Various formulae22, 23, 24 have been proposed to link the fracture toughness to the indenter type, crack geometry, load, and the properties of the materials under test. For indentation with a Berkovich tip, the mode I fracture toughness was obtained using Eq. (2), 24:
SEM micrographs showing radial cracks around the imprints of bulk specimens after sintering and indentation at different load (a) 1200C and 5 N, (b) 1100C and 500 mN, (c) 1000C and 2 N, and (d) 900C and 20 N. Note that the trenches in (c) were machined by FIB.
crack x plane 9 20
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For a given load (500 mN), much longer cracks are found in a 1200C sintered film (b) with 15% porosity, compared to 1100C sintered bulk (a) with even much larger porosity (29%). Note that both images are at the same magnification and the indent sizes (triangles) are very close.
The results of this study raise two general issues regarding the application of indentation to measure the toughness of porous bulk materials and porous films on substrates. The first relates to the appropriateness of the analysis used for indentation of dense materials when applied to porous bulk materials. The present results show that the indentation toughness results are in good agreement with the SEVNB results for the porous bulk materials. This is at first unexpected because the permanent deformation mechanisms in dense and porous materials are different and some contribution from disruption of the particle networks and crushing densification might be expected under the indenter. Since this does not conserve volume, it would result in a smaller plastic zone for a given indenter depth and lower residual stresses on unloading. This would tend to lead to an overestimate of the toughness using Eq. (2) because a larger load would be required to give the same crack length. However, the FIB/SEM analysis under the indent shows no clear evidence of consolidation under the indenter. Furthermore, if anything the indentation toughness is consistently slightly smaller than the SEVNB toughness. Therefore, it is concluded that the permanent deformation in these porous materials does not involve significant consolidation and behaves macroscopically as if it were a dense effective material.
The second issue relates to the use of indentation to measure toughness for porous films on dense substrates. Despite the similar crack patterns exhibited in both LSCF bulk and films, it is worth noting that the toughness equation was originally developed for bulk materials and therefore its applicability for assessing film fracture toughness remains questionable25 even though it has been used to assess dense thin film fracture toughness by a number of researchers.15, 38 In the present work, we have found that indentation did not produce cracks in films sintered at 1100C and below. This cannot be attributed to a low toughness. Furthermore, much longer cracks in films sintered at 1200C compared with bulk specimen of similar porosity. It is therefore more likely that the standard analysis of the indentation method is not valid for these supported films. It might be expected that in general the toughness of a film on a substrate is inherently different from that of a bulk specimen of equal porosity.
Pseudo-ductility in hybrid composites is examined with a PA 6,6 veil at the hybrid interface. A PA 6,6 veil at the hybrid interface increases the fibre area without impacting the modulus of the hybrid composite. Previous authors found that veil interlaying reduced crack propagation during dynamic loading showing an improvement in stiffness degradation [15] and compression after impact behaviour [16] of composites, by mitigating the damage area. Rubber particulates are shown to reduce the storage modulus [17], which is detrimental to load bearing structures. The randomly oriented polyamide (PA 6,6) nanofibre veil interlayering is used to address the shortcomings of delamination through changing failure mechanics by improving Mode I+II and improving interfacial toughness between fibres. The flexural modulus of the hybrid composites is highly dependent on the stacking sequence of carbon fibre and glass fibre layering [11,18,19]. A significant negative hybrid effect for flexural strength occurs when carbon fibre is under compression due to premature buckling [11].
Compression failure caused by out of plane buckling is the most common type of premature failure in the interlayer hybridisation of carbon/glass fibres [8,9,14,20]. In unidirectional fibres, kink band formation after 99% structural loading tends to occur with a z to x direction tilt of the fibre direction [21]. Typically, the failure area of fibre kinking is localised compared with longitudinal cracking causing buckling [22]. The kink banding represents a gradual failure with the formation of the kink band being the final failure. Gradual failure creates nonlinear failure mechanics leading to higher strain making the composite pseudo-ductile [21]. Previously, nanofibres have been used in compression after impact applications reducing failure area, in carbon fibre composites [23,24,25,26,27] and glass fibre [28]. This research will aim to create a more stable failure with veil interlayering to promote more ductile composite, by localising the failure area. The addition of a veil nanofibre improves the fracture toughness of the CF/GF interface, to examine the effect on pseudo-ductility and further investigate the mechanism of failure [29].
Czel et al. [29] determined load bearing compression/tension plies need to have a higher strength than the applied loading after initial fibre/matrix failure. Additionally, the interlaminate fracture toughness (ILFT) at the interface is required to be stronger than the critical release energy. Increasing the ILFT through veil interlayering allows a thin fibrous layer to surround the failure area localizing crack propagation. Shekar et al. [30] showed that glass fibre reinforced plastic (GFRP) improves the Mode I ILFT behaviour with PA 6,6 nanofibre leading to a 68% increase in flexural strength. Tsotsis et al. [16] showed that PA 6 veil toughening in the interlaminar region, causes Mode II shear failure rather than Mode I crack opening. PA 6,6 veil by Beckermann et al. [31] showed an increase in CFRP fracture toughness for Mode I crack opening by 150% and Mode II by 50% with 4.5 g/m2 veils. The addition of nanofibres increases the energy required for crack propagation, as the matrix can tolerate higher stress concentrations. Mohammadi et al. [32] showed matrix cracking is reduced by 92% for CFRP, using electrospun PA 6,6 as an interlayer toughener. Blythe et al. [15] showed that PA 6,6 veil nanofibre localised crack propagation during flexural cyclic loading led to reduced stiffness degradation but was insufficient to prevent high loading amplitude catastrophic failure.
Figure 5, shows a typical stress strain curve for CF and CFT, in which the flexural strength maximum is to be within the standard deviation, suggesting that veil has minimal impact on the flexural characteristics for carbon fibre. The GF sample shown in Figure 6, has a wider crack head than the CF, CFT and GFT suggesting a more gradual failure.
The pseudo-ductile response ends when the glass fibre ply forms a kink band. The stress strain curve of the CTGGTC suggests that gradual matrix cracking occurred until the kink band formed, greatly reducing load bearing. This gradual formation is pseudo-ductile as the linear section of the stress strain curve changes after minor matrix cracking, leading to a failure strain exceeding 0.03 mm/mm and initial failure yielding at 0.013 mm/mm.
In Figure 17B, the CTGGTC sample shows a change in failure mechanism with the inclusion of veil instances of buckling that changed to fibre kinking. Veil toughened composites reduced the out-of-plane buckling caused by compression, resulting in shear kinking being the dominant failure mechanism. In veil toughened samples, there is a lower flexural strength; however, past the yield point, there is a pseudo-ductile response.
Figure 18A,B both show layer-by-layer failure, with CGCG featuring interlaminar failure and CTGTCTG crack propagation, though of the fibre rather than the matrix. When undergoing compression, the carbon fibre failed from delamination of the top layer, causing crack development through the glass to the second carbon fibre layer. Tension failure of the CGCG sample was observed as the elastic buckling resulted in flexural failure to the second carbon ply. The cracks propagated around the cross stitching towards the second carbon later, resulting in ply splitting and tension side delamination. On average for CGCG, elastic buckling resulted in a sharp drop in flexural strength, only increasing at higher strain when the final glass fibre ply began tension loading. Compared to the toughened CTGTCTG, the inter-fibre failure resulted in less fibre failure and more matrix fibre debonding.
As the veil enhances the matrix toughness, the CTGTCTG in CTGTCTG in Figure 18B shows a significant failure on a micro-level between fibres rather than within the matrix. The failure mechanism indicates that the crack propagation has been deflected from the matrix into the fibres. The delamination on Figure 18B shows that the veil traps crack propagation as marked D, this shear failure shows matrix fibre debonding at the interface. The tension side delamination was observed with both CGCG and CTGTCTG where at higher strain the glass fibre over elongates.
Figure 19A CCGG depicts elastic buckling, this was induced when the carbon fibre undergoes compression in the presence of crimping fabric. The crimping of the fabric in the E zone, indicates waviness of the fabric induces buckling. The carbon fibres are shown to undergo compression failure in the alternating sample, where the elongation of the crack head resulted in the failure overlapping with the glass undergoing tension. The overlap of the glass and carbon fibre peaks then resulted in pseudo-ductile behaviour, whereby the material behaves nonlinearly. Additionally, shear failure of the second carbon fibre ply resulted in additional kink band formation and out-of-plane deformation, resulting in extensive fibre failure offset from the centre. 2ff7e9595c
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