Structure, Sintering, and Crystallization Kinetics of Alkaline-EarthAluminosilicate Glass–Ceramic Sealants for Solid Oxide Fuel Cells

Ashutosh Goel,z,y Dilshat U. Tulyaganov,z Anna Maria Ferrari,z Essam R. Shaaban,J Andreas Prange,ww

Federica Bondioliy and Jose M. F. Ferreiraw,z

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

yDipartimento di Ingegneria dei Materiali e dell’Ambiente, Facolta di Ingegneria,Universita di Modena e Reggio Emilia,41100 Modena, Italy

zDipartimento di Scienze e Metodi dell’Ingegneria, Facolta di Ingegneria, Universita di Modena e Reggio Emilia,42122 Reggio Emilia, Italy

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

wwDepartment of Glass and Ceramic Composites, RWTH Aachen University, Mauerstrasse 5,D-52064 Aachen, Germany

The most common approach to seal an electrolyte or a ceramicanode with a metallic interconnect in solid oxide fuel cells ofplanar configuration (pSOFC), in order to achieve a hermeticand stable cell, is to use rigid glass or glass–ceramic (GC) seals.The properties of these sealants can be tailored specifically foruse in SOFCs through variation of the glass composition. In thepresent manuscript, we aim to study the structure, sintering be-havior, and crystallization kinetics of a new series of alumino-silicate GC sealants derived from substitution via scheme0.1(Ca

211Si41)20.1(La

311Al31) in a pure CaMgSi2O6

(diopside) system. The substitution of BaO for CaO and theaddition of fixed amounts of B2O3, Cr2O3, and NiO were at-tempted in order to achieve desirable traits that qualify the in-vestigated GCs for sealing application in pSOFCs. Theinvestigated glasses showed a steady increase in density, coeffi-cient of thermal expansion, and molar volume with addition ofCr2O3 and increasing BaO contents while no significant varia-tion in glass transition temperature (Tg) was observed. Sinteringand crystallization behavior of the glasses were investigated us-ing hot-stage microscopy and differential thermal analysis, re-spectively. The microstructure and the properties of GCs wereinvestigated under nonisothermal heat treatment conditions(8001 and 8501C; 1 h). Finally, Rietveld–R.I.R. technique wasused to quantify the amount of crystalline and amorphous con-tent in the GCs. It was observed that while addition of Cr2O3

improved the sinterability of glass-powder compacts and de-creased the amorphous content in the resultant GCs, substitutionof BaO for CaO led to an increase in amorphous content in theGCs.

I. Introduction

THE planar configuration of solid oxide fuel cells (SOFC) issuperior to other configurations in terms of efficiency and

power density, but requires hermetic seals to prevent fuel–oxi-dant mixing and provide electrical insulation to stacks.1 Sealsrequired for a planar SOFC can be classified as metal–metal,

ceramic–ceramic, and metal–ceramic seals. Among these seals,metal–ceramic seals are particularly challenging because of theirsevere functional requirements as well as difficulty in the selec-tion of material and the associated processing optimization.2

Glasses and glass–ceramic (GCs), in principle, meet most of therequirements of an ideal sealant by suitably choosing the com-ponents of the glasses and their stoichiometric proportion.1

In pursuit of the development of a compliant glass/GC seal-ant, it is necessary to understand the sintering process and crys-tallization kinetics of the glasses. Also, because the nature andamount of various crystalline phases alter the coefficient of ther-mal expansion (CTE) and the mechanical properties of the finalGCs while the remaining glassy phase acts as a major factor fordeciding the reaction kinetics during chemical interaction be-tween the GC sealant and SOFC components, therefore, qual-itative and quantitative analysis of crystalline and amorphouscontent in the GCs is of crucial importance. Lahl et al.3 inves-tigated the crystallization kinetics of a series of alkaline earthaluminosilicate glasses for potential application as sealants inSOFC. It was concluded that the activation energy of crystal-lization (Ec) increases with the replacement of BaO by CaO orMgO and by the addition of nucleating agents like Cr2O3 and Niin the glasses. In another study, Lara et al.4,5 studied the sinte-ring and crystallization behavior of GC sealants in the systemZnO–BaO–SiO2 and MgO–BaO–SiO2. Furthermore, they in-vestigated the nonisothermal crystallization kinetics of glasscompositions (mol%) 50SiO2–30BaO–20ZnO and 55SiO2–27BaO–18MgO using Kissinger, Marseglia, and Matusita–Sakka equations.6 Similarly, Bansal and Gamble investigatedthe non-isothermal crystallization kinetics of a potential glasssealant with composition (mol%) 35BaO–15CaO–5Al2O3–10B2O3–35SiO2 (BCAS).7 The Ec for BCAS glass was calcu-lated to be 259 kJ/mol with a two-dimensional growth of BaSiO3

crystalline phase. Recently, Zhang et al.8 studied the isothermalcrystallization kinetics of a glass with composition (mol%)19.2CaO–18.5SrO–13.2ZnO–1.9B2O3–2.9Al2O3–2.0TiO2–42.2SiO2

and calculated using DTA, the variation extent of the crystallizationfraction in GC with increasing time. However, they did not calcu-late the amount of different crystalline phases.

The present study is aimed toward investigating the structure,sintering behavior, and crystallization kinetics of their parentaluminosilicate glasses derived from substitution via scheme0.1(Ca211Si41)20.1(La311Al31) in a pure CaMgSi2O6 (diop-side) system. The main objective of this work was to optimizethe sealing process of the glasses at temperatures close to work-ing temperatures (8001–8501C) of SOFC on the basis of the

J. Stevenson—contributing editor

This study was financially supported by the University of Aveiro, CICECO, and FCT,Portugal (SFRH/BD/37037/2007).

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

Manuscript No. 26456. Received July 17, 2009; approved October 19, 2009.

Journal

J. Am. Ceram. Soc., 93 [3] 830–837 (2010)

DOI: 10.1111/j.1551-2916.2009.03503.x

r 2010 The American Ceramic Society

830

results obtained during the present study in conjunction with theresults published in Goel et al.9 An attempt will be made to zerodown on the potential sealing GC composition(s) in the conclu-sion and underline the modifications/improvements furtherrequired in order to develop a suitable sealant for SOFC.

II. Experimental Procedure

As mentioned in Section I aluminosilicate glasses were derivedfrom substitution via scheme 0.1 (Ca211Si41)20.1(La311Al31) in a pure CaMgSi2O6 (diopside) system. Becauseit has been widely reported that BaO increases the CTE of sil-icate/borosilicate glasses,1 partial substitution of CaO by BaOwas attempted in the present study as depicted in Table I. Fur-ther addition of 1 wt% of NiO and 2 wt% of B2O3 was carriedout in all the investigated glasses in order to improve the adhe-sion behavior of the sealant to metal and decrease the viscosityand glass transition temperature (Tg), respectively.

10 Further-more, 0.5 wt% of Cr2O3 was added in the parent glass in orderto reduce the surface tension of the glass and decrease the tem-perature of crystallization. Moreover, addition of minoramounts of Cr2O3 is known to promote the formation of mono-mineral GCs in clinopyroxene-based glass compositions, whichfurther help in stabilizing the CTE and the mechanical proper-ties of the sealant during long-term SOFC operations.11 Table Ipresents the compositions of all the glasses investigated. Pow-ders of technical-grade SiO2 (purity 499.5%) and CaCO3

(499.5%), and of reactive-grade Al2O3 (Sigma Aldrich, Dorset,U.K.; � 98%), H3BO3 (Merck, Darmstadt, Germany; 99.8%)MgCO3 (BDH Chemicals, U.K., 499%), BaCO3 (Sigma Ald-rich, 991%), La2O3 (Sigma Aldrich, 99.9%), Cr2O3 (SigmaAldrich, � 98%), and NiO (Sigma Aldrich, 99%) were used.Homogeneous mixtures of batches (B100 g), according to TableI, obtained by ball milling, were preheated at 9001C for 1 h fordecarbonization and then melted in Pt crucibles at 15501C for1 h in air. Glasses in bulk form were produced by pouring themelts on preheated bronze molds followed by annealing at5501C for 1 h. The samples of the glass-powder compactswere produced from glass frits, which were obtained by thequenching of glass melts in cold water. The frits were dried andthen milled in a high-speed agate mill, resulting in fine glasspowders with mean particle sizes of 10–20 mm (determined usingthe light scattering technique; Coulter LS 230, Beckman Co-ulter, Fullerton CA; Fraunhofer optical model).

Infrared spectra for the glass powders were obtained using anInfrared Fourier spectrometer (FT-IR, model Mattson GalaxyS-7000, Madison, WI). For this purpose, each glass powder wasmixed with KBr in the proportion of 1/150 (by weight) andpressed into a pellet using a manual press.

Dilatometry measurements were performed with prismaticsamples with a cross section of 4 mm � 5 mm (Bahr ThermoAnalyse DIL 801 L, Hullhorst, Germany; heating rate 5 K/min).Nonisothermal crystallization kinetics were studied using differ-

ential thermal analysis (DTA-TG, Setaram Labsys, Setaram In-strumentation, Caluire, France) of glass powders carried out inair from room temperature to 10001C with different heatingrates (b) of 5, 20, 30, and 40 K/min. Three samples for each glasscomposition were analyzed to obtain the dilatometry and DTAdata.

The sintering process of the glass sample was investigated byusing a hot-stage microscope (HSM). A side-view HSMEM 201equipped with an image analysis system and an electrical furnace1750/15 Leica was used. The measurements were conducted inair with a heating rate of 5 Kmin. The cylindrical-shaped sam-ples with height and diameter of �3 mm were prepared by cold-pressing the glass powders. The cylindrical samples were placedon a 10 mm� 15 mm� 1 mm alumina (499.5 wt% Al2O3)support. The temperature was measured with a Pt/Rh (6/30)thermocouple contacted under the alumina support. The tem-peratures corresponding to the characteristic viscosity points(first shrinkage, maximum shrinkage, softening, half ball, andflow) were obtained from the photographs taken during the hot-stage microscopy experiment following Pascual et al.12

Rectangular bars with dimensions of 4 mm� 5 mm� 50 mmwere prepared by uniaxial pressing (80 MPa) of the glass pow-ders. The bars were then sintered under non-isothermal condi-tions for 1 h at 8001 and 8501C, respectively. A slow heating rateof 2 K/min was maintained to prevent deformation of the sam-ples. The amorphous nature of the glasses along with qualitativeand quantitative analysis of the crystalline phases in the GCs(crushed to particle size o25 mm) were determined by X-raydiffraction (XRD) analysis using a conventional Bragg–Brent-ano diffractometer (Philips PW 3710, Eindhoven, the Nether-lands) with Ni-filtered Cu-Ka radiation. The quantitative phaseanalysis of the GCs was carried out using the combined Riet-veld-reference intensity ratio (RIR) method.13 A 10 wt% of co-rundum (NIST SRM 674a, annealed at 15001C for 1 day toincrease the crystallinity to 100%) was added to all the GCsamples as an internal standard. The mixtures, ground in anagate mortar, were side loaded in an aluminum flat holder inorder to minimize preferred orientation problems. Data wererecorded in 2y range5 5–1401 (step size 0.021 and 6 s of count-ing time for each step). The phase fractions extracted by Riet-veld-RIR refinements, using GSAS software14 and EXPGUI15

as the graphical interface, were rescaled on the basis of the ab-solute weight of corundum originally added to the mixtures asan internal standard, and therefore internally renormalized. Thebackground was successfully fitted with a Chebyshev functionwith a variable number of coefficients depending on its com-plexity. The peak profiles were modeled using a pseudo-Voigtfunction with one Gaussian and one Lorentzian coefficient. Lat-tice constants, phase fractions, and coefficients corresponding tosample displacement and asymmetry were also refined.

Archimede’s method (i.e., immersion in diethyl phthalate)was used to measure the apparent density of the bulk-annealedglasses and sintered GCs.Molar volume (Vm) and excess volume(Ve) were calculated from the density data for the bulk glasses as

Table I. Batch Compositions of the Glasses

Glass MgO CaO BaO SiO2 Al2O3 La2O3 B2O3 Cr2O3 NiO

10 wt% 17.27 21.63 0.00 48.93 2.18 6.98 2.00 0.00 1.00mol% 25.01 22.51 0.00 47.52 1.25 1.25 1.68 0.00 0.78

10A wt% 17.19 21.52 0.00 48.67 2.17 6.95 2.00 0.50 1.00mol% 24.96 22.46 0.00 47.42 1.25 1.25 1.68 0.19 0.78

10B wt% 16.48 18.34 6.27 46.67 2.08 6.66 2.00 0.50 1.00mol% 24.93 19.94 2.49 47.37 1.25 1.25 1.75 0.20 0.82

10C wt% 16.14 16.85 9.21 45.73 2.04 6.53 2.00 0.50 1.00mol% 24.92 18.69 3.74 47.34 1.25 1.25 1.79 0.20 0.83

10D wt% 15.83 15.41 12.04 44.82 2.00 6.40 2.00 0.50 1.00mol% 24.90 17.43 4.98 47.32 1.25 1.25 1.82 0.21 0.85

7-2B10 wt% 16.56 18.44 6.30 46.91 2.09 6.70 2.00 0.00 1.00mol% 24.98 19.99 2.50 47.47 1.25 1.25 1.75 0.00 0.81

March 2010 Alkaline-Earth Aluminosilicate Glass–Ceramic Sealants for SOFC 831

discussed in our previous study.16 The mechanical propertieswere evaluated by measuring the three-point bending strength ofrectified parallelepiped bars (3 mm� 4 mm� 50 mm) of sinteredGCs (Shimadzu Autograph AG 25 TA, Columbia, MD; 0.5mm/min displacement). The mean values and the standard de-viations presented for density, and bending strength were ob-tained from (at least) 10 different samples.

Microstructure observations were carried out for polished(mirror finishing) samples. The etching of GCs was performedby immersion in 2 vol% of HF solution for 2 min while theunetched 8YSZ/GC and Crofer22 APU/GC interfaces (crosssectioned) were observed using field-emission scanning electronmicroscopy (Hitachi S-4100, Tokyo, Japan; 25 kV accelerationvoltage, beam current 10 mA) under the secondary electronmode.

III. Results and Discussion

(1) Casting Ability and Physical Appearance of Glasses

For all the investigated compositions (Table I), heating at15501C for 1 h was adequate to obtain bubble-free, homoge-nous, transparent glasses. Absence of crystalline inclusions wasconfirmed by XRD and SEM analyses. The glass 10 (Cr2O3-free) was dark brown in color (likely due to the presence of NiO)while the remaining four glasses had dark brown color with agreenish hue. The origin of this brown coloration in the glasses isdue to the presence of nickel in five-coordination while thegreenish tinge in Cr2O3-containing glasses implies the presenceof chromium in Cr31 state.17

(2) Density and Dilatometry

The experimental results showed that the addition of Cr2O3 ledto a slight increase in the density (r of the investigated glasses incomparison with its parent glass 10 (Table II) as the density ofCr2O3 (5.2 g/cm3) is significantly higher in comparison with theother constituents of glass (except BaO, La2O3, and NiO). Fur-ther, the density of glasses increased with a partial substitutionof CaO by BaO due to a higher density of BaO (5.72 g/cm3) incomparison with CaO (3.34 g/cm3), thus resulting in the highestdensity for the glass 10D.

The Vm of the glass 10B is significantly higher than that of theB2O3-free glass with similar composition as investigated in ourprevious study (Table II).10 This may be attributed to the for-mation of directional bonds in the structural skeleton due to theaddition of B2O3. Vm increases while Ve decreases with increasein BaO content in the glasses. However, Ve for the glass 10B ishigher than for BaO-free glass 10A owing to an increase in di-rectional bonding in the structure of glass due to low ionic fieldstrength of BaO in comparison with CaO. The decrease in Ve

with increasing BaO content in the glasses implies that eventhough the introduction of BaO in the glass structure (glass 10B)increased the polymerization (in comparison with glass 10A),further increase in BaO content depolymerized the glass net-work, as discussed in detail in Section III(3).

The CTE of the glasses obtained from the dilatometry is pre-sented in Table II. In general, the highest value of CTE was

obtained for glass 10D while the lowest was for glass 10. TheCTE increased slightly with the addition of Cr2O3 to the parentglass (glass 10), while it decreased slightly with the introductionof BaO in the glasses (10B). The decrease in CTE with the in-troduction of BaO implies an increasing polymerization in thestructure of glass 10B, as suggested by Vm data and will also bediscussed during the FTIR analysis in Section III(3). However,CTE increased with further increase in BaO for glasses 10C and10D due to decrease in Ve in the structural skeleton of the glass.

(3) FTIR Analysis

The room-temperature FTIR transmittance spectra of all thefive glasses investigated are shown in Fig. 1. All spectra exhibitfour broad transmittance bands in the region of 300–1500 cm�1.This lack of sharp features is indicative of the general disorder inthe silicate network, mainly due to a wide distribution of Qn

(polymerization in the glass structure, where n denotes the num-ber of bridging oxygens) units occurring in these glasses. Themost intense bands lie in the 800–1300 cm�1 region, the nextbetween 300–600 and 1350–1500 cm�1, while the least intensivelies between 650 and 800 cm�1. The broad bands in the 800–1300 cm�1 are assigned to the stretching vibrations of the SiO4

tetrahedron with a different number of bridging oxygen atoms,while the bands in the 300–600 cm�1 region are due to bendingvibrations of Si–O–Si and Si–O–Al linkages.18,19 In the presentstudy, the band for glass 10 in 800–1300 cm�1 region was reg-istered at an almost similar position as observed in our previousinvestigation for La2O3-containing diopside-based glasses.10

However, it is noteworthy that unlike the previous study,10 glass10 in the present study is BaO free. No considerable effect wasobserved in the position of the bands due to the addition ofCr2O3. The wave number values for transmittance bands in the800–1300 cm�1 region for glass 10 and 10A (B993 cm�1) areindicative of a major amount of Q2-type silicate species presentin the glass network. The introduction of BaO (glass 10B)shifted this wave number to a higher region (B1018 cm�1) im-plying an increasing polymerization in the silicate glass network.Further, a small shoulder was observed for glass 10B at B1110cm�1, which corresponds to the presence of Q3-type silicate spe-cies along with Q2, thus pinpointing toward an increasing po-lymerization in the silicate glass network (Fig. 1). The increasingconnectivity with the introduction of BaO content implies thatonly a part of BaO acts as network modifier. The remainingBaO might be acting as a network former and existing as inter-connected BaO4 tetrahedra.20 Similar network-forming actionof MgO due to the formation of MgO4 tetrahedron has beenreported for MgO–Na2O–B2O3 glasses.

21 With further increasein BaO, the wave number in the 800–1300 cm�1 region de-creased implying depolymerization of the silicate glass network.No band was resolved at B1100 cm�1 for glasses 10C and 10Dindicating the presence of a major amount of Q2-type silicatetrahedra in the glass network and depolymerized network incomparison with glass 10B. The transmittance bands in the 650–800 cm�1 region are related to the stretching vibrations of theAl–O bonds with Al31 ions in fourfold coordination.17 The

Table II. Properties of the Glasses

10 10A 10B 10C 10D

r (g/cm3) 2.9770.004 2.9870.001 3.0670.006 3.1070.003 3.1470.002Vm (cm3/mol) 19.6870.03 19.6770.01 19.9470.04 20.0870.02 20.1770.01Ve (cm

3/mol) 0.3370.02 0.3070.01 0.3270.04 0.3270.02 0.2970.01Tg (75) (1C) 756 756 755 756 758(CTE70.02) � 106 K�1 (2001–6001C) 8.39 8.43 8.40 8.71 9.04Tc (72) (1C) 885 885 904 906 914Tp (77) (1C) 920 920 937 946 955n 1.8070.005 1.9570.004 1.9770.007 2.0770.009 2.1570.006Ec (kJ/mol) 48774 47372 45974 39573 41276

Tg, Tc, and Tp values are for b5 20 K/min; CTE values were obtained at b5 5 K/min. Tg, glass transition temperature; Tc and Tp, crystalization phases.

832 Journal of the American Ceramic Society—Goel et al. Vol. 93, No. 3

transmittance band in the region 1350–500 cm�1 corresponds toB–O vibrations in the [BO3] triangle.

10

(4) Crystallization Kinetics by DTA

Figure 2 presents DTA thermographs of all the glasses investi-gated at a heating rate of 20 K/min. The Tg, obtained from themidpoint of the endothermic dip observed before the onset ofcrystallization, remained almost constant for all the investigatedglasses irrespective of any additions or substitutions made in theparent composition, as presented in Table II. These results are ingood agreement with the results obtained by Sohn et al.22 andMeinhardt et al.,23 according to whom, Tg and softening tem-perature (Ts) are independent of the BaO content in the glassesand are proportional to the B2O3/SiO2 ratio, which are alsovalid for the glasses under investigation in the present study(Table I). However, these results are in contradiction with theresults of Ghosh et al.,24 as they report a decrease in Tg withincreasing BaO content in the glasses. It is noteworthy thatalong with an increase in the amount of BaO, a considerabledecrease in SiO2 content in the glasses was also made by Ghoshet al.,24 while the exact amount of B2O3 in the glasses was notdisclosed. In the present investigation, the DTA plots of glassesexhibited single exothermic effects (Fig. 2) at all the heatingrates, which shifted toward higher temperatures with increase inthe heating rate. This signifies that the GC is formed either as aresult of single-phase crystallization or of an almost simulta-neous precipitation of different crystalline phases. The Tc and Tp

values for glasses 10 and 10A were almost similar (Table II);they shifted toward higher temperature with increase in BaOcontent in the glasses with glass 10D showing the highest valueof Tp among all the five glasses. These results are in good agree-ment with our previous study,25 while they contradict the resultsof Ghosh et al.24 The reason for this contradiction is similar tothat described for the discrepancies observed in Tg values.

The crystallization kinetics of the glasses were studied using theformal theory of transformation kinetics developed by Johnson

andMehl26 and Avrami,27–29 for nonisothermal processes that hasalready been obtained in our previous work30,31:

lnT2p

b

!¼ Ec

RTp� ln q ¼ 0 (1)

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

1nK0, respectively and the maximum crystallization rate by

the relationship:

dx

dt

����p

¼ 0:37bEcn RT2p

� ��1(2)

which makes it possible to obtain, for each heating rate, a value ofthe kinetic exponent, n. In Eq. (2), w corresponds to the crystal-lization fraction and dw

dt

��pcorresponds to the crystallization rate,

which may be calculated by the ratio between the ordinates of theDTA curve and the total area of the crystallization curve. Thevalues of Ec and Avrami parameter n for all the four glasses arelisted in Table II. An intermediate crystallization mechanism (bothsurface and bulk) was observed in all the glasses investigated. Theactivation energy of crystallization (Ec) slightly decreased with theaddition of Cr2O3. Lahl et al.

3 studied the crystallization behaviorof MgO–Al2O3–B2O3–SiO2 glasses and observed an increase inthe Ec value due to the addition of 0.6 mol% of Cr2O3 in theglasses. This increase in Ec value was explained on the basis ofdual role of Cr2O3 (i.e., nucleating agent and glass-network for-mer) in the glasses. However, it is noteworthy that they did notcalculate the mechanism of crystallization and assumed it to besurface crystallization, thus considering n51, which provides anuncertainty to their results. The addition of BaO in glasses 10Band 10C further led to a decrease in the value of Ec while it slightly

Fig. 2. Differential thermal analysis (DTA) of fine powders of theglasses investigated, heated at a rate of 20 K/min.

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

March 2010 Alkaline-Earth Aluminosilicate Glass–Ceramic Sealants for SOFC 833

increased with further increase in BaO content in glass 10D. TheEc values for all the glasses investigated in present study comparewell with those reported by Zhang et al.8 and Lahl et al.3 for theirsealing glasses.

(5) Sintering and Crystallization Process of Glasses

Sealing is usually applied on the surface (ceramic or metallic) tobe sealed using powder glass mixed with a binder. The GC for-mation involves the sintering of glass powders, followed bycrystallization at a higher temperature. During the sintering of aglass-powder compact with a size distribution of glass particles,small particles get sintered first as shown by Prado et al.32 Thus,sintering kinetics at first shrinkage (TFS) is dominated by theneck formation among the smallest particles by viscous flowand is best described by the Frenkel model of sintering.33 Max-imum shrinkage (TMS) is achieved when larger pores (poresformed from cavities among larger particles) have disappeareddue to the viscous flow that reduces their radii with time. Thisregion of sintering kinetics may be described by the Mackenzie–Shuttleworth model of sintering.34 However, various physicalprocesses (entrapped insoluble gases, crystallization) occurringat the very end of the sintering process might affect the densi-fication kinetics.

In the present investigation, sintering preceded crystallizationin all the five glass compositions as depicted in Fig. 3. Figure 4presents the photomicrographs of all the five glass compositionsdepicting the variation of sample dimensions with increase intemperature. As it is evident from the photomicrographs, dens-ification started at B8001C. The sintering initiated (TFS) atB7751–8001C in all the compositions while maximum shrinkage(TMS) was achieved well before the onset of crystallization (Tc)(Fig. 3). Similar results were obtained from the sintering and

crystallization experiments of the glass-powder compacts (rect-angular bars) heat treated at 8001 and 8501C for 1 h, respec-tively. Well-sintered, dense but amorphous glass-powdercompacts were obtained after heat treatment at 8001C for allthe compositions investigated. A considerable increase in me-chanical strength and density was observed for composition 10Aafter sintering at 8001C due to an introduction of Cr2O3 (TableIII) in the parent composition (glass 10). This shows that addi-tion of Cr2O3 improved the sinterability of the parent glass.Further, with the initial substitution of BaO for CaO (glass10B), the linear shrinkage and mechanical strength of the glass-powder compacts decreased while they increased with furtherincrease in BaO content for composition 10C after sintering at8001C. The highest value of linear shrinkage and mechanicalstrength after sintering at 8001C were observed for compositions10A (14.95%) and 10C (B115 MPa), respectively.

In accordance with HSM and DTA results (Figs. 3 and 4),well-sintered but crystallized glass-powder compacts were ob-tained after heat treatment at 8501C for 1 h. The photomicro-graphs of glass-powder compacts at 8501C (Fig. 4) showconsiderable shrinkage in their dimensions signifying densificat-ion. The evolution of the crystalline regime of the glass-powdercompacts (rectangular bars) heat treated at 8501C for 1 h isdemonstrated in the X-ray diffractograms, as shown in Fig. 5.Augite (Ca(Mg0.85Al0.15)((Si1.70Al0.30)O6), ICDD card: 01-078-1391) crystallized as the primary phase in all the investigatedcompositions along with diopside (CaMgSi2O6; ICDD card: 01-078-1390) as a minor phase (except GC 10A). It is noteworthythat augite is an intermediate member of the solid solutions inthe Diopside-Ca-Tschermak join.35 The quantitative analysis ofcrystalline phases in the GCs (Fig. 6, Table IV) reveals that theaddition of Cr2O3 increases the degree of crystallinity in GC 10Aand suppresses the crystallization of diopside, thus leading to the

Fig. 3. Comparison of differential thermal analysis (DTA) and hot-stage microscope (HSM) curves on the same temperature scale for compositions(a) 10, (b) 10B, and (c) 10D.

834 Journal of the American Ceramic Society—Goel et al. Vol. 93, No. 3

formation of a complete solid solution which in turn leads to anincrease in the mechanical strength of the resultant GC (166MPa). The introduction of BaO in GCs disturbs the formationof solid solution by promoting the formation of diopside alongwith augite while an increase in BaO content further led to thedissolution of augite leading to an increasing amount of amor-

phous glassy phase in the GCs (Table IV). A two-step shrinkagebehavior was observed for BaO-containing compositions as ob-served from the HSM data (Figs. 3(b) and (c)). This second step

Fig. 4. Hot-stage microscope (HSM) images of glasses on alumina substrates at various stages of heating cycle.

Table III. Properties of the Sintered Glass-CeramicsProduced from Glass-Powder Compacts After Heat Treatment

at Different Temperatures for 1 h

Composition 8001C 8501C

Density (g/cm3)10 2.9670.005 3.0270.00210A 2.9870.001 3.0370.00510B 3.0570.005 3.1770.00210C 3.1070.002 3.2270.00210D 3.1370.007 3.2770.003

Flexural strength (MPa)10 97.7710.9 156.076.110A 102.273.4 166.377.810B 90.0711.7 142.679.210C 115.478.6 158.375.110D 89.479.2 130.576.8 Fig. 5. Powder X-ray diffractograms of glass–ceramics sintered at

8501C for 1 h mixed with corundum (internal standard) (C, corundum).

March 2010 Alkaline-Earth Aluminosilicate Glass–Ceramic Sealants for SOFC 835

of shrinkage was not observed in BaO-free glass compositions,10 and 10A (Fig. 3(a)) and may be attributed to the strong ten-dency toward glass-in-glass phase separation as is evident fromthe existence of two glass transition temperatures (Tg1 and Tg2)in DTA curves of BaO-containing glasses (Fig. 3(c)).36 Thus,shrinkage also continued during the crystallization stage ofBaO-containing compositions as can also be seen from photo-micrographs presented in Fig. 4, where BaO-free glass compo-sitions (10 and 10A) depict no significant changes in thedimensions of samples until B12851C while BaO-containingcompositions (10B, 10C, and 10D) reveal a softening and de-formation, thus leading to a considerable shrinkage at B9501C.These results show that addition of BaO improves the flow be-havior of diopside-based GC sealants. Moreover, on the basis ofthe viscosity criteria defined by Pascual et al.,12 the viscosity ofthe BaO-containing sealants investigated is calculated to be inthe range of 106–108 dPa � s at SOFC operation temperature.These are the properties expected from good sealant materials.

Figure 7 presents the microstructure of the glass-powdercompacts sintered at 8001 and 8501C for 1 h. Consistent withXRD results (Fig. 5), the sintering of glasses in close proximityto TFS (B8001C) led to the softening and fusion of glass par-ticles resulting in a dense and amorphous glass-powder compactwith small amount of porosity (Fig. 7(a)). However, sinteringof glasses at 8501C resulted in dense and crystalline micro-structure for all the compositions investigated (Fig. 7(b)). Theseobservations are consistent with the relative importance of thetwo main stages of densification (i.e., first shrinkage and max-imum shrinkage).

IV. Conclusions

The structural characteristics along with sintering behavior andcrystallization kinetics of a new series of diopside-based glasseswere investigated in order to develop a suitable GC sealingcomposition for SOFC. An intermediate mechanism of crystal-lization was obtained for all the glasses while the Ec decreasedwith the addition of Cr2O3 and with an increase in BaO untilglass 10C. Infrared spectroscopy reveals that Si atoms would belikely centered on Q2 units. The introduction of BaO led to anincrease in the polymerization of glass network by the formationof Q3-type silicate species along with Q2. However, further in-crease in BaO depolymerized the silicate glass network.

Sintering preceded crystallization in all the glass compositionsinvestigated. BaO-free glass-powder compacts exhibited single-stage sintering behavior while a two-step shrinkage behaviorwas observed for BaO-containing glasses leading to an improvedflow behavior of the glasses, thus making them suitable for seal-ing applications in SOFC operating in the temperature range of8001–10001C.

The addition of Cr2O3 improved the sintering ability of theglasses, increased the crystalline content, and promoted the for-mation of augite-based solid solution in GC 10A, thus, resultingin highest mechanical strength (166 MPa). On the contrary, ad-dition of BaO led to the dissolution of augite so as to increasethe amount of amorphous phase in the GCs and encouraged theformation of diopside as a secondary phase along with augite,which degraded the mechanical strength of the GCs.

On the basis of results obtained in the present study, glasscompositions 10B and 10C might find application as efficientmetal–ceramic seals. An important evidence for this claim couldbe experimental results from our recent study.9 High thermalshock resistance and absence of any detrimental reaction be-tween the sealant and the metallic interconnect resulting in theformation of BaCrO4 was the main feature of this study. It wasrevealed that the electrical conductivity of GCs decreased withincrease in BaO content in the GCs, thus ensuring the absence ofshort-circuiting in the SOFC stack components. The CTE of theGCs increased with the addition of BaO and showed high sta-bility during long-term heat treatments. However, there is still apossibility of further sealant optimization, in particular via the

Fig. 6. Observed (crosses), calculated (continuous line), and differencecurve from the Rietveld refinement of the glass–ceramic 10B treated at8501C for 1 h in air. Markers representing the phase reflections corre-sponding to corundum, diopside, and augite (from top to bottom).

Table IV. Results of Rietveld-RIR Quantitative Analysis

10 10A 10B 10C 10D

Augite 83.33 (1) 94.08 (1) 92.39 (1) 79.48 (1) 70.26 (1)Diopside 7.78 (5) — 5.09 (3) 5.38 (3) 15.41 (4)Glass 8.89 (6) 5.92 (1) 2.52 (4) 15.14 (4) 14.33 (5)Total 100 100 100 100 100w2 3.77 5.08 3.35 3.55 4.51Rwp 0.059 0.067 0.057 0.056 0.063Rp 0.045 0.050 0.043 0.044 0.049

Fig. 7. Microstructure (revealed via SEM imaging after chemical etching of polished surfaces with 2 vol% of HF solution) of the glass–ceramics (a) 10Aheat treated at 8001C (b) 10B heat treated at 8501C for 1 h, respectively.

836 Journal of the American Ceramic Society—Goel et al. Vol. 93, No. 3

incorporation of nano- or submicrometer-sized crystalline addi-tives such as tetragonal partially stabilized zirconia (TZP).

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