j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 2 0 9 ( 2 0 0 9 ) 1134–1142

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Extruded ZrSiO4 particulate-reinforced LZSA glass–ceramicsmatrix composite

F.M. Bertana, O.R.K. Montedoa, C.R. Rambob,∗, D. Hotzab, A.P. Novaes de Oliveirac

a SENAI/CTCmat, Center of Technology in Materials, Rua General Lauro Sodre, 300, PO Box 3247,Bairro Comerciario, 88802-330 Criciuma (SC), Brazilb Department of Chemical Engineering (EQA), Federal University of Santa Catarina (UFSC), PO Box 476,88040-900 Florianopolis (SC), Brazilc Department of Mechanical Engineering (EMC), Federal University of Santa Catarina (UFSC), PO Box 476,88040-900 Florianopolis (SC), Brazil

a r t i c l e i n f o

Article history:

Received 30 August 2007

Received in revised form

12 March 2008

Accepted 15 March 2008

a b s t r a c t

This work reports on the characterization of ZrSiO4 particulate-reinforced Li2O–ZrO2–

SiO2–Al2O3 (LZSA) glass–ceramic matrix composite added with bentonite as binder and

formed by extrusion. The glass batches and composites were characterized on the point

of view of their typical physical/mechanical and chemical properties. Composition with

60 wt.% ZrSiO4 was preliminary selected, since it showed the best results in terms of bend-

ing strength (190 MPa) and deep abrasion resistance (51 mm3). The same composition as

before but added with 7 wt.% bentonite was selected for further studies since it exhibited

Keywords:

Extrusion

Glass–ceramics

Sintering

the highest plasticity index, which resulted in good billets after extrusion. In this last case,

the extruded samples, after sintering at 1150 ◦C for 10 min, showed a thermal linear shrink-

age of 14% and deep abrasion resistance and bending strength of 51 mm3 and 220 MPa,

respectively.

and can be justified only for large production. A viable alterna-

Crystallization

1. Introduction

Glass–ceramics are relatively new materials specially used duetheir specific properties such as high bending strength, highabrasion resistance, high hardness and wide range of coeffi-cient of thermal expansion (CTE), which yields to high thermalshock resistance and high chemical resistance (according tothe chemical composition). These features basically dependon the nature, size and distribution of the formed crystals aswell as on the residual glassy phase (Strnad, 1996; Duan et al.,1999). The nature of the formed crystalline phases and conse-

quently the final properties can be controlled by modifying thechemical composition of the parent glass and also by adequateselection of the heat-treatment parameters.

∗ Corresponding author.E-mail address: rambo@enq.ufsc.br (C.R. Rambo).

0924-0136/$ – see front matter © 2008 Elsevier B.V. All rights reserved.doi:10.1016/j.jmatprotec.2008.03.018

© 2008 Elsevier B.V. All rights reserved.

Glass–ceramics are interesting not only by their prop-erties but also because of the possibility to produce themusing low cost raw materials like residues from steel indus-try, glass wastes and fly ashes, which can be transformed intoproducts with optimized properties for a given application(Rabinovich, 1985; Yoon and Yun, 2005; Yun et al., 2006). Theclassical fabrication of glass–ceramic materials consists on thepreparation of monolithic glass components followed by heat-treatments for crystallization (McMillan, 1979; Simmons et al.,1982). However, this technology requires great investments

tive could be the production of glass–ceramics processed fromglass powders and consolidated by sintering using the sameequipments of traditional ceramic plants. This technology

t e c h n o l o g y 2 0 9 ( 2 0 0 9 ) 1134–1142 1135

w(paos

rih1sL(ttetMttetbttp

uincmialasntoorpatbm

tSt

2

ASmd

Table 1 – LZSA-frit and ZrSiO4 fractions of the preparedcompositions

Composition LZSA content (%) ZrSiO4 content (%)

A 100 0B 90 10C 80 20D 70 30E 60 40F 50 50

j o u r n a l o f m a t e r i a l s p r o c e s s i n g

as developed in USA (Rabinovich, 1985; Miller, 1975), RussiaRabinovich, 1985), Sweden and Israel (Rabinovich, 1985). Therocess involves the following basic steps: (a) glass meltingnd cooling; (b) pulverization; (c) forming by ceramic technol-gy (unidirectional pressing, extrusion, slip casting, etc.); (d)intering for consolidation and crystallization.

In the last years the technological importance of thiselatively new class of glass–ceramic obtained by sinter-ng and controlled crystallization of parent glass powdersas increased. Oliveira (Novaes de Oliveira et al., 1996,998a,b,c, 2000, 2004; Novaes de Oliveira and Leonelli, 1998)tudied a sintered glass–ceramic material belonging to theZS (Li2O–ZrO2–SiO2) with the same or better propertiesbending strength, abrasion and chemical resistance) thanhose of some commercial glass–ceramics and other tradi-ional ceramic products or natural stones. The propertiesxhibited by the LZS glass–ceramic are due to the formed crys-alline phases (zircon, ZrSiO4 and lithium dissilicate, Li2Si2O5).ontedo et al. (2004) modified the LZS glass–ceramic sys-

em by partially substituting zirconium oxide by alumina sohat the resulted Li2O–ZrO2–SiO2–Al2O3 (LZSA) glass–ceramicxhibited a coefficient of thermal expansion lower thanhat of the LZS glass–ceramic system. A study performedy Novaes de Oliveira and Manfredini (1998) also showedhat is possible to improve the mechanical properties ofhis glass–ceramic system by adding alumina as reinforcingarticles.

In powder technology a forming method that can besed for the production of sintered glass–ceramic materials

s extrusion. Extrusion is a very productive forming tech-ique that is used for mass production of components, whichan weight from 1 ton to few grams, e.g. traditional buildingaterials such as bricks and roof tiles, refractory ceram-

cs, electrical insulators and electronic substrates. Extrusionllows low cost manufacturing of components with thicknessower than 1 mm with complex shapes (Reed, 1995; Malpanind Kumar, 2007). For extrusion the material must exhibitome plasticity, property that glassy and ceramic materials doot show (Barba et al., 1997) and directly affects the quality ofhe extruded product. To provide plasticity to a ceramic body,rganic and inorganic binders can be used. However, the use ofrganic binders may raise the processing costs since debindingequires long heat-treatment steps. Furthermore, the decom-osition of carbon residues may generate some defects suchs internal porosity and cracks. An alternative found to solvehese problems could be the use of inorganic binders likeentonite, which in this case is incorporated to the finishedaterial.This work reports the manufacturing and characteriza-

ion of extruded ZrSiO4 particulate-reinforced Li2O–ZrO2–iO2–Al2O3 glass–ceramic matrix composites added with ben-onite as binder.

. Experimental procedure

glass composition belonging to the system Li2O–ZrO2–iO2–Al2O3 was prepared from commercially available rawaterials. Details of the preparation and characterization are

escribed in a previous publication (Montedo et al., 2004).

G 40 60H 30 70

In order to obtain composite compacts by extrusion, com-mercial ZrSiO4–zircon (average particle size of 4.5 �m) from 10to 70 wt.%, in 10% intervals and 7 wt.% bentonite (average par-ticle size of 2.8 �m), both supplied by Colorminas-Brazil, wereadded to the glass powder and mixed and humidified withwater (23 wt.%). Table 1 shows the Li2O–ZrO2–SiO2–Al2O3-fritand ZrSiO4 fractions of the prepared compositions. Sub-sequently, the mixtures were stored for 12 h to moisturehomogenization and then extruded in a Netzsch extruderMA 01 so that compacted samples with nominal dimensionsof 80 mm (length) × 25 mm (width) × 5 mm (thickness) wereobtained. After 48 h at 20 ◦C extruded samples were dried at110 ◦C and then isothermally sintered in an electric furnaceat a heating rate of 10 ◦C min−1 in air for 10 min at tempera-tures in the range of 500–1300 ◦C. After sintering, samples wereair-quenched to room temperature.

To define the best added reinforcing fraction, using ascriteria densification, deep abrasion resistance and bend-ing strength, glass powder and appropriated zircon amounts(10–70 wt.%) were wet mixed (water content = 9 wt.%) and uni-axially pressed at 40 MPa in a steel die so that compacted sam-ples with nominal dimensions of 100 mm × 50 mm × 10 mmwere obtained. The obtained compacted samples were thendried at 110 ◦C for 2 h.

Thermal linear shrinkage (TLS) and coefficient of thermalexpansion of compacted and sintered samples, respectively,were measured using a dilatometer (Netzsch dilatometerModel DIL 402PC) at 10 ◦C min−1 in air, using alumina as ref-erence material, for rectangular samples of 20-mm length,5-mm width and 5 mm thickness. Three typical and differ-ent illite clays were also tested in the same condition andthe results were compared. The crystallization temperatureof the glass powder and mixtures were measured using dif-ferential thermal analysis (DTA) (Netzsch, STA EP 409) inair at a heating rate of 10 ◦C min−1 using powdered speci-mens of about 30 mg in an alumina sample holder with anempty alumina crucible as reference material. The theoret-ical density (�t) of the sintered samples was measured byHe-picnometry and the apparent density (�ap) was measuredby the Archimedes principle in water immersion at 20 ◦C. Therelative density (�r) was calculated from the relation betweenthe theoretical density and the apparent density, accordingto the expression: �r = �ap/�t. After sintering, samples were

transversally cut, grounded and polished with 1 �m aluminapaste and then etched in 2% HF for 25 s. Subsequently, all sam-ples were coated with a thin Au film for scanning electronmicroscopy (SEM) observations (Model Philips XL-30). To inves-

g t e c h n o l o g y 2 0 9 ( 2 0 0 9 ) 1134–1142

Fig. 2 – Thermal linear shrinkage curves of LZSAglass–ceramic (�), and composites (� ), 10 wt.% ZrSiO4; (� ),

1136 j o u r n a l o f m a t e r i a l s p r o c e s s i n

tigate the crystalline phases formed during heat-treatments,powdered samples were analyzed with a Philips PW 3710 X-ray (Cu K�) powder diffractometer (XRD). Bending strength (�f)of green and sintered samples was performed in an EMIC testmachine (Model DL 2000) according to ISO 10545-4, which con-sisted in a three-point test on six samples with dimensionsof 100 mm × 50 mm × 10 mm at a load rate of 1 MPa s−1. Deepabrasion (Da) of sintered samples was performed (Gabrielli testmachine Model CAP) according to ISO 10545-6. Deep abrasionwas determined through the measure of the volume loss ofmaterial after testing, which consisted in an iron disc (360 A)with diameter of 200 mm and thickness of 10 mm rotatingover the sample with 75 rpm at 150 revolutions, using Al2O3

powder as abrasive media. Chemical resistance (CR) was eval-uated according to JIS procedures, which consisted in theweight loss of the test piece (80 mm × 25 mm × 5 mm) after650 h immersed in a 1% H2SO4 and NaOH solutions, respec-tively at 20 ◦C. In order to determine the hardness of theobtained materials, samples were mounted in epoxy resinand surfaces were ground smooth, and then polished with1 �m alumina paste. Subsequently, microhardness measure-ments were performed with a Vickers automatic hardnesstester equipped with a diamond Vickers indenter at a load of4.903 N. A total time of 15 s was used for each indentation. Eachvalue of hardness is the average of 10 measurements with therespective standard deviation.

3. Results and discussion

3.1. Preliminary study

Differential thermal analysis curves are shown in Fig. 1.According to this curves the crystallization temperature ofthe investigated compositions takes place at approximately780 ◦C. Additionally, as the reinforcing crystalline phase(zircon–ZrSiO4) fraction increased, the crystallization peak

intensity slightly decreased. This occurred due to the decreaseof the glass–ceramic matrix fraction with respect to the rein-forcing phase, which means that lower crystal fractions wereformed from the glass–ceramic matrix. Moreover, it can be said

Fig. 1 – Differential thermal analysis curves of LZSAglass–ceramic (A), and composites (B), 10 wt.% ZrSiO4; (C),20 wt.% ZrSiO4; (D), 30 wt.% ZrSiO4; (E), 40 wt.% ZrSiO4; (F),50 wt.% ZrSiO4; (G), 60 wt.% ZrSiO4; (H), 70 wt.% ZrSiO4.

20 wt.% ZrSiO4; (×), 30 wt.% ZrSiO4; (o), 40 wt.% ZrSiO4; (�),50 wt.% ZrSiO4; (+), 60 wt.% ZrSiO4; (� ), 70 wt.% ZrSiO4.

that the reinforcing crystalline phases did not interfered in thecrystallization peak position.

On the other hand, according to the thermal linear shrink-age curve (Fig. 2), densification is apparently affected by zirconaddition, i.e., as the zircon was added thermal shrinkagedecreased and the shrinkage rate for all studied compositionstends to zero as the temperature increases. This behavior isrelated to the reduction of the formed viscous liquid phase aszircon was added.

Fig. 3 shows SEM micrographs of the Li2O–ZrO2–SiO2–Al2O3

glass–ceramic and composites with different zircon contentssintered at selected temperatures for 10 min. Some poros-ity is observed in all samples. In Fig. 3(d) two phases arepresent: the dark region represents the glass–ceramics matrix,while the light region is related to the zircon particles, asconfirmed by EDX analysis. The ZrSiO4 grains with sizes vary-ing from 1 to 5 �m are homogeneously distributed over theglassy matrix. As expected, the observed fraction of the zir-con phase increase from Fig. 3(b)–(g). The microstructure ofG composition is shown in Fig. 3(g). The pores in this sam-ple are more spherically shaped and uniformly distributed,which could contribute to improve the deep abrasion, as wellas the bending strength. By comparing G and E compositions,it is observed that the porosity of E composition is locatedat the interface between the reinforcing particle (zircon) andthe Li2O–ZrO2–SiO2–Al2O3 glass–ceramic matrix, which mayresult in the reduction of the mechanical properties.

Table 2 shows the total porosity of the sintered samples. Aslight increase of 5% on the porosity is observed with increas-ing ZrSiO4 content. The porosity tends to be constant (around9%) for compositions with ZrSiO4 content between 40% and70%. During sintering of the glassy matrix, gas (bubbles) mustbe released from the bulk material and their outward mobilityis limited by the presence of the ZrSiO4 crystals. Therefore, theamount of total porosity is proportional to the amount of the

ZrSiO4 crystals.

Fig. 4 shows the resistance to deep abrasion (Fig. 4(a)) andbending strength (Fig. 4(b)) in function of ZrSiO4 content inthe glassy matrix. The deep abrasion values are affected by

j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 2 0 9 ( 2 0 0 9 ) 1134–1142 1137

Fig. 3 – SEM micrographs of LZSA glass–ceramic (a) and composites (b) 10 wt.% ZrSiO ; (c) 20 wt.% ZrSiO ; (d) 30 wt.% ZrSiO ;( 0 wtr

tsno

e) 40 wt.% ZrSiO4; (f), 50 wt.% ZrSiO4; (g) 60 wt.% ZrSiO4; (h) 7espectively. Etching: 2% HF, 25 s.

he zircon particle additions decreasing from A to B compo-itions and increasing from B to D compositions. However,o explicit tendency is visible. Abrasion resistance does notnly depend on the hardness of the reinforcement particles,

4 4 4

.% ZrSiO4 sintered at selected temperatures for 10 min,

but on their shape and distribution over the matrix and onthe particle/matrix interaction (Montedo et al., 2004). It canbe also noticed that the bending strength increased from Ato B compositions and decreased from B to C compositions.

1138 j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 2 0 9 ( 2 0 0 9 ) 1134–1142

Table 2 – Porosity of the prepared compositions aftersintering at the respective temperatures (T) for 10 min inair

Composition T (◦C) Porosity (%)

A 725 5.5 ± 0.4B 950 6.4 ± 0.4C 950 7.2 ± 0.2D 950 8.5 ± 0.2E 1100 10.2 ± 0.3F 1125 10.7 ± 0.2

G 1200 8.8 ± 0.2H 1300 9.2 ± 0.2

These deep abrasion and bending strength variations resultedfrom compositional and sintering temperature differences.By comparing B and E compositions, for example, it can beobserved that the deep abrasion resistance decreased evenwith higher zircon content. However, the composition with10% ZrSiO4 (B) exhibits lower porosity than E composition (40%

ZrSiO4), which indicates a higher degree of densification andtherefore, a more homogeneous surface, resulting in lowerabrasion. This occurred because Li2O–ZrO2–SiO2–Al2O3 glassyphase content in B composition is higher, which yields as

Fig. 4 – Resistance to deep abrasion (a) and bendingstrength (b) in function of ZrSiO4 content in the glassymatrix.

Fig. 5 – Microhardness of the composites in function ofZrSiO4 content in the glassy matrix.

consequence higher amount of viscous liquid phase at the sin-tering temperature. Although there is a relative scatter on thedata, the resistance to deep abrasion tends to decrease withincreasing ZrSiO4 fraction (Fig. 4(a)). This is probably due to thehomogeneity of the dispersion of ZrSiO4 crystals. Additionallythe presence of finer grains and lower porosity of G compo-sition improve the interaction between the matrix and thereinforcement, which resulted in a higher strength (Fig. 4(b)).

Microhardness values also increase as the zircon contentincreases (Fig. 5) and it becomes more evident at ZrSiO4 con-tent higher than 60%. This can be explained by the probabilityof the indenter to reach a hard phase, which increases with theamount of this phase. H composition showed the highest value(9.4 GPa) due to the higher zircon content of H composition andalso due to its higher sintering temperature. The more signif-icant results of deep abrasion and bending strength amongthe studied compositions are related to the G compositionsintered at 1200 ◦C.

3.2. Advanced study

Based on the results above and considering some previousadjustments in the water content and plasticity and also theextrusion die design to eliminate extrusion defects, G compo-sition was selected for further study.

In order to obtain composite compacts by extrusion60 wt.% zircon and 7 wt.% bentonite were added to theLi2O–ZrO2–SiO2–Al2O3 frit glass powder and mixed andhumidified with water (23 wt.%). Subsequently, the ceramicwas stored for 12 h to moisture homogenization and thenextruded, following the procedure previously described. Fordrying studies extruded samples (G composition and illiteclays named as A–C) were obtained so that water contentwas determined and dry linear shrinkage, apparent densityand bending strength were measured as shown in Table 3.Clay plasticity increased from A to C. The amount of water

added to the G composition for extrusion was 23 wt.%. How-ever, the moisture content in the G extruded compositionwas 20.5 wt.%, which indicates that a 2.5 wt.% water contentreduction occurred after extrusion. The difference between

j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 2 0 9 ( 2 0 0 9 ) 1134–1142 1139

Table 3 – Moisture content, linear shrinkage, apparent density (�ap) and bending strength (�f) for G composition and A–Cillite clays in the dry state

Composition Moisture (%) Shrinkage (%) �ap (g/cm3) �f (MPa)

G 20.5 1.9 2.12 11.17.1 2.00 8.87.7 1.97 13.79.0 1.85 16.8

tttadtfs

psdfs(FFst

Fcs

Fig. 7 – Thermal linear shrinkage curve of G composition.

Clay A 24.1Clay B 21.9Clay C 26.1

he water content before and after extrusion can be associatedo the vacuum in the extruder chamber and also to tempera-ure variation during the extrusion process. Vacuum promotesir and water elimination from the ceramic body improvingensification. The chamber temperature usually rises by fric-ion between the ceramic body and mechanical parts. Otheractor that contributed for moisture variation in the G compo-ition was its reduced plasticity.

Fig. 6 shows the dry linear shrinkage for G and clay com-ositions. It can be observed in Fig. 6(a) that there is a dryhrinkage difference between G and clay compositions. Thisifference can be explained since G composition was obtainedrom raw materials without any plasticity even if by compari-on with clays that show a natural plasticity, a reduced binderbentonite) amount was used. For this reason, according to

ig. 6(b), in which the dry shrinkage rate (derivative curves ofig. 6(a)) is displayed, drying process is faster for G compo-ition than clays, which indicates a better drying process andherefore, a reduction of the defect generation risk in this step.

ig. 6 – Shrinkage behavior of G composition and A–C illitelays; (a) linear shrinkage curves and (b) derivative linearhrinkage curves.

Fig. 8 – Relative density vs. temperature of G composition.

The extruded G composition after drying process exhibitsan average dry apparent density of 2.12 g cm−3, which is higherthan that determined for the clays (Table 3). This higherdensity of the G composition is due to the zircon phase(4.6 g cm−3). The relative density was about 64% for G com-position, which indicates that raw materials and processingparameters were well adjusted since relative density of uni-axially pressed compacts corresponds to 50–55%. The averagedry bending strength for G composition was 11.1 MPa, whichis lower than those found for B and C clay samples but higherthan that found for A clay sample. These results indicate thatthe addition of 7 wt.% bentonite was sufficient to obtain driedextruded materials with suitable strength for finishing opera-tions in the production line.

Figs. 7 and 8 show the thermal linear shrinkage and rela-tive density curves, respectively. Compared to the results of G

composition without bentonite with 7 wt.% bentonite it can benoticed that the thermal linear shrinkage was reduced from16% to 13.5% and in the same way the temperature relatedto the maximum thermal linear shrinkage decreased from

1140 j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 2 0 9 ( 2 0 0 9 ) 1134–1142

Table 4 – Measured properties of the G composition sintered at 1150 ◦C for 10 min compared with some referencematerials

Proprieties G LZS LZSA M GR NP

�ap (g cm3) 3.04 2.65 2.52 2.69 2.64 2.70Total porosity (%) 5.5 ± 0.2 8.8 5.5 ± 0.4 – – –�f (MPa) 220 ± 10 163 ± 9 60 ± 8 14 ± 4 27 ± 5 50HV500 g (GPa) 7.17 ± 0.4 6.05 ± 0.4a 6.4 ± 0.2 – – –Deep abrasion (mm3) 51 ± 3 35 ± 2 134 ± 5 650 ± 15 142 ± 1 126 ± 1CTE (× 106 ◦C−1) 3.9 8.0–10.0 7.1 19.0 19.8 –

Chemical resistanceTo acids (wt.%) <0.01 <0.01 – 5.74 0.15 0.08To alkalis (wt.%) <0.01 <0.01 – 8.99 0.03 0.05

: neo

LZS and LSZA are sintered glass–ceramics; M: marble; GR: granite; NPa HV200 g.

1200 to 1150 ◦C. This reduction can be associated to bentonitesince it exhibits a relatively high content of alkali and otherlow temperature oxides that contribute to the increase of theamount of viscous liquid phase. In fact, according to Fig. 8relative density at 1150 ◦C was 94.5%, which is higher thanthat for G composition without bentonite addition. Moreover,Fig. 8 shows that densification started at about 640 ◦C and itsrate was reduced at 700 ◦C, probably due to the crystallization

process. According to Montedo et al. (2004) crystallization inthe Li2O–ZrO2–SiO2–Al2O3 system is surface type and it wasexpected since fine and high specific surface area powderswere used. Furthermore, when crystallization starts the vis-

Fig. 9 – SEM micrographs of G composition sintered at 1125 (a), 160 min. Not etched.

paries (sintered glass–ceramic).

cosity of the system increases, which consequently causes adecrease of the densification rate, since glass sintering occursby viscous flow (Shyu and Lee, 1995). After crystallization ther-mal linear shrinkage starts to increase (850 ◦C) by reduction ofthe glassy phase viscosity as the temperature increased. Thisnew shrinkage occurred at 1150 ◦C. A temperature increase,i.e., from 1150 to 1250 ◦C resulted in a material expansioncaused by the melting of the glass–ceramic matrix.

Table 4 shows the properties of the studied G compositesintered at 1150 ◦C for 10 min compared with some referencematerials. As the holding time increased from 10 to 60 minat 1125 ◦C the relative density and the deep abrasion resis-

150 (b) and 1175 ◦C (c) for 10 min and at 1175 ◦C (d) for

j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 2 0 9 ( 2 0 0 9 ) 1134–1142 1141

Fig. 10 – X-ray diffractions patterns for G compositionsamples sintered at 1125 ◦C for 10 min (a), 20 min (b), 30 min(l

tfrbbtib

acmtTm

tALaccsfimwtmFm

papiip(o

ee

Fig. 11 – X-ray diffractions patterns for G compositionsamples sintered at 1125 ◦C (a), 1150 ◦C (b) and 1175 ◦C (c)

c) and 60 min (d), respectively. E, �-spodumene; Z, ZrSiO4,ithium metassilicate; ZO, ZrO2.

ance tend gradually to increase from 93.6% to 94.1% androm 65 to 56 mm3, respectively. The small increase in theelative density promoted a significant improvement on theending strength (from 153 to 220 MPa), probably becausey increasing the holding time at the sintering temperaturehe viscosity of the glass–ceramic matrix decreased provid-ng subsequent densification. As expected, deep abrasion andending strength also decreased.

Fig. 9 shows micrographs of samples sintered at 1125, 1150nd 1175 ◦C for 10 min and at 1175 ◦C for 60 min. The zir-on particles exhibit good interaction with the glass–ceramicatrix, i.e., zircon particles were well wetted by the matrix so

hat the formed interface exhibited low noticeable porosity.he particles are homogeneously distributed over the glassyatrix for all observed samples.Fig. 10 shows X-ray diffraction patterns of G composi-

ion samples sintered at 1125 ◦C for 10, 20, 30 and 60 min.ccording to Montedo et al. (2004) after crystallization thei2O–ZrO2–SiO2–Al2O3 glass–ceramic frit shows preponder-ntly �-spodumene, zircon, lithium dissilicate and �-quartzrystalline phases. The first formed crystalline phases (Gomposition), after heat-treatments, were zircon and �-podumene. The formation of these crystalline phases, withne grains and uniformly distributed in the glass–ceramicatrix, contributes to improve the mechanical strength asell as to decrease the coefficient of thermal expansion. Addi-

ionally, other crystalline phases, in small amounts, as lithiumetassilicate and ZrO2 were detected. It can be verified from

ig. 1 that holding times at 1125 ◦C did not influence the for-ation of other crystalline phases.Fig. 11 shows X-ray diffraction patterns related to G com-

osition samples sintered at 1125, 1150 and 1175 ◦C for 10 minnd at 1175 ◦C for 60 min, respectively. With increasing tem-erature from 1125 to 1150 and 1175 ◦C the peak diffraction

ntensities related to the �-spodumene crystalline phasencreased, while those associated to the zircon crystallinehase decreased. With the increase on the holding time

Fig. 11(d)) no significant variation in the X-ray diffractions was

bserved.

Thus, to obtain G composition with optimized prop-rties with high performance for a given application thextruded compacts must be sintered at 1150 ◦C for 10 min.

for 10 min, respectively and at 1175 ◦C for 60 min (d). E,�-spodumene; Z, ZrSiO4; lithium metassilicate, ZO, ZrO2.

In fact, after sintering, the properties of the compositewith G composition are the same or even better comparedto LZS, Li2O–ZrO2–SiO2–Al2O3 and Neoparies glass–ceramicsand also with natural materials such as marble and gran-ite (Table 4). Porosity is lower than LZS glass–ceramic andbending strength and microhardness are higher than thosefound for the considered reference materials. Deep abra-sion resistance is very high if compared with the referencematerials. Only the LZS glass–ceramic showed a better perfor-mance. This behavior can be attributed to the relatively highhardness of zircon, which is the major crystalline phase inthe LZS system and in particular the composite toughnessand/or stiffness as a consequence of the formed interfaceand the reinforcing particles, which actuate as a barrier tothe crack propagation, increase the material critical stressand therefore, its mechanical resistance. The chemical resis-tance in comparison with the reference materials is superior.The relatively low coefficient of thermal expansion can beexplained since zircon but in particular �-spodumene havea low CTE (4 × 10−6 ◦C−1 and (0.4–2) × 10−6 ◦C−1, respectively)(Ryshkewitch and Richerson, 1985) against 11 × 10−6 ◦C−1 oflithium dissilicate-based glass–ceramics (McHale, 1991).

4. Conclusions

An optimized extruded 60 wt.% zircon (ZrSiO4) particulate-reinforced Li2O–ZrO2–SiO2–Al2O3 glass–ceramic matrix com-posite added with 7 wt.% bentonite as binder and humidifiedwith 23 wt.% water was obtained by sintering and crys-tallization. After crystallization the densification increasedby reduction of the glassy phase viscosity as the temper-ature increased. A further temperature increase resultedin an expansion of the material caused by the melting ofthe glass–ceramic matrix. On heating, the composite com-pacts first crystallize into zircon and �-spodumene andthen into lithium metassilicate and ZrO2. The microstruc-ture consisted of fine crystals uniformly distributed and

randomly oriented through out the glassy phase as well asa residual porosity. The sintered composite exhibited deepabrasion resistance and bending strength of 51 mm3 and220 MPa, respectively. Its chemical resistance was higher than

g t e

r

1142 j o u r n a l o f m a t e r i a l s p r o c e s s i n

the tested reference materials and its relatively low coef-ficient of thermal expansion may lead to a good thermalshock resistance. Extrusion is a potential candidate to pro-duce sintered glass–ceramics matrix composites for severalapplications.

Acknowledgements

The authors are grateful to CAPES and CNPq/Brazil for fundingthis work.

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