Chapter 4

Microstructural, dielectric and spectroscopic properties of Li2O–Nb2O5–ZrO2–SiO2 glass system crystallized with V2O5

Li2O–Nb2O5–ZrO2–SiO2 glasses mixed with different concentrations of

V2O5 were crystallized. The samples were characterized by XRD, SEM and DTA

techniques. The SEM pictures indicated that the samples contain well defined and

randomly distributed crystal grains. The X-ray diffraction studies have revealed

the presence of ZrV2O7, ZrSi24O50, LiV3O8, Li2V2O5, LiVO3, Li2ZrO3, NbVO5,

LiNbO3, Nb6V2O19, Nb2V2O9, Li2SiO3 and ZrSiO4 crystalline phases in these

samples. Optical absorption, ESR and photoluminescence spectral studies on

these samples have indicated that a considerable proportion of vanadium ions do

exist in V4+ state in addition to V5+ state and the redox ratio seems to be

increasing with increase in the concentration of crystallizing agent V2O5. The

infrared spectral studies have pointed out the existence of conventional SiO4,

ZrO4, NbO6, V=O structural units in the glass ceramic network. The study of

dielectric properties suggested a decrease in the insulating character of the glass

ceramics with increase in the crystallizing agent. A.C. conductivity in the high

temperature region seems to be connected mainly with the polarons involved in

the process of transfer from V4+V5+ ions.

Microstructural, dielectric and spectroscopic properties of Li2O–Nb2O5–ZrO2–SiO2 glass system crystallized with V2O5

4.1. Introduction

Niobium mixed lithium silicate glasses containing transition metal ions have been

the subject of an increasing academic and technological interest as mentioned in

the earlier chapter. These glasses find a wide range of applications, such as glass

fibers and optical lenses [1, 2], as electrodes [3], for radioactive waste

immobilization [4], in hermetic sealing of metallic and ceramic materials [5] and

as glass planar optical waveguides fabricated by using the sol gel and dip coating

technique [6]. The addition of Nb2O5 bestows the base glass to possess electro

chromic and electro-optical properties. Earlier investigations on spectroscopic

properties of different niobate glasses indicated that Nb ions participate in the

glass network with NbO4 and NbO6 structural units [7-9]. In the glass ceramics

containing Li2O and Nb2O5, there is a possibility for the formation of LiNbO3

crystal phases which exhibit ferroelectric properties [10] and makes the material

for the potential applications in optoelectronics, acoustic-optics.

The addition of ZrO2 to niobium silicate glasses is expected to increase the

electrical resistivity and chemical inertness. In a number of earlier investigations,

it was also established that the inclusion of ZrO2 in silicate glasses causes a

substantial hike in the refractive index, decreases the cut–off wavelength and

reduces the photochromism of the glass [11, 12]. In view of these qualities ZrO2

containing silicate glasses find variety of applications, such as thermal barrier

161

coating, optical filters, laser mirrors and alternative gate dielectrics in

microelectronics.

When compared with the glass materials, glass ceramics are expected to

have several advantages like good mechanical, electrical and thermal properties,

high chemical durability and low coefficient of thermal expansion with no crack

growth inside. Hence, a considerable interest is attached in this connection to the

studies on crystallization of the glass materials and its bearing on the physical

properties. Catalysts generally used for controlled crystallization processes, giving

rise to enormous numbers of nucleation centers in the original glass are, gold,

silver, platinum or the oxides of transition metals like Ti, Cr, Mn, Ce, V, Fe, Co,

Ni etc. The crystallization of Li2O–Nb2O5–ZrO2–SiO2 glasses with V2O5 is an

added advantage for the simple reason that the presence of vanadium ions makes

the material to exhibit semiconducting behavior with the electrical conductivity

>10−3 to 10−5 (ohm cm)−1 due to electron hopping between V4+ to V5+ ions [13].

V2O5 containing glasses are being extensively used in memory and switching

devices. The crystallization of these materials, further leads to the formation of

Li2V2O5 crystal grains in which vanadium ions present in V4+ state. Presence of

such complexes facilitate to accelerate the rate of hopping of electron between V4+

and V5+ ions which ultimately lead to the enhancement of conductivity. The

process of hopping of the electrons between V4+ and V5+ ions in the presence of

larger concentrations of mobile cations like lithium ions is highly interesting and

162

useful to investigate in the multi component glass ceramic systems like Li2O–

Nb2O5–ZrO2–SiO2 in view of huge technological importance of this material.

In spite of the fact, that there are some studies available on silicate glasses

and glass ceramics containing vanadium ions [14, 15], still there is a lot of scope

to probe this ion in SiO2 based glass ceramics (by means of dielectric studies)

especially when they are mixed with incipient glass formers like ZrO2 and Nb2O5.

Further, the investigation on dielectric properties coupled with spectroscopic

studies not only helps in assessing the insulating character but also expected to

shed the light on many aspects, such as the geometry of various structural units of

glass network, the character of chemical bonds of dopant transition metal ions

[16–19].

The objective of the present investigation is to have a comprehensive

understanding over the topology and valence states of vanadium ions in (30-x)

Li2O–10Nb2O5–5ZrO2–55SiO2: xV2O5 (0 ≤ x ≤ 1.0) glass ceramic network by a

systematic study of dielectric (over a wide range of frequency and temperature)

and spectroscopic properties and to use results of these studies to throw some light

on the structural aspects and the insulating character of the material.

The contents of the paper are as follows: section 1 presents the technology

of the sample preparation and the details of experimental techniques adopted for

the sample characterization. In section 2 the results of characterization studies viz.,

electron microscopy, X-ray diffraction, Energy dispersive spectroscopy (EDS)

analysis and differential thermal analysis (DTA), are presented. Section 3 is

163

devoted to the presentation of the studies on physical properties including

dielectric and spectroscopic properties (viz., optical absorption,

photoluminescence, electron spin resonance (ESR) and infrared (IR) spectra). In

the discussion part, principal influence of valence states of vanadium ion and its

environment on physical properties is presented.

4.2 Brief review of the previous work on the glasses containing vanadium ions

The studies as such on vanadium containing silicate glasses are limited. However

in this review, the studies on different glass systems including some silicate

glasses containing vanadium ions, has been described briefly.

Agarwal et al. [20] have recently reported the electron paramagnetic

resonance studies of vanadyl doped alkali niobium borate glasses. In this study it

was observed that the tetragonality of V4+O6 complex decreases with increasing

concentration of Nb2O5. Lin-Hua et al. [21] have recently reported electron

paramagnetic resonance spectra of vanadyl doped zinc phosphate glass. In this

study the compressed defect structure of V4+ center is discussed in detail. Kerkouri

et al. [22] have reported spectroscopic studies and the structural aspects of V2O5-

CdO-P2O5 glasses. In their study it was reported that the glasses containing more

than 20% of V2O5, the VO4 and VO5 structural units with V–O–V bridges were

formed in the glass network. Gouda et al. [23] have investigated the effect of

replacing vanadium by Cu2+ ion on the dc-electrical conductivity () and I–V

characteristics of (V2O5)0.7(GeO2)0.3(CuO)x glasses. In this report the electrical

164

conduction is interpreted on the basis of electrons hopping from reduced to

unreduced vanadium and/or cupper ions. Behzad et al. [24] have studied the

conductivity of 50P2O5–xV2O5–(50–x)Li2O glass system as a function of

temperature and composition. Isothermal variation of conductivity as a function of

composition of this study showed a minimum for a molar ratio x near 20. Probable

mechanisms for decrease of conductivity with decrease of vanadium oxide

concentration were explained in detail. Ali and Ezz-Eldin [25] have studied some

physical properties of the lithium disilicate (Li2Si2O5) glasses doped with different

ratios of V2O5 before and after gamma-rays irradiation. The observed variations in

the physical properties with the change in the concentration of V2O5 were

correlated with the changes in internal glass network. Abd El-Aal and Afifi [26]

have reported the elastic properties of vanadium tellurite glasses, 65TeO2–(35-

x)V2O5–xCuO, with different compositions of copper at room temperature by

ultrasonic methods. Khattak et al. [27] have studied X-ray photoelectron

spectroscopy (XPS) and magnetic susceptibility studies of vanadium phosphate

glasses. From the analysis of the results of these studies, the authors have

proposed a glass structure model consisting of a mixture of vanadate phosphate

phases that include V2O5, VOPO4, (VO)2P2O7, VO(PO3), and V(PO3)3 with the

abundance of orthophosphate (PO4)3- units increasing with increasing vanadium

content. In another report [28] these authors have also reported the results of X-ray

photoelectron spectroscopy (XPS) of vanadium tellurite glasses. Quantitative

165

analysis of the results of these studies has indicated that there is a change of TeO4

to TeO3 upon V2O5 addition.

Bogomolova et al. [29] have reported electron paramagnetic resonance

studies of V4+ ions in lanthanum-aluminosilicate glasses. From the detailed

analysis of EPR spectra it was concluded that in the samples containing low

content of La2O3 lanthanum acts predominately as modifying ion where as in glass

with a high La2O3 concentration lanthanum gradually occupies a glass forming site

in the network. Saddeek [30] has recently reported FTIR and elastic properties by

ultrasonic methods on MoO3–V2O5–PbO glasses. The observed compositional

dependence of the elastic moduli was interpreted in terms of the effect of MoO3 on

the coordination number of the vanadate units. A good correlation was observed

between the experimentally determined elastic moduli and those computed

according to the Makishima-Mackenzie model. Feng [31] has reported detailed

theoretical studies of the optical and EPR spectra for vanadyl ions in alkaline-earth

aluminoborate glasses. In this study the Optical spectra and electron paramagnetic

resonance (EPR), g and A factors of calcium aluminoborate glasses (CaAB): VO2+

are calculated using a complete diagonalization (of the energy matrix) process.

Good agreement between the theoretical values and experimental results attributed

effectiveness of the CDP method for theoretical studies of optical and EPR spectra

of 3d1 (or V4+) ions in glasses. Shapaan et al. [32] have investigated hyperfine

structure and electric transport properties of vanadium iron phosphate glasses and

interpreted the results with the aid of the data on Mössbauer spectroscopy.

166

Increase of V2O5 content, resulted the increase in dc conductivity while the

activation energy was found to decrease. The observed increase of dielectric

constant () with increasing V2O5 content was attributed to the increase in the

deformation of glass network. Kartashov and Vysloukh [33] have investigated the

propagation of laser beams in SiO2-VO2 nanocomposite waveguides with thermo-

optical nonlinearity. These studies have indicated that the large modifications of

the absorption coefficient as well as notable changes of the refractive index of VO2

nanoparticles embedded into the SiO2 host media that accompany the

semiconductor-to-metal phase transition. At the end they have concluded that such

changes may lead to optical limiting in the near-IR wave range.

Sung et al. [34] have investigated a variety of thermal properties of P2O5–

V2O5–ZnO/B2O3 glasses that include glass transition temperature, dilatometer

softening point and coefficient of thermal expansion, and aqueous durability. From

these studies it was observed that the aqueous durability was improved through the

addition of some additives such as Al2O3 and TiO2. ElBatal et al. [35] have

prepared V2O5-doped sodium phosphate glasses of various compositions and

studied various spectroscopic properties that include UV-vis and infrared, Raman

and electron spin resonance. In this study the changes observed in UV-vis and

infrared spectral data were discussed in relation to the structural evolution caused

by the change in the V2O5 content or glass composition. Kim et al. [36] have

investigated the local structures of the boron and vanadium sites in the ternary

xV2O5–B2O3–yNa2O glass by means of magic angle spinning (MAS) nuclear

167

magnetic resonance (NMR) techniques. In this study it was observed that with

increasing x, the mole ratios of the BO3 and BO4 structures were enhanced, as

were the quadrupole asymmetry parameters for the BO3 structures, while the

quadrupole coupling constants for the sites were reduced. Moawad et al. [37] have

investigated dc conductivity the mixed electronic-ionic conduction in 0.5[xAg2O–

(1–x)V2O5]–0.5TeO2 glasses. In this study it was observed that the mechanism of

dc conductivity changes from predominantly electronic to ionic within the

30Ag2O40 range. Taibi et al. [38] have investigated the influence of the

V2O5/Sb2O3 substitution on the physical properties of the (70–x)Sb2O3–xV2O5–

30K2O glasses. From these studies the authors have suggested that there is a

change in the coordination number of the vanadium cations in relation to the

network topology. Farah [39] has studied the relationship between glass

composition and optical basicity by redox analysis of vanadium in Na2O and CaO

based Al2O3–SiO2 glasses/melts. In this study it was found that the V4+/V5+

equilibrium was more affected by a change in the calculated optical basicity

compared with that of the V3+/V4+ in both sodium and calcium silicate series.

However, Alumina saturation from the crucible did not affect the calculated

optical basicity of the sodium silicate glasses, although there was a change in the

corresponding redox ratios. The results were compared using different

experimental parameters and were found to be useful in glass production and

extractive metallurgy of vanadium.

168

Vedeanu et al. [40] have investigated structural changes induced by CuO

and V2O5 in the phosphate glass network by means of Raman spectroscopy. In this

study it was observed that at higher concentrations of V2O5 a strong

depolymerization of the phosphate network was taking place. Kumar et al. [41]

have reported thermal and electrical properties of tellurium-based glasses doped

with vanadium and vanadium-cobalt oxides. From the thermal properties the

authors have estimated that the thermal stability, fragility and glass-forming

tendency of the glass system. The results of dc conductivity have been analyzed in

the light of Mott's small polaron hopping (SPH) and Mott's and Greave's variable

range hopping (VRH) models. Kundu et al. [42] have investigated effect of V2O5

on structural, physical and electrical properties of bismuth borate glasses. In this

study the dc conduction was found to increase with increase in vanadium content

and the mechanism was explained in terms of polaron hopping. Rao et al. [43]

have reported dielectric dispersion in Li2O–MoO3–B2O3 glass system doped with

V2O5. The observed dielectric relaxation effect was analyzed quantitatively by

pseudo Cole-Cole plot method and the spreading of relaxation times was

established. Li et al. [44] have studied the effect of lanthanum on structure of

vanadate-phosphate glass by means of Fourier infrared spectroscopy, Raman

spectroscopy and nuclear magnetic resonance. The analysis of these results have

indicated that vanadium existed in the glass in the form of (VO3)n single chains,

(V2O8)n zigzag chains, VO4 branches and groups. Rada et al. [45] have

investigated the effect of the introduction of vanadium pentaoxide on structural

169

changes in phospho-tellurite glasses containing gadolinium ions. In these studies it

was found that the addition of V2O5 resulted in gradual depolymerization of the

phosphate chains and formation of short phosphate units in the glass network.

Ardelean et al. [46] reported the EPR study of V2O5–P2O5–Li2O glass system. In

this study it was found that at high V2O5 content, the vanadium hyperfine structure

disappears and only the broad line could be observed in the spectra. Spin

Hamiltonian parameters g||, g⊥, A||, A⊥, dipolar hyperfine coupling parameters, P,

and Fermi contact interaction parameters, K, have been calculated. The

composition dependence of line widths of the first two absorptions from the

parallel band and of the broad line characteristic to the cluster formations was also

discussed in detail.

Al-Assiri [47] has studied electrical properties of vanadium-copper-

phosphate glasses. In this study the dc conductivity was found to increase while

the activation energy is found to decrease with the increase of the V2O5 content.

Further, it was reported that the dc conductivity in these glasses is electronic in

nature and depends strongly upon the average distance, R, between the vanadium

ions. Tian et al. [48] have investigated the effect of V2O5 content on the phase-

transformation and the microstrctural development in the SiO2–MgO–Al2O3–K2O–

V2O5–F glass by XRD, SEM and EPMA measurements. In this study it was found

that the incorporation of V2O5 leads to the precipitation of mullite (Al6Si2O13)

crystalline phases at lower temperatures. Further, it was reported that although the

types of crystalline phases, which are formed as mullite and mica, were less

170

influenced by V2O5 contents, the morphology, volume fraction and sizes of

crystals were dependent sensitively on the V2O5 content. Hager [49] has reported

the region of glass formation of ternary V2O5–BaF2–RF (RF=LiF, NaF and mixed

NaF–LiF) and has also measured a variety of physical parameters that include

density, molar volume, characteristic temperatures, average thermal expansion

coefficient, , and specific heat, Cp. In this study it was observed that the glass-

transition temperature, Tg, decreased by increase of RF while thermal expansion, ,

and specific heat, Cp, near Tg were increased. Jung et al. [50] have carried out

NMR investigations on PbO–B2O3 glasses containing V2O5 and they have

calculated the ratio of BO4 and BO3 units as a function of V2O5 concentration.

Mekki et al. [51] have investigated the magnetic properties of vanadium-sodium

silicate glasses. Ferrari et al. [52] have probed the effect of V2O5 on the

crystallization of CaO–ZrO2–SiO2 glasses. Abd El-Moneim [53] has carried out

DTA and IR spectra of vanadium tellurite glasses. Hoppe et al. [54] have reported

the details on structure of vanadium tellurite glasses by using neutron and X–ray

diffraction. Cozar et al. [55] have studied the IR and EPR studies on some lithium

borate glasses with vanadium ions. Salim et al. [56] have carried out X–ray

photoelectron spectroscopy and magnetization studies of iron–vanadium

phosphate glasses. Khattak et al. [57] have studied X–ray photoelectron

spectroscopy (XPS) and magnetic properties of copper-vanadium phosphate

glasses. Krasowski et al. [58] have measured electrical conductivity of silver

171

vandate glasses. Maria-Camelja et al. [59] have investigated the ionic and

electronic conductivity of P2O5–V2O5–Na2O glasses.

Sudarsan et al. [60] have carried out a study on the structural aspects of

V2O5–P2O5–B2O3 glasses by using MAS NMR and IR spectral studies. Sega et al.

[61] have studied the electrical conduction in V2O5–NiO–TeO2 glasses. Mori et al.

[62] have discussed the conduction phenomenon in V2O5–Sb–TeO2 glasses on the

basis of small polaron hopping model. Rajendran et al. [63] have carried out

ultrasonic investigations in V2O5–PbO glasses containing BaTiO3. Simockova et

al. [64] have investigated complex impedance response of V2O5–P2O5 glasses.

Takahashi et al. [65] have investigated the structure of AgI–Ag2O–V2O5 glasses.

Moustafa et al. [66] have reported the spectroscopic studies of semiconducting

barium vandate glasses doped with iron oxide. Attos et al. [67] have investigated

the structure of borovanadate glasses by Raman spectroscopy. Murawski et al.

[68] has investigated the dielectric relaxation in semiconducting glasses. Seth et al.

[69] have studied the EPR study of vanadyl ion in CoO–PbO–B2O3 glasses.

Prakash et al. [70] have carried out the EPR study of vanadyl ion in CoO-PbO-

B2O3 glasses. Amano et al. [71] have studied the electrical properties of Sb2O3–

CaO–V2O5 glasses and glass ceramics. Bogomolova et al. [72] have investigated

the role of V2O5 on the structure of fluoro germanate glasses using ESR

measurements. Yoko et al. [73] have studied the IR and NMR spectral studies on

lead vanadate glasses. Ghosh et al. [74] have reported the spectral studies of

binary iron vanadate glasses. Gupta et al. [75] reported the influence of V4+ ion

172

concentration on the EPR spectra of vanadate glasses. Adams et al. [76] have

studied the silver ion conductivity during the crystallization of AgI–Ag2O–V2O5

glasses. Dimitrov et al. [77] have analyzed V2O5–GeO2–Bi2O3 glass structure by

IR spectra.

4.3 Characterization

The detailed compositions of the glasses used in the present study are as follows:

V0: 30.0Li2O–10Nb2O5–5ZrO2–55SiO2

V1: 29.9Li2O–10Nb2O5–5ZrO2–55SiO2: 0.1V2O5

V3: 29.7Li2O–10Nb2O5–5ZrO2–55SiO2: 0.3V2O5

V5: 29.5Li2O–10Nb2O5–5ZrO2–55SiO2: 0.5V2O5

V8: 29.2Li2O–10Nb2O5–5ZrO2–55SiO2: 0.8V2O5

V10: 29.0Li2O–10Nb2O5–5ZrO2–55SiO2: 1.0V2O5

4.3.1 Physical parameters

From the measured values of the density and average molecular weight M

of the samples, various other physical parameters such as vanadium ion

concentration Ni, mean vanadium ion separation ri, polaron radius rp in Li2O–

Nb2O5–ZrO2–SiO2: V2O5 glass ceramic samples are computed and presented in

Table 4.1.

Table 4.1

Physical parameters of Li2O–Nb2O5–ZrO2–SiO2: V2O5 glass ceramic samples.

Sample Density (g/cm3)

Molar vol. (cm3/mol)

Vanadium ion conc. Ni (X1020 ions/cm3)

Interionic distance Ri ()

Polaron radius () 3/1

621

=

iP N

Rπ

Refractive Index (n)

Electronic polarizability (10-22)

( ) ( )( )N

nne 3/4

21 22

πα +−=

V0 3.0966 24.14 ---- --- ---- 1.781 4.02

V1 3.0986 24.17 2.49 34.2 13.8 1.784 4.03

V3 3.1041 24.22 7.45 23.8 9.6 1.786 4.05

V5 3.1050 24.32 12.38 20.1 8.1 1.787 4.07

V8 3.1084 24.44 19.72 17.2 6.9 1.789 4.09

V10 3.1141 24.49 24.59 16.0 6.4 1.793 4.12

4.3.2 Scanning electron microscopy

The scanning electron microscopy (SEM) pictures of some of the samples

investigated are presented in Fig. 4.1. The pictures of pre-crystallized glasses

exhibited virtually no crystallinity. The SEM pictures of crystallized Li2O–Nb2O5–

ZrO2–SiO2: V2O5 samples clearly suggests the presence of some structured species

on the surface of the glasses (Figs. 4.1 (a & b)). The chemical makeup of the

samples is analyzed using energy dispersive spectra (EDS); the analysis indicated

the presence of Li, Nb, Zr, Si, V and O elements in various crystalline phases (Fig.

4.1 (c)). The X–ray map for vanadium ions in one of the glass ceramics (viz., V8)

is presented in Fig. 4.1(c). The maps indicate a reasonably uniform distribution of

the vanadium ions in the entire glass ceramic material. The maps of other samples

also exhibited the similar behaviour.

4.3.3 X-ray diffraction

X-ray diffraction studies (Fig. 4.2) indicate the possible formation of

several primary crystalline phases containing the multiple combination of

elements viz., Si, Zr, V, Li, Nb, and O [78]. However, in the context of presence

of mixture of several complex crystalline phases, it is difficult to conclude

quantitatively about the exact species present and the nature of the crystalline

phases. The crystalline phases indicated in Fig. 4.2 are of most possible ones but it

Fig. 4.1. (a) SEM photograph of pre-crystallized glass (V8), (b & c) SEM photographs of Li2O–Nb2O5–ZrO2–SiO2 glasses crystallized with 0.5 and 0.8 mol% of V2O5 (d) EDS of sample V8 with X-ray mappings for vanadium ions.

V8

(d)

requires more qualitative analysis which is beyond the scope of the objective of

this paper.

The V2O5 rich areas in the samples may enhance the reactivity of V2O5 ion

with the other oxides that precipitate as a high density of fine V2O5 rich crystals.

The tiny crystals formed in the samples act as heterogeneous nuclei for the

crystallization of the remaining glass. The diffraction data also indicate the

possibility that, in these samples vanadium exists in V4+ state in addition to V5+

state.

4.3.4 Differential thermal analysis

Differential thermal analysis (DTA) scans for Li2O–Nb2O5–ZrO2–SiO2:

V2O5 glass ceramics are recorded in the temperature region 30–1100 oC; however,

in Fig. 4.3 the scans are presented in the temperature region 400-1100oC. All DTA

traces exhibit typical glass transitions with the inflection point between 560–600

oC; it is interesting that glass transitions temperature (Tg) shows decreasing trend

with increase in the content of nucleating agent (inset of Fig. 4.3). At about 950 oC,

the thermogram of each glass ceramic exhibits well-defined principal exothermic

effect along with weak multiple steps due to crystallization, spreading over a

region of approximately 100 oC. The enthalpy associated with the crystallization

peaks with the concentration of crystallizing agent seems to increase with increase

in the concentration of nucleating agent. In the same figure the thermal treatment

schedule used for crystallization is also presented as inset (a). In the inset (b) of

177

5 10 15 20 25 30 35 40

1

2 58

763

4

910

11

3 LiV3O8

4 Li2V2O5

5 LiVO3

6 Li2ZrO3

7 NbVO5

8 LiNbO3

9 Nb6V2O19

10 Nb2V2O9 12 Li2ZrO3 / Li2Si2O5/ ZrSiO4

11 Li2Si2O31 ZrV2O7

2 ZrSi24O50

12

2θ (degrees)

Pre-crystallized V1

V1

V3

V5

V8

V10

Fig. 4.2. XRD pattern of Li2O–Nb2O5–ZrO2–SiO2: V2O5 glass ceramics.

178

Fig. 4.3, the variation of the difference TC2–Tg (where TC2 is the principal

crystallization peak temperature) with the concentration of nucleating agent are

also presented; this parameter is observed to decrease with the content of V2O5

Thus the results of characterization by XRD, SEM and DTA techniques

unambiguously reveals that the thermal treatment for prolonged times with

different concentrations of V2O5 caused the conversion of Li2O–Nb2O5–ZrO2–

SiO2 glass samples into glass ceramics with entrenchment of fine crystals of

different phases.

4.4 Results

4.4.1 Infrared spectrum

The IR spectrum (Fig. 4.4) of V2O5 free Li2O–Nb2O5–ZrO2–SiO2 glass ceramics

exhibited conventional bands due to Si−O−Si (linkages between SiO4 tetrahedral

units) asymmetric and symmetric vibrations at about 1020 cm−1 and at about 800

cm−1 respectively. More specifically, the symmetrical band is designated as

bending mode of bridging oxygen perpendicular to Si–Si axis within the Si−O−Si

plane [79]. A feeble band at about 980 cm-1 attributed to the vibrations of Si–O–Zr

linkages is also located in the spectrum of this glass [80]; in the same region, the

vibrational frequency due to Si−OH stretchings [6, 81] is also expected. A band

correlated to Si−O−Si / O−Si−O bending modes is also located at about 475 cm−1.

The spectra also exhibited two significant bands at about 530 and 700 cm-1 due to

Zr–O–Zr vibrations of ZrO4 structural units. The niobate groups have also

179

400 500 600 700 800 900 1000 1100Temperature (oC)

V1

V3

V5

V8

V10

Exo

Endo

Tc2Tc1

Tc3

520

610

0.1 0.3 0.5 0.8 1.0410

430

450

Tg

o C

Conc. V2O5 (mol%)

Tc2

-Tg

(b)0

800

0 20 40 60 80

72 h

Time in hours

Tem

p. o C Quenching

temp.

(a)

Fig. 4.3. DTA scans of Li2O–Nb2O5–ZrO2–SiO2: V2O5 glass ceramics. Inset (a) represents the thermal treatment schedule used for crystallization. Inset (b) shows the variation of Tg and Tc2–Tg with the concentration of V2O5.

180

exhibited bands due to Nb=O, Nb−O−Nb stretching vibrations at 860 and 600 cm-1

respectively in addition to the band due to the vibrations of NbO6 (3–mode)

structural units (seemed to have merged with Zr–O–Zr vibrational band) at 530

cm-1 [82]. The IR spectrum of crystalline V2O5 is expected to exhibit vibrational

bands at 970 cm-1 (due to V–O stretching of V=O groups), 815 cm-1 (due to V–O–

V stretchings) and at 600 cm-1 (due to V–O–V bending vibrations) [83]. The

bands observed in the spectra of glass ceramic V1 at 980 cm-1, 800 cm−1 can there

be considered as common vibrational modes due to Si–O–V stretchings where as

the band at 600 cm-1 can be considered as common vibrations due to Nb−O−V

chains. As the concentration of V2O5 is increased in the glass ceramics gradually,

the bands due to asymmetrical vibrations of silicate and other structural units are

observed to grow at the expense of symmetrical bands.

4.4.2 Optical absorption spectra

The optical absorption spectra of Li2O–Nb2O5–ZrO2–SiO2: V2O5 glass ceramics

recorded at room temperature in the wavelength region 400−1200 nm is shown in

the (Fig. 4.5). For the vanadium free glass ceramic the absorption edge is observed

at 368 nm (not shown in the figure). With the gradual inclusion of crystallizing

agent V2O5, the absorption edge exhibited red shift. From the observed absorption

edges, the optical band gaps (Eo) of these glass ceramics are evaluated by drawing

400500

600700

800900

10001100

1200

Si-O-Siassymetric stretching

Si-O-Zr / V=O units

Nb=O units

Si-O-Si Symmetric stretching/

V-O-V chains

Zr-O-Zr units/NbO6 otahedral units

Nb-O-Nb / V-O-V bending

NbO6/Zr-O-Zr

Si-O-Si rocking bands

Wavenum

ber (cm-1)

Transmittance %

V0

V8

V5

V3

V1

V10

Fig. 4.4. IR spectra of L

i2 O–N

b2 O

5 –ZrO

2 –SiO2 : V

2 O5 glass ceram

ics

Urbach plots between ( ) ωωα vs2/1 (Fig. 4.6); from the extrapolation of the linear

portion of these plots, the values of optical band gap (Eo) are obtained and

presented in Table 4.2. The value of Eo is found to be the lowest for the glasses

crystallized with 1.0 mol% of V2O5. The spectrum of glass ceramic sample V1

exhibited two broad absorption bands with the meta-centers at 634 and 1040 nm

attributed to 2B22B1 and 2B2

2E transitions of VO2+ ions [84]. A continuous

increase of absorption under these bands is observed with the content of V2O5 in

the glass ceramic.

Table 4.2

Absorption band positions and optical band gaps of Li2O–Nb2O5–ZrO2–SiO2 glasses crystallized with various concentrations of V2O5

Band positions (nm) Sample

Cut off wavelength

(nm) 2B22B1

2B22E

Optical band gap

(E0) eV

V0 380 - - 3.23

V1 382 634 1041 3.18

V3 385 638 1054 3.13

V5 422 659 1062 2.85

V8 445 661 1069 2.75

V10 455 663 1070 2.65

1.0

3.0

5.0

7.0

400 500 600 700 800 900 1000 1100 1200

Wavelength (nm)

Abs

orpt

ion

coef

fici

ent (

cm-1

)

V0

V1

V3

V5

V8

V10

2B2 2B1

2B2 2E

Fig. 4.5. Optical absorption spectra of Li2O–Nb2O5–ZrO2–SiO2: V2O5 glass ceramics recorded at room temperature.

0.0

2.0

4.0

6.0

8.0

10.0

2.0 2.2 2.4 2.6 2.8 3.0 3.2 3.4

( α ω

)1/2

(eV

cm

-1)

ω (eV)

V0V3V5V8V10

2.5

3.0

3.5

0.0 0.2 0.4 0.6 0.8 1.0Conc. V2O5 (mol%)

OB

G (

eV)

V1

Fig. 4.6. Urbach plots of Li2O–Nb2O5–ZrO2–SiO2: V2O5 glass ceramic samples to evaluate optical band gap.

185

4.4.3 Electron spin resonance

The ESR spectra of Li2O–Nb2O5–ZrO2–SiO2: V2O5 glass ceramics recorded

at room temperature are shown in Fig. 4.7; spectra are observed to be complex

made up of resolved hyperfine components arising from unpaired 3d1 electron of

51V isotope. As the concentration of nucleating agent V2O5 is increased, an

increasing degree of resolution and the intensity of the signal have been observed.

The values of ||

g and g⊥ evaluated from these spectra are observed to increase with

increase in the concentration of V2O5 (Table 4.3).

Table 4.3

Data on ESR spectral studies of Li2O–Nb2O5–ZrO2–SiO2: V2O5 glass ceramics.

Sample g|| g⊥ ∆g||/∆g⊥

V1 1.919 1.930 1.152

V3 1.921 1.931 1.140

V5 1.924 1.933 1.134

V8 1.925 1.934 1.131

V10 1.929 1.935 1.094

186

250 300 350 400 450

V1

V3

V5

V8

V10

Magnetic field (mT)

Firs

t der

ivat

ive

of a

bsor

ptio

n ( a

rb. u

nits

)

.Perpendicular spectrum

Parallel spectrum

Fig.4.7. ESR spectra of Li2O–Nb2O5–ZrO2–SiO2: V2O5 glass ceramics recorded at room temperature.

187

4.4.4 Photoluminescence spectra

Fig. 4.8 shows the photoluminescence spectra of Li2O–Nb2O5–ZrO2–SiO2:

V2O5 glass ceramics recorded at room temperature at the excitation wavelength

corresponding to the transition 2B22B1 in the absorption spectra. The spectrum

of each glass exhibited a broad emission band peaking in the region 750-850 nm;

this band is identified due to 2E2T2 transition of vanadyl ion, the band shows

signs of a slight asymmetry in the higher wavelength region. With the growing

content of crystallizing agent V2O5, the half width of the peak is observed to

increase with shifting of meta-centre towards higher wavelength.

4.4.5 Dielectric properties

The dielectric constant ε' and loss tan δ at room temperature ( 30 oC) of

V2O5 free Li2O–Nb2O5–ZrO2–SiO2 pre-crystallized glasses at 100 kHz are

measured to be 6.45 and 0.018 respectively. Fig. 4.9 represents the variation of

dielectric constant and loss with frequency at room temperature of Li2O–Nb2O5–

ZrO2–SiO2 glasses crystallized with different concentrations of V2O5. The

parameters, ε' and tanδ are observed to increase with a decrease in frequency

exhibiting larger values for the samples crystallized with higher content of V2O5 at

any frequency.

The temperature dependence of dielectric constant of Li2O–Nb2O5–ZrO2–

SiO2 glass and glass ceramic samples containing 0.1 and 1.0 mol% of V2O5

measured at different frequencies is presented in Figs. 4.10 (a) and (b) respectively.

188

700 750 800 850Wavelength (nm)

Lum

inio

us in

tens

ity (a

.u.)

2E 2T2

V1

V3

V5

V8

V10

Fig. 4.8. Photoluminescence spectra of Li2O–Nb2O5–ZrO2–SiO2: V2O5 glass ceramics excited at wavelength corresponding to the transition 2B2

2B1

recorded at room temperature.

189

The close look of these figures clearly suggests that the rate of increase of ε with

temperature is higher for the glass ceramic samples when compared with that of

pre–crystallized glass samples. Further, the comparison of these data for the

samples crystallized with different concentrations of V2O5 (Fig. 4.10(c)) clearly

suggest that for the samples crystallized with higher concentrations of V2O5. ε'

increases with temperature at faster rates. The dielectric properties both constant

and loss were measured in the temperature range 30-300 oC while heating and

cooling. During this process no hysteresis loss is observed.

A comparison plot of variation of tan δ with temperature, measured at a

frequency of 10 kHz for Li2O–Nb2O5–ZrO2–SiO2 glasses crystallized with

different concentrations of V2O5 are presented in Fig. 4.11. The inset of this figure

represents the temperature dependence of tanδ for the sample V3 at different

frequencies. These curves have exhibited distinct maxima; with increase in

frequency the temperature maximum of tanδ shifts towards higher temperatures

and with increase in temperature, the frequency maximum shifts towards higher

frequencies, indicating the relaxation character of dielectric losses in these glass

ceramics. From these curves, it is also observed that the region of relaxation shifts

towards lower temperatures (with broadening of relaxation peaks and increasing

value of (tan δ)max) with increase in the concentration of the nucleating agent. The

effective activation energy Wd for the dipoles evaluated from these plots, is found

to decrease gradually with increase in the concentration of the

190

0.005

0.009

0.013

0.017

0.021

0.025

0.029

0.033

100 1000 10000 1000006

10

14

18

22

Frequency (Hz)ε'

tan

δ

V1

V3

V5

V10

V8V1

V3

V5

V10

V8

Fig. 4.9. Variation of dielectric constant and loss with frequency at room temperature of Li2O–Nb2O5–ZrO2–SiO2 glasses crystallized with different concentrations of V2O5.

191

5

10

15

20

25

30

0 50 100 150 200 250 300

Series3

Series4

Temperature (oC)

0

20

40

60

80ε'

1 kHz

1 kHz

10 kHz

10 kHz

100 kHz

100 kHz

(b)

Pre-crystallized V1

Crystallized V1

1 kHz

10 kHz

100 kHz

ε'

10

30

50

70

0 100 200 300Temperature (oC)

ε'

V1

V3

V5

V10

V8

(c)

(a)

Fig. 4.10. The temperature dependence of dielectric constant of Li2O–Nb2O5–ZrO2–SiO2 glass and glass ceramic samples containing (a) 0.1 and (b) 1.0 mol% V2O5 measured at different frequencies (c) comparison of dielectric constant of Li2O–Nb2O5–ZrO2–SiO2 glass ceramics crystallized with different concentrations of V2O5.

192

crystallizing agent (Table 4.4).

The ac conductivity σac is calculated at different temperatures, using the standard

relation [85] and the plots of log σac against 1/T are shown in Fig. 4.12 for all the

glass ceramics at 100 kHz. From these plots, the activation energy for conduction

in the high temperature region over which a near linear dependence of log σac with

1/T could be observed is evaluated and presented in Table 4.4; this activation

energy is also found to decrease gradually with increase in the concentration of the

crystallizing agent.

Table 4.4

Summary of data on dielectric studies of Li2O–Nb2O5–ZrO2–SiO2: V2O5 glass ceramic samples.

A.E. for conduction (eV)

Sample (tan)max.avg.

Temp. region of relaxation

(oC)

A.E. for dipoles

(eV) From σ vs 1/T

From ∆ε vs 1/T

N(EF) (x1020

eV-1/cm3)

V1 0.038 120-140 2.62 0.48 0.49 2.55

V3 0.039 108-135 2.48 0.41 0.44 2.81

V5 0.041 100-130 2.27 0.37 0.41 3.02

V8 0.043 85-115 2.13 0.34 0.37 3.22

V10 0.045 75-110 2.08 0.32 0.32 3.42

193

0.00

0.04

0.08

0.12

0.16

0 100 200 300

Temperature (oC)

Tan

δ

0.00

0.05

0.10

0.15

0 100 200 300

Temperature (oC)

Tan

δ

100 kHz

10 kHz

1 kHz

V1

V3

V5

V10

V8

Fig. 4.11 Comparison plot of variation of tan δ with temperature, measured at a frequency of 10 kHz for Li2O–Nb2O5–ZrO2–SiO2 glasses crystallized with different concentrations of V2O5. Inset shows the variation of Tan δ with temperature at different frequencies for the sample V3.

194

1.7 2.1 2.5 2.9 3.3

σ ac

( Ω- c

m)-1

1/T(10-3 K -1)

V1

V3

V5

V10

V8

0.1 0.3 0.5 0.8 1.00.2

0.6σ a

c (Ω

- cm

)-1

Conc. V2O5 (mol%)

(b)

2x10-5

10-5

0

A.E

. (eV

)

0.3 0.4 0.5A.E. (eV)

σ ac (

Ω- c

m)-1

(a)10-7

10-5

10-5

10-7

0.65

0.7

0.75

0.0 0.5 1.0

Exp

onen

t, s

Conc. V2O5 (mol%)

(c)

Fig. 4.12 Variation of ac conductivity with 1/T at 100 kHz for Li2O–Nb2O5–ZrO2–SiO2: V2O5 glass ceramics. Inset (a) shows the variation of a.c conductivity with activation energy, (b) shows the variation of a.c conductivity and activation energy with concentration V2O5 and (c) shows the variation of the exponent, s with concentration of crystallizing agent V2O5.

195

4.5 Discussion

Among various constituents of Li2O–Nb2O5–ZrO2–SiO2: V2O5 glass

ceramics, SiO2 is one of the most common glass-formers and participates in the

glass network with tetrahedral [SiO4/2]0 units and all the four oxygens in SiO4

tetrahedral are shared. On addition of modifiers like Li2O and vanadyl ions, the Si–

O–Si linkage is broken and form Si–O− termination. Thus, the structure is

depolymerised or modified. The modification, results in the formation of meta,

pyro and ortho-silicates in the order: [SiO4/2]0, [SiO3/2O]−, [SiO2/2O2]2−,

[SiO1/2O3]3− and [SiO4]

4− which are designated as Q4, Q3, Q2, Q1 and Q0,

respectively [86].

Nb2O5 belongs to the intermediate class of glass forming oxides; it is an

incipient glass network former and as such does not readily form the glass due to

the fact that the Nb–O bonds in Nb2O5 polyhedron, are highly rigid when

compared with those in the conventional glass forming oxides like B2O3, GeO2 but

does participate in the glass forming in the presence of modifiers like Li2O. In the

glasses containing Nb2O5, the niobium ions are often incorporated into the glass

matrix as NbO6 octahedra with well resolved features. This is also quite clear from

the IR spectra. The band observed in the spectral region 850-860 cm-1 is an

indicative of the presence of isolated NbO6 distorted octahedra with one terminal

(non-bridging) Nb–O bond [87]; secondly the band at 700 cm-1 is the suggestive of

formation of chains composed by corner shared NbO6 octahedra and yet, another

band observed at 600 cm-1 reveals the formation of a three-dimensional (3D)

196

network by less distorted corner shared NbO6 octahedra [88]. The close look at

various band positions in the IR spectra points out that the linkages of the type

Nb–O–Zr, Nb–O–V are also quite possible in these glass ceramics. Zirconium ions

in general do participate in the glass network with ZrO4 structural units and

alternate with SiO4 structural units and also forms the linkages with NbO6 units as

mentioned above.

The entry of Li+ ions, may modify these linkages as per the following relations:

Nb-O-Zr + Li2O Nb-O-Li+ + Zr-O-Li+

Si-O-Zr + Li2O Si-O-Li+ + Zr-O-Li+.

As a consequence we expect, there is a disruption in the SiO4 and ZrO4 tetrahedral

linkages with the creation of number of dangling bonds and non-bridging oxygens

in the glass ceramic.

Vanadium ions are expected to exist mainly in V5+ states in the Li2O–

Nb2O5–ZrO2–SiO2: V2O5 glass ceramic. However, during the melting, annealing

and crystallization processes it is quite certain for the following redox equilibrium

to take place:

2V5++O2- 2V4++1/2 O2

The V5+ ions take part network forming positions with VO5 structural units

(in fact in the IR spectra a band due to the vibrations of V=O groups in VO5

trigonal bipyramids is observed at about 978 cm-1) and form linkages of the type

V–O–Si (with SiO4) and V–O–Zr (with ZrO4 structural units). The V4+ ions form

vanadyl complexes (VO2+), act as modifiers and distort the glass network similar

197

to lithium ions. In this type of glass ceramics there is also a possibility to reduce a

small fraction of V5+ ions to diamagnetic V3+ ions [89].

The progressive introduction of crystallizing agent V2O5 caused a slight

increase in the density (or decrease in molar volume) of the glass ceramic. The

degree of structural compactness, the modification of the geometrical

configuration of the glassy network, the size of the micro-crystals formed, changes

in the coordination of the glass forming ions and the fluctuations in the dimensions

of the interstitial holes are the some of the factors that are responsible for the

observed fluctuations in the density.

The formation of Li2V2O5, Nb6V2O19, Nb2V2O9 crystalline phases detected

from the XRD studies emphasizes the presence of vanadium ions in V4+ state in

addition to V5+ state in these glass ceramics. The relative increase in the intensity

of the diffraction peaks due to these crystallites also indicates an increasing

concentration of vanadyl complexes with the increase in the concentration of

nucleating agent in the glass matrix.

The appearance of peaks due to multiple exothermic effects in the DTA

pattern advocates the presence of different phases of crystallization in the samples.

The crystallization in the glass samples may take place following the surface and

bulk nucleations. The general shape of the crystallization peak in DTA curves

reflects the variation of enthalpy. The observed increase in value of enthalpy (area

under the crystallization peaks) with increase in the nucleating agent suggests that

198

the crystallization starts initially inside the material and expands to the surface

gradually [90].

On the basis of energy level scheme provided by Ballhausen and Gray [84],

for molecular orbitals of VO2+ (3d1) ion in a ligand field of C4v symmetry situated

in an octahedral field with tetragonal distortion, three bands corresponding to the

2B22B1 (∆ ), 2B22E (∆||) and 2B22A1 transitions were predicted in the

absorption spectra. In the spectra of the investigated glass ceramics, only the first

two bands are observed. As the content of the nucleating agent V2O5 continues to

increase, a gradual growth of these bands with a slight shift of the meta-centre

towards higher wavelength could clearly be seen; this observation is an evocative

of increase in the rate of reduction of V5+ ions to VO2+ (vanadyl) ions in the glass

ceramic. Further, the optical activation energy associated with 2B2→2B1 is

decreased from 1.95 (sample V1) to 1.85 eV (sample V10) with the increase in the

content of crystallizing agent; this is clearly a characteristic signal of inter valence

transfer or a polaronic type of absorption. This is possible when the associated

electrons are trapped at shallow sites within the main band gap with smaller wave-

function radii; in terms of polaronic perception, this kind of situation is only

possible if the local potential fluctuation is small as compared to the transfer

integral, j. A small overlap between electronic wavefunctions (corresponding to

adjacent sites) due to strong disorder is contributive to polaron formation. So from

the polaronic viewpoint, the electron delivered by the impurity atom at the V5+ site

converts this into a lower valence state V4+, and at the next stage, the trapped

199

electron at this V4+ site is transferred to the neighboring new V5+ site by absorbing

a photon energy. Thus the optical absorption in the glass ceramic samples is

dominated by polaronic transfer between the V4+ and V5+ species [91].

The VO2+ ions are expected to participate in the depolymerisation of the

glass ceramic network, create more bonding defects and non-bridging oxygens

(NBOs) as mentioned earlier. With the increase in concentration of vanadyl ions in

the glass ceramic network, a large number of donor centers are created, and

subsequently, the excited states of localized electrons originally trapped on V4+

sites begin to overlap with the empty 3d states on the neighboring V5+ sites, and as

a result, the impurity or polaron band becomes more extended into the main band

gap. This new polaronic development might have shifted the absorption edge to

the lower energy (Table 4.2) which leads up to a significant shrinkage in the band

gap as the concentration of V2O5 is increased. Further, it may be worth mentioning

here that the interactions between the ligand preliminary 2pO delocalized states

and the highly localized d-levels of the vanadium ions might also contribute to

such charge transfer.

Excitation of Li2O–Nb2O5–ZrO2–SiO2: V2O5 glass ceramic samples with

the wavelength corresponding to 2B22B1, resulted a broad emission band (Fig.

4.8). Since the wavelength of this band is close to the maximum of the band

2B22E, this band is attributed to 2E2T2 transition of V4+ ions; the emission

band is relatively broad and structure less. With increase in the concentration of

V2O5, the intensity of the peak is observed to increase with a red shift. The shift of

200

this PL peak, the shape and the structured nature of the PL emission band are a

signature of shallow levels with an electron–phonon coupling. The distortion of

the luminescence band in the lower energy side is probably due to the reabsorption

by V3+ ions if any in the glass network [92]. Overall the increase in the intensity of

the band with increase in the concentration of V2O5 suggests the increasing

presence of vanadyl ions in the glass ceramic samples.

The ESR spectra observed for the titled glass ceramics by and large

resembles the spectra of previously studied glasses [13, 14]. The spectra consist of

the well-resolved hyperfine structure with eight components of the electron-

nuclear interactions (corresponding to eight values of nuclear magnetic quantum

numbers: MI = –7/2, –5/2, . . . , +7/2 in accordance with selection rules ∆MI = 0

and ∆Ms = ±1. The variations in the ESR line-shape and the intensity in the other

glass systems have been previously explained by a number of investigators on the

basis of variations in the concentration of V4+ ions. In the present case also, the

change in the V4+/Vtot ratio seems to be one of the reasons for the variations in the

line-shape. Additionally other factors mentioned below will also contribute for

such variations in the spectra. In the glass ceramic samples the jumping frequency

of the charge carriers (V4+V5+) is proportional to exp(–W/kT) [93], where

W=1/2 WD + WH; in this, WD is the mean energy difference between adjacent

vanadium sites due to the disordered nature of the glass ceramic and WH is the

activation energy for the, hopping process of the polarons between two identical

sites. In the case of samples containing low content of V2O5, it is possible to

201

suppose that the leading term is WD, and that the jumping rate of the polaron is

low; this fact, together with the relative low concentration of the paramagnetic

species accounts for the weak ESR signal in samples V1 and V2. It may be noted

here that though the jumping frequency of polarons depends on temperature, this

dependence is hardly expected in the present case [94]. The quantitative analysis

of ESR results indicates that the ratio ⊥

∆∆ gg||

is observed to decrease (Table 4.3)

gradually with concentration of V2O5, indicating an increasing degree of distortion

(elongation) of the VO6 octahedron.

Recollecting the data on dielectric properties of Li2O–Nb2O5–ZrO2–SiO2:

V2O5 glass ceramics, with the gradual increase of V2O5, the values of ', tan δ and

σac are found to increase at any frequency and temperature and the activation

energy for a.c conduction is observed to decrease; this observation indicates an

increase in the space charge polarization owing to the enhanced degree of disorder

in the glass network. In other words, as the concentration of V2O5 is increased in

the glass network, there is a growing presence of V4+ ions in the glass network (as

is evidenced from the spectroscopic studies); these ions similar to Li+ ions disrupt

the glass network by creating dangling bonds and non-bridging ions The defects

thus produced create easy path ways for the migration of charges that would build

up space charge polarization leading to the increase in the dielectric parameters as

observed [95, 96].

202

The variation of ' with temperature can be connected to frequency through

modified Debye equation as reported earlier [97]. A plot of log (') against 1/T

(where is the difference between dielectric constant at any temperature T and

that at room temperature) at 1 kHz for all the glasses in the high temperature

region is shown in Fig. 4.13; the graphs obtained are straight lines in the high

temperature region. The computed activation energies are furnished in Table 4.4;

this value is practically the same as the activation energy for ac conduction in the

same temperature region for these glass ceramics. This seems to suggest that the

charge carriers responsible for change in ' and ac with temperature in this

temperature region are the same.

The variation of dielectric loss with the temperature for Li2O–Nb2O5–ZrO2–

SiO2: V2O5 glass ceramics exhibited the relaxation character. Earlier studies on

variety of glasses containing d1 ions like W5+, Cr5+, Mo5+, Ti3+ etc., showed that

these ions contribute to the dielectric relaxation effects [15–18]; hence, the

observed relaxation effects in the present glass ceramic samples can safely be

attributed to V4+ ions. The increase in the breadth and the intensity of the

relaxation peaks also supports the view point that there is a gradual increase in the

concentration of vanadyl ions (which participate in the relaxation effects) with

increase of content of V2O5 in these glass ceramics. Such variations of loss tangent

with the temperature also indicate that there is a spreading of relaxation times of

dipoles; this may be understood due to experience of an approximately random

potential energy on diffusing through the distorted structure of Li2O–Nb2O5–

203

ZrO2–SiO2: V2O5 glass ceramic network by the dipoles [98]. The decreasing value

of activation energy for dipoles with increase in the content of the crystallizing

agent V2O5 (Table 4.4), suggests an increasing degree of freedom for dipoles to

orient in the field direction.

1

10

100

1.75 1.85 1.95 2.05 2.15

1/T(10-3 , K-1)

∆ε'

V1

V3

V5

V10

V8

Fig. 4.13 Plot of log (') vs 1/T.

204

When a plot is made between log σ(ω) vs activation energy for conduction (in the

high temperature region) a near linear relationship is observed (inset (a) of Fig.

4.12); this observation suggests that the conductivity enhancement is directly

related to the thermally stimulated mobility of the charge carriers in the high

temperature region. The increase of conductivity and decrease in activation energy,

with V2O5 content (inset (b) of Fig. 4.12) is observed for the studied glass

ceramics; a number of independent earlier studies have confirmed that such

behavior represents the hopping process of conduction. In the present context, the

decreasing electronic hopping distance between V4+ and V5+ ions can be assumed

to be responsible for increase of conductivity with the content of V2O5. However,

the contribution from the movement of lithium ions to the increase in the

conduction should also be taken into consideration. This type of composition

dependence of conductivity is quite conventional in the glasses containing

monovalent cations like Li+ and transition metal ions like vanadium. The

conductivity is observed to increase slowly with increase in the concentration of

V2O5 up to 0.3 mol% and beyond this concentration the conductivity is observed

to increase at a faster rate (inset (b) of Fig. 4.12). The polarons involved in the

process of transfer from V4+ to V5+, are attracted by the oppositely charged Li+

ions. This cation-polaron pair moves together as a neutral entity. As expected, the

migration of this pair inhibits the net displacement of the charge and hence a lower

magnitude of increase of conductivity could be observed in the samples containing

low concentration of V2O5.

205

The frequency response of real part of ac conductivity is normally

described by power law dependence with s as exponent:

10],)/(1[)(' <≤+= sscdc ωωσωσ (4.1)

ωC is the characteristic macroscopic relaxation frequency

Within the framework of the linear-response theory, the frequency-dependent

conductivity can be related to

dtetrkTHnq ti

R

c ωωωσ −∞

−=0

222

)(6

)( (4.2)

where q is the charge, nc is the mobile ion density, )(2 tr is the mean square

displacement of the mobile ions and HR is the Haven ratio (lies in between 0.2 and

1.0) [99] which represents the degree of correlation between successive hops. A

more detailed description of the frequency response of ac conductivity can be

found in Ref [100].

At short times, when the mean square displacement )(2 tr of ions is small then it

is ~ t1-s, the ion transport is characterized by the non-random forward-backward

hopping process, under these conditions Eq. (4.1) modifies to

σ (ω) ∝ ωs

The variation of the exponent (obtained by plotting log σ (ω) vs ω) is slightly

increased with increasing content of crystallizing agent V2O5 (inset (c) of Fig.

4.12). Such increase suggests that dimensionality of conduction space increases in

proportion to the content of V2O5 in the glass ceramic [101, 102].

206

The a.c conductivity in the low temperature region (where the conductivity

is nearly temperature independent) can be understood based on quantum

mechanical tunneling model. Based on Austin and Mott’s model (quantum

mechanical tunneling model) [103], the density of the energy states near the Fermi

level, N(EF), at nearly temperature independent region of the conductivity (low

temperature) is evaluated using

σ(ω) = (π/3)e2KT [N(EF)]2 α-5ω [ ln(νo/ω) ]4 (4.3)

where α is the electronic wavefunction decay constant (obtained by plotting log

σac against inter ionic distance Ri) and νph ~ 5x1012 Hz, ν0 is the phonon frequency,

and presented in the Table 4.4. The value of N(EF) is found to increase gradually

from the sample V1 to V10, indicating a growing degree of disorder with increase

in the content of V2O5 in the glass network.

Our observations on dielectric parameters of Li2O–Nb2O5–ZrO2–SiO2:

V2O5 glass ceramics as mentioned earlier, indicate, the rate of increase of

tan (which is inversely proportional to breakdown strength [85]) with

temperature is increasing with the content of V2O5. Thus the experiments on

dielectric properties of these glass ceramics also reveal that there is a decrease in

the dielectric breakdown strength with increase in the concentration of

crystallizing agent V2O5. These revelations are also consistent with the view that

there is a gradual increase in the concentration of V4+ ions that act as modifiers

and decrease the insulating strength of the glass ceramic samples.

207

4.6 Conclusions

Multi component Li2O–Nb2O5–ZrO2–SiO2 glasses have been crystallized

with different concentrations of V2O5 as nucleating agent. The characterization of

the samples by SEM, XRD and DTA techniques have indicated that the samples

contain well defined and randomly distributed grains of different crystalline

phases. The IR spectral studies indicated that the glass ceramic samples contains

various structural units with the linkages of the type Si-O-Si, Nb−O−Nb, Zr–O–Zr,

Si–O–V, Nb−O−V; the increasing content of V2O5 in the glass ceramics seemed to

have weakened such linkages. The analysis of the results of optical absorption,

ESR and photoluminescence spectra of the studied glass ceramics have indicated

that a considerable proportion of vanadium ions do exist in V4+ state in addition to

V5+ state, and the redox ratio increases with increase in the concentration of

crystallizing agent V2O5. The analysis of the results of dielectric studies has

suggested a decrease in the insulating character with increase in the crystallizing

agent of these samples. A.C. conductivity in the high temperature region seems to

be connected mainly with the polarons involved in the process of transfer from V4+

to V5+. The low temperature (or the nearly temperature independent) part of

conductivity could be explained on the basis of quantum mechanical tunneling

model.

208

References

[1] A.A. Lipovskii, Y. Kaganovskii, V.G. Melehin, D.K. Tagantsev, O.V. Yanush, J. Non–Cryst. Solids 354 (2008) 1245.

[2] C.H. Song, M. Kim, H.W. Choi, Y.H. Kim, Y.S. Yang, J. Kor. Phys. Soc. 55 (2009) 846.

[3] T. Honma, D. Oku, T. Komatsu, Solid State Ionics 180 (2009) 1457.

[4] I.W. Donald, B.L. Metcalfe, R.N.J. Taylor, J. Mater. Sci. 32 (1997) 5851.

[5] T.L. White, W.D. Bostic, C.T. Lson, C.R. Schaich Work Shop on Vitrification of Low Level Waste: The Process and Potential, San Antonio, TX, USA, 1995.

[6] R.R. Goncalves, J.J. Guimaraes, J.L. Ferrari, L.J.Q. Maia, S.J.L. Ribeiro, J. Non–Cryst. Solids 354 (2008) 4846–4851.

[7] K. Fukumi, S. Sakka, J. Mater. Sci. 23 (1988) 2819.

[8] L.F. Santos, L. Wondraczek, J. Deubener, R.M. Almeida, J. Non-Cryst. Solids 353 (2007) 1875.

[9] N. Krishna Mohan, G. Sahaya Baskaran, N. Veeraiah, Phys. Stat. Solidi (a) 203 (2006) 2083.

[10] M.A. Valente, C.C. Silva, A.S.B. Sombra, M.P.F. Graça, J. Non-Cryst. Solids 356 (2010) 800.

[11] S. Banijamali, B.E. Yekta, H.R. Rezaie, V.K. Marghussian, Thermochim. Acta 488 (2009) 60.

[12] R. R. Gonçalves, Y. Messaddeq, A. Chiasera, Y. Jestin, M. Ferrari, S.J.L. Ribeiro, Thin Solid Films 516 (2008) 3094.

[13] I. Nicula, E. Culea, I. Lupsa, J. Non-Cryst. Solids 79 (1986) 325.

[14] Y. Li, K. Liang, J. Cao, B. Xu, J. Non-Cryst. Solids 356 (2010) 502.

[15] S. Mukherjee, A.K. Pal, J. Phys. : Cond. Matt. 20 (2008) 255202.

[16] A. Gajovi, A. Šanti, I. Djerdj, N. Tomaši, A. Moguš-Milankovi, D. Sheng Su, J. Alloys Compd. 479 (2009) 525.

209

[17] K. Sambasiva Rao, M. Srinivasa Reddy, V. Ravikumar, N. Veeraiah, Mater. Chem. Phys. 111 (2008) 283.

[18] J. Byung-Hae, H. Seong-Jin, K. Hyung-Sun, J. Eur. Cer. Soc. 25 (2005) 3187.

[19] G. Murali Krishna, B. Anila Kumari, M. Srinivasa Reddy, N. Veeraiah, J. Solid State Chem. 180 (2007) 2747.

[20] A. Agarwal, A. Sheoran, S. Sanghi, V. Bhatnagar, S.K. Gupta, M. Arora, Spectrochim. Acta - Part A 75 (2010) 964.

[21] X. Lin-Hua, Z. Guo-Ping, Q. Min, Physica B 405 (2010) 2213.

[22] N. Kerkouri, M. Et-Tabirou, A. Chahine, A. Mazzah, M.C. Dhamelincourt, M. Taibi, J. Optoelectr. Adv. Mater. 12 (2010) 1030

[23] M.E. Gouda, H. Khodair, M.G. El-Shaarawy, Mater. Chem. Phys. 120 (2010) 608.

[24] H. Behzad, M.H. Hekmatshoar, M. Mirzayi, M. Azmoonfar, Ionics 15 (2009) 647.

[25] S.A.E., Ali, F.M. Ezz-Eldin, Nucl. Instr. Meth. Phys. Res., Sec. B 268 (2010) 49.

[26] N.S. Abd El-Aal, H.A. Afifi, Archiv. Acoust. 34 (2009) 641.

[27] G.D. Khattak, A. Mekki, L.E. Wenger, J. Non-Cryst. Solids 355 (2009) 2148.

[28] A. Mekki, G.D. Khattak, L.E. Wenger, J. Electr. Spectr. Rel. Phen. 175 (2009) 21.

[29] L.D. Bogomolova, V.A. Zhachkin, T.K. Pavlushkina, Glass Ceram. 66 (2009) 168.

[30] Y.B. Saddeek, Phil. Mag. 89 (2009) 2305.

[31] W.L. Feng, Phil. Mag. 89 (2009) 1391.

[32] M. Shapaan, E.R. Shabaan, A.G. Mostafa, Physica B 404 (2009) 2058.

[33] Y.V. Kartashov, V.A. Vysloukh, Opt. Lett. 34 (2009) 1228.

[34] W. Sung, J. Won, J. Lee, H. Kim, Mol. Cryst. Liq. Cryst. 499 (2009) 234.

210

[35] F.H. ElBatal, Y.M. Hamdy, S.Y. Marzouk, Mater. Chem. Phys. 112 (2008) 991.

[36] S.H. Kim, O.H. Han, J.P. Kang, S.K. Song, Bull. Korean Chem. Soc. 30 (2009) 608.

[37] H.M.M. Moawad, H. Jain, R. El-Mallawany, J. Phys. Chem. Solids 70 (2009) 224.

[38] Y. Taibi, M. Poulain, R. Lebullenger, L. Atoui, M. Legouera, J. Optoelectr. Adv. Mater. 11 (2009) 34.

[39] H. Farah, J. Am. Ceram. Soc. 91 (2008) 3915.

[40] N. Vedeanu, O. Cozar, I. Ardelean, B. Lendl, D.A. Magdas, Vibr. Spectr. 48 (2008) 259.

[41] M.P. Kumar, T. Sankarappa, A.M. Awasthi, Physica B 403 (2008) 4088.

[42] V. Kundu, R.L. Dhiman, D.R. Goyal, A.S. Maan, J. Optoelectr. Adv. Mater. 10 (2008) 2765.

[43] L. Srinivasa Rao, M. Srinivasa Reddy, M. Rami Reddy, N. Veeraiah, J. Alloys Compd. 464 (2008) 472.

[44] C. Li, Y. Huang, Z. Cui, X. Gao, Z. Gu, J. Chin. Ceram. Soc. 36 (2008) 1288.

[45] S. Rada, M. Culea, M. Rada, E. Culea, J. Mater. Sci. 43 (2008) 6122.

[46] I. Ardelean, O. Cozar, N. Vedeanu, D. Rusu, C. Andronache, J. Mater. Sci. 18 (2007) 963.

[47] M.S. Al-Assiri, Physica B 403 (2008) 2684.

[48] Q.B. Tian, Y. Wang, X.T. Yue, Y.S. Yin, Mater. Sci. Tech. 16 (2008) 543.

[49] I.Z. Hager, Mater. Chem. Phys. 109 (2008) 365.

[50] J.K. Jung, S.K. Song, T.H. Noh, O.H. Han, J. Non-Cryst. Solids 270 (2000) 97.

[51] A. Mekki, G.D. Khattak, D. Holland, M. Chikhota, L.E. Wenger, J. Non-Cryst. Solids 318 (2003) 20.

211

[52] A.M. Ferrari, C. Leonelli, G.C. Pellacani, C. Siligardi, J. Non-Cryst. Solids 315 (2003) 77.

[53] Abd EI-Moneim, Mater. Chem. Phys. 73 (2002) 318.

[54] U. Hoppe, E. Yousef, C. Russel, J. Neuefeind, A.C. Hannon, Solid State Commun.,123 (2002) 273.

[55] O. Cozar, I. Ardelean, I. Bratu, V. Simon, C. Craciun , L. David, C. Cefan, J. Mol. Struct. 563 (2001) 421.

[56] M.A. Salim, P.S. Fodor, L.E. Wenger, J. Non-Cryst. Solids 85 (2001)195.

[57] G.D. Khattak, M.A. Salim, L.E. Wenger, A.H. Gilani, J. Non-Cryst. Solids 262 (2000) 66.

[58] K. Krasowski, M. Wasiucionek, Phys. Stat. Solidi (a) 181 (2000)157.

[59] Maria-Cameljaungureanu, Michri Levy, Jeal-louis Souquet, J. Ceramics. Silikaty 44 (2000) 81.

[60] V. Sudarsan, S.K. Kulshreshtha, J. Non-Cryst. Solids 258 (1999) 20.

[61] K. Sega, H. Kasai, H. Sakata, Mater. Chem. Phys. 53 (1998) 28.

[62] H. Mori, H. Matsuno, H. Sakata, J. Non-Cryst. Solids 276 (2000) 78.

[63] V. Rajendranan, N. Palanivelu, D.K. Modak and B.K. Chaudhuri, Phys. Stat. Solidi (a) 180 (2000) 467.

[64] J. Simockova, P. Miklos, V. Saly, J. Acta Phys. Slovaca, 50 (2000) 685.

[65] H. Takahashi, Y. hakai, Y. Morii, Solid State Ionics 90 (1996) 125.

[66] Y.M. Moustafa, I.A. Gohar, A.A. Megahed and E.L. Mansour, Phys. Chem. Glasses 38 (1997) 92.

[67] O. Attos, M. Massot, M. Aspomoza, J. Non-Cryst. Solids 210 (1997) 163.

[68] L. Murawski, R.J. Barczynski, J. Non-Cryst. Solids 196 (1996) 275.

[69] V.P. Seth, D. Prakash, P. Chand, J. Non-Cryst. Solids 204 (1996) 46.

[70] D. Prakash, V.P. Seth, I. Chand, Prem Chand, J. Non-Cryst. Solids 204 (1995) 46

[71] M. Amano, K. Suzuki, H. Sakata, J. Mater. Sci. 32 (1997) 4325.

212

[72] L.D. Bogomolva, N.A. Krasilnikova, V.L. Bogdanov, V.D. Khalilev, J. Non- Cryst. Solids 188 (1995) 130.

[73] T. Yoko, S. Hayakawa, S. Sakka, J. Non-Cryst. Solids 183 (1995) 73.

[74] A. Ghosh, S. Mandal, S. Hazra, D. Dass, J. Non-Cryst. Solids 183 (1995) 317.

[75] S. Gupta, A. Mansingh, J. Non-Cryst. Solids 181 (1995) 58.

[76] S. Adams, K. Hariharan, J. Maier, Solid State Ionics 75 (1995) 193.

[77] V. Demitrov, Y. Dimitriev, A. Montenero, J. Non-Cryst. Solids 180 (1994) 51.

[78] Powder Diffraction File, Alphabetical Index, Inorganic Compounds 2003, Published by JCPDS – International Centre for Diffraction Data, Newtown Square, PA. 19073.

[79] M. Nakamura, Y. Mochizuki, K. Usami, Y. Itoh, T. Nozaki, Sol. State Commun. 50 (1984) 1079.

[80] T. Uma, M. Nogami, J. Membr. Sci. 334 (2009) 123.

[81] H. Aguiar, J. Serra, P. González, B. León, J. Non-Cryst. Solids 355 (2009) 475.

[82] Y. Yu, X. Wang, Y. Cao, X. Hu, Appl. Surf. Sci. 172 (2001) 260.

[83] Y. Gandhi, N. Venkatramaiah, V. Ravikumar, N. Veeraiah, Physica B 404 (2009) 1450.

[84] J. Ballhausen, H.B. Gray, Inorg. Chem. 1 (1962) 111.

[85] B. Tareev, Physics of dielectric Materials, (Mir Publishers, Moscow, 1979).

[86] K.J. Rao, Structural Chemistry of Glasses (Elsevier, Amsterdam, 2002).

[87] T.V. Bocharova, A.N. Vlasova, G.O. Karapetyan, A.M. Mironov, Inorg. Mater. 46 (2010) 74.

[88] T. Cardinal, E. Fargin, G. Le Flem, S. Leboiteux, J. Non-Cryst. Solids 222 (1997) 228.

[89] C.H. Chung, J.D. Mackenzie, J. Non-Cryst. Solids 42 (1980) 357.

213

[90] A. Marotta, A. Buri, F. Branda, J. Mater. Sci. 16 (1981) 341.

[91] O. Cozar, D.A. Magdas, I. Ardelean J. Non-Cryst. Solids 354 (2008) 1032.

[92] W. Ryba-Romanowski, S. Golab, G. Dominiak-Dzik, M. Berkowski J. Alloys Compd. 288 (1999) 262.

[93] N.F. Mott, J. Non-Cryst. Solids 1 (1968) 1.

[94] F. Momo, A. Sotgiu, E. Baiocchi, M. Bettinelli, A. Montenero J. Mat 17 (1982) 3221.

[95] Y. Gandhi, K.S.V. Sudhakar, T. Satyanarayana, N. Veeraiah, Mater. Chem. Phys. 120 (2010) 89.

[96] T. Satyanarayana, I.V. Kityk, M. Piasecki, P. Bragiel, M.G. Brik, Y. Gandhi, N. Veeraiah, J. Phys. – Cond. Matt. 21 (2009) 245104.

[97] D.K. Durga, N. Veeraiah, J. Mater. Sci. 36 (2001) 5625.

[98] S.R. Elliott, Physics of Amorphous Materials (Longman, Essex, 1990).

[99] H. Kahnt, J. Non-Cryst. Solids 203 (1996) 225.

[100] K. Funke, R.D. Banhatti, S. Brückner, C. Cramer, C. Krieger, A. Mandanici, C. Martiny and I. Ross, Phys. Chem. Chem. Phys. 4 (2002) 3155.

[101] D.L. Sidebottom, Phys. Rev. Lett. 83 (1999) 983.

[102] S. Bhattacharya, A. Ghosh, Phys. Rev. B 70 (2004) 172203.

[103] I.G. Austin, N.F. Mott, Adv. Phys. 18 (1969) 41.