Effect of Al2O3/SiO2 on the structure and properties of blast furnace slag 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.%).

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:

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.

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.

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.