Journal of Non-Crystalline Solids 356 (2010) 1070–1080

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Journal of Non-Crystalline Solids

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

Development and performance of diopside based glass-ceramic sealantsfor solid oxide fuel cells

Ashutosh Goel a, Dilshat U. Tulyaganov b, Maria J. Pascual c, Essam R. Shaaban d, Francisco Muñoz c,Zhe Lü e, José M.F. Ferreira a,*

a Department of Ceramics and Glass Engineering, University of Aveiro, CICECO, 3810-193 Aveiro, Portugalb Turin Polytechnic University in Tashkent, 100174 Tashkent, Uzbekistanc Instituto de Cerámica y Vidrio (CSIC), Kelsen 5, Campus de Cantoblanco, 28049 Madrid, Spaind Physics Department, Faculty of Science, Al-Azhar University, Assuit 71542, Egypte Center for Condensed Matter Science and Technology, Harbin Institute of Technology, Harbin 150001, PR China

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

Article history:Received 15 May 2009Received in revised form 24 December 2009Available online 10 February 2010

Keywords:CrystallizationGlass ceramicsX-ray diffractionScanning electron microscopyAluminosilicatesBorosilicatesSolid oxide fuel cellSealantSinteringNMR, MAS-NMR and NQR

0022-3093/$ - see front matter � 2010 Elsevier B.V. Adoi:10.1016/j.jnoncrysol.2010.01.012

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

We report on the development and performance of diopside (CaMgSi2O6) based glass-ceramic (GC) seal-ants for solid oxide fuel cells (SOFCs). Ten glass compositions were prepared by various additions andsubstitutions of La2O3, Cr2O3 and B2O3 in diopside system. The structure of the glasses has been investi-gated by employing 29Si, 27Al and 11B MAS-NMR spectroscopy. The crystallization kinetics of the glasseshas been investigated using differential thermal analysis (DTA) while X-ray diffraction (XRD) in conjunc-tion with Rietveld-RIR technique has been employed to quantify the amount of crystalline and amor-phous phases in the GCs. Hot-stage microscopy (HSM) has been used to investigate the influence ofLa2O3 and B2O3 on the sintering behavior of the glass powders. The joining and chemical interaction ofthe B2O3 and La2O3-containing glasses with Crofer22 APU and Sanergy HT (metallic interconnects forSOFC), has been investigated in humidified reducing atmosphere at SOFC operating temperature whilethe leak rate of GC sealants has been investigated using oxygen pump leak rate method. The suitable sin-tering and crystallization behavior of the GCs along with good adhesion to interconnect alloys and lowleak rate indicate that the investigated GCs are suitable candidates for further experimentation as SOFCsealants.

� 2010 Elsevier B.V. All rights reserved.

1. Introduction ity and improve the chemical and mechanical durability of the

Sealant materials based on glasses or GCs are critical in thedevelopment of planar type SOFC to provide reliable seals betweendifferent SOFC components in the stack. In addition to the gas-tightness, the essential requirements for the sealant material aresealing temperatures compatible with the working temperatureof the cell (800 �C), chemical stability in oxidizing and reducingatmosphere, low chemical activity to the cell components withwhich they come in contact, high insulating properties, matchingcoefficient of thermal expansion (CTE) and stress relaxation abilityduring operation.

Sealing is usually applied on the surface (ceramic or metallic) tobe sealed using powder glass mixed with a binder. The GC forma-tion involves the sintering of glass powders, followed by crystalli-zation at a higher temperature. In order to obtain a good sealing,the sintering stage should precede crystallization as dense andlow-porosity materials are desired for obtaining a gas-tight GC seal[1,2]. Further, crystallization is needed to increase the seal viscos-

ll rights reserved.

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sealant, which has to maintain the bulk stability and not flow dur-ing operation at high temperature. Therefore, for the developmentof a suitable GC sealant, it is necessary to understand the sinteringand crystallization behavior of the glass system [3].

In our previous studies, we proposed a La2O3-containing diop-side based GC seal for SOFC with very low amounts of BaO/SrOand B2O3 [4,5]. The GC seal demonstrated a sufficiently high CTE(�10 � 10�6 K�1) after heat treatment at SOFC operating tempera-ture for 300 h in air along with the absence of detrimental mono-clinic celsian phase. High electrical resistance and activationenergy of total conductivity, in combination with strong adhesion,and high chemical stability with metallic interconnects and 8YSZwere its other attributes. However, as mentioned in our previousinvestigation [5], we felt the need of making a detailed investiga-tion on the structural aspects, sintering behavior and crystalliza-tion kinetics of the parent glasses and the resultant GC sealantsin order to gain better understanding of the relationships betweenmolecular-level glass structure and crystallization parameters.Therefore, this sequel to our previous studies [4,5], is an attemptto fulfill that lacuna by investigating the structure, sinteringbehavior and crystallization kinetics of the glasses studied in our

A. Goel et al. / Journal of Non-Crystalline Solids 356 (2010) 1070–1080 1071

previous work using magnetic angle spinning-nuclear magneticresonance (MAS-NMR), HSM and differential thermal analysis(DTA), respectively. Further, chemical interaction and adhesion be-tween GC sealants and metallic interconnects has been investi-gated in humidified reducing atmosphere. Also, in our previousstudy [5], we had proposed Sanergy HT to be a better candidatefor interconnect in SOFC in comparison to Crofer22 APU alloydue to low Si content and CTE of the former. In the present article,we aim to shed some more light on the behavior of the two inter-connect alloys in competition with each other under humidifiedreducing atmosphere. Further, leak rate of the GC sealants whenin contact with YSZ (zirconia stabilized by yttria) has been mea-sured using oxygen pump leak rate method.

2. Experimental

2.1. Glass preparation

A total of 10 glasses were prepared in Pt-crucibles using melt-quenching technique. The detailed substitutions and experimentalmethod for preparation of glasses are reported in our previousworks [4,5]. Table 1 summarizes the compositions of all the inves-tigated glasses.

Glasses in bulk form were produced by pouring the melts onpreheated bronze molds followed by annealing at temperaturesnear glass transition (Tg). The samples of the glass-powder com-pacts were produced from glass frits, which were obtained byquenching of glass melts in cold water. The frits were dried andthen milled in a high-speed agate mill resulting in fine glass pow-ders with mean particle sizes of 10–15 lm (determined by lightscattering technique; Coulter LS 230, Beckman Coulter, FullertonCA; Fraunhofer optical model).

2.2. Structural characterization of glasses

29Si MAS-NMR spectra were recorded on a Bruker ASX 400 spec-trometer operating at 79.52 MHz (9.4 T) using a 7 mm probe at aspinning rate of 5 kHz. The pulse length was 2 ls and 60 s delaytime was used. Kaolinite was used as the chemical shift reference.27Al MAS-NMR spectra were recorded on a Bruker ASX 400 spec-trometer operating at 104.28 MHz (9.4 T) using a 4 mm probe ata spinning rate of 15 kHz. The pulse length was 0.6 ls and 4 s delaytime was used. Al(NO3)3 was used as the chemical shift reference.11B MAS-NMR spectra were recorded on a Bruker ASX 400 spec-trometer operating at 128.36 MHz (9.4 T) using a 4 mm probe ata spinning rate of 12 kHz. The pulse length was 3.6 ls and 2 s delaytime was used. H3BO3 was used as the chemical shift reference.

2.3. Thermal analysis of glasses

The Tg was obtained by dilatometry measurements on prismaticsamples with a cross-section of 4 � 5 mm2 (Bahr Thermo Analyze

Table 1Batch compositions of the glasses (wt%).

Glass MgO CaO BaO SrO SiO2

7A 16.90 18.82 6.43 – 47.87B 14.52 17.95 6.14 – 43.27A-Cr 16.82 18.72 6.40 – 47.67B-Cr 14.44 17.86 6.11 – 43.07-2B 16.56 18.44 6.30 – 46.97-5B 16.05 17.87 6.11 – 45.47-10B 15.20 16.92 5.78 – 43.07-15B 14.34 15.97 5.46 – 40.67(Sr)-2B 16.92 18.83 – 4.35 47.97(Sr)-5B 16.40 18.25 – 4.22 46.4

DIL 801 L, Hüllhorst, Germany; heating rate 5 K min�1). A mini-mum of three samples for each composition were analyzed. DTAof fine powders was carried out in air (DTA-TG, Setaram Labsys,Setaram Instrumentation, Caluire, France). In order to investigatethe non-isothermal crystallization kinetics, 50 mg of glass powderswere heated up to 1000 �C with different heating rates (b) of 5, 20,30, and 40 K min�1.

A side-view hot-stage microscope (HSM) EM 201 equipped withimage analysis system and electrical furnace 1750/15 Leica wasused. The microscope projects the image of the sample through aquartz window and onto the recording device. The computerizedimage analysis system automatically records and analyzes thegeometry changes of the sample during heating. The image ana-lyzer takes into account the thermal expansion of the alumina sub-strate while measuring the height of the sample during firing, withthe base as a reference. The HSM software calculates the percent-age of decrease in height, width and area of the sample images. Themeasurements were conducted in air with a heating rate of5 K min�1. The cylindrical shaped samples with height and diame-ter of �3 mm were prepared by cold-pressing the glass powders.The cylindrical samples were placed on a 10 � 15 � 1 mm alumina(>99.5 wt% Al2O3) support. The temperature was measured with aPt/Rh (6/30) thermocouple contacted under the alumina support.The temperatures corresponding to the characteristic viscositypoints (first shrinkage, maximum shrinkage, softening, half balland flow) were obtained from the photographs taken during thehot-stage microscopy experiment following Scholze’s definitionand [6,7].

2.4. Crystalline phase evolution in glass-ceramics

Rectangular bars with dimensions of 4 � 5 � 50 mm3 were pre-pared for glass powders 7A, 7A-Cr, 7B and 7B-Cr while circular diskshaped pellets with Ø 20 mm and thickness �3 mm were preparedfrom glass powders of B2O3-containing glasses by uniaxial pressing(80 MPa) [4,5]. The pellets were sintered under non-isothermalconditions for 1 h at different temperatures between 800 and900 �C, respectively. A slow heating rate of 2 K min�1 was main-tained in order to prevent deformation of the samples. The amor-phous nature of glasses and qualitative crystalline phase analysisof GCs was made by XRD analysis (Rigaku Geigerflex D/Max, C Ser-ies, Tokyo, Japan; Cu Ka radiation; 2h angle range 10–80�; step0.02�/s). The quantitative phase analysis of GCs (crushed to particlesize <25 lm) was made by XRD analysis using a conventionalBragg–Brentano diffractometer (Philips PW 3710, Eindhoven, TheNetherlands) with Ni-filtered Cu Ka radiation in conjunction withcombined Rietveld-RIR method. A 10 wt% of corundum (NIST SRM674a, annealed at 1500 �C for 1 day to increase the crystallinity to100%) was added to all the GC samples as an internal standard. Themixtures, ground in an agate mortar, were side loaded in alumi-num flat holder in order to minimize the preferred orientationproblems. Data were recorded in 2h range = 5–140� (step size

Al2O3 La2O3 B2O3 Cr2O3 NiO

8 2.14 6.83 – – 18 4.08 13.04 – – 14 2.13 6.80 – 0.5 16 4.06 12.97 – 0.5 11 2.09 6.70 2 – 16 2.03 6.49 5 – 14 1.92 6.14 10 – 12 1.81 5.80 15 – 12 2.14 6.84 2 – 14 2.07 6.63 5 – 1

Figure 1 Leak Leak P(O2)

P0

Solartron SI 1287 Electrochemical Interface

WE CE RE1RE2

Pump

sensor OCV

I

O2-

Fig. 1. Schematic representation of leak rate measurement set up.

1072 A. Goel et al. / Journal of Non-Crystalline Solids 356 (2010) 1070–1080

0.02� and 6 s of counting time for each step). The phase fractionswere extracted by Rietveld-RIR refinements, using GSAS softwareand EXPGUI as graphical interface, were rescaled on the basis ofthe absolute weight of corundum originally added to their mix-tures as an internal standard, and therefore, internally renormal-ized. The background was successfully fitted with a Chebyshevfunction with a variable number of coefficients depending on itscomplexity. The peak profiles were modeled using a pseudo-Voigtfunction with one Gaussian and one Lorentzian coefficient. Latticeconstants, phase fractions, and coefficients corresponding to sam-ple displacement and asymmetry were also refined.

2.5. Joining behavior and chemical interaction between glass-ceramicsealant and metallic interconnect in humidified reducing atmosphere

In order to investigate the adhesion and chemical interaction ofthe glasses with SOFC components, wetting experiments betweenglass powders and metallic interconnect were carried out underdifferent conditions. Two different metallic interconnect materials,namely, Crofer22 APU (Thyssen Krupp, VDM, Werdohl, Germany)and Sanergy HT (Sandvik AB, Sandviken, Sweden) were employedin as received and pre-oxidized (850 �C for 2 h, 10 K min�1) form.The composition of both the interconnect alloys is given elsewhere[5]. Further, a third interconnect material derived from Sanergy HTby deposition of a 200 nm thick Co-layer on it was obtained fromSandvik AB and was tested with the glasses for wetting experi-ments. The deposition of Co-layer was made as an attempt toreduce the phenomenon of Cr-poisoning of cathodes and Cr-diffusion from interconnect to the sealant. The glass powders weredeposited on the interconnect alloys by slurry coating. Heat treat-ment was performed in a tubular furnace, without applying anydead load, in humidified reducing atmosphere (a mixture of 95%Ar + 5% H2 passed through water bath). The diffusion couples wereheated to 850 �C with a relatively slow heating rate (2 K min�1)and kept at that temperature for 1 h. Finally, the temperaturewas brought down to SOFC operating temperature (i.e., 800 �C)and maintained at this temperature for 300 h. Scanning electronmicroscopy (SEM; SU-70, Hitachi, Japan) and energy dispersivespectroscopy (EDS; Bruker Quantax, Germany) was employed tostudy the microstructure and distribution of elements along theinterface of electrolyte–glass-interconnect diffusion couples.

2.6. Experimental and theory of leak rate measurements

In order to measure the leak rate of GC sealants, silver paste(DAD-87 type) was applied on one side of dense YSZ (zirconia sta-bilized by yttria) pellets to form the electrodes, and silver wirewith a diameter of 0.1 or 0.2 mm was attached on each electrode.A pair of YSZ pellets with electrodes and wires adhered togetherwere heated at 150 �C for about 20 min for solidification and firedat 650 �C in order to burn out the organic products in silver pasteand to achieve a good join between silver electrodes and wires.The GC powders were mixed with an organic binder to form seal-ant paste. The sealant paste was painted on the circumjacent junc-ture of YSZ pellets in order to seal them together to form a cell withsmall closed gas chamber. After painting, the cell was baked at 80–90 �C for about 20 min. Such painting and baking process was re-peated for three or four times until there was no obvious defecton the sealing region. The cells were fired in a programmed electri-cal furnace using the following temperature program in order toobtain a cell with both oxygen pump and sensor for the leak ratetesting

20 �C !4 �C=min150 �C� 60 min !4:8 �C=min

870 �C� 120 min

!�1:5 �C=min800 �C !�1 �C=min

750 �C !�1:5 �C=min600 �C !�2 �C=min

50 �C

The schematic representation of the leak rate measuring systemis presented in Fig. 1. The pump-sensor system has four silverwires. Two wires for inner electrodes were connected to WE andRE2 terminals of an electrochemical interface of Solartron SI1287, while the remaining two for outer electrodes were con-nected to CE and RE1, respectively (Fig. 1). Thus, an oxygen pumpwas composed between the electrodes connected to WE and CEand an oxygen sensor was composed between the electrodes con-nected to RE1 and RE2. The current was applied between CE andWE to pump out oxygen from the gas chamber surrounded withYSZ pellets and sealant until the oxygen concentration voltage(OCV) achieved �50 mV (P(O2) � 0.02 atm), or feed oxygen into ituntil OCV equaled 0 V. The current range is from 100 to 500 lA.

When the oxygen in chamber is pumped out, both the oxygenconcentration (mol) and the oxygen partial pressure (P(O2)) inthe chamber decreases. At the same time, oxygen may leak intothe chamber through the sealing region with the driving of pres-sure difference (DP(O2) = P0 � P(O2), P0 is the oxygen partial pres-sure in air, i.e., 0.21 atm). Assuming that leak rate is proportionalto DP(O2), the derivative oxygen concentration (mol) in chamberwith respect to time is related to the pumping current and oxygenleak rate under DP(O2)

dndt¼ �qnðPðO2Þ � P0Þ �

I4F

; ð1Þ

where F is Faraday’s constant and qn is rate of oxygen leakage(mol s�1). According to the gas equation and assuming a constantchamber volume,

PðO2Þ ¼RTV� n; ð2Þ

thus,

dPðO2Þdt

¼ �RTV

qnðPðO2Þ � P0Þ �RT4F� 1V� I: ð3Þ

For a well sealed chamber with low leakage, qn is a small num-ber and it causes the second item is much larger than the first one.At the end of pumping process, P(O2)� P0

dPðO2Þdt

� RTV

qnP0 �RT4F� 1V� I ¼ RT

VqnP0 1� I

IL

� �; ð4Þ

where IL is named limited current and symbolises 4FqnP0. Perform-ing an integral of Eq. (4), we get

PðO2Þ ¼ P01 þRTV

qnP0 1� IIL

� �� t; ð5Þ

A. Goel et al. / Journal of Non-Crystalline Solids 356 (2010) 1070–1080 1073

where t is a short time near the end of pumping time. So the P(O2) isproportional to time and we can calculate the qn and V by linear fit-ting P(O2) � t data to get the slope K

K ¼ RTV

qnP0 1� IIL

� �: ð6Þ

For different pumping current I1 and I2, the slopes are K1 and K2,respectively. Then we can obtain both qn and V by solvingequations:

V ¼ RT4F

I2 � I1

K1 � K2; ð7Þ

qn ¼1

4FP0� K1I2 � K2I1

K1 � K2: ð8Þ

Using the qn data, leak rate data can be calculated using the fol-lowing equation:

Q V ¼ Vmolqn

ls� DP; ð9Þ

where Vmol 22.4 L/mol, ls is the sealing length (about 3.26 cm in thisexperiment), DP is the pressure difference.

3. Results and discussion

3.1. NMR structural characterization

Fig. 2(a) depicts the 29Si MAS-NMR spectra of the glasses 7-2B,7-5B, 7-10B and 7-15B, all showing a broad resonance band from Sitypically found in glasses. The resonance bands of 7-2B and 7-5Bglasses are centered at �82 ppm, while the spectra of 7-10B and7-15B glasses clearly present a little wider resonance bands andslightly shifted down to �85 ppm. These chemical shift valuesare indicative of a major amount of Q2-type silica tetrahedra pres-ent in the glass network. The high field shifting and slight broaden-ing of the resonance bands in 7-10B and 7-15B glasses might beinterpreted as re-polymerization of the silica glass sub-networkas boron content increases in the glasses.

The 11B MAS-NMR spectra of the same series of glasses inshown in Fig. 2(b). Broad resonance bands can be observed within the chemical shift range between 0 and �25 ppm. This higherfield chemical shift could be attributed to the higher ionic fieldstrength (IFS) of the modifier elements, Mg2+, Ca2+ and Ba2+, whencompared to alkaline ones. The spectra of the 7-2B and 7-5B

7-2B

7-5B

7-10B

7-15B

-150001- 521--75-50-25

29Si Chemical shift (ppm)(a)

7-2B

7-5B

7-10B

7-15B

-1001020

11B Chemical s(b)

Fig. 2. (a) 29Si MAS-NMR spectra, (b) 11B MAS-NMR spectra

glasses present resonance bands with typical of quadrupolar boronatoms in threefold coordination, i.e., BO3 triangles, meanwhile theydo not allow to infer any presence of Gaussian-shape bands corre-sponding to fourfold coordinated boron atoms, i.e., BO4 units. How-ever, 7-10B and 7-15B NMR spectra show an increase in theresonance contribution at �20 ppm, Gaussian type, which couldindicate the appearance of a signal due to BO4 units as the boroncontent increased in the glasses.

Fig. 2(c) depicts the 27Al MAS-NMR spectra of glass 7A and otherB2O3-containing glasses. All the spectra show an asymmetric andvery broad resonance bands typical for quadrupolar aluminumatoms, between 100 and �50 ppm, centered at around 60 ppm.At this level of magnetic field the resolution does not allow to dis-criminate between different aluminum coordination. According tothe chemical shift, most of aluminum atoms would be fourfoldcoordinated, i.e., forming AlO4 tetrahedra. A small shoulder canalso be appreciated at around 20 ppm for the glasses 7A, 7-2Band 7-5B, which is not clearly seen in glasses 7-10B and 7-15B,indicating that part of aluminum could be of the type Al(V) forthe 7A glass and glasses with low boron contents, disappearingfor further boron additions.

3.2. Sintering and crystallization behavior by HSM and DTA

During the sintering of a glass-powder compact with a size dis-tribution of glass particles, small particles get sintered first asshown by Prado et al. [8]. Thus, sintering kinetics at first shrinkageis dominated by the neck formation among smallest particles byviscous flow and is best described by the Frenkel model of sintering[9]. Maximum shrinkage is reached when larger pores (poresformed from cavities among larger particles) have disappeareddue to viscous flow that reduces their radii with time. This regionof sintering kinetics may be described by the Mackenzie–Shuttle-worth model of sintering [10]. However, various physical processes(entrapped insoluble gases, crystallization) occurring at the veryend of sintering process might affect the densification kinetics. Acomparison between DTA and HSM results under the same heatingconditions can be useful to investigate the effect of glass composi-tion on sintering and devitrification phenomena. In general, twodifferent trends can be observed related to the sintering and crys-tallization behavior of the glasses [2]. In the first case, the begin-ning of crystallization (Tc) occurs after the final sintering stage.Thus, under such circumstances, sintering and crystallization areindependent processes. However, in other case, Tc appears before

-40-30-20

hift (ppm)

7A

7-2B7-5B

7-10B

7-15B

001- 002-0002 001

27Al Chemical shift (ppm)(c)

, (c) 27Al MAS-NMR spectra of the investigated glasses.

7ATp

Tc

DTA

TMS

TFSHSM

-6

-4

-2

0

2

4

600 700 800 900 1000

Temperature (oC)

Hea

t Flo

w (

μV)

0.6

0.7

0.8

0.9

1

1.1

A/A

0

β = 5Kmin-1

(a)

7B

Tp

Tc

DTA

HSM

TMS

TFS

-8

-6

-4

-2

0

2

4

600 700 800 900 1000

Temperature (oC)

Hea

t Flo

w (

μV)

0.5

0.6

0.7

0.8

0.9

1

1.1

A/A

o

β = 5Kmin-1

(b)

7-2B

Tp

Tc

DTA

TMS

TFSHSM

-10

-8

-6

-4

-2

0

2

600 700 800 900 1000

Temperature (oC)

Hea

t Flo

w (

μV)

0.55

0.65

0.75

0.85

0.95

1.05

A/A

0

β = 5Kmin-1

(c)

7(Sr)-2B

Tp

Tc

DTA

TMS

TFSHSM

-7

-5

-3

-1

1

3

600 700 800 900 1000

Temperature (oC)

Hea

t Flo

w (

μV)

0.6

0.7

0.8

0.9

1

1.1

A/A

0

β = 5Kmin-1

(d)

Fig. 3. Comparison of DTA and HSM curves on the same temperature scale forcompositions (a) 7A, (b) 7B, (c) 7-2B, (d) 7(Sr)-2B.

1074 A. Goel et al. / Journal of Non-Crystalline Solids 356 (2010) 1070–1080

maximum density has been reached. In this case, the crystalliza-tion process starts before complete densification, thus, preventingfurther sintering.

On the basis of the relation between the temperatures mea-sured by HSM and their corresponding viscosities, Scholze [11] de-fined the following characteristic viscosity points whose viscosityvalues were more accurately determined in [7]:

(a) First shrinkage or sintering (TFS): the temperature at whichthe pressed sample starts to shrink (log g = 9.1 ± 0.1), whereg is the viscosity.

(b) Point of maximum shrinkage (TMS): the temperature at whichmaximum shrinkage of the glass-powder compact takesplace before it starts to soften (log g = 7.8 ± 0.1).

(c) Softening point (TD): the temperature at which the first signsof softening are observed. This is generally shown by the dis-appearance or rounding of the small protrusions at the edgesof the sample (log g = 6.3 ± 0.1).

(d) Half ball point (THB): the temperature at which the section ofthe sample observed forms a semicircle on the microscopegrid (log g = 4.1 ± 0.1).

(e) Flow point (TF): the temperature at which the maximumheight of the drop of the molten glass corresponds to a uniton the microscopic scale (log g = 3.4 ± 0.1).

In our previous study, it was observed that glasses withB2O3 > 5 wt% showed poor sinterability and the resultant GCs werehighly porous, thus rendered unfit for sealing applications in SOFC[5]. Therefore, glasses with B2O3 > 5 wt% were excluded from HSMinvestigation in the present study. Also, since the molar concentra-tion of BaO in 7-2B and 7-5B is equal to that of SrO in 7(Sr)-2B and7(Sr)-5B, respectively. Therefore, in order to observe the effect ofreplacement of BaO by SrO on the sintering ability and crystalliza-tion of the glass powders, glass 7(Sr)-2B was chosen and the resultswere compared with those of glass 7-2B. Further, even thoughCr2O3-containing glass compositions showed favorable propertiesfor sealant applications [4], they are not investigated here throughHSM as there detailed analysis will be presented in our forthcom-ing article. Therefore, in the present investigation, HSM data inconjunction with DTA data will be used to study the effect ofLa2O3 and B2O3 (65 wt%) on the sintering ability of diopside basedGC sealants. However, a detailed analysis of crystallization kineticson all the 10 glass compositions will be presented and discussedalong the manuscript.

The variation in the relative area and heat flow with respect totemperature as obtained from HSM and DTA, respectively, forglasses 7A, 7B, 7-2B and 7(Sr)-2B at heating rate (b) = 5 K min�1

is presented in Fig. 3 while Fig. 4 presents photomicrographs ofthe changes in geometric shape of the glasses 7A, 7B, 7-2B and7-5B with respect to temperature, as obtained from HSM. Table 2lists the characteristic temperatures for all the five glasses as ob-tained by dilatometry (Tg), HSM (TFS, TMS, TD, THB, TF) and DTA (Tc,Tp). A/A0 corresponds to the ratio of final area/initial area of theglass-powder compact. A/A0 � 0.65 at TMS (as presented in Table2) implies towards good densification (95–98%) [2].

3.2.1. Effect of La2O3

Fig. 3(a) and (b) presents the DTA and HSM data for glasses 7Aand 7B, respectively, while their characteristic temperatures arelisted in Table 2. It is evident that Tg considerably decreases, whileTc increases with increasing La2O3 content. Therefore, the differ-ence between Tc and Tg (i.e., Tc � Tg) is higher for glass 7B in com-parison to glass 7A implying towards an increase in glass stabilityversus crystallization. Sinterability parameter, as defined by Tc

� TMS [2] is also higher for 7B (Tc � TMS = 64) in comparison to7A (Tc � TMS = 48). These two factors are indicative of a better

sintering-crystallization behavior of 7B composition. These resultsare in accordance with the linear shrinkage values of GCs 7A and

Fig. 4. HSM images of glasses on alumina substrates at various stages of heating cycle.

Table 2Thermal parameters obtained from dilatometry, DTA and HSM for the investigated glasses at b = 5 K min�1.

Glass Tg ± 2 (�C) TFS ± 5 (�C) TMS ± 5 (�C) Tc ± 2 (�C) TD ± 5 (�C) Tp ± 2 (�C) THB ± 5 (�C) TF ± 5 (�C) A/A0

7A 685 783 834 882 883 912 1307 1323 0.687B 630 790 844 908 884 951 1257 1312 0.647A-Cr 675 – – 871 – 895 – – –7B-Cr 672 – – 886 – 916 – – –7-2B 580 776 829 863 885 896 1278 1318 0.647-5B 575 754 809 845 842 874 1239 1308 0.667-10B 530 – – 823 – 859 – – –7-15B 490 – – 819 – 855 – – –7(Sr)-2B 533 769 832 857 850 891 1286 1355 0.677(Sr)-5B 545 – – 840 – 871 – – –

A. Goel et al. / Journal of Non-Crystalline Solids 356 (2010) 1070–1080 1075

7B presented in our previous study [4], where composition 7Bshowed higher linear shrinkage in comparison to 7A. The TD forboth the compositions was almost unaffected by the La2O3 content(Table 2). With further increase in temperature, half ball (THB) wasformed for glass 7B at 1257 �C, which is 50 �C lower than for glass

7A (Fig. 4). This was followed by the complete flow of glass 7B atlower temperature than that for glass 7A. The decrease in THB

and TF for composition 7B may be attributed to the presence oflow melting crystalline phase calcium lanthanum oxide silicate(ICDD: 01-071-1368) along with augite (ICDD: 01-078-1392)

Table 3Activation energy of crystallization (Ec) and Avrami parameter (n) for all theinvestigated glasses.

Glass Ec (kJ mol�1) N

7A 300 1.85 ± 0.0097B 274 2.03 ± 0.0177A-Cr 213 2.16 ± 0.0157B-Cr 198 2.18 ± 0.0037-2B 255 2.22 ± 0.0077-5B 236 1.77 ± 0.0037-10B 207 1.78 ± 0.0047-15B 184 2.21 ± 0.0047(Sr)-2B 266 1.82 ± 0.0037(Sr)-5B 260 1.67 ± 0.003

1076 A. Goel et al. / Journal of Non-Crystalline Solids 356 (2010) 1070–1080

instead of a mono-mineral GC with augite as the only crystallinephase, as was observed for composition 7A [4].

3.2.2. Effect of B2O3

The Tg values for glasses decreased with increase in B2O3 con-tent in the glasses (Table 2). With respect to a comparison betweensintering behavior of glasses 7A, 7-2B and 7-5B, the temperaturefor initiation of sintering (TFS) decreases with increase in B2O3 con-tent in the glasses (Table 2). The temperature for maximum shrink-age (TMS) showed a small decrease with initial addition of B2O3 forglass 7-2B in comparison to glass 7A while it decreased consider-ably with addition of 5 wt% B2O3 in glass 7-5B. In all the threeinvestigated glass compositions, sintering and crystallization actas independent processes. The Tc, Tp, TD and THB values decreasedwith increase in B2O3 content in the glasses (Table 2, Fig. 4). TheTF decreased with addition of B2O3 in the parent glass 7A; however,the decrease was not as significant as it was in case of THB. Whenincreasing the temperature, boron tends to diminish the coordina-tion index forming triangular units from tetrahedral units, provok-ing an additional diminution of viscosity at elevated temperaturesjoined to the thermal effect. This effect is more pronounced whengreater is the proportion of boron oxide in the glass composition.

3.2.3. Effect of replacement of BaO by SrOThe Tg value decreased considerably with replacement of BaO

(7-2B) by SrO (7-5B) which leads to a greater glass stability(Tc � Tg) for the glass with SrO. Similar trend was observed duringinitiation of sintering as sintering started (TFS) in glass composition7(Sr)-2B at slightly lower temperature than 7-2B (Table 2). Sinter-ing preceded crystallization in both the glass compositions. Theseresults are in accordance with the values of linear shrinkage for7-2B and 7(Sr)-2B as obtained in our previous investigation [5],where higher linear shrinkage values were obtained for GCs con-taining SrO in comparison to their BaO containing counterparts.A slight decrease was observed for Tc, Tp and TD due to replacementof BaO by SrO while temperatures for THB and TF for 7(Sr)-2B werealmost similar.

It should be mentioned that sintering preceded crystallizationin all the glass compositions in accordance with wide Tc � Tg rangefor all the five compositions and no considerable change in theshape of the samples was observed in the range of SOFC operatingtemperature, i.e., 850–1000 �C. Moreover, on the basis of the vis-cosity criteria defined by Scholze [11], viscosity of the investigatedsealants at SOFC operation temperature comes out to be in therange of 106–108 dPa s. These are the properties expected fromgood sealant materials. However, still there is need to tailor theflow properties of these GC sealants by reducing half ball and flowtemperatures.

3.3. Crystallization kinetics of glasses

Although numerous sealants have been proposed for SOFC [12],very few studies related with crystallization kinetics of GC sealantshave been reported in literature [3,13–15]. Crystallization processplays an important role in determining the properties and applica-tions of GC sealants. For example, an installation process for theSiemens-SOFC required that the sealing glass should be partiallyviscous at 950 �C for 2–3 h to allow small displacement of the sin-gle stack elements after joining at 1000 �C. This can be achieved byusing a slow crystallizing glass [16]. Also, the viscosity of GC seal-ants is affected by the crystallization of parent glasses. Sakaki et al.[17] observed an increase in seal viscosity at about 950 �C becauseof bulk crystallization of wollastonite (CaSiO3) in CaO–Al2O3–SiO2

glass system. Moreover, as discussed in the previous section, sin-tering should precede crystallization in order to obtain a densematerial suitable for a SOFC seal. Uncontrolled crystallization

during the initial sintering process can adversely affect the SOFCoperation.

In the present investigation, crystallization kinetics of variousGC sealants has been studied using the formal theory of transfor-mation kinetics as developed by Johnson and Mehl [18] andAvrami [19], for non-isothermal processes that has already beenobtained in our previous work [20,21]:

lnT2

p

b

!¼ Ec

RTp� ln q ¼ 0; ð10Þ

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

����p

¼ 0:37bEcnðRT2pÞ�1; ð11Þ

which makes it possible to obtain, for each heating rate, a value ofthe kinetic exponent, n. In Eq. (11), v corresponds to the crystalliza-

tion fraction and dvdt

���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. The values ofEc and n for all the four glasses are listed in Table 3. The correspond-ing mean values may be taken as the most probable value of thequoted exponent. In accordance with our previous results [13], anintermediate crystallization mechanism (both surface and bulk)has been observed in all the investigated glasses (Table 3). Theinconsistency in the results of Ec values with those reported inour previous study [4] is due to the fact that in the present investi-gation the calculations have been made in accordance with the ac-tual crystallization parameters. The glass composition 7A showedthe highest Ec value among all the investigated sealant composi-tions. The Ec of La2O3 containing Di-based GC sealants is lower thanLa2O3-free Di–Ca–Ts GC sealants [13] while Ec decreases with fur-ther increase in La2O3 content. This decrease in Ec with increase inLa2O3 may be due to the network modifying effect and clusteringof La3+ ions (La–O–La linkages) in the glass structure [4]. The Ec de-creased with addition of Cr2O3 in the glasses can be ascribed to thenucleation effect of the former. The Ec values decreased consistentlywith increase in B2O3 amount in the glasses with glass 7-15B show-ing the lowest Ec value. The decrease in Ec with B2O3 addition maybe attributed to the lower activation energy of viscous flow of B2O3

which allows the crystallization to occur at lower temperature. Thereplacement of BaO by SrO led to an increase in the Ec values ofglass 7(Sr)-2B and 7(Sr)-5B in comparison to glasses 7-2B and 7-5B. The quantitative phase analysis was made on four GC sealantcompositions (7-2B, 7-5B, 7(Sr)-2B and 7(Sr)-5B; sintered at850 �C for 1 h) in order to elucidate the effect of B2O3 additionand Ba2+

M Sr2+ substitution on the amount and nature of crystal-

A. Goel et al. / Journal of Non-Crystalline Solids 356 (2010) 1070–1080 1077

line and amorphous phases in the GCs. The results reveal that diop-side based crystalline phases crystallized in all the GCs (Fig. 5, Table4) and all the resultant GCs have high degree of crystallization(glassy phase: 10–40%). The Ec value for glass 7A is higher in com-parison to barium calcium aluminum silicate (BCAS) GC sealantwhile Ec values for glasses 7-2B and 7-5B are lower than BCAS GCsealant [3]. The Ec values for all the glasses investigated in presentstudy are much lower than that reported for 19.2CaO–18.5SrO–13.2ZnO–1.9B2O3–2.9Al2O3–2.0TiO2–42.2SiO2 (mol%) GC sealantproposed by Zhang et al. [14] and by Lahl et al. [22] for their sealingglasses.

3.4. Interfacial adhesion and chemical stability of interconnect/GCsealant diffusion couples after heat treatment in humidified reducingatmosphere

Two SOFC interconnect alloys, Crofer22 APU and Sanergy HT,were investigated for their chemical compatibility with the inves-tigated glasses. All the sealing GCs bonded well to the metallic

Fig. 5. Observed (crosses), calculated (continuous line), and difference curve from the Rithe phase reflections correspond to corundum, Sr-diopside (01-080-0388), and augite (0

Table 4Results of Rietveld-RIR quantitative analysis (wt%).

7-2B

Augite (01-078-1391)Ca(Mg0.85Al0.15)((Si1.70Al0.30)O6)

44.77 (1)

Augite (01-078-1392)Ca(Mg0.70Al0.30)((Si1.70Al0.30)O6)

–

Diopside (01-078-1390)CaMgSi2O6

42.39 (1)

Diopside (01-080-0388)Ca0.75Sr0.2Mg1.05(Si2O6)

–

Glass 12.84 (2)Total 100v2 1.60Rwp 0.092Rp 0.073

RIR, reference intensity ratio.

interconnects and no gaps were observed even at the edges ofthe joints. Fig. 6 shows a SEM micrograph of the cross-section ofinterface between Crofer22 APU/7-2B GC sealant after 300 h ofexposure to wet fuel gas atmosphere. The extent of diffusion ofconstituent elements from alloy interconnects into the glass andvice versa has been observed through EDS element mapping andline scanning. A rather smooth interface between GC sealant andsteel is observed without the presence of iron-rich oxide productseven at the edges. According to thermodynamic data available inthe literature [23], the oxygen partial pressure required to oxidizeiron to wustite (FeO) at 800 �C should exceed 1.71 � 10�19 atm.Therefore, we can expect the oxygen partial pressure to be lowerthan that which is required to oxidize iron. Furthermore, no sub-stantial Ba-, Cr-rich oxide zone was observed at the interface(Fig. 6(a)) as was evident after heat treatment in air atmosphere[4,5]. However, a Mn-rich oxide layer was formed at the steel sideof the interface (Fig. 6(a) and (b)). Fig. 7 presents the SEM micro-graphs of the cross-section of interface between Sanergy HT/GCsealants 7A (Fig. 7(a)), 7-2B (Fig. 7(b)) and 7-5B (Fig. 7(c)), respec-

etveld refinement of the GC 7(Sr)-2B treated at 850 �C for 1 h. Markers representing1-078-1391) (from top to bottom).

7-5B 7(Sr)-2B 7(Sr)-5B

17.42 (4) 23.60 (6) 17.54 (1)

53.19 (2) – 55.06 (5)

– – –

– 47.15 (3) 20.09 (6)

29.39 (6) 29.25 (9) 7.31 (9)100 100 100

3.38 4.615 4.4790.058 0.062 0.0640.045 0.046 0.048

1078 A. Goel et al. / Journal of Non-Crystalline Solids 356 (2010) 1070–1080

tively. On the steel surface, formation of whisker like oxide prod-ucts were formed for all the investigated diffusion couples, whichwere rich in Cr- and Mn-oxides indicating towards the possibleformation of manganese–chromium spinel. The oxygen partialpressure required to oxidize Cr has been reported to be 6.91 �10�28 atm [23] which is much lower in comparison to that requiredfor oxidation of Fe and led to the oxidation of chromium. The bond-ing/wetting behavior of silicate glasses to metals is strongly depen-dent on the nature of the surface of the metal. Metallic surfaces

Mn

0

10

20

30

40

50

0 4Posit

Rel

ativ

e C

once

ntra

tion

(%)

Crofer22 APU

(a)

(b)

Fig. 6. (a) Microstructure (SEM) and EDS element mapping of Ba, Mn, Cr and Fe at interfacMn at the interface between glass 7-2B and Crofer22 APU developed after heat treatment‘unoxidized’ refers to the usage of interconnect material in as received form. The alloy w

often result in poor joining due to poor wetting, while an oxidelayer o the metal surface tends to promote wetting which leadsto improved chemical bonding. An Al-rich oxide layer was ob-served beneath this layer for Sanergy/GC-7A diffusion couple(Fig. 7(a)). Similar Al-rich oxide layer was also observed by Hann-appel et al. [24] at an interface between BCAS GC sealant and chro-mia scale forming ferritic steel under wet fuel gas conditions. Theconcentration of impurities like Al and Si in metallic interconnectplays a significant role in the performance of SOFC stacks. It has

Cr

Ba

8 12ion (μm)

7-2B

e between glass 7-2B and Crofer22 APU (b) EDS line profile for diffusion of Cr, Ba andat 850 �C for 1 h and 800 �C for 300 h in humidified reducing atmosphere. (The termas not pre-oxidized.)

A. Goel et al. / Journal of Non-Crystalline Solids 356 (2010) 1070–1080 1079

been reported that higher amount of Si in metallic interconnectenhances the rate of corrosion attack due to internal oxidation,leading to the degradation of SOFC stack. However, no such degra-dation of SOFC stack has been reported due to internal oxidation ofAl [24]. In the present investigation, no such Al-rich layer could beseen in any of the other diffusion couples else than Sanergy HT/GC-7A. All the investigated GC sealants showed highly stable interfacewith Sanergy HT and absence of any severe reaction products un-der humidified reducing atmosphere. No severe effect of increasingB2O3 content in the glass compositions could be observed on thereaction mechanism at the glass–steel interface in reducing condi-tions as was observed in air atmosphere [5]. Although, the degra-

Fig. 7. SEM micrographs of the interface between (a) glass 7A and Sanergy HT, (b)glass 7-2B and Sanergy HT, (c) glass 7-5B and Sanergy HT developed after heattreatment at 850 �C for 1 h and 800 �C for 300 h in humidified reducing atmosphere.(The term ‘unoxidized’ refers to the usage of interconnect material in as receivedform. The alloy was not pre-oxidized.)

dation of B2O3-containing seals under humidified reducingconditions due to the formation of volatile species has been welldocumented in literature [25], no such effects were observed inthe present study owing to the very low concentration of B2O3 inthe glasses. Further, since no chromate formation was observedin any of the investigated diffusion couples, therefore, negligiblechromium depletion can be expected at the three phase boundary(metal–GC–gas). As a consequence, no swelling of the metal withconsequent bulging was observed. No significant differences wereobserved on the chemical interaction between sealant and steelafter replacement of BaO by SrO in the GC sealant (not shown) asa layer of about 1 lm thickness, rich in Cr- and Mn-oxides was ob-served at interface between GC sealant and metallic interconnect.

Similar results (not shown) were also observed for the inter-faces between pre-oxidized Crofer22 APU and pre-oxidized Saner-gy HT interconnect and GC sealants. Also, no significant differencewas observed in the results obtained for the chemical interactionbetween the GC sealants and Sanergy HT coated with 200 nm thickCo.

3.5. Leak rate measurements

The leak rate experiments were performed on glass composi-tions 7A and 7-2B. Glass compositions 7B and 7(Sr)-2B were ex-cluded as they exhibited low CTE after long term heat treatments[4,5]. Fig. 8 shows the OCV (E) vs. time under different pumpingcurrents. P(O2) in the chamber can be calculated using the equation

PðO2Þ ¼ P0 exp4FERT

� �: ð12Þ

0 20 40 60 80 100 120 140 160-0.05

-0.04

-0.03

-0.02

-0.01

0.00

7-2B, T=750oC 100μA 200μA 300μA 400μA 500μA

OC

V (

V)

t (s)

Fig. 8. The relationship of oxygen concentration voltage (OCV) of the sensor andtime under different pump current for composition 7-2B.

0 25 50 75 100 125 150 1750.0

5.0x103

1.0x104

1.5x104

2.0x104

2.5x104

7-2B, T=750oC 100μA 200μA 300μA 400μA 500μA

P(O

2)(P

a)

t (s)

ig. 9. Oxygen partial pressure vs. time plot under different pumping current foromposition 7-2B at 750 �C.

Fc

Table 5Leak rate values for GCs 7A and 7-2B at different temperatures.

Sample Temperature(�C) qn (mol s�1) V (lL) DP(atm)

QV

(sccm cm�1)

7A 600 9.3 � 10�15 12 0.2 7.8 � 10�5

650 8.6 � 10�15 13 7.2 � 10�5

700 8.8 � 10�15 14 7.4 � 10�5

750 8.8 � 10�15 15 7.3 � 10�5

800 1.04 � 10�14 14 8.7 � 10�5

7-2B 600 2.6 � 10�15 14 0.2 2.2 � 10�5

650 2.7 � 10�15 14 2.2 � 10�5

700 2.6 � 10�15 14 2.2 � 10�5

750 3.0 � 10�15 15 2.5 � 10�5

800 2.9 � 10�15 16 2.4 � 10�5

1080 A. Goel et al. / Journal of Non-Crystalline Solids 356 (2010) 1070–1080

Then the plot of P(O2) vs. time is obtained as shown in Fig. 9. Itshows linear relation under low P(O2) and high pumping currentconditions, so we can use the data to calculate qn and V accordingto Eqs. (7) and (8). The calculated results for compositions 7A and7-2B are listed in Table 5. Both compositions showed an acceptableleak rate [26] in order to qualify them as compliant sealants forSOFCs.

4. Conclusions

The structure, sintering behavior and crystallization kinetics ofLa2O3-containing diopside based GC sealants has been studiedand the interfacial adhesion along with chemical stability betweenGC sealants and metallic interconnects has been investigated un-der humidified reducing conditions.

29Si MAS-NMR reveals that Si atoms would be likely centered onQ2 units. No significant change on the structure of silicate networkwas observed with the addition of B2O3 in the glasses. 11B MAS-NMR showed that boron exists predominantly in BO3 units whilean increase in BO4 units was found with an increase in B2O3 con-centration in the glasses. Al was found to be dominantly coordi-nated in AlO4 units.

Sintering preceded crystallization in all the five GC sealants (7A,7B, 7-2B, 7-5B and 7(Sr)-2B), thus, rendering them suitable for fur-ther experimentation as sealants in SOFC. The addition of La2O3

improved the sintering behavior of the glasses while decreasingthe activation energy of crystallization. An increase in B2O3 de-creased the half ball, flow temperature and activation energy ofcrystallization of all the glasses. The viscosity of all the GC sealantsat SOFC operation temperature is appropriate for sealant applica-tions. The addition of Cr2O3 in the glasses decreased the activationenergy of crystallization while it increased with replacement ofBaO by SrO.

No detrimental interfacial products were observed at the inter-face between interconnects/GC sealant after a continuous exposurefor 300 h in humidified reducing atmosphere at 800 �C. Cr- andMn-rich oxide layers were formed at the steel side of the interfacein all the GCs, indicating towards the presence of thermodynami-cally stable chromium–manganese spinel. The GCs 7A and 7-2Bshow acceptable leak rate and therefore, are promising candidatesfor sealing applications in SOFC. Thus, further experimentation onthese GCs for sealing applications has to be continued.

Acknowledgements

This study was financially supported by University of Aveiro,CICECO, FCT, Portugal (SFRH/BPD/65901/2009). Part of this workwas also funded by the research project CICYT (MAT2006-04375).

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