Journal of Non-Crystalline Solids 355 (2009) 193–202

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Journal of Non-Crystalline Solids

journal homepage: www.elsevier .com/locate / jnoncrysol

Effect of BaO on the crystallization kinetics of glasses alongthe Diopside–Ca-Tschermak join

Ashutosh Goel a, Dilshat U. Tulyaganov b, Ishu K. Goel a, Essam R. Shaaban c, José M.F. Ferreira a,*

a Department of Ceramics and Glass Engineering, University of Aveiro, CICECO, Campus Santiago, 3810-193 Aveiro, Portugalb State Committee of Geology and Mineral Resources, Centre of Remote Sensing and GIS Technologies, 11-A, Shevchenko Street, 100060 Tashkent, Uzbekistanc Physics Department, Faculty of Science, Al-Azhar University, Assuit 71542, Egypt

a r t i c l e i n f o a b s t r a c t

Article history:Received 14 June 2008Received in revised form 11 October 2008Available online 11 December 2008

PACS:61.43.Fs64.70.ph64.70.dg68.37.Hk

Keywords:CrystallizationGlass–ceramic sealantsNucleationX-ray diffractionSOFCFTIR measurementsAluminosilicatesGlass transition

0022-3093/$ - see front matter � 2008 Elsevier B.V. Adoi:10.1016/j.jnoncrysol.2008.10.016

* Corresponding author. Tel.: +351 234 370242; faxE-mail address: jmf@ua.pt (J.M.F. Ferreira).

We report on the effect of BaO on the crystallization kinetics of glasses in the diopside (CaMgSi2O6)–Ca-Tschermak (CaAl2SiO6) system. Partial substitution (i.e. 5%, 10% and 20%) of Ba2+ for Ca2+ was attemptedin composition CaMg0.8Al0.4Si1.8O6, in three different glasses while partial substitution of B3+ for Al3+ wasmade in the fourth glass. Structural investigations on the glasses have been made by density measure-ments, molar volume and Infra-red spectroscopy (FTIR). Non-isothermal crystallization kinetic studieshave been employed to study the mechanism of crystallization in all the four glasses. The Avrami param-eter for the glass powders is �2, indicating the existence of intermediate mechanism of crystallization.Crystallization sequence in the glasses has been followed by X-ray diffraction (XRD) analysis, scanningelectron microscopy (SEM) and FTIR. Augite crystallized out being the dominant phase in all the glass–ceramics, while different polymorphs of BaAl2Si2O8 were present as secondary or minor phases.

� 2008 Elsevier B.V. All rights reserved.

1. Introduction

The pyroxenes are a group of important rock-forming silicateminerals found in many igneous and metamorphic rocks. Theyshare a common structure comprised of single chains of silica tet-rahedra and they crystallize in the monoclinic (clinopyroxenes)and orthorhombic systems. Pyroxenes have the general formulaXY(Si,Al)2O6 (where X represents calcium, sodium, iron(II) andmagnesium and more rarely zinc, manganese and lithium and Yrepresents ions of smaller size, such as chromium, aluminum,iron(III), magnesium, manganese, scandium, titanium, vanadiumand even iron(II)). Although aluminum substitutes extensively forsilicon in silicates such as feldspars and amphiboles, the substitu-tion occurs only to a limited extent in most pyroxenes [1]. Diopside(CaMgSi2O6, hereafter briefly designated as Di) and Ca-Tschermak(CaAl2SiO6, hereafter briefly designated as Ca-Ts) are typical repre-sentatives of the clinopyroxene group, where the former has only

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silicon in the tetrahedral sites while the latter has got these sitesoccupied by both Si and Al [2]. The structural features of glass–ceramics (GCs) based on clinopyroxenes, and particularly on solidsolutions of Di and Ca-Ts, allow good control of their propertiesdue to which they are potential candidates for various functionalapplications [3–5].

In our previous study [6], four glasses were derived from thecomposition CaMg0.8Al0.4Si1.8O6 (Di/Ca-Ts = 80/20 mol.%) with par-tial substitution (i.e. 5%, 10% and 20%) of Ba for Ca in three glassesand B for Al in the fourth glass. These glasses were studied for theircrystallization behavior and properties in order to investigate theirsuitability as sealants for solid oxide fuel cells (SOFC). The powderprocessing route was employed to study the sintering and crystal-lization behavior of the respective glasses. Well sintered GCs com-prising of Augite (Ca (Mg0.85Al0.15)((Si1.70Al0.30)O6)) along with twodifferent polymorphs of BaAl2Si2O8 i.e. monocelsian (hereafter des-ignated as C), hexacelsian (hereafter designated HC) were obtainedafter sintering between 850 (1123 K) and 1000 �C (1273 K). Augitethat was obviously Fe-free, acted as an intermediary member be-tween Di and Ca-Ts, representing a mineral midway between these

Table 2Properties of the glasses.

BaCMAS1 BaCMAS2 BaCMAS3 BaCMAS4

Kth 0.595 0.597 0.600 0.599q (g cm�3) 2.90 ± 0.007 2.94 ± 0.003 3.03 ± 0.001 3.01 ± 0.002Vm (cm3 mol�1) 20.16 ± 0.04 20.31 ± 0.03 20.33 ± 0.01 20.40 ± 0.01Vo (cm3 mol�1) 13.48 ± 0.03 13.58 ± 0.02 13.67 ± 0.01 13.64 ± 0.01Ve 0.82 ± 0.04 0.75 ± 0.03 0.51 ± 0.01 0.54 ± 0.01Tdg (K) 927 950 954 915Tc

a (K) 1172 1180 1181 1157Tp

a (K) 1197 1204 1203 1177Avrami parameter, <n> 2.03 ± 0.04 2.12 ± 0.02 2.10 ± 0.05 2.11 ± 0.02Ec (kJ mol�1) 437.12 426.43 378.228 400.477

a b = 5 �C/min.

194 A. Goel et al. / Journal of Non-Crystalline Solids 355 (2009) 193–202

two minerals along this series, where Al3+ occupies both octahedral(AlO6) and tetrahedral (AlO4) positions in the structure. The den-sity of the GCs increased and the bending strength decreased withthe increase of BaO content. The electrical conductivity was shownto increase with the BaO content; however, all the GCs exhibitedgood insulating properties. In conjunction with the above men-tioned properties, the low level of oxygen permeation fluxes underan air/(H2 + H2O + N2) gradient, the negligible interfacial reactionwith 8YSZ and the good matching of coefficient of thermal expan-sion (CTE) indicated that the investigated GCs had the potential forfurther experimentation as sealants in SOFC.

The development of technologically useful glass compositions isbased on an understanding of the relationships between themolecular-level glass structure, crystallization kinetics and impor-tant physical properties [7,8]. Therefore, in order to develop a suit-able GC sealant, it is necessary to understand the structuralbehavior and crystallization kinetics of the glass system. Moreover,despite the significant effort already made to study the crystalliza-tion kinetics in various pyroxene systems [9–13], aluminous clino-pyroxene based glasses have received relatively poor attention. Tofill this gap the present work aims to investigate the effect of BaOon the physical properties, structural behavior and non-isothermalcrystallization kinetics of the glasses in the Di–Ca-Ts system.

2. Experimental

Four experimental glass compositions designed in our previouswork [6] as BaCMAS1 (Ca0.95Ba0.05Mg0.8Al0.4Si1.8O6), BaCMAS2(Ca0.9Ba0.1Mg0.8Al0.4Si1.8O6), BaCMAS3 (Ca0.8Ba0.2Mg0.8Al0.4Si1.8O6),and BaCMAS4 (Ca0.8Ba0.2Mg0.8Al0.2B0.2Si1.8O6) were prepared (Table1) and investigated. Note that BaCMAS4 derives from BaCMAS3 by50% substitution of B3+ for Al3+. Powders of technical grade SiO2

(purity >99.5%) and CaCO3 (>99.5%), and of reactive grade Al2O3,MgCO3, BaCO3, H3BO3 and NiO were used. Homogeneous mixturesof batches (�100 g), obtained by ball milling, were preheated at900 �C (1173 K) for 1 h for calcination and then melted in Pt-cruci-bles at 1580 �C (1853 K) for 1 h, in air. Glasses in bulk form wereproduced by pouring the melts on preheated bronze moulds fol-lowed by annealing at 550 �C (823 K) for 1 h. The annealing tem-perature of the glasses was chosen in accordance with thedilatometric glass transition temperature (Tdg) obtained in our pre-vious study (Table 2) [6]. Glasses in frit form were also obtained byquenching of the melts in cold water. The frit was dried and thenmilled in a high-speed agate mill. The particle size of the glasspowders was determined by light scattering technique (BeckmanCoulter LS 230, CA USA; Fraunhofer optical model). Archimedes’method (i.e. immersion in diethyl phthalate) was employed tomeasure the apparent density of the bulk annealed glasses. Themean values and the SD presented for density have been obtainedfrom (at least) 10 different samples.

Table 1Batch compositions of parent glasses (1 wt.% NiO was added to the batches).

Glass MgO C

BaCMAS1 wt.% 14.54 2Ca0.95Ba0.05Mg0.8Al0.4Si1.8O6 mol.% 21.05 2

Mol ratio 4.00 4BaCMAS2 wt.% 14.23 2Ca0.9Ba0.1Mg0.8Al0.4Si1.8O6 mol.% 21.05 2

Mol ratio 4.00 4BaCMAS3 wt.% 13.64 1Ca0.8Ba0.2Mg0.8Al0.4Si1.8O6 mol.% 21.05 2

Mol ratio 4.00 4BaCMAS4 wt.% 13.83 1Ca0.8Ba0.2Mg0.8Al0.2B0.2Si1.8O6 mol.% 21.05 2

Mol ratio 8.00 8

The glass powders were granulated (by stirring in a mortar) in a5 vol.% polyvinyl alcohol solution (PVA, Merck; the solution of PVAwas made by dissolution in warm water) in a proportion of97.5 wt.% of glass powder and 2.5 wt.% of PVA solution. Rectangu-lar bars with dimensions of 4 � 5 � 50 mm3 were prepared by uni-axial pressing (80 MPa). The bars were sintered under non-isothermal conditions for 1 h at 850 (1123 K), 900 (1173 K), 950(1223 K), and 1000 �C (1273 K). A slow heating rate of 2 K/minaimed to prevent deformation of the samples.

Infrared spectra for the glass powders and GCs (after crushingthem into powder form) were obtained using an infrared Fourierspectrometer (FT–IR, model Mattson Galaxy S-7000, USA) in therange of 300–1500 cm�1. For this purpose each sample was mixedwith KBr in the proportion of 1/150 (by weight) for 15 min andpressed into a pellet using a hand press. 64 scans for backgroundand 64 scans per sample were made with signal gain 1. The reso-lution was 4 cm�1.

Crystallization kinetics of the glasses was studied using differ-ential thermal analysis (DTA-TG, Setaram Labsys, France). The glasspowder weighing 50 mg was contained in an alumina crucible andthe reference material was a-alumina powder. The samples wereheated in air from ambient temperature to 1000 �C (1273 K) at dif-ferent heating rates (b in the range of 5–40 K/min. The value of theglass crystallization onset temperature, Tc and peak temperature ofcrystallization, Tp were determined by using the microprocessor ofthe thermal analyzer.

The crystalline phases were determined by X-ray diffraction(XRD) analysis (Rigaku Geigerflex D/Max, C Series; CuKa radiation;2h angle range 10–80�; step 0.02 deg/s). The phases were identifiedby comparing the experimental X-ray patterns to standards com-plied by the International Centre for Diffraction Data (ICDD).Microstructure observations were done at polished (mirror finish-ing) and then etched (by immersion in 2 vol.% HF solution for 20 s)surfaces of sintered GCs by field emission scanning electronmicroscopy (FE-SEM, Hitachi S-4100, Japan; 25 kV acceleration

aO BaO SiO2 B2O3 Al2O3

4.03 3.46 48.78 – 9.205.00 1.32 47.37 – 5.26.75 0.25 9.00 – 1.002.27 6.77 47.73 – 9.003.68 2.63 47.37 – 5.26.50 0.50 9.00 – 1.008.98 12.98 45.76 – 8.631.05 5.26 47.37 – 5.26.00 1.00 9.00 – 1.009.24 13.15 46.40 2.98 4.371.05 5.26 47.37 2.63 2.63.00 2.00 18.00 1.00 1.00

0

1

2

3

4

0.01 0.1 1 10 100

Particle Diamater (μm)

Vol

ume

(%)

BaCMAS1

BaCMAS2

BaCMAS3

BaCMAS4

Fig. 1. Particle size distributions of milled glass powders used to prepare the glass-powder compacts.

736

511

BaCMAS1

731

511

BaCMAS2

1012

732

512BaCMAS3

BaCMAS4

742

507

1427

987

990

300 600 900 1200 1500Wavenumber (cm-1)

Tra

nsm

ittan

ce (

%)

Fig. 2. FTIR spectra of the investigated glass powders (The spectra are notsmoothened or normalized).

A. Goel et al. / Journal of Non-Crystalline Solids 355 (2009) 193–202 195

voltage; beam current 10 lA; magnification � 6.00 k) under sec-ondary electron mode.

3. Results

3.1. Characterization of the glasses

For all the investigated compositions (Table 1), melting at1580 �C (1853 K) for 1 h was adequate to obtain bubble-free, trans-parent glasses with dark brown color. Absence of crystalline inclu-sions was confirmed by XRD analysis. The mean particle size of theglass powders used in the present investigation lies between 2.6–3.6 lm and their particle size distribution is presented in Fig. 1.The dilatometric glass transition temperature (Tdg) increasedslightly with increase in BaO content from glass BaCMAS1 to BaC-MAS3, and further, reduced considerably with the addition of B2O3

in glass BaCMAS4 (Table 2) [6].Along the series of the four investigated compositions of Table

1, the properties values summarized in Table 2 reveal the followinggeneral features:

3.1.1. Theoretical optical basicity (Kth)The theoretical optical basicity of the four glasses, represented

by Kth, was calculated from the composition of the glass by consid-ering the proportion of oxygen (-II) atoms each oxide contributes(this is called as equivalent fraction, X), and also taking into ac-count the optical basicity values of individual oxides (KRO; RO des-ignates various oxides in Eq. (1)) [14]. The optical basicity values ascalculated using Eq. (1):

Kth ¼X

XROKRO ð1Þ

showed that Kth increases with increase in BaO content. However, aslight decrease was observed with the partial substitution of Al3+ byB3+ in glass BaCMAS4. The Kth value for all the four BaO containingglasses is higher than its BaO-free parent glass composition CaM-g0.8Al0.4Si1.8O6 (Kth = 0.593) [13].

3.1.2. Density and molar volumeThe experimental results showed that substitution of CaO by

BaO caused an increase in the density of the glasses. However,the density of all the glasses under investigation is higher thanthe density of the ZnO-containing Di–Ca-Ts glasses and is lowerthan La2O3 containing Di glasses reported in our previous studies[5,13].

The molar volume (Vm), oxygen molar volume (Vo) and excessmolar volume (Ve) were calculated using the apparent density datafor the bulk glasses using following relations:

Vm ¼Mq

ð2Þ

where M is the molar mass of the glass and q is the apparent densityof the bulk glasses. Similarly, excess molar volume of the glassescan be expressed as:

Ve ¼ Vm �X

i

xiVmðiÞ ð3Þ

Here, xi is the molar concentration of every oxide and Vm(i) isthe molar volume of every oxide [5]. Oxygen molar volume ofthe glasses was calculated using the following relation:

Vo ¼

Pi

xiMi

qP

inixi

ð4Þ

where Mi is the molar weight of the oxide, i and ni is the oxygencontent in the ith oxide, respectively. Vm and Vo increased with in-crease in BaO content in the glasses (Table 2). Vo decreased slightlyfor glass BaCMAS4 in comparison to glass BaCMAS3, with substitu-tion of Al3+ by B3+. The highest value of Ve was calculated for glassBaCMAS1. The value of Ve for glass BaCMAS2 showed a marginal de-crease in comparison to glass BaCMAS1 while a sharp decrease in Ve

was observed with further substitution of BaO for CaO in glass BaC-MAS3. Ve showed a small increment for glass BaCMAS4 in compar-ison to glass BaCMAS3. The values of Vm and Vo for the glasses underinvestigation are higher than the BaO-free parent glass composition(Di/Ca-Ts = 80/20 mol.%), ZnO-containing Di–Ca-Ts glasses andLa2O3 and Cr2O3 containing Di glasses [5,13].

3.1.3. FT–IR analysisThe room temperature FTIR transmission spectra of all the four

glasses are shown in Fig. 2. The spectra of all the B-free glasses ex-hibit three broad transmittance bands in the region of 300–1300 cm�1. The most intense transmission bands lie in the 800–1300 cm�1 region, the next one between 300 and 600 cm�1 and,

5K/min

10K/min

20K/min

40K/min

800 1000 1200 1400Temperature (K)

Hea

t Flo

w (

μV)

Exo

BaCMAS2

Fig. 4. Differential thermal analysis (DTA) thermograph of the glass powderBaCMAS2 at different heating rates (b).

196 A. Goel et al. / Journal of Non-Crystalline Solids 355 (2009) 193–202

while the least intensive lies between 650 and 800 cm�1. The broadbands in the 800–1300 cm�1 are assigned to the stretching vibra-tions of the SiO4 tetrahedron with a different number of bridgingoxygen atoms. A shift towards higher wave number has been reg-istered in this band with the increase in BaO content. The bands inthe 300–600 cm�1 region are due to bending vibrations of Si–O–Siand Si–O–Al linkages [15,16]. The transmission bands in the 650–800 cm�1 region are related to the stretching vibrations of the Al–Obonds with Al3+ ions in four-fold coordination [15]. In addition tothe three bands observed in B-free glasses, the IR spectra of glassBaCMAS4 exhibited one additional transmittance band in the re-gion 1300–1500 cm�1 corresponding to B–O vibrations in [BO3] tri-angle. Also, the band in the region 800–1300 cm�1 in glassBaCMAS4 shifted towards lower wave number with introductionof B2O3 in the glass. Although, no other band could be resolvedin the spectra of the glasses, the presence of [BO4] tetrahedron inthe glass structure of BaCMAS4 can not be neglected. Since the IRband for [BO4] tetrahedron (about 1000 cm�1) overlaps with thatof stretching vibrations of SiO4, therefore, it could not be observedin the present investigation.

3.1.4. Crystallization kinetics by DTAThe DTA plots of all the glasses feature two endothermic dips

(Tg1 and Tg2) before the onset of crystallization (Tc), indicatingthe existence of two glass transition temperatures, which are fol-lowed by an exothermic crystallization curve. Fig. 3 is the DTAthermograph of glass BaCMAS2 at b = 40 K/min, which shows theexistence of two glass transition temperatures, (Tg1 and Tg2), Tc

and peak temperature of crystallization, Tp.The DTA plots of all the four glasses exhibit single exothermic

effects at all the heating rates which shifted towards higher tem-peratures with increase in heating rate (Fig. 4). A slight increasein the peak temperature of crystallization (Tp) was observed forglass BaCMAS2 in comparison to glass BaCMAS1, while no furtherincrease was observed with increase in BaO content in glass BaC-MAS3 (Table 2). The Tp for glass BaCMAS4 decreased considerablywith the addition of B2O3 and is the lowest among all the fourinvestigated glasses. It is noteworthy that Tp of all the BaO-contain-

600 800 1000 1200 1400

Temperature (K)

Hea

t Flo

w (

μV)

Tg1 Tg2Tc

Tp

Exo

β = 40K/min

BaCMAS2

Fig. 3. Differential thermal analysis (DTA) thermograph of the glass powderBaCMAS2 at b = 40 K/min.

ing (but B-free) glasses is higher than their BaO-free parent glasscomposition, ZnO containing Di–Ca-Ts glasses and La2O3-contain-ing Di based glasses [5,13].

The crystallization kinetics of the glasses was studied using theformal theory of transformation kinetics as developed by Johnsonand Mehl [17] and Avrami [18], for non-isothermal processes thathas already been obtained in our previous work [19,20]:

InT2

p

b

!¼ Ec

RTp� In q ¼ 0 ð5Þ

which is the equation of a straight line, whose slope and interceptgive the activation energy, Ec, and the pre-exponential factor,q ¼ Q

1nK0, respectively, and the maximum crystallization rate by

the relationship:

dxdt

����p¼ 0:37bEcnðRT2

pÞ�1 ð6Þ

which makes it possible to obtain, for each heating rate, a value ofthe kinetic exponent, n. In Eq. (6), v corresponds to the crystalliza-tion fraction and dv

dt

���p

corresponds to the crystallization rate, which

may be calculated by the ratio between the ordinates of the DTAcurve and the total area of the crystallization curve. It may be ob-served that the dv

dt

���p

value increases as well as the heating rate

(Fig. 5), which has been discussed in literature [21]. The values ofEc and n for all the four glasses are listed in Table 2 and Fig. 6.The corresponding mean values may be taken as the most probablevalue of the quoted exponent. In all the B2O3-free glasses, Ec de-creases with the increase in BaO-content. The introduction ofB2O3 in glass BaCMAS4 led to a higher Ec in comparison to glassBaCMAS3, however, still it is lower than glass BaCMAS1 and BaC-MAS2. The Avrami parameter for all the investigated compositionsvaries between 2.03 and 2.12, pointing towards bulk crystallizationwith constant number of nuclei and two dimensional growths ofcrystals [22]. These results are in accordance with those reportedfor pure diopside glass by Branda et al. [11]. The Ec values for glassBaCMAS1 and BaCMAS2 are higher than BaO-free Di–Ca-Ts glass

0

0.2

0.4

0.6

0.8

1

1150 1200 1250 1300

Temperature (K)

Cry

stal

lizat

ion

Frac

tion

(χ)

5K/min

20K/min

30K/min

40K/min

0

0.004

0.008

0.012

0.016

0.02

1150 1200 1250 1300

Temperature (K)

d χ/d

t (se

c-1)

a

b

Fig. 5. (a) Crystallization fraction (v) versus temperature (T) for the crystallizationcurve of glass BaCMAS2. (b) Crystallization rate (dv/dt) versus temperature (T) forthe crystallization curve of glass BaCMAS2.

14

15

16

17

0.78 0.8 0.82 0.84 0.86

1000/Tp (K)

ln(T

p2 /β)

BaCMAS1

BaCMAS2

BaCMAS3

BaCMAS4

Fig. 6. Plot of activation energy for crystallization (Ec) for all the four investigatedglasses (The lines represent the slope in accordance with the straight line equationy = mx+c).

A. Goel et al. / Journal of Non-Crystalline Solids 355 (2009) 193–202 197

(405 kJ/mole), while Ec value for glass BaCMAS3 and BaCMAS4 arelower than the parent glass composition [13]. Further, the Ec valuesof all the four glasses are considerably lower than that reported byBranda et al. [11] for stoichiometric diopside glass (572 kJ/mole).The value of Avrami parameter for all the glasses is in accordancewith the studies reported earlier [11,13].

3.2. Crystal phase development in GCs

3.2.1. XRD of the sintered GCsThe crystalline phase development in all the sintered GCs heat

treated between 850 (1123 K) and 1000 �C (1273 K), as followedby XRD, is presented in Table 3. Augite crystallized out being theonly crystalline phase in composition BaCMAS1 after heat treat-ment at 850 �C (1123 K); however, with further increase in tem-perature, orthorhombic polymorph of hexacelsian (a-HC)appeared along with Augite. These two phases were finally accom-panied by monoclinic polymorph of celsian (C) at 1000 �C (1273 K).With further increase in BaO content, different polymorphs of HC(a-, b-) appeared along with Augite at 850 �C (1123 K) in rest ofthe three glasses and this phase assemblage was stable in all theGCs until 1000 �C (1273 K). However, higher BaO content in glassesBaCMAS3 and BaCMAS4 led to the formation of C at lower temper-ature (950 �C) in comparison to composition BaCMAS1. It is note-worthy that no monoclinic celsian was formed in compositionBaCMAS2 at all the investigated temperatures. The presence ofB2O3 in glass BaCMAS4 seems to retard the tendency for crystalli-zation of the glass and stabilizes C, as low intensity XRD peakswere registered for GCs BaCMAS4 in comparison to GCs BaCMAS3and C was also registered in GC BaCMAS4 heat treated at 1000�C(1273 K), however, it was not detected in its B-free analogue.

3.2.2. Microstructure of the sintered GCsFig. 7 shows the SEM micrographs for the GCs heat treated at

different temperatures. In accordance with the XRD results,Fig. 7(a) shows the development of augite crystals in GC BaCMAS1heat treated at 850 �C (1123 K). With the increase in BaO content,the presence of HC becomes evident (Fig. 7(b)) as dark needleshaped crystals in the amorphous material. Such needle shapedcrystals embedded in the amorphous glass is a typical microstruc-ture of HC and has also been reported in other studies [8,23]. Thedevelopment of well formed augite crystals at 850 �C (1123 K) inFig. 7(b) shows that the crystallinity in the glasses increased withthe increase in BaO content. With the increase in temperature to900 �C (1173 K), the microstructure became denser, showing thesigns of good sinterability, (Fig. 7(c)) and the crystallinity furtherincreased (Fig. 7(d)), causing an increase in the density of the resul-tant GCs [6]. As mentioned above, crystallinity was not highly pro-nounced in GCs BaCMAS4 (Fig. 7(e)). Fig. 7(f) shows the occurrenceof extensive crystallization and precipitation of HC in GC BaCMAS3heat treated at 950 �C (1223 K). This may be the reason for the highCTE of the GC BaCMAS3 sintered at 950 �C (1223 K) [6]. However,lower CTE of GC BaCMAS3 in comparison to GC BaCMAS2 may beattributed to the formation of C in former. Even though, XRD re-vealed the presence of HC in GC BaCMAS4, no typical microstruc-ture of the same could be observed, owing to the lowcrystallinity in these GCs (Fig. 7(e) and (g)).

3.2.3. FTIR of the GCsFor all the samples characterized by XRD methods in the previ-

ous section, infrared spectra were recorded, which are essential fordistinguishing different stages of the thermal history. Those spec-tra reproduced in Fig. 8 are in accordance with the XRD results.The FTIR spectra of all glass powder compacts sintered at 850 �C(1123 K) (Fig. 8(a)) show the splitting of broad bands into a num-ber of intense and sharp bands, thus, indicating the crystalline

Table 3Crystal phase development in the glass powder compacts at different temperatures as determined by XRD.

Glass 850 �C 900 �C 950 �C 1000 �C

BaCMAS1 Aug Aug; a-HC Aug; a-HC Aug; a-HC; CBaCMAS2 Aug; a-HC; b-HC Aug; a-HC Aug; a-HC Aug; a-HCBACMAS3 Aug; a-HC; b-HC Aug; a-HC; b-HC Aug; b-HC; C Aug; a-HC; b-HCBaCMAS4 Aug; a-HC; b-HC Aug; a-HC; b-HC Aug; a-HC; b-HC; C Aug; a-HC; b-HC; C

Aug: Augite, ICDD card: 01-078-1391; b-HC: b-Hexacelsian, 01-077-0185; a-HC: a-Hexacelsian, 00-012-0725; C: Celsian, 01-074-1677.

Fig. 7. Microstructure (revealed after chemical etching of polished surfaces with HF solution) of the glass–ceramics (a) BaCMAS1 heat treated at 1123 K, (b) BaCMAS3 heattreated at 1123 K (c) BaCMAS1 heat treated at 1173 K, (d) BaCMAS3 heat treated at 1173 K, (e) BaCMAS4 heat treated at 1173 K, (f) BaCMAS3 heat treated at 1223 K and (g)BaCMAS4 heat treated at 1223 K for 1 h.

198 A. Goel et al. / Journal of Non-Crystalline Solids 355 (2009) 193–202

nature of the samples. As is evident from Fig. 8(a), the transmit-tance bands for GC BaCMAS1 and BACMAS2 are broader and lessin number in comparison to GC BaCMAS3 and BaCMAS4, thusshowing the effect of BaO on the crystallinity of the glasses. Withan increase in temperature (Fig. 8(b–d)), these broad bands splitand became more intense, thus evidencing the increase in crystal-linity of the GCs.

4. Discussion

In recent years, there has been a significant effort in developingglass and GCs based sealing materials for SOFC [24]. Alkaline earthaluminosilicate based glasses and GCs have been seen with alacrityas potential candidates for this application and have attractedinterest in earlier studies [5,24–27]. However, very few studies

850oC

BaCMAS1

BaCMAS2

BaCMAS3

BaCMAS4

1075

1260

975

928

875

680

675

720

660

638

520

480

400

335

1415

300 600 900 1200 1500

Tra

nsm

ittan

ce (

%)

900oC

BaCMAS1

BaCMAS2

BaCMAS3

BaCMAS4

1075

1258

12501418

975

925

875

675640

520

482

418

390337

300 600 900 1200 1500

Tra

nsm

ittan

ce (

%)

950oC

BaCMAS1

BaCMAS2

BaCMAS3

BaCMAS4

1420

1255

1240

1080975

925

875675640

520

480418

392

337

300 600 900 1200 1500

(c)Wavenumber (cm-1) (d)Wavenumber (cm-1)

(b)Wavenumber (cm-1)(a)Wavenumber (cm-1)

Tra

nsm

ittan

ce (

%)

1000oC

BaCMAS1

BaCMAS2

BaCMAS3

BaCMAS4

1252

12451415

1080973

925

875

675640

520

480

418

337

300 600 900 1200 1500

Tra

nsm

ittan

ce (

%)

Fig. 8. FTIR spectra of all the glass powder compacts heat treated at (a) 1123 K, (b) 1173 K, (c) 1223 K and (d) 1273 K. (The spectra are not smoothened or normalized).

A. Goel et al. / Journal of Non-Crystalline Solids 355 (2009) 193–202 199

have been devoted to study the crystallization kinetics [8] andstructure of the glass and glass–ceramic sealants [9]. As mentionedin the Introduction, the development of technologically usefulglass compositions is based on an understanding of the relation-ships between the molecular-level glass structure, crystallizationkinetics and important physical properties. Therefore, the studyof these parameters becomes highly important for developing aneffective glass–ceramic based sealant.

The present work investigated the physical properties, struc-tural behavior and crystallization kinetics of the glasses and GCsin CaO–MgO–Al2O3–SiO2–(B2O3) system containing BaO inamounts varying with in 1.32–5.26 mol.%. The brown color of thebulk annealed glasses and glass-frits were due to the presence ofNiO in the glass compositions. Nickel is probably one of the ele-

ments that impart oxide glasses the widest range of coloration:green, yellow, brown, purple, and blue. These modifications ofthe glass color have been interpreted in the past as arising fromtwo kinds of sites, octahedral and tetrahedral, the proportion ofwhich varies as a function of glass composition [28]. The originof this brown coloration in the glasses is due to the presence ofnickel in five-coordination, in trigonal bipyramids [29].

Kth serves in the first approximation as a measure of the abilityof oxygen to donate a negative charge in the glasses. It can be usedto classify the covalent/ionic ratios of the glasses since an increas-ing Kth indicates decreasing covalency. According to Duffy et al.[30], who introduced the concept of theoretical optical basicity,the divalent metal ions that are stable in aqueous solution wouldrequire the glass to have an optical basicity close to that of water

200 A. Goel et al. / Journal of Non-Crystalline Solids 355 (2009) 193–202

(i.e. K = 0.40). If the glass is much more basic in comparison toaqua ion, then this basicity is transmitted to the divalent ion whichwill then become progressively less stable. In the present investi-gation, Kth for all the glasses lies in the range 0.595–0.600, whichis close to that of aqua ion. The increase in Kth of the glasses, withincrease in BaO content is due to the higher optical basicity of BaO(K = 1.21) in comparison to CaO (K = 1). According to Faaland et al.[31], in silicates, an increasing amount of basic oxides will give riseto various types of structures, ranging from infinite three-dimen-sional frameworks, as in silica itself (SiO2), to isolated SiO4� tetra-hedra in orthosilicates. A slight decrease in the Kth value of glassBaCMAS4 is due to the presence of highly acidic oxide B2O3 withlow optical basicity (K = 0.43) [14].

The experimental results showed an increase in the apparentdensity of the glasses with increasing BaO contents because BaOhas higher density (q = 5.72 g cm�3) than CaO (q = 3.34 g cm�3).The decrease in density of glass BaCMAS4 is because Al2O3

(q = 3.98 g cm�3) has been substituted by B2O3 (q = 2.55 g cm�3).The higher value of Vm and Vo for BaO-containing glasses in com-parison to BaO-free parent glass and ZnO-containing Di–Ca-Tsglasses [13] may be explained on the basis of higher ionic radii(and lower field strength) of the Ba2+ compared to Ca2+ and Zn2+

[32]. As a result of higher Vm, the Ve for BaO-containing glasses ishigher than BaO-free and ZnO containing glasses [13]. The decreasein Ve with increase in BaO content may be attributed to the increas-ing ionic character in the glass structure, as shown by Kth. Since,the ionic bonds are non-directional in nature, increasing BaO willlead to the collapse of the structural skeleton into a closer packing,thus, decreasing the excess volume of the glass [33]. However, thevalue of Ve for glass BaCMAS4 is higher in comparison to glass BaC-MAS3 which can be explained on the basis of increasing covalentnature of the glass structure with addition of B2O3. Moreover, thedecrease in Vo for glass BaCMAS4 may be attributed to the occur-rence of phase separation in the glass due to the addition of B2O3

[34,35].It was observed that the FTIR spectra of all the investigated

glasses exhibited broad transmittance bands. This lack of sharpfeatures is an indicative of the general disorder in the silicate net-work mainly due to a wide distribution of Qn units (polymerizationin the glass structure, where n denotes the number of bridging oxy-gen) occurring in these glasses. In the FTIR spectra of glasses BaC-MAS1 and BaCMAS2, (Fig. 2) in the 800–1300 cm�1 region, maintransmittance band is centered at about 990 cm�1. These resultsindicate a distribution of Qn units centered on Q2 and Q3, whilethe shifting of this band towards higher wave numbers with the in-crease in BaO content (BaCMAS3) indicates the increase in thehigher polymerized units such as Q3. The increasing connectivitywith increasing BaO content implies that only part of BaO acts asnetwork modifier. The remaining BaO might be acting as a networkformer and existing as interconnected BaO4 tetrahedra [36]. Simi-lar network forming action of MgO by forming MgO4 tetrahedronhas been reported for MgO–Na2O–B2O3 glasses [37]. The increasein glass transition temperature Tg with BaO content, as observedin our previous study [6] may be explained on the basis of increas-ing connectivity in the glass structure. With the addition of B2O3

(BaCMAS4), the transmittance band in the 800–1300 cm�1 regionshifted towards lower wave number. Thus, it is evident that theintroduction of B2O3 breaks up Q3 units and favors the formationof Q2 units. This result is in accordance with that reported by Ojo-van and Lee [38], according to whom, the introduction of boron inthe silicate glass leads to the breaking up of Q3 units and the forma-tion of Q2, Q4 and small amounts of Q1 units. The shift in the inves-tigated glasses towards the higher side from the Q2 band may bedue to the presence of Al2O3 in the glass structure. Al is knownto create Q3 bands at the expense of Q2 and Q4 units in the silicateglasses [39]. Borate glasses show two characteristic bands derived

from the B–O bonds in the [BO3] triangles (about 1300–1500 cm�1)and the [BO4] tetrahedra (about 1000 cm�1) [39]. These bands getshifted under the influence of surrounding cations, the extent andthe direction of this shift depending on the type of cation. FTIRspectra of the glasses under investigation show that in glass BaC-MAS4, boron primarily occurs in the form of [BO3] triangles. How-ever, decrease in CTE of the glass BaCMAS4 due to the addition ofB2O3, as observed in our previous study [6], points towards thepossible existence of boron in 4 coordination state along with 3-coordination state.

The existence of two endothermic dips in DTA scans before Tc,point towards the existence of phase separation in the glasses.The existence of phase separation in the diopside based glasseshas been well documented in literature [40,41]. According to DeVeckey et al. [42], in glasses located in CaO–MgO–Al2O3–SiO2 sys-tem, phase separation is caused by the segregation of calcium andmagnesium ions. The existence of a single crystallization exothermsignifies that either the GC formed as a result of crystallization ismono-mineral or different crystalline phases appear from the glassmatrix almost simultaneously and the crystallization curve is theresultant of all the crystallization curves formed due to the appear-ance of different crystalline phases. The decrease in Ec values of theglasses with increase in BaO content is in accordance with the ear-lier studies [43] and has been attributed to the low ionic fieldstrength of BaO [24]. However, this trend was not observed inour previous study [6]. The inconsistency in the Ec results withour previous investigation [6] is due to the fact that in the presentinvestigation, the calculations have been made in accordance withthe exactly calculated values of crystallization parameters n and m,while in our previous study, we had presumed the value ofn = m = 1, based on the shift in the DTA peak of the crystallizationexotherm with the increase in particle size of the glasses and themicrostructure of the bulk glass heat treated at 1000 �C (1273 K)[6], which showed that crystallization started from the surface ofthe glass. According to Branda et al. [11], the crystallization mech-anism in the diopside based glasses shifts from intermediate typeto surface crystallization with an increase in particle size. This ex-plains the appearance of surface crystallization in the bulk an-nealed glasses, as observed in our previous study [6]. The higherEc for glass BaCMAS4 in comparison to glass BaCMAS3 is due tothe presence of B2O3 in the glass. According to Fergus [24], B2O3

tends to stabilize the amorphous phase in the glass and thus, Ec in-creases with increase in B/Al ratio. Also, as in the case of sodiumborosilicate glasses, alkali ions are distributed mainly in the boratecomponent, the same can be expected for alkaline earth ions toobecause the bond between a bivalent ion and the centers of twoboron–oxygen tetrahedrons should be stronger than the bond be-tween outside oxygens of silicon–oxygen tetrahedrons. Therefore,it is highly probable that the bivalent ions are held more firmlyin the borate than in the silicate medium. The value of Avramiparameter for all the glasses suggests that crystallization did notoccur on the fixed number of nuclei. It is also important to notethat complications appear with the systems where n is a non-inte-ger value and the deviation from the integer value of the Avramiexponent might originate from impurities influencing crystalgrowth and the simultaneous appearance of different growthmechanisms as well as from density of growing phases.

Omori [44] made the analysis of the infrared spectra of the pureDi GC and reported the existence of the absorption bands at1070 cm�1, 965 cm�1 corresponding to the stretching of Si–O(br)(here, br is referred to as bridging and nbr is referred for non-bridg-ing) linkages and at 920 cm�1 and 865 cm�1 for Si–O(nbr) linkages.In the present study, bands at �1075, 975, 925 and 875 cm�1 maybe attributed to the four bands as described by Omori [44], respec-tively. The transmittance band observed at 640 cm�1, confirms theexistence of a pyroxene crystalline structure in the GCs [45],

Table 4Wave numbers and assignments of the transmittance bands corresponding to diopside in the investigated GCs at different temperatures.

Wave number (cm�1) Compositions Assignment

Diopside [44] Di–Ca-Ts GCs BaCMAS1 BaCMAS2 BaCMAS3 BaCMAS4

335 335 – – 850 �C 850 �C O(nbr)–Si–O (nbr) wagging337 900 �C, 950 �C, 1000 �C

366 – – – – – Not assigned395 390–392 900 �C, 950 �C, 1000 �C Chain deformation [O (nbr)–Ca–O (nbr)] bending

400 850 �C– 418 900 �C, 950 �C, 1000 �C 950 �C, 1000 �C Not assigned

470 480 850 �C, 900 �C, 950 �C, 1000 �C Chain deformation [O (nbr)–Mg–O (nbr)] bending510 520630 638–640 850 �C, 900 �C, 950 �C, 1000 �C O(nbr)–Si–O (nbr) bending670 660 – – 850 �C – Not assigned

675 900 �C, 950 �C, 1000 �C 850–1000 �C680 850 �C – –

865 875 850 �C, 900 �C, 950 �C, 1000 �C Si–O (nbr) stretching920 925 900–1000 �C 850–1000 �C Si–O (nbr) stretching965 975 850 �C, 900 �C, 950 �C, 1000 �C Si–O (br) stretching

1070 1075 850 �C, 900 �C – Si–O (br) stretching1080 – 950 �C, 1000 �C

A. Goel et al. / Journal of Non-Crystalline Solids 355 (2009) 193–202 201

corresponds to the bending of O(nbr)–Si–O(nbr) linkages. The bandat 675 cm�1 was also observed in the IR spectra of diopside GC andis still unresolved. Table 4 summarizes the assignment of bandscorresponding to diopside in the investigated GCs at different tem-peratures and compares them with the bands observed in purediopside [44]. The transmittance band at �335 cm�1 for all the fourGCs is due to the wagging of O(nbr)–Si–O(nbr). It is worth notingthat the band at �335 cm�1 was not observed for BaCMAS1 andBaCMAS2 at 850 �C (1173 K) (Fig. 8, Table 4), while it was observedfor GCs BaCMAS3 and BaCMAS4 at same temperature (850 �C). Thisstructural change in the GCs with increase in BaO may be attrib-uted to the decrease in Ec with increasing BaO content, thus, lead-ing to an increase in crystallinity. With further increase intemperature this band became evident for all GC compositions.The band observed at �390–400 cm�1 in the investigated GCs (Ta-ble 4) is due to the chain deformation by the bending of O(nbr)–Ca–O(nbr) linkages. In particular, band lies around 400 cm�1 at850 �C while it shifts towards lower wave number (392 cm�1) withincrease in temperature and approaches the wave number of purediopside (395 cm�1). The band observed at 418 cm�1 in the presentinvestigation was not reported either in pure diopside [44] or forCa-Ts [46]. The bands in the region �480–520 cm �1 are due tothe chain deformation by the bending of O(nbr)–Mg–O(nbr) link-ages. It should be noted that GCs from compositions BaCMAS1and BaCMAS2 sintered at 950 �C (1223 K) and at 1000 �C(1273 K) showed a broad transmittance band in the region�775 cm�1. Recently, this band has been assigned for Ca-Ts [46].In addition to all the transmittance bands as discussed above, anadditional band appeared at �1250 cm�1 in all the GCs except BaC-MAS1. This band evidences the existence of HC in the GCs [47].According to Aronne et al. [47], the IR absorption bands for HClie at 1223, 934, 662, 630, 570, 481 and 460 cm�1. Since all thebands except the one at �1250 cm�1 are either very close or over-lap with the bands of the diopside, therefore, it is difficult to iden-tify the other transmittance bands corresponding to HC in thecollected IR spectra. According to Scanu et al. [48], the band cen-tered at 1250 cm�1 may be ascribed to the intramolecular vibra-tions of AlO4 and SiO4 tetrahedra (stretching of the Al–O and Si–O bonds). Similarly, the bands for monoclinic celsian (C) are cen-tered at 1074, 1024, 958, 737, 720, 664, 615, 600, 535, 505, 472and 432 cm�1 [47]. However, unlikely to the XRD results, no trans-mittance band corresponding to C could be detected in any of theGCs showing that a very small amount of C was precipitated in theGCs or the most intense bands of C overlapped with that ofdiopside.

5. Conclusions

The influence of BaO on the physical properties, structure andcrystallization kinetics of the glasses in the Di–Ca-Ts system hasbeen studied. The molar volume and density of experimentalglasses are higher than for BaO-free and ZnO-containing Di–Ca-Ts analogs studied earlier; excess volume decreases with increas-ing BaO fraction due to increasing ionic nature of the glass struc-ture. FTIR of the glasses reveals the increasing connectivity in theglass network with increasing BaO content and the symmetricand the anti-symmetric stretching modes of the Si–O–Si bonds ofthe Qn units in the glass silicate network are distributed aroundQ2 and Q3. The Avrami parameter varies between 2.03 and 2.12,which shows the existence of two-dimensional growth of crystalswith constant number of nuclei in all the investigated glasses.The lower ionic field strength leads to decreasing Ec values withincreasing BaO content. The addition of B2O3 retards the crystalli-zation rate by stabilizing the amorphous phase in the glass, thus,leading to higher Ec. XRD, SEM and FTIR results are in good agree-ment with each other and confirm the presence of augite as theprimary crystalline phase along with the different polymorphs ofcelsian.

Acknowledgement

This study was financially supported by FCT, Portugal (AshutoshGoel, Project No. SFRH/BD/37037/2007), University of Aveiro andCICECO.

References

[1] N. Morimoto, Can. Mineral. 27 (1989) 143.[2] R.L. Flemming, R.W. Luth, Am. Miner. 87 (2002) 25.[3] S.N. Salama, H. Darwish, H.A. Abo-Mosallam, J. Eur. Ceram. Soc. 25 (2005)

1133.[4] A. Goel, D.U. Tulyaganov, S. Agathopoulos, M.J. Ribeiro, J.M.F. Ferreira, J. Eur.

Ceram. Soc. 27 (2007) 3231.[5] A. Goel, D.U. Tulyaganov, V.V. Kharton, A.A. Yaremchenko, J.M.F. Ferreira, Acta

Mater. 56 (2008) 3065.[6] A. Goel, D.U. Tulyaganov, V.V. Kharton, A.A. Yaremchenko, S. Agathopoulos,

J.M.F. Ferreira, J. Am. Ceram. Soc. 90 (2007) 2236.[7] R.K. Brow, D.R. Tallant, J. Non-Cryst. Solids 222 (1997) 396.[8] N.P. Bansal, E.A. Gamble, J. Power Sources 147 (2005) 107.[9] N.M. Pavlushkin, Principals of Glass Ceramics Technology, second ed.,

Stroiizdat, Moscow, 1979 (in Russian).[10] Z. Strnad, Glass-Ceramic Materials, Elsevier, Amsterdam, 1986.[11] F. Branda, A. Costantini, A. Buri, Thermochim. Acta 217 (1993) 207.[12] A. Goel, E.R. Shabaan, D.U. Tulyaganov, J.M.F. Ferreira, J. Am. Ceram. Soc. 91

(2008) 2690.

202 A. Goel et al. / Journal of Non-Crystalline Solids 355 (2009) 193–202

[13] A. Goel, D.U. Tulyaganov, E.R. Shabaan, R.N. Basu, J.M.F. Ferreira, J. Appl. Phys.104 (2008) 043529.

[14] V. Dimitrov, S. Sakka, J. Appl. Phys. 79 (1996) 1736.[15] S.-L. Lin, C.-S. Hwang, J. Non-Cryst. Solids 202 (1996) 61.[16] J.T. Kohli, R.A. Condratr, J.E. Shelby, Phys. Chem. Glasses 34 (1993) 81.[17] W.A. Johnson, K.F. Mehl, Trans. Amer. Inst. Mining. Eng. 135 (1981) 315.[18] (a) M. Avrami, J. Chem. Phys. 7 (1939) 1103;

(b) M. Avrami, J. Chem. Phys. 8 (1940) 212;(c) M. Avrami, J. Chem. Phys. 9 (1941) 177.

[19] A. Goel, E.R. Shabaan, F.C.L. Melo, M.J. Ribeiro, J.M.F. Ferreira, J. Non-Cryst.Solids 353 (2007) 2383.

[20] A. Goel, E.R. Shabaan, F.C.L. Melo, M.J. Ribeiro, J.M.F. Ferreira, J. Phys.: Condens.Matter. 19 (2007) 386231.

[21] J. Vázquez, C. Wagner, P. Villares, R. Jiménez-Garay, J. Non-Cryst. Solids 235–237 (1998) 548.

[22] I.W. Donald, J. Non-Cryst. Solids 345–346 (2004) 120.[23] A. Arora, E.R. Shaaban, K. Singh, O.P. Pandey, J. Non-Cryst. Solids 354 (33)

(2008) 3944.[24] J.W. Fergus, J. Power Sources 147 (2005) 46.[25] S.-B. Sohn, S.-Y. Choi, G.-H. Kim, H.-S. Song, G.-D. Kim, J. Non-Cryst. Solids 297

(2002) 103.[26] D. Bahadur, N. Lahl, K. Singh, L. Singheiser, K. Hilpert, J. Electrochem. Soc. 151

(2004) A558.[27] Y.-S. Chou, J.W. Stevenson, R.N. Gow, J. Power Sources 168 (2007) 426.[28] W.A. Weyl, Coloured Glasses, Soc. Glass. Tech., UK, 1951, p. 558.[29] G. Calas, L. Cormier, L. Galoisy, P. Jollivet, C. R. Chim. 5 (2002) 831.[30] J.A. Duffy, M.D. Ingram, S. Fong, Phys. Chem. Chem. Phys. 2 (2000) 1829.[31] S. Faaland, M.-A. Einarsrud, T. Grande, Chem. Mater. 13 (2001) 723.

[32] J.E. Shelby, Introduction to Glass Science and Technology (ISBN: 0-85404-533-3), The Royal Society of Chemistry, Cambridge, UK, 1997 (Chapter 7).

[33] J.W. Martin, Concise Encyclopedia of the Structure of Materials (ISBN: 978-0-08-045127-5), Elsevier Science Ltd., UK, 2006.

[34] A. Quintas, O. Majérus, D. Caurant, J.-L. Dussossoy, P. Vermaut, J. Am. Ceram.Soc. 90 (2007) 712.

[35] M.O. Prado, A.A. Campos, P.C. Soares, A.C.M. Rodrigues, E.D. Zanotto, J. Non-Cryst. Solids 332 (2003) 166.

[36] R.K. Mishra, V. Sudarsan, C.P. Kaushik, K. Raj, S.K. Kulshreshtha, A.K. Tyagi, J.Non-Cryst. Solids 353 (2007) 1612.

[37] K.S. Kim, P.J. Bray, Phys. Chem. Glasses 15 (1974) 47.[38] M.I. Ojovan, W.E. Lee, An Introduction to Nuclear Waste Immobilisation (ISBN:

978-0-08-044462-8), Elsevier Science Ltd., UK, 2005 (Chapter 17).[39] L. Stoch, M. Sroda, J. Mol. Struc. 511–512 (1999) 77.[40] W. Höland, G.N. Beall, Glass–Ceramic Technology, The American Ceramic

Society, Westerville, Ohio, 2002.[41] F.J. Torres, J. Alarcón, J. Non-Cryst. Solids 347 (2004) 45.[42] R.C.De Vekey, A.J. Majumdar, Glass. Tech. 15 (1974) 71.[43] N. Lahl, K. Singh, L. Singheiser, K. Hilpert, D. Bahadur, J. Mater. Sci. 35 (2000)

3089.[44] K. Omori, Am. Mineral. 56 (1971) 1607.[45] P.R. Griffths, Recent commercial instrumental developments in Fourier

transform infrared spectrometry, in: R.J.H. Clark, R.E. Hester (Eds.), Advancesin Infrared and Raman Spectroscopy, Heyden Ltd., London, 1987, p. 277.

[46] K. Etzel, A. Benisek, Phys. Chem. Miner. 35 (2008) 399.[47] A. Aronne, S. Esposito, C. Ferone, M. Pansini, P. Pernice, J. Mater. Chem. 12

(2002) 3039.[48] T. Scanu, J. Guglielmi, Ph. Colomban, Solid State Ionics 70–71 (1994) 109.