Correlation of Structure and Electrical Conductivity of CdI2 Doped

ARTICLE pubs.acs.org/JPCC

Correlation of Structure and Electrical Conductivity of CdI2 Doped Silver Borophosphate Glass and Nanocomposite S. Kabi and A. Ghosh* Department of Solid State Physics, Indian Association for the Cultivation of Science, Jadavpur, Kolkata 700032, India ABSTRACT: The structural and electrical properties of some CdI2 doped silver borophosphate glasses and glass nanocomposites of compositions 0.2CdI2
0.8(0.50Ag2O
0.50(xB2O3
(1
x)P2O5)), prepared by the melt quenching method, are reported in this paper. The X-ray diffraction study reveals the amorphous nature for the compositions x = 0
0.80 and partial crystalline nature for the composition x = 1. However, the transmission electron microscopic studies reveal the presence of some crystalline phases such as CdI2, β-AgI, and AgPO3 for x = 0
0.80 and β-AgI, Ag3BO3, and CdI2 for x = 1. The FTIR spectra reveal that the BO4 units are predominant for x = 0.20
0.50. The BO3 unit appears beyond the composition x = 0.20, and its population increases with the increase of B2O3 content. The conductivity of the compositions increases (for x = 0
0.60), and the activation energy decreases with progressive substitution of B2O3 in the glass network. However, the conductivity almost saturates for B2O3 rich compositions (for x = 0.60
1). The ion dynamics of the compositions as well as other physical properties such as density, glass transition temperature, etc., have been correlated with the compositional variation of BO4 and BO3 units.

1. INTRODUCTION The importance of ion conducting glasses stems from the interest due to their widespread application in solid state electrolytes.1
5 Different physical and chemical properties of glasses can be tailored by adjusting their compositions. The composition of a glass can be varied in several ways such as by changing the ratio of network modifier to network former, by varying the amount and type of dopant species, etc. Network modifiers or dopant species influence the microstructure or network structure of the glasses in different ways.6 For instance, the incorporation of modifiers imparts the anionic charge on the network. It changes the polymerization of the glass network depending upon formers.6 Similarly, the introduction of the dopant species increases the free volume in the glass network.7 The conductivity of the glassy electrolytes is controlled by doping halide salts.5 Substitution of oxygen by sulfur and doping with a larger anion (e.g., I
) also increases the conductivity of the glassy electrolytes.8 A different approach of controlling physical properties including the conductivity is the introduction of more than one glass former, which is literally known as the mixed glass former effect.8
12 The mixed glass forming systems are important for both academic and technological interest.13
16 A typical mixed glass forming system with network formers A and B can be written as yM2X þ (1
y)(xA þ (1
x)B). Here, M2X is the modifier [M = univalent metal and X = O or S]. The dc conductivity of the mixed former glasses usually passes through a maximum as a function of mixing ratio x. However, this phenomenon is not universal. In P2O5
TeO2 glasses, more than one conductivity maximum have been observed.17 With the variation of the mixing r 2011 American Chemical Society

ratio x, the local environment of each glass forming unit may or may not remain the same.8 When the local environment is the same as in the GeO2
GeS2 system,16 a mixed barrier model has been developed.8 This model predicts the time temperature superposition principle independent of the mixing ratio x. However, when the local environment of the glass forming units varies depending on x, the compositional dependence of different physical properties becomes complex. One example of this family of glass is borophosphate glass which has several applications including hermetic sealing materials and electrolytes.18
22 The compositional dependence of physical properties as well as ion dynamics for different alkali borophosphate glasses has been widely investigated.6,9,23
26 Advanced dipolar NMR studies and XPS studies25 of borophosphate glasses have provided much useful structural information about the concentration of structural units and interatomic interaction. The incorporation of B2O3 into metaphosphate glasses produces new linkages between phosphate chains through P
O
B bonds.6 The degree of network polymerization increases significantly with the increase of B2O3 content due to the formation of four coordinated boron species (BO4 units) which are linked to the phosphorus through a P
O
B bond.9 It has been observed that the number of bridging oxygen and the glass transition temperature increase in the same fashion as the B2O3 content is increased,9 which suggests that the glass transition temperature is strongly correlated to the network polymerization effect. Received: February 24, 2011 Revised: April 14, 2011 Published: April 26, 2011 9760

dx.doi.org/10.1021/jp2018296 | J. Phys. Chem. C 2011, 115, 9760–9766

The Journal of Physical Chemistry C

ARTICLE

Table 1. Interplanar Spacing (d) for Different Crystalline Phases in 0.2CdI2
0.8(0.50Ag2O
0.50(xB2O3
(1
x)P2O5)) as Obtained from XRD and TEM, and Their Corresponding Values along with Reflecting Planes Given in the JCPDS Datasheet d (Å) d (Å) d (Å) (from composition compound (from XRD) (from TEM) JCPDS) plane x = 0
0.80

CdI2

2.02

2.02

(116)

2.82

2.86

(109)

3.23

3.26

(101)

1.97

1.94

(630)

2.33

2.30

(440)

3.89 3.76

3.98 3.75

(100) (002)

3.06

3.06

(002)

3.23

3.23

(400)

3.07

3.03

2.73

2.70

2.84

2.84

2.15

2.18

3.89 3.68

3.89 3.72

3.98 3.75

(100) (002)

2.82

2.82

2.83

(200)

2.27

2.30

2.30

(440)

1.95

1.97

1.94

(630)

3.13

3.14

3.14

(104)

2.12

2.12

(110)

1.83

1.78

(201)

β-AgI

AgPO3

x=1

β-AgI

Figure 1. X-ray diffraction patterns for different compositions of 0.20CdI2
0.80(0.50Ag2O
0.50(xB2O3
(1
x)P2O5)) glass nanocomposites.

In the present report, we have investigated the mixed former effect in silver borophosphate glasses in the presence of a dopant species like CdI2. We have observed some interesting features regarding the structure and ion dynamics of these glasses which are different from undoped mixed former glasses.

2. EXPERIMENTAL PROCEDURE A series of the samples of compositions 0.2CdI2
0.8(0.50Ag2O
0.50(xB2O3
(1
x)P2O5)), where x = 0, 0.20, 0.40, 0.50, 0.60, 0.80, and 1, were prepared by the melt quenching method. The appropriate amounts of reagent grade powders CdI 2, AgNO3, H3BO3, and NH4H2PO4 (Aldrich Chem. Co.) were thoroughly mixed and preheated in alumina crucibles at 450 °C for 2 h for calcination. The mixtures were then melted at temperatures in the range from 800 to 900 °C depending upon the composition. The melts were equilibrated for 2 h and finally quenched between two aluminum plates. All the samples were kept in a dry chamber. X-ray diffraction (XRD) patterns of the powdered samples were recorded in an X-ray diffractometer (Bruker AXS, model D8 Advance) using Cu KR radiation (1.54 Å wavelength) at a scan rate of 0.02 deg/s at room temperature. For transmission electron microscopic (TEM) studies, the powder samples were sonicated in acetone for 15 min in an ultrasonic bath (model EYELA) and a few drops of the solution were poured in a 300 mesh copper grid. The TEM images were taken using a high resolution transmission electron microscope (JEOL HR-TEM, model JEM 2010). The FTIR spectra in KBr matrix (sample: KBr = 1:100) were recorded in a spectrometer (SHIMADZU, model 8400S) at room temperature. The deconvolution of the FTIR spectra was performed using Microcal Origin software. The absorption bands at different positions of the FTIR spectra were selected. The bands were fitted by a Gaussian distribution function. The peak area corresponding to each band as well as the total spectrum was

Ag3BO3 CdI2

determined. Differential scanning calorimetry was performed in an instrument (Perkin-Elmer, model DSC N-536-0022) under constant nitrogen flow with a heating rate of 10 °C/min. The density of the samples was measured using Archimedes’ principle. For electrical measurements, both sides of the samples were coated with silver paint to serve as the electrode. Electrical measurements, such as capacitance and conductance, were carried out in the frequency range 10 Hz to 2 MHz and in the temperature range 93
373 K using an LCR meter (QuadTech, model 7600) and a laboratory made cryostat. The dc conductivity was obtained from the complex impedance plots.

3. RESULTS AND DISCUSSION 3.1. X-ray Diffraction Studies. The X-ray diffraction patterns of different compositions of the 0.2CdI2
0.8(0.50Ag2O
0.50(xB2O3
(1
x)P2O5)) system are shown in Figure 1. It is observed that the compositions are amorphous in nature for the compositions x = 0
0.80. However, some prominent diffraction peaks superposed on the amorphous pattern are observed for the end composition x = 1. To identify the crystalline phases, the values of interplanar spacing (d) corresponding to the diffraction peaks have been calculated and compared with those given in the JCPDS datasheet (Table 1).27 It is observed in Table 1 that the peaks for the composition x = 1 correspond to β-AgI and Ag3BO3 crystalline phases. It should be mentioned that although the compositions x = 0
0.80 exhibit amorphous diffraction patterns, some nanocrystalline phases embedded in the glass matrix were observed in the TEM studies discussed later. 9761

dx.doi.org/10.1021/jp2018296 |J. Phys. Chem. C 2011, 115, 9760–9766

The Journal of Physical Chemistry C

ARTICLE

Figure 2. TEM micrograph for different compositions of 0.20CdI2
0.80(0.50Ag2O
0.50(xB2O3
(1
x)P2O5)) glass nanocomposites: (a) x = 0; (b) x = 0.50; (c) x = 0.80; (d) x = 1. SAED patterns for respective compositions are shown in the inset.

3.2. Transmission Electron Microscopic (TEM) Studies.

Figure 2 represents typical TEM micrographs for some compositions. The selected area electron diffraction (SAED) patterns for different compositions are also shown in the inset of each figure. Figure 2 shows the distribution of nanoparticles embedded in glass matrix. The interplanner spacing (d) of the crystalline phases was calculated from the SAED patterns, and comparing them with those given in the JCPDS-ICDD datasheet, the corresponding crystalline phases were determined.27 The calculated d-values and the corresponding crystalline phases for different compositions are listed in Table 1. Several crystalline phases such as CdI2, β-AgI, Ag3PO4, and Ag3BO3 dispersed in the glassy matrixes have been observed for compositions x = 0
1. The formation of the β-AgI crystalline phase is due to the exchange reaction between CdI2 and Ag2O (ref 26). The average particle size of the crystalline phases is quite small (10
30 nm) for the compositions x = 0
0.40. However, the average particle size of the crystalline phases increases for the composition x = 0.50
1. The particle size is within 40
100 nm for x = 0.50
0.80. It may be noted that the particle size of the crystallites for x = 1 is quite large (100
200 nm). The absence of diffraction peaks in the XRD patterns for the compositions x = 0
0.80 may be due to the absence of a significant population of large crystallites. Similarly, the absence of a diffraction peak due to CdI2 in the XRD pattern for the end composition x = 1 is due to the absence of a significant population of this crystallite. 3.3. Fourier Transformed Infrared (FTIR) Spectroscopic Studies. The FTIR spectra of the compositions 0.2CdI2
0.8(0.50Ag2O
0.50(xB2O3
(1
x)P2O5)) are shown in Figure 3a. The composition x = 0 shows the vibrational spectra due to the metaphosphate network which consists of PO4

structural units with two bridging and two nonbridging oxygen.28 These units are connected through P
O
P bridges, and a polymer-like chain structure is formed. The band observed near 1243 cm
1 is due to overlapping of stretching modes of the PdO group and (PO2)
fragments.29 The band observed near 1093 cm
1 is due to overlapping of stretching modes of the P
O
group and PO2 group.29 The bands observed near 893 and 700 cm
1 are due to asymmetric and symmetric stretching of the P
O
P group, respectively.29 It is observed in Figure 3a that the modification of these vibrational bands starts from the composition x = 0.20, and different vibrational bands in the spectra for the compositions x = 0.40
1 mainly correspond to the borate network. The symmetric stretching of the P
O
P group is observed only for x = 0 and 0.20, and this band almost vanishes for other compositions. The decrease of P
O
P links with the increase of the boron content has been also observed in the NMR studies.9 The vibrational bands at 1293 and 1093 cm
1 for x = 0 merge into a single band around 1100 cm
1 for the composition x = 0.20. This band as well as the bands observed between 900 and 1100 cm
1 for the compositions x = 0.40
1 are attributed to the B
O stretching vibration of BO4 units.30,31 A new band appears at around 1360 cm
1 for the compositions x = 0.40 and 0.50. This band is due to the B
O stretching vibration of BO3 units.30,32 It is noted that the band due to the BO3 units has been shifted nearly 70 cm
1 for the compositions x = 0.60
1. This occurs due to the presence of nonbridging oxygen in the BO3 units.31 The bending vibrations between 500 and 600 cm
1 of different phosphate or borate groups29
32 have been observed for all compositions. A borate network consists of BO4 and BO3 structural units. The proportion of these units is influenced by the modifier 9762

dx.doi.org/10.1021/jp2018296 |J. Phys. Chem. C 2011, 115, 9760–9766

The Journal of Physical Chemistry C

ARTICLE

Figure 4. Compositional dependence of glass transition temperature (a) and density (b) for 0.20CdI2
0.80(yAg2O
(1
y)P2O5) glass nanocomposites. A typical DSC curve for the composition x = 0.40 is shown in the inset of part a.

Figure 3. (a) FTIR spectra for different compositions of 0.20CdI2
0.80 (yAg2O
(1
y)P2O5) glass nanocomposites. (b) Deconvolution of the FTIR spectra is shown for a composition of x = 0.50. (c) Composition dependence of relative peak areas of BO4 and BO3 units.

content or the presence of another glass former (mixed glass former).33 The normalized peak areas of BO4 and BO3 units have been estimated from the deconvolution of the FTIR spectra. One example of deconvolution is shown in Figure 3b. It should be mentioned that the vibrational bands due to the different phosphate units become weaker or progressively disappear when the B2O3 content in the borophosphate glass is increased.33 Thus, the contribution of different phosphate units in the vibrational bands corresponding to BO4 or BO3 units has been neglected. The variation of the normalized peak areas of these units is shown in Figure 3c. Due to the substitution of P2O5 by B2O3, the formation of BO4 units begins to grow from the composition x = 0.20 and it is observed that the BO4 units are predominant for the compositions x = 0.20
0.50. However, the proportion of this unit starts to decrease from the composition x = 0.40 onward. NMR and Raman studies of several borophosphate glasses reveal that most of the boron atoms in phosphate rich glasses are in the four coordinated state,9 which is similar to our observation. The BO3 units appear from the composition x = 0.40. A sharp increase of BO3 units is observed for the composition x = 0.60 and onward. 3.4. Differential Scanning Calorimetric (DSC) Studies. Figure 4a shows the compositional dependence of the glass transition temperatures (Tg) which were obtained from the

endothermic baseline shift of the DSC curves, shown in the inset of Figure 4a for a composition. The glass transition temperature increases up to x = 0.40 and then decreases for x = 0.50
1. The increase of Tg up to x = 0.40 is mainly attributed to the progressive formation of BO4 units, which are linked to phosphate chains through P
O
B bonds.24 Formation of P
O
B bonds increases the network polymerization,24,34 which in turn increases Tg. The decrease of Tg can be attributed to two facts: First, the cross-linking between borate and phosphate networks is mainly accomplished by BO4 groups through B
(OP)4 linkages, which start to decrease from x = 0.50, reducing the polymerization of the glass network. Second, the particle size of the crystalline phases starts to increase from the composition x = 0.50 as observed from TEM investigation. This fact also affects the polymerization of the glass network and decreases the glass transition temperature.24 The compositional dependence of the density (Figure 4b) and the glass transition temperature is quite similar. The density increases from x = 0 to 0.40 due to increase of network polymerization. The decrease of the density from x = 0.50 onward is due to an increase of BO3 units. The sharp decrease of the density (as well as Tg) is due to a sharp decrease of BO4 units for this composition (Figure 3b). The abrupt increase of the density for x = 1 is due to the increase of crystalline volume for this composition. 3.5. dc Electrical Conductivity. The variation of the dc conductivity with reciprocal temperature for the compositions is shown in Figure 5. It is noted that the conductivity follows the Arrhenius relation σdcT = σ0 exp(
Eσ/kBT). The activation energy (Eσ) for the dc conductivity has been calculated from the least-squares straight line fits of the data. The dependence of the 9763

dx.doi.org/10.1021/jp2018296 |J. Phys. Chem. C 2011, 115, 9760–9766

The Journal of Physical Chemistry C

Figure 5. dc conductivity of different compositions of 0.20CdI2
0.80 (0.50Ag2O
0.50(xB2O3
(1
x)P2O5)) glass nanocomposites shown as a function of reciprocal temperature.

dc conductivity (at 303 K) and the activation energy on the B2O3 content is shown in Figure 6a and b, respectively. The room temperature dc conductivity for x = 0.6
1 was obtained from the extrapolation of the data shown in Figure 5. It is observed that, as the B2O3 content increases, the conductivity increases. However, the rate of increment is not the same for all the compositions. The conductivity increases slowly from x = 0 to 0.40, while the rate of increase is quite sharp from the compositions x = 0.40
0.60. From x = 0.60 to 1, the conductivity tends to nearly saturate. The activation energy decreases with an increase of B2O3 content. However, it decreases sharply for the composition x = 0.60. It has been reported35 that the silver ions in AgI doped silver phosphate glasses may be located in an iodine or oxygen environment. As the amount of CdI2 and Ag2O is fixed in the present glass compositions, the average Ag
I coordination for each composition should not vary to a large extent.36 Some of the Agþ ions may be located in the vicinity of nonbridging oxygen of the BO3 units.35 The presence of a nonbridging oxygen atom offers the hopping sites for mobile ions in the glass forming network.33,35 The formation of nonbridging oxygen also creates a relatively open network structure with large free volume for ion drift.33 The Agþ ions may also be present as a charge compensator of (BO4)
anions which act as the binding site or hopping site for mobile ions.9 Thus, both BO4 and BO3 units can influence the ionic motion in different ways. With the variation of the ratio of the two glass formers, the proportion of the BO4 and BO3 units changes as confirmed from the FTIR studies. Thus, the environment of the Agþ ions changes, which influences the mobility as well as the dynamics of Agþ ions in the mixed glass forming system. The ionic conductivity of the compositions can be explained in terms of the compositional variation of BO4 and BO3 units. We can divide the conductivity behaviors of the compositions into three different regions (Figure 6a).

ARTICLE

Figure 6. Composition dependence of (a) dc conductivity (at 303 K) and (b) activation energy shown as a function of B2O3 content (x) for 0.20CdI2
0.80(0.50Ag2O
0.50(xB2O3
(1
x)P2O5)) glass nanocomposites.

Region 1: From x = 0 to 0.40, the conductivity increases mainly due to the increase of BO4 units (Figure 3b). Region 2: From x = 0.40 to 0.60, the compositional dependence of BO4 and BO3 units has decreasing and increasing trends, respectively (Figure 3b). Thus, the sharp increase in the conductivity in this region is attributed to the increase of nonbridging oxygen in the BO3 units. It should be mentioned that in sodium borophosphate glasses the increase of the conductivity upon increase of B2O3 content was attributed to the increase of BO4 units.9 However, for the present compositions, the BO3 units also play a significant role for the enhancement of the conductivity. Region 3: From x = 0.60 to 1, the conductivity increases slightly. This is due to the significant decrease of BO4 units (Figure 3b) which act as the charge compensator for Agþ ions and also due to the increase of the size of the crystalline phases (as observed in TEM micrographs) which hinders the ionic mobility of these samples. The compositional variation of the activation energy can be explained on the basis of the Anderson
Stuart model.37 According to this model, the activation energy is the sum of two terms:37 Eσ = Eb þ Es, where Eb is the electrostatic binding energy and Es is the strain energy. Eb is required to surmount the electrostatic forces between Agþ ions and the neighboring ions such as oxygen or iodine. As the proportion of three and four coordinated borate units changes with the composition, the environment of silver ions also changes with the composition. This, indeed, affects the binding energy of silver ions. The strain energy depends on the elastic modulus of the glass network and the doorway radius of ion migration.36,37 The four coordinated borate units increase the network connectivity or polymerization which should increase the elastic modulus of the glass network. Thus, the strain energy changes with the composition. As both Eb 9764

dx.doi.org/10.1021/jp2018296 |J. Phys. Chem. C 2011, 115, 9760–9766

The Journal of Physical Chemistry C and Es change with the composition, the compositional dependence of activation energy becomes really complex. However, some correlation of activation energy with the relative population of BO4 and BO3 units has been observed. The activation energy for the composition x = 0
0.50 decreases slowly. The amount of BO4 units in this compositional range is nearly the same. For the composition x = 0.60, a sharp decrease in the activation energy is observed concomitant with the decrease of BO4 units for this composition. This observation is similar to the decrease of glass transition temperature and density for this composition. Thus, it may be predicted that the decrease of activation energy for this composition is mostly due to the decrease of strain energy. The compositional dependence of BO3 has an increasing trend. The increase of BO3 units decreases the number of P
O
B linkages, and the elastic modulus of the network is reduced. This also indicates that the decreasing trend of the activation energy for the compositions except for x = 1 is due to the decrease of strain energy. The increase of activation energy for x = 1 may be due to the increase of crystalline volume. It is noted that the structural feature of the CdI2 doped silver borophosphate glasses is somewhat different from undoped silver borophosphate glasses. In CdI2 doped silver phosphate glasses, a coordination exchange between Cd2þ and Agþ ions occurs according to Pearson’s acid base reaction mechanism.38 As a result, Cd2þ ions are mainly coordinated to oxygen and Agþ ions are coordinated to iodine.36 The degree of exchange reaction between Cd2þ and Agþ ions depends on the amount of CdI2 and Ag2O (ref 39). With the introduction of CdI2, the glass transition temperature of silver borophosphate glasses decreases. Similarly, it has been observed that the glass transition temperature of silver borophosphate glasses also decreases with the addition of AgI into the glass network.40 The maximum conductivity at 303 K for silver borophosphate glass is ∼10
5 Ω
1 cm
1.23 On the other hand, the maximum conductivity for the present CdI2 doped silver borophosphate glass is ∼10
2 Ω
1 cm
1. The latter value is almost the same as the maximum conductivity value achieved for AgI doped silver borophosphate glass.23 The high conductivity value achieved for AgI doped silver borophosphate glass was attributed to the formation of a highly conducting R-AgI crystalline phase into the glass network.23 However, we have observed that the AgI phase dispersed in the amorphous matrix of the present glass system is β-AgI. Thus, the increase of the conductivity for the present glasses is due to the systematic variation of the population of the BO3 and BO4 units, i.e., the structural changes of the glass forming networks.

4. CONCLUSIONS X-ray diffraction of the samples of the compositions 0.2CdI2
0.8(0.50Ag2O
0.50(xB2O3
(1
x)P2O5)) indicates amorphous nature for x = 0
0.80. However, some crystalline phases of β-AgI, AgPO3, and CdI2 have been detected in these compositions from transmission electron micrographs. The β-AgI, CdI2, and Ag3BO3 crystalline phases are dispersed in the glass matrix for the composition x = 1. Due to the progressive substitution of P2O5 by B2O3 in the glass compositions, BO4 and BO3 units are formed. The glass transition temperature of the compositions initially increases with an increase of boron content due to an increase of BO4 units. However, for borate rich compositions, the glass transition temperature decreases due to the decrease of BO4 units as well as due to the increase of the particle size of crystalline phases. The conductivity of the

ARTICLE

compositions increases and the activation energy decreases with progressive substitution of B2O3 in the glass network. However, the conductivity almost saturates for B2O3 rich compositions. The composition dependence of the ionic conductivity and activation energy is significantly influenced by the compositional variation of BO4 and BO3 units.

’ AUTHOR INFORMATION Corresponding Author

*E-mail: [email protected].

’ ACKNOWLEDGMENT We acknowledge the DST for support through the nano initiative program. ’ REFERENCES (1) Dutta, D.; Ghosh, A. J. Phys. Chem. C 2009, 113, 9040. (2) Hayashi, A.; Minami, K.; Tatsumisago, M. J. Non-Cryst. Solids 2009, 355, 1919. (3) Minami, K.; Hayashi, A.; Tatsumisago, M. J. Non-Cryst. Solids 2010, 356, 2666. (4) Tatsumisago, M.; Hayashi, A. J. Non-Cryst. Solids 2008, 354, 1411. (5) Bhattacharya, S.; Ghosh, A. J. Phys. Chem. C 2010, 114, 5745. (6) Elbers, S.; Strojek, W.; Koudelka, L.; Eckert, H. Solid State Nucl. Magn. Reson. 2005, 27, 65. (7) Swenson, J.; B€orjesson, L. Phys. Rev. Lett. 1996, 77, 3569. (8) Schuch, M.; M€uller, C. R.; Maass, P.; Martin, S. W. Phys. Rev. Lett. 2009, 102, 145902. (9) Zielniok, D.; Cramer, C.; Eckert, H. Chem. Mater. 2007, 19, 3162. Banhatti, R. D.; Cramer, C.; Zielniok, D.; Jean Robertson, A. H. Z. Phys. Chem. 2009, 223, 1201. (10) Money, B. K.; Hariharan, K. Solid State Ionics 2008, 179, 1273. (11) Chowdari, B. V. R.; Pramoda Kumari, P. Solid State Ionics 1998, 113
115, 665. (12) Chowdari, B. V. R.; Radhakrishnan, K. J. Non-Cryst. Solids 1989, 108, 323. (13) Nandi, P.; Srinivasan, A.; Jose, G. Opt. Mater. 2009, 31, 653. (14) Munoz-Martín, D.; Villegas, M. A.; Gonzalo, J.; FernandezNavarro, J. M. J. Eur. Ceram. Soc. 2009, 29, 2903. (15) Aitken, B. G.; Youngman, R. E.; Deshpande, R. R.; Eckert, H. J. Phys. Chem. C 2009, 113, 3322. (16) Kim, Y.; Saienga, J.; Martin, S. W. J. Phys. Chem. B 2006, 110, 16318. (17) Coppo, D.; Duclot, M. J.; Souquet, J. L. Solid State Ionics 1996, 90, 111. (18) Qiu, D.; Guerry, P.; Ahmed, I.; Pickup, D. M.; Carta, D.; Knowles, J. C.; Smith, M. E.; Newport, R. J. Mater. Chem. Phys. 2008, 111, 455. (19) Malakho, A.; Dussauze, M.; Fargin, E.; Lazoryak, B.; Rodriguez, V.; Adamietz, F. J. Solid State Chem. 2005, 178, 1888. (20) Petit, L.; Cardinal, T.; Videau, J. J.; Durand, E.; Canioni, L.; Martines, M.; Guyot, Y.; Boulon, G. Opt. Mater. 2006, 28, 172. (21) Lee, S.; Kim, J.; Shin, D. Solid State Ionics 2007, 178, 375. (22) Cho, K.-H.; You, H.-J.; Youn, Y.-S.; Kim, J.-S.; Shin, D.-W. Electrochim. Acta 2006, 52, 1571. (23) Magistris, A.; Chiodelli, G. Solid State Ionics 1983, 9 and 10, 611. (24) Mu~ noz, F.; Montagne, L.; Pascual, L.; Duran, A. J. Non-Cryst. Solids 2009, 355, 2571. (25) Raskar, D.; Rinke, M. T.; Eckert, H. J. Phys. Chem. C 2008, 112, 12530. (26) Zeyer-D€usterer, M.; Montagne, L.; Palavit, G.; J€ager, C. Solid State Nucl. Magn. Reson. 2005, 27, 50. 9765

dx.doi.org/10.1021/jp2018296 |J. Phys. Chem. C 2011, 115, 9760–9766

The Journal of Physical Chemistry C

ARTICLE

(27) JCPDS-ICDD 1993, 11-0641,09-0374,21-1078,22-1192,779089, 40-1469. (28) Brow, R. K. J. Non-Cryst. Solids 2000, 263
264, 1. (29) Efimov, A. M. J. Non-Cryst. Solids 1997, 209, 209. (30) Julien, C.; Massot, M.; Balkansk, M.; Krol, A.; Nazarewicz, W. Mater. Sci. Eng., B 1989, 3, 307. (31) Kamitsos, E. I.; Karakassides, M. A.; Chryssikos, G. D. J. Phys. Chem. 1986, 90, 4528. (32) Kamitsos, E. I.; Patsis, A. P.; Karakassides, M. A.; Chryssikos, G. D. J. Non-Cryst. Solids 1990, 126, 52. (33) Cho, K.; Lee, S. H.; Shin, D. W.; Sun, Y. K. Electrochim. Acta 2006, 52, 1576. Anantha, P. S.; Hariharan, K. Mater. Chem. Phys. 2005, 89, 428. (34) Costantini, A.; Buri, A.; Branda, F. Solid State Ionics 1994, 67, 175. (35) Minami, T. J. Non-Cryst. Solids 1983, 56, 15. (36) Swenson, J.; Matic, A.; Gejke, C.; B€orjesson, L.; Howells, W. S.; Capitan, M. J. Phys. Rev. B 1999, 60, 12023. (37) Anderson, O. L.; Stuart, D. A. J. Am. Ceram. Soc. 1954, 37, 573. (38) Pearson, R. G. J. Am. Chem. Soc. 1963, 85, 3533. (39) Kabi, S.; Ghosh, A. Solid State Ionics 2011, 187, 39. (40) Chiodelli, G.; Magistris, A. Solid State Ionics 1986, 18 & 19, 356.

9766

dx.doi.org/10.1021/jp2018296 |J. Phys. Chem. C 2011, 115, 9760–9766