Ceramic tiles derived from coal fly ash: Preparation and mechanical characterization

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Ceramic tiles derived from coal fly ash: Preparation and mechanical characterization ⁎

Yang Luoa,b, Shili Zhenga, Shuhua Maa, , Chunli Liua,b, Xiaohui Wanga a National Engineering Laboratory for Hydrometallurgical Cleaner Production Technology, Key Laboratory of Green Process and Engineering, Institute of Process Engineering, Chinese Academy of Sciences, Beijing 100190, China b University of Chinese Academy of Sciences, Beijing 100049, China

A R T I C L E I N F O

A BS T RAC T

Keywords: Coal fly ash Alkali activation Ceramic tile Zeolite Hydroxysodalite

Coal fly ash (CFA) accounts for a large fraction of the solid waste produced in China. Hence, there is an urgent need for the effective utilization of CFA, for example, as a raw material for ceramics production. In this study, clay- and feldspar-like materials fabricated by alkali activation pre-treatment of CFA were mixed with untreated CFA (regarded as a quartz-like material) and sintered to prepare fully ash-based ceramic tiles. The obtained tiles exhibited excellent sintering properties, e.g., low firing temperature and a wide sintering range; further, they showed better green strength (due to hydrogen bonding) and post-sintering performance (due to fluxing and mullite skeleton effects) than ceramic tiles produced exclusively from untreated CFA. The fully ash-based ceramic tiles sintered at 1100 °C exhibited optimal post-sintering properties (bulk density, 2.5 g/cm3; rupture modulus, 50.1 MPa; and water absorption, 0%). Thus, the proposed method is well suited for preparing a novel kind of ceramic tiles completely derived from CFA, highlighting its importance in the field of fly ash ceramics.

1. Introduction The rapid development of China's thermal power industry has resulted in a continuous rise of coal fly ash (CFA) emissions [1]. Since CFA pollutes the environment besides piling up on the ground [2], its effective utilization is a top priority for China. Ceramic tiles are produced from mixtures containing clay, feldspar, and quartz [3], typically exhibiting a triaxial formulation comprising ~50% clay, 35% feldspar, and 15% quartz. In these mixtures, clay acts as a binder for other the constituents of the green state, conferring the plasticity required for shaping [4]. On the other hand, the low-melting feldspar phase reacts with other components and lowers the liquid formation temperature of the system [5], and the produced liquid permeates the microstructure to cause densification [6]. Finally, quartz is reasonably stable at industrial firing temperatures and thus reduces distortion and shrinkage [7]. Ceramic tiles exhibit grain- and bond-type microstructures, with filler particles (usually quartz) held together by a fine and almost fully dense matrix composed of mullite crystals and a glassy phase [8]. CFA contains valuable oxides such as SiO2, Al2O3, Na2O, K2O, CaO, and MgO, and it is considered as a low-cost finely granular material in the ceramic tile industry [9,10]. To use CFA wide spreadly in ceramic

tiles, much effort has been devoted to preparing ceramic tiles based on CFA. For example, Aineto et al. [11] added CFA to clays having different plasticities to obtain building ceramics. Dana et al. [12] replaced quartz in a traditional triaxial porcelain composition (containing kaolinitic clay, quartz, and feldspar) by 5, 10, and 15 wt% of fly ash to prepare high-strength porcelains. Ji et al. [13], on the other hand, used CFA as a feldspar substitute to fabricate ceramic tiles. Although these studies verify the feasibility of using CFA as an alternative raw material for ceramic tiles, there is no scientific consensus on its advantages when used in ceramics. Currently, the amount of added CFA is not large enough ( < 30 wt%) to affect the widely used triaxial ceramic system. Although Erol et al. [14,15] reported the preparation of additive-free ceramic materials from CFA, the strength of green compacts and the corresponding sintering mechanism have not been studied in detail. CFA is commonly subjected to hydrothermal alkali activation during alumina extraction [16], cement manufacture [17], and zeolite synthesis [18]. However, alkali activation pre-treatment is seldom used in the field of fly ash ceramics. This activation method was adopted from our earlier research, and alkali-activated CFA was used in combination with traditional materials to prepare commercial ceramic tiles with excellent green strength and post-sintering performance [19].

Abbreviations: CFA, Coal fly ash; XRD, X-ray diffraction; ICP-OES, inductively coupled plasma optical emission spectrometry; SEM, scanning electron microscopy; EDS, energydispersive X-ray spectroscopy; XPS, X-ray photoelectron spectroscopy; TG–DSC, thermogravimetric analysis and differential scanning calorimetry ⁎ Corresponding author. E-mail address: [email protected] (S. Ma). http://dx.doi.org/10.1016/j.ceramint.2017.06.045 Received 22 December 2016; Received in revised form 15 May 2017; Accepted 6 June 2017 0272-8842/ © 2017 Elsevier Ltd and Techna Group S.r.l. All rights reserved.

Please cite this article as: Luo, Y., Ceramics International (2017), http://dx.doi.org/10.1016/j.ceramint.2017.06.045

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d(0.5) = 75.32 µm. Reagent-grade sodium hydroxide was sourced from Xilong Chemical Co., Ltd., and was used as received.

Table 1 Chemical composition of untreated CFA (wt%). Constituent Content

Al2O3 21.47

SiO2 55.57

Fe2O3 6.80

TiO2 0.01

CaO 5.12

MgO 2.97

Na2O 3.42

K2O 1.22

LOI 2.83

LOI: loss on ignition.

2.2. Experimental procedures

Therefore, this work aimed to prepare ceramic tiles based solely on CFA subjected to alkali activation and to study the corresponding reaction mechanism. The use of CFA as a ceramic raw material has dual benefits in terms of reducing environmental load: reducing the need to mine materials, and mitigating the accumulation of fly ash waste.

Alkali activation of CFA was performed under stirring in a 1-L hightemperature reactor with external heating and internal cooling. The system was equipped with an automatic proportional–integral–derivative control system to regulate the heating rate, agitation, and temperature. CFA and NaOH solution (60, 80, 100, 120, or 140 g/L) were added to the reactor, and mixtures with a liquid-to-solid ratio of 5 mL/g were digested at 150 °C for 3 h. After the reaction, the slurry was filtered and washed with hot deionized water five times to minimize the amount of adsorbed sodium ions. The obtained samples were dried in an oven at 80 °C for 12 h and denoted as CFA 60, CFA 80, CFA 100, CFA 120, and CFA 140 depending on the NaOH concentration used. Production of ceramic tiles from CFA involves two stages: green body production and sintering. Prior to preparing the green bodies, untreated CFA (~15 wt%), CFA 80 (~50 wt%), and CFA 140 (~35 wt%) were homogenized using a planetary ball mill with alumina balls as grinding media. The homogenized raw materials were mixed with deionized water in a 10:2 mass ratio, and the mixtures were pressed into 100 × 100 × 4 mm cuboids with the help of a uniaxial tablet presser (Model WE-300B) under 20 MPa. The obtained green compacts were dried at 105 °C for 12 h and then sintered under ambient

2. Experimental 2.1. Materials CFA was sampled from a thermoelectric power plant in Inner Mongolia, China. Untreated CFA (Table 1) has Al2O3 and SiO2 contents of 21.47 and 55.57 wt%, respectively, which are very similar to those of conventional ceramic raw materials. The X-ray diffraction (XRD) pattern shown in Fig. 1a reveals quartz (JCPDS card no. 01-0830539), mullite (JCPDS card no. 00-001-0613), hematite (JCPDS card no. 01-073-0603), and amorphous materials as the main phases. The particle size distribution and morphology are shown in Figs. 1b and c, respectively. CFA particles appear as glossy microspheres with different sizes, exhibiting a normal distribution with a peak at 85.32 µm and

12

Q--- Quartz H--- Hematite M--- Mullite

10

Differential Volume/%

Intensity(a.u.)

Q

Q M

M M HQ Q HQ

Q

HMH

8

6

4

2

Q

0

5

10 15 20 25 30 35 40 45 50 55 60 65 70 75 80 85 90

(a)

1

Diffraction Angle,2θ/degree(CuKα)

(b)

10

100

Particle Diameter/μm

(c) Fig. 1. Untreated CFA characterization: (a) XRD pattern, (b) size distribution, and (c) SEM image.

2

1000

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Coal fly ash

2.3. Characterization techniques

No special treatment

Mild alkali activation

Severe alkali activation

Untreated CFA

CFA 80

CFA 140

The chemical composition of CFA samples was analyzed by inductively coupled plasma-optical emission spectrometry (ICP-OES, Optimal 5300DV, PerkinElmer Instruments) under the following conditions: power, 1300 W; carrier gas flow, 0.08 L/min; and peristaltic pump flow, 1.5 mL/min. Crystalline phases were identified by XRD analysis (X′Pert Pro MPD, PANalytical Company) performed at 40 kV and 30 mA using Cu Kα radiation. The microstructures of the samples were observed by scanning electron microscopy (SEM, Sirion 200, FEI), while the concentrations of individual elements were determined by energy-dispersive X-ray spectroscopy (EDS, INCA ENERGY 300). For morphology and phase assemblage analyses, the fracture surfaces of the fired samples were polished with diamond pastes after initial grinding with nanometer MgO powder and water. The polished surfaces were etched in 20 wt% HF solution for 30 s, sonicated in distilled water and ethanol, dried, and coated with platinum. Secondary electron images were used for microstructural examination, and backscattered electron images were used for phase identification. The particle size distribution of the powders was obtained by a laser particle size analyzer (Beckman Coulter, LS 13320). The thermal behavior was monitored by thermogravimetric analysis and differential scanning calorimetry (TG–DSC, NETZSCH STA 449 C) in air flow, at a heating rate of 20 °C/min. X-ray photoelectron spectroscopy (XPS, ESCALAB 250Xi) was performed using Al Kα radiation (hν = 1486.6 eV), and the binding energies were corrected with respect to that of adventitious carbon (284.8 eV). The bulk densities and water absorption of the ceramic tiles were determined using the Archimedes method according to ASTM C373. The rupture moduli of the green compacts and ceramic samples were measured on an electronic universal testing machine (WDW-20E, Jinan Shidai Shijin Testing Machine Group Co., Ltd.), by a three-point loading test at a span of 80 mm and a cross-head speed of 0.05 mm/ min. The linear shrinkage of sintered samples was determined as follows:

Batching

Molding

Sintering

Fully fly-ash-based ceramic tiles Fig. 2. The fabrication diagram of full fly ash based ceramic tiles.

conditions in a laboratory electrical sintering furnace (Model HTF1400) at 1000, 1050, 1100, 1150, or 1200 °C. All samples were heated from room temperature to the required sintering temperature at a rate of 10 °C/min, held at the maximum temperature for 60 min, and cooled to room temperature at the rate of 5 °C/min. The resulting ceramic samples were denoted as NCRE 1000, NCRE 1050, NCRE 1100, and NCRE 1150 depending on the sintering temperature used. A schematic of the abovementioned sequence is illustrated in Fig. 2. For comparison, two common ash-based ceramic tiles were prepared by the same method, using untreated CFA, at a sintering temperature of 1100 or 1200 °C (samples denoted as OCRE 1100 and OCRE 1200). To study the effect of alkali activation pre-treatment on the green strength, the green compact produced from untreated CFA (~15 wt%), CFA 80 (~50 wt%), and CFA 140 (~35 wt%), denoted as GNCRE, and compacts fabricated from untreated CFA, CFA 60, CFA 80, CFA 100, CFA 120, and CFA 140 (denoted as GOCRE, GCFA 60, GCFA 80, GCFA 100, GCFA 120, and GCFA 140, respectively) were used. In addition, the pressure is in keeping with the prior pressure (20 MPa). For clarity, the nomenclatures used for the different samples are summarized in Table 2.

Linear shrinkage (%) =

L g − Ls × 100 Lg

where Lg and Ls are the side lengths (mm) of the green and sintered samples, respectively, as measured by a sliding gauge. The bulk densities, water absorption, rupture moduli, and linear shrinkages of ceramic samples were averaged over five sets of data in replicate measurements. 3. Results and discussion 3.1. Characterization of different types of alkali-activated fly ash 3.1.1. Mineralogical and microstructural properties The reaction of CFA with NaOH proceeded as described in previous studies [20,21]. First, the amorphous glass, quartz, and mullite

Table 2 Nomenclatures used for the different samples. Nomenclature for sintered sample

Description for sintered sample

Nomenclature for green compact

Description for green compact

NCRE 1000

Made of untreated CFA, CFA 80 and CFA 140; sintered at 1000 °C. Made of untreated CFA, CFA 80 and CFA 140; sintered at 1050 °C. Made of untreated CFA, CFA 80 and CFA 140; sintered at 1100 °C. Made of untreated CFA, CFA 80 and CFA 140; sintered at 1150 °C. Made of untreated CFA; sintered at 1100 °C. Made of untreated CFA; sintered at 1200 °C.

GNCRE GOCRE

Made of untreated CFA, CFA 80 and CFA 140. Made of untreated CFA.

GCFA 60

Made of CFA 60.

GCFA 80

Made of CFA 80.

GCFA 100 GCFA 120 GCFA 140

Made of CFA 100. Made of CFA 120. Made of CFA 140.

NCRE 1050 NCRE 1100 NCRE 1150 OCRE 1100 OCRE 1200

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mainly of zeolite P, such as CFA 60 and CFA 80, denoted as class-I CFAs; (2) products consisting of both zeolite P and hydroxysodalite, e.g., CFA 100; (3) products consisting mainly of hydroxysodalite, e.g., CFA 120 and CFA 140, denoted as class-II CFAs. To further verify the above hypothesis, further intuitive insights were obtained from the SEM images (Fig. 4). Secondary-electron SEM imaging showed that the petaloid zeolite gradually transformed into the granular hydroxysodalite with increasing NaOH concentration (the two substances were identified by EDS point analysis, as shown in Fig. 4-I and Table 4). When the NaOH concentration was increased from 60 to 80 g/L, the zeolite gradually covered the surfaces of class-I CFAs (Fig. 4a-I and Fig. 4b-I), and its lamellar microstructure was gradually refined, as confirmed from high-magnification images (Fig. 4a-II and Fig. 4b-II). The zeolite and hydroxysodalite phases coexisted at the NaOH concentration of 100 g/L (Fig. 4c-I and Fig. 4c-II); however, when the alkali concentration exceeded 100 g/L, the surfaces of class-II CFAs were completely covered by granular hydroxysodalite (Fig. 4d-I and Fig. 4e-I). Gradual fiber widening was observed with continuously increasing NaOH concentration, as confirmed from high-magnification images (Fig. 4d-II and Fig. 4e-II). The combination of grinding/polishing with backscattered-electron SEM imaging reveals the elemental distribution within the sample particles. Images of class-I CFA particles show two shades of grey, the inner cores of the particles being darker and the outer surface-coated layers being brighter (Fig. 4a-III and Fig. 4b-III). In these images, the main factor determining the shade of grey is the atomic number, with lighter shades indicating higher average atomic numbers [22]. Therefore, the average atomic number of outer surface-coated layers is higher than that of the inner core, indicating that class-I CFA particles were successfully coated by the newly generated phases during alkali activation. The corresponding shades of the class-II CFA particle surfaces and interiors are almost uniform, implying similar chemical compositions (Fig. 4d-III and Fig. 4e-III). In addition, CFA 100 is a transitional state and hence exhibits a shade of grey between those of the abovementioned two materials (Fig. 4c-III). EDS line scan analysis (from A to B, Fig. 4a-III and Fig. 4b-III) of class-I CFA particles suggests that the Na contents of the outer surface-coated layers are much higher than those of the inner cores. The Na content difference between the outer surface-coated layer and the inner core is greatly reduced for CFA 100 particles as compared to that for the classI CFA particles (from A to B, Fig. 4c-III), and almost no difference is observed for the class-II particles (from A to B, Fig. 4d-III and Fig. 4eIII). The results of EDS line scan analyses are consistent with those of backscattered-electron SEM imaging, confirming that class-I CFAs possess a reactant coating that is absent in class-II CFAs. This phenomenon can be explained on the basis of backscattered-electron and secondary-electron SEM imaging. The dense microstructure of the petaloid zeolite prevents continuous internal mass transfer, while the granular hydroxysodalite has a highly porous microstructure featuring numerous tiny fibers, with the voids between these fibers providing a favorable environment for mass transfer. To re-stress, based on the results of ICP-OES, XRD, and SEM-EDS analyses, alkali-activated CFAs can be grouped into two fairly different classes, namely, class-I and class-II CFAs.

Table 3 Chemical compositions of alkali-activated CFAs (wt%). Sample

Al2O3

SiO2

Fe2O3

TiO2

CaO

MgO

Na2O

K2O

LOI

CFA CFA CFA CFA CFA

22.57 22.33 21.93 21.57 21.31

34.44 36.00 37.20 40.80 42.13

8.97 8.88 7.89 7.11 6.57

0.65 0.65 0.70 0.70 0.65

5.66 5.50 5.57 5.54 5.43

2.62 2.68 2.68 2.77 2.67

8.24 8.65 10.74 11.12 12.01

0.72 0.58 0.47 0.41 0.34

16.13 14.73 12.82 9.98 8.89

60 80 100 120 140

LOI: loss on ignition.

components of CFA partially dissolved in NaOH solution (Eqs. (1) and (2)), and the produced Na+, Al(OH)4– and H2SiO42– ions reacted to precipitate new solid phases on the surfaces of the CFA particles (Eqs. (3) and (4)). To further investigate the effect of the alkali activator (NaOH) concentration on the physical and chemical properties of fly ash, CFA 60, CFA 80, CFA 100, CFA 120, and CFA 140 were characterized by ICP-OES, XRD, SEM, and EDS. SiO2(quartz

or glass)

+ NaOH(aq) = Na2SiO3(aq) + H2O(l)

(1)

3Al2O3·2SiO2(s) + 10NaOH(aq) = 6NaAlO2(aq) + 2Na2SiO3(aq) + 5H2O(l) (2) 6Na+(aq) + 6Al(OH)4–(aq) + 10H2SiO42–(aq) = Na6Al6Si10O32·12H2O(s) (P zeolite) + 20OH–(aq) (3) 8Na+(aq) + 6Al(OH)4–(aq) + 6H2SiO42–(aq) = Na8Al6Si6O24(OH)2·4H2O(s) (hydroxysodalite) + 10H2O(l) + 10OH– (aq) (4) The chemical compositions of the alkali-activated CFAs are listed in Table 3. As seen from the data, the SiO2 and Na2O contents increase with increasing NaOH concentration, while the loss on ignition shows an obvious downward trend. The Al2O3, Fe2O3, CaO, and K2O contents show a slight decrease, whereas the TiO2 and MgO contents remain unchanged. Under the experimental conditions used, increased NaOH concentrations inhibit desilication and promote the introduction of more Na+ ions into the CFA structure. Fig. 3 shows the XRD patterns of the alkali-activated CFAs, highlighting the phases corresponding to reaction products. As the NaOH concentration is increased from 60 to 140 g/L, the diffraction peaks of quartz (JCPDS card no. 01-083-0539) progressively weaken, while peaks due to zeolite P (JCPDS card no. 01-080-0699) and hydroxysodalite (JCPDS card no. 01-076-1639) emerge and become stronger. These diffraction intensity changes suggest that zeolite P is produced first, followed by hydroxysodalite, implying that zeolite P is an intermediate formed in the alkali activation of CFA. The reaction products can be divided into three types: (1) products consisting

H

H

Intensity(a.u.)

H

Q

H

H

HQ

H

H

P PH P P HQ

P

P--- Zeolite P Η--- Hydroxysodalite Q---Quartz

CFA 140 CFA 120

P

CFA 100

P

CFA 80

P

CFA 60

3.1.2. Thermal behavior Thermal behavior is one of the most important characteristics of ceramic raw materials [23]. To investigate the thermal behavior of untreated and alkali-activated CFAs, TG-DSC and XRD analyses were employed; the results are shown in Fig. 5. The observations indicate that alkali activation pre-treatment significantly affects the thermal behavior of CFAs, and alkali-activated CFAs can be divided into three classes based on the shape of their TG-DSC curves, in line with the above classification. In the case of untreated CFA, the mass loss in the TG curve and the endothermic peaks in the DSC curve are less pronounced (Fig. 5a),

P P

P

5

P

P

P Q

P P Q P

P

10 15 20 25 30 35 40 45 50 55 60 65 70 75 80 85 90

Diffraction Angle,2θ/degree(CuKα) Fig. 3. XRD patterns of the alkali-activated CFAs.

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Fig. 4. SEM images of the alkali-activated CFAs: (a) CFA 60, (b) CFA 80, (c) CFA 100, (d) CFA 120, and (e) CFA 140; I – low-magnification secondary-electron imaging, II – highmagnification secondary-electron imaging, and III – backscattered-electron imaging of the grinded and polished samples with EDS line scan analysis.

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densities, rupture moduli, and linear shrinkages, along with a marginal increase in water absorption, which can be explained by the oversintering phenomenon. On the one hand, oversintering leads to the sharp drop in liquid viscosity and the corresponding dramatical expansion of the tiny pores within the ceramic bodies. On the other hand, the excessively high temperature is associated with the abnormal grain growth which means that the grain boundary movement speed is greater than the pore movement speed, and this phenomenon makes the tiny pores encapsulated within the grains and causes the deterioration of various properties of ceramics. Samples sintered above the peak densification temperature developed fully vitrified surfaces but were not bloated, as shown in Fig. 7. Notably, the post-sintering properties of sintered samples were consistent with the GB/T 4100-2015 standard (Chinese national standard for ceramic tiles, requesting for 35 MPa) over a wide sintering temperature range (1050–1150 °C). This observation could be explained by the fact that the simple Na2O-K2O fluxing system of alkali-activated CFAs has a wider melting temperature range compared to the traditional CaO-MgO-Na2O-K2O fluxing system [25]. Therefore, alkali-activated fully fly-ash-based ceramic tiles not only require a lower firing temperature but also exhibit a wide sintering range, eliminating the need for precise temperature control. Among the sintering conditions tested in this study, treatment at 1100 °C produced the densest and most solid ceramics with a bulk density of 2.5 g/cm3, water absorption of 0%, rupture modulus of 50.1 MPa, and linear shrinkage of 17.6%.

Table 4 EDS point analysis in Fig. 4-I (wt%). Point

OK

Na K

Mg K

Al K

Si K

KK

Ca K

Ti K

Fe K

a b c d e f

42.35 48.70 43.05 44.43 46.83 46.77

7.42 7.47 10.83 8.58 11.17 12.73

0.08 0.13 1.81 0.97 1.26 0.22

17.30 15.26 11.10 12.07 13.56 16.24

29.07 25.62 21.07 23.50 21.97 20.93

0.89 0.74 0.75 0.84 0.39 0.22

1.01 0.80 4.45 4.25 3.15 1.42

0.48 0.40 0.65 0.66 0.37 0.33

1.40 0.89 6.30 4.68 1.31 1.14

indicating inert thermal behavior. A slight mass loss is observed between 50 and 400 °C, which is associated with an endothermic hump in the DSC curve due to the loss of adsorbed and crystal water. Moreover, no obvious exothermic peak is detected in the DSC curve, implying that the amount of residual carbon in untreated CFA is small. The gentle thermic valleys between 700 and 1000 °C can be attributed to the formation of a new anorthite (JCPDS card no. 00-002-0537) phase in this temperature range, as confirmed by XRD analysis. The relatively prominent endothermic peak at ~1239 °C represents the melting point, since it is not accompanied by mass loss. For class-I CFAs (Figs. 5b and c), similar mass losses and intense endothermic peaks at ~145 °C represent the loss of large amounts of adsorbed and crystal water. The exothermic peaks between 700 and 1000 °C are not accompanied by a concomitant mass loss, which is ascribed to the transformation of P zeolite (JCPDS card no. 01-0800699) to nepheline (JCPDS card no. 01-083-2372) and mullite (JCPDS card no. 00-001-0613) phases, as confirmed by XRD analysis. Since the endothermic peaks at ~1130 °C are not accompanied by mass loss, they correspond to the melting points of class-I CFAs, which are ~109 °C lower than that of the untreated CFA due to their higher Na2O contents. Class-II CFAs (Figs. 5e and f) exhibit endothermic water loss and melting peaks at ~136 and ~1035 °C, respectively. Compared to class-I CFAs, class-II CFAs exhibit a less pronounced water loss, as is reflected in the smaller mass losses and endothermic peaks. Furthermore, their melting point is ~95 °C lower than that of class-I CFAs, and this value is the lowest among all samples due to the highest Na2O content of these species. The exothermic peaks between 700 and 1000 °C are attributed to the transformation of hydroxysodalite (JCPDS card no. 01-076-1639) into nepheline (JCPDS card no. 01-083-2372). The thermal behavior of CFA 100 (Fig. 5d) is intermediate between that of class-I and class-II CFAs. Briefly, untreated CFA, class-I CFAs, and class-II CFAs exhibit three different kinds of thermal behavior corresponding to quartz, clay, and feldspar in traditional ceramic raw materials, respectively. Therefore, untreated CFA, CFA 80, and CFA 140 were used as substitutes for these materials to prepare the fully ash-based ceramics discussed in this study.

3.2.2. Mineralogical and microstructural characteristics of sintered samples The effect of sintering temperature on the ceramic samples was investigated by XRD, with patterns obtained for different temperatures shown in Fig. 8. As the temperature was increased, the quartz phase (JCPDS card no. 01-083-0539) gradually transformed into a tridymite (JCPDS card no. 00-016-0152) phase, while the nepheline (JCPDS card no. 01-083-2372) phase gradually transformed into albite (JCPDS card no. 00-020-0572) and mullite (JCPDS card no. 00-001-0613) phases. Notably, the intensity of characteristic mullite peaks markedly increased with temperature. Based on the mullite strength mechanism, the strength of the sintered ceramic is proportional to the amount of mullite contained therein. In agreement with the rupture modulus data in Fig. 6c, the mullite diffraction peak intensity reached its maximum at 1100 °C and then dropped with a further increase in temperature. When the sintering temperature reached 1150 °C, nepheline was replaced by albite, and a part of mullite dissolved in the growing glass phase, leading to decreased strength. Fig. 9 shows the results of SEM-EDS characterization of NCRE 1100 fracture surfaces, indicating that particles with the original morphology are held together by compact and dense bonds (Fig. 9a) comprising needle-like mullite crystals (Fig. 9b) and a glass phase. Some CFA particles are fractured during ball milling, while others retain their shapes even after sintering. Three types of sintered CFA particles exist in the ceramic body: (a) particles not integrated into the surrounding vitreous matrix, acting as filler during sintering and identified as sintered untreated CFA by EDS (Fig. 9c); (b) particles possessing an outer glass shell fused with the vitreous matrix and an inner crystalline core with crystallized fibrous whiskers, identified by EDS as nepheline and mullite, respectively (Figs. 9d and e); (c) particles entirely fused with the surrounding vitreous matrix, identified by EDS as nepheline (Fig. 9f). From the characterization of alkali-activated CFAs (Section 3.1), one can infer that untreated CFA exhibits low reactivity and contributes to the reduction of drying and firing shrinkages. On the other hand, CFA 80 has a fused outer shell and a crystalline inner core, and CFA 140 is fused with the surrounding vitreous matrix. During sintering, untreated CFA acts as a filler, similar to quartz; CFA 80 (class-I CFA) acts as a crystal skeleton, similar to clay; and CFA 140 (class-II CFA) functions as a fluxing agent, similar to feldspar. In addition, consider-

3.2. Characterization of fully fly-ash-based ceramic tiles 3.2.1. Post-sintering properties of sintered samples Bulk density, water absorption, rupture modulus, and linear shrinkage are the four main indices determining the quality of ceramics in a real production process. Fig. 6a–d show that increased sintering temperatures initially result in higher bulk densities, rupture moduli, and linear shrinkages, and reduce the water absorption of the sintered samples. Maximum bulk densities and rupture moduli were obtained at ~1100 °C, which is about 100 °C lower than the optimal industrial temperature [24] and can be explained as follows. For alkali-activated fully fly-ash-based ceramic tiles, the formation of zeolite P and hydroxysodalite resulted in increased Na2O content, which lowered the sample softening temperature and induced low-temperature sintering densification. Further increases in sintering temperature resulted in reduced bulk 6

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3.3. Comparison of tiles fabricated from alkali-activated and untreated ash

ing the results of SEM-EDS analysis of alkali-activated CFAs (Section 3.1.1), the uniform zeolite coating of class-I CFAs and the homogeneous hydroxysodalite structure of class-II CFAs enable the homogeneous dispersion of the fluxing agent, which is difficult to achieve in the current ceramic production process.

3.3.1. Comparison of green bodies Alkali activation pre-treatment was found to significantly increase the strength of green bodies (from almost 0–5.12 MPa, significantly exceeding the value of 2.5 MPa required for actual production) and impart plasticity to CFA, as confirmed by the rupture moduli of different green compacts listed in Table 5. The rupture moduli of

Fig. 5. Thermal analyses (TG-DSC) and XRD spectra (at 700 and 1000 °C) of untreated CFA and the alkali-activated CFAs: (a) untreated CFA, (b) CFA 60, (c) CFA 80, (d) CFA 100, (e) CFA 120, and (f) CFA 140； I – TG-DSC, II – XRD spectra (at 700 and 1000 °C).

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Fig. 5. (continued)

lattice oxygen (denoted as Oα), OH groups (denoted as Oβ), and molecular water (denoted as Oγ), respectively [27]. While Oα is the only oxygen species in untreated CFA, indicative of an oxide nature without crystal water and OH groups, CFA 60 and CFA 80 (class-I CFAs) additionally contain Oγ, suggesting that class-I CFAs are oxides containing crystal water. Furthermore, CFA 100 is a transitional state between class-I CFAs and class-II CFAs, whereas both CFA 120 and CFA 140 (class-II CFAs) contain three types of oxygen species, indicating the presence of crystal water and OH groups. Table 6 shows that upon transformation of untreated CFA into class-I CFAs, the

green compacts increase from GOCRE to GCFA80 and then decrease from GCFA80 to GCFA140, implying a large plasticity increase in going from untreated CFA to class-I CFAs, and a subsequent slight decrease going from class-I CFAs to class-II CFAs. XPS is widely used for determining the surface chemical states of solid materials [26]. O 1 s XPS spectra were recorded for untreated and alkali-activated CFAs (Fig. 10 and Table 6) to explain the green strength increase from the viewpoint of the applied force mechanism. The above spectra reveal three types of oxygen with characteristic peaks at 530.1–530.3, 531.1–531.3, and 532.3–532.7 eV, ascribed to 8

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Fig. 6. Post-sintering properties of the full fly ash based ceramic tiles: (a) bulk density, (b) water absorption, (c) rupture modulus, and (d) linear shrinkage.

Fig. 7. Pictures of ceramic samples sintered at different temperatures.

Oα:Oβ:Oγ ratio changes from 100:0:0 to 80:0:20 and then to 75:0:25, coming increasingly closer to the theoretical Oα:Oβ:Oγ ratio of zeolite P (Na6Al6Si10O32·12H2O). When class-I CFAs transformed into class-II CFAs, the above ratio gradually changed from 75:0:25 to 80:6.67:13.33, which is very close to the theoretical ratio of hydroxysodalite (Na8Al6Si6O24(OH)2·4H2O). The results of these analyses are in good agreement with those of ICP-OES, XRD, and SEM-EDS analyses in Section 3.1.1. Based on the above XPS results, the forces acting between particles can be clearly explained by the applied force mechanism (Fig. 11). Van der Waals and capillary interactions are the two main forces acting on untreated fly ash particles, while alkali-activated particles exhibit additional hydrogen bonding interactions of surface crystal water molecules and surface OH groups. The hydrogen bonding interactions of crystal water molecules are stronger than those of OH groups because a water molecule possesses twice as many hydrogen atoms as an OH group, implying that class-I CFAs have better plasticity than

Fig. 8. XRD patterns of ceramic samples sintered at different temperatures.

class-II CFAs. Actually, these two classes of alkali-activated CFAs act as binders in green bodies, imparting good plasticity and acting analogously to clay, despite the existing mechanistic differences.

3.3.2. Comparison of sintered samples This section compares macro-performances and micro-properties of NCRE 1100, OCRE 1100, and OCRE 1200; images of the sintered ceramic samples are shown in Fig. 12. Although both NCRE 1100 and 9

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Fig. 9. SEM images and EDS analyses of NCRE 1100: (a) cross-section image, (b) needle-like mullite crystals, (c) quartz-like particle, (d, e) clay-like particle, and (f) feldspar-like particle.

Table 5 Rupture moduli of green compacts (MPa). Green compact Rupture modulus

GNCRE 5.12

GOCRE ~0

GCFA60 6.51

GCFA80 6.86

10

GCFA100 6.14

GCFA120 5.21

GCFA140 4.25

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Oα

Oα

Oγ

540

538

536

534

(a)

532

530

528

526

524

522

520

540

538

536

534

(b)

Binding Energy/eV

532

530

528

526

524

522

520

Binding Energy (eV)

Oα

Oα

Oγ Oγ Oβ

540

538

536

534

532

530

528

526

524

522

520

540

Binding Energy (eV)

(c)

538

536

Oγ

536

530

528

526

524

522

520

Oα

Oγ

Oβ

538

532

Binding Energy (eV)

(d)

Oα

540

534

Oβ

534

(e)

532

530

528

526

524

522

520

540

538

(f)

Binding Energy (eV)

536

534

532

530

528

526

524

522

520

Binding Energy (eV)

Fig. 10. O 1s XPS spectra of the untreated fly ash and the alkali-activated CFAs: (a) untreated CFA, (b) CFA 60, (c) CFA 80, (d) CFA 100, (e) CFA 120, and (f) CFA 140. Table 6 O 1s peak ratios for untreated and alkali-activated CFAs (%). Sample

Oα

Oβ

Oγ

Untreated CFA CFA 60 CFA 80 CFA 100 CFA 120 CFA 140

100.00 80.00 75.00 76.00 76.93 80.00

0 0 0 5.25 7.69 6.67

0 20.00 25.00 18.75 15.38 13.33 Fig. 11. Schematic diagram of the force existed between particles before and after alkali activation.

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Fig. 12. Pictures of sintered ceramic samples: (a) the optimized sample – NCRE 1100, (b) comparison of NCRE 1100, OCRE 1100 and OCRE 1200. Table 7 Macro-performances of ceramic tiles. Bulk density (g/cm3)

Sample

NCRE 1100 OCRE 1100 OCRE 1200

Water absorption (%)

Rupture modulus (MPa)

Max

Min

Mean

Max

Min

Mean

Max

Min

Mean

Max

Min

Mean

2.55 1.99 2.45

2.41 1.92 2.33

2.50 1.97 2.41

0 13.33 0

0 11.92 0

0 11.99 0

54.49 9.98 13.32

48.27 6.99 12.10

50.10 8.09 12.70

18.00 6.70 16.33

17.50 6.54 14.99

17.60 6.60 15.32

structure of mullite is beneficial to ceramic strength; therefore, NCRE 1100, has an obviously higher rupture modulus than OCRE 1100 and OCRE 1200. Secondary-electron SEM images of the investigated samples are shown in Fig. 14. A comparison of the polished surfaces reveals that NCRE 1100 and OCRE 1200 have similar compactness and are denser than OCRE 1100, which indicates that a better fluxing effect is achieved for alkali-activated CFAs. A comparison of polished and etched surfaces shows that more needle-like mullite crystals are formed in NCRE 1100, suggesting that a better crystal skeleton effect is achieved for alkali-activated CFAs (class-I CFAs). Therefore, alkali activation not only contributes to CFA plasticity but also improves the fluxing and crystal skeleton effects, leading to enhanced performance of ceramic tile products.

OCRE 1100 (Table 7) were sintered at 1100 °C, the former showed significantly better performance and the latter showed clear signs of under-sintering. Comparison of NCRE 1100 and OCRE 1200 showed that both materials exhibit similar bulk densities and water absorptions, whereas the rupture modulus of NCRE 1100 is four times higher than that of OCRE 1200. The slightly higher linear shrinkage observed for NCRE 1100 has little influence on its performance and can be improved in the future. Therefore, alkali activation pre-treatment seems to benefit the macro-performances of sintered samples. Subsequently, comparisons were carried out at a microscopic level. The XRD patterns of NCRE 1100, OCRE 1100, and OCRE 1200 (Fig. 13) show that the main phases of NCRE 1000 correspond to nepheline (JCPDS card no. 01-083-2372) and mullite (JCPDS card no. 00-001-0613), and are completely different from the main phases, albite (JCPDS card no. 00-020-0572) and quartz (JCPDS card no. 01083-0539), of OCRE 1100 and OCRE 1200. The needle-like micro-

4. Conclusion This study describes two types of alkali-activated CFAs: class-I CFAs are coated by dense zeolite P phases, while class-II CFAs are uniform hydroxysodalite phases with porous microstructure. Untreated, class-I, and class-II CFAs exhibited entirely different thermal behaviors resembling those of quartz, clay, and feldspar, respectively, and were used in place of these materials to prepare fully ash-based ceramics with excellent post-sintering properties. At an optimum sintering temperature of 1100 °C, ceramic tiles with the highest rupture modulus of 50.1 MPa, largest bulk density of 2.5 g/ cm3, and lowest water absorption of 0% were produced. Although a slightly higher linear shrinkage of 17.6% was observed at 1100 °C, it had little influence on the performance of ceramic tiles and is subject to future improvement. In addition, alkali-activated fully fly-ash-based ceramic tiles not only required a lower firing temperature but also exhibited a wide sintering temperature range. In the ceramic sintering process, untreated, class-I, and class-II CFAs acted as quartz, clay, and feldspar, respectively. Compared to ceramic tiles derived solely from untreated CFA, alkali-activated ceramic tiles exhibited superior green and sintered performances. Hydrogen bonding interactions between

Intensity (a.u.) 5

Linear shrinkage (%)

10 15 20 25 30 35 40 45 50 55 60 65 70 75 80 85 90

Diffraction Angle,2θ/degree(CuKα) Fig. 13. XRD patterns of NCRE 1100, OCRE 1100 and OCRE 1200.

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Fig. 14. SEM images of NCRE 1100, OCRE 1100 and OCRE 1200: (a) NCRE 1100, (b) OCRE 1100, and (c) OCRE 1200; I – polished surface, II – polished and etched surface.

Thus, the proposed method for fabricating ceramic tiles from CFA waste opens up new ways of its effective utilization, relieving the associated environmental burden and creating huge economic benefits.

the surface crystal water molecules and OH groups were present in the alkali-activated CFA particles, imparting plasticity. Besides, the mullite skeleton effect of class-I CFAs and the fluxing agent effect of class-II CFAs further enhanced the ceramic tile performance.

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