Fracture energy, interfacial measurement

Yet, for systems A and C, the measured fracture energies remain low compared with the critical fracture energy of the bulk aluminum 10 J Moreover, we do not observe islands of passivation material on the A1 fracture surface and, inversely, we do not observe A1 on debonded surfaces of the passivation films. This suggests that the loss of interfacial adhesion is close to a brittle fracture process despite the influence of plasticity of the A1 substrate and crack blunting at the interface. This sort of brittle mode of interfacial failure, including plastic flow in a ductile material (the substrate), has been observed or discussed for a sapphire/Au interface. ... 

Figure 7.2. The double cantilever beam test for measuring interfacial fracture energy. Two welded polymer bars are driven apart by a razor blade of width 6 and the length of the crack ahead of the blade is measured.

The situation is more delicate when the two materials have different moduli. In this case, if the beams are of identical thickness the failure will no longer be purely mode I. In these circumstances the crack will deviate from the interface into the material with the lower deformation resistance, leading to additional energy dissipation. In these circumstances the measured values of the interfacial fracture energy will be larger than Gic- This problem can be overcome by using an asymmetrical test, in which the thicknesses of the two beams are unequal. At a particular ratio of thicknesses the measured fracture energy will be a minimum and this may be taken as the true value of G c. 

Figure 7.3. Fracture energies of interfaces between polystyrene and poly(/)-methyl styrene) of various relative molecular masses (A, PS 1 250 000 and PpMS 570 000 o, PS 310 000 and PpMS 570 000 and O, PS 862 000 and PpMS 157 000) as functions of their interfacial widths, measured by neutron reflectivity. After Schnell et al. (1998).

Figure 7.9. Interfacial reinforcement of a polystyrene/poly(vinyl pyridine) interface by a high relative molecular mass deuterated styrene-vinyl pyridine block copolymer, with degrees of polymerisation of each block 800 and 870, respectively. Circles (right-hand axis) show the measured interfacial fracture energy as a function of the areal chain density of the block copolymer 2, whereas crosses show the fraction of dPS found on the polystyrene side of the interface after fiacture. The discontinuity in the curves at 2 = 0.03 nm is believed to reflect a transition from failure by chain scission to failure by crazing. After Kramer et al. (1994).

Hartness [100], working with XAS and HMS fibers in a PEEK matrix showed similar behavior. The similarities between PEEK and PP are probably greater than the differences in their crystalline structure. Beaumont [101] has shown that with HMS (treated) fiber, there is almost no pull-out, whereas with HMU (untreated) fiber, extensive debonding and pull-out take place. The pull-out lengths can be measured and using an analytical technique outlined by Phillips [102], values can be obtained for the nylon/fiber interfacial bond strength and fracture energies. 

In some uncomplicated examples, it has been possible to analyse the results of adhesion tests to obtain numerical values associated with interfacial forces. An example of this was the work of E. H. Andrews and A. J. Kinloch, who measured the adhesion of SBR (styrene - butadiene rubber) to different polymeric substrates over a range of temperatures and test rates. Three types of tests including a 90° peel test were used. The results were analysed to evaluate the fracture energy per unit area G. (For the peel test, G was P, the peel energy.)... 

This work shows how the type and magnitude of interfacial forces can be deduced from the results of adhesion tests. For this simple polymer, the adhesive fracture energy is made up of an intrinsic term Go related to the interfacial forces, and a term related to the viscoelastic energy losses during testing. The viscoelastic loss can be expressed as an additive term (see Eqns. 1 and 3) or, for this system, as a multiplicative factor, Eqn. 4. Although for SBR at all normal rates and temperatures of test the viscoelastic term xj/ is much greater than the interfacial term W, W exerts a profound influence on the measured adhesion because of the multiplicative relationship. 

Kalton, A. E, Howard, S. J., Janczak-Rusch, J., Clyne, T. W. (1998) Measurement of interfacial fracture energy by single fibre push-out testing and its application to the titanium-silicon carbide system , Acta Materialia, 46(9) 3175-89. 

L. L. Shaw, B. Barber, E.H. Jordan, M. Gell, Measurement of the interfacial fracture energy of thermal barrier coatings, Scripta Materialia, Vol. 39, No. 10, pp. 1427 1434, 1998... 

Barber, A.H. Cohen, S.R. Kenig, S. Wagner, H.D. Interfacial fracture energy measurements for multi-waUed carbon nanotubes pulled from a polymer matrix. Compos. Sci. Technol. 2004, 64, 2283-2289. 

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