SiO2 ratio on Li2O – Al2O3 – SiO2 – B2O3 – BaO glass-ceramic bonds for vitrified cBN abrasives

Ceramics International 45 (2019) 15358–15365

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Exploring the influences of Li2O/SiO2 ratio on Li2O – Al2O3 – SiO2 – B2O3 – BaO glass-ceramic bonds for vitrified cBN abrasives

T

Jiang Shia,b, Feng Hea,b, Junlin Xiea,b, Hu Yangb, Zichen Guoa,b, Xiaoqing Liuc,∗ a

State Key Laboratory of Silicate Materials for Architectures, Wuhan University of Technology, Wuhan, 430070, China School of Materials Science and Engineering, Wuhan University of Technology, Wuhan, 430070, China c Center for Materials Research and Analysis, Wuhan University of Technology, Wuhan, 430070, China b

A R T I C LE I N FO

A B S T R A C T

Keywords: Glass-ceramic cBN Bending strength Wetting

Li2O–Al2O3–SiO2–B2O3–BaO glass-ceramic bonds for vitrified cBN abrasives with various Li2O/SiO2 ratio were prepared by the traditional melt-quenching method. The structure of parent glass was identified by X-ray diffraction (XRD) and Fourier Transform infrared spectroscopy (FTIR). The reduction in the Li2O/SiO2 ratio increased the connectivity of the glassy network, resulting in a decrease of the crystallization ability for the parent glasses. Besides, a comprehensive understanding of measured and calculated coefficient of thermal expansion (CTE) for glass-ceramic bonds demonstrated that the simultaneous presence of a higher content of LiAlSi2O6 and LiAlSi3O8 along with the absence of SiO2 precipitation were believed to contribute largely to the lower CTE of glass-ceramic bonds. Lastly, the highest bending strength of cBN composites (100.53 MPa) relied on the sufficient wetting of B3 glass-ceramic bonds (Li2O/SiO2 = 0.11) on cBN and the close CTE between them when composites were sintered at 860 °C for 2 h, indicating B3 is a kind of suitable and promising glass-ceramic bonds for cBN abrasives.

1. Introduction Vitrified cBN (Cubic boron nitride) abrasives, which is a sort of superhard composites consisting of cBN, vitrified bonds and pores [1], have become one of the most critical tools for precision machining, because of their high processing efficiency, excellent processing quality, long service life, and environmental friendliness [2]. As a vital part of vitrified cBN abrasives, vitrified bonds, the nature of which is a type of amorphous glass material, play indispensable roles in consolidating the cBN particles into the desired shape during the sintering process of vitrified cBN abrasives. To date, there were a large number of published studies focused on borosilicate glass (R2O-MO-B2O3–SiO2, R and M stand for alkali metal and alkaline earth metal, respectively) [3–6]. Besides, vitrified bonds belonging to the Li2O–Al2O3–SiO2 (LAS) glassceramic had been successfully prepared [7], and they were found more promising than borosilicate glass in the field of vitrified bonds for its excellent performance on adjustable CTE over a large temperature range [8], low liquid-phase formation temperature [9], and high bending strength [10]. The crystalline phase precipitated in the glassy matrix, depending on the glass composition [11–15] and the heat treatments [16–18], can significantly govern the performances of glass-ceramics. While

∗

compared with adjustment of the heat treatments, the modification of the composition for glass-ceramic bonds not only controls the potential phase assemblage but also determines the sintering temperature of the glass-ceramic bonds. However, considering that cBN is too easily oxidized to deteriorate its properties in air at temperature above 900 °C [19], the vitrified bonds are supposed to spread and wet on cBN under 900 °C to tightly consolidate the abrasive grains. Given this fact, the composition of glass-ceramic bonds is critical for the preparation of vitrified cBN abrasives with high performance. Moreover, in an attempt to develop glass-ceramic bonds with better fluidity at high temperature and lower melting temperature, BaO and B2O3 had been considered as one of the components in LAS glass-ceramic bonds [20]. Up to now, some changes in the glass-ceramic composition such as Y2O3 [21], B2O3 [22,23], Na2O/BaO ratio [20] and CaF2 [7] on the microstructure and performances of LAS glass-ceramic bonds had been well discussed. To the best of the authors’ knowledge, however, constituted as the main compositions of LAS glass-ceramic bonds, the relationships between the varying content of Li2O, Al2O3, and SiO2 with the properties of LAS glass-ceramic bonds were not described previously. In fact, Li2O/SiO2 ratio had been proved to offer profound influences on the formation of LiSi2O3 phase in Li2O–SiO2 glass-ceramic system [24]. Thus, it is meaningful for exploring the effects of Li2O/SiO2 ratio on

Corresponding author. E-mail address: [email protected] (X. Liu).

https://doi.org/10.1016/j.ceramint.2019.05.029 Received 18 March 2019; Received in revised form 26 April 2019; Accepted 4 May 2019 Available online 07 May 2019 0272-8842/ © 2019 Elsevier Ltd and Techna Group S.r.l. All rights reserved.

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Table 1 Chemical compositions of different glass-ceramic bonds (wt.%). Id

SiO2

B2O3

BaO

Li2O

Al2O3

CaO

Na2O

MgO

others

Li2O/ SiO2]Y

B1 B2 B3 B4 B5

50.00 52.00 54.00 56.00 58.00

10.00 10.00 10.00 10.00 10.00

9.00 9.00 9.00 9.00 9.00

10.00 8.00 6.00 4.00 2.00

10.00 10.00 10.00 10.00 10.00

2.00 2.00 2.00 2.00 2.00

3.00 3.00 3.00 3.00 3.00

3.00 3.00 3.00 3.00 3.00

3.00 3.00 3.00 3.00 3.00

0.20 0.15 0.11 0.07 0.03

Li2O–Al2O3–SiO2–B2O3–BaO glass-ceramic bonds with low-melting temperature and low CTE in this paper.

2. Experimental 2.1. Materials preparation The parent glass with chemical composition shown in Table 1 was synthesized through melting mixed raw materials at 1500 °C for 2 h and then quenching the acquired melts into ice water. For the convenience of description, the ratio of Li2O/SiO2 was defined as Y, and the five samples were sequentially named B1–B5 with a decreasing Y. For measuring the CTE of glass-ceramic bonds, the ground asquenched glasses were pressed into strips of 40 mm × 6 mm × 5 mm with the assistances of 5% Polyvinyl alcohol (PVA) solution under 30 MPa of pressure before they were sintered at 860 °C for 2 h. Besides, the raw glass-ceramic bonded cBN composites consisted of as-quenched glasses, cBN particles, and PVA solution were also pressed into the rectangular bars with the above-mentioned same dimensions. Then the raw vitrified cBN composites were sintered at 500 °C for 1 h for the shrinkage tests and sintered at 860 °C for 2 h for the assessment of bending strength.

2.2. Materials characterization The amorphous state or crystallization of as-quenched glasses and glass-ceramic bonds were verified by X-ray diffraction (XRD, D8 Advance, Bruker, Germany) at a scanning rate of 2°/min from 10 ° to 80 °with utilizing Cu Kα radiation source. A differential scanning calorimetry (DSC, STA449, NETZSCH, Germany) was employed to analyze the thermodynamic performance of the as-quenched glasses at a heating rate of 10 °C/min from room temperature to 1000 °C in the air. Besides, Infrared spectra of the as-quenched glasses were carried out on a Fourier transform infrared spectrometer (Nicolet Nexus FT-IR, Thermo, USA) over 400 cm−1-1600 cm−1 at room temperature. CTE of sintered strip glass-ceramic bonds was measured by a thermal dilatometer (DIL-402C, NETZSCH, Germany) with a heating rate of 5 °C/min from room temperature to 600 °C. In addition, the shrinkage of the raw glass-ceramic bonded cBN composites was also measured by the DIL402C under heating rates of 5 °C/min, 10 °C/min, and 20 °C/min from room temperature to 880 °C in air. A hot stage microscope (HSM, HM 867, TA Instruments, USA) was employed to observe and measure the contact angle between the glass-ceramic bond and the cBN at 800–880 °C. Besides, bending strength of glass-ceramic bonded cBN composites was conducted using a three-point test method in a universal material tester (AG-IC50KN, Shimadzu, Japan) at a crosshead speed of 0.5 mm/min, a fulcrum span of 25 mm, a loading speed of 9.8 ± 0.1 N/s. A field emission scanning electron microscope (FESEM, ULTRA PLUS, Zeiss) was used to observe the cross-section of glassceramic bonded cBN composites.

Fig. 1. DSC curves of as-quenched glasses.

3. Results and discussions 3.1. Influences of Li2O/SiO2 ratio on the structure of glass-ceramic bonds 3.1.1. DSC analysis As shown in Fig. 1, both the glass transition temperature (Tg) and the peak of crystalline temperature (Tc) increased with a decreasing Y. In addition, the reduction in the specific area of the exothermic crystallization peaks shown with the decreasing Y in Fig. 1 was found to be obvious, especially the absence of this exothermic peak had become almost realized when Y = 0.03. Thus, the above-mentioned phenomenon indicated that the decrease of Li2O/SiO2 ratio led to the weakening of the crystallization ability of the parent glass. A reasonable explanation concerned the content of SiO2 and Li2O. As the network former in glass structure, SiO2 focuses on strengthening the degree of polymerization (DOP) of glassy network, while the Li2O offer the power of depolymerizing the glass structure. Hence, the decrease of Li2O/SiO2 ratio increase the connectivity of the glassy network which can hinder the mobility of the atoms in the glass matrix at high temperature, resulting in a weakening of the crystallization ability of parent glass.

3.1.2. FTIR analysis Five distinct absorption bands were observed in the five FTIR spectra of as-quenched glasses shown in Fig. 2. Bands at 461 cm−1 is ascribed to the bending vibration of Si–O–Si in [SiO4] [25], the presence of a weak band at 715 cm−1 is attributed to the overlapping of bending vibration of B–O–B in [BO3] and symmetric stretching vibrations of Si–O–Al [26,27]. Furthermore, the symmetrical bending vibration of Si–O–Si is also indicated by the presence of a band at 794 cm−1 [28,29]. The appearance of a strong band at 1020 cm−1 is referred to the symmetric stretching vibration of Si–O–Si in [SiO4] [30,31]. Asymmetric stretching vibration of B–O–B in [BO3] is represented by the bands at 1426 cm−1 [32]. To state the differences between the five as-quenched glasses in structure, a deep comparison and understanding of intensity and peak position of each absorption band were applied. It was obvious that the intensity of bands at 461 cm−1 and 794 cm−1 gradually enhanced with a decreasing Y, this was ascribed to that the higher number of Si–O–Si bridging oxygen in the parent glass would be obtained when the amount of SiO2 was rising. Besides, the Si–O–Al bond which originally formed by the interconnection of the [SiO4] and [AlO4] would be replaced by a rising number of Si–O–Si, thus band position at 715 cm−1 shifted to a lower wavenumber with a decreasing Y. Finally, an enhanced intensity of the absorption bands at 1020 cm−1 from B1 to B5

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composites, and the calculation formula of sintering activation energy was shown as followed [34]:

lnβ = −

Fig. 2. FTIR spectra of as-quenched glasses.

which was attributed to the increased number of Si–O–Si was also observed. In addition, the shift of band at 1020 cm−1 to higher wavenumber proved evidence of the increased DOP of as-quenched glasses with decreasing Y. In fact, the absorption bands at 1020 cm−1 is considered as the overlapping of different Si–O stretching vibration derived from Si–O–Si bridges with various numbers of bridging oxygen [33], and a smaller value of Y contributes to the increased content of Si–O absorption band with higher wavenumber. Hence, after finishing the convolution of different Si–O absorption band, a higher wavenumber of 1020 cm−1 would be presented in FTIR with a smaller value of Y. 3.1.3. Calculation of sintering activation energy Fig. 3 shows the linear shrinkage of raw glass-ceramic bonded cBN composites sintered at different heating rates, the corresponding thermodynamic temperature for specific linear shrinkage of 3‰, 5‰ and 7‰ at different heating rates β (5, 10, and 20 K/min) were selected for calculating the sintering activation energy of glass-ceramic bonded cBN

Q +A RT

(1)

Where Q is the sintering activation energy, kJ/mol; R stands for the ideal gas constant, 8.314 J/(mol·K); T means the thermodynamic temperature, K; A denotes the Arrhenius parameters, the unit of which is consistent with the reaction rate constant. For a specific linear shrinkage of cBN composites sintered at different heating rates, the Q was calculated by the slope of the fitting line of lnβ∼1/T which was depicted in Fig. 4, and the final results of Q shown in Fig. 4(f) was determined by averaging the Q calculated under three different specific linear shrinkage. Hence, as illustrated in Fig. 4(f), the sintering activity of the glass-ceramic bonded cBN composites increased when the Li2O/SiO2 ratio decreased. As mentioned earlier, since the reduction of Y results in an increased connectivity of the glassy network, it is beneficial for the generation of the liquid phase which is indispensable for the consolidation of cBN particles into a whole part at sintering temperatures. Thus, the less amount of liquid phase and worse wetting and spreading of liquid phase on cBN particles at sintering temperatures are the results of a higher Q, and it is more difficult for the completion of the sintering process for cBN composites at sintering temperature because of the higher energy barrier for the consolidation of glass-ceramic bond and cBN particles for a higher Q. For reaching of predetermined perfect consolidation of cBN abrasives, a higher sintering temperature is required. In brief, the above findings further confirmes that a large amount of liquid phase is conducive to the preparing of glass-ceramic bonded cBN composites at a lower sintering temperature, which is also consistent with the experimental results on a ceramic-sintering process of Singh. V. K [35] and Yawei Chen [36]. 3.2. Influences of Li2O/SiO2 ratio on the properties of glass-ceramic bonds 3.2.1. XRD analysis The broad halo between 20° and 30° in Fig. 5 indicated their amorphous nature of the five as-quenched glasses. Moreover, no

Fig. 3. Linear shrinkage of cBN abrasives sintered at different heating rates; (a)–(e): B1–B5. 15360

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Fig. 4. Fitting results of lnK∼1/T under different shrinkage rates; (a)–(e): B1–B5, (f) calculated results of Q.

Fig. 6. XRD pattern of glass-ceramic bonds sintered at 860 °C for 2 h.

Fig. 5. XRD pattern of as-quenched glasses.

obvious diffraction peaks except for the broad halo in Fig. 5 confirmed the absence of any crystalline phase in the prepared parent glasses. XRD patterns of glass-ceramic bonds sintered at 860 °C for 2 h are presented in Fig. 6. The patterns showed that the LiAlSi2O6 and LiAlSi3O8 were the predominant phases in the B1–B4 glass-ceramic bonds, while the intensity of the diffraction peak for LiAlSi2O6 in B4 was significantly diminished compared with B1–B3. Meanwhile, the diffraction peak corresponding to SiO2 was inspected in B4 and became so apparent that the SiO2 constituted as the only crystalline phase in B5 glass-ceramic bond. Hence, a decreasing ratio of Li2O/SiO2 in parent glass results in the transformation of the major crystalline phase from LiAlSi2O6 and LiAlSi3O8 to SiO2 after sintering at 860 °C for 2 h. 3.2.2. CTE analysis The CTE between 30 and 300 °C of the glass-ceramic bonds sintered at 860 °C for 2 h is shown in Fig. 7, it demonstrated that the CTE decreased first with the decreasing Y and then increased with the further decreasing Y value, and a minimum value of CTE obtained at the B3 glass-ceramic bonds was 3.23 × 10−6 °C−1, indicating that a high 15361

Fig. 7. CTE of glass-ceramic bonds sintered at 860 °C for 2 h.

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content of SiO2 in the glass composition is not conducive to obtaining glass-ceramic with low CTE. Generally, the CTE of glass-ceramic is determined by the relative content of the residual amorphous phase and the precipitated crystalline phase. Hence, the theoretical calculation of glass-ceramic was carried out to demonstrate the effect of the composition of the glass-ceramic bonds on its CTE after heat treatments. 3.2.3. Theoretical calculation of CTE for glass-ceramic bonds According to reference [11], the theoretical CTE of glass-ceramics can be calculated as followed:

α gc = x g α g + x c1α c1 + x c2α c2 + …

(2)

∑ α iNi ∑ Ni

(3)

Where αg represents the theoretical CTE of the glass, αi is the CTE of oxide shown in the glass compositions (Table 2), and Ni stands for the molar fraction of the corresponding oxide for the chemical composition of glass. As the network former in the glass structure, the calculation of CTE for SiO2 and B2O3 needs to be paid more attention. Generally, the CTE of SiO2 is greatly affected by the glass network connectivity which is determined by the content of network formers and network modifiers in the glass structure. The specific empirical method of calculation is shown in Table 3 (N denotes the molar fraction of SiO2) [38,39]. The calculation of the CTE for B2O3 is slightly more complicated than SiO2, a constant defined as ψ is calculated as followed [38]:

ψ=

∑ NR2O + ∑ NMO − ∑ NAl2O3 NB2O3

(4)

In the above equation, R2O and MO represents alkali metal oxides and alkaline earth metal oxides, respectively. In addition, N denotes the molar fraction of the alkali metal oxides or the alkaline earth metal oxides for the chemical composition of glass. Thus, the CTE of B2O3 is considered as 5.0 × 10−6 °C−1 when ψ > 4, while it is calculated by α = 1.25 × ψ for ψ < 4. According to the above-mentioned method of calculation, the results of calculated CTE for the parent glass were listed in Table 4 which pointed out that the CTE showed a tendency to decrease with a decreasing Y. In general, the CTE of glass is considered to be controlled by the magnitude of the attractive force between the cation and the oxygen in the glass structure, and the nature of the thermal expansion of glass is ascribed to the increasing spacing between the internal particles aroused by the thermal vibration of these particles when glass is heated. Thus, the CTE of glass becomes smaller for the restricted enlargement of spacing between the internal particles caused by a larger attractive force. Actually, the bond strength of Si–O is known to be greater than R–O. Hence, the increased content of Si–O and decreasing content of

α( × 10

−6

−1

°C

)

BaO

Li2O

Al2O3

CaO

Na2O

MgO

P2O5

ZrO2

20

26

−4

13

40

6

14

−10

α( × 10−6 °C−1)

67–100 34–67 0–34

10.2–0.1 × N 6.85 + 0.05 × N 5.2

ID

B1

B2

B3

B4

B5

α( × 10−6 °C−1)

9.39

8.56

7.69

6.77

5.80

R–O are considered as a result of the replacement of Li2O by SiO2 in glass composition, and it is more difficult for the enlargement of spacing between the internal particles in the glass, the CTE of parent glass also consequently decrease. In order to quantify the content of amorphous phase or crystalline phases defined in Eq. (2) for glass-ceramic bonds sintered at 860 °C for 2 h, Whole Pattern Fitting (WPF) Refinement based on the method proposed by H. M. Rietveld [40,41]was carried out using “MDI JADE 6″ software. The WPF Refinement results of Fig. 6 are listed in Table 5. Fig. 8 is transformed from Table 5 for clearly showing the correlation between the phase content of glass-ceramics and Y. As shown in Fig. 8, on reducing the ratio of Li2O/SiO2, the content of amorphous phase became higher, indicating the decreasing ability of crystallization from B1 to B5 which was consistent with the findings observed in Fig. 1. In addition, the content of LiAlSi2O6 increased first and then decreased with the decreasing Y, and the maximum content of the LiAlSi2O6 was obtained at the B3. Besides, LiAlSi3O8 found its minimum content at B3, and it was absent in B5. While the content of Li2SiO3 was much less than that of LiAlSi2O6 and LiAlSi3O8, and it gradually decreased with the reduction of Y. At last, the content of BaAl2Si2O8 was always around 1%, while the SiO2 was absent until Y > 0.07 and became the only precipitation in the B5 glass-ceramic bonds. According to the works of literature, the CTE of LiAlSi2O6 and LiAlSi3O8 are 0.9 × 10−6 °C−1 [42] and 3 × 10−6 °C−1 [43], respectively, and it is 2.2 × 10−6 °C−1 for BaAl2Si2O8 [44], 13.2 × 10−6 °C−1 for SiO2 [45]. In addition, CTE of Li2SiO3 is greater than that of Li2Si2O5 [46]which is 9.0 × 10−6 °C−1 [47]. Hence, for the convenience of calculation, the approximate CTE of Li2SiO3 is considered as 9.0 × 10−6 °C−1. Thus, the content of amorphous phase or crystalline phases shown in Table 5, the CTE of glasses shown in Table 4, and the above-mentioned CTE of crystalline phases are substituted into Eq. (2) to obtain the calculated CTE of glass-ceramics. Fig. 9 illustrates the calculated and measured CTE of glass-ceramic bonds sintered at 860 °C for 2 h. The results indicated that the measured CTE were close to the theoretical calculation results, which proved that the above-mentioned calculation method of CTE was suitable for the studied glass-ceramic bonds in this paper. A comprehensive understanding of Figs. 8 and 9 demonstrated that the variety of CTE for glass-ceramic bonds with decreasing Y was identical exactly alike the content of LiAlSi2O6, that was the higher the content of LiAlSi2O6 was, a lower CTE of the corresponding sintered Table 5 WPF Refinement results of glass-ceramic bonds sintered at 860 °C for 2 h (%).

Table 2 CTE of oxides. Oxide

Mole fraction of SiO2 (%)

Table 4 CTE of B1–B5 parent glasses.

Where αgc represents the CTE value of the glass-ceramics, αg, αc1, and αc2 are the CTE values of residual amorphous phase and crystalline phases in the sintered glass-ceramic bonds, respectively. Besides, xg, xc1, and xc2 denote the content of amorphous phase and crystalline phases, respectively. Eq. (2) indicates that the CTE of glass-ceramics is considered as a weighted calculation of the crystalline phases and the amorphous phase (glass). While the CTE of each precipitated crystalline phase is available from works of literature. Hence, to finish the calculation of CTE for glass-ceramic, a detailed introduction to the theoretical calculation method of CTE for glass is described in the following equation [37]:

αg =

Table 3 Calculation of CTE for SiO2.

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ID

Amorphous phase

LiAlSi2O6

LiAlSi3O8

Li2SiO3

BaAl2Si2O8

SiO2

B1 B2 B3 B4 B5

25.66 29.89 32.99 37.64 46.22

34.12 38.98 48.45 16.09 0

28.99 26.71 15.88 26.38 0

9.37 2.66 1.27 0.75 0

1.86 1.75 1.41 0.31 0

0 0 0 18.83 53.78

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cBN at different temperatures, it was clearly seen that the contact angle increased with the reduction of Y at the same temperature. Naturally, the decreasing ratio of Li2O/SiO2 in the bonds elevated the required energy which breaks the bonding between the internal atoms for the convenience of diffusion of the ions. In addition, a less liquid phase would be generated at the same temperature when Y was gradually decreasing. Thus, the spreading and wetting of glass-ceramic bonds on the cBN became more difficult with a decreasing Y at the same temperature. Besides, the content of liquid phase generated by B5 was too low to finish the wetting of glass-ceramic bonds on cBN at 800–880 °C. As a result, a poor holding of glass-ceramic bonds on cBN particles would occur, and it is obviously not conducive to the preparation of vitrified cBN composites with high bending strength and high service life. In addition, with a rising sintering temperature, the contact angle between glass-ceramic bonds and cBN showed the tendency of great reduction, indicating that the increased sintering temperature is beneficial to the wetting and spreading of the glass-ceramic bonds on the surface of cBN particles. Fig. 8. WPF Refinement results of phase content.

Fig. 9. Calculated and measured CTE of glass-ceramic bonds sintered at 860 °C for 2 h.

glass-ceramic bonds showed. In fact, it is reasonable that a higher content of LiAlSi2O6 results in a lower CTE of glass-ceramics for LiAlSi2O6 shows the lowest CTE among all the precipitations. Besides, it was worth noting that the CTEs of B4 and B5 showed a large increase compared with B3. The reasonable explanations are stated as follows: (1) Table 5 depicts that the content of amorphous phase in B4 and B5 glass-ceramics are higher than that of B1–B3, while the CTE of amorphous phase (parent glass) in B4 and B5 are also higher than that of LiAlSi2O6 and LiAlSi3O8. Thus, a higher content of amorphous phase in B4 and B5 result in a higher CTE of glass-ceramics sintered at 860 °C for 2 h (2) SiO2 which presents the highest CTE among all the precipitations in the studied glass-ceramics constituted as the major crystalline phases in B4 and B5, the higher content of SiO2 in B4 and B5 is another reason for the high CTE of glass-ceramics. In brief, a lower content of the amorphous phase and a higher content of LiAlSi2O6 and LiAlSi3O8 contribute to obtaining glass-ceramic bonds with lower CTE. Hence, to obtain glass-ceramic bonds with lower CTE, it is vital for precipitating enough content of crystalline phase including LiAlSi2O6 and LiAlSi3O8. Besides, avoiding the precipitation of SiO2 is another key point according to the above-mentioned analyses. 3.2.4. HSM analysis Fig. 10 depicts the contact angle between glass-ceramic bonds and

3.2.5. Bending strength of glass-ceramic bonded cBN composites Obviously, the bending strength of glass-ceramic bonded cBN composites sintered at 860 °C for 2 h reached the maximum (100.53 MPa) when Li2O/SiO2 was 0.11 (B3). Coincidentally, the B3 glass-ceramic bonds sintered at 860 °C for 2 h showed the lowest CTE among the five samples according to Fig. 7. As depicted in Figs. 7 and 9, the lower bending strength of glass-ceramic bonded cBN composites was accompanied by a higher CTE of glass-ceramic bonds. Reasons were summarized as followed: (1) The CTE value of B3 sintered at 860 °C for 2 h is 3.23 × 10−6 °C−1, which is closest to the CTE of cBN (3.50 × 10−6 °C−1 at 300–1000 °C [48]) among the five glass-ceramic bonds. The almost equal CTE between cBN and glass-ceramic bonds ensures the least number of microcracks generated due to the thermal mismatch during sintering and cooling of the composites. Thus, the corresponding glass-ceramic bonded cBN composites showed the highest bending strength. (2) The spreading and wetting of glassceramic bonds on cBN particles at the sintered temperature determines its holding power for cBN particles which is vital for the bending strength of glass-ceramic bonded cBN composites [49]. Hence, a good wetting state between glass-ceramic bonds and cBN particles are another guarantee of the high bending strength of glass-ceramic bonded cBN composites. Since the high-temperature fluidity of B4 and B5 on cBN was obviously inferior to that of B1–B3 according to Fig. 10, the lower bending strength of composites for B4 and B5 were reasonable. 3.2.6. SEM analysis Fig. 12 shows the SEM of glass-ceramic bonded cBN composites sintered at 860 °C for 2 h. In Fig. 12, the number of bare cBN abrasive particles which were uncovered by glass-ceramic bonds was nearly equal in B1–B3 composites. While it increased significantly when the glass-ceramic bonds were B4 and B5. Fig. 10 indicated that the contact angles between the glass-ceramic bonds and cBN at 860 °C were: 37° (B1), 45° (B2) and 49° (B3), respectively. While the contact angles between B4 or B5 glass-ceramic bonds and cBN were considered as 90° for B4 and B5 failed to produce enough liquid phase to wet cBN at 860 °C. Therefore, B1–B3 glassceramic bonds had demonstrated sufficient ability to wet and cover cBN particles at 860 °C. However, it was unavailable and impossible for B4 and B5. In fact, it is beneficial for the enhancement of the bending strength of glass-ceramic bonded cBN composites with more cBN particles wetted and covered by glass-ceramic bonds at sintering temperature. Hence, the higher bending strength of cBN composites also relies on the sufficient wetting of glass-ceramic bonds on cBN, and this conclusion is consistent with the above-mentioned analysis of Fig. 11. Summarily, according to the results of CTE, HSM, and bending strength, B3 glass-ceramic bonds is proved to be a kind of suitable and promising vitrified bonds for cBN abrasives.

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Fig. 10. The wetting between glass-ceramic bonds and cBN at different temperature, the number in the upper right corner of each figure indicates the contact angle between cBN and the glass-ceramic bonds.

4. Conclusions

Fig. 11. Bending strength of glass-ceramic bonded cBN composites sintered at 860 °C for 2 h.

The effects of Li2O/SiO2 ratio on the structure and properties of Li2O–Al2O3–SiO2–B2O3–BaO glass-ceramic bonds for vitrified cBN abrasives had been discussed. The reduction in the Li2O/SiO2 ratio would increase the connectivity of the glassy network which would hinder the mobility of the atoms in the glass matrix at high temperature, resulting in a decrease of the crystallization ability from B1 to B5. In addition, the sintering activation energy of consolidating glassceramic bonds and cBN particles increased from 114.52 kJ/mol to 722.70 kJ/mol as the Li2O/SiO2 ratio decreased, indicating that the low Li2O/SiO2 ratio was not conducive to generating a large number of liquid phases at sintered temperature for the completion of the sintering process under a lower temperature. Besides, LiAlSi2O6 and LiAlSi3O8 constituted as the major crystalline phases in B1–B4 glass-ceramic bonds, while SiO2 was the only precipitation in B5 when glass-ceramic bonds were sintered at 860 °C for 2 h, and a higher content of crystalline phase including LiAlSi2O6 and LiAlSi3O8 along with the absence of SiO2 precipitation would contribute in the lower CTE of glass-ceramic bonds sintered at 860 °C for 2 h. Lastly, the highest bending strength of cBN composites (100.53 MPa) relied on the sufficient wetting of B3 glassceramic bonds on cBN and the close CTE between cBN and B3 glass-

Fig. 12. SEM of glass-ceramic bonded cBN composites sintered at 860 °C for 2 h; (a)–(e) B1–B5. 15364

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ceramic bonds when composites were sintered at 860 °C for 2 h, indicating B3 is a kind of suitable and promising glass-ceramic bonds for cBN abrasives. Acknowledgments We express heartfelt thanks to Dr. Pijun Gong (State Key Laboratory of Silicate Materials for Architectures, Wuhan University of Technology, China) for the FE-SEM analyses. References [1] J.H. Liu, Z.J. Pei, G.R. Fisher, Grinding wheels for manufacturing of silicon wafers: a literature review, Int. J. Mach. Tool Manuf. 47 (2007) 1–13. [2] J.F.G. Oliveira, E.J. Silva, C. Guo, F. Hashimoto, Industrial challenges in grinding, CIRP Ann. - Manuf. Technol. 58 (2009) 663–680. [3] F. He, Q. Zhou, J. Xie, Q. Zhang, Characterization of low sintering temperature and high strength SiO2-B2O3-CaO vitrified bonds for diamond abrasive tools, Ceram. Int. 41 (2015) 3449–3455. [4] Z.X. Wang, T.B. Yu, J. Zhao, X.Z. Wang, X. Sun, S.X. Liu, Effect of Li2O, K2O and ZnO on vitrified bond composites for CBN grinding wheels, Mater. Sci. Forum 874 (2016) 193–198. [5] J. Yang, D.Y. Kim, H.Y. Kim, Effect of glass composition on the strength of vitreous bonded c-BN grinding wheels, Ceram. Int. 19 (1993) 87–92. [6] M.J. Jackson, B. Mills, Microscale wear of vitrified abrasive materials, J. Mater. Sci. 39 (2004) 2131–2143. [7] D. Feng, Y. Zhu, F. Li, Z. Li, Influence investigation of CaF2 on the LAS based glassceramics and the glass-ceramic/diamond composites, J. Eur. Ceram. Soc. 36 (2016) 2579–2585. [8] V.O. Soares, O. Peitl, E.D. Zanotto, New sintered-Al2O3-SiO2 Ultra‐low expansion glass‐ceramic, J. Am. Ceram. Soc. 96 (2013) 1143–1149. [9] B. Li, Z. Qing, Y. Li, H. Li, S. Zhang, Effect of CaO content on structure and properties of low temperature co-fired glass-ceramic in the Li2O-Al2O3-SiO2 system, J. Mater. Sci. Mater. Electron. 27 (2016) 2455–2459. [10] S. Huang, Y. Li, S. Wei, Z. Huang, W. Gao, P. Cao, A novel high-strength lithium disilicate glass-ceramic featuring a highly intertwined microstructure, J. Eur. Ceram. Soc. 37 (2017) 1083–1094. [11] B. Li, S. Wang, Y. Fang, Effect of Cr2O3 addition on crystallization, microstructure and properties of Li2O-Al2O3-SiO2 glass-ceramics, J. Alloy. Comp. 693 (2017) 9–15. [12] B. Li, Z. Qing, Y. Li, H. Li, S. Zhang, Effect of CaO content on structure and properties of low temperature co-fired glass–ceramic in the Li2O-Al2O3-SiO2 system, J. Mater. Sci. Mater. Electron. 27 (2016) 2455–2459. [13] A. Arvind, A. Sarkar, V.K. Shrikhande, A.K. Tyagi, G.P. Kothiyal, The effect of TiO2 addition on the crystallization and phase formation in lithium aluminum silicate (LAS) glasses nucleated by P2O5, J. Phys. Chem. Solids 69 (2008) 2622–2627. [14] X. Guo, H. Yang, M. Cao, Nucleation and crystallization behavior of Li2O-Al2O3SiO2 system glass-ceramic containing little fluorine and no-fluorine, J. Non-Cryst. Solids 351 (2005) 2133–2137. [15] Y. Wu, K. Hsu, C. Lee, Effects of B2O3 and P2O5 doping on the microstructure evolution and mechanical strength in a lithium aluminosilicate glass–ceramic material with TiO2 and ZrO2, Ceram. Int. 38 (2012) 4111–4121. [16] X. Zheng, G. Wen, L. Song, X.X. Huang, Effects of P2O5 and heat treatment on crystallization and microstructure in lithium disilicate glass ceramics, Acta Mater. 56 (2008) 549–558. [17] A.X. Lu, Z.B. Ke, Z.H. Xiao, X.F. Zhang, X.Y. Li, Effect of heat-treatment condition on crystallization behavior and thermal expansion coefficient of Li2O-ZnO-Al2O3-SiO2P2O5 glass-ceramics, J. Non-Cryst. Solids 353 (2007) 2692–2697. [18] J.M. Jewell, M.S. Spess, R.L. Ortolano, J.E. Shelby, Effect of heat-treatment temperature on the properties of a lithium aluminosilicate glass, J. Am. Ceram. Soc. 74 (1991) 92–97. [19] W.J. Zhang, Y.M. Chong, I. Bello, S. Lee, Nucleation, growth and characterization of cubic boron nitride (cBN) films, J. Phys. D Appl. Phys. 40 (2007) 6159–6174. [20] J. Shi, F. He, J. Xie, X. Liu, H. Yang, Effects of Na2O/BaO ratio on the structure and the physical properties of low-temperature glass-ceramic vitrified bonds, Ceram. Int. 44 (2018) 10871–10877. [21] L. Cui, X. Hao, X. Hu, A. Lu, Effects of Y2O3 addition on structure and properties of LZAS vitrified bond for CBN grinding tools application, Ceram. Int. 41 (2015) 9916–9922. [22] L. Cui, X. Hao, Y. Tang, Z. Zheng, Y. Yuan, A. Lu, Effect of B2O3 on physical properties of LZAS vitrified bond and mechanical properties of diamond composites, Int. J. Refract. Metals Hard Mater. 52 (2015) 50–54.

[23] X. Hu, L. Cui, T. Liu, Z. Zheng, Y. Tang, A. Lu, Crystallization and properties of B2O3 doped LZAS vitrified bond for diamond grinding tools, J. Non-Cryst. Solids 427 (2015) 69–75. [24] A. Gaddam, H.R. Fernandes, D.U. Tulyaganov, M.J. Ribeiro, J.M.F. Ferreira, The roles of P2O5 and SiO2/Li2O ratio on the network structure and crystallization kinetics of non-stoichiometric lithium disilicate based glasses, J. Non-Cryst. Solids 481 (2018) 512–521. [25] A.M. Efimov, Quantitative IR spectroscopy: applications to studying glass structure and properties, J. Non-Cryst. Solids 203 (1996) 1–11. [26] H. Doweidar, G. El-Damrawi, M. Al-Zaibani, Distribution of species in Na2O-CaOB2O3 glasses as probed by FTIR, Vib. Spectrosc. 68 (2013) 91–95. [27] P. Tarte, Infra-red spectra of inorganic aluminates and characteristic vibrational frequencies of AlO4 tetrahedra and AlO6 octahedra, Spectrochim. Acta, Part A 23 (1967) 2127–2143. [28] A.T.G. Kullberg, A.A.S. Lopes, R.C.C. Monteiro, Effect of ZnF2 addition on the crystallization behaviour of a zinc borosilicate glass, J. Non-Cryst. Solids 468 (2017) 100–104. [29] A. Winterstein-Beckmann, D. Möncke, D. Palles, E.I. Kamitsos, L. Wondraczek, A Raman-spectroscopic study of indentation-induced structural changes in technical alkali-borosilicate glasses with varying silicate network connectivity, J. Non-Cryst. Solids 405 (2014) 196–206. [30] C. Huang, E.C. Behrman, Structure and properties of calcium aluminosilicate glasses, J. Non-Cryst. Solids 128 (1991) 310–321. [31] R.B. Rao, R.A. Gerhardt, Effect of alkaline earth modifier ion on the optical, magnetic and electrical properties of lithium nickel borate glasses, Mater. Chem. Phys. 112 (2008) 186–197. [32] D.P. Mukherjee, A.R. Molla, S.K. Das, The influence of MgF2 content on the characteristic improvement of machinable glass ceramics, J. Non-Cryst. Solids 433 (2016) 51–59. [33] Y. Lai, Y. Zeng, X. Tang, H. Zhang, J. Han, H. Su, Structural investigation of calcium borosilicate glasses with varying Si/Ca ratios by infrared and Raman spectroscopy, RSC Adv. 6 (2016) 93722–93728. [34] V. John, Introduction to Engineering Materials, third ed., Macmillan, 1992. [35] V.K. Singh, Densification of alumina and silica in the presence of a liquid phase, J. Am. Ceram. Soc. 64 (1981) 133–136. [36] Y. Chen, E. Li, S. Duan, S. Zhang, Low temperature sintering kinetics and microwave dielectric properties of BaTi5O11 ceramic, Acs. Sustain. Chem. Eng. 5 (2017) 10606–10613. [37] H. Scholze, Glass: Nature, Structure, and Properties, Springer Science & Business Media, 2012. [38] J. Hao, Q. Zan, D. Ai, J. Ma, W. Guo, T. Liang, Theoretical calculations of the thermal expansion coefficient of glass-ceramic sealing materials in solid oxide electrolysis cell, J. Non-Cryst. Solids 361 (2013) 86–92. [39] Z. Qin, Research on Properties and Mechanism of Substrate Materials with Low Thermal Expansion Coefficient, Doctoral Thesis University of Electronic Science and Technology of China, 2016. [40] H.M. Rietveld, Line profiles of neutron powder-diffraction peaks for structure refinement, Acta Crystallogr. 22 (1967) 151–152. [41] H. Rietveld, A profile refinement method for nuclear and magnetic structures, J. Appl. Crystallogr. 2 (1969) 65–71. [42] S. Arcaro, B. Moreno, E. Chinarro, M.D. Salvador, A. Borrell, M.I. Nieto, R. Moreno, A.P.N.D. Oliveira, Properties of LZS/nanoAl2O3 glass-ceramic composites, J. Alloy. Comp. 710 (2017) 567–574. [43] M. Laczka, K. Laczka, K. Cholewa-Kowalska, A.B. Kounga, C. Appert, Mechanical properties of a lithium disilicate strengthened lithium aluminosilicate glassceramic, J. Am. Ceram. Soc. 97 (2013) 361–364. [44] F. Ye, J.C. Gu, Y. Zhou, M. Iwasa, Synthesis of BaAl2Si2O8 glass-ceramic by a sol-gel method and the fabrication of SiCpl/BaAl2Si2O8 composites, J. Eur. Ceram. Soc. 23 (2003) 2203–2209. [45] B. Li, W. Li, J. Zheng, Effect of SiO2 content on the sintering kinetics, microstructures and properties of BaO-Al2O3-B2O3-SiO2 glass-ceramics for LTCC application, J. Alloy. Comp. 725 (2017) 1091–1097. [46] I.L. Denry, J.A. Holloway, Effect of post-processing heat treatment on the fracture strength of a heat-pressed dental ceramic, J. Biomed. Mater. Res., Part B. 68B (2003) 174–179. [47] M. Laczka, K. Laczka, K. Cholewa Kowalska, A.B. Kounga, C. Appert, Mechanical properties of a lithium disilicate strengthened lithium aluminosilicate glass‐ceramic, J. Am. Ceram. Soc. 97 (2014) 361–364. [48] I. Bello, Y.M. Chong, K.M. Leung, C.Y. Chan, K.L. Ma, W.J. Zhang, S.T. Lee, A. Layyous, Cubic boron nitride films for industrial applications, Diam. Relat. Mater. 14 (2005) 1784–1790. [49] J. Shi, F. He, J. Xie, X. Liu, H. Yang, Effect of heat treatments on the Li2O-Al2O3SiO2-B2O3-BaO glass-ceramic bond and the glass-ceramic bond cBN grinding tools, Int. J. Refract. Metals Hard Mater. 78 (2019) 201–209.

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