Effect of Al2O3/SiO2 on the structure and properties of blast furnace slag glass-ceramics

doi:10.21203/rs.3.rs-131872/v1

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Effect of Al2O3/SiO2 on the structure and properties of blast furnace slag glass-ceramics

https://doi.org/10.21203/rs.3.rs-131872/v1

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Different Al₂O₃/SiO₂ glass-ceramics were prepared from blast furnace slag by traditional sintering method. The structure and properties of glasses or glass-ceramics were investigated by DSC, XRD, SEM, FTIR, ²⁷Al MAS NMR. The DSC results showed that with the increase of Al₂O₃/SiO₂, the glass transition temperature (T_g) first increased and then decreased, reached the minimum when Al₂O₃/SiO₂ was 0.34. The volume density, bending strength and microhardness of glass-ceramics also showed the same variation rule. The FTIR and ²⁷Al MAS NMR spectra results revealed this phenomenon. When Al₂O₃/SiO₂ was 0.19, a large amount of Si⁴⁺ was added to the glass network to make the structure dense. With Al₂O₃/SiO₂ increased from 0.24 to 0.34, the amount of [AlO₆] in the glass increased while [AlO₄] decreased, and the degree of network polymerization of the glass decreased; as Al₂O₃/SiO₂ further increased from 0.34 to 0.39, [AlO₄] increased and [AlO₆] decreased in the glass, and the degree of network polymerization of the glass increased. The XRD results showed that the crystal phase of the glass-ceramics was composed of gehlenite, diopside and hyalophane. Moreover, with the increase of Al₂O₃/SiO₂, the gehlenite content in the glass-ceramics increased, while the content of diopside and hyalophane decreased.

Ceramics

Architecture, Design and Planning

glass-ceramics

blast furnace slag

Al2O3/SiO2

crystallization

sintering

As a basic industry in China, steel industry produces about 200 billion tons of blast furnace slag (BFS)every year[1, 2]. Part of the BFS is directly used in low value-added products such as cement, slag brick and subgrade materials, however, a lot of BFS is dumped, which not only occupies land resources, but also causes serious pollution and waste[3]. In addition to some high titanium BFS, the main composition of BFS is CaO, MgO, Al₂O₃ and SiO₂, with the proportion of these four oxides exceeding 90 wt.%, which is similar to the composition of CMAS glass-ceramics, therefore, BFS is an excellent raw material for glass-ceramics[1, 3].

Glass-ceramics is a kind of polycrystalline solid material containing a large number of microcrystalline and glass phases, which is made by controlling crystallization of specific basic glass[4]. The special structure of glass-ceramics makes it have a variety of excellent properties, such as good insulation performance, high mechanical strength, low dielectric loss, etc.[5], which is a kind of products with high added value materials[4]. The preparation of glass-ceramics with BFS as raw material provides a new way for the recovery and utilization of BFS, which can alleviate the stacking and pollution of BFS and effectively improve the added value of the products.

In the composition of BFS, the content of Al₂O₃ and SiO₂ is more than 40 wt.%, in which SiO₂ acts as the glass network former, Si⁴⁺ is connected in the shape of silicon-oxygen tetrahedron to form the basic network structure of glass. In the meantime, there are different degree of polymerization and depolymerization with different content of SiO₂ in glass, which is directly related to the heat treatment, microstructure and crystal phase [6]. Al₂O₃ is a network intermediate oxide, on the one hand, Al³⁺ can enter the glass network as [AlO₄], making the structure compact; on the other hand, Al³⁺ can also exist in the glass as [AlO₆], causing the depolymerization of the glass network and making the structure loose[7]. In partial systems, Al³⁺ can also act as the element of partial crystal phases when the content of Al₂O₃ is high[8]. Wentao Zhang et al. studied the influence of B₂O₃/Al₂O₃ ratio and Al/Na ratio on BFS glass-ceramics, and indicated that Al₂O₃ improved the thermal stability of glass, meanwhile, the crystallization activation energy and network structure were also affected with the change of Al₂O₃ content[9, 10]. Lishun Chen et al. also confirmed that, in the study of glass-ceramics with high calcium system, the crystallization activation energy rose and the main crystal phase changed from akermanite to gehlenite with the increase of Al₂O₃[11]. Jinshu Cheng et al. showed that the workability and kinetic fragility of glass were linked with the structure, with the increase of SiO₂/Al₂O₃, the non-bridging oxygen in the glass network decreased, and the structure of the glass network became compact, resulting in the increase of the viscosity of the glass melt, meanwhile, the change of SiO₂/Al₂O₃ ratio were responsible for the increase of fragility[12]. The research of Janusz Partyka et al. also showed that the variation of SiO₂/Al₂O₃ ratio directly determined the type and content of crystal phase[13]. Therefore, it is important to investigate the influence of SiO₂/Al₂O₃ on the structure and properties of BFS glass-ceramics.

In this paper, based on 60 wt.% BFS, the effects of different SiO₂/Al₂O₃ on the crystal phases, structure and properties of glass-ceramics were studied. In order to further reveal the influence of the Al₂O₃/SiO₂ on the glass network, Fourier Transform Infrared Spectrometer and Solid State Nuclear Magnetic Resonance Spectrometer were used to deeply investigate the microstructure of glass and glass-ceramics.

2.1 Raw materials and experimental formula

The BFS used in this experiment was from China Baowu Iron & Steel Group. The chemical composition of BFS was measured by X-ray Fluorescence (XRF, Zetium, PANalytical B. V.). This experiment was based on 60 wt.% BFS and supplemented by pure chemical reagent to prepare glass-ceramics. The high content of alkaline earth metal oxide in BFS leads to the high alkalinity of base glass, which is not conducive to the sintering process due to the fast crystallization[14]. Therefore, the reagents were mainly acid oxides such as SiO₂ and Al₂O₃, and a small amount of oxides such as Na₂O, K₂O, BaO and B₂O₃ were added to adjust the glass melting, clarification, homogenization and sintering process. The oxides were introduced by Na₂CO₃, K₂CO₃, Ba₂CO₃ and H₃BO₃ respectively. The formula and the specific oxide composition of the BFS were shown in Table 1. The SiO₂/Al₂O₃ ratios in the base glass were different, which were 0.19(A1), 0.24(A2), 0.29(A3), 0.34(A4) and 0.39(A5), respectively.

Table 1

Oxide composition of glasses and BFS (wt.%).

	Al₂O₃	SiO₂	CaO	MgO	BaO	Na₂O	K₂O	B₂O₃	others	Al/Si
BFS	15.12	31.92	39.81	8.54	0.12	0.39	0.58	/	3.52	0.47
A1	9.07	48.19	23.89	5.12	3.00	3.00	3.00	2.00	2.73	0.19
A2	11.08	46.18	23.89	5.12	3.00	3.00	3.00	2.00	2.73	0.24
A3	12.87	44.39	23.89	5.12	3.00	3.00	3.00	2.00	2.73	0.29
A4	14.53	42.73	23.89	5.12	3.00	3.00	3.00	2.00	2.73	0.34
A5	16.07	41.19	23.89	5.12	3.00	3.00	3.00	2.00	2.73	0.39

2.2 Preparation of glass-ceramics

In this paper, glass-ceramics were prepared using conventional sintering method. The melting temperature of the base glass was 1450 ℃, the holding time was 1 h. The melted glass liquid was poured into the clear water to obtain the base glass slag. And the glass slag was ground with planetary grinding ball for 30 min and then passed through a 200 mesh sieve to obtain the parent glass powder which was pressed into 4 mm × 40 mm strip samples under a pressure of 50 MPa, and then the samples were sintered in a resistance furnace, the sintering temperature was obtained by thermal analysis, the heating rate was 10 ℃/min, the holding time was 1.5 h.

2.3 Characterizations

Thermal analysis of base glass powder was performed using Differential Scanning Calorimetry (DSC, STA449F3, NETZSCH), the temperature range was 0-1000 ℃ in the air and the heating rate was 10 ℃/min. The crystal phase was measured by X-ray Diffractometer (XRD, D8 Advance, BRUKER AXS),the scanning range was 10–70º.The micromorphology of the glass-ceramics was observed by Field Emission Scanning Electron Microscope (FESEM, Zeiss Ultra Plus, Zeiss Germany), the polished surface was etched with HF solution (4%) for 40 s. The microstructure of both glasses and glass-ceramics was measured by Fourier Transform Infrared Spectrometer (FTIR, Nexus, Thermo Nicolet) at range of 400–2000 cm^− 1. The ²⁷Al NMR spectra of the base glasses were measured by Solid State Nuclear Magnetic Resonance Spectrometer (NMR, AVANCE Ⅲ).

The bending strength (M, MPa) of glass-ceramics was measured by three-point test method, and the calculation formula was as follows:

$$M=\frac{3FL}{2b{h}^{2}}$$

where F is the fracture load (N), L is the span (mm), b is the fracture width (mm), and h is the fracture thickness (mm). The volume density of glass-ceramics was measured by Archimedes-drainage method. The indentation method was used to measure the Vickers hardness of the glass-ceramics, the loading force was 0.98 N, the loading time was 10 s.

3.1 DSC analysis

Figure 1 showed the DSC graph of parent glasses with different Al₂O₃/SiO₂ at heating rate of 10 ℃/min. Table 2 showed the specific characteristic temperature values for A1-A5, where T_g was the glass transition temperature, T_c1 was the first crystallization temperature, and T_c2 was the second crystallization temperature. In Fig. 1, there were two obvious exothermic peaks within the range of 780–900 ℃, among which the first crystal peak shape of A1 parent glass was relatively flat.

Table 2

Characteristic temperature of glasses with different Al₂O₃/SiO₂ in DSC.

	T_g(℃)	T_c1(℃)	T_c2(℃)
A1	613.7	851.8	900.8
A2	594.7	799.0	878.8
A3	568.8	790.8	887.1
A4	413.1	783.7	886.2
A5	565.2	783.5	886.0

The DSC curve of the base glasses showed that the peak of T_c1 became sharp and moved towards low temperature gradually for A1-A5, while T_c2 showed a trend of moving towards high temperature for A2-A5. While the charge of Al³⁺ is less than that of Si⁴⁺, and the self-diffusion coefficient of Al³⁺ is higher than that of Si⁴⁺[15]. Therefore, the higher Al₂O₃ content in the glasses was more conducive to particle diffusion in the crystallization process, which reduced T_c1 for A1-A5[15]; then the crystal phase was precipitated on the surface of the glass particle, which increased the viscosity of the glasses, hindered the further diffusion of the particle, and led to the rise of T_c2[16]. The first crystallization peak of the A1 glass was not obvious and the T_c2 was high, which may be related to the high SiO₂ content in the A1 glass, as Si⁴⁺ can gather the network, increase the viscosity of the glass, hinder the diffusion of particles inside the glass, and increase the crystallization temperature[17]. Figure 1 and Table 2 showed that T_g decreased first from A1 to A5 (reached the minimum in A4) and then increased, which indicated that the thermal stability of glasses deduced first and then rose with the increase of Al₂O₃/SiO₂. The thermal stability of glasses is closely related to the network structure of glass[18], which also indicated that the network polymerization degree of glasses showed a trend of first decreasing and then increasing in series A glasses. In this experiment, according to the DSC curve (Fig. 1), in order to ensure that the crystallization process of each sample was fully carried out, the sintering temperature of the sample was set at 890 ℃ and the holding time was 1.5 h.

3.2 Crystal phase analysis

The XRD graph of glass-ceramics with different Al₂O₃/SiO₂ was illustrated in Fig. 2. It shows that the peaks of A1-A5 glass-ceramics were basically the same. The analysis by Jade 6.5 indicated that the crystal phase of A series glass-ceramics was composed of the main crystal phase gehlenite(Ca₂(Al(AlSi)O₇,PDF#74-1607), the secondary crystal phase diopside(CaMgSi₂O₆༌PDF#74-1607) and hyalophane (K.₆Ba.₄Al_1.42Si_2.58O₈༌PDF#72-1497). Figure 2 also showed that the diffraction peak intensity of the main crystal phase increased while that of the secondary crystal phase decreased gradually for A1-A5.

It means that the content of gehlenite in the glass-ceramics increased while diopside and hyalophane decreased with the increase of Al₂O₃/SiO₂. The degree of polymerization of anionic groups is different between the primary phase (gehlenite) and the secondary phase (diopside and hyalophane). The reaction process is as follows:

2Ca²⁺ + Al³⁺ + (AlSiO₇)⁷⁻ → Ca₂Al(AlSi)O₇

Ca²⁺ + Mg²⁺ + (Si₂O₆)⁴⁻ → CaMgSi₂O₆

0.6K⁺ + 0.4Ba²⁺ + (Al_1.42Si_2.58O₈)^1.42− → K.₆Ba.₄Al_1.42Si_2.58O₈

The Si(Al)/O ratio of anion group [AlSiO₇] is smaller than that of [Si₂O₆] and [Al_1.42Si_2.58O₈], indicating a lower degree of polymerization of [AlSiO₇][19, 20]. In terms of element composition, Al³⁺ directly participated in the formation of the main crystal phase gehlenite, therefore, the increase of Al₂O₃ content was conducive to the precipitation of gehlenite. As a network intermediate, Al³⁺ can not only replace Si⁴⁺ to participate in the glass network and crystal phase constitution with [AlO₄], but also destroy the glass network structure in the form of [AlO₆][7, 8]. On the one hand, with the increase of Al₂O₃/SiO₂, Si⁴⁺ in the glass decreased and Al³⁺ increased, [AlO₆] played a role in breaking the glass network, reducing the degree of polymerization of glass network and facilitating the precipitation of gehlenite with low degree of anionic polymerization; on the other hand, the increase of Al₂O₃ rose the number of Al³⁺ substituted for Si⁴⁺ in the glass network, which was conducive to the formation of [AlSiO₇] and promoted the precipitation of gehlenite. On account of the large amount of precipitation of gehlenite, Si⁴⁺ content in the glasses was decreased, which led to the reduction of the anion group of the secondary crystal phase and reduced the precipitation of diopside and hyalophane.

The SEM graph was showed in Fig. 3. It illustrated that there were a large number of strip or gridded crystals, and the XRD analysis results showed that the crystal phase was gehlenite; a few clusters of granular crystals were seen in A1-A3, which were speculated to be diopside based on the XRD results. In the A4 and A5 samples, no other crystal phases were found, which was mainly attributed to the high precipitation of gehlenite and the low precipitation of secondary phase. In the SEM graph of A1, there were more holes left by HF solution erosion. With the increase of Al₂O₃/SiO₂, the holes decreased gradually, indicating that the increase of crystal evolution led to the decrease of glass phase for A1-A5. When the Al₂O₃/SiO₂ increased to 0.34 (A4), it was seen that the growth of crystal grains was united and arranged in a certain direction, and the defect area without crystal phase existed; when the Al₂O₃/SiO₂ ratio continued to increase to 0.39 (A5), the crystal grains continued to grow into long strips, and the long strips growing in different directions were interspersed and nested together.

3.3 FTIR analysis

FTIR is a common method used to test the internal structure of glass and glass-ceramics. Figure 4 was the FTIR of the base glasses with different Al₂O₃/SiO₂ at 400 cm^− 1-1400 cm^− 1. In Fig. 4, the shape of vibration absorption peak of A1-A5 parent glasses was substantially the same. There were mainly three wide vibration absorption bands, which were in the range of 400–600 cm^− 1, 600–800 cm^− 1 and 800–1200 cm^− 1 respectively. The irregular arrangement of ions (ion clusters) in the glasses and the existence of non-bridging oxygen bonds made the angle and length of Si-O bond change, which made the infrared vibration absorption peak shift to a certain extent and become gentle and broad[21, 22].

In the infrared vibration spectra of glasses, the vibration absorption band in the range of 400–600 cm^− 1 was attributed to the bending vibration of Si–O–Si, Si–O–Al, O–Si–O and O–Al–O[23, 24].

The absorption peak in the range of 600–800 cm^− 1 was the symmetric stretching vibration of Si–O–Al, which was mainly attributed to the stretching vibration of the [AlO₄] and stretching vibration of the bridge oxygen bond in the glass network structure[11]. From A1 to A5, the vibration absorption peak near 720 cm^− 1 showed a small shift towards low wavenumbers, and the absorption intensity was enhanced, indicating that the Si–O–Al bond with lower vibration frequency in the glass network increased[23]. This was due to the fact that with the increase of Al₂O₃/SiO₂, more and more Al³⁺ replaced Si⁴⁺ to form tetrahedron and entered into the glass network, forming Si–O–Al bond with [SiO₄].

The absorption vibration peak in the range of 800–1200 cm^− 1 was mainly related to the [SiO₄], which was attributed to the asymmetric stretching vibration of Si–O–Si, the symmetric stretching vibration and asymmetric stretching vibration of O–Si–O, the absorption was strong and broad, which was caused by the superposition of different peaks[25–27]. The vibration absorption peak in the range of 800–1200 cm^− 1 can be decomposed into stretching vibration of silicon-oxygen tetrahedron with different degree of polymerization by Gaussian function (Fig. 5), whose symbol is Qⁿ, where n is the number of bridging oxygen (O_b) in the silicon-oxygen tetrahedron (n = 0,1,2,3,4,5)[19, 28]. Figure 5 was the Gaussian deconvolution diagram of the infrared vibration absorption peak at the range of 800–1200 cm^− 1 of the series A glasses, and table 3 was the specific peak values of each Qⁿ. The results of Fig. 5 and Table 3 showed that the vibration absorption peaks of A1-A5 glasses at 800–1200 cm^− 1 were all decomposed into 5 vibration absorption peaks corresponding to Q⁰, Q¹, Q², Q³ and Q⁴ respectively. Moreover, the peak positions of Qⁿ were relatively similar for each sample of series A glasses, indicating that the network structure of the base glass did not change significantly with the increase of Al₂O₃/SiO₂.

Figure 6 was the variation distribution chart of peak position of Qⁿ for A1-A5, which indicated that the peak position of Qⁿ was firstly shifted to direction of the low wavenumbers and then to the high wavenumbers, among which the peak position of Qⁿ reached the lowest when the Al₂O₃/SiO₂ was 0.34 (A4). The vibration absorption peak of Qⁿ is closely related to the glass network[29]. When the glass network tends to be close, the vibration peak is shifted to the direction of high wavenumbers; on the contrary, when the glass network structure tends to be loose, the vibration peak moves to the direction of low wavenumbers[29]. The curve in Fig. 6 illustrated that the glass network structure tended to be loose with the increase of Al₂O₃/SiO₂ from 0.19 to 0.34; subsequently, the glass network structure tended to be close with the Al₂O₃/SiO₂ continues to increase from 0.34 to 0.39.

Table 3

Wavenumbers of Qⁿ in glasses with different Al₂O₃/SiO₂.

		A1	A2	A3	A4	A5
Wavenumbers (cm^− 1)	Q⁰	849.7	847.2	847.1	844.1	844.1
	Q¹	894.6	890.6	889.0	883.5	886.2
	Q²	970.2	962.4	956.6	948.4	959.6
	Q³	1069.5	1061.6	1054.5	1049.2	1060.9
	Q⁴	1144.7	1140.1	1136.5	1132.2	1135.9

The infrared vibration absorption spectra of the glass-ceramics sintered at 890 ℃ was showed in Fig. 7. It indicated that the peak range of the glass-ceramics sample was roughly the same as that of the base glasses (Fig. 4), which was still divided into three vibration absorption bands, however, there were some sharp absorption peaks which were not found in the infrared vibration absorption spectra of the base glasses. This was mainly attributed to the generation of crystal phase, which made the originally covered absorption peaks appear[21, 23]. The position and shape of the infrared absorption peaks of A1-A5 were basically the same, indicating that the crystal phase of glass-ceramics with different Al₂O_3/SiO₂ were basically the same, and the XRD analysis also proves this conclusion.

In the range of 400-600cm^− 1, the vibration absorption peak at 470 cm^− 1 was divided into two absorption peaks: 472 cm^− 1 and 424 cm^− 1. The peak at 472 cm^− 1 was attributed to the bending vibration of Si–O–Si in the residual glass phase and diopside phase (a small amount of hyalophane), the peak at 424 cm^− 1 was attributed to the bending vibration of Si–O–Al bond in gehlenite[24, 30]. The intensity of the vibration absorption peak at 424 cm^− 1 gradually increased for A1-A5, which indicated that the content of gehlenite increased in glass ceramics with the increase of Al₂O₃/SiO₂. There were two vibration absorption peaks at 544 cm^− 1 and 579 cm^− 1, which were not found in the base glasses in Fig. 4. This was due to the coupling effect between the bending vibration of O-Si-O in the crystal phase and the stretching vibration of Ca–O[24, 27]. In the range of 600–800 cm^− 1, in addition to the vibration absorption peak near 714 cm^− 1(base glasses also had the same vibration absorption peak), two new vibration absorption peaks appeared at 682 cm^− 1 and 630 cm^− 1. The peak at 714 cm^− 1 and 682 cm^− 1 were attributed to the symmetric stretching vibration of Si–O–Si and the stretching vibration of Si–O–Al respectively; the peak at 630 cm^− 1 was related to the stretching vibration of Al–O in the aluminum-oxygen tetrahedron[22, 27, 30]. In the range of 800–1200 cm^− 1, the vibration absorption peak near 1000 cm^− 1 split into 1028 cm^− 1, 985 cm^− 1 and 939 cm^− 1 due to crystal phase precipitation, which were associated with the stretching vibration of Si–O in [SiO₄] (asymmetric stretching vibration of Si–O–Si, symmetric stretching vibration and asymmetric stretching vibration of O–Si–O)[22, 26, 27].

3.4 ²⁷Al NMR analysis

To further study the coordination of Al³⁺ in glass, the ²⁷Al MAS NMR spectra were used to investigate the glass structure, as shown in Fig. 8, there was an obvious wide peak in the range of -25-100 ppm. In aluminosilicate system, Al³⁺ has three coordination modes: [AlO₄], [AlO₅], [AlO₆]. Therefore[31, 32], ²⁷Al NMR spectra were deconvoluted by Gaussian function, and the results were shown in Fig. 9 and Fig. 10.

The results in Fig. 9 showed that the ²⁷Al NMR spectra deconvolution of series A glasses were similar, which was consisted of three distinct peaks [AlO₄], [AlO₅] and [AlO₆], in which the area % of [AlO_n] (n = 4, 5, 6) was related to its content[7, 9, 10]. In Fig. 10, the area % of [AlO₆] first increased and then decreased for A2-A5, and reached the maximum in A4 (the Al₂O₃/SiO₂ was 0.34); on the contrary, the area % of [AlO₅] and [AlO₄] first decreased and then increased, and reached the minimum in A4. The area % of [AlO₆] in A1 glass was relatively large, while area % of [AlO₅] and [AlO₄] were relatively small, which may be related to the small Al₂O₃/SiO₂ in A1. The high SiO₂ content in the glass made the glass network densified, at the same time, the low Al₂O₃ content made it difficult for Al³⁺ to replace Si⁴⁺ in glass network, hence, more Al³⁺ existed in [AlO₆] for A1.

3.5 Physical and mechanical properties analysis

The physical and mechanical properties of glass-ceramics with different Al₂O₃/SiO₂ were showed in Fig. 11; (a), (b) and (c) were the volume density, the bending strength and the microhardness respectively.

The results of Fig. 8 showed that the volume density of the glass-ceramics was between 2.735–2.770 g/cm³, the bending strength was between 95–120 MPa and the microhardness was between 540–610 Hv. Meanwhile, with the increase of Al₂O₃/SiO₂, the volume density, bending strength and microhardness of glass-ceramics all decreased first (reached the minimum in A4, the Al₂O₃/SiO₂ was 0.34) and then increased. The physical and mechanical properties of glass-ceramics are closely connected to the crystal phase, the glass phase and the combination among them[33, 34]. According to the previous analysis, with the Al₂O₃/SiO₂ gradually raising from 0.19 to 0.34, the precipitation of main crystal phase gehlenite (Ca₂(Al(AlSi)O₇)) increased, while the precipitation of secondary crystal phase decreased, and the remaining glass phase in the glass-ceramics also decreased. With the precipitation and growth of crystal phase, the internal stress of the glass-ceramics rose, and the decrease of liquid phase content was harmful to the sintering process, which led to the insufficient connection between crystal phase and led to more holes and defects in the glass-ceramics. At the same time, more Al³⁺ destroyed the glass network structure with [AlO₆], making the microstructure of the glasses tend to be loose. These results combined to reduce the physical properties (volume density, bending strength, microhardness) of the glass-ceramics with the increase of Al₂O₃/SiO₂ from 0.19–0.34 (A1-A4). With the further increase of Al₂O₃/SiO₂ from 0.34 to 0.39, the proportion of [AlO₄] in the glass network increased, which was beneficial to strengthen the glass network and make the structure tend to be dense. Meanwhile, the crystal phase in the glass-ceramics was further precipitated and grown, and the long strip crystal phase was interspersed with each other, which made the bonding between the crystal phase more complex and compact, thus increasing the density, bending strength and microhardness.

The effects of different Al₂O₃/SiO₂ on the structure and properties of BFS glass-ceramics were studied in this paper. With the increase of Al₂O₃/SiO₂, the gehlenite content in the glass-ceramics increased, while the diopside and hyalophane contents decreased. As Al₂O₃/SiO₂ increased from 0.24 to 0.34, the content of [AlO₆] in the glass network increased and [AlO₄] decreased gradually, the glass network structure became loose and the temperature of T_g decreased, simultaneously, the volume density, bending strength and microhardness of the glass-ceramics also decreased. With further increased Al₂O₃/SiO₂ from 0.34 to 0.39, the content of [AlO₆] in the glass decreased and [AlO₄] increased, the glass network structure became densified and the temperature of T_g increased, meanwhile, the volume density, bending strength and microhardness of the glass-ceramics increased. When Al₂O₃/SiO₂ was 0.19, the network structure of glass was relatively compact as high SiO₂ content, and the performance of glass-ceramics was the best, the volume density was 2.77 g/cm³, the bending strength was 120.50 MPa, the microhardness was 604.21 Hv.

Acknowledgement

We acknowledge Dr. Yumei Li (Center for Material research and analysis, Wuhan University of Technology, China) for the FESEM analyses.

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Effect of Al2O3/SiO2 on the structure and properties of blast furnace slag glass-ceramics

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Abstract

Figures

1. Introduction

2. Experiment

2.1 Raw materials and experimental formula

2.2 Preparation of glass-ceramics

2.3 Characterizations

3. Results And Discussion

3.1 DSC analysis

3.2 Crystal phase analysis

3.3 FTIR analysis

3.4 ²⁷Al NMR analysis

3.5 Physical and mechanical properties analysis

4. Conclusion

Declarations

Acknowledgement

References

Status:

Version 1

Effect of Al2O3/SiO2 on the structure and properties of blast furnace slag glass-ceramics

Status:

Version 1

Abstract

Figures

1. Introduction

2. Experiment

2.1 Raw materials and experimental formula

2.2 Preparation of glass-ceramics

2.3 Characterizations

3. Results And Discussion

3.1 DSC analysis

3.2 Crystal phase analysis

3.3 FTIR analysis

3.4 27Al NMR analysis

3.5 Physical and mechanical properties analysis

4. Conclusion

Declarations

Acknowledgement

References

Status:

Version 1

3.4 ²⁷Al NMR analysis