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Journal of the European Ceramic Society 34 (2014) 1365–1373

Role of cement content on the properties of self-flowing Al2O3 refractorycastables

Cem Gogtas a,1, Hugo F. Lopez a,∗, Konstantin Sobolev b

a Department of Materials Science and Engineering, University of Wisconsin-Milwaukee, 3200 North Cramer Street, Milwaukee, WI 53211, United Statesb Department of Civil Engineering and Mechanics, University of Wisconsin-Milwaukee, 3200 North Cramer Street, Milwaukee, WI 53211, United States

Received 10 June 2013; received in revised form 8 October 2013; accepted 1 November 2013Available online 12 December 2013

bstract

n this work, Al2O3 self-flowing castables (SFCs) were produced based on various cement contents. The SFCs were sintered at 1273 K, 1573 Knd 1773 K and the exhibited properties were experimentally determined. Among the properties determined in this work are bulk density (BD),pparent porosity (AP), water absorption (WA), cold crushing strength (CCS), modulus of rupture (MOR) and fracture toughness (KIC). It is foundhat additions of 5% cement lead to SFCs with maximum MOR and KIC values after firing at 1773 K. Firing at 1573 K leads to a reduction in both,

OR and KIC. In SFC containing 3% cement, maximum KIC values of 3.53 MPa m1/2 were achieved after firing at 1573 K. In the low cement SFCs

1 wt%) after firing at 1773 K the exhibited KIC values were below those obtained in either the SFC-3 or SFC-5, but they were significantly high3.43 MPa m1/2).

2013 Elsevier Ltd. All rights reserved.

eywords: Self flow castables; Alumina refractories; Cement; Modulus of rupture; Fracture toughness

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. Introduction

Refractory castables are classified by The American Societyor Testing Materials (ASTM) according to their lime content asa) conventional castables (CCs) (CaO > 2.5%), (b) low cementastables (LCCs) (2.5% > CaO > 1.0%), (c) ultra low cementastables (ULCCs) (1% > CaO > 0.2%) and (d) zero cementastables (ZCCs) (CaO < 0.2%). CCs contain 15–30% calciumluminate cement which acts as a lubricant in a water suspensionnd as a binder after hardening for strength at room temperature.igh cement CCs are usually porous and open textured, due to

he relatively large amounts of water required. They also exhibit characteristic drop in strength at intermediate temperatures as

luggish sintering does not allow the development of a ceramicond after breaking down the hydraulic bond. The high lime con-ent of these castables favors the formation of low melting phases

∗ Corresponding author. Tel.: +1 414 229 6005; fax: +1 414 229 6958.E-mail addresses: cgogtas@uwm.edu (C. Gogtas), hlopez@uwm.edu

H.F. Lopez), sobolev@uwm.edu (K. Sobolev).1 Tel.: +1 414 229 6005; fax: +1 414 229 6958.

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955-2219/$ – see front matter © 2013 Elsevier Ltd. All rights reserved.ttp://dx.doi.org/10.1016/j.jeurceramsoc.2013.11.004

uch as anorthite and gehlenite at elevated temperatures. Thesehases are known to degrade the refractoriness and corrosionesistance of the CCs.1,2

Further improvements in the properties of CCs have beenchieved through the development of LCCs and ULLCs.1–3 Inhese castables, minimal water contents are needed to achievehe casting consistency. In turn, these castables posses improvedold and hot strength, reduced porosity, increased density andigh corrosion resistance. In addition, in LCCs the calciumluminate cement is replaced by a hydraulic binder such asydraulic alumina.1–5 Due to the reduced water content lowement additions in the castables can induce poor flowabilitynd vibration is commonly applied in order to promote flow. Inurn, this can result in serious space filling limitations in ordero avoid any defect formation.6–9 In the mid-1980s, self-flowingastables (SFCs) were developed (LCCs or ULLCs) with a mix-ng consistency which allows them to flow and degas withouthe application of vibration.7

In addition, in refractory materials one of the important

actors used to control flowability and porosity is the parti-le size distribution (PSD). Accordingly, the modified modelf Andreassen has been widely used to establish particle size

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istributions.7,10 In the Andreassen model, q is known as theistribution coefficient of the PSD and for q values below 0.25,elf flow properties are attained while values of around 0.30 areypical of vibrational castables.10

Alumina based castables have been widely investigated inhe literature as alumina posses high strength, high hardness,ood thermal shock properties and corrosion resistance.2,5–14 Inhe present work, only alumina grades that qualify for struc-ural engineering, particularly with over 90% Al2O3 and nopen porosity are considered. Moreover, the exhibited proper-ies including fracture strength and toughness are investigated as

function of self flow processing and cement additions (1–5%)ncluding silica.

. Experimental

The starting raw materials and composition of the ceramicefractories produced in this work are given in Table 1. It wasound that 5.5% water was enough to obtain a self-flowing wellispersed castable system. Flow measurements were carriedut by pouring the castable suspension into a truncated flowone as described in the ASTM standard C-230. All self-floweasurements were based on the ASTM-C71 standard. In addi-

ion, the ASTM C-20 standard was used for determinations ofpparent porosity (AP), bulk density (BD) and water absorptionWA). The mechanical properties such as cold crushing strengthCCS, ASTM C-133), modulus of rupture (MOR, ASTM C-161) and fracture toughness (KIC, ASTM C 1421) were alsoeasured. The samples were cured for 24 h at room temper-

ture (298 K) prior to testing. After de-molding, the samplesere dried at 383 K for 24 h, then fired at different tempera-

ures (1273 K, 1573 K and 1773 K) for 3 h and furnace-cooled.hermo-gravimetric analysis (TGA) and differential thermalnalysis (DTA) were employed to disclose possible reactionsxhibited by the SFCs as a function of the firing temperatures.

lso, after mechanical testing, the fractured samples were cut

nto 25 mm × 10 mm × 10 mm bars and phase identification wasade using X ray diffraction (XRD). In addition, microstructural

able 1aw materials used in producing self-flow refractories.

SFC-5(wt%)

SFC-3(wt%)

SFC-1(wt%)

abular alumina3–6 mm 17 17 171–3 mm 26 26 260–1 mm 27 27 27−325 mesh 13 14 14

ecar 71 (CA-14M) 5 3 1lphabond 300 (hydratable alumina) – 1 3

1000SG (calcined alumina) 7 7 771U (microsilica) 5 5 5-ADS (accelerator) – 0.9 0.9-ADW (retarder) – 0.1 0.1arvan-811D 0.05 – –itric acid (retarder) 0.05 – –ater 5.5 5.5 5.5

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eramic Society 34 (2014) 1365–1373

haracterization of the fracture surfaces was made using scan-ing electron microscopy (SEM) including energy dispersivepectroscopy (EDS).

. Results and discussion

.1. TGA and DTA

Fig. 1 shows the TGA and DTA results of the processed SFCsn terms of weight loss (%) and the DTA (K/mg) parameter,espectively as a function of temperature. As expected, all theFCs exhibit weight losses with increasing temperatures. In par-

icular, there is a characteristic drop in the weight of the castablesn the temperature range of 473–573 K. This weight loss is asso-iated with the dehydration of the hydrated cement phases AH3nd C3AH6, where C = CaO, A = Al2O3, and H = H2O.

According to the literature1,2 the dehydrating temperatures ofH3 and C3AH6 are between 483–573 K and 573–633 K respec-

ively. At elevated curing temperatures, the high density stableydrates (AH3 and C3AH6) produce significant porosity witharge pore sizes as the vapor species are able to escape. In turn,his promotes a sharp decrease in the weight of castables aseen in the above figure. Notice that weight loss (%) is highern the high cement SFC-5 when compared with either SFC-3r SFC-1. Since both, SFC-3 and SFC-1 contain hydrated alu-ina, they do not release chemically bonded water easily at

levated temperatures due to its gel structure when comparedith cement.1

The DTA results of the castables show endothermic reac-ions coupled with the formation of the hydrated phases (AH3nd C3AH6) at around 473–573 K. Notice that hydration startshen cement is in contact with water (exothermic reaction) and

nternally, the castables can reach temperatures of up to 348 K.As the temperature is increased, the process of hydration is

ollowed by the dehydration process and it continues until allhe phases lose their water of crystallization. TGA results show

gradual increase in the weight loss (%) for the castables in theemperature range of 573–1273 K. In addition, the DTA resultshow an almost linear reduction in (K/mg) for temperaturesbove 700 K.

.2. Processing

During sintering, reaction between the calcium aluminateement, silica and the fine fractions of alumina (reactive andlumina sizes below 20 �m) occurs, whereas the coarse aluminarains remain virtually inactive. The reactive alumina is not only

binding agent, but also the main component in achieving goodheological and flowing properties. Fig. 2 shows X-ray diffrac-ion intensity peaks (XRD) of the refractory castables dried at83 K for 24 h. From this figure, it is evident that the refrac-ory structure is predominantly �-alumina as the X-ray intensityeaks correspond to the corundum phase. The XRD patterns of

he SFCs fired at 1773 K for 3 h are shown in Fig. 3. Under theseonditions, mullite (3Al2O3·2SiO2) forms in all of the SFCsompositions considered in this work. It is found that the X-ay intensity peaks corresponding to the mullite phase decrease

C. Gogtas et al. / Journal of the European Ceramic Society 34 (2014) 1365–1373 1367

TA c

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Fig. 1. Exhibited TGA and D

n intensity as the amount of cement in the SFCs is increased1%, 3% and 5%). This effect is attributed to the relatively largemounts of CACs (calcium aluminate cement)) which are knowno attack mullite and/or inhibit its formation.13

During the firing process it is expected that CaO will reactith SiO2 and Al2O3 to form anorthite (CaO·Al2O3·2SiO2).2 In

urn, this can lead to avoid or severely restrict the formation ofullite (3Al2O3·2SiO2). The formation of the low melting phase

northite was observed in both, SFC-5 and SFC-3 refractoriesut not in the SFC-1 one (see Fig. 3). In particular, it was found

hat anorthite formed at 1273 K in the SFC-5 and at 1573 K inhe SFC-3. Apparently, increasing contents of cement promotehe formation of anorthite at relatively low firing temperatures

te

Fig. 2. XRD pattern of (a) SFC-5, (b) SFC

urves in experimental SFCs.

1273 K). From the ternary CaO–Al2O3–SiO2 phase diagram,16

t is expected that the transformation sequence followed by theFC-1 at all the firing temperatures will only give rise to cristo-alite formation. The formation of mullite and/or anorthite isffectively suppressed as there is not enough SiO2 or CaO toorm either of these phases.

.3. Physical properties

Fig. 4 shows the relationship between cement content andhe self flow values exhibited by the experimental castables. Asxpected, self flow values decrease with increasing hydraulic

-3, (c) SFC-1 fired at 383 K for 24 h.

1368 C. Gogtas et al. / Journal of the European Ceramic Society 34 (2014) 1365–1373

Fig. 3. XRD pattern of (a) SFC-5, (b) SFC

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Fig. 4. Self flow values of SFCs as a function of cement content.

lumina contents. This stems from the need for relatively largemounts of water required by hydraulic alumina when comparedith cement in order to set the castables. Moreover, in order to

ontrol the setting time for the castables small additions (usuallyround 0.5%) of cement may be necessary,1 otherwise, large

mounts of water might be required to maintain flow.

The resultant bulk density (BD) and apparent porosity (AP) ofhe SFCs after drying and firing at elevated temperatures is given

do

Fig. 5. Apparent porosity, AP and bulk density, BD for

-3, (c) SFC-1 fired at 1773 K for 3 h.

n Fig. 5. Notice from this figure that at 1773 K, the BD valueseach a maximum in all the experimental castables. In addition,his corresponds to a reduction in AP to the lowest levels. Thiss likely to be due to the formation of the high density mullitehase (3.16–3.22 g/cm3) combined with amorphous anorthiteCaO·Al2O3·2SiO2) which can easily fill the open aggregatenterspaces. The increases in the BD values are found to stronglyepend on the firing temperatures and on the cement content.t firing temperatures between 1273 K and 1573 K the SFC-5ields the highest AP and lowest BD values. Apparently, the lackf enough mullite formation at these temperatures in the SFC-5s the main cause for the exhibited AP and BD values. Anorthites found to start forming at 1273 K and 1573 K in SFC-5 andFC-3, respectively but not in the SFC-1. Hence, this phase isot likely to play a major role in the experimentally determinedP and BD values exhibited in the SFC-1 refractories.

The increase in the AP of the castables is attributed to a

ominant dehydration process when firing temperatures belowr equal to 1273 K were used. In addition, there is a clear

the SFCs as a function of the firing temperatures.

C. Gogtas et al. / Journal of the European Ceramic Society 34 (2014) 1365–1373 1369

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ig. 6. Water absorption, WA (wt%) for the SFCs as a function of the firingemperatures.

ncrease in the apparent porosity and a reduction in density asater evaporates and leaves the SFCs. Above 1273 K, closuref the developed porosity starts to occur leading to densificationhrough the formation of new phases such as anorthite, cristo-alite and mullite. This is further confirmed by determinationsf water absorption in the SFCs (see Fig. 6). Notice from thisgure that WA is the highest in the high cement castable, SFC-5ll the way to 1573 K. Yet, firing at 1773 K drastically reduceshe amount of WA leading to a maximum BD with minimal AP.n contrast, the amount of WA in the SFC-1 is still relativelyigh. This can be explained by considering that hydrated alu-ina does not easily release the chemically bonded water at

levated temperatures due to its gel structure.1

.4. Mechanical properties

Fig. 7 shows the cold crushing strength (CCS) exhibitedy the castables at the firing temperatures considered in thisork. Notice that up to 1273 K the SFC-5 possesses the high-

st CCS values when compared with either the SFC-3 or theFC-1. Apparently, at decreasing cement contents, the devel-pment of bonding between aggregates is rather limited at theow firing temperatures. Nevertheless, after heating at 1773 Kor 3 h, the SFC-3 exhibits the highest CCS values while the

FC-5 strength values slightly drop down. The apparent reduc-

ion in CCS observed in the SFC-5 can be associated with theormation of the low melting anorthite phase. In addition, the

Fig. 7. CCS of the SFCs as a function of the firing temperatures.

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Fig. 8. The MOR of SFCs as a function of temperature.

CS exhibited by the SFC-3 does not change much at the fir-ng temperatures employed. Alternatively, firing the SFC-1 at573 K leads to a vast improvement in CCS. In turn this can beelated to a strong bond formation through the development ofhe cristobalite phase.

The modulus of rupture (MOR) exhibited by SFCs as a func-ion of the firing temperatures is given in Fig. 8. Notice thatimilar trends are observed in the exhibited MOR values whenompared with the ones for the CCS (Fig. 7) (i.e., up to 1273 Khe highest MOR values are found in the SFC-5 with a highement content (5%)). Once again, the presence of cement inhe SFC-5 promotes the early development of ceramic bondingith the aggregates when compared with the SFC-3 and SFC-1

efractories. Moreover, firing at 1573 K for 3 h in both, SFC-3nd SFC-1 leads to significant improvements in the resultantOR, suggesting the development of ceramic bonding. How-

ver, in the SFC-5 the MOR values drop down to 36.34 MPa. Thisoincides with the formation of the low melting phase anorthitehich is expected to reduce the MOR strength of the SFC-5. At773 K, all the SFCs show gradual improvements in the MORalues probably due to the formation of mullite.

In a related work, Martinovic et al.15 found that in a lowement castable (1.2% CaO) firing at 1873 K, resulted in MOR,A, CCS values that are comparable the ones found in this work.

n their work, the exhibited property improvements are attributedo the formation of calcium hexaaluminate (CaO·6Al2O3) orA6 phase2 which forms at elevated temperatures. Hence, theseastables must be fired at temperatures of 1873 for maximumtrength as the strength drops by over 50% when firing at loweremperatures (1473 K). Nevertheless, the reported castable prop-rties after firing at temperatures below 1873 K are well belowhe ones found in this work, particularly MOR and CCS. Inheir work, firing at 1073 K, gives rise to CCS and MOR valueshat are not significantly different from the ones found at 383 K,robably due to the lack of any significant ceramic reaction. Inheir work, they use a silica-free composition where the onlyhases present up to 1673 K are Al2O3, CA and CA2 phases.rom the published literature,1 the CCS values for CA and CA2re 60 and 25 MPa, respectively, and they are unable to provide

ny strength improvements.

Alternatively, in the present work the formation of anorthitend cristobalite phases at temperatures between 1273 K and

1370 C. Gogtas et al. / Journal of the European C

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Fig. 9. Fracture toughness values of SFCs as a function of temperature.

573 K promote appreciable enhancements in both, the MORnd CCS values for the investigated SFCs. Hence, the propertymprovements found in this work are comparable or superioro the ones reported by Martinovic et al.15 Notice in particularhat in the low cement alumina-silicate castables, high strengthan be achieved by processing at relatively low firing tempera-ures (below 1373 K). In addition, firing temperatures as low as573 K are sufficient to produce high strength castables in theltra-low cement compositions (SFC-1).

Fig. 9 shows fracture toughness (KIC) values for the experi-ental SFCs as a function of the firing temperatures. In all the

xperimental SFCs the exhibited fracture toughness values wereetween 3 and 4 MPa m1/2 after firing at or above 1573 K. Typ-cal KIC values for conventional castables are of the order of.2–1.5 MPa m1/2.2 Hence, it is evident that the SFCs producedn this work can be considered as high toughness refractories.otice that the KIC values in both, SFC-3 and SFC-1 steadily

ncrease with the firing temperatures. It is found that the KICends to increase with the cement content except for the SFC-5hich exhibits a slight reduction in the exhibited KIC values at

agm

ig. 10. SEM fractographs of a SFC-3 after drying at 383 K for 24 h. (a) Debonded agnd (c) a pore and its effect on the exhibited fracture path.

eramic Society 34 (2014) 1365–1373

573 K. This effect is consistent with the exhibited MOR val-es found in the high cement castable and it can be attributedo the formation of anorthite. Firing at 1773 K for 3 h improveshe fracture toughness of the SFC-5 refractory to maximum KICalues of up to 4.08 MPa m1/2. The toughness improvements cane related to the formation of mullite concomitant with a reduc-ion in the amount of anorthite and a relatively high densificationith reduced porosity.

.5. Fracture morphologies

Refractory castables are rather brittle at low firing temper-tures as the ceramic bond has not yet developed. Fig. 10a–chows the fracture surfaces of a castable after drying at 383 Kor 24 h. Notice from this figure that the fracture surfaces con-ist of aggregates debonded from the refractory castable binder.hus, it is evident that no strong bond exists between aggregatesnd no binding system develops at this temperature. Fig. 10bhows a matrix area associated with the debonded aggregates.n addition, porosity of various sizes is observed (see Fig. 10a)nd shapes close to spherical morphology (see Fig. 10c). Fromhese figures, it is evident that the fracture surface of the SFC isather smooth suggesting that it is inherently brittle.

Firing at high temperatures promoted the development ofggregate bonding through various reactions. Fig. 9 shows thexhibited KIC for the various SFC refractories after firing at273–1773 K. The exhibited improvements in KIC values arettributed to the formation of ceramic phases such as anorthite,ristobalite and mullite. Fig. 11a shows the fracture surface of

SFC-3 after firing at 1773 K for 3 h. Notice in this figure that

strong chemical bonding has developed between the aggre-ates and the binding phases formed through sintering. For theost part, the number of fractured aggregates increases with

gregates and porosity, (b) detail of the debonded area in the surrounding matrix

C. Gogtas et al. / Journal of the European Ceramic Society 34 (2014) 1365–1373 1371

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ig. 11. SEM micrographs of SFC-3 after firing at 1773 K for 3 h. (a) Fracturggregate bonded to surrounding matrix.

he firing temperatures. Fig. 11b and c is SEM micrographs of aractured aggregate, showing a strong bond with the surroundingatrix. X-ray intensity peaks, (EDS, Fig. 11c) corresponding to

he fracture surface indicates that the fractured aggregate is alu-ina. Apparently, as the aggregates become strongly bonded in

he surrounding matrix, the elastic energy needed for debondingurpasses that for fracture of the alumina grains and aggregateracture becomes prevalent.

Fig. 12a and b is SEM micrographs corresponding to aggre-ate fractures in SFC-3 fired at 1273 K and 1773 K, respectively.n both cases, the aggregate fracture is strongly influenced byhe crystal structure of the alumina as multiple ledges or steps

evelop along the fracture path. In addition, at the highest firingemperatures the fracture path becomes increasingly tortuouss the length of the steps or ledges developed are rather short

lic

Fig. 12. SEM micrograph of SFC-3 after firin

ace, (b) X-ray peak intensity corresponding to Al and (c) a fractured alumina

ompared with the ones found in the SFC-1. Moreover, appre-iable closed porosity is also observed within the aluminaggregates. All of these factors can account for the exhibitedtrength and KIC values, as well as for their relative differencesn magnitudes. In particular, they can explain the relatively high

IC values found in the SFC-5 refractories.Fig. 13a and b shows the fracture surfaces of SFC-5 and

FC-1 samples after firing at 1773 K for 3 h. From these fig-res, it is evident that there is an increasing number of shortized cracks developing in the SFC-5 when compared with theFC-1 fracture surfaces. The development of multiple cracking

n the SFC-5 at 1773 K can be related to the formation of mul-

ite at these temperatures. As mullite develops there is a volumencrease associated with the transformation which can promoteracking through the development of residual stresses. Other

g (a) at 1273 K and (b) 1773 K for 3 h.

1372 C. Gogtas et al. / Journal of the European Ceramic Society 34 (2014) 1365–1373

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actors contributing to the development of residual stressesnclude elastic anisotropy due to thermal expansion differencesmong alumina aggregates and also among the new developedhases (anorthite and mullite).

Although there are no fracture mechanics models that canuantitatively explain toughening in single phase polycrys-alline ceramics, among the suggested toughening mechanismsre,17 (a) microcrack toughening and (b) crack bridging. Theechanism of microcrack toughening assumes that stable grain

oundary microcracks are nucleated by the high stresses in theicintiy of the macroscopic crack tip. These micro-cracks thenower the stress experienced by the tip. In the second approach,17

lteration of the stress intensity of a macro-crack tip is assumeds the crack interacts with a discrete array of micro-cracks.he proposed toughenning mechanisms are in agreement with

he observations of an increasing density of short size micro-racks found in the SFC-5 castables. In turn, the exhibitedmprovements in the SFC-5 samples can be attributed to theevelopment of discrete arrays of micro-cracks as a result ofelatively high residual stresses. The development of increasingesidual stresses can be related to the contribution of increas-ng phase formation reactions (mullite and anorthite) inducedt aggregate–aggregate boundaries. Notice that in SFC-1 casta-les, cristobalite is the dominant phase that develops at the firingemperatures. In turn, the residual stress development associatedith the aggregates and the cristobalite phases is not effective

nough to promote the development of a high density of shortracks.

. Summary

In this work, low cement self flowing castables (SFCs) wererocessed. A water addition of 5.5% was found to be enougho achieve full self flow properties. It was found that the selfow values decreased as the volume fraction of hydratable alu-ina was increased. Drying at 383 K, produced a fully corundum

atrix structure in all the SFCs. The formation of anorthite was

bserved in SFC-5 and SFC-3 at 1273 K and 1573 K, respec-ively. In contrast, cristobalite was found to be present in SFC-1t 1573 K. At 1773 K, SFCs having 1% and 3% cement content

(b) SFC-1 after firing at 1773 K for 3 h.

eveloped relatively large amounts of mullite as evidenced by-ray diffraction.The cold crushing strength values were found to reach a max-

mum at 1573 K in the SFC-3 and SFC-1 castables and at 1273 Kor the SFC-5 ones. In addition, in the SFC-3 and SFC-1 casta-les, KIC and MOR values were found to consistently increaseith the firing temperatures as a result of the development of

eramic bonding. In contrast, a slight drop in KIC and MORalues was observed in the SFC-5 at 1573 K, probably due tohe formation of anorthite. Yet, the exhibited KIC in the SFC-5reatly improved after firing at 1773 K. The highest values of BDorresponded to the lowest values of AP and they were attainedfter firing at 1773 K in all of the castables.

cknowledgements

The authors would like to acknowledge the support of Steveowalski and Dale Zacherl from Almatis Corporation by pro-iding all of the materials used in this research work.

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