3.1 Effects of ball-milling time on microstructure and dispersity of fibers
Fig. 1 shows the morphology of dispersed alumina fibers after the first-step ball-milling process in different ball-milling time. Most of the original fibers intertwine together and form cross-linked stable structure (Fig. 1 (a)), which results from electrostatic attraction between fibers. The dispersion of alumina fibers after the first-step ball-milling process for 5 minutes is improved obviously (Fig. 1 (b)), confirming that the short-time friction unties the cross-linked structure and disperses fibers effectively. As deteched in Fig. 1 (c), alumina fibers treated for 10 minutes lose their original fibrous structure, which may attribute to the drastic and long-time physical friction, promoting the cleavage of the inner fiber grains. Fig. 1 (d, e, f) shows the dispersion of alumina fibers in matrix after the second-step ball-milling process in different milling time. It can be found that ceramic powders obviously agglomerated distribute in-homogeneously around dispersed alumina fibers (Fig. 1 (d)), this is because original ceramic powders are too agglomerating to infiltrate evenly into the inner cross-linked space. Ulteriorly, the numerous inter-space between uneven matrix and alumina fibers may lead to weak bonding between fiber/matrix interface, which can't transfer the internal stress from matrix to fibers in time while carrying external loads [20-21]. Besides, the alumina fibers treated for 1 minute in the second-step ball-milling process are uniformly distributed in ceramic powders (Fig. 1 (e)), revealing the two-steps ball-milling process can achieve a well dispersion mixture of fibers and ceramic powders. Inversely, the fibrous structure of fibers is destroyed absolutely over 3 minutes (Fig. 1 (f)), which is consistent with the analysis in Fig. 1 (c).
Fig. 2 (a) shows the average length-diameter ratio of fibers after the first-step ball-milling process in different ball-milling time. The average length-diameter ratio of fibers decreases tardily from 83.8 to 58.9 with the ball-milling time from 0 minutes to 5 minutes, and then declines drastically from 58.9 to 4.6 with the time from 5 minutes to 10 min., illustrating that the friction between fibers and ballstone has a top priority to disperse the cross-linked structure, and then chiefly damages the original fibrous structure of fibers, decreasing the average length-diameter ratio. Notably, the alumina fibers with little average length-diameter ratio by long-time friction are difficult to deflect micro-crack paths effectively, which is adverse to promoting the mechanical property of the thin architectural ceramic plate [22]. In comparison, the average length-diameter ratio of fibers after the second-step ball-milling process decreases sharply from 58.9 to 7.6 with the time from 0 minutes to 3 minutes (Fig. 2 (b)). Besides, the drop rate of it is much faster than that in Fig. 2 (a), which is corresponding to the synergistic drastic friction of matrix, fibers and ballstone. Based on above results, the first-step ball-milling for 5 minutes followed by the second-step ball-milling for 1 minute is proved to be the best two-steps ball-milling process.
3.2 Effects of alumina fiber contents on the thin ceramic plate
Fig. 3 shows the optical photos of the as-prepared thin architectural ceramic plate. Meanwhile, the XRD patterns of samples with different alumina fibers contents are shown in Fig. 4, it is observed that the diffraction peaks of 25.7°, 35.2°, 38.1°, 43.5°, 52.6° and 57.9° agree well with the α-Al2O3 (PDF no. 50-0741), illustrating the main crystal phase of fibers transforms from θ-Al2O3 to α-Al2O3 by fast-sintering method [23]. In comparison, the alumina fibers of α-Al2O3 have higher tensile strength than the fibers of θ-Al2O3, further confirming that the fast-sintering method with short holding time is beneficial to promoting the mechanical properties of alumina fibers as the reinforcement instead of damaging them unilaterally. Besides, the diffraction peaks of 16.4°, 26.3°, 33.3°, 40.8° and 60.9° can be indexed to the mullite (PDF no. 15-0776), attributing to the appropriate phase ratio of alumina and silicon in glass phase. Furthermore, the mass ratio of alumina to silicon of the fibers is 7.85 which is much higher than that of the matrix of 1.51, the fused alumina phase in fibers may infiltrate into the glass phase in fast-sintering process, which collectively increases the alumina content in liquated glass phase and promotes the crystallization of mullite and whiskers [24,25].
Fig. 5 shows the cross-section SEM images of the as-prepared samples with different alumina fiber contents. It can be noticed that the particles without obvious coarse grains of the cross-section sintered at 1200℃ for 30 minutes is compact (Fig. 5 (a)), suggesting that the fast-sintering system is suitable. Meanwhile, the alumina fibers are perpendicular to the cross-section and embed tightly in the matrix (Fig. 5 (b)), this is because fibers and surrounding matrix are subjected to tensile and compressive stress respectively during the fast-sintering process, which results from thermal expansion coefficient difference between alumina fiber of 9.2×10-6/℃ and matrix of 4.8×10-6/℃ [26]. Namely, the existence of internal stress is beneficial to offset the external stress loads and improve the bending strength of the thin architectural ceramic plate. However, it is the matrix that mainly bears the loads at this moment limited by only 3 wt% fibers contents. As shown in Fig. 5 (c, d), the number of pores in matrix increases obviously with the fiber content more than 5 wt% and fibers in the cross-section of the thin ceramic plate disperse uniformly. Close observation reveals that there are numerous white substances in the pores of the matrix, which is presumed to be the in-situ fibrous mullite and whiskers conformed by the analysis in Fig. 4.
Fig. 6 shows the volume density and water absorption of the thin architectural ceramic plate with different alumina fiber contents. On the one hand, the volume density of the thin plate increases slowly from 2.722 g·cm-3 to 2.733 g·cm-3 with the fiber content from 0 wt% to 5 wt%, which mainly ascribes to the density of alumina fibers for 6.22 g·cm-3 is much higher than that of matrix for 2.722 g·cm-3. On the other hand, the volume density of the thin plate decreases from 2.782 g·cm-3 to 2.587 g·cm-3 with the fiber content from 5 wt% to 15 wt%, attributing to the inner pores increase in matrix. Simultaneously, the water absorption of the thin architectural ceramic plate decreases tardily from 0.172% to 0.152% with the fiber content from 0 wt% to 5 wt%, further confirming that the in-situ mullite and whiskers during the fast-sintering method can fill the inner pores and increase the compactness of ceramic matrix.
Fig. 7 shows the bending strength of the thin architectural ceramic plate with different alumina fiber contents. The bending strength of 0 wt%, 1 wt%, 3 wt%, 5 wt%, 7 wt%, 11 wt% and 15 wt% fiber contents are 86.2 MPa, 93.8 MPa, 122.4 MPa, 146.8 MPa, 136.2 MPa, 80.3 MPa and 51.5 MPa, respectively. Moreover, the maximum bending strength of the thin architectural ceramic plate reaches 146.8 MPa that is 70.31% higher than the blank sample of 86.2 MPa with 5 wt% fiber contents, which is corresponding to the strengthening of alumina fibers and the formation of mullite crystallization of the matrix during the fast-sintering process.
3.3 Microstructure analysis
Fig. 8 shows the cross-section SEM images of the as-prepared samples with 5 wt% alumina fiber content. Apparently in Fig. 8 (a, b), the expanding micro-cracks which are supposed to be single orientation are deflected to be discontinuous or winding orientation by inner particles and fibers, indicating that the reinforcements can slow down the prolongation speed of micro-cracks in matrix and retard the stress concentration in the micro-crack tip owing to the Griffith micro-crack theory [27,28]. Combined with the EDS analysis of spot 1 with the mass contents of Al, Si, O and Y are 36.95 wt%, 8.5 wt%, 41.02 wt% and 9.26 wt% (Fig. 8 (e)), illustrating that there are alumina fibers embed tightly in the matrix. In addition, the fiber/matrix interface has clear boundaries after the fast-sintering method and there is plenty of white substance in it, suggesting that the fast-sintering method is appropriate and it can prevent the fiber/matrix interface from excessive thermodynamics damage. To clarify the components of the white substance in fiber/matrix interface, the EDS analysis of spot 2 is analyzed in Fig. 8 (f), proving that it is the mullite with the mass contents of Al, Si and O are 41.22 wt%, 11.45 wt% and 42.86 wt%, severally. Fig. 8 (b, c and d) shows the distribution of reinforcements along the fracture micro-cracks. It is found that there are not only thick fibers but also finer fibrous substance at the micro-cracks fracture, both of them are bridge-pulled out of the matrix. Ulteriorly, as shown in Fig. 8 (g, h), the element mass contents of the thick fiber at spot 3 are 32.47 wt% Al, 10.58 wt% Si, 39.68 wt% O, 11.21 wt% Y and the finer fiber at spot 4 are 38.32 wt% Al, 14.18 wt% Si and 45.17 wt% O, respectively. It can be accurately inferred that the thick fiber is alumina fiber and the finer fiber is fibrous mullite that is precipitated out by crystallization in matrix. Hence, the reinforcements of alumina fibers, particles and in-situ fibrous mullite in matrix and fiber/matrix interface are beneficial to deflecting the paths of the micro-cracks (Fig. 9), which can synergistically slow down the prolongation speed of inner micro-cracks and enhance the mechanical property of the large-sized thin architectural ceramic plate.