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Mechanical behavior of fibrous ceramics with a bird’s nest structure
Composites Science and Technology 100 (2014) 92–98
Contents lists available at ScienceDirect
Composites Science and Technology journal homepage: www.elsevier.com/locate/compscitech
Mechanical behavior of fibrous ceramics with a bird’s nest structure Xue Dong, Guofa Sui, Jiachen Liu, Anran Guo, Sue Ren, Mingchao Wang, Haiyan Du ⇑ School of Materials Science and Engineering of Tianjin University, 92 Weijin Road, Nankai District, Tianjin 300072, PR China
a r t i c l e
i n f o
Article history: Received 14 October 2013 Received in revised form 1 June 2014 Accepted 5 June 2014 Available online 14 June 2014 Keywords: A. Short-fiber composites B. Mechanical properties B. Fracture C. Elastic properties
a b s t r a c t Inspired by the structure of the bird’s nest, a new fibrous ceramics with mullite fibers as the matrix and SiO2–B2O3 phase as the high-temperature binder was designed and synthesized. The most important structure feature of this fibrous material was that the binder only bonded the fibers at the crossing points. The effects of sintering temperature on the ceramic properties were studied. The fibrous ceramics exhibited significant higher elasticity and pseudoductility compared to the ceramic matrix composites reinforced with continuous fibers. The fracture mechanism of this ceramics under compression was discussed. A high porosity (74.2–78.3%), a low thermal conductivity (0.231–0.248 W/m K), a relatively high compressive strength (1.3–3.2 MPa) and a high rebound-resilience (90–98%) measured from samples indicate that this fibrous ceramics is a potential material for the high-temperature sealing. Ó 2014 Elsevier Ltd. All rights reserved.
1. Introduction The porous fibrous materials are considered to be a very interesting type of porous materials. On one hand, they are soft, porous, and voluminous. On the other hand, they show a relatively high resistance to mechanical deformation [1]. In addition, the fibrous porous materials require much less material (fiber) to form a stable structure because of the arrangement of the fibers, whereas the granular porous media require much more material (granules) [1–4]. Therefore, they are applied in numerous and diverse aspects such as catalyst supports, hot-gas filters, composite reinforcement, biomaterials, acoustic and thermal insulation and sealing materials [5]. Contrary to the monolithic ceramics which exhibit a brittle fracture, porous ceramics present a dissipative damage tolerant behavior [6–8]. Numerous research articles were published focusing on forming porous ceramics with special 3-D skeleton structures, such as the foams and cellular materials [9–12]. Besides, another branch of research works is devoted to developing new structures such as continuous fiber-reinforced matrix and multilayer laminates which could provide a high strength along the fiber direction but a weak strength across the fiber direction [13–16]. However, none of these materials satisfies the quasiplastic and elastic requirement of sealing materials due to the inherent brittleness [17,18]. The aim of our study is to design a new fibrous material with both high strength and attractive compression–rebound property [19–21]. ⇑ Corresponding author. Tel.: +86 13502170511; fax: +86 02227408244. E-mail address: [email protected] (H. Du). http://dx.doi.org/10.1016/j.compscitech.2014.06.001 0266-3538/Ó 2014 Elsevier Ltd. All rights reserved.
Highly porous fiber network materials are abundant in nature and in man-made environments. One convincing example is the bird’s nest which is made of randomly arranged tree branches. Inspired by this fact, the idea of designing a fiber matrix porous ceramic with fibers as the skeleton structure bonded by proper binders at the crossing points was brought. Mullite fiber has a good flexibility at relatively high stress, making it a promising candidate for preparing the sealing material [22,23]. In this research a new fibrous ceramic with the framework structure of polycrystalline mullite fibers was fabricated by infiltration method. During the preparation process, the organic and inorganic binders were impregnated in the fiber framework in order to bond the fibers with each other at room temperature and high temperature, respectively. The elasticity of FCFMF sintering at 1200 °C was investigated [24]. The study showed that the sample possessed a high degree of rebound resilience (98%) under a compression stress of 2 MPa. The present study is a continuous work of this previous study [24] and the effects of the sintering temperatures on the microstructure and physical and mechanical properties of the samples were analyzed. 2. Experiment 2.1. Raw materials Commercially available polycrystalline mullite refractory fiber (PMF, 99.5%, Zhejiang Hongda Crystal Fiber Co., Ltd., China) was used as the starting material in this study. The organic binder (OB) was prepared by mixing sodium dodecyl benzene sulfonate (SDBS) with polyacrylamide (CPAM) in water with the weight ratio
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SDBS:CPAM:H2O = 1:1:100. The inorganic binder (mixed sol) was made of silica sol and boric acid, with a final molar ratio of SiO2 to B2O3 of 10:1. The silica sol was produced with the use of tetraethylorthosilicate (TEOS, AR grade, Tianjin Kewei Chemical Co., China) by one-step catalytic method. The weight ratio was TEOS:H2O:ethanol:hydrochloride = 1:4:1:7.5 10 4. The boric acid was obtained by dissolving the solid H3BO3 in the water at room temperature with a weight ratio of H3BO3:H2O = 1:18. Fig. 1(a) shows the micrograph of the mixed gels used as inorganic binder. The diameter of the mixed gels is around 5 lm and the shape is approximately ellipse. Fig. 1(b) presents the micrograph of the mullite fibers. It can be seen that the fiber possessed a diameter in the range of 8–15 lm and a length in the range of 300–500 lm.
carried out at room temperature on an electro-universal testing machine (Instron 5569, USA) in accordance with GB/T 1964– 1996. During the test a set of loads were applied to the samples at a loading speed of 0.05 mm/min, and removed at an unloading speed of 0.05 mm/min. The compressive ratio and the rebound resilience were determined by the following equations: compressive ratio = [(t0 t1)/t0] 100%, rebound-resilience = [(t2 t1)/ (t0 t1)], where t0, t1 and t2 are the height of preloading, loading and unloading, respectively. Thermal conductivity at room temperature was measured by the thermal-conductivity instrument (C-3000, Xian Xiaxi Electric Co., Ltd., Shanxi, China). The dimensions of measured samples were 30 mm in diameter and 5 mm in height. Each value represented an average of five measurements of five different specimens.
2.2. Experimental procedure
3. Results and discussion
Fig. 2 shows the processing steps of ceramic composites preparation and the schematic diagram of the structure of the FCFMF. First, the organic binder (30 g) and PMF (6 g) were mixed together by stirring (Fig. 2(a)). The fibers formed into a fiber block by infiltration with the help of a certain amount of organic binder coated on the fiber surface (Fig. 2(b)). After infiltration, the fiber surface became negatively charged due to the formation of the negatively charged organic coating. Then the mixed sols were impregnated into the fiber block. Since the fiber surface and the mixed sols were both negatively charged, there was a great repulsion between the fiber surface and the mixed sols. Therefore, the mixed sols coated on the fiber surface can be easily taken away by infiltration. However, the mixed sols at the crossing points of the fibers were forced to stay in the original place after infiltration because of the big obstacles at the crossing points of the fibers (Fig. 2(c)). The green bodies were sintered at different temperatures from 1100 to 1400 °C for 2 h assisted by holding at 600 °C for 0.5 h to decompose the organic phase. After sintering, the organic binder was burned out and the mixed sols melt into continuous phase at the crossing points of the fibers, consequently acting as a high-temperature binder (Fig. 2(d)). The mullite fibers with the bonding points constituted a special 3-D skeleton structure.
3.1. The typical structure of the FCFMF
2.3. Characterization Phases were analyzed via X-ray diffraction (XRD, D/Max-2500 Rigaku, Japan). Microstructure of the sintered samples was observed by scanning electron microscope (SEM, XL-30Philips, Japan). Open porosities and densities of the sintered samples were determined by Archimedes method namely the water-immersion technique. The linear shrinkage of samples during drying and sintering process was determined by the following equation: shrinkage = [(la lb)/la] 100%, where la and lb were the diameter of initial samples and dried or sintered samples, respectively. Compression ratio and rebound resilience tests of the FCFMF were
Fig. 3(a) shows the typical SEM image of the FCFMF. After sintering, the organic binder was removed completely, and the fibers lapped with each other forming a loose skeleton structure. It could be clearly seen that the fibers were randomly arranged in the block providing the ceramics with isotropic mechanical properties. Meanwhile, the adjacent fibers were bonded by the surrounding melted SiO2–B2O3 binders at the crossing points which endowed the fibrous ceramics with relatively high strength. Fig. 3(b–d) shows the typical bonding points in the fiber block. This special bird’s nest structure provides this material with unique properties, which were discussed below. 3.2. Phase characterization The sintering temperature plays an important role in the state of SiO2–B2O3 among fibers and the wettability between the SiO2–B2O3 and the mullite fiber, thus determining the properties of the samples. Fig. 4 shows the XRD patterns of the green body and samples sintered at different temperatures. A wide peak was observed in the green body and the peak of mullite was not very strong, which suggest that the fibers were covered by the organics. When the temperature was below 1200 °C, the patterns of the sintered bodies composed of mullite characteristic peaks and a broad background which indicates that the silica–boron gel transformed to the glassy phase. However, when the sintering temperature was above 1200 °C, the patterns of the sample exhibited the a-cristobalite characteristic peaks besides the mullite characteristic peaks. With the temperature increasing from 1200 to 1400 °C, the intensity of cristobalite increased gradually. The existence of the a-cristobalite characteristic peaks proved that silica-based glassy phase crystallized into the b-cristobalite at 1300 °C and then the b-cristobalite underwent a b ? a displacive phase transition when
Fig. 1. The micrograph of (a) the silica–boron gels used as binder and (b) the fibers.
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Fig. 2. The processing steps of ceramic composites preparation and schematic diagram of the structure of the FCFMF.
Fig. 3. Typical SEM image of (a) the FCFMF and (b–d) the fiber framework bonded by melted SiO2–B2O3.
3.3. Porosity, bulk density, linear shrinkage and microstructure
Fig. 4. XRD patterns of green body and samples sintered at different temperatures. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
the temperature decreased to room temperature [25]. This transition was often accompanied by a 3–5% volume change and might generate thermal stresses, resulting in a cracking of the binder, which might affect the compressive strength of the sample [25].
As the fibers were bonded by the inorganic binders and organic binders at the crossing points of the fibers, the green body with the strength of 1.0 MPa was strong enough to hand and machine before sintering. Fig. 5 shows the variation of linear shrinkages, porosities and bulk densities of the FCFMF with the increasing of sintering temperature. It was found that the shrinkage of all the samples sintered at high temperature were lower than 5% which is quite different from that of the traditional ceramics. This small shrinkage can be attributed to the formation of the rigid fibrous skeleton structure constructed by the randomly arranged mullite fibers [26,27]. In addition, as the sintering temperature increased from 1100 to 1400 °C, the open porosity decreased from 78.3% to 74.2%, and the bulk density increased from 0.66 to 0.73 g/cm3. Fig. 6 shows SEM micrographs of the FCFMF sintering at different temperatures. As shown in Fig. 6(a), the fibers overlapped with each other constructing a 3-D skeleton structure, forming a large amount of large pores in the fibrous ceramic bodies. The adjacent fibers were bonded at the crossing points by the surrounding SiO2–B2O3 phase. When the sintering temperature was 1100 °C, the SiO2–B2O3 phase was still discontinuous and the bonding strength between the fibers and the SiO2–B2O3 phase was very
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became strong, which can be observed from Fig. 6(f). Besides, it can be concluded that the slight decrease of porosity (from 78.3% to 74.2%) was mainly due to the increasing melting of the SiO2–B2O3 phase between the fibers. 3.4. Compressive strength and thermal conductivity
Fig. 5. Variation of linear shrinkage, porosity and density of the FCFMF with sintering temperature.
weak, which can be seen from Fig. 6(e). Further increasing of the sintering temperature could accelerate the melting of the SiO2– B2O3 phase which can be seen from Fig. 6(b–d). When the sintering temperature was 1400 °C, the SiO2–B2O3 phase converted to a continuous phase and the bonding strength between fiber/binder
Fig. 7 shows the effect of sintering temperature on the compressive strength and room-temperature thermal conductivity. It demonstrated that the compressive strength increased from 1.3 to 3.2 MPa when the temperature increased from 1100 to 1400 °C. It can be observed from Fig. 6(e) that the SiO2–B2O3 glassy phase slightly melted and lapped loosely with each other, forming a discontinuous phase. And the bonds between the SiO2–B2O3 phase and the fibers were very weak, resulting in a lower compressive strength (1.3 MPa). When the temperature increased, the SiO2–B2O3 phase further melted and the verification of SiO2–B2O3 phase became more and more obvious. Meanwhile, the adjacent fibers were adhered by the melted SiO2–B2O3 phase (Fig. 6(b–d)). The increase of cohesive force between the melted SiO2–B2O3 phase and the fibers finally lead to the increase of the compressive strength (up to 3.2 MPa). The good bonding interface between fibers and inorganic binder resulted in a rigid contiguous network
Fig. 6. SEM micrograph of fracture surface of the FCFMF sintering at different temperatures: (a) 1100 °C, (b) 1200 °C, (c) 1300 °C, (d) 1400 °C, (e and f) the magnification of samples sintering at 1100 and 1400 °C.
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the thermal conductivity may be caused by the decrease of the porosity of the samples. 3.5. Mechanical behavior
Fig. 7. Variation of compressive strength and thermal conductivity of the FCFMF with sintering temperature.
structure which allowed the FCFMF to sustain a relatively high load. In addition, the result shows that the b ? a displacive phase transition of SiO2 caused by the changing temperature did no harm to the compressive strength of the samples. It can also be seen from Fig. 7 that all samples had rather low thermal conductivities from 0.231 to 0.248 W/m K, which made this material a potential candidate as the high-temperature thermal insulator. The low thermal conductivity was mainly attributed to the unique interconnected pores built by the arbitrarily arranged fibers in the fibrous ceramic bodies. With the sintering temperature increasing from 1100 to 1400 °C, a slight increase of
The unique 3-D skeleton structure of mullite fiber in the FCFMF leads to the special stress transfer process. When the load was imposed perpendicularly to the sample, the stress which brought about the bending of the fibers would firstly develop at both ends of the sample, and then gradually transferred inside in two paths. If the bonding between the fiber/binder was very firm, the stress can effectively transfer inside the sample. However, if the bonding between the fiber/binder was weak, the stress would firstly cause the densification of both ends of the sample and then gradually transferred inside the sample, finally leading to the compaction of the whole sample. Therefore, the fracture mechanisms of the materials would be attributed to the debonding of fiber/binder, binder fracture and fiber break. Fig. 8(a) shows the stress–strain curves of the FCFMF sintering at different temperatures ranging from 1100 to 1400 °C. It can be observed that the curves of all samples sintering at different temperatures could be generally divided into two stages: an approximately linear stage and a compaction stage. At the approximately linear stage, the deformation of the sample belonged to elastic deformation which was mainly caused by the bend of the fibers. The cracks of the fibers and other defects (for example, the defects caused by the b ? a displacive phase transition) in the samples introduced few significant damages. With the strain exceeding the elastic strain, the curve experienced a long compaction stage in which the stress increased in a nonlinear path. At this stage, the cracks in the samples kept developing and finally led to the damage. Fig. 8(b) shows the variation of Young’s modulus of the
Fig. 8. (a) The stress–strain curves and (b) variation of Young’s modulus of the FCFMF sintering at different temperatures.
Fig. 9. The microstructure of the FCFMF sintering at (a) 1100 °C and (b) 1400 °C after compression.
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Fig. 10. (a) Variation of stress and the rebound-resilience of the FCFMF with sintering temperature, (b) compression rebound curve of the FCFMF sintering at 1300 °C with 7 different cycles, and (c) overall view of the FCFMF sintering at 1300 °C after 7 cycles of compressions.
FCFMF with sintering temperature. The Young’s modulus of the samples increased from 8.7 to 30.5 MPa as the sintering temperature increased from 1100 to 1400 °C, which is significantly lower than that of many oxide ceramic fiber–matrix composites [28,29]. For example, it is reported that N610/AS porous-matrix composite which has a porous aluminosilicate matrix reinforced by 3M Nextel 610™ alumina fibers woven possesses a high elastic modulus of 70 GPa. The FCFMF exhibiting low elastic modulus indicates it is a potential elastic material. Fig. 9 shows the microstructure of the FCFMF sintering at different temperatures after compression. The significant fiber/binder interface debonding after compression as shown in Fig. 9(a) indicated that the fiber/binder bonding was relatively weak. It should be pointed out that this relatively weak bonding was beneficial to strength of the sample, since the formation of the cracks between the binder and the fibers might consume the energy. It can be seen from Fig. 9(b) that there was no evident fiber/binder debonding in the matrix, indicating the strong fiber/binder bond strength. However, an extensive binder fracture and fiber breaks were observed after compression. In this case, the possible failure of the sample was attributed to the crack that grew in the fibers and at the bonding points. The fibers between the strong bonding points would fracture firstly, and then the binder between fibers fracture in a brittle way as the crack further propagated. The compression–rebound tests were carried out below the elastic limit and all of the compression rebound data of the FCFMF sintering at different temperatures were tested by setting the compressive ratio as 10%. Fig. 10(a) shows the variation of stress and the rebound-resilience of the FCFMF with sintering temperature. The stress of the samples sintering at 1100 °C, 1200 °C ,1300 °C and 1400 °C with compression ratio of 10% were 1.2 MPa, 1.5 MPa, 2.0 MPa and 2.8 MPa, respectively. The increase of the stress indicated that the strength between the mullite fiber and the binder increased with the sintering temperature. The rebound-resilience decreased from 98% to 90% as the sintering temperature increased from 1100 to 1400 °C, illustrating that the samples were not completely elastomer at the elastic stage. Fig. 10(b) shows the compression rebound curve of the FCFMF sintering at 1300 °C with 7 different cycles. With the stress increased, the strain and the rebound-resilience of the sample increased and decreased, respectively. The higher rebound-resilience of sample was mainly due to the complete resilience of the bended fibers during the loading process. On the contrary, the lower reboundresilience was accounted for the fracture of the bonding points and deflects in the material. Fig. 10(c) shows the overall view of sample sintering at 1300 °C after 7 cycles of compressions. It can be observed that the rebound-resilience was high during the first
six cycles. In this period, the fracture of the bonding points and deflects in the material take up only a small amount. However, at the seventh cycle of compression, the amount of the fracture of the bonding points and the fibers increased and the reboundresilience decreased to 70%. 4. Conclusion A new fibrous ceramic was prepared by infiltration method, in which the mullite fiber was used as the matrix and the SiO2– B2O3 as the high-temperature binder. The unique bird’s nest like structure (a 3-D skeleton structure) was the most important designed structure characteristic of this material in this work. The fibrous structure with fixed bonding points exhibited a high degree of elasticity and pseudoductility compared to the ceramic matrix composites reinforced with continuous fibers. The sintering temperature plays an important part in tailoring the properties of this fibrous ceramics. As the sintering temperature increased from 1100 to 1400 °C, the porosity decreased from 78.3% to 74.2%. Correspondingly, the compressive strength and the thermal conductivity increased from 1.3 to 3.2 MPa, and from 0.231 to 0.248 W/ m K, respectively. The samples showed a high rebound-resilience of above 90% after compression of 10% ratio. Such bird’s nest structure obtained from mullite fibers bonded by the SiO2–B2O3 at the crossing points could therefore be a potential high-temperature sealing material. Acknowledgement This study is supported by the National Natural Science Foundation of China (Project Nos. 51272171 and 51372164) for financial support. References [1] Pradan AK, Das D, Chattopadhyay R, Singh SN. Effect of 3D fiber orientation distribution on transverse air permeability of fibrous porous media. Powder Technol 2012;221:101–4. [2] Jackson GW, James DF. The permeability of fibrous porous media. Can J Chem Eng 1986;64:364–74. [3] Tahir MA, Tafreshi HV. Influence of fiber orientation on the transverse permeability of fibrous media. Phys Fluids 2009;21:083604. [4] Tsay RY, Weinbaum S. Viscous flow in a channel with periodic cross-bridging fibers: exact solutions a Brinkman approximation. J Fluid Mech 1991;226:125–48. [5] Gibson LJ, Ashby MF. Cellular Solids: Structure and Properties. Cambridge: Cambridge University Press; 2001. [6] Levine SR, Opila EJ, Haibig MC, Kiser JD, Singh M, Salem JA. Evaluation of ultrahigh temperature ceramics for aeropropulsion use. J Eur Ceram Soc 2002;22:2757–67.
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[7] Chen IW, Xue LA. Development of super plastic structural ceramics. J Am Ceram Soc 1990;73:2585–609. [8] Ye F, Liu M, Huang LJ. Fabrication and mechanical properties of carbon short fiber reinforced barium aluminosilicate glass–ceramic matrix composites. Compos Sci Technol 2008;68:1710–7. [9] Guo A, Liu J, Wang Y, Xu H. Preparation of porous lamellar mullite ceramics with whisker skeletons by electrospinning and pressure molding. Mater Lett 2012;74:107–10. [10] Jang WY, Kraynik AM, Kyriakides S. On the microstructure of open-cell foams and its effect on elastic properties. Int J Solid Struct. 2008;45:1845–75. [11] Gong L, Kyriakides S, Jang WY. Compressive response of open-cell foams. Part I: Morphology and elastic properties. Int J Solid Struct 2005;42:1355–79. [12] Kanaun S, Tkachenko O. Effective conductive properties of open-cell foams. Int J Eng Sci 2008;46:551–71. [13] Li C, Chen Z, Zhu J, Liu Y, Jiang Y, Guan T. Mechanical properties and microstructure of 3D orthogonal quartz fiber reinforced silica composite fabricated by silica sol-infiltration-sintering. Mater Des 2012;36:289–95. [14] Yang W, Fuso L, Biamino S, Vasquez D, Bolivar CV, Fino P. Fabrication of short carbon fiber reinforced SiC multilayer composites by tape casting. Ceram Int 2012;38:1011–8. [15] Kaya C, Butler EG, Selcuk A, Boccaccini AR, Lewis MH. Mullite (Nextel™ 720) fiber-reinforced mullite matrix composites exhibiting favorable thermomechanical properties. J Eur Ceram Soc 2002;22:2333–42. [16] Stoll E, Mahr P, Kruger HG, Kern H, Thomas BJC, Boccaccini AR. Fabrication technologies for oxide–oxide ceramic matrix composites based on electrophoretic deposition. J Eur Ceram Soc 2006;26:1567–76. [17] Heinrich G, Kluppel M. Recent advances in the theory of filler net working in elastomers. Adv Polym Sci 2002;160:1–44. [18] Pelrine R, Kornbluh R, Pei QB, Joseph J. High-speed electrically actuated elastomers with strain greater than100%. Sci Mag 2000;287:836–9.
[19] Mouchon E, Colomban P. Oxide ceramic matrix/oxide fiber woven fabric composites exhibiting dissipative fracture behavior. Composites 1995;26:175–82. [20] Schmucker M, Grafmuller A, Schneider H. Mesostructure of WHIPOX all oxide CMCs. Compos Part A 2003;34:613–22. [21] Zok FW, Levi CG. Mechanical properties of porous matrix ceramic comporous structures fabricated by freeze-gelcasting: using tert-butyl alcohol posites. Adv Eng Mater 2001;3:15–23. [22] Wannaparhun S, Seal S. Combined spectroscopic and thermodynamic investigation of Nextel-720 fiber/alumina ceramic-matrix composite in air and water vapor at 1100 °C. J Am Ceram Soc 2003;86:1628–30. [23] Chawla KK. Interface engineering in mullite fiber/mullite matrix composites. J Eur Ceram Soc 2008;28:447–53. [24] Dong X, Liu J, Hao R, Guo A, Hou Z, Liu M. High-temperature elasticity of fibrous ceramics with a bird’s nest structure. J Eur Ceram Soc 2013;33:3477–81. [25] Chao C, Lu H. Stress-induced b ? a-cristobalite phase transformation in (Na2O + Al2O3)-codoped silica. Mater Sci Eng, A 2002;328:267–76. [26] Hatch DM, Ghose S. The b-phase transition in cristobalite, SiO2: sym-metry analysis, domain structure, and the dynamical nature of the b phase. Phys Chem Miner 1991;17:554–62. [27] Barr SA, Luijten E. Structural properties of materials created through freeze casting. Acta Mater 2010;58:709–15. [28] Kaya C, Kaya F, Butler EG, Boccaccini AR, Chawla KK. Development and characterization of high-density oxide fiber-reinforced oxide ceramic matrix composites with improved mechanical properties. J Eur Ceram Soc 2009;29:1631–9. [29] Larry PZ, Randall SH, Shin SL, James S. Characterization and high-temperature mechanical behavior of an oxide/oxide composite. J Am Ceram Soc 2003;86:981–90.
Contents lists available at ScienceDirect
Composites Science and Technology journal homepage: www.elsevier.com/locate/compscitech
Mechanical behavior of fibrous ceramics with a bird’s nest structure Xue Dong, Guofa Sui, Jiachen Liu, Anran Guo, Sue Ren, Mingchao Wang, Haiyan Du ⇑ School of Materials Science and Engineering of Tianjin University, 92 Weijin Road, Nankai District, Tianjin 300072, PR China
a r t i c l e
i n f o
Article history: Received 14 October 2013 Received in revised form 1 June 2014 Accepted 5 June 2014 Available online 14 June 2014 Keywords: A. Short-fiber composites B. Mechanical properties B. Fracture C. Elastic properties
a b s t r a c t Inspired by the structure of the bird’s nest, a new fibrous ceramics with mullite fibers as the matrix and SiO2–B2O3 phase as the high-temperature binder was designed and synthesized. The most important structure feature of this fibrous material was that the binder only bonded the fibers at the crossing points. The effects of sintering temperature on the ceramic properties were studied. The fibrous ceramics exhibited significant higher elasticity and pseudoductility compared to the ceramic matrix composites reinforced with continuous fibers. The fracture mechanism of this ceramics under compression was discussed. A high porosity (74.2–78.3%), a low thermal conductivity (0.231–0.248 W/m K), a relatively high compressive strength (1.3–3.2 MPa) and a high rebound-resilience (90–98%) measured from samples indicate that this fibrous ceramics is a potential material for the high-temperature sealing. Ó 2014 Elsevier Ltd. All rights reserved.
1. Introduction The porous fibrous materials are considered to be a very interesting type of porous materials. On one hand, they are soft, porous, and voluminous. On the other hand, they show a relatively high resistance to mechanical deformation [1]. In addition, the fibrous porous materials require much less material (fiber) to form a stable structure because of the arrangement of the fibers, whereas the granular porous media require much more material (granules) [1–4]. Therefore, they are applied in numerous and diverse aspects such as catalyst supports, hot-gas filters, composite reinforcement, biomaterials, acoustic and thermal insulation and sealing materials [5]. Contrary to the monolithic ceramics which exhibit a brittle fracture, porous ceramics present a dissipative damage tolerant behavior [6–8]. Numerous research articles were published focusing on forming porous ceramics with special 3-D skeleton structures, such as the foams and cellular materials [9–12]. Besides, another branch of research works is devoted to developing new structures such as continuous fiber-reinforced matrix and multilayer laminates which could provide a high strength along the fiber direction but a weak strength across the fiber direction [13–16]. However, none of these materials satisfies the quasiplastic and elastic requirement of sealing materials due to the inherent brittleness [17,18]. The aim of our study is to design a new fibrous material with both high strength and attractive compression–rebound property [19–21]. ⇑ Corresponding author. Tel.: +86 13502170511; fax: +86 02227408244. E-mail address: [email protected] (H. Du). http://dx.doi.org/10.1016/j.compscitech.2014.06.001 0266-3538/Ó 2014 Elsevier Ltd. All rights reserved.
Highly porous fiber network materials are abundant in nature and in man-made environments. One convincing example is the bird’s nest which is made of randomly arranged tree branches. Inspired by this fact, the idea of designing a fiber matrix porous ceramic with fibers as the skeleton structure bonded by proper binders at the crossing points was brought. Mullite fiber has a good flexibility at relatively high stress, making it a promising candidate for preparing the sealing material [22,23]. In this research a new fibrous ceramic with the framework structure of polycrystalline mullite fibers was fabricated by infiltration method. During the preparation process, the organic and inorganic binders were impregnated in the fiber framework in order to bond the fibers with each other at room temperature and high temperature, respectively. The elasticity of FCFMF sintering at 1200 °C was investigated [24]. The study showed that the sample possessed a high degree of rebound resilience (98%) under a compression stress of 2 MPa. The present study is a continuous work of this previous study [24] and the effects of the sintering temperatures on the microstructure and physical and mechanical properties of the samples were analyzed. 2. Experiment 2.1. Raw materials Commercially available polycrystalline mullite refractory fiber (PMF, 99.5%, Zhejiang Hongda Crystal Fiber Co., Ltd., China) was used as the starting material in this study. The organic binder (OB) was prepared by mixing sodium dodecyl benzene sulfonate (SDBS) with polyacrylamide (CPAM) in water with the weight ratio
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SDBS:CPAM:H2O = 1:1:100. The inorganic binder (mixed sol) was made of silica sol and boric acid, with a final molar ratio of SiO2 to B2O3 of 10:1. The silica sol was produced with the use of tetraethylorthosilicate (TEOS, AR grade, Tianjin Kewei Chemical Co., China) by one-step catalytic method. The weight ratio was TEOS:H2O:ethanol:hydrochloride = 1:4:1:7.5 10 4. The boric acid was obtained by dissolving the solid H3BO3 in the water at room temperature with a weight ratio of H3BO3:H2O = 1:18. Fig. 1(a) shows the micrograph of the mixed gels used as inorganic binder. The diameter of the mixed gels is around 5 lm and the shape is approximately ellipse. Fig. 1(b) presents the micrograph of the mullite fibers. It can be seen that the fiber possessed a diameter in the range of 8–15 lm and a length in the range of 300–500 lm.
carried out at room temperature on an electro-universal testing machine (Instron 5569, USA) in accordance with GB/T 1964– 1996. During the test a set of loads were applied to the samples at a loading speed of 0.05 mm/min, and removed at an unloading speed of 0.05 mm/min. The compressive ratio and the rebound resilience were determined by the following equations: compressive ratio = [(t0 t1)/t0] 100%, rebound-resilience = [(t2 t1)/ (t0 t1)], where t0, t1 and t2 are the height of preloading, loading and unloading, respectively. Thermal conductivity at room temperature was measured by the thermal-conductivity instrument (C-3000, Xian Xiaxi Electric Co., Ltd., Shanxi, China). The dimensions of measured samples were 30 mm in diameter and 5 mm in height. Each value represented an average of five measurements of five different specimens.
2.2. Experimental procedure
3. Results and discussion
Fig. 2 shows the processing steps of ceramic composites preparation and the schematic diagram of the structure of the FCFMF. First, the organic binder (30 g) and PMF (6 g) were mixed together by stirring (Fig. 2(a)). The fibers formed into a fiber block by infiltration with the help of a certain amount of organic binder coated on the fiber surface (Fig. 2(b)). After infiltration, the fiber surface became negatively charged due to the formation of the negatively charged organic coating. Then the mixed sols were impregnated into the fiber block. Since the fiber surface and the mixed sols were both negatively charged, there was a great repulsion between the fiber surface and the mixed sols. Therefore, the mixed sols coated on the fiber surface can be easily taken away by infiltration. However, the mixed sols at the crossing points of the fibers were forced to stay in the original place after infiltration because of the big obstacles at the crossing points of the fibers (Fig. 2(c)). The green bodies were sintered at different temperatures from 1100 to 1400 °C for 2 h assisted by holding at 600 °C for 0.5 h to decompose the organic phase. After sintering, the organic binder was burned out and the mixed sols melt into continuous phase at the crossing points of the fibers, consequently acting as a high-temperature binder (Fig. 2(d)). The mullite fibers with the bonding points constituted a special 3-D skeleton structure.
3.1. The typical structure of the FCFMF
2.3. Characterization Phases were analyzed via X-ray diffraction (XRD, D/Max-2500 Rigaku, Japan). Microstructure of the sintered samples was observed by scanning electron microscope (SEM, XL-30Philips, Japan). Open porosities and densities of the sintered samples were determined by Archimedes method namely the water-immersion technique. The linear shrinkage of samples during drying and sintering process was determined by the following equation: shrinkage = [(la lb)/la] 100%, where la and lb were the diameter of initial samples and dried or sintered samples, respectively. Compression ratio and rebound resilience tests of the FCFMF were
Fig. 3(a) shows the typical SEM image of the FCFMF. After sintering, the organic binder was removed completely, and the fibers lapped with each other forming a loose skeleton structure. It could be clearly seen that the fibers were randomly arranged in the block providing the ceramics with isotropic mechanical properties. Meanwhile, the adjacent fibers were bonded by the surrounding melted SiO2–B2O3 binders at the crossing points which endowed the fibrous ceramics with relatively high strength. Fig. 3(b–d) shows the typical bonding points in the fiber block. This special bird’s nest structure provides this material with unique properties, which were discussed below. 3.2. Phase characterization The sintering temperature plays an important role in the state of SiO2–B2O3 among fibers and the wettability between the SiO2–B2O3 and the mullite fiber, thus determining the properties of the samples. Fig. 4 shows the XRD patterns of the green body and samples sintered at different temperatures. A wide peak was observed in the green body and the peak of mullite was not very strong, which suggest that the fibers were covered by the organics. When the temperature was below 1200 °C, the patterns of the sintered bodies composed of mullite characteristic peaks and a broad background which indicates that the silica–boron gel transformed to the glassy phase. However, when the sintering temperature was above 1200 °C, the patterns of the sample exhibited the a-cristobalite characteristic peaks besides the mullite characteristic peaks. With the temperature increasing from 1200 to 1400 °C, the intensity of cristobalite increased gradually. The existence of the a-cristobalite characteristic peaks proved that silica-based glassy phase crystallized into the b-cristobalite at 1300 °C and then the b-cristobalite underwent a b ? a displacive phase transition when
Fig. 1. The micrograph of (a) the silica–boron gels used as binder and (b) the fibers.
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Fig. 2. The processing steps of ceramic composites preparation and schematic diagram of the structure of the FCFMF.
Fig. 3. Typical SEM image of (a) the FCFMF and (b–d) the fiber framework bonded by melted SiO2–B2O3.
3.3. Porosity, bulk density, linear shrinkage and microstructure
Fig. 4. XRD patterns of green body and samples sintered at different temperatures. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
the temperature decreased to room temperature [25]. This transition was often accompanied by a 3–5% volume change and might generate thermal stresses, resulting in a cracking of the binder, which might affect the compressive strength of the sample [25].
As the fibers were bonded by the inorganic binders and organic binders at the crossing points of the fibers, the green body with the strength of 1.0 MPa was strong enough to hand and machine before sintering. Fig. 5 shows the variation of linear shrinkages, porosities and bulk densities of the FCFMF with the increasing of sintering temperature. It was found that the shrinkage of all the samples sintered at high temperature were lower than 5% which is quite different from that of the traditional ceramics. This small shrinkage can be attributed to the formation of the rigid fibrous skeleton structure constructed by the randomly arranged mullite fibers [26,27]. In addition, as the sintering temperature increased from 1100 to 1400 °C, the open porosity decreased from 78.3% to 74.2%, and the bulk density increased from 0.66 to 0.73 g/cm3. Fig. 6 shows SEM micrographs of the FCFMF sintering at different temperatures. As shown in Fig. 6(a), the fibers overlapped with each other constructing a 3-D skeleton structure, forming a large amount of large pores in the fibrous ceramic bodies. The adjacent fibers were bonded at the crossing points by the surrounding SiO2–B2O3 phase. When the sintering temperature was 1100 °C, the SiO2–B2O3 phase was still discontinuous and the bonding strength between the fibers and the SiO2–B2O3 phase was very
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became strong, which can be observed from Fig. 6(f). Besides, it can be concluded that the slight decrease of porosity (from 78.3% to 74.2%) was mainly due to the increasing melting of the SiO2–B2O3 phase between the fibers. 3.4. Compressive strength and thermal conductivity
Fig. 5. Variation of linear shrinkage, porosity and density of the FCFMF with sintering temperature.
weak, which can be seen from Fig. 6(e). Further increasing of the sintering temperature could accelerate the melting of the SiO2– B2O3 phase which can be seen from Fig. 6(b–d). When the sintering temperature was 1400 °C, the SiO2–B2O3 phase converted to a continuous phase and the bonding strength between fiber/binder
Fig. 7 shows the effect of sintering temperature on the compressive strength and room-temperature thermal conductivity. It demonstrated that the compressive strength increased from 1.3 to 3.2 MPa when the temperature increased from 1100 to 1400 °C. It can be observed from Fig. 6(e) that the SiO2–B2O3 glassy phase slightly melted and lapped loosely with each other, forming a discontinuous phase. And the bonds between the SiO2–B2O3 phase and the fibers were very weak, resulting in a lower compressive strength (1.3 MPa). When the temperature increased, the SiO2–B2O3 phase further melted and the verification of SiO2–B2O3 phase became more and more obvious. Meanwhile, the adjacent fibers were adhered by the melted SiO2–B2O3 phase (Fig. 6(b–d)). The increase of cohesive force between the melted SiO2–B2O3 phase and the fibers finally lead to the increase of the compressive strength (up to 3.2 MPa). The good bonding interface between fibers and inorganic binder resulted in a rigid contiguous network
Fig. 6. SEM micrograph of fracture surface of the FCFMF sintering at different temperatures: (a) 1100 °C, (b) 1200 °C, (c) 1300 °C, (d) 1400 °C, (e and f) the magnification of samples sintering at 1100 and 1400 °C.
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the thermal conductivity may be caused by the decrease of the porosity of the samples. 3.5. Mechanical behavior
Fig. 7. Variation of compressive strength and thermal conductivity of the FCFMF with sintering temperature.
structure which allowed the FCFMF to sustain a relatively high load. In addition, the result shows that the b ? a displacive phase transition of SiO2 caused by the changing temperature did no harm to the compressive strength of the samples. It can also be seen from Fig. 7 that all samples had rather low thermal conductivities from 0.231 to 0.248 W/m K, which made this material a potential candidate as the high-temperature thermal insulator. The low thermal conductivity was mainly attributed to the unique interconnected pores built by the arbitrarily arranged fibers in the fibrous ceramic bodies. With the sintering temperature increasing from 1100 to 1400 °C, a slight increase of
The unique 3-D skeleton structure of mullite fiber in the FCFMF leads to the special stress transfer process. When the load was imposed perpendicularly to the sample, the stress which brought about the bending of the fibers would firstly develop at both ends of the sample, and then gradually transferred inside in two paths. If the bonding between the fiber/binder was very firm, the stress can effectively transfer inside the sample. However, if the bonding between the fiber/binder was weak, the stress would firstly cause the densification of both ends of the sample and then gradually transferred inside the sample, finally leading to the compaction of the whole sample. Therefore, the fracture mechanisms of the materials would be attributed to the debonding of fiber/binder, binder fracture and fiber break. Fig. 8(a) shows the stress–strain curves of the FCFMF sintering at different temperatures ranging from 1100 to 1400 °C. It can be observed that the curves of all samples sintering at different temperatures could be generally divided into two stages: an approximately linear stage and a compaction stage. At the approximately linear stage, the deformation of the sample belonged to elastic deformation which was mainly caused by the bend of the fibers. The cracks of the fibers and other defects (for example, the defects caused by the b ? a displacive phase transition) in the samples introduced few significant damages. With the strain exceeding the elastic strain, the curve experienced a long compaction stage in which the stress increased in a nonlinear path. At this stage, the cracks in the samples kept developing and finally led to the damage. Fig. 8(b) shows the variation of Young’s modulus of the
Fig. 8. (a) The stress–strain curves and (b) variation of Young’s modulus of the FCFMF sintering at different temperatures.
Fig. 9. The microstructure of the FCFMF sintering at (a) 1100 °C and (b) 1400 °C after compression.
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Fig. 10. (a) Variation of stress and the rebound-resilience of the FCFMF with sintering temperature, (b) compression rebound curve of the FCFMF sintering at 1300 °C with 7 different cycles, and (c) overall view of the FCFMF sintering at 1300 °C after 7 cycles of compressions.
FCFMF with sintering temperature. The Young’s modulus of the samples increased from 8.7 to 30.5 MPa as the sintering temperature increased from 1100 to 1400 °C, which is significantly lower than that of many oxide ceramic fiber–matrix composites [28,29]. For example, it is reported that N610/AS porous-matrix composite which has a porous aluminosilicate matrix reinforced by 3M Nextel 610™ alumina fibers woven possesses a high elastic modulus of 70 GPa. The FCFMF exhibiting low elastic modulus indicates it is a potential elastic material. Fig. 9 shows the microstructure of the FCFMF sintering at different temperatures after compression. The significant fiber/binder interface debonding after compression as shown in Fig. 9(a) indicated that the fiber/binder bonding was relatively weak. It should be pointed out that this relatively weak bonding was beneficial to strength of the sample, since the formation of the cracks between the binder and the fibers might consume the energy. It can be seen from Fig. 9(b) that there was no evident fiber/binder debonding in the matrix, indicating the strong fiber/binder bond strength. However, an extensive binder fracture and fiber breaks were observed after compression. In this case, the possible failure of the sample was attributed to the crack that grew in the fibers and at the bonding points. The fibers between the strong bonding points would fracture firstly, and then the binder between fibers fracture in a brittle way as the crack further propagated. The compression–rebound tests were carried out below the elastic limit and all of the compression rebound data of the FCFMF sintering at different temperatures were tested by setting the compressive ratio as 10%. Fig. 10(a) shows the variation of stress and the rebound-resilience of the FCFMF with sintering temperature. The stress of the samples sintering at 1100 °C, 1200 °C ,1300 °C and 1400 °C with compression ratio of 10% were 1.2 MPa, 1.5 MPa, 2.0 MPa and 2.8 MPa, respectively. The increase of the stress indicated that the strength between the mullite fiber and the binder increased with the sintering temperature. The rebound-resilience decreased from 98% to 90% as the sintering temperature increased from 1100 to 1400 °C, illustrating that the samples were not completely elastomer at the elastic stage. Fig. 10(b) shows the compression rebound curve of the FCFMF sintering at 1300 °C with 7 different cycles. With the stress increased, the strain and the rebound-resilience of the sample increased and decreased, respectively. The higher rebound-resilience of sample was mainly due to the complete resilience of the bended fibers during the loading process. On the contrary, the lower reboundresilience was accounted for the fracture of the bonding points and deflects in the material. Fig. 10(c) shows the overall view of sample sintering at 1300 °C after 7 cycles of compressions. It can be observed that the rebound-resilience was high during the first
six cycles. In this period, the fracture of the bonding points and deflects in the material take up only a small amount. However, at the seventh cycle of compression, the amount of the fracture of the bonding points and the fibers increased and the reboundresilience decreased to 70%. 4. Conclusion A new fibrous ceramic was prepared by infiltration method, in which the mullite fiber was used as the matrix and the SiO2– B2O3 as the high-temperature binder. The unique bird’s nest like structure (a 3-D skeleton structure) was the most important designed structure characteristic of this material in this work. The fibrous structure with fixed bonding points exhibited a high degree of elasticity and pseudoductility compared to the ceramic matrix composites reinforced with continuous fibers. The sintering temperature plays an important part in tailoring the properties of this fibrous ceramics. As the sintering temperature increased from 1100 to 1400 °C, the porosity decreased from 78.3% to 74.2%. Correspondingly, the compressive strength and the thermal conductivity increased from 1.3 to 3.2 MPa, and from 0.231 to 0.248 W/ m K, respectively. The samples showed a high rebound-resilience of above 90% after compression of 10% ratio. Such bird’s nest structure obtained from mullite fibers bonded by the SiO2–B2O3 at the crossing points could therefore be a potential high-temperature sealing material. Acknowledgement This study is supported by the National Natural Science Foundation of China (Project Nos. 51272171 and 51372164) for financial support. References [1] Pradan AK, Das D, Chattopadhyay R, Singh SN. Effect of 3D fiber orientation distribution on transverse air permeability of fibrous porous media. Powder Technol 2012;221:101–4. [2] Jackson GW, James DF. The permeability of fibrous porous media. Can J Chem Eng 1986;64:364–74. [3] Tahir MA, Tafreshi HV. Influence of fiber orientation on the transverse permeability of fibrous media. Phys Fluids 2009;21:083604. [4] Tsay RY, Weinbaum S. Viscous flow in a channel with periodic cross-bridging fibers: exact solutions a Brinkman approximation. J Fluid Mech 1991;226:125–48. [5] Gibson LJ, Ashby MF. Cellular Solids: Structure and Properties. Cambridge: Cambridge University Press; 2001. [6] Levine SR, Opila EJ, Haibig MC, Kiser JD, Singh M, Salem JA. Evaluation of ultrahigh temperature ceramics for aeropropulsion use. J Eur Ceram Soc 2002;22:2757–67.
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