Passivating ability of surface film derived from vinylene carbonate on tin negative electrode

Article abstract of DOI:10.1149/1.3561424

The passivating ability of surface film derived from vinylene carbonate (VC) is addressed on tin (Sn) negative electrode after a comparative study on the thickness, film growth pattern, chemical composition, and mechanical flexibility of the surface films generated from VC-free and VC-added electrolytes. The surface film derived from the former electrolyte is enriched by inorganic fluorinated and carbonate species, and shows a poor passivating ability to cause a continued electrolyte decomposition, film growth and eventual electrode failure. In contrast, organic carbon-oxygen species are dominant in the film derived from the VC-added electrolyte. Even if this film is thinner than the other, it passivates the Sn electrode surface more effectively. As a result, the film growth and electrode polarization are less significant. The superior passivating ability of organic-rich surface film has been ascribed to a uniform coverage and higher mechanical flexibility.

Full text of DOI:10.1149/1.3561424

A498
Journal of The Electrochemical Society, 158 (5) A498-A503 (2011)
0
013-4651/2011/158(5)/A498/6/$28.00 VC The Electrochemical Society
Passivating Ability of Surface Film Derived from Vinylene
Carbonate on Tin Negative Electrode
*,z
Sangjin Park, Ji Heon Ryu, and Seung M. Oh
Department of Chemical and Biological Engineering, WCU Program of C2E2, College of Engineering, Seoul National
University, Seoul 151-744, Republic of Korea
The passivating ability of surface film derived from vinylene carbonate (VC) is addressed on tin (Sn) negative electrode after a
comparative study on the thickness, film growth pattern, chemical composition, and mechanical flexibility of the surface films gen-
erated from VC-free and VC-added electrolytes. The surface film derived from the former electrolyte is enriched by inorganic fluo-
rinated and carbonate species, and shows a poor passivating ability to cause a continued electrolyte decomposition, film growth
and eventual electrode failure. In contrast, organic carbon-oxygen species are dominant in the film derived from the VC-added
electrolyte. Even if this film is thinner than the other, it passivates the Sn electrode surface more effectively. As a result, the film
growth and electrode polarization are less significant. The superior passivating ability of organic-rich surface film has been
ascribed to a uniform coverage and higher mechanical flexibility.
VC 2011 The Electrochemical Society. [DOI: 10.1149/1.3561424] All rights reserved.
Manuscript submitted October 18, 2010; revised manuscript received February 2, 2011. Published March 18, 2011.
The Li-alloying materials (Si and Sn) have been developed as an
alternative to the present graphitic carbons for lithium-ion batteries
(
ference in passivating ability and mechanical flexibility for two
surface films, the chemical compositions are analyzed by using
depth-profiling x-ray photoelectron spectroscopy (XPS) and Fourier-
transform infrared spectroscopy (FT-IR), whereas the film growth
pattern and thickness are estimated from the cross-sectional scanning
electron microscope (SEM) images of the cycled electrodes.
1
LIBs). The practical use of these high-capacity negative electro-
–4
des is, however, still limited due to their poor cycling behavior,
which may be caused by: (i) cracking/crumbling of the active mate-
rial (Si and Sn), (ii) electrical contact loss by a massive volume
change during alloying/de-alloying, and (iii) continuous electrolyte
decomposition and surface film formation, named solid electrolyte
5
interphase (SEI).
Experimental
The electrolyte decomposition and concomitant film deposition are
inevitable in the present LIBs since they operate far beyond the ther-
modynamic stability window of the common electrolytes. The suc-
cessful commercialization of LIBs has, however, been possible since
this surface film passivates the electrode surface, once it forms, to pre-
vent additional electrolyte decomposition. In order for surface film to
accomplish its own passivating role, the SEI layer should thus be com-
pact and perfectly cover the electrode surface to prevent additional
electrolyte decomposition. In addition, SEI layer should be thin to
The Sn film was electro-deposited on a piece of copper foil in a plat-
ing bath composed of 0.14 M tin (II) sulfate (SnSO ), 0.34 M potassium
4
sulfate (K
2 4 6 11 7
SO ), and 0.42 M sodium d-gluconate (C H O Na). The
–2
electro-deposition was performed at a current density of 5 mA cm for
8
min at room temperature.
Electrochemical performances of the Sn film negative electrodes
were examined using a coin-type cell (2032-type), which was fabri-
cated by inserting a polypropylene separator (20 lm) between the Sn
film electrode and lithium foil (Cyprus Co.). The used electrolyte was
þ
allow a fast Li ion movement. The mechanical flexibility of surface
6
1.0 M lithium hexafluorophosphate (LiPF ) dissolved in ethylene car-
film seems to be another important requirement particularly for the Li-
alloying materials. Namely, inflexible surface film cannot effectively
passivate the electrode surface since it is easily broken or deformed by
the massive volume change evolved with the Li-alloying materials,
leading to an exposure of electrode surface to electrolyte solution.
The SEI formation mechanism on the Li-alloying materials is some-
what different to that on the carbonaceous negative electrodes. On
graphite surface, SEI layer forms within a few initial cycles, whereas
the film formation is continued on the Li-alloying materials since new
electrode surface is generated as a result of continued cracking/crum-
bling. The film growth patterns are much more complicated on the
Li-alloying materials. As a result, the SEI study on the Li-alloying mate-
bonate (EC) and dimethyl carbonate (DMC) (1:1, volume ratio), in
which 5 wt. % of VC was added. Galvanostatic charge/discharge cy-
cling was made at a current density of 400 mA g in the voltage
ꢀ
1
þ
range of 0.01–1.2 V (vs. Li/Li ).
For a post mortem field-emission SEM (Model JSM-6700F,
JEOL), XPS and FT-IR analyses, the cells were dismantled in a glove
box and the electrodes were washed with dimethyl carbonate. A her-
metic vessel was used to transfer the samples from the glove box to
the instrument chamber. The XPS data were collected in an ultra-high
vacuum multipurpose surface analysis system (Sigma probe, Thermo,
ꢀ
10
UK) that operates at a base pressure of <10
mbar. The photoelec-
6,7
trons were excited by Al Ka (1486.6 eV) radiation at a constant
power of 150 W (15 kV and 10 mA); the x-ray spot size was 400 lm.
During data acquisition, a constant-analyzer-energy mode was used at
a pass energy of 30 eV and a step of 0.1 eV. The depth-profiling XPS
rials is still limited, while those on graphite electrodes are numerous.
The prime concern in this work is to assess the passivating abil-
ity of surface film derived from vinylene carbonate (VC), which is
8
,9
one of the common SEI formers. To this end, a comparative study
is made on the surface films derived from the VC-added and VC-
free electrolyte, from which the passivating ability of VC-derived
surface film is addressed. In details, the passivating ability is
assessed by measuring the reduction current associated with the
electrolyte decomposition on an exposed Sn electrode surface. To
estimate the mechanical flexibility, a bending experiment is devised,
in which the Sn electrode is bent after surface film deposition. An
expectation here is that, if the surface film is not mechanically flexi-
bility, it is deformed by bending to generate new Sn surface, onto
which electrolyte decomposition is induced. To account for the dif-
þ
measurement was made by a continued Ar ion sputter etching and
the chemical composition was estimated by using the atomic sensitiv-
10,11
ity factor.
For the cross-sectional images, the electrode samples
þ
were crosscut by using an Ar ion beam polisher (SM-09010, JEOL)
at a constant power of 0.5 W (5 kV and 0.1 mA) under vacuum
–
4
(
<2.0 ꢁ 10 Pa). The infrared spectra were obtained using the atte-
nuated total reflection (ATR) mode (Nicolet 5700 FT-IR spectrome-
ter). The cycled electrodes were mounted on an internal reflection
crystal (Ge) facing the surface layer toward the crystal. For the bend-
ing experiment, a pouch-type cell was fabricated with the Sn film
electrode (3ꢁ 3 cm), polypropylene separator and lithium foil. In this
report, the lithiation (alloying) is expressed as discharging, whereas
the de-lithiation (de-alloying) charging on the basis of the standard
lithium-ion cell configuration.
*
z
Electrochemical Society Active Member.
E-mail: seungoh@snu.ac.kr
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Journal of The Electrochemical Society, 158 (5) A498-A503 (2011)
A499
Results and Discussion
In the first cycle, the Sn negative electrode shows a comparable
charge-discharge behavior in two electrolytes with a lithiation capacity
of ca. 680 mA h g and coulombic efficiency of >85%. From the
second cycle, however, the Sn electrode delivers a different cycle per-
formance in two electrolytes. A severe capacity fading is observed in
the VC-free electrolyte, whereas a much better cycle performance in
the VC-added electrolyte with a de-lithiation capacity amounting to
Figure 1 compares the lithiation and de-lithiation voltage profiles
obtained with the Li/Sn cells in the VC-free and VC-added electrolyte.
–1
>85% of the initial value even after 50 cycles (Fig. 1c). The poor
cycle performance in the VC-free electrolyte seems to be deeply asso-
ciated with the electrode polarization that steadily increases with
cycling (Fig. 1a). Also noted in Fig. 1c is that the Sn film electrode
shows an irreversible capacity in every cycle, implying that the elec-
þ
trolyte decomposition and Li trapping inside the “dead” Sn grains
are continued with cycling. The detailed discussion on this will be
advanced in the later section.
The surface and cross-sectional images of the pristine Sn film
electrode are displayed in Figs. 2a and 2b, respectively. The Sn
grains of 1–2 lm are deposited on the Cu foil (Fig. 2a) and the Sn
layer thickness is ca. 1 lm (Fig. 2b). The cross-sectional images of
the cycled electrodes are also presented in Figs. 2c–2h. The follow-
ing points should be noted. First, both the surface film and Sn layer
Figure 2. (Color online) FE-SEM images of the Sn film electrode before cy-
cling: (a); top view and (b) side view. The cross-sectional images of Sn film
electrodes cycled in the VC-free (left panel) and VC-added electrolyte (right
panel): (c)/(d); after 5 cycles, (e)/(f); after 10 cycles, and (g)/(h); after 20
cycles. The EDS mappings are presented in (g)/(h). The pink and green
points correspond to C and Sn element, respectively. The images in (c)–(h)
were obtained after a full de-lithiation.
Figure 1. Galvanostatic discharge (lithiation, downward) and charge
(de-lithiation, upward) voltage profiles obtained with Li/Sn film cell in: (a);
VC-free electrolyte and (b); VC-added electrolyte. The cycle performance is
ꢀ
1
compared in (c). Current density ¼ 400 mA g . Voltage cut-off ¼ 0.01–1.2
þ
V (vs. Li/Li ).
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A500
Journal of The Electrochemical Society, 158 (5) A498-A503 (2011)
expands to the normal direction with cycling, implying that the elec-
trolyte decomposition and electrode swelling are continued with cy-
cling. Second, the growth of surface film is more prominent in the
VC-free electrolyte (left panel). Third, the population of Sn (green
points) is not confined within the initial Sn layer, rather extended to
the surface film layer as seen on the EDS mappings (Figs. 2g and 2h).
This illustrates that the Sn grains are cracked/crumbled with cycling
and some of them move from the initial Sn layer to the surface film
layer to be embedded into the surface film. A high population of
carbon (pink points) is also noticed in the Sn layer, reflecting that the
carbon-containing film forms not only in the surface film layer but
also within the Sn layer. From these, the following scenario is assumed
for the film formation. The Sn grains expand with alloying but do not
contract to the initial dimension upon de-alloying, leading to an elec-
6,12,13
trode layer expansion with cycling.
During this volume change,
the Sn grains are cracked/crumbled due to mechanical stress, and
some of the fine Sn grains are detached from the Sn layer. Then, addi-
tional electrolyte decomposition and film deposition take place on the
1,8,9,13,14
fresh Sn surfaces that are exposed inside the cracks.
The high
population of carbon in the Sn layer explains this feature. Film deposi-
tion is also possible on the detached Sn grains if they are electrically
connected to the Sn layer and Cu foil. Due to the continued film
growth, which is electrically insulating, the detached Sn grains are
even more electrically isolated from the Sn layer to be inactive
(
“dead”) as the negative electrode. The irreversible capacity appearing
þ
in every cycle (Fig. 1c) can thus be ascribed to the Li trapping inside
these dead Sn grains as well as the electrolyte decomposition on the
newly exposed Sn surface. The capacity fading shown in Fig. 1 can
also be accounted for by the continued film growth and dead Sn for-
mation. That is, the thicker surface film generated from the VC-free
electrolyte leads to the larger electrode polarization due to impeded
þ
Li ion diffusion. Due to this polarization, the lithiation voltage
reaches the lower cut-off earlier during the galvanostatic cycling,
leading to a decrease in the lithiation and de-lithiation capacity. The
substantial capacity fading in the VC-free electrolyte can also be
explained by the generation of dead Sn. As seen in Figs. 2g and 2h,
the initial Sn layer becomes thinner in the VC-free electrolyte, sug-
gesting that a larger amount of Sn grains are detached, such that the
formation of dead Sn is more probable in this electrolyte.
As shown in Fig. 2, the film growth rate differs in two electro-
lytes. Given that film does not grows if surface film effectively pas-
sivates electrode surface, it is likely that the passivating ability is
better for the surface film generated from the VC-added electrolyte.
To ensure the difference in passivating ability, cyclic voltammo-
grams were obtained after the film formation in two electrolytes.
This experiment was inspired by the previous reports that the cata-
lytic reductive decomposition of electrolytes takes place near 1.5 V
Figure 3. Cyclic voltammograms obtained after the film deposition with the
predetermined number of cycling. After the surface film deposition, the elec-
trode potential was moved from 1.2 to 2.0 V (vs. Li/Li ) to convert the lithi-
ated Sn (Li Sn) to metallic Sn and then from 2.0 to 0.8 V to detect the
x
catalytic reduction current.
þ
To gain an insight into the factors affecting the passivating abil-
ity, the chemical composition and thickness of two surface films are
analyzed by using XPS. The C 1s, O 1s and F 1s photoelectron
spectra are represented in Fig. 4. The C 1s spectra were fitted with
þ
15–17
(
detected, the lithiated Sn (Li Sn) should be converted to metallic
vs. Li/Li ) on a fresh Sn surface.
In order for this current to be
x
Sn, and furthermore the as-generated Sn is exposed to the electrolyte
solution since this reduction current evolves only on the Sn surface
exposed to electrolyte solution, not on either the Li-Sn alloy surface
or the Sn surface covered by passivation layer. Two features should
be noted in the cyclic voltammograms shown in Fig. 3. First, the ox-
idation current appears in both electrolytes during the potential
sweep from 1.2 to 2.0 V, which is associated with the de-alloying
four peaks according to the reported binding energies; ROCO Li
2
(CO -like) at 290 eV, O¼C-O at 288 eV, C-O at 286.5 eV, and CH
3
2
12,18–21
at 285 eV.
All these photoelectrons must be emitted from the
surface films since the Sn film electrode does not contain any carbon
species such as polymeric binder and conductive carbon. As seen
in the C 1s spectra, the population of carbon-oxygen species is
higher in the film derived from the VC-added electrolyte. The
enriched population of carbon-oxygen species in this film is further
confirmed from the O 1s spectra. In Fig. 4b, the O 1s spectra are
fitted with three broad peaks centered at 534.4 eV (poly vinylene
carbonate), 533.5 eV (oxygen bound to carbon with a single bond in
organic carbonates) and 532 eV (Li CO and oxygen atoms doubly
1
5–17
reaction of the Li Sn phases.
The larger oxidation current with
the thicker surface films illustrates that a larger amount of Li Sn
x
x
phases still remains at the last de-lithiation period of the film genera-
tion step due to slower Li diffusion through the thicker films. Sec-
þ
ond, the reduction current appears at 1.5 V when the electrode
potential moves from 2.0 to 0.8 V in the VC-free electrolyte (Fig. 3a).
This illustrates that the metallic Sn is restored upon de-alloying up to
2
3
10,18,22
bound to carbon atoms in organic carbonates).
The high pop-
2
.0 V, on which the reductive electrolyte decomposition takes place
ulation of carbon-oxygen species in the VC-derived film is con-
trasted by the enriched deposition of inorganic fluorinated species
(mainly LiF) that must be derived from the salt (LiPF ) in the
due to poor passivating ability of the surface film. In contrast, the
reduction current is not detected with the VC-derived surface films,
even with the thinnest one deposited by five charge/discharge cycles
(Fig. 3b). Clearly, the VC-derived surface film is more passivating
even if it is thinner than that forms in the VC-free electrolyte.
6
VC-free electrolyte (Fig. 4c). The stronger F 1s peaks are observed
in this surface film, which corresponds to LiF (685.2 eV) and Li PF
(687–688 eV).
x
y
10,18,21,23,24
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Journal of The Electrochemical Society, 158 (5) A498-A503 (2011)
A501
Figure 5. The variation of chemical composition as a function of film depth.
The results were obtained by fitting the depth-profiling XPS data. The mea-
surement was made after the first lithiation.
species must come from either the carbonate solvent (EC and DMC
in this work) or VC. Of the two, however, VC seems to be the more
probable source since this unsaturated carbonate is more vulnerable
25–29
to reductive decomposition than the saturated counterparts.
Another important observation in Fig. 5 is that the VC-derived sur-
face film is thinner than the other as evidenced by the earlier expo-
sure of Sn surface in the depth-profiling experiment. Namely, the
atomic percent of Sn is less than 20% even if the film in Fig. 5a is
sputter etched for 300 s, but this value reaches 60% after the same
period of etching on the VC-derived film. The earlier exposure of
Sn in the latter illustrates the presence of a thinner surface film to be
removed on the Sn electrode surface.
Figure 4. XPS spectra obtained with the surface films derived from the VC-
free and VC-added electrolyte. The measurement was made after the first
lithiation. Note that the spectra were obtained before the depth-profiling
measurement, such that the photoelectrons are emitted from the topmost sur-
face layer.
Figure 6 displays the FT-IR spectra obtained with two surface
films, in which a deposition of lithium carbonate (Li
ium alkyl carbonate (ROCO Li) is observed. The peaks indexed by
2 3
CO ) and lith-
2
–1
–1
–1
D [1493 and 1436 cm (mC-O,as), 1096 cm (mC-O,s), and 869 cm
CO3)] can be assigned to the infrared absorptions by Li CO
whereas the peaks indicated by ; (1647 cm (m ) and 1317
C¼O,as
–1
(
d
2
3
,
–1
The two surface films differ in chemical composition not only at
the topmost region (Fig. 4) but also in the deeper region as shown in
the depth-profiling XPS data (Fig. 5), in which the atomic concen-
tration of four elements is plotted as a function of etching time. The
higher population of fluorinated species is noted even in the deeper
region of the film derived from the VC-free electrolyte (Fig. 5a),
whereas the carbon-oxygen species are dominant over the fluori-
nated ones in the VC-derived film (Fig. 5b). The carbon-oxygen
30–32
cm (mC¼O,s)) to those for lithium alkyl carbonate (ROCO
2
Li).
The other vibrational absorptions for lithium alkyl carbonate at
1400 cm (dCH3, CH2) and 1074 cm (mC-O) cannot be located since
they overlap with those for Li CO . Two major differences can be
–1
–1
2
3
noticed in the IR spectra. First, the film derived from the VC-added
electrolyte contains the decomposition products of VC, of which the
–1
absorption peaks are indicated by : in Fig. 6. The peak at 1815 cm
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A502
Journal of The Electrochemical Society, 158 (5) A498-A503 (2011)
2
5,28
alkoxide (C¼COLi) have been proposed.
Second, the popula-
tion of organic carbon-oxygen species is larger in the film derived
from the VC-added electrolyte. That is, a comparison of two spectra
reveals that the absorption peaks belonging to the organic C-O spe-
cies (for instance, ROCO Li and VC-derived ones) are larger in the
2
film derived from the VC-added electrolyte at the expense of those
for the inorganic species (Li CO ).
2 3
By combining the XPS, FT-IR spectra and the film growth pat-
tern shown in Fig. 2, the factors affecting the passivating ability of
surface films may be delineated. An assumption can be made from
the experimental observations that organic-rich surface films have a
stronger passivating ability than inorganic-rich ones, probably due
to a more uniform coverage and higher degree of compactness. An
ensuring feature is that the surface film deposited from VC in the
first cycle is organic-rich (Fig. 5). Even if this film is thinner, the
film growth is less significant from the second cycle (Fig. 2). In con-
trast, the inorganic-rich film is less passivating to allow a continued
film growth (Fig. 2). In order for any thin surface films to have a
strong passivating ability, they must be uniformly covered with a
high degree of compactness. To draw a conclusion that organic-rich
films have a superior passivating ability to inorganic-rich ones, how-
ever, additional experimental data should be collected. Also the rea-
sons why organic species are more passivating should be studied.
Figure 7 presents the schematic illustration for the bending experi-
ment and the current evolution observed after the electrode bending.
The reduction current relevant to the electrolyte decomposition on
Figure 6. FT-IR spectra obtained with the surface films derived from VC-
free and VC-added electrolyte. The measurement was made after the first
lithiation.
can be assigned to the stretching vibration of C¼O bonds in poly
–
1
(
bonds. As the double-bond containing products to be derived from
vinylene carbonate) and that at 1624 cm to the stretching of C¼C
28
Li Sn is larger when the electrode is bent by a larger angle in both
x
VC, lithium vinylene carbonate (C¼COCO Li) and lithium vinylene
electrolytes. This must be due to a higher degree of film deformation
to lead an enlarged exposure of Li Sn surface to the electrolyte. The
2
x
reduction current is, however, smaller for the film generated from VC,
suggesting that the VC-derived surface film is mechanically more flex-
ible. This favorable mechanical property seems to be inherited by the
high population of organic species rather than inorganic ones in the
VC-derived film. Furthermore, the higher passivating ability of this
film seems to be partly originated by the higher flexibility. With this
flexible film, the exposure of electrode surface to electrolyte solution
seems less likely since the film is not easily broken or deformed upon
a massive volume change of the Sn component.
Conclusions
The physicochemical properties of surface films derived from
the VC-free and VC-added electrolyte were examined, from which
the role of VC as a surface-film former has been addressed. The fol-
lowing points are summarized.
1. The surface films steadily grow with cycling in both VC-free
and VC-added electrolyte. However, a thicker film is deposited from
the VC-free electrolyte. The Sn electrode shows a poorer cycle per-
formance in this electrolyte with a larger electrode polarization, which
þ
is caused by a slower Li diffusion through the thicker surface layer.
The VC-derived surface film is, however, thinner and cycle perform-
ance is better in this electrolyte with a less significant polarization.
2
. The Sn grains are crumbled with cycling and some of fine Sn
grains move to the surface film layer. The generation of this “dead”
Sn seems to be an important degradation mechanism prevailing in
this Li-alloying negative electrode.
3
. In the first lithiation, organic-rich surface film is deposited in
the VC-added electrolyte. Even if this film is thin, it effectively pas-
sivates the Sn surface to suppress additional electrolyte decomposi-
tion. In the VC-free electrolyte, however, a surface film enriched by
inorganic fluorinated and carbonates species is generated. This film
shows a poorer passivating ability probably due to non-uniform cov-
erage and poor mechanical flexibility.
Figure 7. (a); Schematic illustration for the bending experiment and (b); the
reduction current observed after the bending. To generate surface films, the
þ
potential of Sn film electrode was fixed at 0.01 V (vs. Li/Li ) for 2 h. The
Acknowledgments
ꢂ
pouch-type cells were then bent by 10 and 30 while the electrode potential
being fixed at 0.01 V. The reduction current associated with electrolyte
This research was supported by WCU program through the
National Research Foundation of Korea funded by the Ministry of
x
decomposition on lithiated Sn (Li Sn) was traced.
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Journal of The Electrochemical Society, 158 (5) A498-A503 (2011)
A503
Education, Science and Technology (R31-10013). The authors also
wish to acknowledge the Research Center for Energy Conversion
and Storage for financial support.
16. J. Hassoun, P. Reale, and S. Panero, J. Power Sources, 174, 321 (2007).
1
1
1
2
7. S. D. Beattie, T. Hatchard, A. Bonakdarpour, K. C. Hewitt, and J. R. Dahn, J. Elec-
trochem. Soc., 150, A701 (2003).
8. L. El Clualtania, R. Dedryvere, J. B. Ledeuil, C. Siret, P. Biensan, J. Desbrieres,
and D. Gonbeau, J. Power Sources, 189, 72 (2009).
Seoul National University assisted in meeting the publication costs of
this article.
9. H. Bryngelsson, M. Stjerndahl, T. Gustafsson, and K. Edstr, J. Power Sources,
174, 970 (2007).
0. R. Dedryvere, L. Gireaud, S. Grugeon, S. Laruelle, J. M. Tarascon, and D. Gon-
beau, J. Phys. Chem. B, 109, 15868 (2005).
References
1
. S. C. Chao, Y. C. Yen, Y. F. Song, Y. M. Chen, H. C. Wu, and N. L. Wu, Electro-
chem. Commun., 12, 234 (2010).
. M. Winter and R. J. Brodd, Chem. Rev., 104, 4245 (2004).
21. A. M. Andersson, M. Herstedt, A. G. Bishop, and K. Edstrom, Electrochim. Acta,
47, 1885 (2002).
22. Y.-C. Lu, A. N. Mansour, N. Yabuuchi, and Y. Shao-Horn, Chem. Mater., 21,
4408 (2009).
23. L. Daheron, R. Dedryvere, H. Martinez, M. Menetrier, C. Denage, C. Delmas, and
D. Gonbeau, Chem. Mater., 20, 583 (2008).
24. M. Stjerndahl, H. Bryngelsson, T. Gustafsson, J. T. Vaughey, M. M. Thackeray,
and K. Edstrom, Electrochim. Acta, 52, 4947 (2007).
25. L. Chen, K. Wang, X. Xie, and J. Xie, J. Power Sources, 174, 538 (2007).
26. S. S. Zhang, J. Power Sources, 162, 1379 (2006).
27. L. B. Chen, K. Wang, X. H. Xie, and J. Y. Xie, Electrochem. Solid-State Lett., 9,
A512 (2006).
2
3
4
5
6
7
8
. R. A. Huggins, J. Power Sources, 81–82, 13 (1999).
. C. J. Wen and R. A. Huggins, J. Electrochem. Soc., 128, 1181 (1981).
. M. Winter and J. G. O. Besenhard, Electrochim. Acta, 45, 31 (1999).
. J. O. Besenhard, J. Yang, and M. Winter, J. Power Sources, 68, 87 (1997).
. J. Yang, M. Winter, and J. O. Besenhard, Solid State Ionics, 90, 281 (1996).
. H. Mukaibo, T. Momma, Y. Shacham-Diamand, T. Osaka, and M. Kodaira, Elec-
trochem. Solid-State Lett., 10, A70 (2007).
9
. I. A. Courtney, W. R. McKinnon, and J. R. Dahn, J. Electrochem. Soc., 146, 59
(
1999).
0. R. Dedryvere, S. Laruelle, S. Grugeon, L. Gireaud, J. M. Tarascon, and D. Gon-
beau, J. Electrochem. Soc., 152, A689 (2005).
1
28. H. Ota, Y. Sakata, A. Inoue, and S. Yamaguchi, J. Electrochem. Soc., 151, A1659
(2004).
1
1
1. J. H. Scofield, J. Electron Spectrosc. Relat. Phenom., 8, 129 (1976).
2. K. K. D. Ehinon, S. Naille, R. Dedryvere, P. E. Lippens, J. C. Jumas, and D. Gon-
beau, Chem. Mater., 20, 5388 (2008).
29. D. Aurbach, K. Gamolsky, B. Markovsky, Y. Gofer, M. Schmidt, and U. Heider,
Electrochim. Acta, 47, 1423 (2002).
30. J.-T. Li, S.-R. Chen, X.-Y. Fan, L. Huang, and S.-G. Sun, Langmuir, 23, 13174
(2007).
31. D. Aurbach, B. Markovsky, A. Shechter, Y. Ein-Eli, and H. Cohen, J. Electro-
chem. Soc., 143, 3809 (1996).
32. S. Oohira, M. Kakihana, Y. Fujii, T. Nagumo, and M. Okamoto, J. Nucl. Mater.,
133–134, 201 (1985).
1
1
1
3. M. Wachtler, J. G. O. Besenhard, and M. Winter, J. Power Sources, 94, 189
(
2001).
4. S. Yang, P. Y. Zavalij, and M. S. Whittingham, Electrochem. Commun., 5, 587
2003).
5. I. T. Lucas, E. Pollak, and R. Kostecki, Electrochem. Commun., 11, 2157 (2009).
(
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