Effect of ZnO content on the physical, mechanical and chemical properties of glass-ceramics in the CaO–SiO2–Al2O3 system

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Effect of ZnO content on the physical, mechanical and chemical properties of glass-ceramics in the CaO–SiO2–Al2O3 system Estefanía Montoya-Quesadaa,∗, Mónica A. Villaquirán-Caicedoa, Ruby Mejía de Gutiérreza, J. Muñoz-Saldañab a b

Composites Materials Group (GMC-CENM), Universidad del Valle, Cali, Colombia Centro de Investigación y de Estudios Avanzados del IPN, (CINVESTAV), Querétaro, Mexico

A R T I C LE I N FO

A B S T R A C T

Keywords: ZnO Glass-ceramics Crystallization Solid wastes Vitrification

The objective of this study was to evaluate the effect of ZnO content on the physical, mechanical and chemical properties of CaO–Al2O3–SiO2 (CAS) glass-ceramics produced from Colombian wastes, such as fly ash, granulated blast furnace slag and glass cullet. The CaO/SiO2 molar ratio of the mixtures was held constant (0.36). ZnO was added to the mixtures in proportions of 4, 7 and 10 wt%. The glass-ceramics were produced by the controlled crystallization of a parent glass. The values of crystallization temperature (Tp) show a fall up to 7 wt% and then shoots up with 10 wt% concentration of ZnO, but in general, ZnO addition lowers the temperature required for the formation of crystalline phases. In general, anorthite (CaAl2Si2O8) is the main phase observed in all heat treated samples, in addition to albite (Na(AlSi3O8)) and labradorite (Na0.45 Ca0.55 Al1.55 Si2.45 O8). The crystalline phases hardystonite (Ca2ZnSi2O7) and willemite (Zn2SiO4) were also identified in the samples with 7 and 10 wt% ZnO. The densities of the glass-ceramics were between 2658 and 2848 kg/m3, and it was found that ZnO helps to increase the density of glass-ceramics. The elastic modulus was in the 100–105 GPa range, the fracture toughness was between 0.45 and 0.64 MPa m1/2, and the Vickers microhardness was between 632 and 653 MPa. With regards to the durability, the weight loss of the glass-ceramics immersed in alkaline solution (NaOH) did not exceed 1.5 wt% after immersion for 6 h at 80 °C. The results of this study confirm that the vitrification process is a favorable option to utilize these industrial wastes.

1. Introduction Glass-ceramics are materials that are processed and formed from glass obtained at high temperatures (1300–1500 °C) and are subsequently converted into crystalline materials through a heat treatment [1]. Usually, a glass-ceramic is not completely crystalline since the microstructure is typically 50–95 vol% crystalline, while the remainder is residual glass [2]. These materials combine properties that are characteristic of conventional ceramics and glass-like materials into a single material [3]. The type of constituents, cooling rate and presence or absence of nucleating agents used to prepare the glass-ceramic will determine whether the material is amorphous and/or crystalline [1], and these parameters can be modified to produce materials that are either transparent or opaque and either have color or are colorless and that have a suitable composition and microstructure design. The most common glass-ceramic materials are based on SiO2–Al2O3 and contain oxide modifiers, such as LiO2, Na2O, K2O, CaO, ZnO and MgO [4]. Glass-ceramics have numerous applications, for example, they can be

∗

used as machinable engineering components, insulating coatings, architectural panels, telescope mirrors, bioactive glasses, etc. [4–7]. To produce glass materials, different varieties of rocks (e.g., aplite, granite, basalt, nepheline syenite, etc.), minerals (e.g., feldspar, kaolin, nepheline, talc, magnesite, petalite, etc.) and/or synthetic chemicals have been used as raw materials [8,9]. Because some mineral reserves are in short supply, new raw materials have been explored, and thus, the use of byproducts and industrial waste as raw materials has been researched [9], as they contribute to reducing the environmental impacts caused by exploiting nonrenewable natural resources [2,10,11]. According to the results published by Karamanov [12], when industrial waste is used as raw materials to produce glass-ceramics, the process is economically justified only if materials with commercial value are obtained. The preparation of glass-ceramic materials using industrial waste, specifically slags, was developed in the former Soviet Union by Kitaigorodski and Pavlushkin in the 1960s; some of these glass-ceramics have been used as either commercial products, such as Slagsitall (which

Corresponding author. E-mail address: [email protected] (E. Montoya-Quesada).

https://doi.org/10.1016/j.ceramint.2019.10.154 Received 12 August 2019; Received in revised form 27 September 2019; Accepted 16 October 2019 0272-8842/ © 2019 Elsevier Ltd and Techna Group S.r.l. All rights reserved.

Please cite this article as: Estefanía Montoya-Quesada, et al., Ceramics International, https://doi.org/10.1016/j.ceramint.2019.10.154

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are crystals of anorthite and wollastonite), Slagceram and Slagkyston, or preindustrial products, such as Silceram (derived from pyroxene devitrification) [12]. Additionally, the versatility of the process used to produce glass-ceramic materials from industrial waste has been reported in several studies; notable wastes used to produce glass-ceramics include slag [1,13–15], fly ash [6,16–20], glass waste [21], sugarcane bagasse waste [22], copper slag [15] and mixtures of different wastes [10,23–27]; these studies report that glass-ceramics produced from industrial waste have properties that are similar or even better than conventional glass-ceramics. However, the chemical composition must be strictly controlled when waste is used as raw materials because the mixtures must contain components that are typical of stable glasses, especially Si and Al oxides and nucleating agents that help the crystallization process. Thus, each residue must contribute an appropriate amount of vitrifying agents (e.g., SiO2 and AI2O3), modifying elements that promote fusion (e.g., Na2O and K2O) and stabilizing agents (e.g., CaO, MgO, ZnO and PbO), and these components produce glasses with appropriate characteristics; note that some of the stabilizing oxides may act structurally as vitrifying agents or as intermediate oxides [16,28,29]. This study investigates the use of a mixture of fly ash, granulated blast furnace slag (GBFS) and glass waste as potential raw materials to synthesize glass-ceramics in the ternary CaO–Al2O3–SiO2 (CAS) system. The incorporation of a fourth stabilizing oxide, ZnO, was evaluated, based on the hypothesis that a fourth oxide helps to lower the melting temperature and viscosity of the glassy phase without affecting the chemical properties of the material. ZnO was added to the materials to result in final mass percentages of 4, 7 and 10% ZnO. The effect of the ZnO content was evaluated by studying properties such as the crystallization temperature and workability of the molten glass, physical properties (e.g., density and porosity), mechanical properties and durability (i.e., the material stability in 10 vol% solutions of HNO3 and NaOH). Additionally, the microstructure of the obtained product was characterized using techniques such as X-ray diffraction (XRD), scanning electron microscopy (SEM) and Fourier-transform infrared spectroscopy (FTIR).

Table 1 Chemical composition (wt%) and density of the raw materials used for glass production.

SiO2 Al2O3 CaO Na2O Fe2O3 ZnO K2O CuO TiO2 NiO MgO P2O5 Loss on ignition at 1000 °C (%) Density (kg/m3)

FA

RV

GBFS

ZnO

62.13 26.31 1.27 0.27 4.88 0.01 0.81 0.01 1.19 0.01 0.25 1.18 1.02 2480

72.27 1.49 11.15 13.37 0.62 0.52 0.08 0.26 2488

38.21 15.88 40.80 0.21 1.87 0.40 0.51 1.32 2920

99.4 5610

GBFS and 10% by weight of RV with a molar ratio of CaO/SiO2 of 0.36 because It was the sample that showed the highest crystallization temperature, therefore, the effect of replacing 0, 4, 7 and 10 wt% of the mixture with ZnO powder was evaluated. The CaO/SiO2 molar ratio was taken into account, as reported by Yang et al. [15], the CaO/SiO2 ratios higher than 0.38 could increase the viscosity of the melt, due to the formation of Ca2SiO4, by the excess of calcium. The following procedure was followed to synthesize the four samples: i) The powders were mixed and dry-homogenized for 10 min. ii) The powder mixture was placed in an alumina crucible for subsequent heating and melting for 2 h at 1450 °C in a Nabertherm electric furnace. iii) The molten mixture was quickly poured into a water bath at room temperature to obtain a glass frit. iv) The resulting frit was ground in a mortar to achieve a particle size smaller than 75 μm (i.e., 200 mesh), later the powder was pressed into a KBr pellet and then characterized using FTIR via a PerkinElmer Spectrum 100 FTIR spectrometer. The crystallization temperature (Tp) was determined by differential scanning calorimetry (DSC) using an SDTQ 600 apparatus (TA Instruments); DSC measurements were performed using an alumina sample holder, and the air flow and heating rate were 100 ml/min and 10 °C/min, respectively.

2. Experimental procedures 2.1. Raw materials The solid wastes used as raw materials for the production of glassceramics were as follows: fly ash (FA) from a Lago Verde brick company boiler, granulated blast furnace slag (GBFS) from Acerías Paz de Río and glass cullet (RV). Zinc oxide powder (ZnO) with a 99.4% purity (J. T. Baker brand) was used as a stabilizing agent. The solid wastes were mechanically conditioned for 2 h in a ball mill with ceramic balls. The particle size distribution was determined using a laser particle size analyzer (MASTERSIZER 2000, Malvern, UK); the average particle size D [3,4] was 43.65 μm for the fly ash, 26.44 μm for the GBFS and 43.12 μm for the RV. The chemical composition of the raw materials, as shown in Table 1, was determined by X-ray fluorescence (XRF) using an AXIOS MAX (PANalytical, USA) spectrometer, and the powder density was calculated using a pycnometer method according to the standard ASTM C329-88 [30]. The XRF results show that the solid wastes used for glass production have significant amounts of network-forming oxides (e.g., SiO2 and Al2O3). Furthermore, GBFS contains a considerable percentage of CaO that acts as a stabilizing agent (40.8%wt). RV contains 13.4 wt% Na2O, which could act as a flux agent, and FA contains 4.9 wt % Fe2O3, which could act as a nucleating agent.

2.3. Preparation and characterization of glass-ceramics The above mentioned frit powder was mixed with 10 wt% H2O and then compacted and pressed at a pressure of 8–9 ton in a PIKE IR hydraulic press (Crush Technologies) until pellets with 13 mm diameters and 3 mm thicknesses were obtained. These pellets were heated from room temperature to the Tp at a rate of 10 °C/min in a Nabertherm electric furnace and maintained at the desired temperature for 2 h, and then, the samples were cooled to room temperature. Fig. 1 shows the glass-ceramics obtained with different ZnO ratios. Samples A and B (Fig. 1A and B), which correspond to 0 and 4 wt% ZnO, respectively, are lighter in color compared to samples C and D (Fig. 1C and D), which correspond to 7 and 10 wt% ZnO. This is attributed to a higher concentration of iron, which tends to give the samples a yellow appearance [15]. It is noted, that the proportion of ZnO was added as replacement of the original mixture (0% ZnO). The crystallized samples were characterized by XRD using an X'Pert MRD PANalytical diffractometer with Cu Kα1 radiation generated at 20 mA and 40 kV, and the specimens were scanned from 10° to 60° 2θ with a step size and time per step of 0.02° and 4.0 s, respectively. Microstructures of sintered samples were determined from backscattered electrons of gold-coated surfaces by using a JEOL JSM6490LV Scanning Electron Microscope (SEM). Additionally, the mechanical properties were evaluated by Vickers hardness (Hv) testing using a 1 N load, and the elastic modulus (E) was obtained by

2.2. Precursor glass preparation and characterization As a result of preliminary experiments, a ternary mixture was selected with a composition of 50% by weight of FA, 40% by weight of 2

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Fig. 1. Glass-ceramics obtained by maintaining the samples at the Tp for 2 h with different ZnO concentrations: A) VC-0%ZnO, B) VC-4%ZnO, C) VC7%ZnO and D) VC-10%ZnO.

instrumented nanoindentation testing (Hysitron TI950 Triboindenter, Bruker) using a Berkovich diamond tip. Each consecutive indentation was separated by 4.5 μm to avoid any residual stress interference from adjacent indentations. Force-displacement curves were used to evaluate the elastic moduli. The E of each indentation was calculated using equation (1), which is based on the standard Oliver and Pharr method [31]: −1

ν2 1 E= (1 − ν 2) ⎡ − i⎤ ⎢ Er E i⎥ ⎣ ⎦

(1)

where ν corresponds to Poisson's ratio, which is estimated to be 0.245 for a CaO–SiO2–Al2O3 glass-ceramic system [32], Er (GPa) is the reduced elastic modulus of the glass-ceramics (102.5 ± 9.0, 100.3 ± 8.3 and 97.5 ± 5.7 GPa for VC-0%ZnO, VC-7%ZnO and VC10%ZnO, respectively), and νi (0.07) and Ei (1141 GPa) are values corresponding to the Berkovich nanoindenter. The applied load was decreased from 10.000 μN to 160 μN. The H values were calculated according to the Oliver and Pharr method and therefore are the averages obtained for each group of indentations. The fracture toughness (KiC) was calculated using equations (2) and (3) [33,34], which are based on the measured length of the cracks emanating from the corners of the Vickers notches: 2/5 −3/2 ⎡ K c ∅ ⎤ ⎛ H ⎞ = 0.129 ⎛ c ⎞ ⎢ H1/2 ⎥ ⎝ E∅ ⎠ a ⎝ ⎠ ⎣ a ⎦

K c = 0.016

E P ⎞ ⎛ H ⎝ c3/2 ⎠

Fig. 2. DSC curves obtained for the frits, which was used to determine the Tp.

5 mol% of ZnO (∼4% wt ZnO). When ZnO acts as a network former, this oxide enters the glass network as structural units of ZnO4 and favors the formation of junction bonds, while when ZnO acts as a network modifier, glass restructuring occurs due to the formation of non-bridge oxygen and this leads to a decrease in the connectivity of the glass network [38]. Note that the increase in the ZnO concentration decreased the viscosity of the glass, which was verified during melting and casting; the DSC results also show that the peak corresponding to the Tp becomes more pronounced and defined. The Tp chosen for the heat treatments were 957 °C (for V-4%ZnO), 925 °C (for V-7%ZnO) and 932 °C (for V-10%ZnO). Fig. 3 shows the FTIR transmittance spectra of the glass samples before and after crystallization. There are significant differences between the FTIR spectra of the glasses compared to those of the glassceramics (Fig. 3). The spectra of all the glasses have the same three transmittance bands (i.e., bands 1, 2, and 3) in the 400–1000 cm−1 region (Fig. 3A). The wider band 1, centered at 991 cm−1, is attributed to the stretching vibration of Si–O–Si; band 2, located at 705 cm−1, is attributed to the vibrations of the silicon and oxygen tetrahedral groups in the silicate, and band 3 located at 484 cm−1, corresponds to the bending vibration modes of Si–O–M, where M = Si or Al [15]. Addtionally, the FTIR spectra of the glass and glass-ceramics samples showed a band between 3000-3700cm-1 region that corresponds to the OH- bond stress vibration of the water that is physically adsorbed or in free molecular form in the samples, which is corroborated by the endothermic peak observed in the DSC curves at T < 100°C [39]. As shown in Fig. 3B shows that new bands develop after the heat treatment, due to the crystallization process; a large band appears in the spectra obtained after the heat treatment and comprises two distinct signals (denoted as signals 1 and 2); signal 1 occurs at 1076 cm−1 and is attributed to the vibration of the Si–O–Si bond, and signal 2 occurs at 947 cm−1 and is attributed to the asymmetric stretching vibration of O–Si–O [40]. The band at 705 cm−1 shifts to wavenumbers higher than 747 cm−1 (i.e., signal 3) after the crystallization process, and this shift is associated with the stretching vibrations of the Si–Si bond, which

(2)

(3)

where a is the mean diagonal indentation length, 2c is the radial crack length and ϕ is a constant equal to 3. For all the measurements, c/a ≥2.5 [35]. The chemical resistance was calculated as the difference between the initial mass and the final mass of the materials after they were immersed in alkaline 10 vol% NaOH and acidic 10 vol% HNO3 solutions at 80 °C for 2, 4 and 6 h. After the desired time elapsed, the samples were washed with distilled water, dried and weighed. The density and water absorption of the resulting materials were determined according to the procedures outlined by ASTM standard C642-13 [36]. 3. Results and discussion Fig. 2 shows the results of the DSC testing, which was used to determine the optimal heat treatment temperatures for each type of mixture. An endothermic peak was present at T < 100 °C in all the samples, which corresponds to the elimination of the adsorbed water or moisture present in the samples [37]. The exothermic peaks corresponding to the crystallization of the frit were observed between 900 and 1010 °C. The Tp of sample V–0%ZnO was 1008 °C, while the Tp of the samples containing zinc oxide slightly decreased due to addition of ZnO, however, as seen in Fig. 2, the V-10% ZnO sample shows a slight increase in the crystallization temperature compared to the V-7% ZnO sample, which coincides with that reported by Balu et al. [38], who found that with additions of more than 10% mol of ZnO (∼8% wt ZnO) in glasses of the B2O3–Na2O–ZnO system, the ZnO acts as a network former, instead of a network modifier and therefore the crystallization temperature increases with respect to samples with additions less than 3

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Fig. 3. FTIR spectra of A) glasses and B) glass-ceramics.

could be attributed to anorthite formation [15]. The band at 684 cm−1 (i.e., signal 4) could be associated with the M–O stretching vibrations of anorthite, where M = Si or Al, while the bands at 620 cm−1 (i.e., signal 5) and 571 cm−1 (i.e., signal 6) are attributed to the bending vibrations of O–Si–O. The band at 539 cm−1 (i.e., signal 7) is caused by the coupling of O–Si–O bending vibrations [40] with the Ca–O stretching vibration of anorthite [15]. Finally, the band at 468 cm−1 (i.e., signal 8) can be attributed to O– Si –O bending and Ca–O stretching, which are associated with the formation of anorthite [15]. Note that, as the ZnO concentration increases, in general, all the bands slightly shift toward lower wavenumbers compared to VC-0%ZnO, because ZnO modifies the network. This result coincides with the studies performed by Partyka [40] and Leśniak [41], who claim that the increase in the amount of zinc ions and the simultaneous decrease in the amount of silicon and aluminum ions systematically shift the bands to lower wavenumbers, because the zinc ions depolymerize the silica-oxygen network. In addition, for the samples containing ZnO, the band located at 571 cm−1 (i.e., signal 6) slightly shifts to a wavenumber higher than the corresponding band in the VC-0%ZnO sample (583 cm−1), which, according to Effendy [42], corresponds to bands associated with the asymmetric stretching vibrations of Zn–O in ZnO4. Fig. 4 shows the X-ray diffraction results of the samples after heat treatment. The resulting diffraction patterns (PDF) were identified with Crystallographica Search-Match software. The four mixtures tested generally showed the presence of the same crystalline phases, which included anorthite (CaAl2Si2O8), albite (Na(AlSi3O8)) and labradorite (Na0.45 Ca0.55 Al1.55 Si2.45 O8). Unlike sample VC-0%ZnO, the samples containing 4, 7 and 10 wt% ZnO did not contain wollastonite (CaSiO3). Although ZnO was originally added to improve the fluidity of the parent glass, ZnO plays an active role in modifying the properties of the parent

glass, as ZnO acts structurally as a network modifier. This result agrees with Chen [43], who found that the amount of the wollastonite phase decreased when the ZnO concentration increased in iron-calcium-aluminosilicate glasses. Additionally, adding ZnO increased the peak intensity of the anorthite phase, compared to the standard sample (VC-0% ZnO), and helps to form hardystonite (Ca2ZnSi2O7) in samples VC-7% ZnO and VC-10%ZnO and an additional willemite phase (Zn2SiO4) in sample VC-10%ZnO. No phases associated with zinc oxide (e.g., willemite) were found in sample VC-4%ZnO, that is, the Zn or compounds containing Zn remained in the glassy phase; therefore, adding small amounts of zinc oxide to CAS glass-ceramics only has a modifying effect on the glass network, which is confirmed by the improvement of the fluidity. The microstructures of the glass-ceramics samples were evaluated using SEM; the samples were grounded by abrasive paper (P# 4001200), and then, the samples were polished with a cloth and MasterMet™2 silica solution to achieve a mirror-smooth surface. Fig. 5 shows distinctive crystals with various sizes and morphologies: the dendritic microstructure with randomly distributed pyroxene crystals corresponds to wollastonite (point 1), the needle-like structures with irregular plate shapes are characteristic of anorthite and albite (point 2) and the presence of darker crystals dispersed in the matrix are associated with the presence of Zn and Fe in the samples (point 3) [2,10,40]. The images also show pores (black dots, point 4), and the amount of pores increases with increasing ZnO concentrations; it is possible that adding water (10 wt%) may have contributed to the porosity that developed during the initial drying process; in addition, during sintering and crystallization, the volume change associated with each phase could have caused additional pores to form, as a result of the different thermal expansion coefficients of the crystalline phases that formed [44]. Meanwhile, the crystal shapes generally become less uniform, and, conversely, the crystal size decreases with increasing ZnO concentrations. As Fig. 5C shows, the crystals become dark gray, which is associated with ZnO. According to Gasek [45], the diffusion of solid ZnO on the surface of SiO2 initiates the formation of the willemite phase, and this occurs in a similar manner as the hardystonite crystals that form on quartz grains [45]. According to Karamanov [43], glass-ceramics are considered nonporous materials, because it is assumed that the crystallization volume variation due to the transformation of the parent glass in a polycrystalline solid is related to some crystallization shrinkage; however, SEM images showed the formation of some micropores in the glassceramic samples (Fig. 5) which can be attributed to the powder preparation and subsequent crystallization performed in this study. To prepare the pellet, water was used as a binder (10 wt%), and in this process, the powders were pressed to ensure that the particles were as compact as possible to facilitate subsequent sintering. It is possible that adding water may have contributed to the porosity that developed during the initial drying process; in addition, during sintering and crystallization, the volume change associated with each phase could

Fig. 4. X-ray patterns of the prepared glass-ceramic samples. 4

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Fig. 5. SEM images of the glass-ceramics: a) VC-0%ZnO, b) VC7% ZnO and c) VC-10%ZnO.

which may be due to smaller crystal sizes (see Fig. 5). This occurs because cracks change their propagation direction when larger amounts of crystals are present, which act as barriers [48]. Additionally, the formation of spherical pores in the center of grains can disrupt the propagation of microcracks [44]; this is also supported by the fracture toughness measured in sample VC-10%ZnO relative to the reference sample (VC-0%ZnO). Furthermore, the sample densities are between 2658 and 2848 kg/m3 and generally increase with increasing ZnO concentrations since zinc oxide has a higher density than the other raw materials used (see Table 1); in addition, the densities are a function of the crystal size and crystal volume in the glass matrix. Both the volume of permeable pores (Vpp) and water adsorption correlated well with each other, except in sample VC-10%ZnO, which had a slightly smaller absorption percentage relative to sample VC-7%ZnO. The materials prepared in this study (see Table 2) have higher elastic modulus values than commercial glass -ceramics such as Slagsitall and Neoparies [12]. Likewise, the elastic modulus values of the prepared glass-ceramics are up to twice as large as those reported for natural stones such as granite and marble, and the Kic values found in this study are similar to those reported for marble. Note that the Kic of sample VC-10%ZnO is 93% of that reported for Neoparies. Chemical durability is not an intrinsic property of the material but the response to a variety of factors, such as the pH of the solution used, and test parameters, including the exposed surface, time, temperature, composition and amount of glass waste, as well as the amount of crystalline phases formed [49]. The chemical resistance of the glassceramic samples was evaluated by immersing the samples in 10 vol% NaOH and 10 vol% HNO3 solutions at 80 °C for 6 h; the results are shown in Fig. 7. These results show that the mass loss occurring in the presence of HNO3 is higher than that occurring in the presence of NaOH, which agrees with studies presented by several authors, such as Hieu [20] and Rawling [2]. Deng [50] claim that crystallization is the main factor causing the durability of the glass-ceramics in an acid solution, since phases such as lithium metasilicate, disilicate, gehlenite

have caused additional pores to form, as a result of the different thermal expansion coefficients of the crystalline phases that formed [43]. The porosity was determined by the Archimedes principle, the values obtained for VC-0%ZnO, VC-4%ZnO, VC-7%ZnO and VC-10%ZnO were 1.03%,1.45%, 2.44% and 1.55%, respectively. The VC-10%ZnO samples had lower porosities than VC-7%ZnO; this may occur because a new denser crystalline phase (i.e., willemite) formed, as shown by the XRD measurement (Fig. 4). The average nanohardness values obtained for each of the glassceramic samples were practically the same: the nanohardness of the VC0%ZnO, VC-7%ZnO and VC-10%ZnO samples were 10.21 ± 1.00 GPa, 9.91 ± 1.86 GPa and 10.06 ± 0.56 GPa, respectively. As shown by Fig. 6, the 2D scanning probe images of the samples show the morphologies of the nanoindented areas, and plastic deformation can be seen without the presence of radial cracks. The spacing between each test on the X and Y axes was 4.5 μm. Hysitron Triboview image processing software was used to determine the roughness of the glassceramics. Table 2 shows the roughness values (RMS), E, Hv, Kic, volume of permeable pores (Vpp), bulk density (ρa) and water absorption by immersion (A) of the samples. The results in Table 2 indicate that VC-0%ZnO has a greater roughness compared to the VC-7%ZnO and VC-10%ZnO samples; the crystalline phases present in the glass matrix caused the surface roughness to increase [46]. The VC-0%ZnO sample was the only sample that showed the presence of wollastonite, and this phase probably contributes to the greater roughness. The E of the samples decreased as the ZnO concentration increased because pores negatively impact mechanical properties (as demonstrated by ceramic materials), even when few pores exist [44,47]. The elastic modulus can be evaluated using the exponential relationship E = E0 exp(-bP), where E is the elastic modulus affected by the porosity, E0 is the elastic modulus without porosity, P is the porosity (0 < P < 1) and b is a constant (3 < b < 5) [47]. However, although the effect of porosity is similar to that of hardness, the HV in the VC-7%ZnO and VC-10%ZnO samples slightly increased,

Fig. 6. Scanning probe microscope images of a nanoindentation matrix: a) VC-0%ZnO, b) VC-7% ZnO and c) VC-10% ZnO. 5

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Table 2 Physical and mechanical properties of glass-ceramic materials and other similar construction materials. Material

Properties RMS (nm)

Prepared glass-ceramics VC-0%ZnO 7.58 VC-4%ZnO VC-7%ZnO 6.08 VC-10%ZnO 6.17 Commercial glass-ceramics Slagsitall Neoparies Natural Stones Marble Granite -

Ref. 3

ρa (kg/m )

Vpp (%) >

A (%)

0.45 ± 0.12 0.64 ± 0.06

2658 2775 2786 2848

0.03 0.02 0.02 0.00

1.41 1.98 2.42 1.83

0.29 0.39 0.68 0.50

70–100 88

1.8 0.7

2660 2700

30–50 40–60

0.5 1.1

2400–2600 2600–2800

E (GPa)

Hv (MPa)

105.8 ± 10.3 103.4 ± 9.3 100.2 ± 6.4

631.9 591.1 655.6 653.3

± ± ± ±

Kic (MPa·m

78.4 60.0 44.1 23.8

1/2

)

± ± ± ±

± ± ± ±

0.65 0.89 1.10 0.17

± ± ± ±

0.19 0.27 0.38 0.11

Authors

(Karamanov, 2009)

>1 -

(Karamanov, 2009; Rincón & Romero, 1996)

4. Conclusions Mixtures of industrial waste (fly ash, granulated blast furnace slag and glass waste) were successfully used to prepare CAS glass-ceramics, with similar or even better properties (e.g., elastic modulus) than some natural stones and commercial glass-ceramics. According to the results found, it was highlighted that: ZnO can act as a modifier and network former in the CAS glass network and it depends on the concentration of ZnO. In general, ZnO improves glass fluidity, reduces the temperature necessary for crystallization and positively impacts the final glass-ceramic mechanical properties. However, adding ZnO increased the porosity of the glassceramic and promoted the formation of new crystalline phases compared with the sample without ZnO, which negatively impacted the chemical properties of the samples VC-7%ZnO and VC-10%ZnO, but, the mass loss of the glass-ceramic samples did not exceed 1.5% when immersed in an alkaline solution (NaOH) and did not exceed 8% when immersed in an acid (HNO3) solution. The glass-ceramics developed in this study achieved an elastic modulus that was 110% larger than that reported for natural marble and comparable to those of commercially available glass-ceramics. Additionally, the Kic obtained for VC-10%ZnO was similar to that reported for marble.

Fig. 7. Chemical resistance of the glass-ceramics in mass loss terms for 6 h.

and anorthite can coexist. These phases decompose in the presence of acids and adversely affect the chemical resistance of the material; conversely, quartz or different pyroxene phases and their solid solutions improve the chemical resistance of materials due to their high resistance to acid attack (i.e., these phases are almost insoluble) [51,52]. Other factors besides the crystalline phases impact the final material properties, and these factors include the porosity, A and Vpp, which may have contributed to transporting liquids through the sample. It is possible that the aqueous solution penetrated the pores, and therefore, chemical attack occurred faster according to a kinetic process that can /RT be expressed as V = k exp(−E ), where k and R represent two cona stants, E is the activation energy of the process and T is the temperature; therefore, a slight increase in the temperature (80 °C in this test) produces a large increase in the reaction rate [53]. Furthermore, according to El-Alaily [54], cation modifiers or metal ions occupy the percolation channels between parts of the network structure, and these channels are the most likely routes for cation diffusion, ion exchange and water entry. It is important to note that the chemical stability is consistent with the progressive formation of a more compact and interconnected glass network as glass formers are added; therefore, the increase in the proportion of SiO2 strengthens the chemical stability of the glasses [55], and in this study, although the CaO/SiO2 ratio was held constant for all of the mixtures, these species could be replaced as ZnO was added. The low mass loss (< 1.5 wt%) of the glass-ceramic synthesized in this study after immersion for 6 h in the NaOH solution is probably caused by the amphoteric nature of ZnO [56].

Declaration of competing interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgements The authors would like to thank the Universidad del Valle (Colombia) for the financing support received. In addition, the authors thank CINVESTAV for its support in characterization tests. References [1] A.A. Francis, Conversion of blast furnace slag into new glass-ceramic material, J. Eur. Ceram. Soc. 24 (2004) 2819–2824, https://doi.org/10.1016/j.jeurceramsoc. 2003.08.019. [2] R.D. Rawlings, J.P. Wu, A.R. Boccaccini, Glass-ceramics : their production from wastes — a review, J. Mater. Sci. 41 (2006) 733–761, https://doi.org/10.1007/ s10853-006-6554-3. [3] W. Höland, G. Beall, Glass-Ceramic Technology, Second, John Wileys & Sons, 2012. [4] E. Le Bourhis, Glass mechanics and technology, Second, Wiley-VCH, Verlag GmbH & Co.KGaA, 2014, p. 77. [5] J.M. Rincón, M. Romero, Los materiales vitrocerámicos en la construcción, Mater. Construcción 46 (1996) 91–106. [6] A.A. Francis, R.D. Rawlings, R. Sweeney, A.R. Boccaccini, Processing of coal ash into glass ceramic products by powder technology and sintering, Glass Technol. 43 (2002) 58–62.

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