Synthesis and characterization of glass-ceramics prepared from high-carbon ferrochromium slag

Abstract

Glass-ceramics have been successfully prepared from high-carbon ferrochromium slag (HCFS) and waste glass (WG), and the microstructural characterization and mechanical properties of the glass-ceramics were subsequently investigated. The development of HCFS-based glass-ceramics involves the nucleation and crystallization stages from the parent glass. With the increase in mass ratio of HCFS and WG (R(H/W)) from 0.60 to 1.67, the number of bridging oxygens of Si in the parent glass is reduced, as shown via Raman spectroscopy. Thus, their degree of polymerization decreases with it, and the temperature of nucleation and crystallization increase, which is consistent with the DSC results. The SEM images and EDS results indicate that the increasing value of R(H/W) decreases the crystal grain size and consequently increases the microhardness of the glass-ceramics. But the porosity simultaneously increases, which makes the bending strength increase at first and subsequently decrease. And the optimum properties of HCFS-based glass-ceramic samples in the present work are obtained when R(H/W) reaches 1.29, that is, a bending strength of 104 MPa and a microhardness of 9860 MPa.

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1. Introduction

Environmental and ecological considerations currently demand the use of industrial waste, such as slag, which is produced in traditional ferroalloy plants, particularly hazardous wastes from the production of ferrochromium. High-carbon ferrochromium slag (HCFS) is a by-product from the production of ferrochrome. Producing one ton of high carbon ferrochrome (HCFC) will generate 1.3 tons of HCFS.^1–4 HCFS is a heavy hazardous waste and a promising secondary resource,^5–7 in which the chromium content is approximately 1–10%. Considering the shortage of chromium resources in China, the comprehensive use of chromium slag becomes meaningful.⁸ The treatment is usually stockpiling or using chromium slag as landfill. Since the 1960s, chromium slag brick,⁹ calcium magnesium phosphate, glass colourant,¹⁰ colour cement made of reduced chromium slag¹¹ and mineral wool products of chromium slag¹² have been experimentally studied. However, such processes do not achieve total recycling of the chromium slag.^13,14

Glass-ceramics are fine-grained polycrystalline materials that are crystalized after heat-treating the glass phase of suitable compositions. Commonly, glass-ceramics are not fully crystalline, that is, they are approximately 50–95 vol% crystalline phase with residual glass. Glass-ceramics have superior mechanical properties compared to the parent glass. At present, glass-ceramic materials have been successfully prepared from solid waste. The commercial exploitation of solid waste, particularly blast furnace slag,^15–19 is well developed because the chemical compositions of blast furnace slag are stable and can be easily made into glass-ceramics.^20–22 The HCFS and blast furnace slag have similar compositions, such as Al₂O₃, SiO₂, CaO and MgO, which consist more than 90% content of the slag, so that the HCFS is reasonably suggested to synthesize glass-ceramics. A review of the phase analysis reveals that HCFS contains too many high-melting phases and requires high temperatures to melt, and therefore, some composition adjustments and sintering aids need to be involved, and complex reactions must be considered. According to the present study, there is limited research about synthesizing glass-ceramics from HCFS. The work of Wang et al.^23,24 preliminarily studied the preparation of glass-ceramics from ferrochrome slag and indicated that when the percentage of slag reached 50 wt% in the raw materials, the slag-based glass-ceramic had a maximum microhardness of 8000 MPa. Thus, this study is motivated to comprehensively use HCFS. The characterizations of HCFS are investigated, and a HCFS-based glass-ceramic is synthesized.

2. Experimental procedure

2.1. Sample preparation

The high-carbon ferrochromium slag (HCFS) used in this study was received from a ferrochromium plant, and the phases of the HCFS were obtained using an X-ray diffraction analysis and are shown in Fig. 1. There are ringwoodite ((Mg, Fe)₂SiO₄), spinel (MgAl₂O₄), and ferrochromium (FeCr₄) phases in the HCFS. The dominating phases are ringwoodite and spinel, whose melting points are above 1800 °C, so they require a special furnace to be melted. Chromium is mainly located in a variety of phases, such as spinel, ferrochrome, and so on.

Fig. 1 X-ray diffraction pattern of HCFS.

To produce glass-ceramics at an appropriate temperature that contains the high melting point of HCFS, the prime problem was to decrease the melting point of the glass-ceramics system. Hence, waste glass (WG) was selected for the composition adjustment, limestone and soda (≤75 μm) were used as the fining agent, and fluorite (≤75 μm) was added as the flux. All of the chemical compositions are shown in Table 1. Five samples with different mass ratios of HCFS and WG (R(H/W)) were investigated, and their compositions are shown in Table 2. The phase diagram of the glass-ceramics was calculated using Factsage 6.2 before its synthesis, and the results are shown in Fig. 2. The target phases were consisted of diopside, nepheline, and a small amount of pyroxene. The calculated melting point of this system is approximately 1500 °C. Hence, it is reasonable to suggest that the melting point of the glass-ceramics system can be decreased by composition adjustment from the thermodynamic calculation viewpoint.

Table 1 Chemical characterization of raw materials (wt%)

Raw materials	SiO2	MgO	Al2O3	CaO	Na2O	Cr2O3	Fe2O3	K2O	CaF2
HCFS	29.61	37.88	22.79	2.06	0.14	4.89	1.69	0.11
WG	75.68	3.40	0.94	8.95	9.73	0.024	0.27	0.68
CaCO3				55.44
Na2CO3					58.37
CaF2									98.5

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Table 2 Compositions of the samples (mass per g)

Sample no.	HCFS	Waste glass	CaCO3	Na2CO3	CaF2 (extra addition)	Mass ratio of HCFS and WG, R(H/W)
1#	30.0	50.0	15.0	5.0	5.26	0.60
2#	35.0	45.0	15.0	5.0	5.26	0.78
3#	40.0	40.0	15.0	5.0	5.26	1.00
4#	45.0	35.0	15.0	5.0	5.26	1.29
5#	50.0	30.0	15.0	5.0	5.26	1.67

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Fig. 2 Phase diagram of the CaO–MgO–SiO₂–Na₂O–Al₂O₃ system glass-ceramics, in which the Na₂O content is 7.83–5.91%, and the Al₂O₃ content is 7.31–11.68%.

Appropriate amounts of HCFS and WG were first ground to 75 μm and subsequently evenly mixed with the other materials, that is, limestone (CaCO₃), soda (NaCO₃) and fluorite (CaF₂). The mixture was placed into a corundum crucible, melted at 1550 °C in a muffle furnace and kept for 2 hours. Then, the samples were poured into a stainless-steel mould that was pre-heated to 500 °C and annealed at 500 °C for one hour in a muffle furnace to produce the parent glass.

The parent glass was subsequently investigated with TG-DSC to determine the nucleation and crystallization temperatures. Then, a two-stage heat treatment was used to prepare the glass-ceramics, which included the nucleation and crystallization processes. First, the parent glass was heated from room temperature to the nucleation temperature with a step of 5 °C per minute and maintained for 2 hours. Afterwards, the samples were continuously heated to the crystallization temperature with steps of 5 °C per minute and subsequently maintained for 2 hours. Finally, the glass-ceramics were obtained by furnace cooling. The nucleation time and crystallization time are both set on the preliminary work. The parent glass have been treated from 30 to 150 min, the highest mechanical properties were achieved after 120 min nucleation and crystallization, and thus the nucleation time and crystallization time are both set at 2 hours in this work, which are also corresponded to the report.^23,24

2.2. Characterization and mechanical-property measurements

All of the chemical compositions of the raw materials were detected using a Shimadzu-XRF-1800. The nucleation temperatures (T_n) and crystallization temperatures (T_g) were detected using Differential Scanning Calorimetry (DSC, NETZSCH-STA409). The microstructural characterizations of the as-cast parent glass and crystallized glass-ceramic samples were performed using a MLA250 scanning electron microscope (SEM) with an energy dispersive (EDS) attachment. X-ray diffraction (XRD) was performed using a TTR3 diffractometer with CuKα radiation at 40 kV and 30 mA. The parent-glass samples were also characterized according to their Raman spectra to analyse the structural variation with various R(H/W) values. For the Raman tests, a laser confocal Raman spectrometer (JY-HR800, Jobin Yvon Company) was used at room temperature with an excitation wavelength of 532 nm, and the light source was a 1 mW semiconductor.

The mechanical properties of the HCFS-based glass-ceramics were determined using the microhardness and 3-point bending tests. Vickers (or Knoop) microhardness tests were conducted using a Leica-VMHT30M tester. Optical mounts were performed using standard metallographic techniques, and a load of 100–300 g was used to indent their surfaces. To obtain reliable statistical data, at least 13 indentations were made on each sample. The empirical relation of Ponton and Rawlings, which uses fracture lines that emanate from Vickers diamond indentation, was used to determine the fracture toughness values.²⁵ Three-point bending tests were performed using a CDW-5, 5 KN tester with the sample size of 36 mm × 3 mm × 4 mm.

3. Results and discussion

3.1. Characterization of the parent glass

The parent glass was prepared from raw materials after melting, casting and annealing in the muffle furnace at 500 °C. To identify the phases of the parent glass, XRD was performed on the samples with different R(H/W) values, and the results are shown in Fig. 3. The results indicate that the glass phase is the main phase for all parent glass, which suggests that the melting point of the system sharply decreased. This result is consistent with the initial experimental design objective, and the results of the phase diagram were calculated using Factsage. However, some tiny peaks appear in the phase diagram, and the intensity of these peaks is proportional to the R(H/W) of the raw materials, as shown in Fig. 3. There are some high-melting-point materials (ringwoodite and spinel) in the parent glass that were not completely melted during the melting process; thus, trace amounts remained in the parent glass. However, these high-melting-point materials can act as nucleation seeds and are beneficial to the following nucleation process.

Fig. 3 X-ray diffraction patterns of the HCFS-based parent-glass samples with different R(H/W) values.

To analyse the molecular structures of the parent glass with different R(H/W), the samples were characterized by Raman spectra, and the results are shown in Fig. 4. The structural units in the silicate network are commonly described by the nomenclatures of Qⁱ(Si), where i represents the number of bridging oxygen atoms per Si atom. It can be seen from Fig. 4(a) that all parent glass have three regions: 200–600 cm⁻¹, 600–800 cm⁻¹, and 800–1200 cm⁻¹.

Fig. 4 Raman spectra and deconvolved results of HCFS-based parent-glass samples with different R(H/W) values: (a) Raman spectra of the parent glass; (b) R(H/W) = 0.60; (c) R(H/W) = 0.78; (d) R(H/W) = 1.00; (e) R(H/W) = 1.29; (f) R(H/W) = 1.67.

The peak at approximately 350 cm⁻¹, which is associated with the bending motion of Si–O–Si,^26,27 did not change with the increase in R(H/W), which reveals that the increasing content of HCFS does not affect the bending motion of Si–O–Si. The 600–800 cm⁻¹ band, which is attributed to the presence of Si–O stretching linkages,^28–31 became pronounced with the increase in R(H/W) because of the tiny increase in content of ringwoodite ((Mg, Fe)₂SiO₄) in the parent glass, as shown in Fig. 3. Meanwhile, the centre of this envelope curve shifted to a higher spectral frequency, which generally indicates a lower degree of polymerization of silicate structures. To further investigate the silicate structure, the Raman curves were deconvolved in the range of 800 to 1200 cm⁻¹ according to the rules proposed by Mysen et al.²⁷ It is assumed that the Raman curves are abided by Gaussian functions and that the bands are only fitted in the regions where obvious shoulders or peaks were observed or strictly proven by previous studies. Many studies^32–34 have investigated the Raman spectra of silicate glasses and melts and assigned the peaks at approximately 870 cm⁻¹, 960 cm⁻¹, 990 cm⁻¹, and 1050 cm⁻¹ to the stretching vibrations of Q⁰(Si), Q¹(Si), Q²(Si), and Q³(Si), respectively, each of which corresponds to a Gaussian fitting function. Then, four Gaussian functions were used in this study to fit the Raman curves of the parent-glass samples, and the fitting results are shown in Fig. 4(b)–(f).

According to the studies of Mysen and Frantz,^35,36 the calculation of the mole fractions of different structural units, i.e., Qⁱ(Si) (i = 0, 1, 2, and 3), could be calibrated by the Raman scattering coefficient, which is only related to Qⁱ(Si) species. Therefore, the following expression can be used to calculate the mole fractions:

X_i = θ_iA_i, i = 0, 1, 2, and 3 (1)

where X_i, θ_i, and A_i denote the mole fraction of Qⁱ(Si), Raman scattering coefficient, and band area of Qⁱ(Si), respectively. The reciprocal of θ_i, which is S_i, was obtained by carefully analysing various silicate systems.³⁷ With the assumption that S₀ is equal to 1, S₁, S₂, and S₃ can be separately calculated as 0.514, 0.242, and 0.09, respectively. These constants, which are consistent with those proposed by Frantz,³⁵ were used in this study. Then, the mole fractions of different Qⁱ(Si) species were calculated as follows:

The variation of the mole fractions of different Qⁱ(Si) species with different R(H/W) values is shown in Fig. 5. The stretching vibrations of Q²(Si) and Q³(Si) decrease with the increase in R(H/W), but the stretching vibrations of Q¹(Si) becomes higher, so that the number of bridging oxygen atoms of Si reduces. The number of bridging oxygen atoms of Si is the main factor that affects the degree of polymerization of the silicate structures in the melting glass-ceramics: fewer bridging oxygen atoms of Si in the system corresponds to a lower degree of polymerization. Thus, when R(H/W) increases, the degree of polymerization of the parent glass decreases, which may affect the nucleation and crystallization process and, particularly, their temperatures.

Fig. 5 Mole fraction variation of various Qⁱ(Si) (i = 0, 1, 2 and 3) with different R(H/W).

3.2. Nucleation and crystallization temperature detection

Five parent glasses with different mass ratios of HCFS and WG (R(H/W)) were subsequently investigated using Differential Scanning Calorimetry (DSC) to determine the nucleation and crystallization temperatures. Normally, the parent-glass transformation into glass ceramics involves two steps: nucleation and crystallization. Nucleation is the initial process to form a new phase for the two-step heat treatment. This step requires an activation energy to overcome the resistance from the parent phase, and it appears as an endothermic peak in the DSC curve. On the contrary, the crystallization process ensures the attainment of the lowest energy of the system and is an exothermic process. Thus, it is visualized as an exothermic peak in the DSC curve. The DSC results of the parent-glass samples are shown in Fig. 6. Fig. 6(a) presents the entire range of DSC curves and shows that there are endothermic and exothermic peaks in the DSC curves that correspond to the nucleation and crystallization temperatures of the parent glass, respectively.

Fig. 6 DSC curves of the HCFS-based parent-glass samples with different R(H/W): (a) DSC curves of the HCFS-based parent glass samples ranging 0–1000 °C, (b) the nucleation temperature of the samples ranging 580–780 °C, (c) the crystallization temperature of the samples ranging 760–920 °C.

To determine the nucleation (T_n) and crystallization (T_g) temperatures, the partially enlarged DSC curves are shown in Fig. 6(b) and (c). As shown in Fig. 6(b), a series of small endothermic peaks that start from 630 °C are observed in each curve. The marks on each line are the starting temperature for the nucleation of glass-ceramics and defined as the T_n temperature to consider the lowest energy cost. There is a sharp exothermal peak in each DSC curve shown in Fig. 6(c), and the temperature of the peak is the best efficiency point of the crystallization, so it is defined as the T_g temperature. Thus, the optimum nucleation and crystallization temperatures of each sample are gathered and shown in Table 3. The nucleation and crystallization temperatures increase with the increase in R(H/W), which is attributed to the molecular structure of the parent glass. When R(H/W) increases, the number of bridging oxygen atoms of Si in the glass phase of the parent glass decreases according to the Raman spectra, as shown in Fig. 5. Thus, the heat and mass transfers become difficult, that is, the parent glass will require much more energy for nucleation and crystallization. Hence, the nucleation and crystallization temperatures must increase.

Table 3 Nucleation and crystallization temperatures of the samples

R(H/W)	0.60	0.78	1.00	1.29	1.67
Nucleation temperature/°C	635	642	648	652	658
Crystallization temperature/°C	835	838	842	855	859

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Considering the small-scale variation of the nucleation and crystallization temperatures and the large scale of content changing of HCFS in the raw materials of 30–50%, the occurrence of high-melting-point phases, such as ringwoodite and spinel, and the heat-treatment condition of the glass-ceramics in this study are obviously beneficial to large-scale industrial production.

3.3. Characterization of glass-ceramics

The five parent-glass samples were separately heat-treated considering their respective nucleation and crystallization temperatures (Table 3), and the glass-ceramics were finally prepared. The phase analysis of the HCFS-based glass-ceramic samples was performed using XRD, and their results are shown in Fig. 7. Several new phases are formed in the glass-ceramics during the heat-treatment of the parent glass: diopside (CaMg(SiO₃)₂), nepheline ((Na, K)AlSiO₄), pyroxene (Ca(Mg, Fe)Si₂O₆, and Ca(Mg, Al, Fe)(Al, Si)₂O₆).

Fig. 7 XRD patterns of the HCFS-based glass-ceramic samples with different R(H/W) values.

Fig. 8 presents the backscattered electron images of the glass-ceramic samples with different R(H/W) values, and the EDS results of some randomly selected points are shown in Table 4. In Fig. 8, there are two phases in the glass-ceramics: white and grey phases. According the EDS results in Table 4, the white phase corresponds to CaMg(SiO₃)₂ and the grey phase is (Na, K)AlSiO₄. A comparison of the XRD and SEM results shows that CaMg(SiO₃)₂ and (Na, K)AlSiO₄ are the main phases, whereas Ca(Mg, Fe)Si₂O₆ and Ca(Mg, Al, Fe) (Al, Si)₂O₆ are the minor parts of the glass-ceramics. The grain size was counted from Fig. 8, and the results are shown in Fig. 9. With the increase in R(H/W), that is, the increasing content of HCFS in raw materials, the grain sizes of the white and grey phases decrease. This result is attributed to the increase in content of nucleation agent with the increase in the content of HCFS. Thus, the crystallization of the parent glass will become easier and quicker, which subsequently causes smaller crystal grain size.

Fig. 8 SEM images of glass-ceramic samples with different R(H/W) values: (a) R(H/W) = 0.60; (b) R(H/W) = 0.78; (c) R(H/W) = 1.00; (d) R(H/W) = 1.29.

Table 4 EDS analysis of the HCFS-based glass-ceramic samples

Point no.		Element norm. C (wt%)
		O	Mg	Al	Ca	Na	Cr	Fe	Si	Phase
1	Grey phase	39.76	4.04	14.93	3.41	19.07			18.97	(Na, K)AlSiO4
2		38.31	4.48	16.35	2.08	19.54			19.24
3		39.42	4.76	14.74	4.38	18.14			18.57
4	White phase	37.53	11.54	4.82	20.90	0.90	0.80	1.26	22.26	CaMg(SiO3)2
5		37.40	11.71	4.63	20.07	2.14	0.75	1.14	22.16
6		36.53	11.48	4.46	21.70	1.29	0.65	1.27	22.62

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Fig. 9 Grain size distributions of the glass-ceramic samples with different R(H/W) values.

Similarly, when the nucleating agent increases with the increase in the content of HCFS, the average crystal size decreases and crystallization accelerates. This phenomenon makes some gases unable to completely escape and become trapped inside the sample, which leads to a high degree of porosity, as shown in Fig. 10. In Fig. 11, the porosity of the glass-ceramics initially increases with the increase in R(H/W), but when R(H/W) is above 1.29, the porosity of the glass-ceramics suddenly increases.

Fig. 10 Porosity measurement of glass-ceramic samples with different R(H/W) values: (a) R(H/W) = 0.78; (b) R(H/W) = 1.00; (c) R(H/W) = 1.29; (d) R(H/W) = 1.78.

Fig. 11 Porosity distribution of the glass-ceramic samples with different R(H/W) values.

3.4. Phase evolution analysis

Fig. 1 and Table 1 indicate that the raw materials consist of ringwoodite ((Mg, Fe)₂SiO₄), spinel(MgAl₂O₄), and ferrochromium(FeCr₄), silicate, limestone (CaCO₃), and soda (NaCO₃). When they are melted, most compounds in the raw materials, such as CaCO₃ and NaCO₃, are decomposed to oxides, as shown in eqn (3) and (4). Ferrochromium is oxidized as shown in eqn (5).

CaCO₃ = CaO + CO₂ (3)

Na₂CO₃ = Na₂O + CO₂ (4)

4CrFe₄ + 11O₂ = 16FeO + 2Cr₂O₃ (5)

2MgO + 2Cr₂O₃ = 2MgCr₂O₄ (6)

Chromium sesquioxide would further react with MgO and form the chrome spinel phase (eqn (6)). Then, the trace of high melting-point materials, that is, (Mg, Fe)₂SiO₄, MgAl₂O₄, Mg(Al, Cr)₂O₄, and Cr₂O₃, remains in the parent glass during the melting process and acts as nucleation sites. During the crystallization stage of the heat treatment, Fig. 7 and Table 4 show that CaO, MgO, FeO and SiO₂ further react and form CaMg(SiO₃)₂ and Ca(Mg, Fe)Si₂O₆, as shown in eqn (7) and (8). The reactions among Na₂O, K₂O, Al₂O₃ and SiO₂ are described in eqn (9) and form (Na, K)AlSiO₄. These substances constitute the main phases of the glass-ceramics. (Na, K)AlSiO₄ may also react with sylvite or sodium salt and generate a low-melting solid solution, which is likely the main source of the glassy phase.

CaO + MgO + 2SiO₂ = CaO·MgO·2SiO₂ (7)

CaO + MgO(FeO) + 2SiO₂ = Ca(Mg, Fe)Si₂O₆ (8)

Na₂O(K₂O) + Al₂O₃ + 2SiO₂ = 2(Na, K)AlSiO₄ (9)

3.5. Mechanical properties measurement

In addition to the aforementioned microscopy investigations, the macroscopic mechanical properties can also reflect different characteristics of the samples. Hence, a combination of mechanical tests, that is, the microhardness and 3-point bending mechanical-property measurements, was performed.

Fig. 12 is the representation of Vickers microhardness values of the HCFS-based glass-ceramic samples. The hardness values clearly continuously increase with the increase in R(H/W), possibly because of the increasing concentration of the nucleation agent. When R(H/W) increases, the average crystal size correspondingly decreases, as shown in Fig. 9, and further leads to the increase of the amount of grain boundary, which enhances the resistance to indentation. When the value of R(H/W) increases to 1.67, the microhardness of the glass-ceramic sample reaches a maximum of 9860 MPa, which is higher than that of other glass-ceramics made from metallurgical slags with a maximum value of 7000 MPa.²⁵

Fig. 12 Microhardness of the HCFS-based glass-ceramic samples with different R(H/W) values.

Fig. 13 shows an error-bar representation of the bending strength of the HCFS-based glass-ceramic samples. Here, with the increase in R(H/W), the bending strength of the samples initially increases to the top (approximately 110 MPa) and subsequently slightly decreases. This phenomenon is attributed to the increase in the porosity of the specimens when the HCFS content increases, as shown in Fig. 10. When R(H/W) is beyond 1.29, the bending strength starts to decrease, whereas the porosity starts to sharply increase. Thus, a moderate porosity, that is, when the porosity is below 1.5%, has a notably small effect on the bending strength. However, an abundant porosity will deteriorate it. When the value of R(H/W) reaches 1.29, the bending strength of glass-ceramic sample reaches the maximum of 104 MPa. The bending strength of glass-ceramics that are made from metallurgical slag is generally 56.6–97 MPa,^25,38−41 but some kind of the slag-based glass-ceramics that just focused on the enhancement of the bending strength has reached 136 MPa.^25,42 Hence, the HCFS-based glass-ceramics in this study show a comparable bending strength with the normal slag-based glass-ceramics. In addition, HCFS-based glass-ceramic materials are produced from the toxic substances and wastes from the normal procedure. Thus, it is a good reference to recycle this kind of material because of its low-cost raw material, good properties and notably effective manner in treating toxic substances.

Fig. 13 Bending strength of the HCFS-based glass-ceramic samples with different R(H/W) values.

4. Conclusions

In this paper, HCFS-based glass-ceramics are prepared from HCFS, waste glass, limestone, soda and fluorite using the two-stage heat treatment, and the microstructural characterization and mechanical properties of the glass-ceramics are investigated. With the increase in R(H/W), the Raman spectra of the parent glass indicates that the number of bridging oxygens of Si in the glass phase and the degree of polymerization of the silicate structures in the parent glass decrease, which makes heating and mass transfer difficult. Thus, the nucleation and crystallization temperature increases, which is consistent with the DSC results. The SEM images and EDS results of the glass-ceramics indicate that with the increase in R(H/W), the increasing content of nucleation agent results in the decrease of the average crystal size and rapid crystallization, which increases the porosity of the glass-ceramics. Moreover, the increase in R(H/W) directly affects the mechanical properties of the glass-ceramics. The microhardness and bending strength continuously increase with the increase in R(H/W), but when the value of R(H/W) exceeds 1.29, the bending strength starts to decrease. The optimal bending strength and microhardness of HCFS-based glass-ceramics are 104 MPa and 9860 MPa, respectively.

Acknowledgements

The National Science Foundation of China (No. 51372019, 51272025, 50874013) and the National Science and Technology Supporting Program (No. 2013AA032003) are duly acknowledged for their financial support.

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