit. The composite,
copolymerization, and polymer blend of polymers are needed as the metallic and
inorganic materials had done. The multicomponent polymer will satisfy enough
demanded performance by a suitable combination from many existing polymers.
Polymer blend composed of two immiscible polymers(169, 170)
A co-continuous morphology is a non-equilibrium morphology that is
generated during mixing of two polymers. As such, it is an unstable morphology,
and it starts changing through filament break-up and retraction as soon as the fluid
blend comes out of the mixer. However, the blend may remain co-continuous, if it
is frozen fast. Considerable attention has been given to the conditions that make
co-continuous morphologies possible in blends during the mixing. It has been
generally suggested that co-continuity occurs at the phase inversion point.
Existing empirical relations(171-173) and theories(174) give a volume fraction for phase
inversion as a function of the viscosity ratio, as shown in Fig. 1.4.1. However,
many experimental results(170) cannot be described with these relations (Fig. 1.4.1).
By basing the phase inversion point on the viscosity ratio of the components only,
these relations neglect the influence of the blending conditions and material
properties, e.g. the interfacial tension. Moreover, they do not take into account any
requirements as to the shape of the dispersed component necessary to obtain
co-continuity. Especially at low volume fractions, a co-continuous network can
only exist if the minor blend component consists of structures with an extended
shape(170). These structures can be formed and remain so only under appropriate
blending conditions. For this reason, it is to be expected that the existence of a
co-continuous morphology in the blend will be strongly dependent both on the
processing conditions, e.g. the stress levels, and the processing properties of the
blend components, e.g. the matrix viscosity, viscosity ratio and interfacial tension.
In the literature(170), an equation has been derived that describes the critical volume

Page 35
Chapter 1
27
fraction of the minor phase for complete co-continuity as a function of the matrix
viscosity, interfacial tension, shear rate and phase dimensions, and it was shown
that a high viscosity of the matrix component was favorable for co-continuity over a
broad composition range in blends of commercial grades of polyethylene and
polystyrene.
Figure 1.4.1.Phase inversion as a function of the viscosity ratio, p = ηd/ηm, according
to several empirical relations summarized. * and ●: points found for PE/PS systems
in Ref. (170).

Page 36
Chapter 1
28
Polymer blend composed of a di-block copolymer (175)
Block copolymers have recently received much attention not only thanks to the
scale of the microdomains (tens of nanometers), their various chemical and physical
properties but also due to the convenient size and shape tunability of microdomains
afforded by simply changing their molecular weights and compositions. Many
potential uses of block copolymers for different nanotechnologies have been proposed
based on principally their ability to form interesting patterns. However, the main
challenge of using block copolymers lies with control of microstructure. Achievement
of precise microdomain location, orientation, and elimination of various defects requires
introduction of external fields during the processing step. A variety of mechanical,
electrical, magnetic biases and surface interactions have been proposed to manipulate
and guide the microstructures of block copolymers.
Block copolymers consist of chemically distinct polymer chains covalently
linked to form a single molecule. Owing to their mutual repulsion, dissimilar blocks
tend to segregate into different domains, the spatial extent of the domains being limited
by the constraint imposed by the chemical connectivity of the blocks. Area
minimization at the interface (the IMDS) of two blocks takes place to lower the
interfacial energy. From an entropic standpoint, the molecules prefer random coil
shapes but the blocks are stretched away from the IMDS to avoid unfavorable contacts.
As a result, of these competing effects, self-organized periodic microstructures emerge
on the nanoscopic length scale. Various microdomain structures are achieved,
depending on relative volume ratio between blocks and chain architecture as well as the
persistence lengths of the respective blocks.
In the simplest case of non-crystalline flexible coil AB di-block copolymers, the
composition of the AB di-block (i.e. the volume fraction f of block A) controls the
geometry of the microdomain structure. As shown in Fig. 1.4.2, for nearly symmetric
di-blocks (f ~ 1/2) a lamellar phase occurs. For moderate compositional asymmetries, a
complex bi-continuous state, known as the double gyroid phase, has been observed in
which the minority blocks form domains consisting of two interweaving
threefoldcoordinated networks. At yet higher compositional asymmetry, the minority
component forms hexagonally packed cylinders and then spheres arranged on a
body-centered cubic lattice. Eventually, as f → 0 or 1, a homogeneous phase results.(176)

Page 37
Chapter 1
29
Figure 1.4.2. Schematic phase diagram showing the various ‘classical’ block
copolymer morphologies adopted by non-crystalline linear diblock copolymer. The
blue component represents the minority phase and the matrix, majority phase
surrounds it.(177)

Page 38
Chapter 1
30
1.4.2. Liquid crystals(178)
Polyoxyethylene surfactants are widely used as emulsifying agents and
detergents. Like ionic surfactants they form micelles above a critical concentration
in water (the c.m.c.) with liquid crystals frequently occurring at higher
concentrations. They are unusual in having a lower consolute temperature, termed
the cloud point. For many years this was attributed to the presence of giant
micelles. However, it is now thought to correspond to the phase separation of a
concentrated surfactant solution containing small micelles from a more dilute
solution, perhaps due to the presence of significant intermicellar
attractions.(179-181) Some surfactants even show a ‘double cloud point’, where a
1 % aqueous solution goes cloudy then clears and clouds a second time when the
temperature is increased.(182)
Over years a number of theories to describe micelle shapes have been
published.(183, 184)
These involve a balance of alkyl-chain/water repulsions and
repulsion between adjacent head groups within the micelle, together with surface
curvature and limitations due to alkyl-chain packing. Most lyotropic liquid crystals
consist of ordered micelles.(185)
It should be possible to extend the theories of
micelle shapes to account for mesophase formation by addition of appropriate terms
for intermicellar forces. Parsegian(186) has already successfully described the
hexagonal/lamellar transition of ionic surfactants using Poisson-Boltzmann theory
to account for electrostatic repulsions. Wennerstram and co-workers(187, 188)
have developed this approach further. For non-ionic surfactants there is no
quantitative theory to describe intermicellar repulsions. However, with
poly(oxyethylene) alkyl ethers [n-CnH2n+1(OCH2CH2)mOH,CnEOm] it is possible to
change systematically the surfactant chemical structure by alteration of m and n,
hence varying head-group interactions and micelle size. By determining
water/surfactant phase diagrams of a range of compounds, the validity of this
theoretical description and the importance of each contributing factor (alkyl-chain
conformations, curvature, head-group area etc.) for the structure and stability of
mesophases can be assessed.
The shape of micelles just above the c.m.c. is highly dependent upon
surfactant type and solution conditions (concentration, electrolyte level,
temperature). It appears that spherical micelles, rod micelles and oblate spheroid
(bilayer) micelles all occur, but under different circumstances. The structures of

Page 39
Chapter 1
31
normal hexagonal (H1), lamellar (Lα) and reversed hexagonal (H2) lyotropic
mesophases are well known.(185, 189)
The lamellar phase consists of equidistant
parallel surfactant bilayers separated by water layers. In the hexagonal phases
very long normal or reversed rod micelles are packed in a hexagonal array. Two
different classes of cubic phases have been discovered,(185, 190) one occurs at
compositions between micellar solution (L1) and H1 phases, consisting of ordered
spherical micelles packed in body- or face-centred arrays (labelled I1), and the
second occurs between Lα and H1 (labelled V1) or H2 (labelled V2). While the
structures of these phases are not fully known, the best candidates(185, 190) are
regular ‘bicontinuous’ networks where the chain/water interface has both positive
and negative curvatures.
The forces responsible for micellar shape and mesophase structure can be
divided into intramicellar and intermicellar contributions. The former determine
the shape just above the c.m.c. while the latter take account of intermicelle
interactions at higher concentrations. The description given below is a summary
and development of several treatments.(183-185, 190, 191)
At the c.m.c. micelle shape is determined by the surface area per molecule
(a) at the alkyl-chain/water interface and the interface curvature. The possible
micelle shapes are spheres, rods (prolate spheroids) or discs (oblate spheroids,
bilayer micelles). For a spherical micelle, the volume of the surfactant alkyl chain
(v), micelle radius (r) and a are related by
3
ar
v =
........................................................................................................(1)
Since r cannot extend beyond the ‘all-trans’ length of alkyl chain (lt), while v is
invariant, then a cannot fall below a certain limit (as
c). From the known densities of
alkanes, assuming a micelle radius of 1.25 Ǻ per C-C bond this is estimated to be as
c
= 70 Ǻ 2. Similar arguments apply to rod micelles, with the limit (ar
c ) being ar
c =
47 Ǻ 2. Thus if a ≥ 70 Ǻ, the possible micelle shapes are spheres, rods or discs. With
a in the range 70 > a/ Ǻ 2 ≥ 47, only rod or disc micelles can form. With a < 47 Ǻ 2,
rod micelles are excluded and only bilayer micelles (or reversed phases) can form.
This limitation of micelle type is generally referred to as the ‘alkyl-chain packing
constraint’.
The detailed shapes adopted by rod and disc micelles are not known.
Prolate and oblate ellipsoids, or rods with hemi-spherical ends and bilayers with
hemi-cylindrical edges, are both popular pictorial representations. It has been

Page 40
Chapter 1
32
argued that only ellipsoids with low eccentricity can occur because alkyl chains are
unable to fill up completely the highly curved regions at the edge of ellipsoids with
high eccentricity.(183)
This conclusion involves an assumption that the alkyl-chain
axis is, on average, normal to the interface. There seems to be no physical reason
why a few chains at the ellipsoid–micelle edge should not be tilted to allow packing
into the highly curved region, and thus the restraint is too severe.
Figure 1.4.3 Schematic illustration of forces at the micelle surface (see text). Plane x
represents alkyl-chain/water interface, inter-head-group repulsions act in plane y.
Given that the value of a is sufficiently large for any particular shape to
occur, the actual form adopted is determined by surface curvature. The forces
present at the chain/water interface (intramicellar forces) can be represented as
occurring in different planes, as shown schematically in Fig. 1.4.3. The
chain/water interface is at x. Interactions between hydrated head groups
(electrostatic, solvation, steric) usually result in a repulsive force in some plane y,
further into the water than x. The minimization of hydrocarbon-chain/water
interactions gives an attractive force in plane x. These are dominant contributions.
However, depending on the value of a, there will also be an apparent
repulsive force in some plane z within the alkyl-chain region. This arises from the
unfavorable entropy change caused by restrictions on the conformations of the
chains with small a values. While no calculations of the magnitude of this effect
have been published, it is obvious that it will increase with temperature and with
alkyl-chain length, since the number of available conformations is proportional to
3n-2 (n = chain carbon number). The effect will also depend on the shape of the
micelle. Without detailed model calculations of chain conformations it is not
possible to estimate this contribution to the stability of spheres, rods or bilayers at
large a value (a > as
c). However, any restriction on the chain taking the all-trans
conformation will increase the limiting values as
c, and ar
c. Moreover, with a > as
c,
the average chain length within a binary is < lt/3. This will involve numerous
y
x
z

Page 41
Chapter 1
33
gauche conformations which have an unfavorable enthalpy, so disfavoring the
lamellar phase. For low values of a, spheres and rods cannot pack. On comparing
reversed phases and the lamellar phase as a approaches the limit of the all-trans
cross-sectional area within the lamellar phase (ab
c), the reversed phases will be
favored at sufficiently high temperatures because they have more conformational
states available to the chains. Thus in practice the alkyl-chain curvature effect is
likely to transform bilayers into reversed structures when a is too small for rods or
spheres to occur.
The contributions listed above can be represented by an equation of the
form
g
a
a
rC +
+
=
)(
0
γ
μ
..........................................................................................(2)
where μo is the energy per surfactant molecule, and C(r) is a curvature free-energy
expression, being a function of micelle radius; the alkyl-chain/water repulsion is
given by γa, where γ represents surface tension, and all other contributions are
included in g. One could employ expressions involving higher powers of a but
these do not qualitatively alter the description,(183) nor does inclusion of an
allowance for the area of the chain/water interface occupied by the head groups.(184)
To summarize, the balance of forces indicated in Fig. 1.4.3 determines the
aggregate shape within packing constraints. A large repulsion in plane y gives
water continuous phases while small repulsions give reversed micelles. With large
repulsions spherical micelles occur, where a (sphere) < a (bilayer) [provided that a
(sphere) ≥ as
c]. When a is just too small for spheres, one expects rod micelles
(prolate ellipsoids), and as a is reduced further, oblate spheroids (bilayers) with a <
ar
c. The particular values of a where shape changes occur will depend on the
detailed nature of the forces in planes x and y (and z). Reversed phases are
expected at still smaller values of a.
There are two major effects of intermicellar forces. First, if we consider
the micelles formed at the c.m.c. as hard-core particles, then as the volume fraction
of micelles is increased, we will observe order/disorder transitions. Spherical
micelles will close-pack into a regular cubic array. Long rod micelles will form a
hexagonal array, while a lamellar phase will form with bilayers. The limiting
volume fractions are 0.74 for face-centred cubic and 0.91 for a hexagonal phase.
Theoretically, a lamellar phase can occur without water being present, i.e. limiting

Page 42
Chapter 1
34
volume fraction = 1. (In practice it is well known that anhydrous soaps form a
lamellar phase at high temperatures.(185, 189)
The actual volume fractions for
order/disorder transitions appear to be in the range 0.7-0.8 of the close-packed
volume values(192) (i.e. ca. 0.5 and ca. 0.7 for spheres and rods, respectively). For
lamellae, two large bilayers in any volume will be aligned to some extent because
they cannot pass through each other. Thus the lowest theoretical volume fraction
for a lamellar phase is just above the c.m.c. However, van der Waals attractions
between bilayers can be sufficiently large to overcome entropy and interbilayer
repulsions (193) causing phase separation of a lamellar phase where the water layers
do not swell indefinitely. (Similar considerations apply to disordered oblate micellar
solutions, hence the occurrence of the cloud point.(180, 185)) Thus instead of a dilute
solution of large bilayers, a lamellar-phase + water dispersion occurs just above the
c.m.c. Starting with various different micellar shapes, one would expect the
following phase sequences (the numbers refer to the volume fraction of micelles
required for the phase transition):
A hexagonal phase occurs at volume fractions above the close-packing limit of
spherical micelles because the rods can relieve some of the surface strain due to
curvature better than the lamellar phase. In the second and third cases a lamellar
phase occurs only when the close-packing volume of rods is exceeded. In practice,
the limit of mesophase formation is often determined by the stability of a crystalline
surfactant phase, particularly with ionic or zwitterionic compounds. This
description has nothing to say about the crystalline state.
Of course, the repulsions between micelles are not of the hard-sphere type,
but occur at a range of distances from the micelle surface. This soft-core repulsion
can extend to many micelle diameters if ionic surfactants with low concentrations of
added electrolyte are present. For zwitterionic surfactants (lecithins) the existence
of an apparent ‘hydration’ repulsion force has been demonstrated.(193)
The force
can be thought of as arising from interactions between hydrogen-bonded water
network structures on adjacent micelle surfaces. Theoretical considerations(194)
oblate spheroid → lamellar phase
(disc micelles)
prolate spheroid → hexagonal phase → lamellar phase
(rod micelles)
spherical phase → cubic phase → hexagonal phase → lamellar phase
0-1.0
0.7
0.91
0.5
0.74
0.91

Page 43
Chapter 1
35
suggest that it could have an exponential fall-off with increasing distance from the
surface. For polyoxyethylene surfactants we expect that this force will be
accompanied by another repulsive force due to limitations on the conformation of
EO groups arising from steric hindrance when micelles are close together.
One possible consequence of the soft-core repulsions is the formation of
cubic structures other than the face-centred variety. Recent results suggest that
primitive or body-centred structures can occur in addition to face-centred cubic.
Thus the minimum volume fraction for cubic structures is further reduced. On
increasing volume fraction, one might expect to observe the sequence: primitive
cubic → body-centred cubic → face-centred cubic → other phases, with the volume
fraction for these transitions being in the ratio 0.52 : 0.68 : 0.74, respectively.
However, the actual structures observed could vary from this sequence according to
the form of the intermicellar repulsions.
The soft-core repulsion between micelles has a second effect: it causes the
value of a to decrease with increasing volume fraction (unless a is limited by
alkyl-chain packing). This leads to an increase in micelle size. More importantly,
a micelle shape transition of the type sphere → rod → bilayer can occur with a > ac,
where the unfavorable curvature energy is balanced by a contribution from micelle
interactions. This type of transition will be accelerated when a approaches the
packing limits of spheres or rods. If the repulsions between the new shapes are
sufficiently large, then an ordered phase can form immediately. Thus one could
observe the sequences:
What is observed in practice will depend on the relative magnitudes of the forces
and their detailed dependence on area (a) or distance (micelle separation, i.e.
concentration). An illustration of the possible phase sequences at different
curvatures and as a function of volume fraction is given in Fig. 1.4.4. Note that the
micelle shape transitions in disordered solutions occur over a range of volume
fractions (dotted lines) while transitions involving mesophases occur at constant
volume fraction. Also, it was emphasized that mesophases occur from interactions
between micelles. If interactions are absent then ordered phases cannot form.
spherical micelles → hexagonal phase → lamellar phase
(disordered)
rod micelles→ lamellar phase
(desordered)

Page 44
Chapter 1
36
So far the effects of entropy were ignored. This is because there was no obvious
model to allow for the change caused by forming small micelles from large ones
which includes the effects of alkyl-chain conformations and changes in
water/head-group orientations. However, any effect is likely to scale as kT/N,
where N is the aggregation number of the small micelle. For a given area per
molecule, N will be a function of r2, where r is the micelle radius, which in turn is
proportional to the alkyl-chain length. Thus the entropy contribution will decrease
rapidly with increasing alkyl-chain length.
Figure 1.4.4 Schematic illustration of mesophase structures as a function of
surfactant volume fraction and increasing curvature. Dotted lines indicate micelle
shape transitions. Mesophase regions are cubic (I1), hexagonal (H1) and lamellar
(Lα).

Page 45
Chapter 1
37
1.5. Summary
The energy properties of metallic lithium are the most advantageous
because its redox potential is the highest and its atomic mass is small. However
when the lithium metal anode is charge and discharged several times, it is often
deposited as a loose, inhomogeneous dendrite layer containing badly with the
operating electrode. Moreover, some of the deposited metal may simply fall away
from the electrode surface to the bottom of the cell. Thus, in order to solve this
problem for achievement of a reversible lithium battery, a suitable electrolyte and/
or anode must be developed.
In the approach from electrolytes, the possibility of internal shorting, leaks,
and producing combustible reaction products at the electrode surfaces, existing in
the liquid electrolytes, is eliminated by the presence of a solid polymer electrolyte.
The polymers, in general, being light-weight and non-combustible materials can be
fabricated to requirements of size and shape, thus offering a wide range of designs.
Nevertheless, the polymer electrolytes should exhibit ionic conductivities, at least,
of the order of 10-3 to 10-2 S/cm at room temperature and play the role of a separator,
played by the liquids. The polymer should also allow good cycle lives, low
temperature performances, and good thermal and mechanical strengths in order to
withstand internal temperature and pressure buildup during the battery operation.
These problems were solved at some level by the application of gel electrolyte
system. Nevertheless these techniques need expensive materials, complicated
process, and/or hazardous solvent. Further improvements are required for
industry.
In the approach from anodes, the lithium ion secondary batteries which
indicate high power and high capacity were commercialized due to the development
of carbon anodes. Nevertheless, higher capacity is required to anodes of lithium
ion secondary batteries, whereat alloy anodes, i.e. Sn, Si, have been investigated.
The problem of these anodes, that is the capacity degradation due to the volume
change during charge-discharge cycles, are being improved by alloying of inactive
metals with lithium. Now the interests on alloy anode are next stage, that is
improvement for high rate property.
Finally the phase separated structures described in this chapter have a
tremendous amount of potential for application of electrochemistry, because these
materials can be selected as electrolyte.
In this thesis, the improvements of the materials for lithium secondary
batteries by applying the phase separated structures were discussed.

Page 46
Chapter 1
38
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Page 55
Chapter 2:
PEO/PS polymer blend gel electrolyte

Page 56

Page 57
Chapter 2
49
2.1. Introduction
As it has been introduced in chapter 1, polymer electrolytes for lithium
secondary batteries have been studied for years for safety issues in future lithium
batteries, and gel polymer electrolytes are one of the most promising candidates to
be considered in the coming first stage because of its higher ionic condudtivity.
However, gel polymer electrolytes have two major problems to be realized in
practical use that relates to the nature of the polymer-solvent interaction. For
example, polymers with weak solvent interaction such as PVdF
(polyvinylidenefluoride)(1-5) and PAN (polyacrylonitrile)(6-12) form unstable gels, and
are likely to loose the plasticizer from the gel matrix, even though they exhibit good
mechanical strength. On the other hand, polymers with strong solvent interaction
such as PEO (polyethylene oxide)(13-18) and PMMA (polymethyl-methacrylate)(6, 19-26)
form stable gels with the plasticizer and show higher ionic conductivity than that of
solid polymer electrolytes, while the gels in the system have poor mechanical
strength. It is indeed difficult to achieve both properties of high ionic conductivity
and mechanical strength.
However, recently, many excellent studies were undertaken to tackle these
problems, i.e., a co-polymerization of the host polymer, a three-dimensional
cross-link of the host polymer, or a polymerization of room temperature molten salt
(ionic liquid)(27-33). In these studies, S. Passerini et al. proposed a polymer blending
method to get the novel polymer matrix of gel electrolyte for lithium secondary
batteries(28, 29) with low-cost, green, and easy processes. This proposed method was
to use the simple hot blending method for combining common polymers. Two
immiscible polymers, PEO and polystyrene (PS), were blended to obtain the
electrolyte demonstrating both merits of ionic conduction and mechanical strength.
The PEO was plasticized to provide the ionic conductivity whereas the PS as the
framed network to yield the membrane mechanical strength. This blended
polymer system showed co-continuous morphology. This morphology was formed
with the IPN (interpenetrated polymer network) of the two polymers. Figure 2.1.1
illustrates a schematic model for a co-continuous polymer blend system, where
lithium ion transfers through the ion conductive phase. The phase of frame of the
film keeps the mechanical strength when the ion conductive phase is plasticized.
Figure 2.1.2 shows the photographs comparing the PEO film with and without PS,
i.e., PEO film (a) and IPN film (b) after the plasticizing process. As seen in Fig.
2.1.2, the IPN film kept its shape, however the PEO films with plasticizing process
could not keep its shape due to its soft and viscous texture. This plasticized IPN

Page 58
Chapter 2
50
electrolyte was expected to posses a high ionic conductivity and a good mechanical
strength, while there are some concerns to loose the ion conductive phase by flowing
out of the soft plasticized PEO phase from PS frame with the large amount of
plasticizer. In the practical use, a higher conductivity is requested, and one of the
solutions would be to increase the amount of plasticizer for gel electrolytes.
The aim of this chapter is to improve the retention of ionic conductive phase on
the PEO/PS polymer blend gel electrolyte system by the application of PEO-LiBF4
complex to the polymer blend gel electrolyte. There are some reports on that a
complex is formed when PEO and LiX are mixed(34, 35). The lithium ion serves as
the bridging of PEO chains as shown in Fig. 2.1.3. The PEO-LiX complex should
keep high viscosity in spite of impregnation with the plasticizer. Therefore, the
formation of PEO-LiX complex is to give the potential to prevent that conductive
phase runs out from the frame phase in the IPN. Characteristics of the gel
polymer electrolyte made with IPN film of PEO-LiBF4 complex and PS, and the
possibility of the electrolyte formed with this IPN matrix were examined for the
future Li metal secondary batteries.
Figure 2.1.1 Schematic model for co-continuous polymer blend interpenetrated
polymer network system.
PS : The frame of the film
PEO-LiX : Ion conductive phase
PS : The frame of the film
PEO-LiX : Ion conductive phase
Enabling ionic conduction
Imparting mechanical strength
Enabling ionic conduction
Imparting mechanical strength

Page 59
Chapter 2
51
Figure 2.1.2 Photographs of the films of PEO (a) and PEO-PS IPN (b) films which
are plasticized with LiBF4/EC-PC.
Figure 2.1.3. Schematic illustration of the formation of PEO-LiX complex.
(a)
(b)
PEO polymer chain
Li+
Li+
O
O
O
O
O
O
O
O
PEO polymer chain
Li+
Li+
Li+
Li+
O
O
O
O
O
O
O
O

Page 60
Chapter 2
52
2.2. Experimental
Sample preparation and characterization
PEO (POLYOX WSR N-300, Union Carbide; Mw=400,000), PS (STYRON 656D,
Dow Chemical; Mw=30,000) and LiBF4 (Stella Chemifa, Japan) were dried under
dry airflow (dew point below -60 oC). PS was porphyrized before hot-blending.
The mixture of weighted PEO, PS and LiBF4 powders was hot-blended with blend
equipment (Small Amount Blend Machine, Imoto Machinery Co. LTD, Japan). The
blending conditions were as follows, PEO: PS = 48 : 52 by weight., O : Li = 6 : 1 by
mol, 180 oC, 30 min., 107 rpm. The blended samples were quenched in liquid N2.
Adequate amounts of the obtained blends were hot-pressed at 120 oC under 30 MPa
with sandwiching between aluminum foils to form thin films. The pressing time
was reduced to a minimum by setting the hold-under-pressure time to zero. The
thickness of the obtained films was c.a. 100 μm.
To investigate the morphology of the obtained film by hot-blending, the
PEO-LiBF4 phase in the IPN films was removed by dissolution of PEO-LiBF4 in
water. After the dissolution, the samples were rinsed with water and dried under
vacuum. The sample was sputtered by Pt-Pd for observation with Scanning
Electron Microscope (SEM). SEM images were taken with a Hitachi, S-2500CX.
The formation of PEO-LiBF4 complex was confirmed by means of Raman
spectroscopy. Raman spectra were recorded using JASCO spectrometer equipment
(RFT-800, FT/IR-800), with the condition as follows, excitation laser: He-Ne, 1064
nm, 10 mW, resolution: 4 cm-1, scanning rate: 0.5 mm/s.

Page 61
Chapter 2
53
Electrochemical characterization
To make the obtained film into the gel electrolyte, a LiBF4/EC (ethylene
carbonate)-PC (propylene carbonate) (1:1 vol.) was used as a plasticizer.
The PEO-LiBF4 phase was plasticized with the solution containing the desired
concentration of Li+, which was quickly absorbed into the PEO-LiBF4 phase, by
wetting the surfaces of the IPN film. SUS / IPN-gel electrolyte / SUS cell was
assembled and the ionic conductivity of IPN gel electrolyte was evaluated by a.c.
impedance measurements (FRA 5080, NF, Japan, GPIB Potentiostat / Galvanostat
HA-501G, Hokuto Denko, Japan). The impedance spectrum was measured in the
constant potential mode by sweeping the frequencies from 20 kHz to 1 Hz range at
a.c. amplitude of 10 mV. And the temperature dependency of the conductivity of
the gel electrolyte was also examined in the temperature range of 25 to 70 oC in the
same way with Ni / IPN-gel electrolyte / Ni cell.
Li / IPN-gel electrolyte / Li cell was assembled, and the chemical stability at the
interface between the IPN gel electrolyte and lithium metal was evaluated by a.c.
impedance measurements with the conditions of a.c. amplitude voltage of 10 mV
and of the frequency range of 20 kHz - 100 mHz.
Li / IPN-gel electrolyte / Au cell was assembled to evaluate the oxidation stability
of the IPN gel electrolyte by linear sweep voltammetry (LSV) (Function Generator
HB-104, Hokuto Denko, Japan, Potentiostat / Galvanostat HA-501G, Hokuto Denko,
Japan) at the scan rate of 20 mV/s in the potential ranging from the open circuit
potential (OCP) to 6.0 V vs. Li/Li+.
The charge-discharge test was conducted to evaluate the performance of the IPN
gel electrolyte as an electrolyte for lithium secondary batteries. Li / IPN-gel
electrolyte / LiCoO2 cell was assembled with the LiCoO2 supplied from Japan
Chemical Industry. The charge-discharge test was performed (HJ1010mSM8,
Hokuto Denko, Japan) under constant current-constant voltage (CC-CV) mode at
C/5 rate of loading current and the voltage upper limit of 4.2 V for charging and CC
mode for discharging with C/5 rate of current with cut-off voltage of 3.0 V.

Page 62
Chapter 2
54
2.3. Results and Discussion
2.3.1. Confirmation of phase separated structure
There are some reports that a blended polymer composed of PEO and PS
indicated the phase separated structure(28, 29). Therefore, to confirm the structure
of the polymer blend film applied PEO-LiBF4 complex, it was observed with SEM.
Because PEO and LiBF4 can be dissolved in H2O, the polymer blend film was dipped
in distilled water to confirm the matrix formations, where the PS phase remains
after PEO and LiBF4 are dissolved into H2O. The remained PS phase of the
polymer blend film was observed with SEM, shown in Fig. 2.3.1. The pore in the
image is the part occupied by PEO-LiBF4 phase. From the morphology of the PS
phase in the resulting film, the IPN structure was also demonstrated in the
PEO-LiBF4/PS polymer blend system.
Fig. 2.3.1 SEM image of PS phase of polymer blend film. The PEO-LiBF4 phase was
removed by rinsing the polymer blend with H2O
50 μm50
μm50
μm

Page 63
Chapter 2
55
2.3.2. Confirmation of PEO-LiBF4 complex formation
In order to confirm the formation of PEO-LiBF4 complex expected during
hot-blending, Raman spectra of the IPN film were measured and are shown in Fig.
2.3.2. In Fig. 2.3.2 the spectra of four samples are indicated, i.e., hot-blended
PEO-LiBF4, cooled down PEO melted with the heat treatment of PEO powder,
LiBF4 powder, and mixture of PEO powder with LiBF4 powder. The spectra of the
heated and cooled PEO without LiBF4 and the mixture of PEO powder and LiBF4
powder show similar peaks except the peak at 800 cm-1 that appeared in the
spectrum of the mixture of PEO powder and LiBF4 powder. In contrast, both of the
spectra of LiBF4 powder and the mixture of PEO powder and LiBF4 powder has the
peak at 800 cm-1. From these results, it is concluded that the peak at 800 cm-1 is
due to LiBF4, and the operation of melting PEO powder without LiBF4 does not
alter the solid-state properties of PEO. On the other hand, the peak around 800
cm-1 in the spectrum of the hot-blended PEO-LiBF4 shifted to lower wave number,
777cm-1. Generally, the peak of BF4
- anion is observed at 777cm-1, whereas the
peak of LiBF4 powder was observed at 800 cm-1. This might be the effect of Li+
cation. It might be considered that the peak of BF4
- anion changed back into the
original peak by the formation the complex of PEO and Li+. This result suggests
that the complex of PEO and LiBF4 may be formed as reported in the literatures(36,
37). The formation of PEO-LiBF4 complex in the matrix is thus confirmed.
Furthermore, it is expected that the ion conducting phase would keep high viscosity
by the interaction between PEO and LiBF4, when large amount of plasticizer is
added.

Page 64
Chapter 2
56
Fig. 2.3.2 Raman spectra of four samples of hot-blended PEO-LiBF4, PEO melted
once and cooled from PEO powder, LiBF4 powder, and mixture of PEO powder +
LiBF4 powder. The excitation laser: He-Ne, 1064 nm, 10 mW, resolution: 4 cm-1,
scanning rate: 0.5 mm/s.
100
400
700
1000
1300
1600
Wavenumber / cm
-1
In
ten
sity / a.u.
LiBF4 Powder
PEO Powder + LiBF4 Powder
PEO Chunk
PEO-LiBF4 Complex
600
700
800
900
1000
Wavenumber / cm
-1
Intensity / a.u.
LiBF4 Powder
PEO Powder + LiBF4 Powder
PEO Chunk
PEO-LiBF4 Complex

Page 65
Chapter 2
57
2.3.3. Electrochemical properties of IPN gel electrolyte
Figure 2.3.3 shows the variation of the ionic conductivity of IPN gel electrolytes
as function of the concentration of the supporting electrolyte in the plasticizer. For
the sake of comparison, the ionic conductivity of gelled PEO-LiBF4/PS polymer
blend film and PEO/PS polymer blend film were also measured at room
temperature. Uptake of the plasticizer is calculated according to the following
formula (1).
100
][
][
][
][
[%]
4
×
+
+
=
gr
Plasticize
g
LiBF
g
PEO
gr
Plasticize
r
plasticize
the
of
Uptake
(1)
The ionic conductivity of PEO-LiBF4/PS is higher than that of PEO/PS in the range
of the investigated concentration. This is presumably due to the anion of LiBF4
added during the hot-blending process. The cation Li+ is used to form the complex,
while the anion BF4
- is comparatively free in the conductive phase in the IPN.
Besides, the ionic conductivity of PEO-LiBF4/PS has a peak at 0.5 M, and that of
PEO/PS has a peak at 0.7 M. This is presumably due to the high concentration of
the supporting electrolyte in the PEO-LiBF4/PS. LiBF4
added during the
hot-blending process would increase the concentration of LiBF4
in the
PEO-LiBF4/PS gel electrolyte. However, accurate concentrations of Li+ and BF4
-
ion were unidentified, because the percent of LiBF4 that was used for the formation
of PEO-LiBF4 complex and the degree of LiBF4 dissociation in the matrix are
unknown. Therefore, the increase of concentration of LiBF4 causes solvated ions,
ternary ions and other ion complexes to start forming at lower concentration of the
supporting electrolyte in the plasticizer. Owing to these multiple factors, the ionic
conductivity as a function of the supporting electrolyte in the plasticizer showed the
phenomena in Fig. 6. In the following investigation, 0.5 M LiBF4 / EC-PC would be
used to plasticize the PEO-LiBF4/PS films, and the 0.7 M LiBF4 / EC-PC films would
be used to plasticize PEO/PS.

Page 66
Chapter 2
58
Fig. 2.3.3 Room temperature ionic conductivity of IPN gel electrolyte as a function
of the concentration of the supporting electrolyte in plasticizer. Uptake of the
plasticizer was 50 wt%. The impedance spectrum was measured in the constant
potential mode by sweeping the frequencies from 20 kHz to 1 Hz range at a.c.
amplitude of 10 mV.
0
0.2
0.4
0.6
0.8
1
0
0.4
0.8
1.2
1.6
Concentration of the supporting electrolyte / M
Ionic conductivity / m
S
cm
-
1
PEO-LiBF4 / PS
PEO / PS

Page 67
Chapter 2
59
The temperature dependency of the conductivity of the of the PEO-LiBF4/PS
polymer blend gel electrolyte is shown in Fig. 2.3.4. The figure shows a straight
line, indicating that the conductivity of the PEO-LiBF4/PS polymer blend gel
electrolyte obeys Arrhenius law. This implies that the conductive environment of
Li in the polymer gel electrolyte is liquid like and remains unchanged in the
investigated temperature region. From Fig. 2.3.4, it is clear that the conductivity
increases evidently with increasing temperature and the conductivity at 25 oC was
calculated to be 1.14 mS / cm. The activation energy (Ec) of conductivity for this
film can be calculated from Arrhenius equation
⎛ -
=
RT
Ec
exp
0
σ
σ
where T is temperature on the Kelvin scale and σ0 is a proportional constant. Ec
value of this polymer film is 20.64 kJ/mol and this value is lower than the value of
PEO gel electrolyte reported in the literature of 30.8kJ/mol(38).

Page 68
Chapter 2
60
Fig. 2.3.4 Arrhenius plot of ionic conductivity for gel electrolyte containing 0.5 M
LiBF4/EC-PC in the temperature range of 25 to 70 oC.
-3
-2.5
-2
2.80
2.90
3.00
3.10
3.20
3.30
3.40
1000/T / K
-1
L
og (cond./S
cm
-1
)

Page 69
Chapter 2
61
Figure 2.3.5 shows the Cole-Cole plots of the Li / IPN-gel electrolyte / Li cell
investigated by a.c. impedance measurement. For the sake of comparison,
PEO-LiBF4/PS and PEO/PS gel electrolytes were applied. Assuming these results
due to the interfacial resistance derived from the charge transfer and the surface
film resistance on lithium electrode, the equivalent circuit was constructed as
shown in Fig. 2.3.6, with which the fitting were carries out on these results. The
fitted data of the Cole-Cole plots on PEO-LiBF4/PS and PEO/PS gel electrolytes are
shown in Table 2.3.1 and Table 2.3.2 respectively. From the fitted data of the
Cole-Cole plots, the semicircles observed in higher and lower frequency regions
should be ascribed due to the interfacial resistance and the surface film resistance
on lithium electrode respectively, because the value of the electric double layer
capacitance between gel electrolyte and lithium electrode was the order of μF/cm2.(39,
40)
Here the interfacial resistance was focused on because its increase was more
remarkable. The diameter of the semi-circle observed in higher frequency region is
assigned to be the sum of interfacial resistances of the two interfaces of the gel
polymer electrolyte and lithium metal in the sandwiched type cell. Figure 2.3.7 (a)
shows the change in the interfacial resistance of the gel polymer electrolyte and
lithium metal as a function of the time, and the normalized change by the
interfacial values of both resistances are shown in Fig.2.3.7 (b). The interfacial
resistance was calculated by the radius of fitting curve for Fig. 2.3.5 divided by 2
with the assumption of the interfaces of both side of gel electrolyte is almost the
same. The resistance of PEO-LiBF4/PS gel electrolyte shows lower than that of
PEO/PS gel electrolyte, while both of the normalized change of PEO-LiBF4/PS and
PEO/PS gel electrolyte are the same. This is presumably due to that a small
amount of plasticizer covered the Li surface in contact with the PEO-LiBF4/PS gel
electrolyte. Because PEO-LiBF4/PS gel electrolyte contains large amount of
plasticizer compared to the PEO/PS. The Li surface covered with plasticizer has less
chance to contact with PEO or PS polymer molecules and the solid electrolyte
interface (SEI) formed with the reaction between the plasticizer and Li might have
smaller interfacial resistance than that formed by the reaction between Li and PEO
or PS used this study. This suggests that application of PEO-LiBF4 complex to IPN
system improved the interfacial properties of the gel polymer electrolyte and
lithium metal.

Page 70
Chapter 2
62
Fig. 2.3.5 Cole-Cole plot of Li / IPN-Gel / Li cell. The IPN film was plasticized with
plasticizer containing electrolyte. (a) PEO-LiBF4/PS gel electrolyte, (b) PEO/PS gel
electrolyte. The impedance spectrum was measured in the constant potential mode
by sweeping the frequencies from 20 kHz to 100 mHz range at a.c. amplitude of 10
mV.
0
1000
2000
3000
4000
0
1000
2000
3000
4000
Zreal / Ω cm
2
-Z
im
a
g
/ Ω
cm
2
1h
4h
36h
100h
385
915h
0
5000
10000
15000
20000
0
5000
10000
15000
20000
Zreal / Ω cm
2
-Z
im
a
g
/ Ω
cm
2
1h
4h
36h
100h
385h
915h
(b)
(a)

Page 71
Chapter 2
63
Fig. 2.3.6 Assumed equivalent circuit for the fitting of impedance spectra. Rs:
electrolyte resistance, Ri: interfacial resistance, Rf: surface film resistance, Cdl:
electric double layer capacitance, Cf: capacitance due to surface film
Table 2.3.1 Parameters of determined from fitting the experimental impedance
spectra of PEO-LiBF4/PS gel electrolyte shown in Fig.2.3.5 (a) to the equivalent
circuit of Fig. 2.3.6.
Time [hour] Rs [Ωcm2]
R1 [Ωcm2]
C1 [μF/cm2]
R2 [Ωcm2]
C2 [μF/cm2]
0
49
125
1.8
36
8990
1
39
153
2.2
90
9000
4
34
230
2.2
126
9000
36
34
547
2.4
283
1000
100
32
727
2.4
398
1800
385
32
1625
2.8
647
2000
915
32
2415
2.8
802
2000
Table 2.3.2 Parameters of determined from fitting the experimental impedance
spectra of PEO/PS gel electrolyte shown in Fig. 2.3.5 (b) to the equivalent circuit of
Fig. 2.3.6.
Time [hour] Rs [Ωcm2]
R1 [Ωcm2] C1 [μF/cm2]
R2 [Ωcm2] C2 [μF/cm2]
0
500
931
1.6
770
30000
1
184
839
1.4
770
30200
4
186
1145
2.0
768
30600
36
163
1845
2.6
1100
30000
100
152
2729
2.6
1200
30000
385
150
7027
2.4
2003
30000
915
148
11700
3.0
3000
30000
R
s
R
i
C
dl
C
f
R
f
R
s
R
i
C
dl
C
f
R
f

Page 72
Chapter 2
64
Fig. 2.3.7 Interfacial resistance change of Li / various electrolytes over time during
static storing. The interfacial resistances in this figure were calculated from
impedance data. (a) variation per hour, (b) change per hour. Uptake of the
plasticizer is 80 wt%.
0
1000
2000
3000
4000
5000
6000
1
10
100
1000
Time / hour
In
terfacial restan
ce / Ω
cm
2
PEO-LiBF4 / PS (0.5M LiBF4 / EC-PC)
PEO / PS (0.7M LiBF4 / EC-PC)
0
200
400
600
800
1000
1200
1400
1600
1800
1
10
100
1000
Time / hour
R
ate of ch
an
ge of R
i
/ %
PEO-LiBF4 / PS (0.5M LiBF4 / EC-PC)
PEO / PS (0.7M LiBF4 / EC-PC)
C
h
an
ge of R
i
/ %
0
1000
2000
3000
4000
5000
6000
1
10
100
1000
Time / hour
In
terfacial restan
ce / Ω
cm
2
PEO-LiBF4 / PS (0.5M LiBF4 / EC-PC)
PEO / PS (0.7M LiBF4 / EC-PC)
0
200
400
600
800
1000
1200
1400
1600
1800
1
10
100
1000
Time / hour
R
ate of ch
an
ge of R
i
/ %
PEO-LiBF4 / PS (0.5M LiBF4 / EC-PC)
PEO / PS (0.7M LiBF4 / EC-PC)
C
h
an
ge of R
i
/ %
(a)
(b)

Page 73
Chapter 2
65
Figure 2.3.8 shows LSV curves of Li / IPN-gel electrolyte / Au cells. For the
sake of comparison, PEO-LiBF4/PS gel electrolyte (Fig.2.3.8 (a)) and PE
(polyethylene) separator with electrolytic solution (Fig.2.3.8 (b)) were also
investigated. The same solution was used as plasticizer of the polymer blend film
and electrolytic solution. From this result, both of the curves in PEO-LiBF4/PS gel
electrolyte and 0.7 M LiBF4 / EC-PC stand about 4.5 V vs. Li/Li+. The oxidation
current starting at 4.5 V vs. Li/Li+ is ascribed to the oxidation of 0.7 M LiBF4 /
EC-PC. This result suggests that PEO-LiBF4/PS can be applied to Li / LiCoO2 cell
as the electrolyte in the range between 0 and 4.2 V.

Page 74
Chapter 2
66
Fig. 2.3.8 Linear sweep voltammogram of Li / electrolyte / Au cell with (a)
PEO-LiBF4 / PS gel electrolyte, (b) polyethylene separator with electrolyte solution
at the scan rate of 20 mV/s.
C
u
rren
t d
en
sity / m
A
cm
-2
Cell voltage / V vs. Li
0.1
0.2
0.3
0.4
0
1
2
3
4
5
6
0
C
u
rren
t d
en
sity / m
A
cm
-2
Cell voltage / V vs. Li
0.1
0.2
0.3
0.4
0
1
2
3
4
5
6
0
0.1
0.2
0.3
0.4
0
1
2
3
4
5
6
0
0.1
0.2
0.3
0.4
0
1
2
3
4
5
6
0
C
u
rren
t d
en
sity / m
A
cm
-2
Cell voltage / V vs. Li
0
1
2
3
4
5
6
0.1
0.2
0.3
0.4
0.5
C
u
rren
t d
en
sity / m
A
cm
-2
Cell voltage / V vs. Li
0
1
2
3
4
5
6
0.1
0.2
0.3
0.4
0.5
0
1
2
3
4
5
6
0
1
2
3
4
5
6
0.1
0.2
0.3
0.4
0.5
0.1
0.2
0.3
0.4
0.5
(b)
(a)

Page 75
Chapter 2
67
Figure 2.3.9 shows the discharge capacity per gram of cathode. For the sake of
comparison, PEO-LiBF4/PS polymer blend film and PE separator with electrolytic
solution were used. The capacity of the cell using PEO-LiBF4/PS is a little lower
than that of the cell using PE separator with electrolytic solution. Besides, there
are some cycles with deathly low capacity. However, the degradation rate is
unimportant. The capacity per gram of cathode keeps 130 mAh/g over 30 cycles.
Fig. 2.3.9 Charge-discharge test of Li / (PEO-LiBF4/PS) gel electrolyte / LiCoO2 cell
and Li / (0.7 M LiBF4 / EC-PC) / LiCoO2 cell. The charge-discharge rate was C/5 in
the CC-CV charge mode and CC discharge mode, and the voltage range was 3.0 –
4.2 V.
0
20
40
60
80
100
120
140
160
180
0
5
10
15
20
25
30
35
Cycle number / -
D
isch
arge cap
acity / m
A
h
g
-1
Discharge Capacity (PEO-LiBF4 / PS)
Discharge Capacity (electrolytic solution)

Page 76
Chapter 2
68
2.4.
Summary
In this chapter, the possibility of the PEO-LiBF4/PS polymer blend as electrolyte
for lithium secondary batteries. A PEO-LiBF4/PS polymer blend with three-
dimensionally IPN structure was obtained by hot-blending of the powders of PEO,
LiBF4, and PS. PEO-LiBF4 complex was confirmed to be formed in IPN prepared
by the polymer blend method.
The ionic conductivity of PEO-LiBF4/PS gel electrolyte was ~ 1 mS/cm as the
maximum value. The chemical stability of the interface between the IPN gel
electrolyte and lithium metal was successfully improved by applying PEO-LiBF4
complex to IPN. PEO-LiBF4/PS gel electrolyte demonstrated the stability against
oxidation below 4.5 V vs. Li/Li+. This result revealed that LiCoO2 cathode can be
adapted to the IPN gel electrolyte. With the charge-discharge rate of C/5, the
capacity per gram of cathode was found to be over 130 mAh/g over 30 cycles.
Consequently, these results should suggest that PEO-liBF4/PS polymer blend can be
applied to lithium secondary battery as the electrolyte.

Page 77
Chapter 2
69
References
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F. Croce, G. B. Appetecchi, S. Slane, M. Salomon, M. Tavarez, S. Arumugam, Y.
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C. Capiglia, Y. Saito, H. Kataoka, T. Kodama, E. Quartarone and P. Mustarelli,
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A. Manuel Stephan, S. Gopu Kumar, N. G. Renganathan and M. Anbu
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B. Huang, Z. Wang, L. Chen, R. Xue and F. Wang, Solid State Ionics, 91, 279 (1996).
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O. Bohnke, G. Frand, M. Rezrazi, C. Rousselot and C. Truche, Solid State Ionics, 66,
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O. Bohnke, C. Rousselot, P. A. Gillet and C. Truche, Journal of the Electrochemical

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F. Croce, S. D. Brown, S. G. Greenbaum, S. M. Slane and M. Salomon, Chemistry of
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T. Iijima, Y. Toyoguchi and N. Eda, Denki Kagaku, 53, 619 (1985).
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J. Vondrak, M. Sedlarikova, J. Velicka, B. Klapste, V. Novak and J. Reiter,
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Page 79
Chapter 3:
PEO-PS diblock polymer gel electrolyte

Page 80

Page 81
Chapter 3
73
3.1. Introduction
In chapter 2, the improvement of the retention of ionic conductive phase on
the PEO/PS polymer blend gel electrolyte system(1-3) by the application of PEO-LiX
complex was discussed. And the potential of the application of PEO-LiX complex
to the PEO/PS polymer blend gel electrolyte system was demonstrated.(4)
In this chapter, to improve the retention of ionic conductive phase on the
polymer gel electrolyte, another approach which is different from the approach
discussed in chapter 2 is investigated. A di-block copolymer, composed of PEO
chain and PS chain covalently bonded, was applied to polymer gel electrolyte for the
restraint of the ionic conductive phase running out from the frame phase in the IPN,
and also to achieve both properties of high ionic conductivity at RT and excellent
mechanical strength. There are some reports on solid electrolyte of block
copolymer(5-8), however a study on gel electrolyte of block copolymer has not been
reported. It is expected that the gel electrolyte formed by self-assembling of the
di-block copolymer has micro phase separation structure(9), that is IPN structure.
The characteristic and possibility of the prepared PEO-PS di-block copolymer gel
electrolyte were examined for the future lithium secondary batteries.

Page 82
Chapter 3
74
3.2. Experimental
Sample preparation and characterization
PEO-PS di-block copolymer (P1614-SEO, Sowa Science Co., Mn (PEO /PS) =
29,000 / 51,000) was dried under dry airflow (dew point below -60 oC). LiBF4 /
ethylene carbonate (EC) - propylene carbonate (PC) (1:1 vol.) and ethylhexylacetate
were used as plasticizer for PEO and PS respectively. The plasticizer for PEO was
prepared by dissolving LiBF4 (Stella Chemifa corporation, Japan) into EC-PC
(Kishida Chemical Co., Ltd.) under dry air condition. As the plasticizer for PS(10),
ethylhexylacetate (Kanto Chemical Co., Inc.) was used as received. The PEO-PS
di-block copolymer, LiBF4 / EC-PC, and ethylhexylacetate were dissolved into
toluene. The mixture solution was cast into a glass fiber separator in Teflon cup.
The solvent was allowed to evaporate at room temperature, covered with a Petri
dish to prevent rapid evaporation of toluene under dry air atmosphere overnight.
Then, the Petri dish was removed for further toluene evaporation. The thickness
of the IPN film was ~ 700 µm, due to the quantity of casted mixture solution and the
thickness of the glass fiber separator. The prepared gel electrolyte was dipped into
the plasticizer, whose Li+ concentration was the same as that used its preparation,
and wiped the residual plasticizer before electrochemical measurements. In this
study, a glass fiber separator was applied just as supporting film so as not to crack
easily when treated coarsely; the glass fiber separator should not affect any
electrochemical properties of polymer gel electrolyte. The polymer gel electrolyte
was enough strong mechanically to keep its shape without the glass fiber separator.
To confirm the phase separated structure of the obtained film, the surface of film
was observed with Atomic Force Microscope (AFM, Veeco, Multimode Dimension
3100). The film for AFM observation was prepared on a Cu substrate by
spin-coating, because the film prepared by casting in Teflon cup was too rough to
observe with AFM. The morphology difference of a prepared IPN film between the
upper side exposed to dry air and the bottom side contacted with Teflon sheet
during evaporating process were investigated with Scanning Electron Microscope
(SEM, Hitachi, S-2500CX).

Page 83
Chapter 3
75
Electrochemical characterization
The ionic conductivity of the polymer gel electrolyte was measured by a.c.
impedance measurement (FRA 5080, NF, Japan, Automatic Polarization System,
HZ-3000, Hokuto Denko, Japan), using Ni / polymer gel electrolyte / Ni cell.
The impedance spectrum was measured in the constant voltage mode at open
circuit potential (OCP) by sweeping the frequencies range from 10 kHz to 100 mHz
with a.c. amplitude of 10 mV. And the temperature dependency of the conductivity
of the gel electrolyte was also examined in the temperature range of 25 to 100 oC in
the same way.
Li half cell with the polymer gel electrolyte was assembled to investigate the
effect of suppressing dendritic deposition, and the chemical stability of polymer gel
electrolyte against Li metal at its static storing. To investigate the effect of
suppressing dendritic deposition of Li, lithium deposition and dissolution were
cycled 21 times at current density of 2.5 mA / cm2 (1st cycle for pretreatment) and
0.5 mA / cm2 (2nd -21st cycle), Li deposition and dissolution time of 2 hours with
Battery Charge / Discharge System (HJ1010mSM8, Hokuto Denko, Japan), then,
the Li / polymer gel electrolyte interfaces were observed.
The chemical stability at the interface between the polymer gel electrolyte and
Li metal was evaluated in light of the interfacial resistance shift at static storing
with time measured by a.c. impedance measurement with the conditions of a.c.
amplitude voltage of 10 mV and of the frequency range of 10 kHz - 100 mHz.
Li / polymer gel electrolyte / Ni cell was assembled to evaluate the oxidation
stability of the polymer gel electrolyte by normal pulse voltammetry (NPV) with
Automatic Polarization System at the scan rate of 50 mV / s in the potential rang
from OCP to the positive potential with the pulse period and width are 0.5 and 0.1
sec respectively.
The charge-discharge test was conducted to evaluate the performance of the
polymer gel electrolyte as an electrolyte for lithium secondary batteries. Li /
polymer gel electrolyte / LiFePO4 cell was assembled. LiFePO4 cathode was
composed of LiFePO4 (battery grade, Aldrich), acetylene black (AB), and PVdF-HFP
gel. The charge-discharge test was performed with Battery Charge / Discharge
System under CC-CV (constant current-constant voltage) mode at C/5 rate of
loading current and the voltage upper limit of 3.6 V with 2 hours for voltage holding
for charging and CC mode for discharging with C/5 rate of current with cut-off
voltage of 2.0 V.

Page 84
Chapter 3
76
3.3. Results and Discussion
3.3.1. Confirmation of phase separated structure
For confirmation of the phase separated structure of PEO-PS di-block copolymer
film, a spin-coated PEO-PS di-block copolymer film with LiBF4 / EC-PC and
ethylhexylacetate, which are plasticizer for PEO and PS respectively, on Cu
substrate was observed with AFM. Figure 3.3.1 shows AFM image of the film,
bright and dark region are PS and PEO phase respectively; it was confirmed from
expanding bright region of the film prepared by the addition of a PS polymer to the
casting solution. It suggested that double gyroid or double diamond structure, that
is interpenetrated polymer network (IPN) structure, was formed in the film
prepared from PEO-PS di-block copolymer, LiBF4 / EC-PC, and ethylhexylacetate.
Fig. 3.3.1 AFM image of spin-coated PEO-PS di-block copolymer film (PEO content
of 36 wt %), with LiBF4 / EC-PC (1:1 vol.) and ethylhexylacetate as plasticizer for
PEO and PS respectively, on Cu substrate.
250nm
0.0
nm
100.0
nm
0.0
nm
100.0
nm
250nm
0.0
nm
100.0
nm
0.0
nm
100.0
nm

Page 85
Chapter 3
77
3.3.2. Electrochemical properties of IPN gel electrolyte
The ionic conductivity of IPN gel electrolyte at room temperature was evaluated.
Figure 3.3.2 shows the variation of the ionic conductivity of interpenetrated
polymer network gel electrolytes as function of the Li+ concentration in LiBF4 /
EC-PC used as the plasticizer for PEO. The ionic conductivity of IPN gel
electrolyte had a peak when 0.84 M LiBF4 / EC-PC was used as the plasticizer; the
maximum ionic conductivity was 2.65 mS/cm. The phenomenon of indicating
maximum as a function of the Li+ concentration of the plasticizer for PEO on the
IPN gel electrolyte is presumably similar to that of liquid electrolyte(11, 12). The
phenomenon in this PEO-PS interpenetrated polymer network gel electrolyte would
be explained by the deficiency of carrier ion, Li+, at lower Li+ concentration in the
plasticizer for PEO and due to the limited mobility of Li+ by forming ternary ions
and other ion complexes at higher Li+ concentration in that. Therefore the ionic
conductivity of IPN gel electrolyte had a peak at a Li+ concentration of the
plasticizer. In this study, the IPN gel electrolyte added 1.12 M LiBF4 / EC-PC was
used in the following electrochemical measurements.

Page 86
Chapter 3
78
Fig. 3.3.2 Ionic conductivity of interpenetrated polymer network gel electrolytes as a
function of Li+ concentration in LiBF4 / EC-PC used as plasticizer at room
temperature. The impedance spectrum was measured in the constant voltage mode
at open circuit potential (OCP) by sweeping the frequencies range from 10 kHz to
100 mHz with a.c. amplitude of 10 mV.
0.1
1
10
0
1
2
3
Li
+
concentration of LiBF4/EC-PC as plasticizer / M
Ionic conductivity / m
S
cm
-1

Page 87
Chapter 3
79
The temperature dependency of the conductivity of the of the PEO-PS di-block
copolymer gel electrolyte is shown in Fig. 3.3.3. The figure shows a straight line,
indicating that the conductivity of the PEO-PS di-block copolymer gel electrolyte
obeys Arrhenius law. This implies that the conductive environment of Li in the
polymer gel electrolyte is liquid like and remains unchanged in the investigated
temperature region. From Fig. 3.3.3, it is clear that the conductivity increases
evidently with increasing temperature and the conductivity at 25 oC was calculated
to be 0.54 mS / cm. The activation energy (Ec) of conductivity for this film can be
calculated from Arrhenius equation
⎛ -
=
RT
Ec
exp
0
σ
σ
where T is temperature on the Kelvin scale and σ0 is a proportional constant. Ec
value of this polymer film is 33.57 kJ/mol and this value is almost same as PEO gel
electrolyte reported in the literature(13).

Page 88
Chapter 3
80
Fig. 3.3.3 Arrhenius plot of ionic conductivity for gel electrolyte containing 1.12 M
LiBF4/EC-PC in the temperature range of 25 to 100 oC.
-3.5
-3
-2.5
-2
-1.5
2.60
2.80
3.00
3.20
3.40
1000/T / K
-1
L
og (cond./S
cm
-1
)

Page 89
Chapter 3
81
The chemical stability of electrolytes is important factor to assemble the cell
with Li anode. To investigate the chemical stability of interpenetrated polymer
network gel electrolyte against Li metal, Li / IPN gel electrolyte / Li cell was
assembled and a.c. impedance measurement was carried out. Figure 3.3.4 shows
the Cole-Cole plots of the Li / IPN gel electrolyte / Li cell measured with the
symmetric cell. Assuming these results due to the interfacial resistance derived
from the charge transfer and the surface film resistance on lithium electrode, the
equivalent circuit was constructed as described in chapter 2, with which the fitting
were carries out on these results. The fitted data of the Cole-Cole plots on PEO-PS
di-block copolymer gel electrolyte and 1.12 M LiBF4/EC-PC as a reference are shown
in Table 3.3.1 and Table 3.3.2 respectively. From the plots fitting and analysis of
the Cole-Cole plots, the diameter of the semi-circle observed in higher frequency
region is assigned to be the sum of interfacial resistances of the two interfaces of gel
polymer electrolyte and Li metal in the sandwiched type cell as described in chapter
2. The interfacial resistance was calculated from the radius of fitting curve for Fig.
3.3.4 divided by 2 with the assumption of the interfaces of both side of gel electrolyte
is the same in this symmetrical cell. Figure 3.3.5 shows the change in the
interfacial resistance between PEO-PS di-block copolymer gel electrolyte and Li
metal analyzed from the Cole-Cole plots in Fig. 3.3.4 as a function of the time. For
comparison, those using PEO/PS polymer blend film(3) and liquid electrolyte, LiBF4 /
EC-PC, in a separator were also analyzed. The interfacial resistance of PEO-PS
di-block copolymer gel electrolyte was lower than that of PEO/PS polymer blend gel
electrolyte; both of those of PEO-PS di-block copolymer gel electrolyte and liquid
electrolyte, LiBF4 / EC-PC, were almost the same. The difference between the
interfacial resistances of PEO/PS polymer blend gel electrolyte and liquid
electrolyte of LiBF4 / EC-PC is existence or nonexistence of PEO, and the difference
between those of PEO-PS di-block copolymer and PEO/PS polymer blend gel
electrolyte is whether PEO phase is fixed or not to PS matrix by covalent bonding.
Therefore the difference of chemical stability between PEO-PS di-block copolymer
and PEO/PS polymer blend gel electrolyte should be derived from the diffusivity of
PEO; in the case of PEO/PS polymer blend gel electrolyte, as the result of the
reaction of solved PEO and Li, the interfacial resistance should be increased. On
the other hand, in the case of PEO-PS di-block copolymer gel electrolyte, the
increase of the interfacial resistance should be able to be suppressed by fixing PEO
phase to PS matrix absolutely. Consequently, the PEO-PS di-block copolymer gel
electrolyte exhibited excellent chemical stability against Li.

Page 90
Chapter 3
82
Fig. 3.3.4 Change in impedance spectra of a half cell during static storage. The half
cell of Li / interpenetrated polymer network gel electrolyte (with 1.12 M
LiBF4/EC-PC used as plasticizer). Inset is enlarged view of the high frequency
domain of the impedance spectra. The data were obtained under the following
conditions: frequency range of 10 k - 0.1 Hz, amplitude voltage of 10 mV, and room
temperature.
0
50
100
150
0
50
100
150
Zreal
/ Ωcm
2
-Z
im
ag
/ Ω
cm
2
1 hr
2 hr
4 hr
8 hr
12 hr
0
10
20
0
10
20
30
40
50

Page 91
Chapter 3
83
Table 3.3.1 Parameters of determined from fitting the experimental impedance
spectra of PEO-PS di-block copolymer gel electrolyte.
Time [hour]
Rs [Ωcm2]
R1 [Ωcm2] C1 [μF/cm2] R2 [Ωcm2] C2 [μF/cm2]
0
30
45
2.4
201
1224
1
35
76
4.2
301
1286
2
40
113
5.0
346
1286
4
40
151
5.2
403
1428
6
40
179
5.4
422
1560
8
40
221
5.2
446
1628
12
40
257
5.4
525
1768
23
40
284
5.4
553
2068
59
40
505
5.2
822
2262
466
40
1264
5.0
1647
2294
Table 3.3.2 Parameters of determined from fitting the experimental impedance
spectra of 1.12 M LiBmeF4/EC-PC electrolyte as reference.
Time [hour]
Rs [Ωcm2]
R1 [Ωcm2] C1 [μF/cm2] R2 [Ωcm2] C2 [μF/cm2]
0
9
29
4.0
176
1768
1
11
58
6.4
295
1900
2
11
82
6.4
378
1894
4
11
128
6.8
481
2016
6
11
160
7.4
525
2252
8
11
193
7.4
566
2356
12
11
241
7.4
660
2356
23
11
332
7.0
662
2722
59
11
589
7.0
929
3276
466
12
1419
7.4
1915
3400

Page 92
Chapter 3
84
Fig. 3.3.5 Interfacial resistance change of Li / various electrolytes over time during
static storing. The interfacial resistances in this figure were calculated from
impedance data.
0
500
1000
1500
2000
2500
3000
1
10
100
1000
Time / hour
Interfacial resistance / Ω
cm
2
diblock polymer gel
polymer blend gel
liquid electrolyte

Page 93
Chapter 3
85
Electrolytes are also required to suppress the growth of dendritic deposition of Li
for lithium secondary batteries. To confirm the suppression effect of the PEO-PS
di-block copolymer gel electrolyte against dendritic deposition Li, Li / IPN gel
electrolyte half cell was assembled and dissolving-deposition reactions were cycled
with a constant current. Figure 3.3.6 shows the voltage profile during lithium
dissolving-deposition reactions cycling. At 1st cycle as pretreatment, the current
of 2.5 mA/cm2 was flown to remove impurity on Li electrode surface by anode
stripping. The larger overvoltage was observed at both of first dissolving and
deposition process. These should be derived from removing impurity on Li
surfaces. In 2nd-21st cycles, the current of 0.5 mA/cm2 was flown to form the
dendritic lithium. Stable overvoltage was observed during 2nd-21st cycles,
indicating that dissolving-deposition reactions occurred stably. After 21 cycles of
dissolving-deposition reactions, the Li / IPN gel electrolyte half cell was
disassembled to observe the surfaces of IPN gel electrolyte contacted with Li. The
dendritic deposition of Li was not observed on one surface as shown in Fig. 3.3.7 (a),
however that was observed on the other surface as shown in Fig. 3.3.7 (b). This
difference can be explained with the SEM images indicating the dense side of IPN
gel electrolyte prepared on Teflon sheet (Fig. 3.3.8 (a)) and the porous side of that
prepared facing air (Fig. 3.3.8 (b)). Though the dendritic deposition of Li grew into
the voids of PEO-PS copolymer at its porous side, the dendritic deposition of Li was
not able to grow by IPN gel electrolyte at its dense side. Consequently, it was
suggested that PEO-PS di-block copolymer gel electrolyte was able to suppress the
dendritic deposition of Li at charge-discharge cycles.

Page 94
Chapter 3
86
Fig. 3.3.6 Voltage profile during lithium dissolving-deposition reactions cycling with
a half cell of Li / interpenetrated polymer network gel electrolyte (1.12 M
LiBF4/EC-PC was used as plasticizer). The conditions of the dissolving-deposition
reactions were as follows: constant current flow of 2.5 mA / cm2 at the 1st cycle and
0.5 mA / cm2 during the 2nd - 21st cycles.
-400
-200
0
200
400
0
20
40
60
80
Time / hour
O
vervoltage / m
V

Page 95
Chapter 3
87
Fig. 3.3.7 Micrographs of interpenetrated polymer network gel electrolyte (1.12 M
LiBF4/EC-PC was used as plasticizer) after 21 cycles with a constant current flow of
2.5 mA / cm2 at the 1st cycle and 0.5mA / cm2 at the 2nd – 21st cycles. (a) Dense side
of the gel electrolyte prepared in a Teflon cup, (b) Porous side of the gel electrolyte
prepared facing air.
Fig. 3.3.8 SEM images of interpenetrated polymer network gel electrolyte (1.12 M
LiBF4/EC-PC was used as plasticizer) before dissolving-deposition cycles of Li. (a)
Dense side of gel electrolyte prepared in a Teflon cup, (b) Porous side of gel
electrolyte prepared facing air.
80μm
80μm
(a)
(b)
(a)
(b)
5 mm
5 mm

Page 96
Chapter 3
88
The electrochemical stability of electrolytes against oxidation is important to
achieve high power batteries. To investigate the electrochemical stability against
oxidation, Li / IPN gel electrolyte / Ni cell was assembled and NPV was carried out,
shown in Fig. 3.3.9. In Fig. 3.3.9, the oxidation current started to flow from 3.0 V
vs. Li/Li+, and then its increasing rate changed around ca. 3.9 V vs. Li/Li+. From
this result, the IPN gel electrolyte should be indicated to stand about 3.9 V vs. Li/Li+
by extrapolating the NPV curve measured with IPN gel electrolyte. The oxidation
current observed 2.0-3.9 V vs. Li/Li+ was presumably due to a decomposition of
impurity. This should be able to explain from the charge-discharge efficiency of the
full cell using IPN gel electrolyte, shown in Fig. 3.3.10. Figure 3.3.10 indicates the
discharge capacity per gram of cathode and charge-discharge efficiency. The
LiFePO4 cathode was selected on the bases of the assumption that the IPN gel
electrolyte is stable under the potential of 3.9 V. The charge-discharge efficiency
except for 1st cycle was ca. 99 %, though that of 1st cycle was 88 %. This result
implies that impurity was almost decomposed at 1st cycle. Therefore it was
suggested that IPN gel electrolyte can be applied to Li / LiFePO4 cell as the
electrolyte in the range between 0 and 3.6 V.

Page 97
Chapter 3
89
Fig. 3.3.9 Normal pulse voltammogram of Li / interpenetrated polymer network gel
electrolyte (1.12 M LiBF4 /EC-PC was used as plasticizer) / Ni cell. The electrode
potential was scanned from open circuit potential to positive potential at a scan rate
of 50 mV / s, and the pulse period and width were 0.5 and 0.1 sec, respectively.
0
0.5
1
1.5
3
3.5
4
4.5
5
5.5
Voltage / V vs. Li/Li
+
C
urrnt dencity / m
A
cm
-2

Page 98
Chapter 3
90
The discharge capacity of the cell using IPN gel electrolyte kept 124 mAh/g with
~ 99 % of coulombic efficiency over 30 cycles. This discharge capacity was lower
than theoretical capacity of LiFePO4, due to IR drop derived from the electrolyte
resistance. However, its improvement should be achieved by thinning the
electrolyte film.
Fig. 3.3.10 Discharge capacity and coulomb cycle efficiency of Li / interpenetrated
polymer network gel electrolyte (1.12 M LiBF4/EC-PC was used as plasticizer) /
LiFePO4 cell. The charge-discharge rate was C/5 (195 μA / cm2) in the CC-CV charge
mode and CC discharge mode, and the voltage range was 2.0 - 3.6 V
.
80
90
100
110
120
130
140
0
10
20
30
Cycle number / -
D
ischarge capacity / m
A
h g
-
1
50
60
70
80
90
100
110
C
ycle efficiency / %

Page 99
Chapter 3
91
Summary
In this chapter, the possible application of a PEO-PS di-block copolymer gel
electrolyte for lithium secondary batteries was investigated. Since the PEO-PS di-block
copolymer gel electrolyte was too stiff to handle, a plasticizer for PS, ethylhexylacetate,
was added to the electrolyte for the casting process. Indications were that the structure
of the electrolyte was a co-continuous interpenetrated polymer network, consisting of
PS domains with a columnar or spherical shape. The IPN gel electrolyte had a high
ionic conductivity (2.65 mS / cm at maximum value) and also good mechanical strength.
It was confirmed visually that the dendritic deposition at the interface between the Li
electrode and the gel electrolyte was suppressed as a result of the electrolyte’s good
mechanical strength. The electrolyte also showed excellent chemical stability when it
contacted the Li electrode, which probably is the result of the covalent bonding between
the PEO chain and the PS chain. Furthermore, the IPN gel electrolyte showed oxidation
stability below 3.9 V vs. Li / Li+. The electrolyte, moreover, was applied to a Li / LiFePO4
cell. With a charge-discharge rate of C/5, the capacity per gram of cathode was found to
be over 120 mAh / g over 30 cycles. This suggests that the PEO-PS di-block copolymer
gel electrolyte can be successfully applied as the electrolyte in lithium secondary
batteries.

Page 100
Chapter 3
92
Reference
1.
F. Croce, L. Persi, F. Ronci and B. Scrosati, Solid State Ionics, 135, 47 (2000).
2.
S. Passerini, F. Alessandrini, T. Momma, H. Ohta, H. Ito and T. Osaka,
Electrochemical and Solid State Letters, 4, A124 (2001).
3.
S. Passerini, M. Lisi, T. Momma, H. Ito, T. Shimizu and T. Osaka, Journal of the
Electrochemical Society, 151, A578 (2004).
4.
T. Momma, H. Ito, H. Nara, H. Mukaibo, S. Passerini and T. Osaka,
Electrochemistry, 71, 1182 (2003).
5.
G. Liu, M. T. Reinhout and G. L. Baker, Solid State Ionics, 175, 721 (2004).
6.
T. Niitani, M. Shimada, K. Kawamura, K. Dokko, Y. H. Rho and K. Kanamura,
Electrochemical and Solid State Letters, 8, A385 (2005).
7.
T. Niitani, M. Shimada, K. Kawamura and K. Kanamura, Journal of Power Sources,
146, 386 (2005).
8.
D. R. Sadoway, B. Y. Huang, P. E. Trapa, P. P. Soo, P. Bannerjee and A. M. Mayes,
Journal of Power Sources, 97-8, 621 (2001).
9.
C. Park, J. Yoon and E. L. Thomas, Polymer, 44, 6725 (2003).
10.
T. A. Martin and D. M. Young, Polymer, 44, 4747 (2003).
11.
Y. Matsuda, M. Morita and T. Yamashita, Journal of the Electrochemical Society,
131, 2821 (1984).
12.
M. Ue and S. Mori, Journal of the Electrochemical Society, 142, 2577 (1995).
13.
T. C. Wen and W. C. Chen, Journal of Applied Polymer Science, 77, 680 (2000).

Page 101
Chapter 4:
Mesoporous Sn Anode electrodeposited
with lyotropic liquid crystals

Page 102

Page 103
Chapter 4
95
4.1. Introduction
There is a strong incentive to develop and characterize non-carbonaceous
materials for use as anodes for lithium ion secondary batteries that deliver higher
capacities than carbon. A Sn anode has been reported to have higher theoretical
capacity (994 mAh/g) than that of carbon (372 mAh/g).(1-3)
However, the Sn anode
has a problem that its cycle life is unsatisfactory because of Sn disintegration due to
the volume change during the charge and discharge cycling of Li-ions. To solve
this problem, Sn-based alloys with inactive elements against lithium such as Ni(4-7),
Fe(8-12), Cu(13-16), Mn(9, 17), and Co(9) have been investigated. Recently, an amorphous
ternary Sn-Co-C anode has been introduced to a practical use.(18)
On the other hand, the study of the synthesis, properties, and possible uses of
mesoporous materials has been devoted, since the discovery of FSM-16(19) and
MCM-41(20), both silica and aluminumsilicate. Already in 1993 it was suggested, on
the basis of mechanistic ideas, that it should be possible to synthesize non-siliceous
materials.(21)
The first examples have been reported already in 1994.(22)
However, for
these materials it had not been possible to remove the template and thus no mesoporous
materials, but only mesostructured materials, could be obtained. The first mesoporous
nonsiliceous frameworks were reported in 1995-1996,(23, 24)
from which rapid
development started. Since the first report by Attard et al.(25), several mesoporous
metals including Sn(26-28) have been prepared by the reduction of the corresponding
metal ions in the presence of lyotropic liquid crystals (LLC)(29, 30) made of nonionic
surfactants(29, 31) which is described in Section 1.4.2. The mesoporous metals have been
mainly prepared by electrodeposition for the reduction of metal ions to form thin films
on conductive substrates. Thus, many mesoporous materials with specific structural
features (e.g., uniform mesopore size and high surface area) have extensively been
investigated. In particular, mesoporous metals with high electroconductivity are very
promising for various electrochemical applications.
In this chapter, the cycle and rate properties of mesoporous Sn anode
prepared from a LLC bath were investigated for lithium secondary batteries. The
mesoporous structure should suppress the influence of the volume change during
the charge and discharge cycling of Li-ions, leading to a longer life. Also, the
mesoporous structure with high surface area can provide a low diffusion resistance
of Li-ions into the electrode, which should contribute to an increase of the reaction
rate.

Page 104
Chapter 4
96
4.2. Experiments
Sample preparation and characterization
A mesoporous Sn was prepared on a copper foil by electrodeposition with a
LLC bath. The LLC including Sn ions was prepared by mixing 0.300 g nonionic
surfactant, octaethyleneglycol monohexadecyl ether (C16EO8), and an aqueous Sn
solution (0.300 g) which was prepared by dissolving both 1.69 g SnCl2・2H2O and
9.15 g of H2SO4 aq. (18.3 mol/L) into water and fixing the volume to 50 mL. The LLC
takes a 2D-hexagonal symmetry, as is confirmed by low angle XRD analysis (Figure
4.2.1). The electrodeposition was carried out at room temperature and H2O
saturated atmosphere with a constant voltage of -100 mV till the current of 1.5
C/cm2 was passed. Sn plate was used as the counter electrode.
And a Sn anode, called a “dense Sn” in this thesis, was electrodeposited on a
copper foil without nonionic surfactant to evaluate its electrochemical properties as
control. The bath used for this preparation is shown in Table 4.2.1(32). The
electrodeposition was conducted under room temperature with current density of 5
mA/cm2 for 166 sec, constant agitation by magnetic stirrer. Pt wire was used as
the counter electrode and an Ag/AgCl electrode was used as reference electrode.
The thickness of the mesoporous and dense Sn anode were 1.2 and 1.0 μm,
respectively, which were confirmed from cross-sectional scanning electron
microscope (SEM, HITACHI S2500CX) image.
The crystalline structure analysis of both Sn anodes was conducted with a
X-ray diffractometer (XRD). The XRD apparatus (Rigaku, Rint-TTR3) with Cu Kα
radiation was equipped with the voltage and current of 50 kV and 300 mA,
respectively. The nano-structure and the electron diffraction (ED) pattern of both
samples were evaluated using a transmission electron microscope (TEM: JEOL,
JEM-2010), with the powdery samples which were scratched from the substrate and
dispersed in ethanol and mounted on a TEM microgrid.

Page 105
Chapter 4
97
Figure 4.2.1 Small-angle X-ray diffraction patterns of lyotropic liquid crystal bath
containing a Sn solution.
Table 4.2.1 Bath composition of a “dense Sn” anode.(32)
Chemical
Concentration
SnSO4
40 g/L
H2SO4
60 g/L
ο-Cresolsulfonic acid (CH3C6H3(OH)SO3H
40 g/L
Polyethylene glycol (Polymerization degree: 20,000)
((CH2CH2O-)n)
100 ppm
1
2
3
4
5
6000
5000
4000
3000
2000
1000
0
2θ / degree
In
te
n
s
ity
/ a
.u
.
3.6 3.8 4.0 4.2 4.4
200
150
100
50
0
2θ / degree
In
te
n
s
ity
/ a
.u
.

Page 106
Chapter 4
98
Electrochemical characterization
Cycle properties were evaluated using conventional glass cells with two
pieces of lithium foil as counter and reference electrodes, and 1 M LiClO4 / ethylene
carbonate (EC) + propylene carbonate (PC) (1:1 vol. %) (as purchased from Kishida
Chemical Co.,Ltd.) as the organic electrolyte. All the cells were sealed under Ar
atmosphere. For evaluation of the cycling, both charge and discharge of Li-ions were
carried out by constant current (CC) mode at current density of 0.994-16.9 A/g
(1-17C rate)in the potential range of 0.01 to 1 V vs. Li/Li+, with Battery Charge /
Discharge System (HJ1010mSM8, Hokuto Denko, Japan).
Cyclic voltammetry was carried out with Automatic Polarization System,
HZ-3000, Hokuto Denko, Japan), in the potential range of open-circuit potential
(OCP) → 0.01 to 1 V vs. Li/Li+, at scan rate of 0.2 mV/sec.

Page 107
Chapter 4
99
4.3. Results and Discussion
4.3.1. Confirmation of mesoporous structure
Morphological and structural differences of Sn anodes electrodeposited with
or without nonionic surfactant were investigated with TEM. Figure 4.3.1 indicates
the bright-field TEM images and selected ED patterns in a 100 nm region of
as-prepared mesoporous Sn and dense Sn. Figure 4.3.1 (a) shows the TEM image
of the as-prepared sample with nonionic surfactant and its selected-electron
diffraction (ED) pattern in a 100 nm region. Though the nonionic surfactant
template was ordered with 2D-hexagonal symmetry of a d spacing of 5.8 and 2.8 nm,
confirmed by low angle XRD analysis (Fig. 4.2.1), a disordered mesoporous
structure was observable over the entire area (Fig. 4.3.1 (a)). This reason should
be contributed by the dendritic Sn deposition; this dendritic deposition of Sn should
dislodge the nonionic surfactant template. The bright and dark areas indicate
pores and deposited Sn, respectively. The ED patterns are assignable to (111)
diffraction of the α-Sn structure and (200), (101), (220), and (211) diffractions of the
β-Sn structure. The ring pattern indicates that the deposited Sn is composed of
very minute polycrystals. On the other hand, Figure 4.3.1 (b) shows the TEM
image of the as-prepared sample without nonionic surfactant and its selected-ED
pattern in a 100 nm region. A mesoporous structure was not observed on the
as-prepared sample without nonionic surfactant. The ED patterns are assignable to
(111), (220), and (311) diffraction of the α-Sn structure.
Thus, the introduction of mesoporous structure to Sn anode with nonionic
surfactant, and the crystallinity of both samples were confirmed.

Page 108
Chapter 4
100
Figure 4.3.1 Bright-field TEM images and selected ED patterns in a 100 nm region.
(a) As-prepared mesoporous Sn. (b) As-prepared dense Sn. The sample were
scratched from the substrate. The powdery samples were dispersed in ethanol and
mounted on a TEM microgrid.
20 nm
(a)
β-Sn (220)
β-Sn (101)
α-Sn (111)
β-Sn (211)
β-Sn (200)
20 nm
(b)
α-Sn (311)
α-Sn (220)
α-Sn (111)

Page 109
Chapter 4
101
4.3.2. Electrochemical properties of mesoporous Sn anode
To evaluate the influence of the introduction of mesoporous structure to Sn
anode, the charge-discharge tests were carried out. Figure 4.3.3 and Figure 4.3.4
show charge-discharge curves of the mesoporous Sn anode, and dense Sn anode is
also shown for comparison, at initial stage (1-ca.5th cycle) and second stage or later
(ca. 5-100th cycle), respectively.
In Fig. 4.3.3, the plateaus at 0.58, 0.72, and 0.78 V vs. Li/Li+ on discharge
curves; assigned to the reactions of Li7/3Sn → LiSn, LiSn → Li0.4Sn, and Li0.4Sn →
Sn, respectively(33), were observed. The potentials of these plateaus of mesoporous
Sn anode at discharging were ca. 20 mV lower than that of dense Sn. This
difference should be contributed by the reduction of iR drop due to the higher
surface area of mesoporous Sn anode compared with dense anode. This discussion is
also supported by the cyclic voltammogram of both anodes, shown in Fig. 4.3.5. The
current density of the mesoporous Sn anode was ca. 3 times as high as that of the
dense Sn anode; this indicates that the active site of the mesoporous Sn with Li was
larger than that of the dense Sn.
Figure 4.3.6 shows the discharge capacity of the mesoporous Sn anode and
dense Sn anode. In Fig. 4.3.6, the discharge capacities of both the mesoporous and
dense Sn anodes, increased at the initial stage (0-5th cycle). This reason is
described later. At the second stage (ca. 5-20th cycle), the discharge capacity of the
dense Sn anode was dramatically decreased due to electrical disconnection. The
disintegration of the dense anode from the Cu foil was visually observed. On the
other hand, in the case of the mesoporous Sn anode, such a phenomenon was not
confirmed. The mesoporous structure suppresses the influence of the volume
change during the cycling to some extent, which can prevent the disintegration of
the Sn anode from the substrate. This reason is also described later. At the
following stage (over 20th cycle), the discharge capacities of both Sn anodes were
maintained with small changes. After the 100th cycle, the discharge capacity of
the mesoporous anode was 425 mAh/g, while that of the dense anode was only 46
mAh/g. For comparison, when nanosized carbon particles were cycled at a
constant current of 100 mA/g, the discharge capacity after 100th cycle was 290
mAh/g.(34)
Consequently, the introduction of the mesoporous structure is effective
to improve the cycle durability.

Page 110
Chapter 4
102
Figure 4.3.3 The charge-discharge curves of (a) mesoporous Sn, (b) dense Sn at
initial stage. The charge-discharge current density was 0.994 A/g (1 C rate) in the
potential range of 0.01 to 1 V vs. Li/Li+.
0.0
0.2
0.4
0.6
0.8
1.0
1.2
-1000 -800
-600
-400
-200
0
200
400
600
800
Capacity / mAh g
-1
P
otential / V
vs. L
i/L
i+
0
0.2
0.4
0.6
0.8
1
1.2
-1000 -800 -600 -400 -200
0
200
400
600
800
Capacity / mAh g-1
P
o
tential / V
vs. L
i/L
i+
(a)
(b)
1st charge
1st discharge
5th charge
3rd discharge
3rd charge
5th discharge
1st charge
1st discharge
6th charge
3rd discharge
3rd charge
6th discharge

Page 111
Chapter 4
103
Figure 4.3.4 The charge-discharge curves of (a) mesoporous Sn, (b) dense Sn at
second stage or later. The charge-discharge current density was 0.994 A/g (1C
rate) in the potential range of 0.01 to 1 V vs. Li/Li+.
0.0
0.2
0.4
0.6
0.8
1.0
1.2
-800 -600 -400 -200
0
200
400
600
800
Capacity / mAh g
-1
P
otential / V
vs. L
i/L
i+
100th 30th 20th 15th 5th
0
0.2
0.4
0.6
0.8
1
1.2
-800 -600 -400 -200
0
200
400
600
800
Capacity / mAh g
-1
P
otential / V
vs. L
i/L
i+
(a)
(b)
5th 15th 20th 30th 100th
100th 40th 20th 10th 6th
6th 10th 20th 40th 100th

Page 112
Chapter 4
104
Figure 4.3.5 The cyclic voltammogram of the mesoporous Sn (black line) and the
dense Sn anodes (gray line) at 1st cycle. The potential range was open-circuit
potential → 0.01 to 1 V vs. Li/Li+, at a scan rate of 0.2 mV/sec.
-1.0
-0.5
0.0
0.5
1.0
0
0.5
1
1.5
Potential / V vs. Li/Li+
C
urrent D
ensity / m
A
cm
-
2

Page 113
Chapter 4
105
Figure 4.3.6. The discharge capacities of the mesoporous Sn (●) and the dense Sn
anodes (○). The charge-discharge current density was 994 mA/g (1 C rate) in the
potential range of 0.01 to 1 V vs. Li/Li+.
0
200
400
600
800
0
20
40
60
80
100
Cycle number / -
D
is
c
h
a
rg
e
c
a
p
a
c
ity
/ m
A
h
g
-1

Page 114
Chapter 4
106
Figure 4.3.7 shows the discharge capacity at the 1st cycle with various C
rates. With increasing C rate, the discharge capacity decreased for both anodes.
At lower C rate, the difference between mesoporous and dense Sn anodes was
hardly observed. However at higher C rate, a remarkable difference was
observed; only the mesoporous Sn anode operated at as high as 17 C rate. And
the discharge capacity at the 1st cycle with 10 C rate was 352 mAh/g, which is
almost the theoretical capacity of carbon anodes. This result is ascribable to the
alleviation of voltage loss due to the moderation of actual current density and
should be ascribed to the shortening of the solid state diffusion distance of Li. The
alleviation of voltage loss was confirmed in Figure 4.3.8. The plateaus of the
reaction (LiSn → Li7/3Sn) on both mesoporous and dense Sn anode were observed
at ca. 0.40 V vs. Li/Li+ when they were charged with 1 C rate. However the plateau
on the dense Sn anode was drastically shifted to ca. 0.12 V vs. Li/Li+ when it was
charged with 7 C rate, though the plateau on the mesoporous Sn anode was shifted
to as small as 0.26 V vs. Li/Li+ when it was charged with 10 C rate. As the result,
rate property was improved by introducing mesoporous structure to the Sn anode.

Page 115
Chapter 4
107
Figure 4.3.7. The discharge capacity at the first cycle of the mesoporous Sn anode
(♦) and the dense Sn anode electrodeposited with a conventional bath (◊). The
charge-discharge current density was varied from 0.994 to 1.69 A/g (1 – 17 C rate) in
the potential range of 0.01 to 1 V vs. Li/Li+.
C rate / -
D
is
c
h
a
rg
e
c
a
p
a
c
ity
/ m
A
h
g
-1
0
200
400
600
800
0
5
10
15
20

Page 116
Chapter 4
108
Figure 4.3.8. The charge-discharge curves of (a) mesoporous Sn, (b) dense Sn at 1st
cycle. The charge-discharge current density was varied from 0.994 A/g (1 C rate) to
16.9 A/g (17 C rate) in the potential range of 0.01 to 1 V vs. Li/Li+.
0
0.2
0.4
0.6
0.8
1
1.2
-1000 -800 -600 -400 -200
0
200 400 600 800
Capacity / mAh g
-1
P
otential / V
vs. L
i/L
i+
1C
2C
3C
5C
10C
15C
17C
0
0.2
0.4
0.6
0.8
1
1.2
-1000 -800 -600 -400 -200
0
200 400 600 800
Capacity / mAh g
-1
P
otential / V
vs. L
i/L
i+
1C
2C
3C
7C
(a)
(b)

Page 117
Chapter 4
109
4.3.3. Morphology and Structure change with cycle
Figure 4.3.7 shows the surface morphology of the mesoporous and the dense
Sn anode, observed with SEM. The ditch aligned in constant direction was
observed on the both the mesoporous and the dense Sn anode as deposited. These
ditches were derived from the surface of Cu substrate. After a few cycles, both the
mesoporous and the dense Sn anode formed cracks like island due to the volume
change with cycles. These clacks contributed to the increase the electrochemically
active surface area, so that the discharge capacities were increased at fist stage,
described in Fig. 4.3.6.(35)
Figure 4.3.8 shows the wide-area surface morphology of
the mesoporous and the dense Sn anode after 3 cycles. The clacks of the
mesoporous Sn anode after 3 cycles were smaller than that of the dense Sn anode.
This smaller clack formation should contribute to the suppression of the stress in
the electrode. As the result, the discharge capacity degradation, which was
observed at second stage indicated in Fig. 4.3.6, by the disintegration of the anode
from the Cu substrate should be controlled on the mesoporous Sn anode.

Page 118
Chapter 4
110
Figure 4.3.7 SEM images of mesoporous Sn anode and dense Sn anode (as-deposited,
after 1st cycle, and after 3rd cycle at 1 C rate in the potential range of 0.01 to 1 V vs.
Li/Li+.).
Figure 4.3.8 Wide-area SEM images of (a) mesoporous Sn anode and (b) dense Sn
anode after 3rd cycle shown in Fig. 4.3.7.
As Deposited
After 3 cycling
Mesoporous
Sn anode
After 1 cycling
12μm
12μm
12μm
12μm
12μm
12μm
Dense
Sn anode
12μm
12μm
12μm
12μm
12μm
12μm
80 μm
80 μm
(a)
(b)

Page 119
Chapter 4
111
Figure 4.3.9 shows the TEM image of the mesoporous and dense Sn anode
after the 100 cycles and its selected-ED pattern. From Fig. 4.3.9 (a), the formation
of the Sn grains of 10 nm in size was observable, indicated by the arrow (Fig. 4.3.9
(a)). The ring pattern including the intense spots (Fig. 4.3.9 (a)) shows the
presence of both very minute polycrystalline phase and crystal Sn grains. From
Fig. 4.3.9 (b), the formation of the Sn grains was not observed in this area, however
the formation of the Sn grains of 20-50 nm in size was observed from the dense Sn
after the 30 cycles, hence the Sn grains could be formed in the other area. There
was the report that an amorphous-like phase should be effective to improve the
cycle durability of Sn anode.(36)
Therefore, the very minute polycrystalline phase
should contribute to the improvement of the cycle durability of the mesoporous Sn
anode. When the dense Sn anode was used, larger Sn crystals grow easily, after
the charge-discharges of Li ions are repeated for several times.(35)
On the contrary,
the mesoporous Sn anode still retains its very minute polycrystalline phase even
after the 100th cycle. Therefore, the presence of the mesoporosity should prevent
the larger Sn crystal growth during cycling, leading to a longer life.

Page 120
Chapter 4
112
Figure 4.3.9 Bright-field TEM images and selected ED patterns in a 100 nm region.
(a) the mesoporous Sn after 100 cycles. (b) the dense Sn after 100 cycles. The
samples were scratched from the substrate. The powdery samples were dispersed in
ethanol and mounted on a TEM microgrid.
20 nm
20 nm
(a)
(b)

Page 121
Chapter 4
113
4.4. Summary
In summary, a mesoporous Sn anode was prepared from LLC bath by
electrodeposition. The cycle durability and charge-discharge property with high
current density were improved. The discharge capacity of the mesoporous anode
was 425 mAh/g after the 100th cycle, while that of the dense anode was only 46
mAh/g. The mesoporous structure suppressed the influence of the volume change
during the charge-discharge cycling of Li-ions, and provided the high cycle
durability to the Sn anode. The high cycle durability should be affected by a very
minute polycrystalline phase. Moreover, the improvement of the charge-discharge
property with high current density at the first cycle was demonstrated by the
introduction of the mesoporous structure. Only the mesoporous Sn anode operated
at as high as 17 C rate; the discharge capacity at the 1st cycle with 10 C rate was
352 mAh/g, which is almost the theoretical capacity of carbon anodes. This result
is ascribable to the alleviation of voltage loss due to the moderation of actual
current density and should be ascribed to the shortening of the solid state diffusion
distance of Li. Thus, it was demonstrated that the introduction of mesoporous
structure was effective to improve the cycle and rate properties of the Sn anodes.

Page 122
Chapter 4
114
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Page 124

Page 125
Chapter 5:
General Conclusion

Page 126

Page 127
Chapter 5
119
The aim of this dissertation is to investigate the potentials of the
introduction of phase separated electrolytes to preparation process of materials for
lithium secondary batteries and suggest the availability of phase separated
electrolytes for improvement of the materials.
In this chapter, the results obtained from the research described in Chapter
2, 3, and 4 were summarized and a possibility of phase separated electrolytes for a
future work.
I. The polymer blend gel electrolyte composed of PEO-LiBF4 complex and PS.
(Chapter 2)
The PEO-LiBF4/PS polymer blend gel electrolyte indicated the ionic conductivity
of 0.9 x 10-3 S/cm as the maximum value by optimizing the Li+ concentration in the
plasticizer, even though it possessed mechanically strong enough to handle, because
it had co-continuous three-dimensionally interpenetrated polymer network (IPN)
structure. In addition, the chemical stability of the interface between the IPN gel
electrolyte and lithium metal was improved compared with the PEO/PS polymer
blend gel electrolyte, because of the retentivity improvement of the ionic conductive
phase due to PEO-LiBF4 complex. The stability against the oxidation was enough
to be applied to LiCoO2
cathode which is 4 V class cathodes. With the
charge-discharge rate of C/5, the capacity per gram of cathode was found to be over
130 mAh/g over 30 cycles.
II. The di-block copolymer gel electrolyte composed of PEO and PS. (Chapter 3)
The ionic conductivity of the PEO-PS di-block copolymer gel electrolyte
demonstrated to be 2.65 mS/cm at maximum value, even though it was
mechanically strong enough to suppress the dendritic deposition of lithium at the
interface between Li electrode and the gel electrolyte. The excellent chemical
stability of the PEO-PS di-block copolymer gel electrolyte with Li electrode was
revealed. This result was led by fixation of the PEO with PS matrix by covalent
bond. The stability against the oxidation was found to be enough for application of
LiFePO4 cathode. With the charge-discharge rate of C/5, the capacity per gram of
cathode was found over 120 mAh/g over 30 cycles.

Page 128
Chapter 5
120
III. Mesoporous Sn anode electrodeposited with a lyotropic liquid crystal bath.
(Chapter 4)
The significant improvements of Sn anode on the cycle durability and
charge-discharge property with high current density were discovered by
introduction of mesoporous structure. The discharge capacity of the mesoporous
Sn anode was 425 mAh/g after the 100th cycle, while the anode without mesoporous
structure was only 46 mAh/g. The mesoporous structure suppressed the influence
of the volume change during the charge and discharge cycling of Li-ions, and the
high cycle durability was realized. The high cycle durability should be affected by
a very minute polycrystalline phase. Moreover, the improvement of the
charge-discharge property with high current density at the first cycle was
demonstrated by the introduction of the mesoporous structure. Only the
mesoporous tin anode operated at as high as 17 C rate; the discharge capacity at the
1st cycle with 10 C rate was 352 mAh/g, which is almost the theoretical capacity of
carbon anodes. This result is ascribable to the alleviation of voltage loss due to the
moderation of actual current density and should be ascribed to the shortening of the
solid state diffusion distance of Li.

Page 129
Chapter 5
121
These materials indicated potential performances as described above,
moreover the benefits are expected as follows.
The polymer blend gel electrolyte can be prepared from PEO and PS, which
have been widely used industrially, by simple hot-blending method without volatile
organic solvents, which are used for most preparation of polymer electrolytes.
The interfacial stability between the di-block copolymer gel electrolyte and
Li metal anode was established by its covalent bond, enable to suppress the
performance degradation at long storage. And the large film of block copolymer has
come to be prepared recently.
The mesoporous Sn anode which has high performances can be prepared by
electrodeposition, hence it can be prepared on large substrate. It is beneficial for
mass production.
The author believes that these results demonstrate the potential of phase
separated electrolytes, and has visions that anode and electrolyte would be
prepared all together three-dimensionally; the surfactant itself act as polymer
electrolyte by controlling the surfactant which is used as template at
electrodeposition of anode in future. If this process is brought to realization, some
benefit can be received; the degradation of Sn anode due to its volume change at
charge-discharge cycles can be suppressed by polymer matrix, and the organic
solvent for removal of surfactant as template becomes to unnecessary.

Page 130

Page 131
Appendix
“Numerical Simulation of DMFC-Capacitor Hybrid Power Supply
System for Small Electric Devices”

Page 132

Page 133
Appendix
125
1. Introduction
In recent years, mobile electronic devices including cellular phone, laptop computer,
and PDA, which are generally equipped with Li ion batteries, require large power output of the
electric power supply used to operate those devices. Besides the development of high energy
density, rechargeable Li ion batteries and small fuel cells are proposed as the electric power
supply for small electric devices of the future. Fuel cells attract many researchers because of
conceivable advantages, such as an enhanced energy density and a shorter period of time
required for recharging the energy consumed. Fuel cells, especially direct liquid-fuel fuel cells
such as direct methanol fuel cell (DMFC), are being developed to fulfill the need for high energy
density power supplies [1,2]. The direct liquid-fuel fuel cells have a high energy density, and
much efforts are being expended to improve their output power density to reduce the
production cost and the size of the fuel cells.
In designing a power supply for electric devices, both maximum and mean power
requirements should be taken into account together with the length of time during which a
large power consumption occurs. To satisfy the requirement of the electric devices, a power
supply consisting of fuel cells only should be designed and fabricated to provide maximum
power output. However, hybrid power supply systems, such as DMFC connected with another
power supply, capacitors [1, 4, 5] or secondary batteries [6-10], level the requirement of DMFC
output into the mean value in a short period of time by supplying the pulsed output from
another power supply. Especially capacitors are advantageous for use with a pulsed load,
because of its character, short charge-discharge time and extensive cycle durability.
The simulation of hybrid power supply systems is important, because, for constructing
such a system, it is useful to know the required power output of DMFC and capacitor for the
specific electric device without performing actual fabrication. The DMFC and capacitor can
then be designed for the required power output which has been calculated by the simulation.
In the present work, an internal resistance of DMFC was approximated by a simple
resistance based on a current-voltage profile of a DMFC. The simple numerical simulation of a
hybrid power supply system consisting of a DMFC and a capacitor connected in parallel was
attempted and the obtained results were discussed with the current profile of hybrid power
supply set up. In addition, the possibility for constructing a smaller hybrid power supply
system was confirmed for the system consisting of our DMFC and micro electrochemical
capacitor (MECC).

Page 134
Appendix
126
2. Modeling
As electric devices used for the modeling study, two Japanese cellular phones (2G;
KDDI au, 2004 product, and 3G; NTT Docomo, 2006 product) were chosen to obtain power
consumption profiles. The power consumption profiles, when the constant voltage of 3.7 V was
applied to the cellular phone with removal of lithium battery pack, was determined by recorded
from the current flow to the cellular phone from the DC power supply (potentiostat /
galvanostat, HA-301, Hokuto Denko Corp.), every 20 and 50 msec with a voltage recorder
(Voltage Data Logger, VR-71, T&D corporation) as shown in Fig. 1 (a) and (b), respectively. The
power consumption profile shown in Fig. 1 (c) is described in simulation and discussion section.
The current profiles were assumed as merely a pulsed load, though a capacitor might have
been built in the 3G cellular phone.

Page 135
Appendix
127
Fig. 1. Power consumption profiles of a cellular phone. (a) when calling with 2G, recorded every
20 msec, (b) when watching TV with 3G, recorded every 50 msec, and (c) when watching TV
and changing the channel with 3G, recorded every 20 msec. Note the power consumption
profile of (c) was derived from the sum of the current flow passed through the DC power supply
and capacitor connected in parallel.
0
0.2
0.4
0.6
0
20
40
60
80
100
Time / sec
I lo
ad
/ A
0
0.2
0.4
0.6
0
100
200
300
400
Time / sec
I lo
ad
/ A
0
0.2
0.4
0.6
0
10
20
30
40
Time / sec
I lo
ad
/ A
(a)
(b)
(c)

Page 136
Appendix
128
A parallel connection of a DMFC and a capacitor was considered in formulating the
model of hybrid power supply in this work. The capacitor in this system is required to supply a
pulsed high power when demanded by the electric device. The DMFC supplies the power
leveled by the capacitor to the electric device and charges the capacitor.
In this study, the DMFC was fabricated in a same way presented by Shimizu et al. [3],
to obtain the current-voltage and current-power curves in its operation as shown in Fig. 2. The
operation range of the power supply for cellular phone was considered in the current range
lower than the value at which the maximum power was obtained. As many factors:
electrochemical reaction, methanol crossover, and water management, have to be considered
for a DMFC performance, the current-voltage data actually measured as shown in Fig. 2 should
be containing all factors affecting its performance; while the current-voltage profile shown in
Fig. 2 can be ascribed linearly aligning as long as in the current range of 0.025 to 0.7 A where
the DMFC will be operated. Accordingly, it would be possible to simulate with a simple
resistance which approximated an internal impedance of a DMFC. While the electric double
layer of the electrode in the DMFC is considered to be take into account, in this simulation the
electric double layer capacitance is omitted because of it’s negligibly small charge compared
with the generated current during the fuel cell operation; actually the simulation results
indicated in Fig. 4 showed good agreement with the current profile in the real world. Then,
with the assumptions listed below, an equivalent circuit of the hybrid type power supply
system shown in Fig. 3 was constructed.
Fig. 2. Polarization (■) and power density (□) curves. of DMFC (25 cm2 area).Catalysts:
unsupported Pt–Ru anode (6.4 mg cm−2 of Pt–Ru) and unsupported Pt cathode (3.9 mg cm−2 of
Pt), current collector SUS/Au, 2 M methanol solution, air passive, room temperature, and
Nafion 117.
0
0.1
0.2
0.3
0.4
0.5
0.6
0
0.2
0.4
0.6
0.8
1
IFC / A
V
F
C
/ V
0
0.1
0.2
0.3
P
o
w
er / W
Voltage
Power

Page 137
Appendix
129
1.
The internal impedance of the fuel cell, which is a combination of ohmic resistance,
charge transfer reaction resistance, mass transfer resistance and double layer
capacitance, is represented as a simple ideal resistor in the operation range. DMFC
generates the power output with a voltage which is linearly related to the output
current with the ideal resistance described above. The virtual open circuit voltage is
the value of intersection found by extrapolating the current-voltage plots of the
operation current range to the voltage axis. The value of intersection is lower than the
real open circuit voltage of DMFC by the voltage drop caused by the charge-transfer
resistance.
2.
Capacitor is described as a serial connection of ideal capacitance and ideal resistance,
that is equivalent series resistance (ESR), which includes resistances of electrolyte,
electrodes, and lead wire in a capacitor.
The parameters in Fig. 3 are shown in Table 1. The resistance and the virtual open circuit
voltage of DMFC were calculated from the current-voltage profile of DMFC shown in Fig. 2;
this graph yields the resistance value of 0.27 Ω, and the virtual open circuit voltage of 0.485 V.
As the maximum output of the DMFC was about 0.2 W shown in Fig. 2, the DMFC stack of 10
DMFCs connected in series was designed in this study to modeling for operation of cellular
phone with its consumption of 2 W estimated from Fig. 1. RFC and VFC_ocv were then decided
from the resistance value and the virtual open circuit voltage for 10 DMFCs connected in series
for the operation of the cellular phone. The values of Rcap and C were measured by the a.c.
impedance method. For the equivalent circuit, the following numerical expressions were used
with the assumption that these values are constant during a length of time (the finite element
method). The parameters in below expressions are described in Table. 1.

Page 138
Appendix
130
Fig. 3. Equivalent circuit of hybrid power supply system. The parameters are listed in Table. 1.
Table 1. Parameters in the equivalent circuit of hybrid power supply system.
VFC_ocv
Virtual open circuit voltage
V
Voltage of hybrid power supply system
Iload
Current flow through the electric devices
IFC
Current flow through DMFC
Icap
Current flow through capacitor
RFC
Inner resistance of DMFC
Rcap
Inner resistance of capacitor, ESR
Q
Quantity of charged electricity
C
Capacitance
V
Load
Iload
C
VFC_ocv
FC
Rcap
IFC
Icap
C
R
I

Page 139
Appendix
131
0
)(
)(
)(
=
+
+
m
m
m
t
load
t
cap
t
FC
I
I
I
(1)
FC
t
FC
t
ocv
FC
R
I
V
V
m
m
×
=
-
)(
_
(2)
cap
t
cap
t
t
R
I
VCQ
m
m
m
×
=
-
)(
(3)
cap
I
dt
dQ
-
=
(4)
)(
_
m
m
m
t
load
cap
FC
cap
FC
ocv
FC
cap
FC
cap
t
cap
FC
FC
t
I
R
R
R
R
V
R
R
R
C
Q
R
R
R
V
×
+
×
+
×
+
+
×
+
=
(5)
(1), (2), (3))
)
(
1
)
(
1
1
-
-
×
-
=
-
-
m
m
t
cap
t
t
tt
I
Q
Q
m
m
m
(6)
(4))
where tm indicates time.
With this model, the numerical simulation of current flow was performed with Microsoft Excel
by assigning the power consumption profiles as Iload (tm) shown in Fig. 1. As the results, V(tm),
IFC (tm), Icap (tm) and Q(tm) were calculated. The resolution time of the simulation was used the
same value at recording of the power consumption profiles. At the starting point the capacitor
was assumed to be made at the fully charged state by the fuel cell after a period of no
operation.

Page 140
Appendix
132
3. Simulation and Discussion
To confirm the appropriateness of the proposed numerical simulation, the magnitudes
of currents actually measured were compared with currents calculated. For actual
measurement, a hybrid power supply was constructed with parallel connection of the DC power
supply with a series resistor, which represents the internal resistance of DMFC, and two
laminate type series capacitors, which are EDLCs, for achieving durability at cellular phone
operation voltages. The power consumption profile shown in Fig.1 (c) was determined from the
sum of the current flow passed through the DC power supply and electric double layer
capacitor (EDLC) connected in parallel, which are recorded every 20 msec with ampere meter
(zero shunt ammeter, HM-104, Hokuto Denko Corp.) and the voltage recorder; the parallel
connection of the DC power supply and EDLC were assumed to be hybrid power supply.
Though the power consumption profile shown in Fig.1 (c) was measured in another way, it can
be dealt as the power consumption profile shown in Fig. 1 (a), (b). The actual currents flowing
through the DC power supply and the capacitor to the cellular phone were shown by gray lines
in Fig. 4. The simulated currents based on the load current shown in Fig.1 (c) were shown by
black lines in Fig. 4. The simulated currents were calculated by using the values of Rcap, C, RFC,
and VFC_ocv. RFC is the value of the resistor connected with the DC power supply, and VFC_ocv is
the value applied with the DC power supply. The current profiles, i.e., values of measured IFC,
Icap and simulated IFC, Icap, are compared in Fig. 4. The small difference between measured and
simulated profiles may be attributable to the resistance of the circuit. This result demonstrates
that our numerical simulation method is applicable to the case using the hybrid power supply.

Page 141
Appendix
133
Fig. 4. Profiles of current flowing through DMFC and capacitor. Simulated current profiles
(black line) and measured current profiles (gray line).
-0.2
0
0.2
0.4
0.6
0
20
40
60
80
100
Time / sec
C
urrent / A
Current flowing through capacitor
Current flowing through DMFC

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134
Figure 5 shows examples of simulated results for load current with two laminate type
capacitors in series, 1.92 F and 0.42 Ω, and with a virtual capacitor, with 1.92 F and 4.2 Ω, the latter
being 10 times as larger as the laminate type capacitor resistance, that is ESR. The load leveling rate
was calculated using the formula defined as follows:
100
[%] rate
leveling
Load
max
_
max
_
max
_
×
-
=
load
FC
load
I
I
I
(7)
where Iload_max, and IFC_max are maximum load current and current flow through the DMFC in the
range of measuring time, respectively. The load leveling rate were calculated with various load
currents and capacitances. The load leveling rate of (a), (b), (c), (d), (e) and (f) in Fig. 5 were 62.7,
27.3, 22.1, 38.9, 14.7, and 12.9 % respectively. These results indicate the importance of ESR,
implying that the higher ESR disables the capacitor in the hybrid power supply system for charging
and discharging. Therefore, the influence of capacitor with several values of capacitance and ESR
to the load leveling rate was simulated. Figure 6, (a), (b), (c) show the load leveling rate for the load
indicated in Fig. 1, (a), (b), (c), respectively. These results indicated that the load leveling rate is
converged with increasing ESR, even if the capacitor in the hybrid power supply system has a high
capacitance. Consequently, it was demonstrated that ESR is a very important factor to consider,
when a hybrid power supply system is constructed. In addition, this load leveling effect suggests the
possibility of downsizing the surface area of the DMFC electrode. In the case of (a) in Fig. 5, only
ca. 40 % of the DMFC electrode area is necessary for the cellular phone operation, compared with
the area required without the laminate type capacitors.

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135
Fig. 5. Profiles of current flowing through DMFC (black line) and capacitor (gray line) with
different capacitances; (a) - (c) 1.92F, 0.42 Ω, and (d) - (f) 1.92F, 4.2 Ω, (a) and (d): when making
a call with 2G; (b) and (e): when watching TV with 3G; (c) and (f) when watching TV and
change channels with 3G.
Fig. 6. Load leveling rate vs. ESR for capacitors with various capacitances, calculated on the
load current profiles; (a) when making a call with 2G, (b) when watching TV with 3G, (c) when
watching TV and change channels with 3G.
(a)
(d)
(b)
(e)
(c)
(f)
-0.2
0
0.2
0.4
0.6
C
u
rre
n
t / A
0
20
40
60
80
100
-0.2
0
0.2
0.4
0.6
0
10
20
30
40
Time / sec
C
u
rre
n
t / A
0
100
200
300
400
Time / sec
0
20
40
60
80
100
Time / sec
(a)
(b)
(c)

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Appendix
136
The possibility of constructing a smaller hybrid power supply system was also examined
for our μDMFC [11-14] and micro electrochemical capacitor (MECC) [15, 16] for the 2G cellular
phone. The parameters of μDMFC were determined from the current-voltage curve [17]. The
resistance and the value of virtual open circuit voltage were 361 Ω and 0.333 V respectively. Values
of RFC and VFC_ocv were derived from the resistance and the virtual open circuit voltage of 15
μDMFCs, note that μDMFCs connected in series is called a “μDMFC-stack” in this paper,
assembled in series for the cellular phone application. To obtain a sufficiently large current, these
components of the stacks were assumed to be connected in parallel. The parameters of MECC used
were that announced by Takasu et al. [15], and MECCs were also assumed to be connected in
parallel. In this case, 1000 μDMFC-stacks and 3000 MECC units were connected in parallel. Fig. 7
shows voltage and current profiles for μDMFC and MECC. From Fig. 7 (a), (b), the minimum
voltage is 3.37 V, which is large enough to operate the cellular phone, in the range of measurement
time. The load leveling rate is 48.5 %, which indicates the MECCs functioned effectively. Here, the
volume of each cell was 1 cm
3
. Therefore the volume of this hybrid power supply system was
estimated to be 18000 cm
3
, which is smaller than 25000 cm
3
, the volume of the power supply
system without MECC for operating the cellular phone.

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137
Fig. 7. Voltage (a), (c) and current (b), (d) profiles of the hybrid power supply system consisting
of μDMFC and MECC. The profiles of (c), (d) refer to the case of the voltage raised with a
DC-DC converter (efficiency; 80 %). Profiles of current flowing through μDMFC and MECC
indicated by black line and gray line.
0
2
4
6
V
o
ltag
e / V
0
4
8
12
0
10
20
30
40
Time / sec
V
o
ltag
e / V
(a)
(b)
(c)
(d)
-0.2
0
0.2
0.4
0
10
20
30
40
Time / sec
C
u
rren
t / A
-0.2
0
0.2
0.4
C
u
rren
t / A

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Appendix
138
In addition, a DC-DC converter that is charge pump assumed to be adapted to this system,
which enables to boost the voltage of the power supply, with reducing the current flowing, so that
the influence of the IR drop can be alleviated. The calculation with a DC-DC converter was simply
conducted; when it raises the voltage of the system, the current of the system is reduced
commensurately and is reduced in consideration of the loss by its efficiency. Figure 7 (c), (d) shows
the voltage and current profiles in case that a DC-DC converter (efficiency was assumed to be
80 %) which decuple the voltage of the system was adapted to this system which is composed of
4500 μDMFC-stacks, each consisting of 3 μDMFCs in series, and 1000 MECC units were
connected in parallel. The total volume of this hybrid power supply system was estimated to be
14500 cm
3
. Thus the DC-DC converter enables to decrease the volume by about 20 %. In spite of
the decrease in the volume, the load leveling rate is as high as 59.9 %. Consequently, it was
suggested that with the use of a DC-DC converter, further downsizing was possible. The μDMFC
and MECC used in this simulation were so of prototype that the electrode areas of μDMFC and
MECC were considerably smaller against the cell size. Therefore, those electrode areas would be
able to be expanded by a factor of 50 by the accumulation of the electrodes. When the electrodes of
μDMFC and MECC are accumulated, the volume of the hybrid power supply system would got
smaller by a factor of 50, i.e., to 290 cm
3
. Moreover, downsizing of μDMFC and MECC should be
achievable by increasing the effective electrode area, improving electrode performances, and
suppressing methanol crossover. Improvement of these factors may enable our hybrid power supply
system to become comparable to the contemporary standard system developed by the industry.

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Appendix
139
4. Conclusion
Numerical simulation of a hybrid power supply system consisting of a DMFC and a capacitor
connected in parallel was simply investigated based on a known current-voltage data. The
possibility of designing a smaller hybrid power supply system was demonstrated for use with our
μDMFC and MECC. The simulated profiles of current flowing through the DMFC and the
capacitor demonstrated an excellent agreement with experimental profile. The ESR was revealed to
be a very important factor in the hybrid power supply system of DMFC and capacitor. The
numerical simulation helps to find improvements required of hybrid power supply system for
practical applications.

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Appendix
140
Reference
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Nakamura, and Y. Kubo, Electrochemistry, 70, 966 (2002).
2) T. Shimizu, M. Mohamedi, T. Momma, and T. Osaka, Electrochemistry, 72, 637 (2004).
3) T. Shimizu, T. Momma, M. Mohamedi, T. Osaka and S. Sarangapani, J. Power Sources, 137,
277 (2004).
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11) S. Motokawa, M. Mohamedi, T. Momma, S. Shoji, and T. Osaka, Electrochemistry
communications, 6, 562 (2004).
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Electrochemistry, 5, 352 (2005).
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(2005).
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JAPANESE JOURNAL OF APPLIED PHYSICS PART 1-REGULAR PAPERS BRIEF
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Meeting of The Electrochemical Society of Japan, Abstract #1L28 (2006).
16) W. Sugimoto, K. Ohuchi, K. Yokoshima, J. Park, T. Momma, T. Osaka, and Y. Takasu, 208th
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17) S. Tominaka, H. Obata, T. Momma, J. Park, and T. Osaka, EMNT2006 (2006).

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List of Achievement
141
List of Achievements
1. Original Articles
“Characteristics of Interpenetrated Polymer Network System made of
Polyethyleneoxide -LiBF4 Complex and Polystyrene as the Electrolyte for Lithium
Secondary Battery”
Toshiyuki Momma, Hiroaki Ito, Hiroki Nara, Hitomi Mukaibo, Stefano Passerini,
Tetsuya Osaka
Electrochemistry, 71, 1182 (2003).
“Numerical Simulation of DMFC-Capacitor Hybrid Power Supply System for Small
Electric Devices”
Toshiyuki Momma, Hiroki Nara, Tetsuya Osaka
Electrochemistry (Accepted)
“Feasibility of an Interpenetrated Polymer Network System Made of Diblock
Copolymer Composed of Polyethylene oxide and Polystyrene as the Gel Electrolyte
for Lithium Secondary Batteries”
Hiroki Nara, Toshiyuki Momma, Tetsuya Osaka
Electrochemistry (Accepted)
“Cycle and Rate Property of Mesoporous Tin Anode for Lithium Ion Secondary
Batteries”
Hiroki Nara, Yoshiki Fukuhara, Hitomi Mukaibo, Yusuke Yamauchi, Toshiyuki
Momma, Kazuyuki Kuroda, and Tetsuya Osaka.
Chemistry Letters, 37, 142 (2008).
“Mechanical Analysis and In-situ Structural and Morphological Evaluation of Ni-Sn
Alloy Anodes for Li Ion Batteries”
J. Chen, S. J. Bull, S. Roy, H. Mukaibo, H. Nara, T. Momma and T. Osaka, Y.
Shacham-Diamand
J. Phys. D: Appl. Phys., 41, (2008) 025302 (13pp).

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List of Achievement
142
2. Oral Presentations
A.International Conferences (Presented by H. Nara, et al.)
“Cycle Property of Mesoporous Sn Anode for Lithium Ion Secondary Batteries”
212th Meeting of The Electrochemical Society, Washington, DC, USA, October
2007.
“Application of Diblockpolymer Gel Electrolyte Having Micro Phase Separation
Structure to Lithium Secondary Batteries”
The 56th Annual Meeting of International Society of Electrochemistry, Busan,
Korea, September 2005.
“Gel Electrolyte Having Micro Phase Separation Structure for Lithium Secondary
Batteries”
2nd International Conference on Polymer Batteries and Fuel Cells, Las Vegas,
USA, June 2005.
“Preparation of Gel Electrolyte Using Self-assembling Diblockpolymer for Lithium
Secondary Battery”
The Electrochemical Society (ECS) 2004 Joint International Meeting, Honolulu,
Hawaii. October 2004.

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143
B.Domestic Conferences (Presented by H. Nara, et al.)
“Effects of Dissolved O2 on Charge-Discharge Performance of Lithium Metal
Electrode in Organic Electrolytes with TFSI Anion Based Ionic Liquids as an
Additive”
The 48th Battery Symposium in Japan, Fukuoka, Japan, November 2007.
“Improvement of NiSn Alloy Powder for Lithium Ion Secondary Batteries by
Surface Oxide Control”
The 48th Battery Symposium in Japan, Fukuoka, Japan, November 2007.
“Cycle Performance of Sn Anode with Mesoporous Structure for Lithium-ion
Secondary Batteries”
The Electrochemical Society of Japan, Tokyo, Japan, September 2007.
“Improved Mechanical Property of Diblock-polymer Gel Electrolyte for Lithium
Secondary Battery and Battery Performance”
The chemical society of Japan, Osaka, Japan, March 2007.
“Voltage Loss Profile within the Electrolyte Membrane Caused by Proton Transfer
for Numerical Simulation of DMFC-Capacitor Hybrid Power Supply”
The 47th Battery Symposium in Japan, Tokyo, Japan, November 2006.
“Load Leveling Simulation of Direct Methanol Fuel Cell-Capacitor Hybrid System”
The Electrochemical Society of Japan, Tokyo, Japan, April 2006.
“Application of Diblockpolymer Gel Electrolyte to Lithium Secondary Batteries”
The Electrochemical Society of Japan, Kumamoto, Japan, April 2005.
“Preparation of Gel Electrolyte Using Self-assembling of Diblockpolymer for
Lithium Secondary Battery”
The chemical society of Japan, Hyogo, Japan, March 2004.

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List of Achievement
144
“Characterization of the Polymer Blend with PEO-LiX Complex as the Electrolyte
for Lithium Secondary Battery”
The Electrochemical Society of Japan, Tokyo, Japan, April 2003.
“Polymer Blend Gel Electrolytes by Complex of PEO-LiX and PS”
The Electrochemical Society of Japan, Kanagawa, Japan, September 2002.

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List of Achievement
145
3. Patent
Application No. 2007-293318
“DENKAIEKI, RICHIUMUNIJIDENCHI, OYOBI RICHIUMU NIJIDENCHI NO
SEIZOHOHO”
Tetsuya Osaka, Toshiyuki Momma, Chika Tatsumi, Hiroki Nara, Kunihiko Kojima,
Kentaro Tada, Satoshi Obara, Hitoshi Iwaki
Waseda University, Toyo Gosei Co., Ltd, November 2007.
Application No. 2007-119848
“HUKUGODENGEN”
Tetsuya Osaka, Toshiyuki Momma, Hiroki Nara, Masahiko Usuda, Masataka
Takeuchi
Waseda University, Showa Denko K.K, April 2007.

Page 154

Page 155
147
Acknowledgement
I would like to express my sincerely gratitude to Professor Dr. Tetsuya
Osaka for his supervision and continuous encouragement. I wish to express my
hearty gratitude to Professor Dr. Takayuki Homma for his advice and
encouragement. I would express my sincere appreciation to Professor Dr.
Kazuyuki Kuroda, chapter 4 in my thesis would not be achieved without his
cooperation. I would appreciate Professor Dr. Bruno Scrosati, University "La
Sapienza" Rome for valuable discussion and encouragement. I would extend my
sincere appreciation to Professor Dr. Yoshiyuki Sugawara, Professor Dr. Hiroyuki
Nishide, and Professor Dr. Yosi Shacham-Diamand for their valuable suggestions
and encouragement. And I would like to thank Associate Professor Dr. Toshiyuki
Momma for his valuable suggestions and discussion, my thesis would not be
reached completion without his advice. I would like to express my appreciation to
Professor Dr. Toru Asahi, Professor Dr. Kiyoshi Kanamura, Tokyo Metropolitan
University, Professor Dr. Masayoshi Watanabe, Yokohama National University,
Professor Dr. Hideki Masuda, Tokyo Metropolitan University, Professor Dr. N.
Furukawa, Dr. J.-E. Park and Professor Dr. Soo-Gil Park, Chungbuk National
University, for their valuable advice and discussions. It is a great pleasure to
express my deepest gratitude to Dr. Y. Okinaka for his advice, correction and
encouragement. I would tender my acknowledgments to Dr. Y. Yamauchi, Dr. H.
Mukaibo, Mr. Y. Fukuhara, Ms. A. Takai,and Mr. M. Komatsu, chapter 4 in my
thesis would not be accomplished without their cooperation. My appreciation goes
to all the people who have shared their time and talent with me in my everyday
battery life: Dr. S Motokawa, Mr. S. Kamei, Mr. H. Ito, Dr. T. Shimizu, Ms. J. Kim,
Mr. T. Sumi, Mr. N. Chihara, Ms. M. Hosoda, Mr. H. Kitoh, Mr. S. Tominaka, Mr. Y.
Wada, Ms. M. Iwasaki, Mr. M. Ueda, Mr. H. Obata, Ms. K. Saigusa, Mr. N.
Akiyama, Mr. R. Sebata, Mr. K. Takada, Mr. Y. Fukuhara, Ms. C. Tatsumi, Mr. S.
Ohta, Mr. K. Goto, Mr. S. Yamagami, Mr. H. Kawase, Mr. T. Kawano, and Mr. Y.
Nakamura. My colleagues, Mr. H. Iida, Ms. M. Matsunaga, Ms. K. Sakata and all
the other members of applied physical chemistry lab. So many of you have made me
feel at home and sometimes work hard. Finally I express greatest thanks to my
parents, Mr. Shinji Nara and Ms. Sayoko Nara, for deep affection and constant
financial support.
Hiroki Nara

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