Study of Crystallization Kinetics in Glasses along theDiopside–Ca-Tschermak Join

Ashutosh Goel,z Essam R. Shaaban,y Dilshat U. Tulyaganov,z,z and Jose M. F. Ferreira*,w,z

zDepartment of Ceramics and Glass Engineering, University of Aveiro, CICECO, Aveiro 3810-193, Portugal

yPhysics Department, Faculty of Science, Al-Azhar University, Assuit 71542, Egypt

zScientific Research Institute of Space Engineering, Tashkent 700128, Uzbekistan

We report on the thermal stability and crystallization kinetics ofthe glasses in the diopside (CaMgSi2O6)–Ca-Tschermak(CaAl2SiO6) system. Four glasses with compositions corre-sponding to different diopside/Ca-Tschermak ratio were stud-ied. Structural investigations on the glasses have been made byemploying Infrared spectroscopy (FTIR). Activation energiesfor structural relaxation and viscous flow have been calculatedusing the data obtained from differential thermal analysis. Theexistence of glass-in-glass phase separation was observed in allthe glasses. Kinetic fragility of the glasses along with otherthermal parameters have also been calculated. Nonisothermalcrystallization kinetic studies have been employed to study themechanism of crystallization in all the four glasses. The Avramiparameter for the glass powders is B2, indicating the existenceof intermediate mechanism of crystallization. Crystallizationsequence in the glasses has been followed by X-ray diffractionanalysis, scanning electron microscopy, and FTIR.

I. Introduction

PYROXENES are important rock-forming mineral in severalmetamorphic and basic to ultra basic igneous rocks, also

found in meteorites. The chain silicate structure of pyroxenesenables solubility of various cations in their structure. Diopside(CaMgSi2O6, hereafter briefly designated as D) is typical repre-sentative of the pyroxene group that has only silicon in the tet-rahedral sites. These sites are occupied by both Si and Al in theCa–Al-bearing monoclinic pyroxene, known as Ca-Tschermak(CaAl2SiO6, hereafter briefly designated as Ca-TS).1 The D–Ca-TS solid solutions have a general formula of Ca[Mg1�xAlx](Si(1�x/2)Alx/2)2O6, where the square brackets and the parenthe-ses denote the octahedral and the tetrahedral sublattices, respec-tively.1 The properties of glass–ceramics (GCs) based on solidsolutions of Ca-TS and D, can be highly controlled due to theirparticular structural features.2–5

In our previous study,6 four glasses with nominal compositionscorresponding to D/Ca-TS mole ratios of 80/20 (DT1),75/25 (DT1a), 70/30 (DT1b), and 65/35 (DT1c) were investigatedfor further experimentation as candidates for SOFC sealants (Ta-ble I). The powder-processing route was employed to study thesintering and crystallization behavior of the respective glasses.The densification was complete at 9001C (1173 K) and precededcrystallization. The experimental results obtained by studyingglass crystallization between 8501C (1123 K) and 10001C (1273

K) suggested a solubility limit of Ca-TS in D at about 30 mol%TS. Augite (hereafter designated as Au), which was obviously Fe-free, acted as an intermediary member between D and Ca-TS,representing a mineral midway between these two minerals alongthis series, where Al31 occupies both octahedral (AlO6) and tet-rahedral (AlO4) positions in the structure. The limit of Ca-TSsolubility was justified by the behavior of the glass DT1c (i.e., D/Ca-TS565/35) where, along with Au, anorthite (designated asAn, CaAl2Si2O8) and akermanite (designated as Ak, Ca2Mg-Si2O7) also precipitated as secondary crystalline phases at tem-peratures 8501C (1123 K) and 9001C (1173 K), respectively.

Despite the significant effort already made to study the crys-tallization kinetics in various pyroxene systems,2,3,7–12 alumi-nous clinopyroxene-based glasses have received relatively poorattention. To fill up this gap the present work aims to studynonisothermal crystallization kinetics of the glasses in theD–Ca–TS system so as to gather detailed information on theircrystallization mechanisms and thermal stability.

II. Experimental Procedure

Four experimental glass compositions designed in our pre-vious work6 as DT1 (CaMg0.8Al0.4Si1.8O6), DT1a (CaMg0.75Al0.5Si1.75O6), DT1b (CaMg0.7Al0.6Si1.7O6), and DT1c (CaMg0.65Al0.7Si1.65O6) were prepared (Table I) and examined.Powders of technical grade SiO2 (purity 499.5%) and CaCO3

(499.5%), and of reactive grade Al2O3, MgCO3, and NiO wereused. Homogeneous mixtures of batches (B100 g), obtained byball milling, were preheated at 9001C (1173 K) for 1 h forcalcination and then melted in alumina crucibles at 15801C(1853 K) for 1 h, in air.

Glasses in bulk form were produced by pouring the melts onpreheated bronze molds followed by annealing at 5501C for 1 h.The annealing temperature for the glasses was chosen in accor-dance with dilatometric glass transition temperature obtained inour previous study.6 In order to observe crystallization mecha-nism in the glasses, bulk annealed glasses were cut into smallpieces and heat treated in an electric furnace in air at 9001C (1173K) and 10001C (1273 K) for 1 h. The heating rate was 5 K/min.

Glasses in frit form were also obtained by quenching of themelts in cold water. The frit was dried and then milled in a high-speed agate mill. The obtained powders had specific surfaceareas in the range of 1.26–1.37 m2/g (measured by BET tech-nique, Micromeritics, Gemini II 2370, Micromeritics InstrumentCorporation, Norcross, GA). The glass powders were granulated(by stirring in a mortar) in a 5 vol% polyvinyl alcohol solution(PVA, Merck, VWR International-Material de LaboratOrio,Lisboa, Portugal; the solution of PVA was made by dissolutionin warm water) in a proportion of 97.5 wt% of glass powder and2.5 wt% of PVA solution. Rectangular bars with dimensions of 4mm� 5 mm� 50 mm were prepared by uniaxial pressing (80MPa). The bars were sintered under nonisothermal conditions for1 h at 8501C (1123 K), 9001C (1173 K), 9501C (1223 K), and

D. Johnson—contributing editor

*Member, the American Ceramic SocietyThis work was financially supported by the University of Aveiro and FCT, Portugal and

CICECOwAuthor to whom correspondence should be addressed. e-mail: jmf@ua.pt

Manuscript No. 24253. Received January 23, 2008; approved April 18, 2008.

Journal

J. Am. Ceram. Soc., 91 [8] 2690–2697 (2008)

DOI: 10.1111/j.1551-2916.2008.02495.x

r 2008 The American Ceramic Society

2690

10001C (1273 K). A slow heating rate of 2 K/min aimed to pre-vent deformation of the samples.

Infrared spectra for the glass powders and GCs (after crush-ing them into powder form) were obtained using an infraredFourier spectrometer (FT-IR, model Mattson Galaxy S-7000) inthe range of 300–1500 cm�1. For this purpose each sample wasmixed with KBr in the proportion of 1/150 (by weight) for 15min and pressed into a pellet using a hand press.

Crystallization kinetics of the glasses was studied using differ-ential thermal analysis (DTA-TG, Setaram Labsys, SetaramInstrumentation, Caluire, France). The glass powder weighing50 mg was contained in an alumina crucible and the referencematerial was a-alumina powder. The samples were heated in airfrom ambient temperature to 10001C (1273 K) at different heat-ing rates (a in the range of 5–40 K/min). The value of the glasstransition temperature Tg, crystallization onset temperature, Tc

and peak temperature of crystallization, Tp were determinedusing the microprocessor of the thermal analyzer.

The crystalline phases were determined by X-ray diffraction(XRD) analysis (Rigaku Geigerflex D/Max, C Series; CuKa radi-ation; 2y angle range 101–801; step 0.021/s). The phases were iden-tified by comparing the experimental X-ray patterns to standardscomplied by the International Centre for Diffraction Data. Mi-crostructure observations were done at polished (mirror finishing)and then etched (by immersion in 2 vol%HF solution) surfaces ofglasses (etched for 10 s) and sintered GCs (etched for different timedurations varying between 20 and 120 s); and on fractured bulkGCs (etched for 20 s) by field emission scanning electron micros-copy (FE-SEM, Hitachi S-4100, Tokyo, Japan; 25 kV accelerationvoltage, beam current 10 mA) under secondary electron mode.

III. Results and Discussion

(1) Glass Production and Characterization

For all the investigated compositions, melting at 15801C for 1 hwas adequate to obtain bubble-free, transparent glasses with adark color of honey (brown). The brown color of the glasses wasdue to the presence of NiO. 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.13 The origin ofthis brown coloration in the glasses due to NiO has been ex-plained by Calas et al.,14 according to which this color arises

from the presence of nickel in five-coordination, in trigonalbipyramids.

The room temperature FTIR transmittance spectra of all thefour glasses are shown in Fig. 1. All spectra exhibit three broadtransmittance bands in the region of 300–1300 cm�1. This lackof sharp features is indicative of the general disorder in thesilicate network mainly due to a wide distribution of Qn unitsoccurring in these glasses. The broad bands in the 800–1300cm�1 are assigned to the stretching vibrations of the SiO4 tet-rahedron with a different number of bridging oxygen atoms.15,16

For all the four glasses, this band is around 1020–1040 cm�1

indicating the distribution of Qn units centered around the Q2

andQ3 unit. Al is known to createQ3 bands at the expense ofQ2

and Q4 units in the silicate glasses.17 In particular case with thechange in D/Ca–Ts ratio from 80/20 to 70/30 this band shiftedtoward lower wave number while further increase in Ca-TScontent at the expense of D in composition DT1c (D/Ca-TS5 65/35) resulted in higher wave number (Fig. 1). The low-ering of wave number in FTIR spectra with the increase inCa-TS content up to 30 mol% suggests that gradual occupancyof tetrahedral sites by Al resulted in depolymerization of thesilicate glass network.

The transmittance bands in the 300–650 cm�1 region (Fig. 1)are due to bending vibrations of Si–O–Si and Si–O–Al link-ages.18,19 This band is centered around 505–513 cm�1 for all thefour glasses and shifted to lower values when Ca-TS amountincreases from 20 to 30 mol%. However, a slight increase is re-vealed for glass DT1c when D/TS ratio is 65/35.

The transmittance bands in the 650–800 cm�1 region (cen-tered at B730 cm�1) are related to the stretching vibrations ofthe Al–O bonds with Al31 ions in fourfold coordination.18

However, in glass DT1b, this band is very sharp and is centeredat 713 cm�1 seemingly indicating distinct structural arrangementin the Al–O bonds.

The deviation of the IR spectra of glass DT1c from theobserved trends in the glasses DT1, DT1a, and DT1b signifiesthe occurrence of some rearrangements in the glass structurewhen amount of Ca-TS exceeded 30 mol%. This result supportsour previous finding that the threshold dissolution of Ca-TSin D to be ca. 30 mol % between 8501C (1123 K)–10001C(1273 K).6

The DTA plots of all the glasses feature two endothermic dips(Tg1 and Tg2) before the onset of crystallization (Tc), indicating

Table I. Compositions of the Parent Glasses (1 wt% NiO wasAdded to the Batches)

MgO CaO SiO2 Al2O3

DT1, CaMg0.8Al0.4Si1.8O6 (D/Ca-TS5 80/20)wt% 14.87 25.86 49.87 9.40mol% 21.05 26.32 47.37 5.26mol ratio 4 5 9 1

DT1a, CaMg0.75Al0.5Si1.75O6 (D/Ca-TS5 75/25)wt% 13.93 25.85 48.47 11.75mol% 19.99 26.67 46.67 6.67mol ratio 3 4 7 1

DT-1b, CaMg0.7Al0.6Si1.7O6 (D/Ca-TS5 70/30)wt% 13.00 25.84 47.06 14.09mol% 18.92 27.03 45.95 8.11mol ratio 2.33 3.33 5.67 1

DT-1c, CaMg0.65Al0.7Si1.65O6 (D/Ca-TS5 65/35)wt% 12.07 25.83 45.66 16.44mol% 17.81 27.40 45.20 9.59mol ratio 1.86 2.86 4.71 1

DT1

1037730513

DT1a

1024

730503

DT1c

1020730505

990

713501

DT1b

300

Wavenumber (cm−1)

Tra

nsm

ittan

ce (

%)

15001200900600

Fig. 1. FTIR spectra of the investigated glass powders.

August 2008 Diopside-Ca-Tschermak Glass-Ceramics 2691

the existence of two glass transition temperatures, which is fol-lowed by an exothermic crystallization curve. Particularly, DTAthermograph for the glass DT1 at b5 40 K/min (Fig. 2) showsthe existence of two glass transition temperatures, (Tg) and peaktemperature of crystallization, Tp.

The existence of two Tg values suggests the possible occur-rence of phase separation in the glasses. The SEM images ofthe glasses DT1a and DT1b shown in Fig. 3 reveal that somesegregation has occurred, thus supporting our claim for phaseseparation. The existence of phase separation in D-based glasseshas been well documented in literature.4,12 According to DeVeckey et al.,20 in glasses located in the CaO–MgO–Al2O3–SiO2

quaternary system, the phase separation is caused by the segre-gation of calcium and magnesium ions. Because the glass tran-sition took place during its first Tg, Tg1 will be considered as thetransition temperature in the following discussion of the presentresults.

The glass transition temperatures Tg (Table II) for all theglasses vary between 7271C (1000 K) and 7471C (1020 K) forb5 5 K/min. The value of Tg shifted toward higher side with theincrease in heating rate for all the glasses (Fig. 4) and is highestfor the glass DT1. This may be attributed to the stronger con-nectivity in the glass network structure as is evident from theFTIR spectra (Fig. 1). Further, it diminished sharply with theinitial decrease in D/Ca-TS ratio from 80/20 (DT1) to 75/25(DT1a). The values of Tg for glass DT1a and DT1b are almostsimilar. However, Tg increases for glass DT1c most probablydue to the increasing connectivity in the glass structure. The ac-tivation energy for the structural relaxation (Erelax) of the glassoccurring around the glass transition was calculated by using thefollowing equation:21,22

lnb ¼ �Erelax

RTg(1)

DT1

700Temperature (K)

Hea

t Flo

w (

µV)

Exo

β = 40K/min

Tg1 Tg2

Tp

Tc

13001100900

Fig. 2. Differential thermal analysis thermograph of the glass powderDT1 at b5 40 K/min.

Fig. 3. Scanning electron microscopy images of the glasses DT1a andDT1b (revealed after chemical etching of polished surfaces with HFsolution for 10 s).

Table II. Thermal Parameters for the Glasses Obtained fromDTA (b 5 5 K/min)

Glass

Tg

(K)

Tc

(K)

Tp

(K)

Erelax.

(kJ/mol)

EZ

(kJ/mol) F DT (K)

KP

(min�1)

DT1 1020 1180 1199 669 652 33.40 160 0.184DT1a 1000 1183 1213 428 411 21.47 183 0.202DT1b 1002 1204 1220 391 374 19.48 202 0.179DT1c 1012 1212 1229 473 456 24.42 200 0.188

DTA, differential thermal analysis.

DT1

5K/min

20K/min

40K/min

30K/min

700Temperature (K)

Hea

t Flo

w (

µV)

Exo

13001100900

Fig. 4. Differential thermal analysis thermograph of the glass powderDT1 at different heating rates (b).

2692 Journal of the American Ceramic Society—Goel et al. Vol. 91, No. 8

While the activation energy for viscous flow, EZ as calculatedfrom the following equation:

lnðT2g=bÞ ¼ EZ=RTg þ const (2)

where b is the heating rate and R is the gas constant. It is evidentfrom Fig. 5 and Table II that both the calculated activation en-ergies are in good agreement with each other and the highestvalues of Erelax and EZ have been obtained for glass DT1, whilethe lowest ones were obtained for glass DT1b. The Erelax and EZdecrease with decreasing the D/Ca-TS ratio from D/Ca-TS5 80/20 until D/Ca-TS5 70/30. Erelax and EZ increase with fur-ther decreasing this ratio to D/TS5 65/35 (DT1c). It has alreadybeen reported by various researchers23,24 that both Erelax and EZare responsible for the molecular motion and rearrangement ofthe atoms around Tg and the glass with lowest activation energyvalues is the most stable one.

From the point of view of technological applications, theglass should be thermally stable. This thermal stability of glasseshas been ascertained on the basis of thermal analysis usingDTA. Usually, unstable glasses show crystallization peak closeto the glass transition temperature. Therefore, difference be-tween onset temperature of crystallization, Tc and glass transi-tion temperature, Tg i.e., DT5Tc�Tg is a good indication ofthermal stability because the higher the value of this difference,more will be the delay in the nucleation process.25 For furtherconfirmation of this result, the crystallization rate factor KP cor-responding to the temperature at which the crystallization rate is

maximum has been obtained using the following equation:

bEc

RKPT2P

¼ 1 (3)

where Tp is the crystallization peak temperature and Ec is acti-vation energy for crystallization (calculated in the section III-2).The physical significance of the crystallization rate factor is thatits minimum value gives an indication of the retardation of thecrystallization while its higher value diminishes the glass-form-ing ability.26 Table II presents the values of DT andKP for all thefour investigated compositions at b5 5 K/min and it revealsthat the highest value of DT and lowest value of KP occurs atD/Ca-TS5 70/30 (DT1b). It has been widely reported22,24,25

that the glasses with higher DT and minimum KP are thermallymost stable.

Another important parameter for determining the glass-forming ability is the kinetic fragility index, F, because it is ameasure of the rate at which the relaxation time decreases withincreasing temperature around Tg and is given by the followingequation: 27,28

F ¼ EZ

RTg ln 10(4)

According to Vilgis,29 glass-forming liquids that exhibit anapproximate Arrhenius temperature dependence of their relax-ation times are defined as strong and specified with a low valueof F (F�16), while the limit for fragile glass-forming liquids ischaracterized by a high value of F (F�200).30 The values for allthe samples have been obtained at all the heating rates andTable II presents the values of F at b5 5 K/min. Because allthese values of F are within the above-mentioned limit, it is rea-sonable to state that all the glasses are obtained from strongglass-forming liquids. However, the value of F is maximum forglass DT1 and minimum for glass DT1b, confirming higherthermal stability of the latter.

(2) Crystallization Kinetics of the Glasses

The DTA plots of all the four glasses exhibit single exothermiceffects at all the heating rates which shifted toward higher tem-peratures with increase in heating rate.6 The existence of a singlecrystallization exotherm signifies that either the GC formed as aresult of crystallization is mono-mineral or different crystallinephases appear from the glass matrix almost simultaneously andthe crystallization curve is the resultant of all the crystallizationcurves formed due to the appearance of different crystallinephases. The peak temperature of crystallization, Tp, increasesfrom D/Ca-TS5 80/20 (DT1) to D/Ca-TS5 65/35 (DT1c) aspresented in Table II. XRD results of the sintered GCs (TableIII) reveal that a complete solid solution exists between D andCa-TS until D/Ca-TS5 70/30 and Au is the only crystallinephase precipitated in the as-formed GCs until 10001C (1273 K).6

However, An and Ak appear as secondary phases in GC DT1cat 8501C (1123 K) and 9001C (1173 K), respectively. Figure 6shows the SEM images of the sintered GCs. It is clearly visiblethat the crystal formation starts at 8501C (1123 K) in glass DT1a(Fig. 6(a)), and reaches its maximum at 9001C (1173K)

1

2

3

4

0.951000 / Tg (K)

lnβ

DT1

DT1a

DT1b

DT1c

(a)

(b)

10

11

12

13

1000 /Tg (K)

ln(T

g2/β

)

DT1

DT1a

DT1b

DT1c

1.010.990.97

0.95 1.010.990.97

Fig. 5. (a) Plot for determination of activation energy for structuralrelaxation (Erelax) for the investigated glasses. (b) Plot for determinationof activation energy for viscous flow (EZ) for the investigated glasses.

Table III. Crystal Phase Development in the Glasses atDifferent Temperatures as Determined by XRD

Glass 1123 K 1173 K 1223 K 1273 K

DT1 Amorphous Au Au AuDT1a Amorphous; Au Au Au AuDT1b Amorphous; Au Au Au AuDT1c Amorphous; Au;

AnAu; An;Ak

Au; An;Ak

Au; An;Ak

Au, Augite, ICDD card: 01-078-1392; An, Anorthite, 00-041-1486;

Ak, Akermanite, 00-035-0592; XRD, X-ray diffraction.

August 2008 Diopside-Ca-Tschermak Glass-Ceramics 2693

(Fig. 6(b)). The disk-shaped crystals of Au can be seen for GCDT1a (Fig. 6(b)) and GC DT1b (Fig. 6(c)). However, SEM im-ages of GC DT 1c (Fig. 6(d)) revealed some dendrites along withthe Au phase may be due to the appearance of An and Ak assecondary phases (as confirmed also by XRD data).

The FTIR spectra of all the crushed glass powder compactssintered at 8501C (1123 K) (Fig. 7(a)) show the existence of threebroad bands similar to those observed for the parent glass pow-ders, thus evidencing the highly amorphous nature of all thecompositions. Evaluation of phases in glasses heated at different

temperatures according to XRD results is summarized inTable III. All compositions except composition DT1 exhibitedtraces of crystals at 8501C (1123 K). However, the amorphousnature of the glass predominated, with the crystalline phaseshowing the maximum intense peak for composition DT1a.6

The careful observation of FTIR spectra of composition DT1asintered at 8501C (1123 K) (Fig. 6(a)) revealed the broad trans-mittance bands in the regions 650–800 and 800–1300 cm�1,which can be split into three parts, respectively, thus giving theevidence of initiation of crystallization. This finding is in a good

Fig. 6. Microstructure (revealed after chemical etching of polished surfaces with HF solution) of the glass–ceramics (a) DT1a heat treated at 1123 K,(b) 1173 K (c) DT1b heat treated at 1173 K and (d) DT1c heat treated at 1223 K 1 h.

300

Wave number (cm−1)

Tra

nsm

ittan

ce (

%)

850°C

725

995

DT1

495

881 1080

987

DT1a

652

688

731

480

490987

DT1b

723

490984

DT1c

717

(a)

15001200900600

Wavenumber (cm−1)

Tra

nsm

ittan

ce (

%)

900°C

DT1

935

DT1a

DT1b

DT1c

947

1082

982

889

677

642

625

542

528

485

423

340

(b)

300 15001200900600

1000°C

DT1

DT1b

DT1a

DT1c1081982

890

679

690

644

532

550

487

425

343

Wavenumber (cm−1)

Tra

nsm

ittan

ce (

%)

(c)

300 15001200900600

Fig. 7. FTIR spectra of all the glass powder compacts heat treated at (a) (1123 K), (b) 1173 K, and (c) 1273 K. Spectra were recorded by crushing thesintered GCs to powder form and mixing the GC powder with KBr.

2694 Journal of the American Ceramic Society—Goel et al. Vol. 91, No. 8

correlation with XRD and SEM data as presented in Table IIIand Fig. 6(a) that evidence precipitation of Au phase in DT1a at8501C (1123 K). In general, the IR bands in the region 800–1300and 300–650 cm�1 shifted toward lower wave numbers in com-parison with the parent glasses, suggesting some structural re-arrangements in the silicate glass network on heat treatment.With an increase in temperature to 9001C (1173 K), the threebroad transmittance bands split into a number of intense andsharp bands, thus, showing an increase in the crystallinity of theGCs. Omori31 made the analysis of the infrared spectra of thepure diopside GC and reported the existence of the absorptionbands at 1070 and 965 cm�1 corresponding to the stretching ofSi–O(br) (here, br is referred to as bridging and nbr is referredfor nonbridging) linkages and at 865 cm�1 for Si–O(nbr) link-ages. In the present study, bands at B1080, 980, and 880 cm�1

may be attributed to the three bands as described by Omori,31

respectively. The transmittance band observed at 652 cm�1 forglass DT1a at 8501C (1123 K) shifted to lower wave numberwith the increase in temperature and appeared at B642 cm�1

while the same band shifted further lower for DT1c (Fig. 7(b)).Griffith32 assigned this band to the pyroxene phase, while ac-cording to Omori,31 this band is due to the bending of O(nbr)–Si–O(nbr) linkages. The transmittance band at B340 cm�1 forall the four GCs at 9001C (1173 K) and 10001C (1273 K) is dueto the wagging of O(nbr)–Si–O(nbr) while the bandB425 cm�1

may be due to the chain deformation by the bending of O(nbr)–Ca–O(nbr) linkages. The bands in the region B485 and 530cm�1 are due to the chain deformation by the bending ofO(nbr)–Mg–O(nbr) linkages. The latter band (530 cm�1) shiftedto higher wave numbers in case of GC DT1c (Fig. 7(b) and (c)).This deviation from the trends followed by other three compo-sitions, point toward the occurrence of some different structuralrearrangements. X-ray data, in conjunction with the FTIR data,reveal that the solid solution between D and TS no more existsfor composition DT1c and An and Ak precipitate out along

with Au at 9001C (1173 K) and 10001C (1273 K), respectively.The appearance of Au at lower temperature followed by An inGC DT1c may be attributed to the structure of the two crys-talline phases. Au is an inosilicate of simpler structure with re-spect to tectosilicate An. Consequently, the formation of Aurequires lesser structural rearrangements.

The crystallization kinetics of the glasses was studied usingthe formal theory of transformation kinetics as developed byJohnson and Mehl33 and Avrami,34–36 for nonisothermal pro-cesses that has already been obtained in our previous work:37,38

lnT2p

b

!¼ Ec

RTp� ln q ¼ 0 (5)

which is the equation of a straight line, whose slope and inter-cept give the activation energy, Ec, and the preexponential fac-tor, q ¼ Q

1nK0, respectively, and the maximum crystallization

rate by the relationship:

dwdt

����p

¼ 0:37bEcn RT2p

� ��1(6)

which makes it possible to obtain, for each heating rate, a valueof the kinetic exponent, n. In Eq. (6), w corresponds to the crys-tallization fraction and dw

dt

��pcorresponds to the crystallization

rate, which may be calculated by the ratio between the ordinatesof the DTA curve and the total area of the crystallization curve.It may be observed that the dw

dt

��pvalue increases as well as the

heating rate (Fig. 8), which has been discussed in literature.39

The values of Ec and n for all the four glasses are listed in TableIV and Fig. 9. The corresponding mean values may be taken asthe most probable value of the quoted exponent. Besides, fromthe mean value of the kinetic exponent, n, it is possible to pos-tulate a crystallization reaction mechanism for the D–Ca-TSglasses. A value of n close to 3 supports bulk or three dimen-

0

0.2

0.4

0.6

0.8

1

1150Temperature (K)

Cry

stal

lizat

ion

frac

tion

(χ)

5K /min20K/min30K/min40K/min

(a)

130012501200

0

0.005

0.01

0.015

0.02

Temperature (K)

dχ/d

t (se

c−1)

5K/min20K/min30K/min40K/min

(b)

1150 130012501200

Fig. 8. (a) Crystallization fraction (w) versus temperature (T) for thecrystallization curve of glass DT1. (b) Crystallization rate (dw/dt) versustemperature (T) for the crystallization curve of glass DT1.

Table IV. Crystallization Parameters for the Glasses asObtained from DTA

Glass Ec (kJ/mol) /nS

DT1 439 2.11170.032DT1a 493 2.10670.019DT1b 444 2.08170.047DT1c 473 1.96370.112

DTA, differential thermal analysis.

14

15

16

17

0.761000/Tp (K)

ln(T

p2/β

)

DT1

DT1a

DT1b

DT1c

0.840.820.80.78

Fig. 9. Plot of activation energy for crystallization (Ec) for all the fourinvestigated glasses.

August 2008 Diopside-Ca-Tschermak Glass-Ceramics 2695

sional crystal growth and a value close to 1 indicates surfacegrowth. Intermediate values of n between 1 and 3 might resultwhen surface and internal crystallization occur simultaneously.

As is evident from Table IV, the trend followed by the acti-vation energy of crystallization, Ec, for all the glasses is consis-tent with our previous results.6 The small differences in the Ec

values relative to our previous work are due to the fact that inthe present investigation, the calculations have been made inaccordance with the exactly calculated values of crystallizationparameters n and m, while in our previous study, we had pre-sumed the value of n5m5 1, based on the shift in the DTApeak of the crystallization exotherm with the increase in theparticle size of the glasses.40 The value of the kinetic exponent,so called, Avrami parameter for all the glasses suggests that thecrystallization did not occur on the fixed number of nuclei and

with the increase in Ca-TS content, the tendency for surfacecrystallization increases. It is also important to note that com-plications appear with systems where n is a noninteger value andthe deviation from the integer value of the Avrami exponentmight originate from impurities influencing crystal growth andthe simultaneous appearance of different growth mechanisms aswell as from the density of the growing phases.11 The Ec valuesfor D–Ca-TS glasses are higher than those reported for theirBaO-containing derivatives41; however, they are lower thanLa2O3-containing diopside-based glasses.42

In order to visualize the effects of crystallization mechanismin the bulk glasses, bulk annealed glass DT1 was heat treated at9001C (1173 K) and 10001C (1273 K) for 1 h at a heating rate of5 K/min in a box furnace. Figure 10 shows the SEM images ofthe fractured GCs. Au crystals can be seen in Figs. 10(a) and (c)from two different views while Fig. 10(b) shows that the crys-tallization initiates from the edge or surface of the glass. Theseresults are in accordance with the study of Branda et al.,10 ac-cording to whom, crystallization mechanism in diopside glassshifts from surface to bulk nucleation with a decrease in theparticle size. However, considering that in the present investi-gation we are not dealing with stoichiometric diopside glass,such a deviation from /nS5 3 to /nS5 2 in the powderedglasses is expected due to the influence of Ca-TS content. More-over, the glass powders in present study have much lower par-ticle size in comparison with those studied by Branda et al.,10

which will further alter the crystallization mechanism in theglasses.

IV. Conclusions

The thermal stability, structure, and crystallization kinetics ofthe glasses in the system D–Ca-TS has been studied. The infra-red spectroscopy reveals that the symmetric and the antisym-metric stretching modes of the Si–O–Si bonds of the Qn units inthe glass silicate network are distributed around Q2 and Q3. Thephenomenon of glass-in-glass phase separation exists in theglasses due to the presence of Ca21 and Mg21 ions. The valuesof various thermal parameters around glass transition revealthat the glass composition corresponding to D/Ca-TS5 70/30(DT1b) is thermally most stable. The Avrami parameter for theglass powders decreases from 2.11 to 1.96 with decrease in D/Ca-TS ratio, which shows that the crystallization did not occuron the fix number of nuclei and the crystallization mechanism isintermediate (i.e., both bulk and surface crystallization) in all theglasses, which shifts more toward surface crystallization withdecrease in D/Ca-TS ratio.

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