J. Mater. Sci. Technol., 2011, 27(7), 659-667.

Fatigue Behavior of HDPE Composite Reinforced with Silane Modified TiO2 C.X. Dong1,2) , S.J. Zhu1)† , Mineo Mizuno3) and Masami Hashimoto3) 1) Department of Intelligent Mechanical Engineering, Fukuoka Institute of Technology, Fukuoka, 811-0295, Japan 2) Nanomaterials Laboratory, Beihua University, Jilin, 132013, China 3) Japan Fine Ceramics Center, Nagoya, 456-8587, Japan [Manuscript received September 26, 2010, in revised form March 27, 2010]

The composite of high density polyethylene reinforced with silane-modified TiO2 particles (silane-TiO2 /HDPE) is a potential bone substitute biomaterial. The structure, bioactivity, and mechanical properties of silaneTiO2 /HDPE are analogous to those of natural bone, correspondingly. In order to investigate the effect of silane connection and saline solution on fatigue behaviors, flexural fatigue tests with this composite were carried out in both air and saline solution. Saline solution was found to have different effect on fatigue life. In saline solution, the fatigue life could be improved at stress levels lower than 30 MPa, while the fatigue life could be reduced at stress levels higher than 30 MPa. After analyzing the fracture morphologies, different failure mechanisms were proposed, and the important role of silane connection in the composite during the fatigue process was discussed. Silane connection cannot only support the loading stress but also hinder the failure process under loading effectively. For dry specimens, no interfacial failure between the filler and matrix was found. For wet specimens, it is inferred that the synergetic effect of saline solution and high concentrated stress at high stress level could easily destroy the silane connection, which accelerated the fracture process, whereas the synergetic effect of saline solution and silane connection at low stress level could promote the formation of more microcracks on sample surface, which hindered the final fracture. KEY WORDS: Fatigue; Composites; Polymers; Fracture; Interfaces

1. Introduction Bone biomaterials are mainly used in the therapy or replacement of damaged bones. Most of them are bioactive ceramics[1,2] and metal alloys[3] . However, because of the intrinsic properties of those biomaterials and the diversity of natural bones, there are still many limitations for their widespread application. For example, because of the high toughness and high bioactivity, glass-ceramic A-W has been used to replace bones subjected to high loads, such as vertebrae and intervertebral discs[4] . Nevertheless, glassceramic A-W still cannot replace damaged tibial and femur, because its fracture toughness is lower, and its † Corresponding author. Tel.: +81 92606 4265; Fax: +81 92606 0747; E-mail address: [email protected] (S.J. Zhu).

elastic modulus is higher than those of these natural long bones. Metal alloy biomaterials have high Young’s moduli, and can be used in the replacement of bones bearing heavy load. However, they can result in the resorption of surrounding bones, because of the stress shielding[5,6] . The composite with inorganic fillers and organic matrix is a choice for the design of bone biomaterials, because bones are mainly composed of inorganic apatite crystals and organic collagen fibrils. Based on this idea, a composite of hydroxyapatite particles with high-density polyethylene (HAPEX) was developed by Bonfield[7] , in the early 1980s. Because its mechanical properties are analogous to those of some natural bones, HAPEX has already been clinically used in the replacement of artificial middleear bones[8] . However, the fracture toughness and

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elastic modulus of HAPEX are lower than those of heavy load bearing bones. Because TiO2 has higher elastic modulus than the hydroxyapatite, Hashimoto et al.[9] developed a biomaterial TiO2 /high density polyethylene (TiO2 /HDPE), which shows high bioactivity. The bending strength and Young0 s modulus of TiO2 /HDPE were found to be in the range of 28 to 54 MPa and 1.4 to 7.6 GPa, respectively, depending on the content of TiO2 . Nevertheless, the weak mechanical adhesion at interfaces was still a main reason for the mechanical failure. In order to improve the interfacial adhesion, a silane-coupling agent was used to modify the surface of TiO2 particles[10] . The obtained composite is silane-TiO2 /HDPE, in which the surfaces of TiO2 are connected tightly with HDPE by chemical bonds of Ti-O-Si. The chemical connection between TiO2 and HDPE greatly improves the mechanical properties. For example, the bending yield strength and Young’s modulus could be increased from 49 to 65 MPa and 7.5 to 10 GPa, respectively, depending on hot-pressing pressure[10] . Compared with HAPEX, the microstructure of silane-TiO2 /HDPE is more similar to that of natural bone, with silane connection in silane-TiO2 /HDPE corresponding to the acid phosphoprotein bond, which connects non-calcified collagen fibril bands to adjacent apatite crystals. As a potential biomaterial used in clinical applications, it is necessary to study its fatigue properties, because of the presence of cyclic loading in the potential clinical applications[11,12] , and the fatigue fracture, a main problem associated with implant failure[13] . Meaningful problems in the course of developing new biomaterials or polymer composites are: whether the similarity in structure between silane-TiO2 /HDPE and natural bone yields similar fatigue properties; and what is the effect of silane connection on mechanical properties. Although silane-coupling agents are widely used to modify the surface of inorganic fillers in polymer composites[14,15] , to the best knowledge of the authors, no article has been previously published on the fatigue behavior of the polymer composites with silane connection between inorganic fillers and organic matrix. Only a few articles investigate the fatigue properties of polymer composites without silane connection[16–18] . For example, Kultural et al.[17] reported the fatigue behaviors of calcium carbonate filled polypropylene under high frequency loading with all the samples experiencing thermal fatigue failure. Bonfield et al.[18] studied the fatigue behaviors of apatite-reinforced polyethylene composite (HAPEXTM ) in saline solution, and compared the fatigue damage of different load modes on the composite. In this study, the bending fatigue behavior of silane-TiO2 /HDPE was investigated both dry and in saline solution. In comparison with dry samples, the surrounding saline solution improves the fatigue life at stress below 30 MPa, whereas it decreases fatigue

life at stress higher than 30 MPa. The silane connection between TiO2 and HDPE was hypothesized to play an important role in the improvement of fatigue life in saline solution. It is inferred that because of the presence of silane connections, the formation of more numerous microcracks on the specimen surface resulted in a longer fatigue life in saline solution at low load stress, which is in good agreement with results on natural bones[19] . 2. Experimental 2.1 Materials Silane-TiO2 /HDPE composite was fabricated in Japan Fine Ceramics Center (Nagoya, Japan). The ratio of TiO2 to HDPE is 40 vol.%, with anatase mean particle size of 535 nm (Ishihara Sangyo Co. Ltd., Osaka, Japan). The TiO2 powder was treated with the silane-coupling agent of γ-MPS (γ[(methacryloxy)propyl]trimethoxysilane) (Shin-Etsu Chemical Co. Ltd., Tokyo, Japan), and mixed with high density polyethylene (Japan Polyolefins Co. Ltd., Tokyo, Japan), which has a number average molecular weight (Mn) of 1.21×104 , weight average molecular weight (Mw) of 7.67×104 , and z-average molecular weight (Mz) of 47.6 ×104 . The detailed manufacturing process was described in literature [10]. First, TiO2 powders were modified with the silane-coupling agent as follows: 1.1 g of γMPS, 1.6 g of ethanol, and 0.2 g of ion-exchanged distilled water were stirred with a magnetic stirrer for 10 min. The solution containing the silane-coupling agent was added to 110 g of TiO2 powder, and mixed in the shaker at 25◦ C for 1 h. The mixtures were then dried and heated at 130◦ C for 5 min. Second, the manufacturing of silane-TiO2 /HDPE by kneading and compression molding was carried out as follows: HDPE was dried at 80◦ C for 8 h and then kneaded at 210◦ C in a batch kneader PBV 0.3 (Irie Shokai, Ltd., Tokyo, Japan). Modified TiO2 particles with silane-coupling agent were added slowly into the melted HDPE with kneading at 210◦ C. After adding modified TiO2 , silane-TiO2 /HDPE compound was kneaded at 25 r/min for 30 min. The obtained compounds were molded at 230◦ C for 1 h, and then hot-pressed at 5 MPa. From the Fourier transform infrared spectroscopy (FTIR) analysis, it was checked that the surfaces of the TiO2 particles are connected with HDPE through formation of Ti-O-Si bonds. 2.2 Mechanical testing Four-point bending test was employed. 16 specimens were cut to a desired shape, and polished to a size of 40 mm×4 mm×3 mm. A fatigue testing machine, Model MMT-101NV-10 (Shimadzu Co. Ltd., Japan), was used to apply load with an outer span 30 mm and an inner span 10 mm. The surfaces measuring 40 mm×4 mm were

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70

Stress / MPa

60 50 40 30 20 10

In air In saline solution

103

104

105

106

107

Cycles to failure / Cycle

Fig. 1 S-N curves (maximum stress as a function of number of cycles to failure) for four-point bending fatigues under different conditions

directly loaded downward. The experiments were performed on dry specimens or on specimens in 0.9% NaCl saline solution at 25 ◦ C. The room temperature was 25±2◦ C. The fatigue tests were carried out by sinusoidal loading at a frequency of 10 Hz under load control with stress ratio of 0.1. In saline solution, the specimens were immersed in saline solution for at least half an hour prior to testing. 2.3 Morphology observation A scanning electron microscope (SEM) Hitachi E1010 and an optical microscope Keyence VHX-600 were used for observation of fracture surface on the same specimens after mechanical testing. For SEM observation, the specimens were cleaned ultrasonically and dried, and the fracture surfaces were sprayed with Au to ensure clear images. The optical pictures were obtained combining two adjacent photos to show the whole image of a sample. 3. Results and Discussion 3.1 Fatigue properties It was known[10] that after hot-pressing at 5 MPa, bending strength and Young0 s modulus of silaneTiO2 /HDPE of dry specimen were 65 MPa and 10 GPa, respectively. The S-N curves (maximum stress as a function of number of cycles to failure) of flexural fatigue for dry and wet specimens (Fig. 1) show that the fatigue life is prolonged through a decrease of maximum stress. The curve obtained on the wet specimens has a bilinear behavior, while the curve obtained on the dry specimens is linear. At stress or higher than 30 MPa, the fatigue life for wet specimen is much smaller than that for dry specimen, whereas at stress lower than 30 MPa, the fatigue life of wet specimen is longer. Therefore the effect of saline solution on fatigue life depends on stress magnitude. It should be noted that the difference in fatigue

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lives between dry and wet specimens at low stress level is not resulted from hysteretic heating, which can be understood from two respects. On the one hand, the movement of molecular chains during fatigue testing in silane-TiO2 /HDPE is difficult due to the structure with intensive silane connection between the filler and matrix[20] . Necking and plastic deformation were not seen on the tested specimens especially at low stress magnitude, which is consistent with the structure. In polymer PI/SiO2 without chemical connection between filler and matrix, the hysteretic heating can also be neglected at the same frequency of 10 Hz[21] . On the other hand, even if some local segments of molecular chains moved against each other, the heat rise could not affect the temperature of bulk specimens, which plays an important role in the movement of molecular chains. It is reported that in neat polymer without inorganic filler like chlorinated polyvinyl chloride, the hysteretic heating at 10 Hz and room temperature caused a heat rise only at the crack tip rather than in the bulk specimen, which can result in a lower FCP (fatigue crack propagation) rate[22] . The lower FCP rate is conductive to the prolongation of fatigue life of dry specimens, while the liquid medium cools down the heat rise at the crack tip and has an opposite effect on fatigue life. Therefore, the prolongation of fatigue life at low stress level in liquid medium is not related to the hysteretic heat at all. The real reasons for the prolongation of fatigue life of wet specimens at low stress level will be discussed in later passages. 3.2 Fractography 3.2.1 Fracture morphologies of dry samples Figure 2 shows the optical morphology of the fracture surface of dry samples tested at 30 and 60 MPa. The fracture processes started at the top of the image and the fracture patterns differ. At 30 MPa, the starting of the crack is evident and most of the surface is flat with shearing edges at the bottom of the fracture surface. At 60 MPa, it is difficult to distinguish the starting of the crack and all surface is covered with dimples of different sizes. Frequently shearing edges at the ends of fracture surfaces are present. The fracture of dry silane-TiO2 /HDPE turns from brittle to ductile with the increase of the applied stress from 30 MPa to 60 MPa. In order to investigate the failure mechanisms, SEM observations at various magnifications were carried out. The SEM morphologies of the fracture surface of a dry sample tested at 30 MPa are shown in Fig. 3. In Fig. 3(a), crack starting is focused to a small area. The crack propagation zone around the crack starting is flat without obvious voids on it. Figure 3(b) shows the intact structure of the zone just around the crack source. No fibrils of HDPE, no exposed TiO2 particles, and no voids can be seen on it. However, on the flat zone far away from the edge of

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Fig. 2 Optical fracture surface morphologies of dry samples at different mechanical stresses: (a) 30 MPa, (b) 60 MPa. To show the complete morphology from the fracture beginning to the end, two pictures are combined in the same sequence as they occur

Fig. 3 SEM fracture surface morphologies of dry specimen tested at 30 MPa: (a) failure surface with crack starting at low magnification, (b) microstructure of adjacent zone around crack starting at high magnification, (c) microstructure of flat zone far away from the crack starting

fracture beginning in Fig. 3(c), HDPE fibrils and voids of different sizes can be seen, revealing the obvious ductile deformation of HDPE. At 60 MPa, the morphologies are totally different from those obtained at 30 MPa. The fracture surface of dry sample tested at 60 MPa is covered with dimples of different sizes, as shown in Fig. 4(a). At higher magnification, the details of the dimples differ. Some areas are full of fibrils and voids, while other areas exhibit almost intact structure, as shown in Fig. 4(b). The two different structures at higher magnification are shown in Fig. 4(c) and Fig. 4(d), respectively. 3.2.2 Fracture morphologies of samples tested in saline solution The optical morphologies of fracture surfaces for

the samples tested in saline solution at 30 and 60 MPa are exhibited in Fig. 5. From the top down, the fracture surfaces consist of three parts: black crack starting zone, flat area, and shearing zone. The zone adjacent to the crack source is flat, while the shearing zone is full of dull fibrous fracture faces and shearing edges. Although the flat perimeter and rough region in the fracture surfaces are similar, the degrees of coarseness of the upper edge differ, enhanced with the increasing mechanical stress. For the samples tested at 60 MPa, large dimples can be clearly seen at the two sides of the black crack initiation. In addition, the shapes of the black crack starting differ. It is a long line on the upper edge of the sample tested at 30 MPa, while the one obtained at 60 MPa is suborbicular. With an increased mechanical stress from 30 to 60 MPa, the

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Fig. 4 SEM fracture surface morphologies of dry specimen tested at 60 MPa: (a) fracture surface at low magnification, (b) fracture surface at high magnification, (c) microstructure of flat zone, (d) microstructure of voids

Fig. 5 Optical fracture surface morphologies of specimens tested in saline solution at different mechanical stresses: (a) 30 MPa, (b) 60 MPa. To show the complete morphology from the fracture beginning to the end, two pictures are combined in the same sequence as they occur

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Fig. 6 SEM fracture surface morphologies of sample tested in saline solution at 30 MPa: (a) fracture surface with crack starting at low magnification, (b) microstructure of adjacent zone around crack starting at high magnification, (c) microstructure of flat zone far away from the crack starting

Fig. 7 SEM fracture surface morphologies of sample tested at 60 MPa in saline solution: (a) fracture surface with crack starting at low magnification, (b) microstructure of the crack starting at high magnification, (c) microstructure of the zone between crack starting and flat zone, (d) microstructure of flat area far away from crack starting

black zone becomes focused. The variation of fracture surface with mechanical stress in saline solution differs magnificently from that of dry samples. Figure 6 shows the SEM fracture morphology in saline solution for the sample tested at 30 MPa. Figure 6(a) shows a long line on the edge of sample, corresponding to the black line in Fig. 5(a). Figure 6(b) and 6(c) are the morphologies of the adjacent zone near the crack starting, and other flat zone far away from the crack starting, respectively. These morphologies are very similar to those obtained from dry sample at 30 MPa, showing that the effect of saline solution on fracture process is very small at 30 MPa. Figure 7 shows the SEM fracture morphology

of the sample tested at 60 MPa in saline solution. The crack origin at high magnification is shown in Fig. 7(b). Clearly, large blocks of the composite with exposed TiO2 particles on them can be seen at the crack origin. Deformed HDPE fibrils cannot be seen. The morphology between crack origin and flat zone is exhibited in Fig. 7(c), where the interface between the two zones is very distinct, because of their different microstructures. In the adjacent flat zone, no block of HDPE can be found. The morphology of the flat zone at higher magnification is exhibited in Fig. 7(d). Deformed HDPE fibrils and microvoids can be observed. In conclusion, the morphologies of fracture surfaces are different both between different regions

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within the same specimens and between the same regions of dry and wet specimens at high loading stress. These variations in fracture morphologies are mainly resulted from the different stress concentrations caused by loading within the specimens. Although the loading stress is definite, the stress concentrations within different parts of the specimens are different. In four-point bending test, when the load is applied downward, the lower surface of the specimens supports the biggest stress, while the top surface bears the lowest stress. During fatigue process, stress concentrations destroy the weak regions on the lower surface and propagate up to the upper. Different regions on the fracture surface experience different stresses, which resulted in the different fracture morphologies within the specimens. Similarly, the variations of fracture morphologies in the same regions between dry and wet specimens at high loading stress level were also resulted from the different stress concentrations essentially, which can be understood after the description of fracture mechanisms in the following sections. 3.3 Failure mechanisms 3.3.1 Dry samples The morphology of dry samples shows that the fracture process differs as function of stress. At 30 MPa, the fracture is due to brittle failure, which includes microcrack initiation, propagation, and fast failure process. Microcracks originated from the microvoids or micropores in the matrix of HDPE on sample surface. At this low stress level, although it was difficult for the concentrated stress to destroy the interfacial adhesion, HDPE experienced plastic deformation around the microvoids to relieve stress concentrations, which promoted the formation of microcracks. At 60 MPa, the sample undergoes void coalescence. First, the high stress breaks HDPE molecule chains around some of the microvoids, which are manifested by the small intact areas on fracture surface. Then, these microvoids become bigger voids and big voids coalesce to final failure of sample. The coalescence process was realized mainly by plastic deformation and destruction of HDPE. 3.3.2 Samples tested in saline solution Each sample was placed into saline solution for 42 h before testing. At different time intervals, the sample was picked out of saline solution, dried by absorbing paper, and weighed with an electronic precision balance. Results showed that the weight of the sample remained exactly the same, proving the waterproofing quality of silane-TiO2 /HDPE. Namely, without mechanical stress, the surrounding saline solution can not influence the specimens. The different fatigue behavior in saline solution should be due to the synergetic effect of saline solution on mechanical stress.

[23]

Guild and Bonfield have shown by computer modeling that there are two regions of stress concentrations: one is the pole of the particle, and the other is the interface. These stress concentrations resulted in failures at weak interfaces of many materials, such as HA/PE[7] and untreated TiO2 /HDPE. It is still possible for the interfaces between silanatedTiO2 particles and HDPE matrix to fail at high stress. In order to illustrate the variation of interface between TiO2 and HDPE under loading, the sketch map of microcrack initiation on sample surface of silaneTiO2 /HDPE was plotted (Fig. 8). The original intact structure without loading is shown in Fig. 8(a). In addition to the chemical adhesion due to the formation of Ti-O-Si bonds, there is physico-mechanical adhesion between TiO2 and HDPE. Under loading, stress concentrations at the interface may destroy the mechanical adhesion first, and strengthen the chemical bonds of Ti-O-Si with the reorientation and deformation of HDPE (Fig. 8(b)). When they are not enough to destroy the Ti-O-Si bond, the stress concentrations can only be relaxed by the further deformation and destruction of HDPE. When the stress concentrations are high enough, they can break the chemical bonds Ti-O-Si. The two ways of relaxing concentrated stress coexist, and complement each other. It is well known that TiO2 is hydrophilic, with the ability to form Ti-OH bond when water molecules are present. It is the formation of Ti-OH groups that induced the growth of apatite, a sign of bioactivity, on the surface of bone-biomaterials in simulated body fluid[24] . Even when Ti-O-Si has been broken by stress concentrations in dry specimens, the formation rates of Ti-OH and Si-OH (Fig. 8(c)) are relatively small. In comparison with the formation rates of TiOH and Si-OH, the deformation and destruction of HDPE happen more easily and faster. Therefore destruction of HDPE becomes the main way to relax stress concentrations. After the breakage of Ti-O-Si bonds in wet specimens, the formation rate of Ti-OH and Si-OH is higher, and subsequently, water can be adsorbed by the hydroxyl groups by forming hydrogen bonds. The steric hindrance of the introduced -OH groups and water molecules can in turn increase the stress concentrations, and break additional Ti-OSi bonds (Fig. 8(d)). Therefore, interfacial failure of silane-TiO2 /HDPE augments increasingly faster in saline solution. After the microcracks at the interfaces grow to a certain extent, a main crack forms, resulting in the final quickly generated fracture. The fracture due to interfacial failure occurs faster than the failure due to plastic deformation and destruction of HDPE. Therefore, it is concluded that it is the synergetic effect of saline solution on high load stress that accelerated the fatigue failure in saline solution at high stress level. At stress of lower than 30 MPa, the stress is not high enough to destroy the Ti-O-Si bonds on the surfaces of both dry and wet specimens. However,

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Fig. 8 Surface crack formation illustration of silane-TiO2 /HDPE: (a) without mechanical stress, (b) under mechanical stress and intact Ti-O-Si bonds, (c) Ti-O-Si bonds broken in air, (d) Ti-O-Si bonds broken in saline solution

because of the hydrophilicity of TiO2 surface, the hydrostatic force of saline solution could promote the destruction of the physico-mechanical adhesion between TiO2 and HDPE on the sample outer surface, and subsequently enhance the formation of more microcracks. The formation of more microcracks could relieve the destruction of the inner portion of the sample, and result in a longer fatigue life. Although the most outer layer on the sample surface experienced the destruction of chemical bonds Ti-O-Si, the effect of interfacial failure was negligible. Therefore, the fracture morphologies of sample tested in saline solution at 30 MPa (Fig. 6) are almost the same as those of the dry sample at the same stress level (Fig. 3). It should be noted that when the stress levels were 30 MPa or lower (Fig. 1), both the dry and wet samples experienced the same failure mechanism. That is, in the range of low stress level, all samples experienced the microcrack fracture mechanism originating within the HDPE matrix rather than at the interface between TiO2 and HDPE. At 60 MPa, the fracture mechanism in saline solution differs from that of dry specimens. The exposed TiO2 particles and the blocks of HDPE at the crack initiation site show that the interfacial failure between TiO2 and HDPE played an important role in the fracture process. On the sample surface, large concentrated stress not only could destroy weak mechanical adherence but also could break Ti-O-Si quickly. Once one Ti-O-Si bond between fillers and matrix was broken in the sample, additional Ti-O-Si bonds were

easily subsequently destroyed, because of the steric hindrance effect of water molecule. Simultaneously, the microvoids at the interface easily grew first to microcracks, then to a main crack, and finally resulted in the complete failure. The fracture due to the interfacial failure was faster than the plastic fracture process of HDPE. The dimples beside the crack origin in Fig. 5(b), and the existence of voids in Fig. 7(d) exhibits the competition between the void coalescence process and microcrack fracture process from interfacial failure. From the failure processes, it is known that silane connection between TiO2 and HDPE played different roles in fatigue behaviors in saline solution. At high stress, the effect of silane connection was negligible, and the hydrophilicity difference between the fillers and matrix was still a main factor affecting the fatigue behaviors, while at low applied stress, silane connection could not only act as an obstacle of destruction as in the dry specimens, but also further improve fatigue life, because of its synergetic effect with saline solution. The variations of the fracture mechanisms between dry and wet specimens at high loading stress level resulted in the differences in the fracture morphologies. At the beginning of loading, the synergetic effect of saline solution with silane connection focuses the stress mainly onto the interfaces between TiO2 filler and HDPE matrix and destroys the interfaces in wet specimen, which is different from the direct destruction of HDPE matrix in dry specimen. The altered

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location of destruction and improved stress concentrations at interfaces changed the morphologies of fracture surfaces in wet specimen in high stress level. 3.4 Enhancement of fatigue life in saline solution at low stress levels From the above discussion, it is known that the variation process of silane chains during loading is almost the same as that of its counterpart in real bone. In natural bone, it was hypothesized that because the non-mineralized collagen fibrils would tend to straighten during fatigue, the acid phosphoprotein bonds between the non-calcified collagen fibril bands and the adjacent apatite crystals tended to weaken and eventually break[19] . That means the ultra-structure of bone is highly resistant to lesion propagation, which brought about the result that microcracks were increased in number more than in length during loading[19] . The formation of more microcracks was proposed to be able to improve the fatigue life of natural bones[25] . In this study, the effect of silane connection in silane-TiO2 /HDPE is similar to that of the acid phosphoprotein bond in real bone. That is, the ultra-structure of silane-TiO2 /HDPE is also highly resistant to lesion propagation. However, because the synergetic effect of load stress and saline solution differs at different stress levels, the hindrance effect of silane connection on fatigue destruction behaved differently in saline solution. At low stress level, with the help of hydrostatic power of saline solution, more microcracks could be formed on the sample surface, resulting in the longer fatigue lives than those of dry samples at the same stress level. To sum up, during fatigue loading, silane chains between TiO2 and HDPE can not only support the loading stress but also hinder the failure process from interfaces, which has been proved in the study of creep behaviors of this material[26] . 4. Conclusion Through the flexural fatigue tests of silaneTiO2 /HDPE, the function of silane connection in supporting load was studied. The research showed that silane connection is very effective in supporting load, and can result in different fatigue behaviors under different test conditions. It is very effective in improving the interfacial compatibility of dry specimens, and can avoid the interfacial failure during flexural fatigue. At low stress, it can further improve fatigue life of wet specimens compared with dry ones, while at high stress level, when the concentrated stress breaks the chemical bond of silane connection, the easy interfacial failure between TiO2 and HDPE is still the main reason for fatigue fracture of wet specimens. The introduction of silane connection between filler and matrix was proved to be efficient in resisting the lesion propagation during fatigue process at low stress range.

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