A significant number of modelling studies have shown the effects of stiffness gradients upon strain energy release rate, G, in bodies containing a stiffness gradient for cracks which initiate both perpendicular and parallel to the gradient direction [1]. Crack deflection has also been considered where cracks initiate perpendicular to the gradient direction [2] and also observed experimentally, even in shallow stiffness gradients [3]. The majority of these studies however ignore two important aspects of fracture: (i) the effect of spatially changing fracture resistance, R, and (ii) the effect of an R-curve, which is an inevitable feature of many of the composite materials used in FGMs. It has been shown that crack growth resistance behaviour may have a significant effect upon G for a crack which propagates parallel to the gradient direction from a brittle to a ductile composition or reverse [4, 5]. In this case no crack deflection occurs. When a crack initiates perpendicular to the gradient direction both mode I and II stress intensity factors, KI and KII, result in a material containing a stiffness gradient. This leads to crack deflection. In many cases, it would appear advantageous for the crack to propagate towards the tougher material where, depending upon component shape, it may halt prior to ultimate failure. Crack bridging, however may cause significant reductions in mode I crack-tip stress intensity factors, and presumably has an effect upon mode II ones also. It is the purpose of this work to assist in elucidating the effects of R-curve behaviour, resulting from ductile phase bridges, upon crack deflection of a crack which initiates perpendicular to the gradient. EXPERIMENTAL WORK An experiment was devised which modelled the mixed-mode crack-tip loading in a gradient material. A piece of polyurethane foam was placed at one end of a mould into which an alumina slip was poured, penetrating the foam. This was then dried and sintered, during which time the foam pyrolised leaving a section at one end of the ceramic with interconnecting porosity. This was then placed into another mould and polyester resin poured in, penetrating the porous section of the ceramic. The sample was then machined into a single-edge tensile sample with a sharp-tipped notch placed perpendicular to the loading direction as shown in Fig. 1. The sample was ~40 mm wide by ~10 mm thick with the composite 'interface' region being ~10 mm high. Notches were placed at varying positions across the gradient with a notation of 0 closest the polymer and 1closest the ceramic, as shown in Figure 2. The composite region was measured to contain ~45vol% ceramic and was ~11% porous due to incomplete penetration of the polyester resin. Figure 1: Schematic representation of single-edged tensile sample The sample was then loaded to fracture. Figure 2 shows the nature of crack propagation and deflection, observed after failure. It was found that the fracture energy/area uncracked region was in the order of 650900 J/m2, depending upon initial crack position, compared to 55 J/m2 for a similar bimaterial interface. Crack deflection occurred due to mixed mode loading resulting from the stiffness gradient along the sample and across the crack plane. The crack deflection angle, q, was invariably positive and in the direction of the more compliant polymer. As intrinsic toughness was constant across the interface, this deflection could not be due to changing intrinsic fracture resistance. The deflection angle as a function of crack propagation in the initial crack plane was determined from a digitised post-failure image and results are shown in Figure 3. P P Introduced notch Polymer Ceramic Composite
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