Solid State Ionics 239 (2013) 50–55
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Solid State Ionics
journal homepage: www.elsevier.com/locate/ssi
Electrical transport in Li2SO4–Li2O–P2O5 ionic glasses and glass–ceramic composites:
A comparative study
Munesh Rathore, Anshuman Dalvi ⁎
Physics Department, Birla Institute of Technology and Science, Pilani, RJ 333031, India
a r t i c l e
i n f o
Article history:
Received 12 December 2012
Received in revised form 20 February 2013
Accepted 11 March 2013
Available online xxxx
Keywords:
Ionic conductivity
Crystallization
Li+ ion glasses
Glass–ceramics
Impedance spectroscopy
a b s t r a c t
A comparative investigation carried out between glasses and glass–ceramics, in the system (Li2SO4)x–(LiPO3)1 − x,
reveals interesting results. The conventionally melt-quenched compositions for x ≤ 60 mol% were found to be
purely glassy in nature. The glass–ceramic composites obtained by crystallization of the glassy samples were
found to be composed of fine nanocrystallites of LiPO3 and Li2SO4 embedded in the glass matrix. The electrical conductivity, in both glasses as well as glass–ceramics, increases with Li2SO4 content and found to be maximum for a
composition with 60 mol% of Li2SO4. Scaling of the conductivity spectra reveals that the relaxation dynamics of Li+
ions is independent of temperature and composition for glasses as well as glass–ceramics. Further, the cyclic
voltammetry investigations suggest a relatively better stability of glass–ceramic samples at least up to 300 °C.
© 2013 Elsevier B.V. All rights reserved.
1. Introduction
The ionic conductivity of a solid electrolyte plays an important
role in the functioning of the all-solid-state battery/cell. For example,
the higher the ionic conductivity the better will be the power output
of the ionic device [1]. Further, to realize high energy density batteries
for future ionic devices, development of fast ionic solids has been a
major concern for the ionics community [2]. Moreover, highly ion
conducting and thermal stable solid electrolytes are also desirable
for device miniaturization.
Fast ionic glasses have wide applications as electrolytes in solid state
ionic devices because of their high ionic conductivity, liquid-like structure and negligible electronic conductivity [1–6]. Ever since the discovery of Li2O–MxOy (MxOy = P2O5, SiO2, B2O3, Al2O3 etc.) glassy ionic
systems, various Li+ ion conducting glasses have been developed and
thoroughly characterized [3–11]. Ionic conductivity enhancement in
these glasses was achieved by various routes in the last two decades.
It was soon realized that salt incorporation in the binary system further
enhances ionic conductivity [11–20]. Addition of salts, e.g., Li2SO4
[11–18] and LiCl [19,20] in oxide glass matrix (Li2O–MxOy) was found
to influence the conductivity significantly. The enhancement in the
ionic conductivity was attributed to salt incorporation in the macromolecular network, as suggested by various workers [11–18].
Despite the high ionic conductivity, glasses have a tendency to crystallize above a certain temperature (Tc) and hence the direct use of
glassy electrolytes may affect the performance of the solid state battery,
⁎ Corresponding author. Tel.: +91 1596 515498; fax: +91 1596 244183.
E-mail address: [email protected] (A. Dalvi).
0167-2738/$ – see front matter © 2013 Elsevier B.V. All rights reserved.
http://dx.doi.org/10.1016/j.ssi.2013.03.022
especially where extreme and abrupt changes in temperatures are inevitable (e.g., in satellites or defense related applications) [21,22]. Thus, in
view of better expected thermal stability, Li+ ion conducting glass–
ceramic composites, produced by controlled crystallization of the glasses,
have recently got considerable attention [23–30]. Though most of these
glass–ceramic composites exhibit relatively poor conductivity than the
host glasses, in many systems notably higher ionic conductivity is also
observed in the former [27–30]. Although there are some interesting
studies on glass–ceramics and their thermal stability issues, a systematic
comparative electrical transport study with those of the host glasses
further needs emphasis. Furthermore, highly Li+ ion conducting sulfide
glass–ceramics have been explored thoroughly, and are also used in
commercially available devices, but their stability in air is not reliable
[3–6,10,31]. Some of these also exhibit poor stability when in contact
with lithium metal [31].
On the other hand, although the commercially applied solid polymer
electrolytes (SPEs) in Li+ ion batteries exhibit high ionic conductivity,
they are stable only up to ~60 °C and hence are not suitable for high
temperature battery applications [1,29]. Thus, oxide (Li2O) based glass–
ceramics should also be explored and revisited because of their potential
use in future ionic devices, especially, for high temperature applications
[1,23–30]. In order to develop suitable glass–ceramic composites for
high temperature ionic devices, a thorough and systematic understanding of the crystallization process is found to be very useful. Crystallization
and glass transition kinetics studies on various fundamental [33–36] and
applied [37–42] ion conducting glasses have been carried out. Recently,
Li+ ion glasses have also been investigated using crystallization kinetics
and thermal stability parameters have been calculated and analyzed by
different theoretical models [39]. There are limited but interesting studies in which the conductivity (σ)–temperature (T) cycles have also been
M. Rathore, A. Dalvi / Solid State Ionics 239 (2013) 50–55
investigated crossing the glass transition (Tg) and crystallization (Tc)
temperatures. Those have again provided better understanding of various interesting phenomena, for example, phase transformation, structural relaxation at Tg and effect of precipitation of new compounds on
the ionic conductivity of various samples at Tc [33–42]. Souquet et al.
[32] measured the electrical conductivity of some Na+ ion glasses
(Na2O–SiO2) slightly above the Tg and found a significant rise of the
conductivity deviating from Arrhenius behavior that was attributed to
the increase of the free volume of the glass matrix at Tg. In another
approach, St. Adams et al. [33,34] measured the conductivity above
the Tg as well as Tc on various Ag+ ion conducting glasses where the
anomalous rise in the conductivity at Tg was attributed to the glass–
crystallite interface effect and its fall at Tc to crystallization.
In order to develop better glass–ceramic composites as electrolytes
for high temperature battery applications, the present work attempts
a comparative investigation between glasses and the glass–ceramics
in a well known Li2SO4–Li2O–P2O5 system [13,18]. The effect of crystallization on the electrical transport in Li2SO4–Li2O–P2O5 fast ionic glasses
is emphasized. The fundamental properties of the glass–ceramics are
compared with those of the parent ionic glasses with emphasis on electrical characterization.
2. Experimental
The glasses of compositions xLi2SO4–(100 − x)(0.5Li2O–0.5P2O5)
were chosen for the present investigation and abbreviated as xLSLP
(e.g. x = 10 mol% is referred to as 10LSLP). The composition for
x = 0 (i.e. 0.5Li2O–0.5P2O5) is abbreviated as LP. Glasses were synthesized using high-purity powders of Li2CO3, NH4H2PO4, and Li2SO4
by conventional melt-quenching. Appropriate mixtures of these
materials were taken in porcelain crucibles and heated to 450 °C for
2 h (to remove CO2, NH3 and H2O). The mixtures were then melted
at 900–950 °C for 30 min and the subsequently pressed between
two copper plates to obtain transparent glasses. The glass–ceramic
samples were prepared by heating/annealing the glass above the
crystallization temperature followed by slow cooling. The structure
of the glasses and glass–ceramics was studied by X-ray diffraction
(Rigaku MiniFlex II) with Cu-Kα radiation. Differential scanning calorimetry (Shimadzu-60, DSC) was performed to confirm the glassy
nature. The glassy/glass-ceramic samples were ground to fine powder
and pelletized and used for electrical characterization. Graphite (C)
paint was used to achieve better electrode–electrolyte contact. The
electrical conductivity (σ)–temperature (T) cycles were obtained
using computer interfaced HIOKI 3532-50 LCR meter at a controlled
heating rate of 1 °C/min. Impedance spectroscopy measurements
(42 Hz–5 MHz) were also performed on the glasses and the glass–
ceramics. To compare the electrochemical stability of the glasses and
glass–ceramics, cyclic voltammetry investigations were carried out on
the cells in two electrode configuration of the types C|electrolyte|C and
C|electrolyte|LiCoO2 using a Princeton 263A potentiostat/galvanostat.
Fig. 1. X-ray diffraction patterns for as prepared glassy samples.
found to be suppressed, whereas, that of Li2SO4 is facilitated. The average particle size of these precipitated crystallites calculated using the
Debye–Scherrer formula [45] is found to be ~20–30 nm.
Furthermore, to understand the sequence of precipitation of the compounds, pristine 10LSLP glassy sample is annealed at two different temperatures and the XRD patterns are shown in Fig. 3. Annealing at 350 °C
(a temperature close to first crystallization, revealed by differential scanning calorimetry) results in the precipitation of two compounds, viz.,
LiPO3 and Li2SO4 in little amounts. On further annealing of the sample
at 420 °C significant growth of LiPO3 peaks is evident. Thus it may be
suggested here that crystallization has essentially two stages: (i) mainly
Li2SO4 comes out of the glass matrix at the first crystallization and (ii) at
a higher temperature second stage major precipitation of LiPO3 is
observed and (iii) overall crystallization of LiPO3 is reduced with Li2SO4
content in the glass matrix.
3.2. Differential scanning calorimetry (DSC)
The DSC patterns for 10 and 60LSLP glassy samples at a typical
heating rate of 20 °C/min are shown in Fig. 4. The pattern for 10LSLP
sample exhibits a single broad exothermic peak with onset at ~390 °C
and width of ~100 °C that may be due to massive crystallization of
mainly LiPO3 as also suggested by XRD (Fig. 2). Whereas, for the
60LSLP sample (Fig. 4b), in addition to the peak at 396 °C (Tc2), one
more sharp and significant exothermic peak is observed at ~288 °C
(Tc1). These two successive peaks may correspond to the precipitation
of Li2SO4 (at Tc1) and LiPO3 at (Tc2), respectively. Crystallization of two
compounds may further suggest two types of surroundings for Li+
3. Results and discussion
3.1. X-ray diffraction
The X-ray diffraction (XRD) patterns of the as prepared glassy
samples are shown in Fig. 1. The complete absence of diffraction
peaks (Fig. 1a–d) suggests that the samples with ≤60 mol% Li2SO4
are indeed glassy in nature. Above x ≥ 60 mol% of Li2SO4 content,
there appear tiny peaks (Fig. 1e) that suggests the partially crystalline
nature of the as prepared samples.
The XRD patterns of the glass–ceramic samples shown in Fig. 2
indicate the crystallization of various compounds. For the LP sample
(0 mol% of Li2SO4) there is a precipitation of LiPO3 alone. Further,
addition of Li2SO4 in the glass matrix leads to crystallization of LiPO3
along with Li2SO4. It is also evident that the crystallization of LiPO3 is
51
Fig. 2. XRD patterns for glassy–ceramic samples.
52
M. Rathore, A. Dalvi / Solid State Ionics 239 (2013) 50–55
Fig. 3. XRD patterns for 10LSLP glass annealed at (a) 350 °C and (b) 420 °C.
ions, viz., Li (1) and Li (2) ions may be surrounded by P–O and SO4−2
units, respectively.
3.3. Electrical conductivity
At the outset, in order to observe the thermal events and microstructural changes during crystallization the electrical conductivity
(1 kHz) is measured as a function of temperature at a typical heating
rate of 1 °C/min. The conductivity for the pristine glassy samples is
plotted as a function of Li2SO4 content for two different temperatures
in the inset of Fig. 5A. Interestingly, the conductivity increases significantly and exhibits a maximum (of ~ 10 −6 Ω −1 cm −1 at 100 °C) for
60LSLP and further drops because of the increased crystallinity in the
samples (x ≥ 70).
The σ–T cycles for the as prepared glasses for clarity are divided in
three different regions, viz., I–III. In the thermally stable region I, the
electrical conductivity exhibits apparent Arrhenius behavior with
reversible cycles. The activation energy, which decreases with Li2SO4
content, is found to be in the range ~0.56–0.80 eV. In the thermally
unstable region II, initially the electrical conductivity deviates from
linearity and exhibits a remarkable rise that may be attributed to an
increase in the free volume because of the presence of the glass transition temperature (Tg) of the respective samples [32]. The Tg values calculated from σ–T cycles are found to be decreasing with Li2SO4 content.
Similar trend of the composition dependence of Tg is also observed in
the DSC patterns (Fig. 4).
On further heating (region II), the conductivity exhibits a peak
followed by a drastic fall. The fall and saturation in the conductivity
may be attributed to massive crystallization and its completion,
Fig. 4. DSC scans at a heating rate of 20 °C/min for (a) 10LSLP and (b) 60LSLP glassy
samples.
Fig. 5. A: Electrical conductivity verses temperature cycles for the following glassy
samples: (a) LP; (b) 10LSLP; (c) 30LSLP; (d) 50LSLP; (e) 60LSLP and (f) 70LSLP. Inset: conductivity versus Li2SO4 content for glassy samples. B: σ–T cycles on an extended scale in
the thermally unstable region. Fall in the conductivity is mainly due to precipitation of
LiPO3.
respectively. Further in region III, the conductivity increases
Arrheniusly. It is also noted that for the LP glass, a hump with single
peak is observed at which the conductivity falls prominently (Fig. 5A)
that may be due to precipitation of LiPO3 alone. Similar hump in the
conductivity appeared with two successive peaks for the samples with
low Li2SO4 content (e.g. 10LSLP). The conductivity drop at the two
peaks may be attributed to the precipitation of Li2SO4 (Tc1) and LiPO3
(Tc2), respectively. Evidently, the magnitude of conductivity drop at
Tc2 decreases as the Li2SO4 content increases. Further, a plateau like
behavior is apparent in the σ–T cycles that may suggest a gradual and
slow crystallization of LiPO3 and Li2SO4 compounds. For more clarity,
on an extended scale, the region II for LP and 60LSLP is once again
shown in Fig. 5B where the fall in the conductivity is clearly suppressed
after incorporation of Li2SO4 in the glass matrix. Thus, the σ–T cycles are
in good agreement with those of the XRD results.
The samples are slowly cooled to room temperature and second
heating cycles are performed. Now the samples transformed into
glass–ceramics and the σ–T cycles for these are separately shown in
Fig. 6. Firstly, the σ–T cycles for the all the samples exhibit Arrhenius
behavior with reversible nature up to ~ 500 °C. Unlike the behavior
seen for pristine glasses, the σ–T cycles are fairly linear suggesting
good thermal stability of the glass–ceramic samples. The conductivity
is again found to be increasing with Li2SO4 (inset of Fig. 6) content
M. Rathore, A. Dalvi / Solid State Ionics 239 (2013) 50–55
and exhibits a maximum (~ 10 −8 Ω −1 cm −1 at 100 °C) for the sample with 60 mol% Li2SO4. It is also interesting to note that the conductivity of the 60LSLP glass–ceramic is almost comparable to that of LP
glass in its thermally stable region. It may be further noted that in
the binary glassy systems (e.g., Li2O–MxOy) the rise in the electrical
conductivity is mainly due to Li + ions supplied by Li2O in the glass
matrix. In the present study, however, we have kept the Li2O–P2O5
ratio constant to ensure the formation of the LiPO3 compound alone
and its expected two phase composite formation with Li2SO4 during
crystallization Thus, these glass–ceramic samples may be assumed
to have a composition of (Li2SO4)x–(LiPO3)1 − x, in which LiPO3 may
exist typically in (i) glassy/amorphous as well as (ii) crystalline
form. Thus it is the Li + ion content from Li2SO4 whose increasing fraction may be responsible for the rise of the conductivity in glass–
ceramic samples.
To further compare electrical transport, impedance spectroscopy
investigations have also been performed on all the samples and
shown here for best conducting 60LSLP glass and its corresponding
glass–ceramic.
The Cole–Cole or Nyquist plots (for T ≪ Tc) for glassy samples is
shown in Fig. 7. In the case of the glasses, depressed semicircles are
observed followed by an inclined straight line that may be modeled
as a parallel combination of a resistance R and constant phase element (CPE1) for bulk in series with another element CPE2 (interface).
The low frequency spur is most likely due to the polarization at the
interface and suggests the absence of grain boundaries in the glasses.
However, in the case of the glass–ceramics, the Nyquist plots (Fig. 8)
appear to be noticeably dissimilar. In this case, the two consecutive
depressed semicircles are visible in the plot in which each of the semicircles may be modeled as a parallel combination of a resistance R and a
constant phase element CPE. In view of the fact that the glass–ceramic
samples may be visualized as tiny crystallites surrounded by glass
tissue, the electrical transport is expected within the grains as well as
through the grain boundary. Thus the two semicircles at higher and
low frequencies may be attributed to transport through bulk response
and grain boundary, respectively.
The dissipation factor tan δ versus ω for 60 mol% of Li2SO4 content
(Fig. 9a, b) also exhibits relaxation which is typically associated with
dielectric loss due to conduction. The nature of the relaxation peaks is
found to be similar for the glass and glass–ceramics and found to be
shifting towards higher frequency with an increase in the temperature.
Hopping frequency for ionic conduction (ωp) is obtained from the
peaks and plotted as a function of the inverse of the temperature in
Fig. 10. Interestingly, ln ωp verses 1/T plot is very much linear for
Fig. 6. σ–T cycles for glassy–ceramic samples, (a) LP, (b) 10LSLP (c) 30LSLP, (d) 50LSLP,
(e) 60LSLP and (f) 70LSLP. Inset: compositional dependence of conductivity at 100 °C.
53
Fig. 7. Nyquist plots for 60LSLP glassy sample.
glass and glass–ceramics and the activation energy of hopping transport, obtained from the slope, is also found to be comparable to their
respective values obtained from the σ–T plots. Thus it may be suggested here that the conductivity is predominately due to the hopping transport of Li + ions in glasses as well as in glass–ceramics.
To further understand the conductivity relaxation process in glasses
and glass–ceramics the conductivity spectra is plotted as a function of
frequency at different temperatures. The scaling behavior is studied
according to the model suggested by Ghosh et al. and others [43,44],
n ω
σ ðωÞ ¼ σ dc 1 þ
ωc
ð1Þ
which is the sum of the dc conductivity, σdc, and frequency-dependent
dispersive conductivity that exhibits a power law behavior with exponent n. Here ωc is a characteristic cross over frequency from dc to dispersion region and the same can also be obtained from the relaxation peak
Fig. 8. Nyquist plots for the 60LSLP glass–ceramic sample at (a) 450 °C, (b) 500 °C and
(c) 550 °C.
54
M. Rathore, A. Dalvi / Solid State Ionics 239 (2013) 50–55
3.4. Cyclic voltammetry
Fig. 13 shows cyclic voltammograms for most conducting composition
of glass obtained at 200 °C as well as for glass–ceramics (at 300 °C) with
different electrode configurations. Fig. 13a and c represents the CV curves
for samples sandwiched between two inert graphite (C) electrodes. Interestingly, the glassy sample exhibits a stability window of at least ~4 V.
The glass–ceramic samples however show relatively poor current density,
nevertheless, they do exhibit stability comparable to that of the glassy
samples even at higher temperatures. The current for both systems does
not exhibit any abrupt change, rather, a systematic variation with applied
voltage with a little hysteresis is observed in cycles between −4 to 4 V.
Further, Fig. 13b and d represents the CV curves for samples sandwiched
between LiCoO2 and inert graphite (C) electrodes to realize a Li+ ion
battery configuration. Again the overall stability window of at least ~4 V
is apparent with no significant change in the current. It is also observed
that repeated cycling hardly affects stability. Thus it may be suggested
here that the glass and glass–ceramic solid electrolytes are stable
under battery conditions. Since the glass–ceramic samples show stability
up to relatively higher temperatures, with proper engineering of the
electrode thickness of electrolytes and proper compositional variations,
glass–ceramic samples are suitable candidates for high temperature
applications.
4. Conclusions
Comparative study reveals interesting results. Addition of Li2SO4
affects the electrical as well as the thermal properties of glasses and
glass–ceramics. It is seen that a large amount of (60 mol%) Li2SO4 is
found to be completely dissolved in Li2O–P2O5 glassy matrix. In both
Fig. 9. Temperature dependence of tan δ versus ω for 60LSLP (a) glassy and (b) glass–
ceramic samples.
from the tan δ–ω plots. For glassy and corresponding glass–ceramic samples, σ/σdc versus ω/ωc is plotted for different temperatures as well as for
compositions and shown in Figs. 11 and 12, respectively. As apparent, for
60LSLP glassy as well as the glass–ceramic samples, perfect scaling is
seen that clearly indicates the (i) temperature and (ii) composition independent process of relaxation for both samples. Thus, it may be suggested that for the chosen range of frequency (42 Hz–5 MHz), the
relaxation dynamics of ions in glass and glass–ceramics appears to be
similar. In addition, the glass–ceramic formation affects the conductivity
but not the conductivity relaxation process.
Fig. 10. Log ωP vs 1000/T plot for 60LSLP (a) glass and (b) glass–ceramic samples.
Fig. 11. Scaling behavior at different temperatures for 60LSLP (a) glassy and (b) glass–
ceramic samples.
M. Rathore, A. Dalvi / Solid State Ionics 239 (2013) 50–55
55
Acknowledgments
This work is supported by DST-FIST and UGC Special Assistance
Programme, Government of India. We sincerely thank Professor S. A.
Hashmi, Delhi University, India for fruitful discussions on cyclic
voltammetry.
References
Fig. 12. Scaling behavior for different compositions of (a) glasses (200 °C) and
(b) glass–ceramic samples (400 °C).
cases, the ionic conductivity is found to be increasing with the addition
of Li2SO4 content that may be attributed to an increase in the mobile Li+
ions in the Li2O–P2O5 glass matrix. Interestingly, the electrical conductivity of the 60LSLP glass–ceramic is found to be comparable to that of
the pristine glass with no Li2SO4 content (LP). Addition of Li2SO4 also
leads to multiple crystallizations in the Li2O–P2O5 glassy system though
the overall crystallization is found to be suppressed. Scaling behavior
results suggest that the relaxation process is independent of temperature
and composition in both cases. The glass–ceramic samples are thermally
stabler and thus suitable for high temperature battery applications.
Fig. 13. Cyclic voltammograms for the 60LSLP G (glass) and GC (glass–ceramic) samples in different electrode configurations.
[1] M. Park, X. Zhang, M. Chung, G.B. Less, A.M. Sastry, J. Power Sources 195 (2010)
7904.
[2] K. Shahi, Phys. Status Solidi 41 (1977) 11.
[3] H.L. Tuller, D.P. Button, D.R. Uhlmann, J. Non-Cryst. Solids 40 (1980) 93.
[4] T. Minami, A. Hayashi, M. Tatsumisago, Solid State Ionics 136 (2000) 1015.
[5] M. Duclot, J.L. Souquet, J. Power Sources 97 (2001) 610.
[6] P. Knauth, Solid State Ionics 180 (2009) 911.
[7] R. Collongues, A. Kahn, D. Michel, Ann. Rev. Mater. Sci. 9 (1979) 123.
[8] J. Kawamura, R. Asayama, N. Kuwata, O. Kamishima, in: T. Sakuma, H. Takahashi (Eds.),
Physics of Solid State Ionics, Research Signpost, Kerala, India, 2006, pp. 193–246.
[9] D.P. Button, L.S. Mason, H.L. Tuller, D.R. Uhlmann, Solid State Ionics 9 (1983) 585.
[10] K. Minami, F. Mizuno, A. Hayashi, M. Tatsumisago, Solid State Ionics 178 (2007)
837.
[11] In: K.J. Rao, M. Ganguli, M.A.Z. Munshi (Eds.), Handbook of Solid State Batteries
and Capacitors, World Scientific, Singapore, 1995.
[12] M. Ganguli, K.J. Rao, J. Non-Cryst. Solids 243 (1999) 251.
[13] M. Ganguli, M.H. Bhat, K.J. Rao, Solid State Ionics 122 (1999) 23.
[14] P.R. Gandhi, V.K. Deshpande, K. Singh, Solid State Ionics 36 (1989) 97.
[15] A.V. Deshpande, V.K. Deshpande, Solid State Ionics 177 (2006) 2747.
[16] C.N. Reddy, R.P.S. Chakradhar, IOP Conf. Ser.: Mater. Sci. Eng. 2 (2009) 012053.
[17] J.L. Souquet, A. Kone, M. Ribes, J. Non-Cryst. Solids 38 (1980) 307.
[18] V.K. Deshpande, S.G. Charalwar, K. Singh, Solid State Ionics 40 (1990) 689.
[19] A.V. Deshpande, V.K. Deshpande, Solid State Ionics 154 (155) (2002) 433.
[20] V.K. Deshpande, IOP Conf. Ser.: Mater. Sci. Eng. 2 (2009) 012011.
[21] A.D. Robertsona, A.R. Westa, A.G. Ritchieb, Solid State Ionics 104 (1997) 1.
[22] G.L. Henriksen, D.R. Vissers, J. Power Sources 51 (1994) 115.
[23] X. Xu, Z. Wen, Z. Gu, X. Xu, Z. Lin, Solid State Ionics 171 (2004) 207.
[24] T. Minami, A. Hayashi, M. Tatsumisago, Solid State Ionics 177 (2006) 2715.
[25] M. Tatsumisago, F. Mizuno, A. Hayashi, J. Power Sources 159 (2006) 193.
[26] Y. Inda, T. Katoha, M. Babab, J. Power Sources 174 (2007) 741.
[27] K. Minami, A. Hayashi, S. Ujiie, M. Tatsumisago, J. Power Sources 189 (2009) 651.
[28] J.S. Thokchom, B. Kumar, J. Power Sources 195 (2010) 2870.
[29] J.W. Fergus, J. Power Sources 195 (2010) 4554.
[30] S. Soman, Y. Iwai, J. Kawamura, J. Solid State Electrochem. 16 (2012) 176.
[31] C.-H. Lee, K.H. Joo, J.H. Kim, S.G. Woo, H.-J. Sohn, T. Kang, Y. Park, J.Y. Oh, Solid
State Ionics 149 (2002) 59.
[32] J.L. Souquet, Annu. Rev. Mater. Sci. 11 (1981) 211.
[33] St. Adams, K. Hariharan, J. Maier, Solid State Ionics 75 (1995) 193.
[34] St. Adams, K. Hariharan, J. Maier, Solid State Ionics 86/88 (1996) 503.
[35] A. Dalvi, A.M. Awasthi, S. Bharadwaj, K. Shahi, Mater. Sci. Eng. B 103 (2003) 162.
[36] N. Gupta, A. Dalvi, J. Therm. Anal. Calorim. 102 (2010) 851.
[37] B.K. Money, K. Hariharan, Solid State Ionics 179 (2008) 1273.
[38] B.K. Money, K. Hariharan, Solid State Ionics 192 (2011) 672.
[39] B.K. Money, K. Hariharan, J. Phys. Condens. Matter 21 (2009) 115102.
[40] R.S. Gedam, V.K. Deshpande, Solid State Ionics 177 (2006) 2589.
[41] J.L. Nowiński, M. Mroczkowska, J.R. Dygas, J.E. Garbarczyk, M. Wasiucionek, Solid
State Ionics 176 (2005) 1775.
[42] K. Moricova, E. Jóna, A. Plško, S.C. Mojumdar, J. Therm. Anal. Calorim. 100 (2010)
817.
[43] A. Ghosh, A. Pan, Phys. Rev. Lett. 84 (2000) 2188.
[44] G. Govindaraj, C.R. Mariappan, Solid State Ionics 147 (2002) 49.
[45] B.D. Cullity, Elements of X-ray Diffraction, 2nd ed. Addison-Wesley Publishing
Company Inc., Reading, MA, USA, 1978.
Solid State Ionics 263 (2014) 119–124
Contents lists available at ScienceDirect
Solid State Ionics
journal homepage: www.elsevier.com/locate/ssi
Effect of conditional glass former variation on electrical transport in
Li2O–P2O5 glassy and glass-ceramic ionic system
Munesh Rathore, Anshuman Dalvi ⁎
Physics Department, Birla Institute of Technology and Science, Pilani, RJ 333031, India
a r t i c l e
i n f o
Article history:
Received 11 March 2014
Received in revised form 21 May 2014
Accepted 23 May 2014
Available online 11 June 2014
Keywords:
Ionic conductivity
Crystallization
Li+ ion glasses
Glass-ceramics
a b s t r a c t
Giving emphasis to electrical transport in the thermally unstable region, a conditional glass former based system 50Li2O–(50-x)P2O5–xMoO3 is investigated. Though glass forming region is narrow, the electrical conductivity exhibits significant rise up to x ≤ 15 mol%. Scanning electron microscopy investigations suggest existence of
tiny crystallites well separated by glass tissues for higher MoO3 content samples. It is therefore revealed that addition of MoO3 improves the thermal stability. Electronic conductivity in this system is found to be fairly low and
suggests phonon assisted polaron hopping. Electrical conductivity is found to be comparable to glass and glassceramic samples.
© 2014 Elsevier B.V. All rights reserved.
1. Introduction
Lithium ion sulfide (Li2S based) and oxide (Li2O based) glasses have
drawn considerable attention in the last three decades due to their potential candidature as solid electrolytes [1–7]. Lithium ion oxide glasses
though exhibit better thermal stability, their ionic conductivity at room
temperature is generally poor. On the other hand, Li+ ion sulfide glasses
exhibit high ionic conductivity near room temperature, however, their
thermal stability is poor and applicability to low dimensional ionic
devices is always questionable [3–6].
In view of possible applications to thin film as well as high temperature solid ionic devices, Li2O based glasses are promising, provided that
ionic conductivity could be significantly improved. Therefore, mainly
two approaches have been used to increase the ionic transport, viz.,
(i) addition of salt to provide mobile ions and (ii) use of mixed glass
formers. Salt addition (Li2SO4, LiCl etc.) [7–9] leads to an increase in
the number of charge carries, whereas, use of mixed glass formers
(e.g. P2O5–B2O3) essentially increases the free volume of the glass matrix
[11–23]. Typical glass formers e.g. SiO2, P2O5, and B2O3form a glassy state
naturally, following the Zachariasen rules [10]. However, there is another type, viz., conditional glass formers (transition metal based oxide
e.g. V2O5, MoO3) which can form a glassy state only in the presence of
other compounds [10]. Use of one of these conditional glass formers
certainly enhances the ionic conductivity. Even though, their addition
also incorporates electronic conductivity surprisingly. This is essentially
⁎ Corresponding author at: Physics department, BITS Pilani, RJ 333031, India. Tel.: +91
1596 515640; fax: +91 1596 244183.
E-mail address: [email protected] (A. Dalvi).
http://dx.doi.org/10.1016/j.ssi.2014.05.018
0167-2738/© 2014 Elsevier B.V. All rights reserved.
due to the tendency of these oxides to loose oxygen during molten state
that results into lower valence state for some of the transition metal
ions. This further leads to availability of electrons in the glass matrix.
Therefore, in view of search of novel glassy and glass-ceramic cathode
materials, conditional glass former based systems have got importance
[11–14].
Numerous studies are available on the mixed glass former effect
using substitution of conditional glass formers with emphasis on electrical
transportand structural and thermal aspects. For example, Lee et al. reported the enhancement in the ionic conductivity with the addition of
V2O5 into the binary Li2O–B2O3 glassy system at constant Li2O content
[15]. It was demonstrated by Jozwiak et al. [16] that the electrical conductivity of Li2O–V2O5–P2O5 glasses changes from predominant electronic to
ionic with compositional variations. Recently, Gedam et al. observed a rise
in the electrical conductivity in Li2O–B2O3–V2O5 glassy system with addition of V2O5[17]. Further, Dabas et al. [18] observed that addition of
Nb2O5 on binary lithium phosphate glass increases the thermal stability.
Effect of MoO3 addition in Li2O–P2O5–MoO3 glassy system was also
studied by Chowadri et al. [19], but mainly in the thermally stable region (T ≪ Tg). Such a substitution leads to enhancement in electronic
as well as ionic conductivity. In another investigation, Montani et al.
studied Li2O–V2O5–MoO3–TeO2 glassy system and revealed separate regions for predominant electronic and ionic transports[20]. It was found
that addition of Bi2O3 in Li2O–P2O5 glassy matrix leads to an enhancement in the electrical conductivity and ionic nature of the system due
to the presence of Li+ ions and BiO6 octahedra in the glass matrix
[21]. Similarly, it was also found that incorporation of SeO2 into binary
Li2O–B2O3 glassy matrix leads to significant enhancement in electrical
conductivity [22].
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M. Rathore, A. Dalvi / Solid State Ionics 263 (2014) 119–124
In spite of several important studies on conditional modifier based
systems, there are certain aspects which are either missing in the literature or have not gotten attention. For example, (i) crystallization and
its effect on electrical transport, (ii) mechanism of glass-ceramic formation and (iii) estimation of comparative electronic transport in glasses
and glass-ceramics with compositional alterations. These studies are
important in view of development of solid electrolytes for new generation solid ionic devices. Therefore, the present study emphasizes the
effect of MoO3 substitution in Li2O–P2O5 matrix on thermal, ionic and
electronic transports in Li2O–P2O5–MoO3. Thus the glass and glassceramic composites in this system are investigated.
2. Experimental
The glassy samples of the composition 50Li2O–(50-x)P2O5–xMoO3
(x = 5–20 mol%) were prepared by conventional melt quenching. A
sample with x = 0 i.e. 50Li2O–50P2O5 is abbreviated as LP, whereas,
compositions x = 5–20 as 5–20MLP, respectively. The glassy/
amorphous nature of the samples was examined by X-ray diffraction
(Rigaku, MiniFlex II). Further, to investigate the thermal events, differential scanning calorimetry (DSC, Simadzu DSC-60) measurements were
performed on the as prepared glassy samples. Field emission type scanning electron microscopy (FEI-Nova NanoSEM 450) was performed on
the annealed flakes of glasses to visualize the microstructural changes
during crystallization.
The melt-quenched glassy samples were thoroughly ground and
pelletized. The cylindrical shaped pellets, with conductive graphite
paint pasted on both sides, were used for conductivity measurements.
Impedance spectroscopy was carried out in the frequency range
42 Hz–5 MHz using HIOKI 3532-50 LCR meter. Further to avoid the expected water absorption by the samples which may lead to the proton
conduction, samples were annealed prior to each measurement. To
study electronic transport dc polarization technique was used. To determine the electrochemical stability of these glassy samples, cyclic
voltammetry (CV) was performed using Princeton 263A potentiostat/
galvanostat on cells of type C|glass|C.
3. Results and discussion
3.1. X-ray diffraction (XRD)
The XRD patterns of as prepared glassy and glass-ceramic samples
for x = 5 (5MLP) and 15 (15MLP) are shown in Fig. 1. The absence of
significant peaks in case of both the samples (Fig. 1a and c) confirms
Fig. 1. XRD patterns for glass and glass-ceramic samples: (a) 5MLP as prepared glass;
(b) 5MLP glass-ceramic; (c) 15MLP as prepared glass; and (d) 15MLP glass-ceramic sample.
their glassy nature. For x N 15 mol% the samples are found to be partially
crystalline (Fig. 1a) in nature. Thus the glass forming region is found to
be narrow.
The samples were further annealed above their respective crystallization temperatures (as later revealed by DSC) and slowly cooled. The
XRD patterns for these glass-ceramic samples are shown in Fig. 1b
and d. Appearance of tiny peaks strongly suggests crystalline domains
of LiPO3 and possibly another compound Li(MoP2O5) embedded in the
glass matrix. The crystallite size for 5MLP and 15MLP is found to be in
the range of 30–70 nm and 20–40 nm, respectively. Estimation of area
under the peaks suggests that in case of 5MLP, precipitation of LiPO3 is
in appreciable amount, whereas, in case of 15MLP, crystallization of
Li(MoP2O8) is facilitated and that of LiPO3 is suppressed. Therefore, it
may be suggested that (i) the addition of MoO3 suppresses the crystallization of LiPO3 and (ii) samples are complex in nature.
Further to understand the sequence of precipitation of compounds
during crystallization, the samples were annealed at two distinct temperatures of 400 °C and 500 °C and slowly cooled to room temperature.
The XRD results for these samples are shown in Fig. 2. On annealing at
400 °C, significant peaks corresponding to Li(MoP2O8) grow as shown
in Fig. 2a, hence the first major precipitation confirmed is most likely of
Li(MoP2O8) compound. On further heating, at 500 °C, area under the
peaks corresponding to LiPO3 increases, along with subtle increase for
peaks corresponding to Li(MoP2O8). Thus the first (at low temperature)
and second (at high temperature) crystallizations may correspond to
major precipitation of Li(MoP2O8) and LiPO3 compounds, respectively.
3.2. Differential scanning calorimetry (DSC)
Further, to confirm the glassy nature of as prepared melt-quenched
samples, differential scanning calorimetry (DSC) scans (10 °C/min)
were performed on LP, 5MLP and 15MLP samples (Fig. 3). The appearance of endothermic smooth dip followed by exothermic peak corresponds to the glass transition temperature (Tg) and crystallization
temperatures (Tc), respectively.
In case of LP glassy sample (Fig. 3a), the Tg appears at 320 °C followed
by a board exothermic peak with onset at 365 °C. As investigated previously, this peak may correspond to massive crystallization of LiPO3[24].
Whereas in case of 5MLP and 15MLP (Fig. 3 b and c) the crystallization
peaks are apparently merged-up, board and spanning over ~ 60 and
90 °C, respectively. These merged up peaks suggest at least two crystallizations. The first and second exothermic peaks may correspond to major
precipitation of (i) Li(MoP2O8) and (ii) LiPO3, respectively as suggested
Fig. 2. XRD patterns for the 15MLP glassy sample after annealing at (a) 400 °C and
(b) 500 °C.
M. Rathore, A. Dalvi / Solid State Ionics 263 (2014) 119–124
Fig. 3. DSC scans at a heating rate of 10 °C/min for the glassy samples.
by XRD results. Interestingly, the Tg and Tc values shift towards higher
temperatures which further suggests improved thermal stability with
MoO3 addition.
3.3. Electrical transport
To understand the effect of MoO3 addition on electrical transport in a
complete temperature range, the σ–T cycles (1 kHz) (Fig. 4) were
carried out on glass as well as glass-ceramic samples, firstly at a controlled heating rate of 1 °C/min essentially to observe thermal events.
Since in the whole region of temperature, the conductivity at 1 kHz
lies in the plateau region of σ–ω plot [8], it is assumed as dc conductivity
Fig. 4. Electrical conductivity–temperature cycles at a heating rate of 1 °C/min for the
samples' symbols denoting: (●) glass (G) and (○) glass-ceramics (GC) samples.
121
exhibiting bulk response. Fig. 4a, c and e represents the σ–T cycles for
pristine glassy compositions. In the thermally stable region (i.e. below
Tg), the conductivity shows reversible Arrhenius behavior for all the
three compositions. Thus the samples are thermally stable at least up
to ~330 °C. Whereas, in the thermally unstable region (T ≥ Tg), a deviation from linearity at Tg is observed that may be due to an increase in
free volume of the glass matrix, as also seen in Ag+ and Li+ ion oxide
glasses [25]. Such a deviation is more prominent for low MoO3 content
samples. On further heating, as the temperature approaches Tc, crystallization begins and the σ–T cycles exhibit two kinks as shown in Fig. 4.
Firstly, the fall in the conductivity is more prominent for second kink
(Tc2) which may correspond to crystallization of LiPO3. Secondly, such
a fall in the conductivity at Tc2 is more apparent and massive in low
MoO3 content sample and gradually becomes less prominent for samples with high MoO3 content (15MLP). Thus it is once again evident
that MoO3 stabilizes the glassy phase effectively suppressing the crystallization of LiPO3.
It may be further suggested that during crystallization also the sample
remains predominantly ionic. In a previous study, on a novel salt free
V2O5–P2O5 glassy system [12], a sudden rise in the conductivity was
witnessed during the crystallization due to predominately electronic
conduction.
Fig. 4b, d and f represents the σ–T cycles obtained for the corresponding glass-ceramic samples. The conductivity of these samples
exhibits an Arrhenius behavior with no thermal events in the whole
temperature range, except neat Tc2 above which the conductivity increases with a little higher activation energy. Such behavior further suggested better thermal stability of glass-ceramics. It may be emphasized
here that the conductivity of the glass-ceramics as compared to glasses
at 100 °C is low only by few factors, whereas, in our previous investigation on Li2SO4–Li2O–P2O5[8] system the conductivity in glass-ceramics
was found to be significantly low (at least ~ 2 orders of magnitude)
than that of the corresponding glasses.
Further, the conductivity increases with the addition of MoO3 for
both glasses as well as glass-ceramic samples, and as expected opposite
Fig. 5. Activation energy and conductivity as function of MoO3 content for (a) glasses and
(b) glass-ceramics.
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M. Rathore, A. Dalvi / Solid State Ionics 263 (2014) 119–124
trend is seen in the activation energy (Fig. 5). For x ≥ 15 mol% conductivity drops that may be due to partially crystalline nature of the samples.
Results are in good agreement with a previous study on xMoO3–(1-x)
[0.5Li2O–P2O5] glassy system [19]. However, in the present case Li+ ion
content is fixed and MoO3 substitutes P2O5. Thus the thermally stable
ionic system is obtained, without compromising the Li+ ion content.
As suggested earlier [19] the enhancement in conductivity in the
glassy phase may be attributed to substitution of a relatively bigger
ion (Mo in place P) that, in turn, leads to an increase in free volume of
the glass matrix. The reason for conductivity enhancement in glassceramic due to MoO3 substitution is further scrutinized. To understand
the glass-ceramic formation and its correlation with electrical transport,
scanning electron microscopy (SEM) is performed on the samples 5MLP
and 15MLP annealed well above crystallization temperature and the results are shown in Fig. 6. In both the samples, glass-ceramic formation is
evident. The low MoO3 content sample (5MLP) is evidently a dense
glass-ceramic with relatively larger average grain size of ~500 nm due
to which massive crystallization in this sample is observed in σ–T cycle.
On the other hand, the 15MLP sample is found to be consisted of well separated, relatively smaller grains of average size ~50–100 nm surrounded
by connective glass tissues. The SEM results may be correlated with electrical transport in the glass-ceramics. It may be suggested that the massive fall of the conductivity during the crystallization seen in case of
5MLP may be attributed due to unavailability of connective glass tissues
due to dense glass-ceramic formation, schematically also shown in
Fig. 6c. Thus a relatively poor conductivity of crystallites results into a
drastic fall during crystallization [26]. Such a fall is not prominent for
15MLP due to available glass tissue due to smaller crystallite size. As
shown in Fig. 6d, ions can bypass the crystallites and move comfortably
through these tissues during electrical transport. These connective tissues
are stable and thus 15MLP glass-ceramic is more glass-like and exhibits
higher conductivity. Thus the SEM results are in good agreement with
those of the σ–T cycles and also compliment the X-ray results.
It is relevant to mention here that in some phosphate and silicate
glasses [27] there is a possibility of proton conduction especially when
measurements are done in open atmosphere. Since these samples are
prepared by melt-quenching route and measurements are performed
above 150 °C, such possibility is not expected in the present case. To estimate the contribution of ions in total electrical transport impedance
spectroscopy measurements (42 Hz–5 MHz) were carried out at different temperatures and for different compositions. The Nyquist or cole–
cole plots (Z″ vs Z′) obtained for 15MLP glassy sample are shown in
Fig. 7a. Further, the same plot is also shown for different compositions
at temperature 250 °C (Fig. 7b). In both cases, the plot exhibits depression semicircles followed by an inclined line that may be represented by
parallel combination of resistance (R) and constant phase element
(CPE1) for bulk in series with another element CPE2 (interface). The
low frequency inclined line is most likely due to polarization at the
electrode–electrolyte interface and strongly suggests predominantly
ionic nature of the samples. The depressed semicircles represent the
bulk response and suggest distribution of relaxation times.
To further estimate the electronic contribution to the conductivity, a
dc potential (1 V) was applied to the sample and transient current (It) as
function of time was measured (Fig. 8). As apparent, initially the transient current drops rapidly and saturates subsequently to a notably
small value (Is). For all the samples, probably due to faster response of
Li+ ions to the applied electric field, it was not possible to measure Io
(current at exactly t = 0) accurately. Nevertheless, for glass as well as
glass-ceramic samples apparently a drastic fall in the transient current
confirms their predominant ionic nature. In the present case (Fig. 8),
the saturation current (Is) is essentially due to electronic contribution.
Thus the electronic conductivity for the sample is obtained by the following relation:
σ¼
lI s
V pA
ð1Þ
where Vp is the applied voltage, l the thickness and A is cross-sectional
area of the sample. To examine temperature dependence of the electronic transport, the electronic conductivity was carefully measured in
the thermally stable region. For one of the samples i.e. 15MLP the temperature dependence of conductivity is shown for glass (Fig. 9a) and its
Fig. 6. SEM images (2 μm resolution) of glass-ceramic samples for (a) 5MLP and (b) 15MLP. Schematic representation of possible motion of ions in: (c) 5MLP and (d) 15MLP. The arrow
indicates likely pathways for ions after crystallization.
M. Rathore, A. Dalvi / Solid State Ionics 263 (2014) 119–124
123
Fig. 9. Total and electronic conductivity vs inverse of temperature for 15MLP in the
thermally stable region for (a) glass and (b) glass-ceramic.
Fig. 7. Nyquist plots for (a) 15MLP glassy sample at three different temperatures and
(b) for different compositions at a fixed temperature.
corresponding glass-ceramic (Fig. 9b) sample. Interestingly, the conductivity increases exponentially with temperature for both cases.
The electronic transport in disordered solids can be explained using
a theory proposed by Mott [28,29]. According to this model which was
essentially proposed for semiconducting glasses, for very low temperatures T ≪ θD (θD = debye temperature) the contribution of phonons
is essentially suppressed and electronic transport follows the well
Fig. 8. Transient current (It) as function of time for (a) glass and (b) glass-ceramic samples.
known Mott's T1/4 law which is also known as variable range hopping
mechanism.
On the other hand, when temperature exceeds θD, significant phonon contribution is predominant and conductivity follows Arrhenius
dependence of temperature according to the following equation:
E
− kTp
σ ¼ σ 0e
:
ð2Þ
The activation energy (Ep) in the above expression corresponds to
polaron hopping and is mainly due to two important contributions,
viz., (i) the binding energy for polaron formation and (ii) the energy difference in the initial and final states due to the difference in the ion
coordination.
Arrhenius type conductivity behavior over a wide range of temperature in the present case strongly suggests phonon assisted polaron hopping as also suggested in earlier investigations [14,16] in the glass and
glass-ceramics.
It may be emphasized that in similar to previous investigations [19]
MoO3 addition increases the electronic conductivity notably (Fig. 9).
The electrons are introduced to the system from MoO3 during annealing
of the sample in its molten state. These electrons contribute to electrical
transport through hopping conduction from a site near Mo+5 to a Mo+6
ion. Thus electrons behave like a charge placed in polarizable medium
and the induced polarization surrounding the electrons accompanies
electrons during transport. The activation energy for polaron hopping,
calculated from the slope, is found to be slightly higher for glassceramic (0.61 eV) than that of glassy sample (0.50 eV) for 15MLP. This
further suggests that the glassy state appears to be more favorable for
the polaron hopping. Furthermore, the nano-crystallites that precipitated during crystallization hardly facilitate electronic conduction. It
should also be emphasized that the total activation energy (Fig. 5)
which is predominantly due to ionic contribution is found comparable
to that of polaron hopping. This further suggests that the transport of
Li+ ions and polaron is equally facilitated in the matrix.
Previously in 15Li2O–15P2O5–70MoO3 system [18], predominately
electronic nature was seen due to high MoO3 content. In the present
124
M. Rathore, A. Dalvi / Solid State Ionics 263 (2014) 119–124
content samples. The glass and glass-ceramic samples are predominately ionic. The CV measurements confirm an appreciable stability even at
higher temperatures and compliment the ionic nature of these samples.
The present investigation reveals that highly conducting glass-ceramic
with prominent ionic nature can be obtained using conditional glass
former substitution as well.
Acknowledgments
This work is supported by DST-FIST (SR/FST/PSI-150/2010)
and UGC Special Assistance Programme (F-530/3/DRS/2009/SAP-I),
Government of India. Authors would like to thank the late Professor
Suresh Chandra of Banaras Hindu University, India for his suggestions
and fruitful discussions.
References
Fig. 10. Cyclic voltammograms for the glassy samples for two compositions at 200 °C
(a) 10MLP and (b) 15MLP.
work, due to low MoO3 content ionic transport predominates and electronic conductivity is notably low.
3.4. Cyclic voltammetry
Cyclic voltammograms (CV) at 200 °C for two of the glassy compositions, viz., 10MLP and 15MLP, sandwiched between two inert graphite
(C) electrodes, are shown in Fig. 10. Unlike systems in which existence
of large amount of conditional glass former contributes to pure ohmic
behavior of CV curves, a hysteresis is observed which may be attributed
to accumulation of charge carries at electrode–electrolyte interface [30].
As apparent, the behavior of current is not abrupt but very systematic
that further indicates a stability window of at least ±2 V even at higher
temperatures. These samples therefore are promising candidates as
solid electrolyte for high temperature ionic device applications.
4. Conclusions
Addition of MoO3 in the Li2O–P2O5 glass matrix reveals interesting
results. Such a substitution in place of P2O5 improves the thermal stability as well as ionic conductivity, in turn confirmed by SEM, XRD, DSC
and electrical transport. Interestingly, for these samples, the conductivity of glass and corresponding glass-ceramic is almost comparable. This
is essentially due to suppressed crystallization that leads to availability
of connective glass tissues, as more evidently seen for higher MoO3
[1] P. Knauth, Solid State Ionics 180 (2009) 911.
[2] M. Park, X. Zhang, M. Chung, G.B. Less, A.M. Sastry, J. Power Sources 195 (2010)
7904.
[3] K. Otto, Phys. Chem. Glasses 7 (1966) 29.
[4] K. Minami, F. Mizuno, A. Hayashi, M. Tatsumisago, Solid State Ionics 178 (2007) 837.
[5] H.L. Tuller, D.P. Button, D.R. Uhlmann, J. Non-Cryst. Solids 40 (1980) 93.
[6] K. Homma, M. Yonemura, T. Kobayashi, M. Nagao, M. Hirayama, R. Kanno, Solid
State Ionics 182 (2011) 53.
[7] M. Ganguli, K.J. Rao, J. Non-Cryst. Solids 243 (1999) 251.
[8] M. Rathore, A. Dalvi, Solid State Ionics 239 (2013) 50.
[9] M. Ganguli, M.H. Bhat, K.J. Rao, Solid State Ionics 122 (1999) 23.
[10] W.H. Zachariasen, J. Chem. Soc. 54 (1932) 3841.
[11] T.K. Pietrzak, M. Wasiucionek, I. Gorzkowska, J.L. Nowiński, J.E. Garbarczyka, Solid
State Ionics 251 (2013) 40.
[12] T.K. Pietrzak, J.E. Garbarczyk, M. Wasiucionek, I. Gorzkowska, J.L. Nowinskia, S.
Gierlotk, Solid State Ionics 192 (2011) 210.
[13] K. Nagamine, T. Honma, T. Komatsu, J. Power Sources 196 (2011) 9618.
[14] J.E. Garbarczyk, M. Wasiucionek, P. Jozwiak, J.L. Nowinski, C.M. Julien, Solid State
Ionics 180 (2009) 531.
[15] Y.I. Lee, J.H. Lee, S.H. Hong, Y. Park, Solid State Ionics 175 (2004) 687.
[16] P. Jozwiak, J.E. Garbarczyk, Solid State Ionics 176 (2005) 2163.
[17] R.S. Gedam, V.K. Deshpande, Bull. Mater. Sci. 32 (2009) 83.
[18] P. Dabas, K. Hariharan, Solid State Ionics 225 (2012) 641.
[19] B.V.R. Chowadri, K.l. Tan, W.T. Chia, R. Gopalakrishnan, Solid State Ionics 40/41
(1990) 684.
[20] R.A. Montani, M.A. Frechero, Solid State Ionics 177 (2006) 2911.
[21] S. Rani, S. Sanghi, A. Agarwal, N. Ahlawat, J. Mater. Sci. 44 (2009) 5781.
[22] C.H. Lee, K.H. Joo, J.H. Kim, S.G. Woo, H.-J. Sohn, T. Kang, Y. Park, J.Y. Oh, Solid State
Ionics 149 (2002) 59.
[23] L. Bih, M.E. Omar, J.M. Reau, M. Hadded, D. Boudlich, A. Yacoubi, A. Nadiri, Solid State
Ionics 132 (2000) 71.
[24] B.K. Money, K. Hariharan, J. Phys. Condens. Matter 21 (2009) 115102.
[25] J.L. Souquet, Annu. Rev. Mater. Sci. 11 (1981) 211.
[26] N. Gupta, A. Dalvi, Indian J. Pure Appl. Phys. 51 (2013) 328.
[27] M. Nogami, Y. Abe, Phys. Rev. B 55 (1997) 12108.
[28] N.F. Mott, Philos. Mag. 19 (1969) 853.
[29] I.G. Austin, N.F. Mott, Adv. Phys. 18 (1969) 41.
[30] P. Machowski, M. Opallo, J.E. Garbarczyk, M. Wasiucionek, Solid State Ionics 157
(2003) 287.
Journal of Non-Crystalline Solids 402 (2014) 79–83
Contents lists available at ScienceDirect
Journal of Non-Crystalline Solids
journal homepage: www.elsevier.com/ locate/ jnoncrysol
Crystallization in Li2SO4–Li2O–P2O5 glassy ionic system: An assessment
through electrical transport
Munesh Rathore, Anshuman Dalvi ⁎
Department of Physics, Birla Institute of Technology Science, Pilani, RJ 333031, India
a r t i c l e
i n f o
Article history:
Received 27 January 2014
Received in revised form 30 April 2014
Accepted 12 May 2014
Available online xxxx
Keywords:
Ionic glasses;
Crystallization;
Thermal stability;
Ionic conductivity
a b s t r a c t
The electrical transport during the crystallization of ternary Li2SO4–Li2O–P2O5 glassy ionic system reveals interesting results. Electrical conductivity isotherms recorded at the crystallization temperature are found to be of immense importance, especially, in the understanding of phase transformation. The electrical conductivity falls
rapidly during crystallization of LiPO3 compound and saturates at its completion. Such studies on conductivity
isotherms further confirm the suppression of crystallization with addition of Li2SO4 in the glass matrix. A systematic variation of the Cole–Cole plots and conductivity spectra during crystallization further suggests a slow and
predetermined process of glass–ceramic formation. The mechanism of electrical transport during crystallization
is also explained using a crystallite bypass model. It is further revealed that the samples remain purely ionic during and after crystallization.
© 2014 Elsevier B.V. All rights reserved.
1. Introduction
Fast ionic glasses have wide applications as electrolytes in solid ionic
devices due to their liquid like structure, high thermal stability, wide
glass forming region and negligible electronic conductivity [1–3]. Consequently, the last one decade has witnessed considerable attention
on all solid state batteries based on Li+ ion glassy electrolytes [3–5].
In spite of relatively wide thermal stability window, ionic glasses still
have a tendency to crystallize above a certain temperature (Tc) and
hence their direct use may affect the performance of the ionic device,
especially during high temperature applications [6,7]. Therefore, in
view of better expected thermal stability, ion conducting glass–ceramic
composites attracted the solid state ionics community [5–13]. In general, the glass–ceramics are produced by controlled crystallization of
glasses and exhibit lower conductivity than the parent glasses owing to
poor ionic conductivity of the precipitated crystallites. Nevertheless, in
many cases, they are found to be exhibiting higher conductivity than
the pristine glasses of the same compositions. For example, in Ag+ ion
conducting AgI–Ag2O–MoO3 glasses, repeated heating–cooling cycles
near the crystallization temperature (Tc) lead to the formation of highly
conducting embedded superionic crystals in the glass matrix [9]. In another work, heat treatment on Li+ ion conducting 80Li2S–20P2S5 glassy powders at around 220 °C leads to the formation of dense glass–ceramics, viz.
Li7PS6 and Li3PS4, with appreciable high Li+ ion conductivity of 10−3
Ω−1 cm−1 at 300 K [10]. Furthermore, oxide glass–ceramics obtained
from the Li2O–Al2O3–TiO2–P2O5 [11] glassy system were found to be
⁎ Corresponding author at: Physics Department, BITS Pilani, RJ 333031, India. Tel.: +91
1596 515640; fax: +91 1596 244183.
E-mail address: [email protected] (A. Dalvi).
http://dx.doi.org/10.1016/j.jnoncrysol.2014.05.012
0022-3093/© 2014 Elsevier B.V. All rights reserved.
exhibiting better conductivity than the host glass matrix. Such a notable
high conductivity is attributed to formation of the LiTi2(PO4)3 (LTP)
phase during crystallization [11]. Efforts have also been to stabilize high
temperature superionic α-phase of various ionic compounds in the
glass matrix. The very first successful attempt was stabilization of α-AgI
in the AgI–Ag2O–B2O3 glassy system by rapid quenching [12]. Exceptionally high Ag+ ion conductivity of 10−1 Ω−1 cm−1 was achieved at room
temperature in these α-AgI frozen glass–ceramics.
Since performance of the ionic device is strongly correlated with the
ionic conductivity of the electrolyte, the electrical transport in these
glasses has been widely studied, but mostly below glass transition temperature, i.e., T b Tg [14–20]. To develop better glass–ceramics in terms
of ionic conductivity and thermal stability, focus subsequently shifted
to the thermally unstable region of the glasses. There are few interesting
investigations in which the conductivity (σ)–temperature (T) cycles
were studied above the glass transition (Tg) and even crystallization
(Tc). These studies have developed immense understanding of various
interesting phenomena, e.g. phase transformation, structural relaxation
at Tg and effect of precipitation of new compounds on the ionic conductivity of various samples during crystallization. Souquet et al. [20,21] for
the first time analyzed the electrical conductivity of some Li+, Na+ and
K+ ion based oxide binary glasses above the Tg and found a significant
deviation from Arrhenius behavior. The unusual rise in conductivity at
Tg was attributed to increase in the free volume due to structural relaxation. Adems et al. [22,23] systematically measured the σ–T cycles in the
unstable region for a series of AgI–oxysalt glassy systems and observed
apparent thermal events at Tg and Tc. In this study the anomalous rise in
the conductivity at Tg was attributed to glass–crystallite interface effect,
whereas, the sudden and significant fall in its value to crystallization
(Tc). Such thermal events in the σ–T cycles were also observed in
80
M. Rathore, A. Dalvi / Journal of Non-Crystalline Solids 402 (2014) 79–83
mechanochemically synthesized Ag+ ion oxide glasses [24–27]. In a
recent study, it was demonstrated that substitution of CuI in CuIx–AgI1
− x–Ag2O–V2O5 [28] superionic glasses leads to suppression of crystallization which was in turn confirmed during σ–T cycles in the thermally
unstable region. Further, in the Li+ ion conducting Li2O–P2O5 glassy system, crystallization of LiPO3 [29] was found to affect σ–T cycles
significantly.
For a better understanding of the conductivity–structure correlationship, and in order to develop new generation fast ionic and
thermally stable glass–ceramics, understanding of crystallization is
inevitable. Therefore, the present work on fundamental glassy
ionic system is undertaken. The effect of structural changes during
crystallization (especially at T = Tc) on the electrical transport in
Li2SO4–Li2O–P2O5 fast ionic glasses is investigated.
2. Experimental
The system xLi2SO4–(100 − x) (0.5Li2O–0.5P2O5) for two important
compositions, having low and high content of Li2SO4, viz. x = 10 and
60 mol.% of Li2SO4 were chosen and abbreviated as 10LSLP and
60LSLP, respectively. The high-purity raw materials (Li2CO3,
NH4H2PO4, and Li2SO4) in appropriate amount were taken in alumina
crucibles and heated to ~ 500 °C for 2 h (to remove CO2, NH3 and
H2O). The mixture was then melted at 900–950 °C for 30 min and
then subsequently pressed between two copper plates to obtain transparent glassy flakes [30]. The melt-quenched glassy films were ground
to fine powder for further studies and pelletized in a steel die for the
electrical characterization. To achieve better electrode–electrolyte contact, high purity graphite paint was uniformly pasted on both sides of
the cylindrical pellets. The electrical conductivity (σ)–temperature (T)
cycle measurements were carried out using a computer interfaced
HIOKI 3532-50 LCR meter and a programmable PID temperature controller (Libratherm PRC 309). The electrical conductivity isotherms
(42 Hz–5 MHz) were measured by keeping the sample carefully near
the onset of the crystallization temperatures (T–Tc). Further, differential
scanning calorimetry (Shimadzu, DSC-60) was performed to identify
the thermal events.
3. Results and discussion
Fig. 1. Electrical conductivity (at 1 kHz) versus inverse of temperature obtained at 1 °C/min
for (a) 10LSLP and (c) 60LSLP. For comparative study DSC scans are also shown.
conductivity exhibits a significant deviation from Arrhenius behavior at
glass transition (Tg) which is most likely due to an increase in the free volume due to structural relaxation [21]. For 10LSLP the electrical conductivity exhibits a significant fall at temperature Tc that is attributed to a major
precipitate of LiPO3 into the glass matrix. Further for the 60LSLP sample at
temperature Tc1 instead of a drop, the conductivity exhibits a plateau-like
region and a relatively less prominent fall only at Tc2, which in turn confirms the suppressed crystallization of LiPO3.
To further understand the crystallization and its effect on electrical
transport, complete heating and cooling σ–T cycles at 1 kHz were also
obtained for the 10LSLP and 60LSLP samples (Fig. 2). The conductivity,
as expected, exhibits Arrhenius behavior with reversibility for T ≤ Tg
3.1. Thermal stability and electrical transport
At the outset, as-prepared samples x ≤ 60 mol.% are found to be essentially glassy in nature as suggested by X-ray measurements reported
in our previous work [30]. As discussed previously, XRD patterns for
glass–ceramic samples confirm the major precipitation of LiPO3 (according to the reaction Li2O + P2O5 → 2LiPO3) [30]. However, it is
also seen that addition of Li2SO4 in the glass matrix suppresses the crystallization of LiPO3 and facilitates that of Li2SO4 itself. Further, for high
Li2SO4 content, major precipitation of Li2SO4 at Tc1 and that of LiPO3 at
Tc2 is confirmed [30]. Further, in the samples with low Li2SO4 content,
precipitation of LiPO3 is predominant during crystallization [30].
The DSC scans (Fig. 1a and c) along with conductivity (σ)–
temperature (T) cycles obtained at 1 kHz (Fig. 1b and d) are plotted to observe and correlate the thermal events. As seen in the DSC plots for
10LSLP, the Tg appears at ~310 °C followed by a broad exothermic peak
with a span of 140 °C, which corresponds to major precipitation of
LiPO3 compound. On the other hand, in the case of 60LSLP, Tg appears at
a slightly lower temperature of ~260 °C followed by two well separated
exothermic crystallization peaks. At ~288 °C (Tc1) a sharp and significant
major peak is witnessed in addition to the similar board peak with onset
at ~370 °C. These two temperatures may correspond to precipitation of
Li2SO4 and LiPO3, respectively [30]. It is further noted from our previous
results [30] and the present DSC patterns that the crystallization of
Li2SO4 is facilitated and that of LiPO3 is suppressed when Li2SO4 is
added to the glass matrix. Apparently, for both the samples the electrical
Fig. 2. Temperature dependence of electrical conductivity (at 1 kHz) during heating and
cooling cycles for (a) 10LSLP and (b) 60LSLP.
M. Rathore, A. Dalvi / Journal of Non-Crystalline Solids 402 (2014) 79–83
81
Fig. 3. Crystallite bypass model (assuming growth dominated mechanism) to understand the conductivity behavior during crystallization.
for both cases. Above the crystallization (T N Tc) i.e. after completion of
the crystallization, the behavior of conductivity is again found to be
Arrhenius that further suggests the ionic nature of the glass–ceramic
system. From T ≫ Tc the samples were gradually cooled to room temperature and cooling cycle is obtained. The value of room temperature
conductivity for the 10 and 60LSLP samples is found to have dropped
by a factor of ~12 × 103 and ~4 × 103, respectively, due to crystallization.
As evident, the σ–T cycle was found to be irreversible and exhibited typical Arrhenius behavior in complete cooling cycle as well. Moreover, no
thermal event is witnessed during cooling that confirms appreciable
stability of the glass–ceramics.
The mechanism of electrical transport during crystallization can be
understood by a crystallite bypass model previously proposed for fundamental AgI–oxysalt glasses [32]. According to which, at a temperature
well above Tg, stable nucleation begins and tiny crystallites precipitate
in the glass matrix (Fig. 3a). The Li+ ions contributing to the conductivity still comfortably move through the glass matrix through the glass tissue. However, when the glass is further heated to a temperature ~Tc, the
size of these crystallites starts growing (Fig. 3b). The ions participating
in the conductivity still move through the glass matrix by squeezing
their surroundings. When the sample is further heated, at T ≥ Tc due
to significant growth of the size of crystallites, the ion pathways in the
glass matrix are eventually blocked (Fig. 3c). The volume fraction of
the glassy state reduces drastically. The ions are therefore bound to
move through the precipitated crystallites. The conductivity of crystallites (σXL) therefore decides the nature of electrical transport in thermally unstable region. The conductivity undergoes a drastic fall when
σXL is smaller than that of the glass matrix (σG). Whereas, if these two
are comparable, no appreciable change in conductivity should be observed during crystallization, though a deviation from Arrhenius behavior is likely due to different activation energies of ionic conduction. In
the present case, poor conductivity of LiPO3 is responsible for such a notable drop. On the other hand, conductivity of Li2SO4 may not be very
different from that of the glass matrix and thus a plateau-like region is
observed.
and its corresponding impedance/conductivity at a given time. Therefore, the following two important sources of errors are considered:
(i) Error in identifying the exact temperature of crystallization
where the precipitation of LiPO3 is massive. Thus the Tc2 is identified
with an error of ±2 °C after several careful observations.
(ii) Since the electrical transport is extremely sensitive to temperature, an inaccuracy of ±2 °C in identifying crystallization certainly
affects the accurate conductivity/impedance measurement. Thus,
an error in conductivity/impedance for each sample is calculated
and mentioned in the figure captions (Figs. 4–7).
The variation of Cole–Cole (Z′–Z′′) plots (in a wide frequency range
of 42 Hz–5 MHz) with progressive crystallization for 10LSLP glassy samples is shown in Fig. 4. Interestingly, for initial (Fig. 4a) and later times
(Fig. 4b) depressed semicircles followed by an inclined straight line
are observed that may be modeled as a parallel combination of a resistance (R) and constant phase element (CPE1) in series with another
3.2. Crystallization and electrical transport
Since the crystallization of LiPO3 affects the conductivity significantly, pristine glassy samples were isothermally kept at Tc2 and the in situ
crystallization event is studied by measuring impedance spectroscopy
parameters [31] as a function of time. The time when Tc2 is achieved is
taken to be t = 0. After each 5–10 min impedance spectrum is recorded
using a computer interfaced LCR meter. The experiments were performed for a long time (~7 h) and data for complete crystallization process was recorded. Thus large numbers of measurements were
recorded, but for clarity only few with a definite interval are shown.
It is important to mention here that the source of error in the measurements is essentially in estimation of crystallization temperature
Fig. 4. Cole–Cole plots for glassy samples during major crystallization of LiPO3 in 10LSLP at
Tc ~ 345 ± 2 °C for (a) initial and (b) later times. The impedance values at Tc are reported
with a maximum error of 5%.
82
M. Rathore, A. Dalvi / Journal of Non-Crystalline Solids 402 (2014) 79–83
element CPE2 (interface). The low frequency inclined line is most likely
due to polarization at the interface. Further, the interface polarization
initially may be at the electrode–electrolyte boundary; at later times,
however, the polarization may also build up at the glass–LiPO3 or
glass–Li2SO4 interface, but only after sufficient growth of crystallites
and at low frequencies. At later times (Fig. 4b), the initially rapidly
growing diameter of the semicircles almost reaches to a constant
value that may refer to completion of the crystallization.
On the other hand, the nature of the time dependence of Cole–Cole
plot for the 60LSLP is found to be dissimilar (Fig. 5). As apparent, since
the sample is more conductive due to high Li+ ion content, the high frequency semicircle in the Cole–Cole plot is almost missing from the beginning of the crystallization. In this case the system may be modeled
as a combination of bulk resistance (R1) in series with a constant
phase element (CPE1) during initial phase due to interfacial polarization. However, at later times, the slope of the incline line tends to decrease and a second semicircle at low frequencies begins to appear
that may be due to beginning of transport through the grain boundaries.
However, the expected second semicircle is not seen prominently at
later times, most probably due to (i) degraded graphite paint contact
at the interface, and (ii) experimental limitations. Thus it may be
noted that salt addition significantly affects the electrical transport during crystallization.
Bulk conductivity obtained from Cole–Cole plots (Fig. 6) is plotted as
a function of time for both the samples in Fig. 6. For both compositions,
the conductivity falls initially significantly and stabilizes subsequently,
that may be attributed to the beginning and subsequently the completion of crystallization of LiPO3. Apparently, the fall in the conductivity
is more significant at Tc2 for the sample with low content of Li2SO4
Fig. 5. Cole–Cole plots for glassy samples during crystallization of LiPO3 in 60LSLP at Tc2
~ 380 ± 2 °C for (a) initial and (b) later times. The impedance values at Tc are reported
with a maximum error of 2%.
Fig. 6. Variation of dc conductivity (obtained from Cole–Cole plots) during crystallization
of LiPO3 for (a) 10LSLP at ~345 ± 2 °C and (b) 60LSLP at ~380 ± 2 °C. The conductivity at
respective Tc is reported with maximum errors of 2% and 4% for 10 LSLP and 60 LSLP,
respectively.
that once again confirms stabilization of unstable region with salt
addition.
Jonscher's power law (JPL) predicts [33–35] the behavior of conductivity with frequency for the ionic glasses according to the following
relation.
0
σ ¼ σ dc þ Aω
n
ð1Þ
where, σdc is the dc conductivity, A is the pre-exponential constant and
n is the power law exponent, where 0 b n b 1. According to Eq. (1), the
dispersion of the conductivity at higher frequencies is also termed as JPL
behavior. To scrutinize this JPL behavior, the real part of the conductivity
(σ′) is plotted as a function of angular frequency (ω) at different crystallization times for 10LSLP and 60LSLP in Fig. 7. For 10LSLP, the conductivity is almost independent of frequency for lower values of ω followed by
a significant dispersion at higher frequencies (Fig. 7a). As the crystallization time progresses, the plateau (σdc) region shrinks and the JPL region
becomes significantly visible.
The σ–ω behavior of 60LSLP is significantly different (Fig. 7b). In this
case the deviation from the plateau region at low frequencies is more
prominent that may be attributed to interfacial polarization. It is further
evident that such a drop in the conductivity at low frequencies gradually
Fig. 7. Behavior of frequency dependence of conductivity during crystallization of LiPO3 for
(a) 10LSLP at ~345 ± 2 °C and (b) 60LSLP at ~380 ± 2 °C glassy samples. The conductivity
at respective Tc is reported with maximum errors of 2% and 4% for 10 LSLP and 60 LSLP,
respectively.
M. Rathore, A. Dalvi / Journal of Non-Crystalline Solids 402 (2014) 79–83
becomes less significant as the crystallization progresses due to suppressed interfacial polarization of free Li+ ions at the interface. Moreover, for the higher frequency region conductivity shows significant
dispersion (JPL) only at larger crystallization times.
From the above discussion it may be concluded that:
(i) Large amount of Li2SO4 stabilizes the thermally unstable region.
(ii) The overall mechanism of ionic transport during crystallization is
apparently different for low and high Li2SO4 content samples. However, the conductivity parameter variation during crystallization
process is not random but very systematic. Thus, the crystallization
process appears to be much more defined and occurs throughout
the glassy matrix uniformly.
(iii) The conductivity behavior also suggests that samples remain
completely ionic during and after crystallization.
4. Conclusions
Crystallization studies in Li2SO4–Li2O–P2O5 ionic glasses using electrical conductivity isotherms revealed interesting results. The electrical
conductivity does not exhibit any appreciable fall at crystallization of
Li2SO4 due to its comparable conductivity with that of the connective
glass tissue. On the other hand, falls in conductivity are significant during precipitation of LiPO3 due to its relatively poor conductivity. Thus,
the behavior of conductivity during the precipitation of Li2SO4 and
LiPO3 could be very well explained by crystallite bypass model. It is
also evident that addition of Li2SO4 in the Li2O–P2O5 glassy system not
only enhances the ionic conductivity, but also suppresses the crystallization. The electrical transport parameters, viz. Cole–Cole plots, dc conductivity and frequency dependence of conductivity change very
systematically during the phase transformation. This further suggests
a systematic and predetermined (not a sporadic) crystallization process.
Therefore, by controlling the crystallization via appropriate composition
alterations, degradation of ionic devices at high temperatures can be
immensely controlled.
Acknowledgments
This work is supported by DST-FIST (SR/FST/PSI-150/2010) and UGC
special assistance program, Government of India.
83
References
[1]
[2]
[3]
[4]
[5]
[6]
[7]
[8]
[9]
[10]
[11]
[12]
[13]
[14]
[15]
[16]
[17]
[18]
[19]
[20]
[21]
[22]
[23]
[24]
[25]
[26]
[27]
[28]
[29]
[30]
[31]
[32]
[33]
[34]
[35]
C.A. Angell, Solid State Ionics 9/10 (1983) 3–16.
T. Minami, J. Non-Cryst. Solids 73 (1985) 273–284.
P. Knauth, Solid State Ionics 180 (1) (2009) 911–916.
M. Tatsumisago, Solid State Ionics 175 (2004) 13–18.
K. Minami, F. Mizuno, A. Hayashi, M. Tatsumisago, Solid State Ionics 178 (2007)
837–841.
A.D. Robertson, A.R. West, A.G. Ritchie, Solid State Ionics 104 (1997) 1–11.
G.L. Henriksen, D.R. Vissers, J. Power Sources 51 (1994) 115–128.
M. Tatsumisago, F. Mizuno, A. Hayashi, J. Power Sources 159 (2006) 193–199.
M. Tatsumisago, T. Saito, T. Minami, Chem. Lett. 8 (2001) 790–791.
M. Tatsumisago, S. Hama, A. Hayashi, H. Morimoto, T. Minami, Solid State Ionics 154/
155 (2002) 635–640.
S. Soman, Y. Iwai, J. Kawamura, A.R. Kulkarni, J. Solid State Electrochem. 16 (2012)
1761–1766.
M. Tatsumisago, Y. Shinkuma, T. Minami, Nature 354 (1991) 217–218.
A. Hayashi, K. Noi, A. Sakuda, M. Tatsumisago, Nat. Commun. 3 (2012) 856, http://
dx.doi.org/10.1038/ncomms1843.
G.D.L.K. Jayasinghe, M.A.K.L. Dissanayake, P.W.S.K. Bandaranayake, J.L. Souquet,
Solid State Ionics 121 (1999) 19–23.
G.D.L.K. Jayasinghe, M.A.K.L. Dissanayake, M.A. Careema, J.L. Souquet, Solid State
Ionics 93 (1997) 291–295.
M. Ganguli, K.J. Rao, J. Non-Cryst. Solids 243 (1999) 251–267.
V.K. Deshpande, S.G. Charalwar, K. Singh, Solid State Ionics 40 (1990) 689–692.
M. Ganguli, M.H. Bhat, K.J. Rao, Phys. Chem. Glasses 40 (1999) 297–304.
N.K. Karan, B. Natesan, R.S. Katiyar, Solid State Ionics 177 (2006) 1429–1436.
J.L. Souquet, M.L.F. Nascimento, A.C.M. Rodrigues, J. Chem. Phys. 132 (2010) 034704.
J.L. Souquet, Annu. Rev. Mater. Sci. 11 (1981) 211–231.
St. Adams, K. Hariharan, J. Maier, Solid State Ionics 75 (1995) 193–201.
St. Adams, K. Hariharan, J. Maier, Solid State Ionics 86/88 (1996) 503–509.
A. Dalvi, K. Shahi, Solid State Ionics 148 (2002) 431–436.
A. Dalvi, K. Shahi, Indian J. Phys. 79 (2005) 727–732.
A. Dalvi, K. Shahi, Solid State Ionics 159 (2003) 369–379.
M. Mroczkowska, T. Czeppe, J.L. Nowinski, J.E. Garbarczyk, M. Wasiucionek, Solid
State Ionics 179 (2008) 202–205.
N. Gupta, A. Dalvi, A.M. Awasthi, S. Bhardwaj, Solid State Ionics 180 (2010)
1607–1612.
B.K. Money, K. Hariharan, J. Phys. Condens. Matter 115102 (2009) (10 pp.).
M. Rathore, A. Dalvi, Solid State Ionics 239 (2013) 50–55.
E. Barsoukov, J.R. Macdonald, Impedance Spectroscopy: Theory, Experiment, and
Applications, second ed. John Wiley & Sons, New Jersey, 2005.
N. Gupta, A. Dalvi, Ind. J. Pure Appl. Phys. 51 (2013) 328–330.
B. Roling, Solid State Ionics 105 (1998) 185–193.
D.P. Almond, G.K. Dunkan, A.R. West, Solid State Ionics 8 (1983) 159–164.
A.K. Jonscher, J. Phys. D. Appl. Phys. 32 (1999) R57–R70.
Indian Journal of Pure & Applied Physics
Vol. 51, May 2013, pp. 372-375
Electrical transport in Li2SO4-Li2O-B2O3 glass-ceramic composites
Munesh Rathore & Anshuman Dalvi*
Physics Department, Birla Institute of Technology and Science, Pilani 333 031, Rajasthan
*E-mail: [email protected]
Received 10 January 2013; revised 25 February 2013; accepted 3 April 2013
Lithium ion conducting glass-ceramic composites have been synthesized in Li2SO4-Li2O-B2O3 system by annealing the
glass above its crystallization temperature. The electrical, structural and thermal characterization of these glass-ceramics
reveals interesting results. The conductivity of the glass-ceramic increases with Li2SO4 content and exhibits a maximum of
~ 10−4 at 200°C interestingly for a composition 1Li2SO4-99(0.67Li2O-0.33B2O3). The glass-ceramic samples are found to be
thermally more stable than those of the glassy samples.
Keywords: Glass-ceramics, Ionic conductivity, Crystallization, Lithium borate glasses
1 Introduction
Fast ion conducting glasses have drawn
considerable attention due to their potential
application as solid electrolyte in solid state ionic
devices1-4. A glassy structure is preferable for the fast
ion transport in solids mainly due to (i) high ionic
conductivity owing to liquid like structure,
(ii) isotropic nature and (iii) highly non-crystalline
structure that enables negligible electronic
conductivity5. As a result, large number of
superionic glasses have been developed and
also used successfully as electrolytes in solid state
batteries6.
The lithium ion conducting glasses have the most
promising candidature as solid electrolytes in all solid
state lithium batteries7-9.To enhance the ionic
conductivity in these glasses, two strategies have been
used. The first one is by using mixed glass formers
where more than one glass former is normally used to
increase the mobility. In the second one, Li+ ion salt
(Li2SO4, LiI) is added to the glass matrix to increase
the number of mobile lithium ions10,11.
Several investigations have been carried out on
lithium silicate, lithium phosphate and lithium borate
glasses to which lithium halides and lithium oxy-salts
have been added10-12. Thus, there are many interesting
studies on salt- Li2O-MxOy type system (salt = LiCl,
Li2SO4 etc). Addition of LiCl is found to affect
electrical and structural properties of Li2O-SiO2-B2O3
glassy sysytem10-12. In a system xLi2SO4- 42.5Li2O(57.5-x)B2O3, notable rise in the conductivity11 is
observed till x = 15 m/o. The incorporation of Li2SO4
in the macromolecular network is responsible for such
a conductivity rise and decrease of glass transition
temperature11-15.
Though glasses exhibit high ionic conductivity,
their stability is limited due to the glass transition and
crystallization. Thus glass-ceramics have recently
drawn considerable attention due to their thermal
stability16,17. These are expected to be thermally more
stable than the glasses and thus are likely to be better
candidates as electrolytes in bulk and thin film solid
state batteries.
In the present study, we focus on the preparation
and characterization of Li+ ion oxide glass-ceramics.
In 67Li2O-33B2O3, the effect of Li2SO4 addition on
fundamental structural, thermal, and electrical
properties is discussed.
2 Experimental Details
The glasses of compositions xLi2SO4-(100-x)
[0.67Li2O-0.33B2O3] with x varying from 0 to 5 m/o
were synthesized by conventional melt-quenching.
Glasses were obtained using high-purity of powders
Li2CO3, H3BO3, and Li2SO4. Appropriate mixture of
these materials was taken in porcelain crucibles and
heated to 450ȠC for 2 h to remove CO2 and H2O from
Li2CO3, H3BO3. The mixture then melted at
900-950ȠC for 30 mins and subsequently pressed
between two copper plates. From these glasses, the
glass-ceramics were obtained by annealing at 600°C
for 6 h followed by slow cooling at room temperature.
For the structural characterization, the X-ray
diffraction
(RIGAKU
MiniFlex
II,
X-Ray
diffractometer) and scanning electron microscopy
(Ziess ultra-60 Field emission type) were used. Glass
RATHORE & DELVI: ELECTRICAL TRANSPORT IN GLASS-CERAMIC
transition and crystallization temperatures were
obtained by Differential scanning calorimetry
(DSC-60, Shimadzu). The electrical conductivity was
measured using computer controlled HIOKI LCR
meter model 3532-50 and Libratherm programmable
PID temperature controller.
3 Results and Discussion
The XRD patterns for glasses as well as glassceramics were obtained for various compositions.
Figure 1 shows the XRD patterns for sample with
x = 1 and 5 m/o Li2SO4 content. Fig. 1 (a and c)
shows the XRD results for as prepared melt quench
glassy samples. The absence of any prominent Bragg
peak in Fig.1 (a) confirms the amorphous/glassy
nature of the sample, whereas for the sample with 5
m/o there exist prominent peaks correspond to Li2SO4
that may have remained undissolved in glass matrix.
Further, Fig. 1(b and d) shows the XRD patterns of
glass-ceramic samples (annealed at 600°C).
Appearance of significant peaks confirms the
existence of crystallites viz., Li4B2O5 and Li2SO4 in
the glass matrix. Using debye-scherrer relation, size
of crystallites is calculated to be ~ 22-30 nm.
To understand the microstructural changes during
crystallization, scanning electron microscopy (SEM)
images of the glasses and glass-ceramics were
obtained for various compositions. Figure 2 shows
SEM images of 67Li2O-33B2O3 composition for (a) as
prepared glassy film and (b) glass-ceramic sample at a
resolution of 200 nm. For the glassy film, a good
homogeneity is seen on the surface. However, for the
373
glass-ceramic, apparent growth of rod shaped textured
crystallites surrounded by connective tissues of the
glass matrix (Fig. 2b) is noticed.
In order to understand the thermal stability of the
glasses and glass-ceramics, differential scanning
calorimetry (DSC) scans have been performed at a
typical heating rate of 10°C/min and shown in
Figs 3 and 4, respectively.
As apparent in Fig. 3(a), the glassy sample (with no
Li2SO4 content) exhibits a glass transition at ~ 350°C.
Fig. 2 — SEM images of 67Li2O-33B2O3 composition (a) glass,
and (b) glass-ceramic composite
Fig. 1 — X-ray diffraction patterns for samples: (a) as prepared
glass of composition 1Li2SO4-99(0.67Li2O-0.33B2O3), (b) glassceramic composite of composition 1Li2SO4-99(0.67Li2O0.33B2O3), (c) as prepared 5Li2SO4-95(0.67Li2O-0.33B2O3) and
(d) glass-ceramic composite of composition 5Li2SO495(0.67Li2O-0.33B2O3)
Fig 3ʊ DSC scans at a heating rate of 10°C/min for composition
67Li2O-33B2O3: (a) glass; (b) glass-ceramic composite
374
INDIAN J PURE & APPL PHYS, VOL 51, MAY 2013
Interestingly, the sample exhibits multiple (at least
two) crystallization at 396° and 442°C. Thus, the
presence of Tg and Tc in DSC scans suggests that
sample is purely glassy in nature and may have
different surroundings (at least two) about Li+ ions.
Fig. 3(b) shows DSC pattern for the glass-ceramic
sample of the same composition. Apparently, subtle
peaks appear at Tp1~ 385°C and Tp2 ~ 425°C.
Fig. 4 shows the similar DSC patterns for glass and
glass-ceramic samples with composition 1Li2SO499(0.67Li2O-0.33B2O3). In Fig. 4(a), the Tg and Tc
followed by two merged-up crystallization peaks are
seen at 287° and 330°C, respectively, which confirms
the glassy nature. In case of the glass-ceramic sample
of the same composition, no significant thermal event
is observed [Fig. 4(b)] till 500°C.
From Figs. 3 and 4, it may be suggested that
(i) addition of very little amount of Li2SO4 reduces the
Tg and Tc of the glass significantly (ii) thermal events
are not prominently seen in glass-ceramics and thus
they appear to be thermally more stable and (iii) XRD
and DSC results are found to be in good agreement
with each other.
Further, to investigate the thermal events,
conductivity-temperature cycles are obtained at
controlled heating rate of 1 ȠC/min and scrutinized in
the thermally unstable region, i.e, T > Tg and T > Tc.
Fig. 5 shows the electrical conductivity as a function
of temperature for glasses as well as for glass-ceramic
samples of composition xLi2SO4-(100-x)(0.67Li2O0.33B2O3) with x = 0 and 1.
For clarity, the whole ı-T cycles are divided in
three regions, viz. I, II and III. Figure 5 (a and b)
shows ı-T plot for glassy samples. The conductivity
exhibits an Arrhenius behaviour for both the glassy
compositions as shown (region I) reversible up to
~ 300°C. Interestingly, in region II, an anomalous rise
Fig. 4 — DSC scans at a heating rate of 10°C/min for the
composition 1Li2SO4-99(0.67Li2O-0.33B2O3): (a) glass, and
(b) glass-ceramic composite
in conductivity is observed at temperature 318° and
304°C, respectively, for both samples. On further
increase in temperature, the conductivity falls
drastically due to massive crystallization in multiple
stages. In region III, the crystallization completes and
the conductivity again increases linearly as a function
of temperature which once again confirms the ionic
nature of the glass-ceramic samples formed during
crystallization.
The conductivity temperature behaviour for glassceramic samples of the two compositions is shown in
Fig. 5 (c and d). For both the samples, conductivity
exhibits an Arrhenius behaviour as can be seen in
region I. In addition, both the samples exhibit a small
but notable fall in the conductivity at 370°C possibly
due to crystallization of the left over glassy state in
glass-ceramic samples. However, the fall in conductivity
at crystallization is drastically suppressed.
Fig. 5 — Temperature dependence of the conductivity samples:
(a), (b) glasses; (c), (d) glass-ceramic composite
Fig. 6 — Electrical conductivity (200°C) and activation energy for
glass and glass-ceramic as function of Li2SO4 content
RATHORE & DELVI: ELECTRICAL TRANSPORT IN GLASS-CERAMIC
For both, glasses as well as glass-ceramics, the
highest conductivity and minimum activation energy
is obtained for the composition 1Li2SO4-99(0.67Li2O0.33B2O3) as shown in Fig. 6. For x > 1 m/o of Li2O4
the samples are partially crystalline (as suggested by
XRD) and hence, conductivity decreases.Thus, it may
be suggested that low Li2SO4 content samples do
receive Li+ ions form the added salt. Further added
Li2SO4 does not dissolve in glass matrix, as a result
conductivity decreases. Thus, it is also expected that
the overall disorder in glass-ceramics also decreases
for x > 1 m/o of Li2SO4 and thus similar trend is seen
in the composition dependence of conductivity of
glass-ceramics.
4 Conclusions
The glass-ceramic samples exhibit pure ionic
transport and good thermal stability. The addition of
very small amount of Li2SO4 in 67Li2O-33B2O3
system enhances the ionic conductivity. The highest
conductivity for the glassy sample as well as glassceramic sample is found to be 10−4 and 10−6 Ω−1cm−1,
respectively, at 200°C for the composition 1Li2SO499(0.67Li2O-0.33B2O3). The XRD results do infer
that, annealing the glassy samples results into
precipitation of tiny crystallites in glassy matrix. SEM
results confirm the composite nature of the glassceramic samples. Li2SO4 addition also found to be
suppressing the crystallization. These glass-ceramics
may be promising candidate for high temperature
electrolytic applications with engineering of the
battery components. Efforts are on to further enhance
the ionic conductivity of Li+ ion conducting glassceramics in our laboratory.
375
Acknowledgement
This work is supported by DST FIST and UGC
Special assistance programme. We sincerely thank
Dr Dinesh Deva and Barkha Awasthi of Nanoscience
laboratory (IIT-Kanpur, India) for scanning electron
microscopy measurements.
References
1 Tuller H L, Button D P & Uhlmann D R, J Non-Cryst Solids,
40 (1980) 93.
2 Souquet J L, Ann Rev Mater Sci, 11 (1981) 211.
3 Hayashi A, J Ceram Soc of Japan, 115 (2007) 100.
4 Knauth P, Solid State Ionics, 180 (2009) 911.
5 Tatsumisago M, Solid State Ionics, 175 (2004) 13.
6 Quartarone E & Mustarelli P, Chem Soc Rev, 40 (2011)
2525.
7 Machida N, Yamamoto H, Asano S & Shigematsu T, Solid
State Ionics, 176 (2005) 473.
8 Hirai K, Tatsumisago M & Minami T, Solid State Ionics, 78
(1995) 269.
9 Souquet J L, Duclot M & Levy M, Solid State Ionics, 105
(1998) 237.
10 Deshpande A V &V. K. Deshpande Solid State Ionics, 154
(2002) 433.
11 Gandhi P R, Deshpande V K & Singh K,Solid State Ionics,
36 (1989) 97.
12 Deshpande A V & Deshpande V K, Solid State Ionics, 177
(2006) 2747.
13 Ganguli M, Bhat M H & Rao K J, Physics and Chemistry of
Glasses-European Journal of Glass Science and Technology
Part B, 40 (1999) 297.
14 Deshpande V K, Charalwar S G & Singh K, Solid State
Ionics, 40 (1990) 689.
15 Chryssikos G D, Kamitsos E I & Patsis A P, J NonCrystalline Solids, 202 (1996) 222.
16 Inda Y, Katoh T & Baba M, J Power Sources, 174 (2007)
741.
17 Xu X, Wen Z, Gu Z, Xu X & Lin Z, Electrochemistry
Commun, 6 (2004) 1233.