Sintered Glass-Ceramics and Glass-Ceramic Matrix Composites from CRT Panel Glass

Journal

J. Am. Ceram. Soc., 88 [7] 1886–1891 (2005) DOI: 10.1111/j.1551-2916.2005.00380.x

Sintered Glass–Ceramics and Glass–Ceramic Matrix Composites from CRT Panel Glass E. Bernardow Dipartimento di Ingegneria Meccanica, settore Materiali, Universita` di Padova, 35131 Padova, Italy

F. Andreola, L. Barbieri, and I. Lancellotti Dipartimento di Ingegneria dei Materiali e dell’Ambiente, Universita` di Modena e Reggio Emilia, 41100 Modena, Italy

erties. The parent glass, as a consequence, needs to be refined at high temperatures, in order to remove the eventual gas bubbles from chemical reactions upon vitrification. This problem is particularly pressing with waste glasses, since heavy metals may cause a dark coloration of the glass, resulting in a low thermal conductivity (and long fining times). In addition, long treatments at high temperatures may cause heavy metals to vaporize from the glass melt, thus compromising the sealing of hazardous pollutants. Sintered glass–ceramics, established since the 1960s, have the potential for solving the above mentioned problems. The parent glass is formed, then finely ground; a short time vitrification treatment is allowed, with no fining. Sintering of finely powdered glass occurs together with crystallization.12 Such phenomenon depends on the surface crystallization of glass: as reported in the literature12–15 free glass surfaces are preferred sites for devitrification. Catalysts are not needed because of the enhanced nucleating activity of the high surface area of small grains.12 Commercial examples of this approach are the wollastonitebased ‘‘Neoparies’’, developed in Japan since 1970s.4,5,16 Other sintered glass–ceramics, from cheaper and more accessible raw materials has been developed, in Bulgaria17 and in Italy,18–20 since the early 1990s. An unique advantage of sintering may be found in the incorporation of a reinforcing phase, thus leading to glass–ceramic matrix composites in a very simple and cost effective way. It should be noted that glass–ceramic matrix composites are generally manufactured by applying a nucleation/crystal growth treatment to previously obtained glass matrix composites.21–26 A typical reinforcement for sintered glasses is represented by low-cost Al2O3 platelets, commercially available as abrasive material for polishing applications. The improvement in mechanical properties is mainly related to crack deflection, caused by the introduction of residual stresses in the matrix, because of the thermal expansion mismatch with the reinforcing phase. Al2O3 platelets have been used successfully to reinforce hot-pressed glass matrix composites for the last 10 years.27,28 Some work has been carried out on the manufacturing of analogue composites by cold-pressing and pressure-less viscous flow sintering,29–31 yielding a cost effective processing route. A few studies, however, have been dedicated to Al2O3 platelet reinforced glass– ceramic matrix composites.32,33 In this paper, we describe the production of sintered glass– ceramics and glass–ceramic matrix composites based on glass from dismantled cathode ray tubes (CRTs). In a CRT, different types of heavy metal containing glasses, lead- or barium-based, are employed.31,34,35 The usage of heavy metal oxides in the chemical formulation of CRT glasses is essential for the absorption of UV- and X-radiation, escaping from the electron gun. The recent Decision of the European Communities 200/532/ EC, in agreement with the United States Environmental Protection Agency (EPA), classifies waste of electrical and electronic equipment (WEEE) as hazardous waste. WEEE are constituted

Sintering with simultaneous crystallization of powdered glass represents an interesting processing route for glass–ceramics, especially originating from wastes. Highly dense glass–ceramic samples may be obtained from a simple and short treatment at a relatively low temperature. In addition, glass–ceramic matrix composites may be obtained by mixing glass with suitable reinforcements. In this work sintered nepheline glass–ceramics, based on panel glass from cathode ray tubes, are illustrated. A limited addition of Al2O3 platelets caused a significant improvement in the mechanical properties (elastic modulus, bending strength, microhardness, fracture toughness), already remarkable for the un-reinforced glass–ceramic, compared with traditional nepheline glass–ceramics.

I. Introduction

T

vitrification treatment, since the 1970s, has been found to be a promising way to seal and make inert a number of industrial wastes. These wastes may be hazardous because of the presence, for example, of heavy metals or organic pollutants, so that they can be stored only in selected and controlled landfills.1 After being dissolved in a glass, however, the wastes can be easily stored, since the obtained ‘‘waste glass’’ generally possesses a high chemical stability. The enormous quantity of potentially treated wastes and daily recovered recycled glasses leads to a reconsideration of waste glasses as a new kind of raw material for engineering applications. An undoubtedly well developed and widespread field of application for waste glasses is that of glass–ceramics. The first and most important example was certainly that of Russian Slagsitalls, developed, as early as the 1960s, by using several metallurgy slags and wastes from mining and chemical industries, and intended for building applications.2–5 It should be noted that the large market for building materials allows large quantities of waste glasses to be absorbed. This approach is currently followed for a number of wastes6–11 and it may lead to materials with remarkable properties. Conventional glass–ceramic production follows a two-step thermal treatment of nucleation and crystal growth. Although widespread, this route presents some disadvantages. Firstly, the nucleation/crystal growth treatment is slow and expensive, so that some catalysts in the glass formulation (TiO2, Cr2O3) must be employed.4,5 Secondly, defects in glass articles, like pores, may remain in the glass–ceramic, causing poor mechanical propHE

V. M. Sglavo—contributing editor

Manuscript No. 11335. Received September 17, 2004; approved February 3, 2005. w Author to whom correspondence should be addressed. e-mail: enrico.bernardo@ unipd.it

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II. Experimental Procedure The starting glass for the production of sintered glass–ceramics and glass–ceramic composites was obtained by mixing panel glass with alumina and dolomite (CaMg(CO3)2), inexpensive materials from the ceramic industry. A preliminary study on glass–ceramics from panel glass40 demonstrated that the optimum formulation consisted of 50% (by weight) panel glass, 20% alumina, and 30% dolomite. The chemical composition of the raw materials is shown in Table I. The glass-forming mixture was melt at 15001C for about 5 h, in an electric furnace. The molten glass was poured into cold water, yielding a glass frit. The frit was dried, dry ball milled and sized to grains o37 mm. The powdered glass was subjected to differential thermal analysis (DSC g404, Netzsch Gera¨tebau GmbH, Selb, Germany), with a heating rate of 101C/min, in order to identify glass transition and crystallisation temperatures. a-alumina monocrystals (platelets) were chosen as the reinforcement (Microabrasives Co., Westfield, MA). The platelets were hexagonal-shaped, with major axes between 5 and 10 mm and axial ratio (thickness/average diameter)
0.2. The density of a-alumina was considered to be 3.99 g/cm3.41 Alumina platelets in 5%, 10%, 15% concentrations by volume were added to the frit, and mixed in a dry ball mill for 1 h. Also powders of pure frit were considered. The powders were

Table I. Chemical Composition (in wt%) of the Raw Materials for the Preparation of the Starting Glass

SiO2 Al2O3 CaO MgO Na2O K2O Fe2O3 TiO2 ZrO2 ZnO BaO PbO SrO NiO CoO MnO P2O5 w

Panel

Alumina

Dolomite

Glassw

57.9 3.8 – – 12.9 7.3 0.2 0.4 1.4 0.6 8.0 0.02 8.5 0.03 0.01 – –

– 99.5 – – – – – – – – – – – – – – –

0.1 0.1 30.8 21.6 – – 0.02 – – – 0.02 – – – – 0.01 0.03

33.5 25.4 10.8 7.6 7.5 4.2 0.1 0.2 0.8 0.4 4.6 0.01 4.9 0.01 – – –

Results from the chemical analysis of the obtained glass.

uniaxially pressed in a steel die at room temperature, by using an hydraulic press operating at 40 MPa, without any binder. The obtained green tiles were sintered in air at 9301C for 3 h, with a heating rate of 101C/min, leading to dense samples with a white coloration. The bulk density of the sintered compacts was measured by the Archimedes’ principle. At least ten fragments were analysed for each sample. The true density of the composite materials was evaluated by means of a gas pycnometer (Micromeritics AccuPyc 1330, Norcross, GA). The theoretical densities were calculated from the density of the constituents by applying the rule of mixtures. The thermal expansion coefficient of the matrix was determined by a dilatometric analysis (Netzsch Gera¨tebau) from a fragment of un-reinforced sintered compact. Beam samples of about 3 mm  2 mm  30 mm, for bending strength determinations were cut from larger sintered tiles. All samples were carefully polished to a 6 mm finish, by using abrasive papers and diamond paste. The edges of the bars were bevelled by using fine abrasive papers and diamond paste. Four point bending tests (24 mm outer span, 8 mm inner span) were performed by using an Instron 1121 UTS (Instron, Danvers, MA), with a crosshead speed of 0.2 mm/min. Each data point represents the average of at least 10 individual tests. The Young’s modulus (E) was measured by means of non-destructive resonance frequency testing (GrindoSonic Mk5, Leuven, Belgium). Polished samples were employed for Vickers indentation tests, which yielded the hardness (Hv), at low load (5 N), and the indentation fracture toughness (KIC), at high load (20 N), of the investigated materials. The fracture toughness was calculated by using the well known equation of Anstis et al.42 starting from the measured length of the cracks emanating from the corners of the Vickers indents, as follows: KIC ¼ xðE=Hv Þ0:5 ðP=c1:5 Þ

(1)

where P is the applied load, c is the length of the emanated cracks and x is a calibration factor (x 5 0.01670.004). The polished and fracture surfaces of sintered samples were characterized by scanning electron microscopy (SEM; Philips XL 40 SEM and Philips XL 30 ESEM, Eindhoven, The Netherlands). Powdered samples were investigated by X-ray diffraction (Philips PW 3710), employing CuKa radiation (0.15418 nm).

III. Results and Discussion (1) Crystallization and Densification The differential thermal analysis plot of the parent glass is shown in Fig. 1. The transition temperature Tg at 7091C is 25 Tmelting=1138°C

20 15 10 DTA (µV)

of about 80% by TV and computers, which contain CRTs made for the 85 wt% of glass. The waste stream of electrical and electronic equipment has been identified as one of the fastest growing waste streams in the European Union (EU). The EU produces around 7.5 million tons of electrical waste every year, among which the amount of end of life CRTs in Western Europe is in the range of 300 kton/year.36 Today these glasses constitute the 4% of the municipal waste, and is assumed to increase by 16%–28% every 5 years.37 The investigation reported in this work was focused on panel glass, that corresponding to the front part. In this glass more expensive barium and strontium oxides are currently employed, instead of lead oxide, in order to prevent the darkening of the screens because of the precipitation of metallic clusters from easily reducible oxides (like PbO) induced by high-energy electrons.38 Panel glass from dismantled CRTs is actually a waste material, since it is employed in the manufacturing of new CRTs only in low percentages (15 wt%):39 in fact, only high-quality pristine glass, from selected raw materials and with no lead contaminations, is desirable.

Oxide

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endo

5

Tg=708.7°C

0 −5

−10 −15 −20 −25 200

Tcryst=934.6°C 400

600 800 1000 Temperature (°C)

1200

1400

Fig. 1. Differential thermal analysis plot of the parent glass.

Journal of the American Ceramic Society—Bernardo et al.

well-distinguished. Large exothermic and endothermic effects, which correspond, respectively, to crystallization and melting processes, are clearly evident at 9351 and 11381C. All the samples were sintered at a temperature of 9301C, practically identical to the crystallization temperature, as reported in previous experiences.33 The proposed temperature is significantly lower than those reported in most works on the simultaneous sintering and crystallization of glass,18–20 in which the processing temperatures were much higher than crystallization temperature. This choice was suggested by the fact that the glass was fine-grained (o37 mm), so that the sintering and nucleating activities were enhanced. Moreover, relatively low processing temperatures are desirable from an economic point of view. The density of the sintered glass–ceramic sample (that corresponding to glass powders without any reinforcement) was 2.7970.02 g/cm3. As illustrated by Fig. 2(a), the sample appeared dense, with a low residual porosity, the true density being 2.8070.01 g/cm3, thus confirming the efficiency of the abovedescribed approach. The obtained true density value constituted the approximate reference for the sintered glass–ceramic matrix in the calculation of the theoretical density of the Al2O3 platelet reinforced composites. The XRD analysis conducted on the matrix composition revealed the presence of nepheline ((Na,K)AlSiO4), as the main crystalline phase (Fig. 3). This phase may have various compositions resulting from the formation of solid solutions by substitution of ions such as K1, Ca11 and Ba11 and it is a stuffed derivative of tridymite. When glass compositions contain BaO, barium aluminumsilicate BaAl2Si2O8, i.e. celsian crystalline phase, can be separated.5 A confirmation of this behavior is represented by the effective development of celsian in the glass– ceramic studied, because of the content of BaO in the panel glass

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Intensity (a.u.)

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5

10 15 20 25 30 35 40 45 50 55 60 65 70 2θ (degree)

Fig. 3. XRD spectrum of un-reinforced sintered glass–ceramic.

(7.95 wt%). Other crystalline phases, such as akermanite (Ca2Mg(Si2O7)) and wollastonite (CaSiO3), derived from dolomite, and aluminumsilicate corresponding to Al0.5Si0.75O2.25, derived from alumina, were identified. The microstructure of un-reinforced and reinforced samples, resulting from SEM analysis, (Figs. 4 and 5) appeared slightly different, although no differences in the mineralogical characteristics were detected. In the un-reinforced material (Fig. 4) the three main crystalline phases identified by XRD analysis were evident, while the reinforced glass-ceramic showed a beehive-like structure with only two clearly visible different phases (Fig. 5).

Fig. 2. Scanning electron micrographs of the fracture surfaces of sintered products: (a) un-reinforced sintered glass–ceramic; (b) glass–ceramic matrix composite with 5 vol% Al2O3; (c) 10 vol% Al2O3; (d) 15 vol% Al2O3; the roughness of the fracture surfaces increases with increasing Al2O3 platelets volume fraction.

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As shown by Table II, the actual density of the glass–ceramic matrix composites was close to the predicted value. In practice, it was found that the introduction of non-sintering inclusions like Al2O3 platelets did not alter significantly the densification. Although some pores are visible in the micrographs of the composites (Fig. 2(b)–(d)), the deviation form full density was found to be much less remarkable than in other experiences reported in the literature for glass matrix composites30 and for other sintered glass–ceramics.33 Non–sintering inclusions, as reported by Boccaccini,43 in fact, may retard densification by dramatically increasing the viscosity of the system; the observed densification behavior is somewhat anomalous. One likely reason, to the authors’ opinion, may be the particular chemical composition of the glass employed, with a low content of glass formers, hence a low viscosity upon sintering. The XRD spectra of the sintered composites are reported in Fig. 6; there is no significant difference with the spectrum of the un-reinforced matrix except for the presence of the reinforcement (which corresponds, as previously mentioned, to a-Al2O3).

Fig. 4. Scanning electron micrographs of un-reinforced sintered glass– ceramic revealing different crystal phases with relative energy dispersive X-ray analysis spectra (1 5 acicular white crystals, 2 5 grey crystals, and 3 5 darker crystals).

The crystals evident in the un-reinforced sample were distinguished by means of EDS analysis. Acicular white crystals containing the heavier element barium and grey crystals enriched of Ca and Si, were revealed; these two types of crystals can be attributed to celsian and akermanite, respectively. Finally, darker crystals rich in Na, K, and Al, showing hexagonal prism shape typical of nepheline, were found. The microstructure showed the growth of celsian around the edge of each nepheline crystals suggesting that nepheline is not only the principal crystalline phase but also that formed before with respect to celsian. This behavior has been verified also with other crystalline phases such as kalsilite.4 All the crystals have very small dimensions, for example elongated white crystals are in the range of 3–7 mm, and dark nepheline crystals show width in the range of 1–3 mm and height of 2–5 mm detectable only for higher magnification. Further, grey crystals are with difficulty singularly measurable since, because of their lower dimensions, are very interlocked. This very fine microstructure shows that the thermal treatment ensured the formation of large amount of tiny crystals uniformly distributed throughout the bulk of the glass. This microstructure can positively influence the mechanical properties of the materials.

Fig. 5. Scanning electron micrograph of the 15% reinforced glass–ceramics.

(2) Mechanical Properties The mechanical properties of the investigated materials are summarized in Table II. It should be noted, firstly, that the un-reinforced sintered glass–ceramic exhibited notable characteristics. In particular, the bending strength, being 120713 MPa, represents a remarkable increase compared with analogue glass–ceramics from a nucleation/crystal growth treatment, which is generally much longer and performed at higher temperatures than the proposed sintering treatment (the bending strength of commercial nepheline-based glass–ceramics, as reported in the literature,4 varies from 58 to 91 MPa). The addition of Al2O3 platelets brought an overall improvement of the mechanical properties. The composites exhibited a gradual increase of Young’s modulus with increasing Al2O3 content; the measured moduli are in a good agreement with those predicted (from the elastic modulus of the matrix and the elastic modulus of alumina) by Eq. (2):44 ( " p 1
 1  Ecomp ¼ Em 1  A 9 1 þ ð1:99=BÞ Em =Ep  1 1
   3 1 þ ð1:68=BÞ Em =Ep  1 #) 1
 9 5 1 þ ð1:04=BÞ Em =Ep  1

(2)

where Ecomp, Em, Ep are, respectively, the moduli of the composite, the matrix and the inclusion phase (Em being 380 GPa).44 A and B are factors depending on the axial ratio and the orientation of the reinforcement. The equation refers to the hypothesis of full densification. The fact that the Young’s modulus of the composites, despite some residual porosity, agreed with that predicted for full densification, might be attributable to the slightly different phase distribution in the composites, compared with un-reinforced sintered glass–ceramics. The glass–ceramic composites exhibited remarkable bending strength values, particularly when referred to the limited amount of reinforcing phase. Even if some pores may be recognized especially with the highest volume fractions of Al2O3 platelets (as illustrated by Fig. 2), the composites reached an excellent value of 163714 MPa (with a 15 vol% Al2O3, corresponding to a 30% increase compared with the un-reinforced matrix). The reinforcing ability of the employed platelets, in practice, far exceeded the weakening effect of residual porosity. Finally, Al2O3 platelets had a notable effect on both Vickers’ microhardness and indentation fracture toughness. The Vickers’ microhardness exhibited an almost linear increase with increasing platelet content. The achieved hardness enhancement (Hv410 GPa) by platelets addition made glass matrix composites suitable for applications requiring good wear resistance.

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Journal of the American Ceramic Society—Bernardo et al. Table II. Density and Mechanical Properties of the Obtained Sintered Products Density (g/cm3)

Al2O3 vol. fraction%

0 5 10 15

Bulk density

True density

Young’s modulus (GPa) Predicted density

2.7970.02 2.8070.01 2.8470.01 2.8570.01 2.86 2.9070.02 2.9270.01 2.92 2.9570.04 2.9770.01 2.98

Predicted value

Measured value

105.0 109.1 113.2 117.5

111.9 116.1 122.4

Bending strength (MPa)

Hv (GPa)

KIC (MPa  m0.5)

120713 140720 143716 163714

6.670.3 7.870.4 9.170.3 10.470.3

1.570.1 1.670.1 1.870.1 1.970.1

nepheline glass–ceramics. The illustrated approach could be useful to the manufacturing of a number of similar sintered glass–ceramics and glass–ceramic matrix composites from several waste glasses.

α-Al2O3(corundum)

Intensity (a.u.)

15% Al O

Acknowledgments The authors would like to thank Dr. Sandro Hreglich of Stazione Sperimentale del Vetro (Murano, Venice, Italy) for supplying panel glass and for the help in the preparation of the glass–ceramic composition. E. B. acknowledges Dr. Claudio Furlan for experimental assistance during SEM investigation of the fracture surfaces.

10% Al O

5% Al O

References matrix

10

15

1

20

25

30

35

40

45

50

55

60

65

70

2θ (degree) Fig. 6. XRD spectra of the Al2O3 platelet reinforced glass–ceramic composites; the spectra differ from that of un-reinforced sintered glassceramic only for the peaks corresponding to the reinforcement (marked with the symbol).

Like in previous experiments on Al2O3 platelet reinforced glasses,30,31 although the volume fractions of the reinforcement were limited, the improvements in the fracture toughness were substantial. KIC reached, for 15 vol% Al2O3 platelets, 1.970.1 MPa  m0.5, with more than a 25% increase compared with the value for the matrix (1.570.1 MPa  m0.5). The improvement of the Young’s modulus could justify partially the improvement of fracture toughness. Applying Eq. (1), by considering only the improvement in the elastic modulus, thus neglecting any variation of microhardness and crack extension, leads to a limited improvement of fracture toughness (B8% of the total). An additional contribution to fracture toughness could be attributed to crack deflection, owing to the thermal expansion coefficient mismatch between the matrix (12.5  1061C1, in the range 01–3001C, with a full correspondence to that of nephelinebased glass-ceramics)4 and the alumina reinforcement (8.9  106 1C1).41 An indication of crack deflection may be found in the increasing roughness of the fracture surfaces of the studied composites with increasing Al2O3 platelet content (see Fig. 2).

IV. Conclusions A powder technology approach for the manufacturing of sintered glass–ceramic and alumina platelet reinforced glass–ceramic matrix composites from wastes was illustrated. It was demonstrated the feasibility of highly dense sintered nepheline-based glass–ceramic from a treatment of short duration and relatively low temperature (compared with traditional processing routes for glass–ceramics). The introduction of reinforcing Al2O3 platelets did not modify substantially the densification nor the crystallization of the matrix, but caused a significant improvement in the mechanical properties (elastic modulus, bending strength, Vickers’ microhardness, fracture toughness), which were already remarkable for the matrix itself, compared with traditional

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