Phase evolution and properties of glass ceramic foams prepared by bottom ash, fly ash and pickling sludge

1.   Introduction

Rapid urbanization and industrialization in China have accompanied a rapid increase in the consumption of various valuable and non-renewable resources, along with the production of numerous solid wastes [1]. Solid waste management is often associated with a fairly high cost: the global average fee is $205.4 per ton and is predicted to increase by five times by 2025–2030 [2]. Although the management of solid waste has caught the attention of the Chinese government, the complex composition of solid waste and the possible harmful components, such as chromium, lead, arsenic, and fluoride, make it difficult to treat properly. Therefore, identifying a way to fully use solid waste can be conducive to reducing environmental pressure and the cost of solid waste management [3].

Glass ceramic foams are porous materials with a high percentage of pores and needle-shape crystals distributed uniformly in the main phase of glass [4]. These foams are widely used in building materials, sound-absorbing materials, and catalyst carriers, because of their superior physical properties [5]. The preparation of glass ceramic foams using solid waste containing abundant Si- and Al-based compounds, instead of natural resources, is popular [6–7]. Additionally, heavy metal (such as Pb, Cd, and Cr) can be immobilized in the glass network structure to make products harmless [8]. Furthermore, a variety of solid wastes, such as waste glass [9], fly ash (FA) [10], and vitrified bottom ashes (BA) [11], are used as raw materials for the production of glass ceramic foams.

BA composed of metal, glass, and non-combustible substances, is the most common solid waste (accounting for approximately 80%–90% of the total weight of residues) produced by municipal solid waste (MSW) incineration [12–13]. Incineration can reduce the amount of MSW, and the generated waste heat generated can be used for heating and power generation [14]. Urbanization in China will be accompanied by a significant increase in MSW, and the proportion of BA in solid waste will continue to increase. FA is a solid waste with major emissions. The World Coal Association estimates that the world coal industry produces about 750 million tons of FA annually, along with a continuous increase in the output of FA with the development of industrial countries [15]. Landfill and stockpiling are major methods to treat BA and FA; however, they require numerous land resources that have a negative impact on the environment. BA and FA are rich in SiO₂ and Al₂O₃, which are major raw materials for glass-ceramics. However, the high melting point of these oxides makes it difficult to form a liquid phase, which is an inevitable problem. Fortunately, pickling sludge (PS) containing generous amounts of CaF₂ and Fe can promote the formation of a liquid phase and crystallization [16]. PS is a complex composition waste produced during stainless steel production, whose rate of comprehensive utilization is quite low. The synthesis of glass ceramic foam by BA, FA, and PS has good practical value and potential environmental performance.

Pore morphology has a great impact on the physical properties of glass ceramic foams [9]. The formation of uniform pores plays a core role in the successful preparation of glass ceramic foams by sintering. Scholars are currently focusing on adjusting the process parameters, such as foaming temperature, time, and agent, while the study of the relationship between the glass structure and pore morphology is relatively limited [4,7]. There is a close relationship between the content of SiO₂ and Al₂O₃ in a silicate system glass and the crystallization ability of glass ceramic foams. SiO₂ and Al₂O₃ are used as the skeleton component of glass ceramic foams, mainly in the form of Si–O tetrahedra and Al–O tetrahedra. An increase of their content will improve the activation energy of the crystallization of glass ceramic foams and slow the crystallization. In this paper, glass ceramic foams with uniform pore structure were successfully prepared with BA, FA, and PS as major raw materials, and their effect on the pore morphology, glass structure unit, phase composition, and crystallization in the pore were evaluated under the different amount of BA. Furthermore, based on the changes in the glass structure units, the effects of SiO₂ and alkali metal oxides in the silicate system on the connection degree of the silica framework and the state of the foams in the molten glass were evaluated.

2.   Experimental

2.1   Raw materials

BA was derived from a waste incinerator in Beijing, China. FA, Class F according to ASTM C618, was obtained from the Shenhua Power Station, Taicang, Jiangsu Province, China. PS was collected from a steel plant in Shandong, China. Calcium carbonate (CaCO₃, ≥99%) and boric acid (H₃BO₃, ≥99.5%) were purchased from Sinopharm Chemical Reagent Co., Ltd., China. Glass ceramic foams were fabricated by BA, FA, and PS. The chemical compositions and mineralogical phases are listed in Table 1 and Fig. 1, respectively.

Table  1.  Chemical composition of the raw materials wt%

System	CaO	MgO	SiO2	Al2O3	Fe2O3	Cr2O3	Cl	TiO2	CaF2	Na2O	K2O	LOI
BA	25.89	3.67	42.14	8.59	5.68	0.12	0.76	0.98	—	3.86	1.90	6.41
FA	10.14	1.14	53.99	18.41	9.47	0.04	—	1.35	—	1.03	2.74	1.42
PS	—	1.17	9.30	2.72	25.52	5.01	0.12	0.12	45.71	1.74	0.25	8.34
Note: LOI—loss on ignition.

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Fig. 1.  XRD patterns of the raw materials: (a) BA; (b) FA; (c) PS.

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2.2   Samples preparation

PS was fixed at 20wt%. BA and FA were mixed in different proportions. In all samples, 15wt% calcium carbonate and 7wt% boric acid were added. After ball milling at 250 r/min for 4 h (QXQM-6, Ten Can Powder Co., Ltd., Changsha, China), the raw material mixtures (~10 g) were compacted into a cylindrical shape (diameter = 25 mm) by uniaxial pressing at 20 MPa and holding for 30 s. The samples were named D1, D2, D3, D4, and D5, respectively. The proportion of each raw material, mass ratio of BA/FA, and the chemical composition of the mixture are shown in Table 2.

Table  2.  Proportion of each raw material wt%

Constituents	CaO	MgO	SiO2	Al2O3	Fe2O3	Cr2O3	TiO2	P2O5	Na2O	SO3	Cl	K2O	CaF2	LOI
D1	12.05	1.78	42.09	12.82	11.79	1.05	0.26	0.99	1.88	1.49	0.21	1.98	9.14	2.47
D2	12.85	1.91	41.50	12.33	11.59	1.06	0.29	1.18	2.02	1.54	0.25	1.94	9.14	2.40
D3	13.62	2.03	40.90	11.83	11.40	1.06	0.34	1.36	2.16	1.59	0.29	1.90	9.14	2.38
D4	14.41	2.16	40.31	11.34	11.21	1.06	0.39	1.55	2.30	1.64	0.33	1.86	9.14	2.30
D5	15.20	2.29	39.72	10.85	11.01	1.07	0.44	1.74	2.44	1.68	0.36	1.81	9.14	2.25

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The samples were preheated in a muffle furnace first to prevent fracture but uneven heating (400°C for 20 min) [4]. Samples were then heated to 1180°C for 30 min at a heating rate of 10°C/min, annealed at 500°C for 40 min to stabilize the bubbles, and then cooled to room temperature in the furnace.

2.3   Characterization techniques

X-ray fluorescence spectroscopy (XRF-1800, RIGAKU, Japan) was employed to determine the chemical compositions of raw materials. The crystalline phase analysis of the raw materials and sintered samples were conducted using an X-ray diffraction diffractometer (XRD, Ultima IV, Japan). Differential scanning calorimetry (DSC, DSC204F1, German) measurements were performed for the powdered specimens (D1–D5) at a heating rate of 10°C/min from room temperature to 1200°C. The pore morphology and regional composition of the samples were analyzed by scanning electron microscopy (SEM, ULTRA 55, Zeiss, German) and energy-dispersive X-ray spectroscopy (EDS), respectively. In order to observe the sample morphology, before SEM observation, all samples were corroded in a 1vol% HF solution for 45 s, and then washed with deionized water and dried at 100°C for 5 h in a drying oven (DHG-9075A, Yi Heng Technology Co., Ltd., Shanghai, China). Digital photos of the samples were analyzed by an image analyzer (Nano Measurer, China) to determine the pore size distribution [17]. The compressive strength test was conducted with a universal testing machine (WDW-100, Keʼxin Co., Ltd., Changchun, China) at a compression speed of 0.5 mm/min. The Fourier transform infrared (FTIR, Nicolet-is10, Japan) spectra of the glass ceramic foams were measured in the range of 400–2000 cm⁻¹. The Raman spectra were performed by LabRAM HR Evolution (Horiba Scientific, France) with 10 mW power. The bulk density was calculated as the ratio of mass and volume, and the porosity was calculated according to the Archimedes method in water [18].

The water absorption was measured by the waterlogged method. The sintered samples were immersed in a beaker filled with deionized water, and the beaker was placed in a vacuum drying box and vacuumed for 1 h (DZF-6050, Yi Heng Technology Co., Ltd., Shanghai, China). Water absorption (W) of the sample was calculated from the weights of the dry samples (${m}_{1}$) and soaked samples (${m}_{2}$) using Eq. (1) [19]:

3.   Results and discussion

3.1   Characterization of the raw materials

The chemical composition of raw materials is shown in Table 1. The chemical composition of BA and FA are similar, with main oxides of SiO₂, CaO, Al₂O₃, Fe₂O₃, and MgO. The contents of Ca and Si compounds in BA are relatively high, at 25.89wt% and 42.14wt%, respectively. Compared with BA, much higher contents of SiO₂ and Al₂O₃ were in FA, the total amount of these two oxides reached 72.40wt%. SiO₂ and Al₂O₃ are conducive to the formation of a stable glass network structure during the sintering process. Alternatively, Na₂O, K₂O, CaO, and MgO in the BA and FA are glass network modifiers that help decrease the sintering temperature of the samples [20]. PS contains 45.71wt% CaF₂, 25.52wt% Fe₂O₃, and 5.01wt% Cr₂O₃. A higher content of CaF₂ is beneficial to reducing the melting point of raw materials, while Fe₂O₃ and Cr₂O₃ can promote nucleation and crystallization of glass ceramic foams.

The XRD patterns of raw materials are shown in Fig. 1. BA mainly consists of quartz (SiO₂, PDF#05-0490), calcite (CaCO₃, PDF#24-0027), gismondite (CaAl₂Si₂O₈·4H₂O, PDF#02-0096), and natrolite (Na₂Al₂Si₃O₁₀·4H₂O, PDF#02-0117). The mineral phases in FA are quartz (SiO₂, PDF#05-0490) and mullite (Al₆Si₂O₁₃, PDF#01-0613), which is consistent with the XRF result. Five main crystalline phases of fluorite (CaF₂, PDF#77-2245), quartz (SiO₂, PDF#05-0490), albite ((Na,Ca)(Si,Al)₄O₈, PDF#09-0457), anorthite (CaAl₂Si₂O₈, PDF#09-0464), and riebeckite (Na₂(Fe,Mg)₃Fe₂Si₈O₂₂(OH)₂, PDF#02-0074) are identified in PS.

3.2   Pore morphology of the prepared samples

The CaO and MgO content in the mixture increases gradually with the increase in the BA/FA ratio. During the formation of a melt, alkaline oxide releases free oxygen ions, which promotes the fracture of [SiO₄] [21]. Alkali metal cations also play an anti-polarization role on Si–O bonds in tetrahedrons [SiO₄] and destroy the integrity of glass networks. The contents of SiO₂ and Al₂O₃ in the mixture decrease gradually, which limits the construction of the glass network. Both factors lead to a decrease in the melt viscosity. Thus, it is necessary to achieve an appropriate viscosity to maintain the stability of the interface between the melt and gas phase in order to prepare a uniform foam structure.

The pore structure and the pore size distribution of the prepared samples with different BA/FA mass ratios are shown in Fig. 2. The BA/FA mass ratio has a significant effect on the pore morphology and pore size. When the mass ratio increases from 5/11 to 3/5, the average pore size increases from 125 μm (D1, BA content 25wt%) to 148 μm (D2, BA content 30wt%). However, Fig. 2(a) and (b) also show abnormally larger pores in the form of irregular polygons. This shape change is mainly because of the difference in the liquid-phase regional viscosity, which results in different resistance to bubble growth. With the rise in alkalinity of D2, the viscosity of the liquid phase decreases, more spherical pores are generated, and the shape of the pore changes to a typical circle and ellipse. When the BA content further increases to 35wt%, the average pore size decreases to 108 μm, and the size distribution of pores smaller than 200 μm reaches 91% (Fig. 2(c)). Additionally, the pore structure gradually transformed from irregular polygons to uniformly spherical. This phenomenon indicates that the D3 composition can form a uniform liquid phase during sintering, which is beneficial to the formation of pore [22]. When the BA content is 40wt%–45wt%, most pore structures form irregular polygons because of pore connectivity and the tiny pores disappear. Some large pores (1800–2200 μm) are observed in Fig. 2(d) and (e) as the pore distribution is no longer homogenous. Notably, the average pore size (D_ave) increases significantly from 261 to 623 μm, and a smooth non-porous zone appears in D5. A possible explanation for this phenomenon is that the viscosity of liquid phase is too low. The mass ratio of BA/FA continues to increase, which leads to bubble expansion and coalescence, as well as an increase in the bubble rise, resulting in the formation of connected pores or bubbles escaping. Thus, macropores and non-pore areas are formed in the sintering process.

Fig. 2.  Pore structure and pore size distribution of (a) D1, (b) D2, (c) D3, (d) D4, and (e) D5 samples with different BA/FA mass ratios.

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Considering the change in composition may affect the glass transmission and phase precipitation temperatures of the materials, the thermal behavior analysis of D1–D5 samples was performed as shown in Fig. 3. The results show that the glass transmission temperature (T_g) of the material decreases from 899 to 889°C with increasing BA/FA mass ratio. The material with lower T_g is prone to generate a liquid phase, in which the viscosity will more easily decrease with an increase in sintering temperature. The phase precipitation temperature of D3 is 1173°C, which is the highest among D3–D5, and the area of the peak that represents the heat released during the crystallization process corresponding to the phase precipitation temperature of D3 is larger than that of D4 and D5. However, the exothermic peak of D1 and D2 disappears in the range of 1100–1200°C, which may be caused by the excessively high crystallization temperature.

Fig. 3.  Differential scanning calorimetric curves of the powdered raw materials.

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SEM–EDS was performed on the glass surface (Area A) and inner wall of a hole (Area B) of D1 to analyze the different elements in the matrix and on the pore wall. As shown in Fig. 4, neither the matrix nor pore wall of D1 contained elements such as Si, Ca, Al, O, Na, Fe, or P. The main components of Si, Ca, and Al in Area A (glass phase) are 17.64wt%, 19.17wt%, and 8.22wt%, respectively. In contrast, the main components of the glass network in Area B (pore interior), such as Si and Al, decreased to 8.60wt% and 4.17wt% respectively, while the Ca content significantly increased to 39.83wt%, more than 31.23wt% of Si content. The Ca element is enriched in the area where pores are formed. CaO, a network modifier [23], provides free oxygen to interrupt Si–O in the glass network, which reduces the vitreous body in the area and promotes the formation of pores.

Fig. 4.  Microstructure and EDS spectra of D1.

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3.3   Glass structural units of the prepared samples

The FTIR spectra between 400 and 2000 cm⁻¹ is shown in Fig. 5. There are five main absorption bands in the spectra of the prepared samples and the band assignments are summarized in Table 3. The bands at 468 and 588 cm⁻¹ are assigned to the bend vibrations of O–Si–O and Si–O–Si [24–25]. The 700–720 cm⁻¹ band is ascribed to the bending of the B–O–B linkage in [BO₃] units and symmetrical stretching vibration of Si–O–Al, while the region at 1027–1035 cm⁻¹ is the stretching vibration of Si–O in [SiO₄] [26]. The band near 1400 cm⁻¹ originates from the asymmetric stretching vibration of B–O in [BO₃] [27]. The intensity of the absorption peak at 468 cm⁻¹ increases first and then decreases with an increase in the BA/FA mass ratio. As shown in Table 3, the CaO and Na₂O contents of D1–D5 increase gradually. In D1–D3, CaO and Na₂O provide the free oxygen required by the glass network, facilitating the formation of Si–O in [SiO₄] [28]. However, the decrease in the absorption peak intensity of D4–D5 may be because the excessive alkali metal (Na⁺), alkaline earth metal (Ca²⁺) cations, and free oxygen can destroy the integrity of the glass network and reduce the amount of Si–O–Si [29]. At the absorption peaks near 588 and 715 cm⁻¹, strength and peak position deviations occur because of the gradual decrease in the SiO₂ and Al₂O₃ content in the mixture, while the absorption peak at 1030 cm⁻¹ gradually narrows. The absorption peak of D4–D5 near 1400 cm⁻¹ shifts to a strong wavenumber and the intensity increases, which may be caused by the disproportionation reaction of CaO and [BO₃] units [26].

Fig. 5.  FTIR spectra of the prepared specimens.

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Table  3.  FTIR vibrational bands of the prepared samples

Wavenumber / cm−1					Assignments	References
D1	D2	D3	D4	D5
468	468	468	468	468	Bend vibrations of O–Si–O and Si–O–Si bond	[24]
588	588	588	588	588	Bend vibrations of Si–O–Si and O–Si–O bond	[25]
723	721	717	715	713	Bending vibration of B–O–B linkage in [BO3] and symmetrical stretching vibration of Si–O–Al	[26]
1027	1028	1031	1035	1035	Stretching vibration of Si–O bond in [SiO4]	[26]
1396	1398	1396	1427	1423	Asymmetric stretching vibration of B–O bond in [BO3]	[27]

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The Raman spectra and peak-splitting fitting results of the prepared samples are shown in Fig. 6. The pattern in the range of 800–1100 cm⁻¹ can be used to analyze the connection mode between Si–O and bridging oxygen in glass. The bands in the ranges of 840–880, 900–920, 950–1000, 1050–1100 cm⁻¹ correspond to the characteristic frequencies of monomer (Q⁰Si), dimer (Q¹Si i.e. [SiO₄] connecting one bridging oxygen), chain (Q²Si, [SiO₄] connecting two bridging oxygen), and sheet (Q³Si, [SiO₄] connecting three bridging oxygen), respectively [30]. A Gaussian function was used to deconvolute the characteristic peaks of the concentration of SiO₄. The content ratio of silicon–oxygen structural units is linearly correlated with the characteristic peak area ratio. Therefore, the following formula can be used to calculate the content of QⁿSi in structural elements [30].

Fig. 6.  Raman spectra (a) and deconvoluted Raman spectra (b–f) of the glasses using a Gaussian-type function.

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where ${C}_{{\mathrm{Q}}_{\mathrm{S}\mathrm{i}}^{n}}$ is the relative content fraction of the structural unit QⁿSi and ${A}_{{\mathrm{Q}}_{\mathrm{S}\mathrm{i}}^{n}}$ is the area ratio of its characteristic peak.

The fitting results of peaks are shown in Table 4. With the change in the BA/FA mass ratio, the relative content of Q³Si in [SiO₄] reduces gradually, which means that the connectivity of the glass network and the degree of polymerization decrease. Fig. 6 shows that the absorption near 1000 cm⁻¹ gradually moves to a lower wavenumber; when the BA content exceeds 35wt%, the absorption peaks in this range enter the characteristic frequency Q²Si. The lowest Q²Si content ratio in D3 is 47.63%, while the moderate Q³Si content ratio indicates that the pore feature may be related to the glass structural unit. The higher the Q³Si content ratio, the more the complete glass network connection, which means that more resistance must be overcome in the process of bubble growth. The presence of Q⁰Si, Q¹Si, and Q²Si indicates that some glass network structures are incompletely aggregated. These structures may have less binding force on the bubbles, thus causing gas escape or connected holes.

Table  4.  [SiO₄] structural unit content of the samples %

Sample	Q0Si	Q1Si	Q2Si	Q3Si	${C}_{{\mathrm{Q}}_{\mathrm{S}\mathrm{i}}^{0}}$	${C}_{{\mathrm{Q}}_{\mathrm{S}\mathrm{i}}^{1}}$	${C}_{{\mathrm{Q}}_{\mathrm{S}\mathrm{i}}^{2}}$	${C}_{{\mathrm{Q}}_{\mathrm{S}\mathrm{i}}^{3}}$
D1	8.09	18.17	38.39	10.84	10.72	24.07	50.85	14.36
D2	9.53	17.30	40.08	10.43	12.32	22.37	51.82	13.48
D3	11.68	18.34	36.63	10.26	15.19	23.85	47.63	13.34
D4	12.21	22.10	58.04	7.65	12.21	22.10	58.04	7.65
D5	10.01	15.61	68.04	6.34	10.01	15.61	68.04	6.34

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3.4   Phases and crystallization of the prepared samples

The XRD patterns of the glass ceramic foams prepared by different BA/FA mass ratios are shown in Fig. 7. A typical amorphous structure appears between 20° and 40° in all samples, which is mainly because of the formation of a glassy phase in the sintering process of silica. There are three main crystalline phases in glass ceramic foams, namely ferroaxinite (FeCa₂Al₂BSi₄O₁₅OH, PDF#70-1856), chromite (FeCr₂O₄, PDF#24-0512), and fluorapatite ((CaF)Ca₄(PO₄)₃, PDF#02-0845). With an increase in the BA/FA mass ratio, the diffraction peaks of chromite and fluorapatite increase most notably, and the chromite phase has a higher hardness (5.5–6 GPa), which can improve the mechanical strength of the samples. The appearance of fluorapatite and chromite proves that CaF₂ and Cr₂O₃ in the PS can promote crystallization [31]. Precipitation of a siliceous mineral phase can reduce the content of amorphous silica, thus destroying the integrity of the glass network and reducing the viscosity of the melt. However, excessively low amorphous silica content is a disadvantage of uniform pore formation.

Fig. 7.  XRD patterns of the prepared glass ceramic foams.

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As shown in Fig. 8(a–c), the crystalline grain shapes of D1–D3 are mostly columnar or acicular. Long columnar and acicular crystals grow gradually and change to equiaxed crystals with a further increase in the BA content, which suggests that the growth mode of the crystals change from one-dimensional to three-dimensional. The amount of the glass network regulators (CaO and MgO) increases with the mass ratio of BA/FA. As a result, the bridging oxygen bonds in the glass network decrease, the degree of glass polymerization decreases, and the grains are refined. Different samples have different crystallization ability, which may also affect the formation of pores. Fig. 8(a) and (b) show that the grains precipitated from the pore of D1–D2 grow in one dimension with weak crystallization ability and uneven crystallization in some regions. Therefore, resistance of bubbles in the glass phase is different, which is not suitable for the formation of even pores. In Fig. 8(c), there is Region C, where a glass phase coexists with acicular crystal grains. Given the D3 sample has uniform foaming and regular spherical pore, it is an appropriate glass phase and the crystallization ability can maintain steady growth of bubbles. However, the proportion of residual glass in D4 and D5 decreases, the crystallization ability increases, and the grains begin to aggregate and grow. This phenomenon may reduce the resistance of bubbles and promote pore coalescence. However, the grains in D4 and D5 begin to aggregate and grow, and the crystallization ability is enhanced. This phenomenon may reduce the resistance of bubbles and promote the coalescence of pores and bubble escape.

Fig. 8.  Microstructure of the inner walls in (a) D1, (b) D2, (c) D3, (d) D4, and (e) D5 samples.

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3.5   Physical and mechanical properties of the prepared samples

The bulk density, porosity, and mechanical properties of the prepared samples are present in Fig. 9. The porosity (58.65% to 51.25%) and bulk density (1.65 to 1.95 g/cm³) in D1–D4 showed opposite trends (Fig. 9(a) and (b)). Because of CO₂ escaping the glass melt, densification continuously proceeded, leading to lower porosity and a high bulk density of D5. When the mass ratio of BA/FA is greater than 1, the content of CaO and MgO can further increase, leading to a decline in the melt viscosity and an increase in surface tension [23]. Notably, an appropriate material composition is necessary for uniform pore formation. The porosity and bulk density depend on the number and size of the pores. When the dense glass area in the sample increases, the bulk density has a tendency to increase. The water absorption (7.14% to 2.77%) of sintered specimens with different ratios of BA/FA is shown in Fig. 9(c). The water absorption of D3 is relatively low because of the small pore size and less connected pores. Fig. 9(d) shows the compressive strength (6.10–16.23 MPa) of different samples. D3 (bulk density of 1.76 g/cm³, porosity of 56.01%) has better compressive strength with 16.23 MPa because of its uniform and low pore size.

Fig. 9.  (a) Porosity, (b) bulk density, (c) water absorption, and (d) compressive strength of samples with different BA contents.

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3.6   Formation mechanism of glass ceramic foams

The forming mechanism of glass ceramic foams is shown in Fig. 10. Because of the influence of heat transfer, the raw materials are heated unevenly, and the surface material whose temperature is greater than the internal material first reaches the glass transmission temperature and the decomposition temperature of calcium carbonate such that the outer layer of the material first appears in the melting bubble zone. With extended heat treatment time, heat is transferred inside the material, forming a large softened zone. Concurrently, the number of bubble cores increases, and some of the bubbles converge and overflow the melt. If the heat treatment time is insufficient, there will be no holes inside (as shown in the sintering sample in Fig. 10). The raw material composition has an important influence on the formation of pores, phases, and viscosity. Al₂O₃ and SiO₂, as the main structures of the glass network, provide a skeleton structure for the glass ceramic foams. The increase in alkaline earth oxides (CaO and MgO) and alkaline oxides (Na₂O and K₂O) will promote the fracture of bridging oxygen bonds, destroying the integrity of the glass network resulting in a reduced viscosity [20]. The addition of H₃BO₃ as a fluxing agent increases the content of B₂O₃, and the appropriate amount of B₂O₃ can reduce the viscosity of the melt, while excessive B₂O₃ will lead to the formation of [BO₄] and increase the viscosity [9]. An appropriate viscosity is imperative to the growth and stability of the bubbles. If the viscosity is too high, the resistance of the bubbles in the melt will increase, which will prevent the formation of uniform and regular pores. However, if the viscosity is too low, bubbles will fuse and escape. A certain amount of nucleating agent can promote the crystals of glass nucleation and growth, which is similar to a key factor to control the crystallization of glass. In the process of heat treatment, CaF₂ and Cr₂O₃ in PS can be dissolved in a glass, and non-uniform nucleation can be promoted by phase separation or direct precipitation of crystals, resulting in the formation of ferroaxinite, chromite, and fluorapatite. Therefore, a crystalline phase and glass phase are formed.

Fig. 10.  Glass ceramic foam formation mechanism.

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4.   Conclusion

Glass ceramic foams were successfully prepared using BA, FA, and PS. The pore structure and pore size distribution of glass ceramic foams were greatly affected by the BA/FA mass ratio. When the BA/FA mass ratio was 7/9, glass ceramic foams with uniform pore size and a spherical structure were obtained because the viscosity of the liquid phase was conducive to maintaining equilibrium of the melt and gas under this component. The appropriate content of the glass phase and crystallization conditions is conducive to the formation of pores. Alkaline oxides and alkaline earth oxides in the D1–D5 mixture increases, while the glass network (SiO₂ and Al₂O₃) decreases gradually. FTIR and Raman analysis shows that the change in composition reduces the bending vibration and symmetrical stretching vibration of Si–O in [SiO₄], which indicates that the quantity of Si–O–Si decreases as well as the glass integrity. The pore structure has an important influence on the porosity and other physical properties. Glass ceramic foams with increasing mass ratios of BA/FA at 1180°C for 30 min, porosity, bulk density, water absorption, and compressive strength obtained from 58.65% to 51.25%, 1.65 to 1.95 g/cm³, 7.14% to 2.77%, 6.10–16.23 MPa, respectively. Samples with different properties can be obtained by adjusting the components in the preparation of glass ceramic foams.

Acknowledgements

This work was financially supported by the National key R&D projects (Nos. 2019YFC1907101, 2019YFC1907103, 2017YFB0702304), the Key R&D project in Ningxia Hui Autonomous Region (No. 2020BCE01001), the National Natural Science Foundation of China (No. 51672024), the Xijiang Innovation and Entrepreneurship Team (No. 2017A0109004), the Program of China Scholarships Council (No. 201806465040), the Fundamental Research Funds for the Central Universities (Nos. FRF-IC-19-007, FRF-IC-19-017Z, FRF-MP-19-002, FRF-TP-19-003B1, FRF-GF-19-032B, and 06500141), the State Key Laboratory for Advanced Metals and Materials (No. 2019Z-05), and the Integration of Green Key Process Systems MIIT. The authors would like to thank the editor for editing of the manuscript and the anonymous reviewers for their detailed and helpful comments.

Conflict of Interest

The authors declare no potential conflict of interest.