stable glass-ceramic sealants for solid oxide fuel cells: influence of bi2o3 doping

i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n en e r g y 3 5 ( 2 0 1 0 ) 6 9 1 1e6 9 2 3

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Stable glass-ceramic sealants for solid oxide fuel cells:Influence of Bi2O3 doping

Ashutosh Goel a, Maria J. Pascual b, Jose M.F. Ferreira a,*aDepartment of Ceramics and Glass Engineering, University of Aveiro, CICECO, 3810-193 Aveiro, Portugalb Instituto de Ceramica y Vidrio (CSIC), Kelsen 5, Campus de Cantoblanco, 28049 Madrid, Spain

a r t i c l e i n f o

Article history:

Received 12 March 2010

Received in revised form

17 April 2010

Accepted 18 April 2010

Available online 21 May 2010

Keywords:

Solid oxide fuel cell

Glass-ceramic sealant

Diopside

Sintering

Interconnect

X-ray diffraction

* Corresponding author. Tel.: þ351 234 37024E-mail address: [email protected] (J.M.F. Ferreira

0360-3199/$ e see front matter ª 2010 Profedoi:10.1016/j.ijhydene.2010.04.106

a b s t r a c t

Diopside (CaMgSi2O6) based glass-ceramics in the system SrOeCaOeMgOe

Al2O3eB2O3eLa2O3eBi2O3eSiO2 have been synthesized for sealing applications in solid

oxide fuel cells (SOFC). The parent glass composition in the primary crystallization field of

diopside has been doped with different amounts of Bi2O3 (1, 3, 5 wt.%). The sintering

behavior by hot-stage microscopy (HSM) reveals that all the investigated glass composi-

tions exhibit a two-stage shrinkage behavior. The crystallization kinetics of the glasses has

been studied by differential thermal analysis (DTA) while X-ray diffraction adjoined with

Rietveld-R.I.R. analysis have been employed to quantify the amount of crystalline and

amorphous phases in the glass-ceramics. Diopside and augite crystallized as the primary

crystalline phases in all the glass-ceramics. The coefficient of thermal expansion (CTE) of

the investigated glass-ceramics varied between (9.06e10.14) � 10�6 K�1 after heat treat-

ment at SOFC operating temperature for a duration varying between 1 h and 200 h. Further,

low electrical conductivity, good joining behavior and negligible reactivity with metallic

interconnects (Crofer22 APU and Sanergy HT) in air indicate that the investigated glass-

ceramics are suitable candidates for further experimentation as sealants in SOFC.

ª 2010 Professor T. Nejat Veziroglu. Published by Elsevier Ltd. All rights reserved.

1. Introduction In planar design of SOFC, which involves stacking of tens

Glass-ceramics (GCs) combine the generally superior proper-

ties of crystalline ceramicswith the ease of processing of glass.

Major attributes of GCs include more refractory behavior and

superior mechanical properties, relative to glasses as well as

ceramics. Undoubtedly, one of themajor qualities, however, is

an ability to tailor their thermal expansion characteristics.

This makes GCs ideal candidates where compatible thermal

expansions are necessary. Most recently, there has been

a dramatic revival of interest in both glass- and GC- to metal

seals [1], particularly, for newapplications including SOFC [2,3]

and high temperature sensors [4,5].

2; fax: þ351 234 370204.).ssor T. Nejat Veziroglu. P

of repeating unit cells (anode/electrolyte/cathode) separated

by metallic interconnect plates, seal is required to prevent

fuel leakage and air mixing at high temperature

(800e1000 �C) along with to seal the electrolyte against the

metallic body of the device, in order to create a hermetic

rugged and stable stack. Any leakage of fuel into the air (or

air into the fuel) will lead to direct combustion of fuel and

may cause local overheating (hot spots) and sometimes may

burst. Therefore, the seals must be stable in a wide range of

oxygen partial pressure (air and fuel) and be chemically

compatible with other fuel cell components, while mini-

mizing thermal stresses during high temperature operation

ublished by Elsevier Ltd. All rights reserved.

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i n t e rn a t i o n a l j o u r n a l o f h y d r o g e n en e r g y 3 5 ( 2 0 1 0 ) 6 9 1 1e6 9 2 36912

which creates a major challenge in the development of

planar SOFCs [2,3].

As mentioned above, glasses and GCs are ideal candidates

for the job of sealing in SOFCs due to the flexible and

compliant nature of glass at temperatures above glass tran-

sition, which leads to decrease in mechanical stresses caused

by the difference in CTE between the sealing material and

SOFC component(s). Moreover, controlled crystallization of

glass seal leads to an increase in the mechanical strength and

electrical resistivity of the GC, while tailoring the CTE of the

final product with respect to the crystalline phases formed.

Majority of the glass/GC sealants developed so far are

either BaO-based [6] or Na2O-based [7,8] aluminosilicates.

However, due to the stringent requirements most sealants are

not practicable because of the drawbacks concerning either

thermal expansion mismatch or due to reactions with SOFC

components [9,10]. Further, significant content of BaO may

also promote interaction with water vapor, leading to slow

sealant degradation under SOFC operating conditions. A

remedy to these problems lies in the development of BaO- and

Na2O-free GC sealant which exhibits good CTE matching and

low/negligible reactivity with other SOFC components, in

particular with metallic interconnect.

An attempt in this direction has been made by various

research groups. Ley et al. [11] studied the glass and GC system

of SrOeAl2O3eLa2O3eSiO2eB2O3. The CTE values of the as-

made materials were in the range of (8e13) � 10�6 K�1, while

the long term stability was not reported. Recently, Brochu

et al. [12] compared the performance of the BaO- and SrO-

based borate glass-composites for sealing materials in SOFCs

and reported the formation of low CTE crystalline phase,

BaZrO3, on interaction with 8YSZ (ZrO2 stabilized by 8 mol%

Y2O3), for BaO-containing glass-composites. However, in case

of SrO-based glass-composites, formation of strontium zirc-

onates was observed, which has CTE similar to 8YSZ. Maha-

patra et al. [13,14] studied the structure and thermophysical

properties and devitrification behavior of the glasses in the

system (25�X )SrOe20La2O3e(7 þ X )Al2O3e40B2O3e8SiO2

(mol.%) (X ¼ 0e10). Similarly, Kumar et al. [15] studied the

influence of substituting La2O3, Y2O3 andAl2O3 on thermal and

physical properties of a glass with composition (mol.%)

30SrOe40SiO2e20B2O3e10A2O3 (A: La, Y, Al) and studied their

chemical interactionwith bismuth vanadate based electrolyte

material. However, high boron seals proposed in earlier

studies [11e15] are apt to eventually corrode under humidified

hydrogen environments (common in fuel cell operation) over

Table 1 e Batch compositions of the glasses.

Glass MgO CaO SrO SiO2

9 wt.% 14.53 15.72 12.45 45.72

mol.% 22.47 17.48 7.49 47.44

9-Bi1 wt.% 14.38 15.56 12.32 45.25

mol.% 22.43 17.45 7.48 47.36

9-Bi3 wt.% 14.08 15.24 12.07 44.31

mol.% 22.36 17.39 7.45 47.20

9-Bi5 wt.% 13.78 14.91 11.81 43.37

mol.% 22.28 17.33 7.43 47.04

time. Glasses with B2O3 as the only glass former have shown

up to 20% weight loss in the humidified H2 environment and

extensive interactions with cell component materials both in

air and wet fuel gas [16].

Therefore, in the light of abovementioned perspective, a SrO-

based aluminosilicate GC composition has been formulated in

the primary crystallization field of diopside (CaMgSi2O6) via

substitution scheme 0.2Ca2þ þ 0.1Mg2þ 4 0.3Sr2þ and 0.1

(Ca2þ þ Si4þ)4 0.1(La3þ þ Al3þ) in pure CaMgSi2O6 system, thus

resulting in a theoretical composition Sr0.3Ca0.7Mg0.9Al0.1La0.1-Si1.9O6. Further, the influence of Bi2O3 addition (1, 3 and 5 wt. %)

on the sintering and crystallization behavior, flow properties of

glasses along with CTE, electrical properties and chemical

interactionof resultantGCswithmetallic interconnectshasbeen

investigated. The deliberate addition of Bi2O3 has beenmadedue

to its low melting point (817 �C) which might be helpful in

tailoring the flow properties of sealants. Also, Bi2O3 is a major

component of bismuth vanadate based electrolyte materials for

SOFC [15]. However, the amount of Bi2O3 has been kept �5 wt.%

(<1 mol%) because if present in higher concentration it might

exhibit reducing behavior (Bi3þ / Bi0) in hydrogen rich envi-

ronment on the anode side of SOFC [17]. It is noteworthy that

even though minor amounts of PbO in GC sealant leads to rapid

andmassive internal oxidation and iron oxide formation on the

metallic interconnect surface at theair side of SOFC stack [18]; no

such results pertaining to Bi2O3 addition have been reported to

the best of our knowledge. Furthermore, 1 wt.% NiO and 2 wt.%

B2O3 have been added to all the investigated glass compositions

in order to improve adhesion behavior of GCs to metal and

decrease the viscosity and glass transition temperature (Tg),

respectively. Table 1 presents the compositions of all the inves-

tigated glasses.

2. Experimental

2.1. Synthesis of glasses

Homogeneous mixtures of batches (w100 g) in accordance

with glass compositions presented in Table 1 were prepared

by ball milling of powders of SiO2 (purity >99.5%), CaCO3

(>99.5%), Al2O3 (Sigma Aldrich, �98%), H3BO3 (Merck, 99.8%),

MgCO3 (BDH chemicals, UK, >99%), SrCO3 (Sigma Aldrich,

99þ%), La2O3 (Sigma Aldrich, 99.9%), Bi2O3 (Sigma Aldrich,

99.9%) and NiO (Sigma Aldrich, 99%) and calcination at 900 �Cfor 1 h. The glass batch was melted in Pt crucibles at 1550 �C

Al2O3 La2O3 B2O3 Bi2O3 NiO

2.04 6.53 2.00 e 1.00

1.25 1.25 1.79 e 0.83

2.02 6.46 2.00 1.00 1.00

1.25 1.25 1.81 0.13 0.84

1.98 6.32 2.00 3.00 1.00

1.24 1.24 1.84 0.41 0.86

1.94 6.19 2.00 5.00 1.00

1.24 1.24 1.87 0.70 0.87

3.16

3.20

3.24

mcg

(yti

3-)

20.2

20.3

20.4

mc(e

mu3

lom

1-)Density

Molar volume

i n t e r n a t i o n a l j o u r n a l o f h y d r o g e n en e r g y 3 5 ( 2 0 1 0 ) 6 9 1 1e6 9 2 3 6913

for 1 h, in air. Glasses in bulk form were produced by pouring

the melts on preheated bronze moulds followed by annealing

at 550 �C for 1 h while glass frits were obtained by

quenching of glass melts in cold water. The frits were dried

and then milled in a high-speed agate mill resulting in fine

glass powders with mean particle sizes of 10e20 mm (deter-

mined by light scattering technique; Coulter LS 230, Beckman

Coulter, Fullerton CA; Fraunhofer optical model).

3.07

3.12

0 1 2 3 4 5

Bi2O3 (wt.%)

sneD

20.0

20.1

lovralo

M

Fig. 1 e Influence of Bi2O3 on density and molar volume of

the glasses.

2.2. Density and dilatometry

Archimedes’ method (by immersion in diethyl phthalate) was

employed to measure the apparent density of the bulk

annealed glasses. The obtained density values were further

employed along with composition of glasses to calculate their

molar volume and excess volume.

The glass transition temperature (Tg) and softening

point (Ts) of glasses along with CTE of glasses and GCs,

respectively were obtained from dilatometry measure-

ments which were carried out on prismatic samples with

a cross section of 4 mm � 5 mm (Bahr Thermo Analyze DIL

801 L, Hullhorst, Germany; heating rate 5 K min�1). The

dilatometry measurements were made on a minimum of 3

samples from each composition and the standard devia-

tion for the reported values of CTE are in the range

�0.1 � 10�6 K�1.

Table 2 e Thermal parameters of glasses obtained fromdilatometry, DTA and HSM at b [ 5 K minL1.

9 9-Bi1 9-Bi3 9-Bi5

Tg (�5) (�C) 555 535 520 520

Ts (�3) (�C) 710 695 695 705

(CTE � 0.05) � 106 K�1 (200e500 �C)

8.28 8.38 8.57 8.33

TFS (�5) (�C) 760 758 755 755

TMS1 (�5) (�C) 811 807 807 805

Tc (�2) (�C) 875 871 880 877

TD (�5) (�C) 861 850 850 850

Tp (�2) (�C) 904 898 905 904

THB (�5) (�C) 1217 1211 1208 1203

TF (�5) (�C) 1239 1237 1235 1230

Sc ¼ Tc-TMS 64 64 73 72

A/A0 0.64 0.65 0.64 0.67

n � 0.005 1.93 1.86 1.96 2.05

Ec (kJ mol�1) 483 (0.9992)a 513 (0.9996) 495 (0.9994) 473 (1)

a The values in parenthesis correspond to square regression

coefficient (r2) obtained from the slope of 4 points.

2.3. Sintering and crystallization kinetics by HSM andDTA

The sintering behavior of the glass powders was investigated

by using a hot-stage microscope (HSM). A side-view HSM EM

201 equipped with an image analysis system and electrical

furnace 1750/15 Leica was used. The microscope projects the

image of the sample through a quartz window and onto the

recording device. The computerized image analysis system

automatically records and analyzes the geometry changes of

the sample during heating. The measurements were con-

ducted in air with a heating rate of 5 K min�1. The cylindrical

shaped samples with height and diameter of w3 mm were

prepared by cold-pressing the glass powders and were placed

on an alumina support. The temperature was measured with

a Pt/Rh (6/30) thermocouple contacted under the alumina

support. The temperatures corresponding to the character-

istic viscosity points were obtained from the photographs

taken during the hot-stage microscopy experiment following

Ref. [19].

The differential thermal analysis (DTA-TG, Setaram Lab-

sys, Setaram Instrumentation, Caluire, France) of glass

powders was carried out in air from room temperature to

1000 �C with different heating rates (b) of 5, 10, 20 and

30 K min�1. The glass powders (mean particle size:

10e20 mm) weighing 50 mg were contained in an alumina

crucible and the reference material was a-alumina powder.

The crystallization kinetics of the glasses was studied using

the formal theory of transformation kinetics as developed by

Johnson and Mehl [20] and Avrami [21e23], for non-

isothermal process that has already been obtained in our

previous work [24]:

ln

T2p

b

!¼ Ec

RTp� lnq ¼ 0 (1)

which is the equation of a straight line, whose slope and

intercept give the activation energy, Ec, and the pre-expo-

nential factor, q ¼ Q1/nK0, respectively and the maximum

crystallization rate by the relationship:

dxdt

jp¼ 0:37bEcn

�RT2

p

��1

(2)

which makes it possible to obtain, for each heating rate,

a value of the kinetic exponent, n. In Eq. (2), c corresponds to

the crystallization fraction and dcdtjp corresponds to the crys-

tallization rate, which may be calculated by the ratio between


the ordinates of the DTA curve and the total area of the

crystallization curve.

2.4. Isothermal and non-isothermal heat treatments ofglass powder compacts

The circular disc shaped pellets with Ø 20 mm and thickness

w3 mm were prepared from glass powders by uniaxial

pressing (80 MPa) and were sintered under non-isothermal

conditions for 1h 800and850 �C, respectively, at a slowheating

rate of 2 K min�1. Further, in order to study the crystalline

phase assemblage in GC sealants after a prolonged usage in

SOFC stack, the green glass powder compacts were initially

sintered at 850 �C for 1 h at a heating rate of 2 Kmin�1 and then

the temperature was brought down to SOFC operating

temperature i.e. 800 �C. Finally, the glass powder compacts

were heat treated at SOFC operating temperature for 200 h.

2.5. Joining behavior and chemical interactions betweeninterconnect-seal-interconnect diffusion couples

Two differentmetallic interconnectmaterials, namely, Crofer22

APU (Thyssen Krupp, VDM,Werdohl, Germany) and Sanergy HT

(SandvikAB, Sandviken, Sweden)were employed for joining and

interaction experiments with the glasses. The chemical

composition of the two interconnect materials has been pre-

sentedelsewhere [25]. The joinedCrofer/glass/CroferorSanergy/

glass/Sanergy samples were obtained by deposition of glass

0.5

0.7

0.9

1.1

600 700 800 900 1000

Temperature (oC)

A/A

0

-40

-30

-20

-10

Δ(

Tμ

)V

Exo

β = 5 K min-1

9a b

c d

0.5

0.7

0.9

1.1

600 700 800 900 1000

Temperature (oC)

A/A

0

-45

-35

-25

-15

Δ(

Tμ

)V

Exo

β = 5 K min-1

9-Bi3

Fig. 2 e Comparison of DTA and HSM curves on the same tempe

9-Bi5.

powders (mixed with 5 vol.% solution of polyvinyl alcohol (PVA)

prepared by dissolution of PVA in warm water) on metallic

interconnects by slurry coating. Heat treatments were per-

formed in air without applying any dead load. The diffusion

coupleswere heated to 850 �Cwith a relatively slowheating rate

(2 K min�1) and kept at that temperature for 1 h. Finally,

the temperature was brought down to SOFC operating temper-

ature (i.e. 800 �C) andmaintained at this temperature for 200 h.

2.6. Crystalline phase analysis of glass-ceramics

The amorphous nature of glasses and qualitative along with

quantitative analysis of crystalline phases in the GCs (crushed

to particle size <45 mm) was made by XRD analysis using

a conventional Bragg-Brentano diffractometer (Philips PW

3710, Eindhoven, The Netherlands) with Ni-filtered Cu-Ka

radiation. The quantitative phase analysis of GCs was made

by combined Rietveld-R.I.R (reference intensity ratio) method.

A 10 wt.% of corundum (NIST SRM 676a) was added to all the

GC samples as an internal standard. The mixtures, ground in

an agate mortar, were side loaded in aluminum flat holder in

order to minimize the preferred orientation problems. Data

were recorded in 2q range ¼ 5e140� (step size 0.02� and 25 s of

counting time for each step). The phase fractions extracted by

Rietveld-R.I.R refinements, using GSAS software and EXPGUI

as graphical interface, were rescaled on the basis of the

absolute weight of corundum originally added to their

mixtures as an internal standard, and therefore, internally

HSM

DTA

Tc

Tp

TFS

TMS1TMS2

0.5

0.7

0.9

1.1

600 700 800 900 1000

Temperature (oC)

A/A

0

-45

-35

-25

-15

Δ(

Tμ

)V

Exo

β = 5 K min-1

9-Bi1

0.5

0.7

0.9

1.1

600 700 800 900 1000

Temperature (oC)

A/A

0

-50

-40

-30

-20

-10

Δ(

Tμ

)V

Exo

β = 5 K min-1

9-Bi5

rature scale for compositions (a) 9, (b) 9-Bi1, (c) 9-Bi3 and (d)

Fig. 3 e HSM images of glasses on alumina substrates at various stages of heating cycle.


renormalized. The background was successfully fitted with

a Chebyshev function with a variable number of coefficients

depending on its complexity. The peak profiles were modeled

using a pseudo-Voigt function with one Gaussian and one

Lorentzian coefficient. Lattice constants, phase fractions, and

coefficients corresponding to sample displacement and

asymmetry were also refined.

Fig. 4 e SEM images of glass powder compacts from compo

2.7. SEM-EDS analysis

Microstructural observations were done on polished GC

samples (chemically etched by immersion in 2 vol.% HF

solution for 2 min) and interconnect/GC/interconnect diffu-

sion couples (un-etched) by scanning electron microscopy

(SEM; SU-70, Hitachi) with energy dispersive spectroscopy

sition (a) 9 and (b) 9-Bi3 heat treated at 800 �C for 1 h.


(EDS; Bruker Quantax, Germany) to study the distribution of

elements in the crystals and also to study the distribution of

elements along the GC-interconnect diffusion couples.

2.8. Electrical conductivity of glass-ceramics

The total conductivity was studied by the AC impedance

spectroscopy (Potentistat/Galvanostat/ZRA, Reference 600,

10 Hze1 MHz; Gamry Instruments, Warminster, PA, USA)

using dense disk-shaped samples (sintered at 850 �C for 1 h)

with porous Pt electrodes and Pt current collectors, in atmo-

spheric air. In the course of impedance measurements, the

Fig. 5 e (a) X-ray diffractograms of glass powder compacts sinte

phase reflections corresponding to corundum have not been m

intensity axes 23,000 cps. (b) Observed (crosses), calculated (con

refinement of the GC 9-Bi1 heat treated at 850 �C for 1 h in air.

corundum, Sr-diopside and augite (from top to bottom).

magnitude of AC voltage was fixed at 1.00 V. The electrical

conductivity experiments were performed on aminimum of 3

samples for each composition in order to confirm the accuracy

of the measurements.

3. Results and discussion

3.1. Density and dilatometry

The density of glasses increased with increase in Bi2O3

content (Fig. 1) due to its highest density (8.9 g cm�3) in

red at 850 �C for 1 h (Di: Diopside; Sr-Di: Sr-diopside; the

arked). The spectra have not been normalized. Full scale

tinuous line), and difference curve from the Rietveld

Markers representing the phase reflections correspond to

Fig. 6 eMicrostructure (revealed via SEM imaging after chemical etching of polished surfaces with 2 vol.% HF solution) of the

GCs (a) 9 and (b) 9-Bi3 after heat treatment at 850 �C for 1 h, respectively.


comparison to other constituents of glasses. The molar

volume of glasses was observed to follow a similar trend

as it increased with increasing Bi2O3 concentration in

glasses (Fig. 1). These results are in good agreement with the

results of Ahlawat et al. [26] and may be explained on the

basis of the fact that Bi2O3 is an unconventional glass

network former and increasing Bi2O3/SiO2 ratio in the

glasses has been reported to increase the covalent character

of Bi3þ atoms and decrease the covalency of Si4þ [27].

However, in the present case, considering the substantial

difference between SiO2 and Bi2O3 concentration in the

glasses, the influence of slight decrease in covalency of Si4þ

might be neglected.

The dilatometric glass transition temperature (Tg)

decreased with increasing Bi2O3 content in the glasses

(Table 2) while no significant impact of Bi2O3 concentration

could be observed on softening temperature (Ts) of glasses.

The CTE values of glasses (200e500 �C) increased slightly with

increasing Bi2O3 content (Table 2), however, no general trend

could be observed in variation of these values. A detailed

structural and thermal investigation on these glasses with

wide variation in Bi2O3/SiO2 concentration will be helpful in

gaining a better insight about the structure-property rela-

tionships of these glasses.

Table 3 e Results of Rietveld-R.I.R quantitative analysis.

850 �C, 1 h

9 9-Bi1 9-Bi3

Diopside (01-078-1390) 32.85 (8) e e

Augite (01-078-1392) 61.59 (1) 9.64 (8) 11.18 (2)

Sr-diopside (01-080-0386) e 89.46 (1) 86.06 (1)

Calcium lanthanum silicate

oxide (04-008-8013)

e e e

Glass 5.56 (9) 0.90 (9) 2.76 (3)

Total 100 100 100

c2 2.18 1.97 1.91

Rwp 0.069 0.068 0.069

Rp 0.050 0.050 0.051

3.2. Sintering behavior and crystallization kinetics

Sealing is usually applied on the surface (ceramic or metallic)

to be sealed using powder glass mixed with a binder. The GC

formation involves the sintering of glass powders, followed by

crystallization at a higher temperature. In order to obtain

a good sealing, the sintering stage should precede crystalli-

zation as dense and low porosity materials are desired for

obtaining a gas-tight GC seal. Further, crystallization is

needed to increase the seal viscosity, CTE and improve the

chemical and mechanical durability of the sealant which has

to maintain the bulk stability and not flow during operation at

higher temperature. In order to assess the sintering and

devitrification behavior of glass system, a comparison

between DTA and HSM thermographs obtained under same

heating conditions can reveal a great deal of information in

this regard.

In the present study, a comparison between DTA and HSM

thermographs of all the investigated glasses obtained at

a heating rate of 5 K min�1 in the temperature range of

25e1000 �C reveals that sintering precedes crystallization in

all the glasses (Fig. 2). Fig. 3 presents the photomicrographs of

all the four glass compositions depicting the variation of

sample dimensions with increase in temperature. The

800 �C, 200 h

9-Bi5 9 9-Bi1 9-Bi3 9-Bi5

e 40.78 (2) e e e

6.23 (3) 31.08 (2) 72.00 (1) 18.15 (1) 23.80 (3)

89.11 (5) e 12.31 (1) 47.87 (1) 36.46 (1)

e e e 3.27 (1) 5.29 (2)

4.66 (8) 28.14 (4) 15.69 (2) 30.71 (3) 34.45 (6)

100 100 100 100 100

3.51 2.93 3.065 2.082 1.88

0.088 0.101 0.114 0.094 0.089

0.058 0.083 0.084 0.073 0.069

Table 5e Electrical properties of the GCs sintered at 850 �Cfor 1 h.

Composition s � 105 (S m�1) EA (kJ mol�1)

775 �C 800 �C

9 3.0 4.4 142 � 3

9-Bi1 2.3 3.4 148 � 4

9-Bi3 2.5 3.9 154 � 4

9-Bi5 2.4 3.8 157 � 2

-4

-2

T)

9 9-Bi19-Bi3 9-Bi5


sintering initiated (TFS: temperature of first shrinkage; log

h ¼ 9.1 � 0.1, h is viscosity; units: dPa s) at w755e760 �C in all

the compositions (Table 2). A two-stage shrinkage behavior

was observed for the investigated glasses where the temper-

ature for maximum shrinkage (TMS1; log h ¼ 7.8 � 0.1) was

achieved well before the onset of crystallization (Tc) (Fig. 2;

Table 2), thus resulting in a well sintered, dense but amor-

phous glass powder compact as represented by the SEM

images of glass compositions 9 (Fig. 4a) and 9-Bi3 (Fig. 4b),

respectively obtained after sintering of their glass powders at

800 �C for 1 h. Further, as is evident from Fig. 2 as well as Fig. 3,

second stage of shrinkage progressed in parallel with the

onset of crystallization and finished before the appearance of

peak temperature of crystallization (Tp) as has been repre-

sented in Fig. 2b by TMS2. This second step of shrinkagemay be

attributed to the possible tendency towards glass-in-glass

phase separation in the investigated glasses as has also been

reported in our recent studies [25,28]. Table 2 lists the values of

sinterability parameter (Sc) [29], where Sc ¼ Tc�TMS. The

parameter Sc is the measure of ability of sintering versus

crystallization: the greater this difference, the more inde-

pendent are the kinetics of both processes. The Sc values

greater than 25 �C, as obtained in the present study (60e75 �C),are related with high final densities, which indicate good

sintering/crystallization behavior. In general, addition of Bi2O3

improved the sintering ability of the investigated glass

compositions with composition 9-Bi3 exhibiting the best sin-

tering ability and flow behavior (Fig. 3) among all the investi-

gated compositions.

The peak temperature of crystallization (Tp) decreased

slightly with an initial addition of Bi2O3 in the parent glass (9)

while it increasedwith further increase in Bi2O3 content in the

glasses as is presented in Table 2 while a vice-versa trend was

observed for activation energy of crystallization (Ec) which

initially increased with 1 wt.% Bi2O3 in the parent glass while

further increase in Bi2O3 content led to a gradual decrease in

the values of Ec (Table 2). The Avrami parameter, n, for all the

investigated compositions vary in the range 1.85e2.0 which

implies towards intermediate (simultaneous occurrence of

both volume and surface nucleation and crystallization)

mechanism of crystallization in all the glasses. The values of

Ec for the investigated glasses are higher in comparison to

BaO-containing diopside based sealants investigated in our

previous studies [28,30].

Further, it is noteworthy that all the investigated compo-

sitions sinter completely before acquiring SOFC operation

temperature and the softening of glass powder compacts (TD;

log h ¼ 6.3 � 0.1; the temperature at which the rounding of

small protrusions or edges of the sample are observed) in all

Table 4e (CTE ± 0.05)3 106 KL1 (200 �Ce700 �C) of the GCsproduced at different conditions.

Composition 850 �C, 1 h 800 �C, 200 h

9 10.06 9.06

9-Bi1 9.88 9.53

9-Bi3 9.56 10.14

9-Bi5 9.83 9.62

the investigated samples occurs around SOFC operating

temperature (800e850 �C) (Table 2, Fig. 3). The half ball

temperature (THB; log h ¼ 4.1 � 0.1) and flow temperature (TF;

log h ¼ 3.4 � 0.1) of all the four compositions were obtained

from the HSM micrographs (Fig. 3), respectively and were

observed to decrease with increasing Bi2O3 content in the

glasses.

3.3. Heat treatment at 850 �C for 1 h: crystallizationbehavior and properties

In accordance with HSM and DTA results, well sintered, dense

and crystallized GCs were obtained after heating glass powder

compacts at 850 �C for 1 h as depicted in Fig. 5a. It should be

mentioned here that closed porosity was observed in all the

sintered glass powder compacts after heat treatment at 800 �Cfor 1 h (Fig. 4) which gradually decreased but did not

completely disappear after heat treatment at 850 �C as can be

seen in Fig. 6. It is noteworthy that appearance of closed

porosity in sintered GCs is a usual phenomenon [31] and does

not lead to enhancement of leak rate of final sealing material.

Table 3 presents the qualitative as well as quantitative anal-

ysis of the crystalline phases present in all the investigated

GCs as obtained from XRD analysis adjoined with Rietveld-R.I.

R technique. Augite (Ca(Mg0.70Al0.30)(Si1.70Al0.30)O6; ICDD card:

01-078-1392) crystallized as a dominant crystalline phase in

parent GC (composition 9) along with diopside (CaMgSi2O6;

ICDD card: 01-078-1390). However, addition of Bi2O3 to the

parent glass promoted the crystallization of Sr-containing

diopside (Ca0.8Sr0.2MgSi2O6; 01-080-0386) as the major

-8

-6

0.85 0.95 1.05 1.151000/T (K)

ln (

Fig. 7 e Plot for determination of activation energy of total

electrical conductivity in the investigated GCs.


crystalline phase in all the Bi2O3 containing GCs along with

augite (01-078-1392) as a minor crystalline phase as depicted

in Table 3. Fig. 5b shows the fit of a measured XRD pattern of

a sintered GC by using the GSAS-EXPGUI software. The

difference plot does not show any significant misfits. The

differences under the main peaks of Sr-diopside and augite

are caused by adjustment difficulties based on crystallinity of

phases. The quantitative analysis of crystalline phases reveal

that the GCs exhibit a high amount of crystallinity (�95 wt.%)

with formation of no detrimental crystalline phases after heat

treatment at 850 �C for 1 h. The microstructure of GCs (Fig. 6)

reveal densely packed crystals of uniform morphology in all

Fig. 8 e (a)X-ray diffractograms of glass powder compacts

(already sintered at 850 �C for 1 h) heat treated at 800 �C for

200 h (Di: Diopside; Sr-Di: Sr-diopside; the phase

reflections corresponding to corundum and calcium

lanthanum silicate oxide have not been marked). The

spectra have not been normalized. Full scale intensity axes

23,000 cps. (b) Microstructure (revealed via SEM imaging

after chemical etching of polished surfaces with 2 vol.% HF

solution) of the GC 9-Bi5 after heat treatment at 850 �C for

1 h and further at 800 �C for 200 h.

the investigated compositions. The addition of Bi2O3 did not

significantly affect the microstructure of sintered GCs

although it changed the crystalline phase assemblage in GCs.

The almost similar microstructure of the investigated GCs

may be due to the fact that all the three different crystalline

phases are derivatives of diopside (augite is Al-containing

diopside; while in Sr-diopside, Sr partially replaces Ca in the

structure of diopside without affecting the over crystal

symmetry) and belong to the family of clinopyroxenes.

The CTE of the investigated GCs varied in the range

(9.56e10.06) � 10�6 K�1. Although no specific trend could be

observed in the variation of CTE values with an increase in the

Bi2O3 content; CTE values decreased slightly with addition of

Bi2O3 in the parent GC. However, still the CTE of investigated

GCs is in good agreement with that of ceramic electrolyte,

8YSZ (w10 � 10�6 K�1) and metallic interconnect, Sanergy HT

(w11 � 10�6 K�1) [25] considering the fact that CTE differences

in seal and SOFC component can be accommodated until

1 � 10�6 K�1 [3].

The electrical conductivity of the GCs varies in the range

(2.3e3.0)� 10�5 S m�1 at 775 �C and increases slightly at 800 �C[(3.4e4.4) � 10�5 S m�1] as presented in Table 5. However, still

the total conductivity of GCs is in good agreement with other

sealants proposed in literature [18,32,33] and also in compar-

ison to the BaO-containing diopside based sealants investi-

gated in our recent study [34]. This level of conductivity

ensures an absence of short circuiting between the SOFC stack

components, especially in the intermediate temperature

range. It should be noted that at temperatures below

775e800 �C, the conductivity of all the studied GCs becomes

similar with in the limits of experimental uncertainty.

Further, since it has been already determined that the

conductivity mechanism in diopside based sealants in

predominantly ionic [34], therefore ion transference numbers

were not obtained in the present study. In general, the elec-

trical conductivity decreased with addition of Bi2O3 in the GC

compositions as listed in Table 5. The activation energy of

conductivity as calculated from Arrhenius equation,

increasedwith increase in Bi2O3 content in the GCs (Fig. 7) and

varied in the range 142e157 kJ mol�1 (Table 5).

Fig. 9 e SEM image of the polished interface between

Sanergy HT/GC 9/Sanergy HT after heat treatment at 850 �Cfor 1 h followed by 800 �C for 200 h.


3.4. Long term thermal stability and chemicalinteraction between sealant and metallic interconnect

The XRD data reveals that prolonged heat treatment of GCs

(already sintered at 850 �C for 1 h) at 800 �C for 200 h caused

a significant variation in their crystalline phase assemblage

Fig. 10 e SEM image and EDS element mapping of Sr, Cr, Fe, Ti

after heat treatment at 850 �C for 1 h followed by 800 �C for 200

(Fig. 8a). An increase in amorphous content in all the GCs

could be observed after prolonged heat treatments from the

differences in peak intensities of their x-ray diffractograms

obtained after heat treatment at 850 �C for 1 h (Fig. 5a) and at

800 �C for 200 h (Fig. 8a). The qualitative XRD results have been

confirmed by the results obtained from quantitative analysis

, Mo and Nb at the interface between GC 9 and Sanergy HT

h.


of crystalline/amorphous ratio in GCs as presented in Table 3.

The Rietveld-R.I.R results reveal that the amount of amor-

phous content in GCs increased considerably after prolonged

heat treatments and varied between 15 and 35 wt.% (Table 3).

It is noteworthy that the remaining glassy phase in GCs acts as

a major factor for deciding the flow behavior of the GC seal

and plays a crucial role in determining the reaction kinetics

during chemical interaction between sealant and SOFC

components. Also, such a high amount of glassy phase

present in the GCs is expected to exhibit self-healing behavior

during SOFC operation. The microstructure of GCs after pro-

longed heat treatment reveals the segregation of heavy metal

oxides (La2O3 and Bi2O3) in amorphous phase around the cli-

nopyroxene crystals as depicted by white color zones in

Fig. 8b. The segregation of La2O3 and Bi2O3 was not observed in

the case of GCs sintered at 850 �C for 1 h. It is quite reasonable

to expect slow diffusion kinetics for the heavy metal ions,

which tend to remain in the glassy phase. Further, the amount

of augite (01-078-1392) increased in all the GCs at the expense

of Sr-diopside (01-080-0386), especially in GC 9-Bi1, where

former crystallized as major crystalline phase. Similarly, in

the parent composition (GC 9), amount of diopside (01-078-

1390) and amorphous phase increased at the cost of augite,

thus, leading to the appearance of diopside as primary crys-

talline phase (Table 3). The increased concentration of Bi2O3

(�3 wt. %) led to the crystallization of calcium lanthanum

oxide silicate [CaLa4(SiO4)3O; ICDD card: 04-008-8013] in the

GCs as minor crystalline phase (Table 3). A slight decrease in

the CTE values due to an increase in the glassy phase was

observed for all the GCs after prolonged heat treatments at

SOFC operation temperature except GC 9-Bi3 as listed in Table

4. The CTE value of GC 9-Bi3 matches well with ceramic

Fig. 11 e SEM image and EDS element mapping of Sr, Cr and Bi a

treatment at 850 �C for 1 h followed by 800 �C for 200 h.

electrolyte (8YSZ) and metallic interconnect (Sanergy HT) of

SOFC. In general, the CTE values of all the GCs matches fairly

well with 8YSZ electrolyte for SOFC.

All the sealing GCs bonded well to metallic interconnects

(Sanergy HT and Crofer22 APU) and no gaps were observed

even at the edges of the joints. Fig. 9 shows the SEM image of

the interface between Sanergy HT/GC 9/Sanergy HT join after

heat treatment at 850 �C for 1 h followed by heat treatment at

800 �C for 200 h in air. As is evident fromFig. 9, the investigated

GCs were successful in making a strong metal-to-metal seal

without appearance of any detrimental reaction products at

the interface between metal and seal. It should be noted that

a number of voids that can be seen in the GC part of SEM

images obtained from metal-seal-metal diffusion couples

(Figs. 9e12) are due to the removal of some amorphous or

crystalline material from the GC samples during mechanical

grinding and polishing of interfaces. Fig. 10 presents the EDS

element mapping along the interface of GC 9 and Sanergy HT.

As is evident fromelementmapping, a rather smooth interface

was obtained between the investigated GC seals and metallic

interconnect Sanergy HT without the presence of iron-rich

oxide products. A very thin layer rich in Cr could be observed at

the interface between GC 9 and Sanergy HT (Fig. 10) indicating

the possible existence of Mn, Cr-rich spinel as revealed by EDS

elemental mapping. No significant differences could be

observed in the chemical interaction between GC sealants and

metallic interconnects due to addition of Bi2O3 as can be seen

in SEM image and elementmapping of interface betweenGC 9-

Bi3 andSanergyHT (Fig. 11). Also, similar resultswereobtained

for the Crofer22 APU/seal/Crofer22 APU diffusion couples.

However, we could not observe the existence of Cr or Mn-rich

zones at the interface between GC seal/Crofer22 APU as is

t the interface between GC 9-Bi3 and Sanergy HT after heat

Fig. 12 e SEM and EDS element mapping of Sr, Ti and Cr at the polished interface between GC9/Crofer22 APU depicting the

formation of titanium rich layer at the interface.


evident fromEDSelementmappingdepicted in Fig. 12. Instead,

the presence of Ti-rich zone near their interface was observed

(Fig. 12). The formation of titanium oxide layer may have

apositive effect as theouter layer containsnoCr, therefore, the

migration of Cr to cathode which poisons its effectiveness

would be expected to be greatly reduced if not eliminated

[34,35]. A similar observation was reported by Jablonski and

Alman [36] for a steel containing 22wt.%Cr and 1wt.%Tiwhen

surface treated by CeO2 while untreated steel samples did not

show formation of Ti-enriched protecting layers. According to

Jablonski and Alman [36], the formation of titanium oxides is

much more favorable from thermodynamic point of view, in

comparison to the other Cr-, Mn-rich oxidizing species.

Nevertheless, it will be interesting to investigate the chemical

interactions between these sealants and metallic intercon-

nects in humidified reducing conditions in order to analyze the

redox behavior of Bi3þ.

4. Conclusions

Glass-ceramic sealants free from BaO and Na2O have been

designed and investigated in the crystallization field of diop-

side (CaMgSi2O6). Further, the influence of Bi2O3 (1e5 wt. %)

addition on the flow properties, sintering and crystallization

behavior along with electrical conductivity and long term

thermal stability of sealants has been investigated. All the

glasses exhibit two-stage shrinkage behavior resulting in well

sintered glass-ceramics with diopside based crystalline pha-

ses. The amount of amorphous character in the GCs increases

considerably during prolonged heat treatments at SOFC

operating temperatures which can be beneficial to provide

self-healing ability to the sealant. Further, highly stable crys-

talline phase assemblage and matching of CTE with SOFC

components and low electrical conductivity are some more

attributes of the investigated glass-ceramic compositions. The

investigated glass-ceramic compositionswere highly effective

in performing metal-to-metal sealing with smooth interface

and negligible interfacial reactions, thus proving them to be

potential sealants for applications in SOFC. However, some

issues including redox stability of Bi3þ in humidified reducing

atmosphere, bond strength between sealant and SOFC

components and leak ratemeasurements of GC sealants, need

to be addressed. Therefore, further experimentation on these

GCs for sealing applications has to be continued.

Acknowledgements

Ashutosh Goel is thankful to FCT-Portugal for research grant

(SFRH/BPD/65901/2009). The support of CICECO is also

acknowledged.

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