Cellular Anorthite Glass–Ceramics: Synthesis, Microstructureand Properties

Mirabbos Hojamberdiev,z,y Jessica D. Torrey,z Marilia Sergio da Silva Beltrao,z and Lothar Wondraczekw,z

zDepartment of Materials Science, University of Erlangen-Nuernberg, 91058 Erlangen, Germany

yUzbekistan Academy of Sciences, Laboratory of High Temperature Materials Chemistry, Institute of General andInorganic Chemistry, Tashkent, Uzbekistan

zDepartment of Materials Science and Engineering, University of Washington, Seattle, Washington 98195-2120

Highly permeable cellular anorthite glass ceramics with porosityof B95% were obtained by a simple replication technique usingnear-stoichiometric glass powders and a polymeric foam as sac-rificial template material. Impacts of sintering conditions and ad-ditions of minor constituents, respectively, on microstructuralevolution and resulting macroscopic properties of the derivedfoams were investigated by X-ray diffraction, scanning electronmicroscopy (SEM), differential thermal analysis and X-ray mi-crocomputer tomography. Apparent activation energies of crys-tallization and the Avrami coefficient were estimated fromnonisothermal crystallization experiments to evaluate the impactof titania and zirconia, respectively, as potential nucleationagents. Consistent with SEM analyses, it was found that crystalgrowth occurs in two dimensions. While TiO2 primarily acts onthe viscosity of the precursor glass and, thus, on the process ofsintering by viscous flow, ZrO2 is found to exhibit at least somenucleation efficiency. As compared with sintering of ceramic pow-ders and solid-state reactions in general, the glass ceramic routeenables significant reduction in sintering time and temperature.

I. Introduction

CELLULAR (macroporous) ceramics with porosity exceeding60 vol% are receiving continuously growing interest for

various applications such as particulate filters, molten metal fil-ters, catalyst supports, acoustic insulators, porous burners, andbiomedical devices.1,2 By designing highly specific microstruc-tures and varying the strut composition, property profiles ofsuch materials can be tailored for a wide range of operatingconditions and applications.

A recent review on different preparation techniques is givenby Studart et al.,3 ranging from replication, direct foaming, andextrusion to the use of sacrificial templates. Among these, sincethe early 1960s,4 the use of sacrificial templates is, at the labo-ratory scale, the widest spread preparation technique. It allowsfor highly specific tailoring of porosity (20–90 vol%) and poremorphology, but requires availability of a suitable template. Itusually consists of five consecutive steps that can be applied al-most universally to any glass or ceramic material: preparation ofa slurry with low viscosity, impregnation of a template structure(e.g., polymer foam), removal of excess slurry, drying and burn-out of the template, and sintering at elevated temperature.

Here, we consider cellular anorthite (CaAl2Si2O8) glassceramics as a potential material for some of the above noted

applications. Anorthite is a calcium-rich aluminosilicate mineraland member of the large plagioclase fieldspar group. It is an im-portant constituent of many igneous and metamorphic rocks andbelongs to the most abundant minerals in earth’s upper crust. Thecrystal structure of anorthite comprises a three-dimensional (3D)array of corner-sharing [AlO4] and [SiO4] tetrahedra, linked to-gether in a framework of [Al2Si2O8]

2� composition in whichcharge-balancing calcium (Ca21) cations occupy four distinct, ir-regular cavities. From technological and economic viewpoints,interest arises from its relatively high melting point (15531C),5 lowcoefficient of thermal expansion (4.82� 10�6 K�1), low dielectricconstant (6.2 at 1 MHz), potentially high mechanical strength atlow specific weight, and also from the abundance of raw mate-rials. Preparation via the glass ceramic route, i.e. synthesis bycontrolled sintering and crystallization of a glass, promises lowerreaction temperatures, significantly shortened reaction times andthe use of recycled or waste materials in the batch.

II. Experimental Procedure

(1) Glass Powder and Slurry Preparation

Powder batches (100 g) of composition (mol%) CaO–Al2O3–2SiO2 (anorthite stoichiometry, sample An), were prepared byproperly mixing analytic grade SiO2, Al2O3, and CaCO3 (Merck,Darmstadt, Germany), adding 2 wt% B2O3 as fluxing agent(H3BO3, Carl Roth GmbH, Karlsruhe, Germany). In some cases,4 wt% TiO2 (Kronos Inc., Houston, TX, sample AnTi) and ZrO2

(Alfa Aesar, Ward Hill, MA, sample AnZr), respectively, wereintroduced to study the impact of these compounds on hetero-geneous nucleation. As-prepared batches were calcined for 1 h at9001C before conventionally melting in Pt crucibles at 16001C, 1h. Glass frit was obtained by quenching glass melts in distilledwater. The frit was dried, crushed, and dry milled in a vibrationdisk mill (Siebtechnik GmbH, Mulheim a.d. Ruhr, Germany) for1 h to obtain fine powders. Particle size distribution was deter-mined by laser-granulometric analysis of the glass powders (Mas-ter Sizer 2000/Hydro 2000S, Malvern Instruments, Malvern,U.K.), specific surface area by the N2 adsorption method(BET, ASAP 2000, Micromeritics, Monchengladbach, Ger-many). Glass transition and crystallization temperatures were de-termined by differential thermal analyses (DTA) (Netzsch STA429, Netzsch-Geratebau GmbH, Selb, Germany) at heating ratesranging from 5 to 20 K/min (ambient to 12001C). Linear coeffi-cients of thermal expansion were measured on samples with a sizeof 6 mm� 6 mm� 15 mm (particle size given Table I) using ahigh-temperature dilatometer (Netzsch DIL 402 C, Netzsch-Geratebau GmbH) operating between room temperature and13001C at a heating rate of 5 K/min in air. Standard properties ofthe glass powders are listed in Table I.

Using these powders, 120 mL of slurry were prepared by thefollowing procedure: ethanol (57 vol%) was mixed with 1.5

P. Colombo—contributing editor

wAuthor to whom correspondence should be addressed. e-mail: lothar.wondraczek@ww.uni-erlangen.de

Manuscript No. 25889. Received March 11, 2009; approved June 9, 2009.

Journal

J. Am. Ceram. Soc., 92 [11] 2598–2604 (2009)

DOI: 10.1111/j.1551-2916.2009.03268.x

r 2009 The American Ceramic Society

2598

vol% of a polyester/polyamine copolymer dispersion agent(10000 g/mol, Uniqema, Emmerich, Germany, Hypermer KD1,1.11 g/cm3) in a plastic bottle for 30 min. Alumina balls (2 mm)were added as grinding media. Thirty six volume percent of glasspowder were added to the solution and thoroughly mixed for 24 hin a heavy-duty shaker (Willy A. Bachofen AGMaschinenfabrik,Basel, Switzerland). Afterwards, 4 vol% PVB-98, polyvinyl but-yral (Solutia, density: 1.10 g/cm3) were added. Finally, 1.5 vol%Santicizer 261A, alkyl benzyl phthalate (Ferro Corp., Cleveland,OH, density: 1.063 g/cm3) were added as binder and plasticizer,respectively. Viscosity measurements were performed in a rheo-meter (UDS 200, Paar Physica Messtechnik GmbH, Stuttgart,Germany). Shear viscosity (0.9 Pa � s at a shear rate of 100 s�1 atroom temperature) was calculated from the storage modulus G0

and loss modulus G00.

(2) Preparation of Anorthite Foams

Commercially available polyurethane foams (Koepp SchaumGmbH, Elchingen, Germany; 4075 pores per in.) were used aspolymer templates. These foams were cut into pieces of 25mm� 25 mm� 25 mm, immersed in the slurry and squeezedmanually to remove excess slurry. Slurry-loaded foams weredried at room temperature for 24 h and at 901C for 10 h. Afterdrying, samples were heated to 4501C at a heating rate of 1 K/min and held at this temperature for 1 h to burn out the poly-mer. Sintering was performed at a temperature between 9001and 13001C for 1 h in air, applying heating and cooling rates of 5K/min. Total sintering shrinkage (%) was calculated from thechange in sizes of the samples (measured with a digital caliper)before and after firing.

(3) Characterization

X-ray diffraction (XRD) spectra were recorded (Diffrac 500,Siemens AG, Mannheim, Germany, CuKa) in order to confirmformation of anorthite as the predominant crystal phase in thesintered foams. Morphology and crystal formation were furtheranalyzed by scanning electron microscopy (SEM) (Quanta 200,FEI, Prague, Czech Republic) and energy-dispersive X-rayspectrometry (EDX). Microcomputer tomography (mCT 40,Scanco Medical AG, Bassersdorf, Switzerland) was used to an-alyze cell morphology of the foam. For that, samples werescanned in x, y, and z directions with an isotropic resolutionof 37 mm. The object was rotated through 3601 with 11 per step.3D reconstructions and histomorphometric parameters (e.g.,cell size distribution, strut thickness) were derived from themCT raw data, using the distance transformation method.6,7

Assuming fully accessible porosity (open porosity), strut den-sity rs was directly determined with an He-pycnometer (Micro-meritics Accupyc 1330, Norcross, GA) on ground samples,dividing the derived sample volume by the sample mass. Totalporosity e was evaluated from the difference of strut density andgeometrical density (geometrical volume Vg of a cube of foam,measured with a micrometer screw, divided by the mass of thecube, m), e5 (Vg�m/rs)/Vg. Obtained values of e correlate wellwith later CT analyses.

To determine compressive strength, cubic samples sintered atdifferent temperatures were subjected to a uniaxial compressiveload in a universal testing machine (Instron 4202, Instron Corp.,Canton, MA, 5 kN, deformation rate of 1 mm/min). Owing tothe complex preparation procedure, only five specimen wereconsidered for each sintering temperature. This specimen num-ber is not sufficient for a complete statistical analysis. Therefore,only average values and standard deviations were determinedfor each set of samples.

Chemical durability was evaluated by determining the loss ofweight after chemical leaching. For each test, five samples wereimmersed in 100 mL of 1% HNO3 and 20% NaOH, respec-tively, at 1001C for 3 h, washed in distilled water, dried at 1001Covernight and weighted.

Table I. Basic Properties of Precursor Glass Powders

Property

Glass compositions

An AnTi AnZr

CaAl2Si2O8 98 94 94ZrO2 — — 4TiO2 — 4 —B2O3 2 2 2Density, 70.01 g/cm3 2.70 2.71 2.74Particle size (d50), mm 11.08 11.76 12.50Specific surface area, 70.5 m2/g 7.8 5.7 9.2Tg, 721C 815 785 812Tc, 721C

5 K/min 965 925 96410 K/min 986 940 98620 K/min 1005 959 1004

Ec, 730 kJ/molKissinger 434 478 432Matusita and Sakka 744 817 741u 2.2 2.6 2.0

CTE @ 9001C, 70.1� 10�6 K�1 6.72 7.05 7.62

Batch compositions are given in wt%.

Fig. 1. Scanning electron microscope images of anorthite-based glass–ceramic sample An prepared at 13001C for 1 h: overview of the foam (a)and triangular hollow strut (b).

November 2009 Cellular Anorthite Glass–Ceramics 2599

III. Results and Discussion

(1) Microstructure

Figure 1(a) shows the cellular structure of an as-prepared anort-hite glass ceramic (sample An): replicating the microstructure ofthe used polymer template, it consists of a 3D network of hollowpolyhedra (cells) with edges randomly oriented in space. Hollowstruts with triangular shape are a result of template-burnoutduring sintering (Fig. 1(b)).2,8–10 The total porosity as evaluatedfrom strut and geometrical densities is B95 vol%.

HT-XRD patterns, recorded in the temperature range of8501–13001C are shown in Fig. 2 (exemplarily for the TiO2

and ZrO2-free material). Similar for all samples, patterns revealcrystallization between 8501 and 9001C, anorthite (PDF CardNo. 41-1486) being the predominant crystal phase. Noteworthy,no intermediate or secondary phase could be detected in theconsidered temperature range in neither of the three samples(with the exception of minor amounts of rutile in AnTi andzirconia in AnZr, respectively, as further confirmed by SEManalyses). For the solid-state route,11 Okada et al. reported theoccurrence of gehlenite (Ca2Al2SiO7) as a dominant intermedi-ate phase in unground samples, and layered CaAl2Si2O8 inground samples.10 These phases, however, were not detectedby Kobayashi and Kato12 when preparing dense anorthite ce-ramics on a low-temperature route.

The crystallization mechanism and kinetics of the glass pow-ders were investigated by DTA. DTA traces were recorded up toa temperature of 12001C, applying different heating rates(Fig. 3). From these curves, the position of the crystallizationpeak (Tc) and glass transition Tg were determined with an ac-curacy of72 K (Table I). Additions of TiO2 were found to leadto decreasing values of these temperatures while ZrO2 had al-most no effect. The fact that both compounds have only littleeffect on the difference between Tg and Tc indicates that theirinfluence is mainly on melt viscosity, and, if any, only to a lesserextend on initiation (nucleation) of crystallization.13

The apparent activation energy of crystallization (Ec) can,in principle, be estimated by the Kissinger equation,14 using

different heating rates. However, as discussed by Donald andcolleagues,12,15,16 this ignores the impact of nucleation and di-mensionality of crystal growth (and, strictly spoken, cannot be applied to DTA data). In the present case, we therefore

Fig. 2. X-ray diffraction patterns of sample An after sintering at theindicated temperature for 1 h.

Fig. 3. Differential thermal analysis scans of samples An (a), AnTi (b),and AnZr (c) for different heating rates.

2600 Journal of the American Ceramic Society—Hojamberdiev et al. Vol. 92, No. 11

rely on Matusita and Sakka’s17 modification of the Kissingerequation:

lnðfm=T2c Þ ¼ �nEc=RT þ const (1)

where Tc is the crystallization peak temperature in a power-compensated DSC experiment, f is the heating rate, R is theideal gas constant, and m and n are parameters that depend onthe process of nucleation and crystallization. For bulk crystal-lization of as-quenched glasses with a heating rate-dependentnumber of nuclei, and assuming a diffusion-controlled crystal-lization mechanism, it is n5 2.5 and m5 1.5.15 Then, apparentactivation energies of crystallization for the three different sam-ples can be calculated from the peak positions in Fig. 3. Data aregiven in Table I. Additional qualitative information on the crys-tallization mechanism can be derived from the Avrami param-eter, u. Using isothermal crystallization data,18 it can be deriveddirectly from the Johnson–Mehl–Avrami–Kolmogorv equa-tion.19 Its deviation from nonisothermal data, however, is con-troversial.16,20,21 Here, we consider the relatively pragmaticapproach of Augis and Bennet16 to obtain a rough estimate of u:

u ¼ 2:5RT2

p

EcFWHM

�����

(2)

where FWHM is the full-width at half-maximum of the ob-served DTA peak at Tp. Values of u, calculated from the scanswith a heating rate of 20 K/min and using Kissinger’s apparentactivation energy are given in Table I. They clearly indicate two-

dimensional crystal growth,15,16 what is consistent with SEManalyses as will be discussed in the following.

Macroscopic properties of glass–ceramics strongly depend onmicrostructure—crystallite size and number density, andamount and viscosity of the residual glass phase—and, there-fore, on the crystallization and nucleation processes. The moreinformation on these mechanisms is available, the better themicrostructure can be controlled by adjusting glass compositionand temperature profile of the conversion process. The evolutionof microstructure is shown in Fig. 4. Each SEMmicrograph wastaken after heat treating for 1 h at the indicated temperature.For low sintering temperature (i.e., below Tc), foams are com-pletely in the glassy state and no significant differences are vis-ible for An, AnZr, and AnTi. At this stage, sintering of theprecursor glass occurred by viscous flow. With increasing tem-perature, the roughness of the strut surface increases (similarlyfor all samples), what is related to the onset of (surface) crys-tallization. If the sintering temperature exceeds 11001C, the typ-ical microstructure of interlocked crystals starts to develop. Inparallel, the impact of TiO2 and ZrO2 additions, respectively,becomes visible. Both TiO2 and ZrO2 exhibit only very limitedsolubility in the plagioclase lattice and therefore tend to precip-itate in secondary phases. After heat treating for 1 h at 13001C,sample An (Fig. 4(a)) exhibits a very rough strut surface withless uniform crystals of an approximate size of 1 mm. Additionsof TiO2 and ZrO2 seem to result in larger crystals (45 mm) aswell as in the precipitation of secondary crystals of ancicularshape (AnTi, Fig. 4(b)) and lesser size (AnZr, Fig. 4(c)), respec-tively. In AnZr, a hail of bright spots covers the primary anort-hite crystals. The SEM contrast indicates that these spots are

Fig. 4. Scanning electron microscopic images of An (a), AnTi (b), and AnZr (c) samples as a function of sintering temperature after 1 h of sintering.

November 2009 Cellular Anorthite Glass–Ceramics 2601

areas that are rich in zirconia—as anorthite crystallizes from thebase glass, ZrO2 first accumulates in the residual glass phase(increasing its liquidus temperature22) and, when the liquidustemperature exceeds the sintering temperature, precipitates in asecondary phase (o0.5 mm crystallite size). For AnTi, the EDSspectrum (Fig. 4(b)) indicates that the acicular crystals are richin TiO2. Taking into account previous observations on the crys-tallization of TiO2-bearing glass powders of anorthite composi-tion,23 it can be concluded that the secondary phase in Fig. 4(b)is comprised of rutile whiskers (confirming XRD observations).

Figure 5 is a 3D reconstruction of a mCT scan of the anorthitefoam (a) and corresponding geometrical data (b). The recon-structed image clearly shows the highly interconnected cell struc-ture as well as the very high porosity (Fig. 1). Mean strut thicknessslightly differs among the three samples in the order An (228 mm)oAnZr (300 mm)oAnTi (325 mm) (710%). Correspondingly,mean cell size decreases with increasing strut thickness, An(1.3070.34 mm)4AnZr (1.1270.38 mm)4AnTi (1.0770.36mm). These differences are mainly associated to differences inthe sintering behavior (viscosity) of the three precursor powders.

During sintering, intrinsic density of the struts increases dueto relaxation of the glass phase, disappearance of residual(closed) porosity during sintering and crystallization. Observa-tion of the macroscopic shrinkage of cellular specimen reveals a

clear transition between two regimes of sintering (Fig. 6): vis-cous flow below B11001C and solid-state sintering at highertemperatures. Differences between the three samples are mainlydue to the presence of ZrO2 in AnZr what leads to an increase indensity in both glassy and crystalline specimen. Clearly, viscousflow is the process that governs the sintering process. Remem-bering the difference between Tg and Tc for the three glass sam-ples, additions of ZrO2 and TiO2 can be used to adjust glassviscosity without significantly changing Tc. Noteworthy, theoverall densification, i.e. the difference in density between theprecursor glass powder and the crystallized strut also differsfrom 1.5% (An) to 1.9% (AnTi) and 2.5% (AnZr). Densificat-ion provides important information on the crystallizationprocess. Here, the differences between the three samples quali-tatively indicate changes in the nucleation process and a tran-sition from purely surface-controlled to parallel surface andbulk crystallization. While surface crystallization often results inthe formation of inner pores (Fig. 1), bulk crystallization gen-erally leads to more homogeneous microstructures and highertotal densification.

Fig. 5. 3D reconstruction (a) strut thickness and (b) cell size distribu-tion of the anorthite-based glass–ceramic foams sintered at 13001C for1 h. Dashed, solid, and dotted lines are related to An, An1TiO2, andAn1ZrO2 samples, respectively.

Fig. 6. Total shrinkage of foams An (circle), AnTi (rectangle), AnZr(diamond) as a function of sintering temperature after sintering for 1 h.

Fig. 7. Compressive strength of foams An (circle), AnTi (rectangle),AnZr (diamond) as a function of sintering temperature after sinteringfor 1 h.

2602 Journal of the American Ceramic Society—Hojamberdiev et al. Vol. 92, No. 11

(2) Physical Properties

Figure 7 is a plot of the measured compressive strength of thefoams that was obtained after 1 h of sintering as a function ofsintering temperature. Compressive strength directly correlateswith strut density and degree of sintering, respectively: thehigher the sintering temperature and, thus, the obtained strutdensity, the higher the compressive strength of the foam. Aquantitative approach to this relation was proposed by Breznyand Green.24 They experimentally confirmed that compressivestrength (sf) and the normalized density ratio (r/rs) are linkedby the (empirical) relation zf 5Csfs(r/rs)

3/2, where C is a con-stant, C B0.65, and sfs is the macroscopic fracture strength ofthe strut material.25 It was suggested that by using this formula,it is also possible to evaluate the compressive strength of indi-vidual struts as a function of the normalized density. Sousaet al.26 further reported on the impact of strut thickness oncompressive strength of cellular ceramics. For example, for sam-ple An (strut thickness of 228 mm, Fig. 5), we find the lowestcompressive strength (0.80 MPa, sintering for 1 h at 13001C).However, also the microstructure of the struts and, as discussedpreviously, the presence of secondary phases (rutile whiskersand zirconia precipitates) alter strut strength. Here, due to therelatively large inner porosity of the struts, the mean compressivestrength was found to be below that of other aluminosilicate ma-terials, e.g. Gibson and Ashby.25 If the extend of internal nucle-ation (bulk crystallization) in relation to surface nucleation can beincreased, also the compressive strength of the foams will in-crease. Results indicate that in this context, particularly additionsof ZrO2 are beneficial. Noteworthy, typical values of the com-pressive strength of commercial ceramic foams lie in the range of0.5–2MPa for porosities between 70% and 90%.27 In this respect,the here observed values of B1 MPa hold significant promise.

Hydrothermal dissolution of anorthite was studied by Fred-erickson and Cox28 They found the ion exchange reactionCa2122 H3O

1 as the first step of dissolution, leading to theformation of a hydrated layer on the anorthite surface. For thepresent case, data on chemical resistance (after sintering for 1 hat 13001C, leaching for 3 h at 1001C) are reported in Table II.Samples readily dissolve in acidic environment (complete disso-lution was observed after leaching in 5%–10% HNO3). Con-trary, they exhibit relative stability against alkali solutions. Thatis, no dissolution (weight loss) could be observed after leachingfor 3 h in 5% and 10%, respectively, NaOH. Additions of ZrO2

and TiO2, respectively, have, if any, an only minor influence onthe corrosion behavior. Besides being highly corrosion resistantby itself, particularly ZrO2 additons seem to positively impactresistance by leading to a higher degree of crystallinity and im-proved homogeneity of the microstructure.

IV. Conclusions

Reported data allow drawing the following conclusions:(1) Cellular anorthite glass–ceramics with high porosity can

be produced by a simple replication method, using polymericfoams as sacrificial template material. When using near-stoic-hiometric anorthite glass powders as precursor, sintering occursat temperatures around 8001 to 10001C, predominantly by vis-cous flow. For example, heat-treating for 1 h at 13001C pro-duces a highly crystalline anorthite foam with a compressivestrength of about 1 MPa for a porosity of 490%. Compared

with conventional sintering of ceramic powders and solid-statereactions in general, both maximum sintering temperature andsintering time can be reduced significantly.

(2) Investigation of the crystallization mechanism and ki-netics of the glass powders (apparent activation energy of crys-tallization and Avrami parameter) indicate that crystal growthprimarily occurs two-dimensional, what is consistent with SEManalyses. Additions of TiO2 and ZrO2 result in decreasing glasstransition temperature and peak crystallization temperature,thus affecting first sintering and, at higher temperature, micro-structural evolution. At elevated temperature, both componentscrystallize in secondary phases. Particularly for TiO2 additions,formation of rutile whiskers can clearly be observed after sinte-ring at temperatures above 11001C. On the other hand, it wasalso found that ZrO2 seems to exhibit at least some efficiency asheterogeneous nucleation agent.

(3) Depending on the used template material, glass ceramicfoams exhibit a highly interconnected cell structure with cell sizeranging from B1.0 to 1.3 mm and strut thickness ranging fromB220 to 330 mm. This results in very high Darcian permeability.

(4) As is typical for anorthite and other calcium-rich mate-rials, the foams exhibit only very little resistance to acidic attack,but decent stability under basic conditions.

Acknowledgments

M. H. would like to thank the Deutscher Akademischer Austausch Dients(DAAD) for the award of a research fellowship under which the present study wascarried out. The Institute of Science and Technology ofMetals of the University ofErlangen-Nuernberg is gratefully acknowledged for providing generous access tothe microcomputer tomography.

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Table II. Chemical Resistance of the Cellular Anorthite GlassCeramic After Sintering at 13001C for 1 h

Samples 1% HNO3 20% NaOH

An 15.00 10.83An1TiO2 14.64 10.49An1ZrO2 14.27 8.65

Weight loss (%) after leaching for 3 h at 1001C. No mass loss could be detected

after leaching for 3 h in distilled water (1001C).

November 2009 Cellular Anorthite Glass–Ceramics 2603

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2604 Journal of the American Ceramic Society—Hojamberdiev et al. Vol. 92, No. 11