3.1. Composite Fibre Volume Fraction
Typical optical micrographs of polished cross sections of the flax and sisal composites are shown in
Figure 1. Using the methods described above, the average fibre volume fraction for the flax and sisal composites were found to be 0.31 and 0.4, respectively.
It can be observed that, unlike conventional composite fibers, the cross-sectional shapes of natural fibers vary widely. Thomason has discussed the influence of the noncircularity of the natural fibres in terms of the challenges of obtaining accurate values from single fibre micromechanical tests; such as, single fibre tensile testing and single fibre adhesion tests for determining interfacial shear strength [
16,
17]. The images in
Figure 1 also highlight the complex internal structure of these natural fibres. The ‘technical’ fibres are actually composite structures consisting of an assembly of many elementary fibres, which have a polygonal cross-section, allowing them to fit closely together. Due to the low cost requirements of many of the applications, where natural fibres are being considered as reinforcement, these fibre bundles are the most prevalent morphology observed. In the case of flax, the technical fibres are located around the stem of the plant and tend to contain between 10 and 40 elementary fibres. Sisal technical fibres originate from within the large leaves of the plant, and contain between 50 and 200 elementary fibres, which tend to be slightly smaller than the flax elementary fibres. Each elementary fibre contains multiple concentric cell walls, with a void in the middle, known as the lumen. It is clear from the micrographs, that the structure of these fibres reveals many potential weak interfaces. Indeed,
Figure 1 reveals that the flax fibre has been damaged during the sample preparation and has failed at a number of internal interfaces. The images in
Figure 1 highlight the need for the determination of the transverse properties of such fibres, as it is clearly unreasonable to expect that the complex structures can be anything other than highly anisotropic.
3.2. Composite and Fibre Modulus
A selection of the DMA storage modulus results for fibre and epoxy are plotted as a function of temperature in
Figure 2. A large reduction in stiffness, due to the epoxy matrix glass transition temperature (T
g), is clearly visible between 70 °C–90 °C. It can be seen over the whole temperature range, that the composites of Young’s modulus decrease as the off-axis loading angle increases. It is also interesting to note that the transverse stiffness of some of the composites loaded at higher off-axis angles is lower than the epoxy resin at temperatures below the matrix T
g. This is an indication of the likely probability of low transverse stiffness of the flax and sisal fibre. The longitudinal fibre modulus E
f1 can be obtained at any temperature from the composite longitudinal moduli (E
C0) and the epoxy matrix modulus (E
m) using the well-known Voigt model (also known as the linear rule of mixtures) using the equation below:
V
f and V
m are the volume fractions of the fibre and the matrix, respectively. It should be noted that the properties of the composites and epoxy matrix change by orders of magnitude in the region of T
g. These large changes can give unexpected results from the micromechanical equations used in this work; and so, we have restricted the temperatures for the micromechanical analysis to the range of −50 °C to +50 °C. The results for the longitudinal fibre moduli of flax and sisal fibres obtained from Equation (1) are compared with the epoxy matrix modulus in
Figure 3a. It can be seen that the axial stiffness of both of these fibres is considerably greater than the stiffness of the epoxy matrix; and hence, as observed in the introduction, it can be expected that these fibres will give a considerable reinforcement effect in the composite longitudinal direction. It is further noted that the values obtained for the room temperature fibre moduli of flax and sisal are in good agreement with the values previously reported from single fibre tensile testing.
The transverse fibre modulus E
f2 can be obtained at any temperature from the composite longitudinal moduli (E
c90) and the epoxy matrix modulus (E
m) using the well-known Reuss model (also known as the inverse rule of mixtures) using the equation below:
Results for the transverse fibre moduli of flax and sisal fibres that were obtained from Equation (2) are compared with the epoxy matrix modulus over the temperature range −50 °C to +50 °C in
Figure 3b. It can immediately be observed that the values of E
f2 for both fibres are significantly lower than the modulus of the epoxy matrix across the whole temperature range. The transverse modulus of the sisal fibres is only approximately 50% of that of the matrix. The flax fibre transverse modulus is even lower with values of only 30% of that of the epoxy matrix.
Figure 3b clearly illustrates a major weakness in the application of these natural fibres as a composite reinforcement, in that they provide no reinforcement of the polymer matrix in the transverse direction. On the contrary, these fibres have an anti-reinforcement effect in the transverse direction and result in a composite with a lower transverse modulus than the matrix polymer alone.
The mechanical anisotropy of these natural fibres is further visualised in
Figure 4, which presents the ratio of E
f1/E
f2 for the flax and sisal fibres, and compares them with some other reinforcement fibres [
22]. It can be seen that both sisal and flax are highly anisotropic in their mechanical performance. The flax fibres have a modulus ratio from 55–80 across the temperature range studied, and can be seen to have a level of anisotropy comparable with pitch carbon fibre. The sisal fibres have a considerably lower level of anisotropy, with a modulus ratio of approximately 17. This is lower than most carbon and aramid fibres, but still highly anisotropic in comparison with glass fibres.
The values obtained for the off-axis of Young’s moduli (E
Cθ) from
Figure 1 and
Figure 2 can also be used to obtain values for the composite shear modulus G
C12 , which can then be used to obtain a value for the fibre shear modulus G
f12. The mechanics concerning the coordinate system transformations can be applied, leading to the well-known relationship between off-axis modulus and the principal properties of a unidirectional composite ply [
12]:
Hence, by measuring the composite modulus as a function of the loading angle, it is possible to obtain a value for G
C12 by using the curve fitting Equation (3) with the experimental results taken from
Figure 2 at various temperatures. An example is shown in
Figure 5 for the flax and sisal composite moduli obtained at 25 °C.
The fibre shear modulus G
f12 can then be obtained at any temperature from the composite moduli (G
C12) and the epoxy matrix shear modulus (G
m), also using the well-known Reuss model:
Results for the shear moduli of flax and sisal fibres obtained from Equation (2) are compared with the epoxy matrix shear modulus in
Figure 6. It can be seen that the shear modulus of the sisal fibres differs a little from the shear modulus of the epoxy polymer and is actually lower, at lower temperatures. The shear modulus of the flax fibres (1.0 GPa–2.1 GPa) is significantly higher than both the sisal fibres and the epoxy matrix. However, as a reference, it can be noted that the shear modulus of E-glass fibres is more than an order of magnitude greater, at around 30 GPa.
3.3. Composite and Fibre Thermal Expansion
Figure 7 illustrates the TMA measured thermal strain of the epoxy matrix and the natural fibre composites with different fibre orientation. The data exhibits some typical elements of TMA analysis of polymer materials. When below the glass transition temperature (T
g), the thermal strain is low but as the temperature increases towards T
g, the thermal strain increases, and above T
g, the thermal strain is significantly high. The temperature region of this change of slope is indicative of the polymer (matrix) T
g which, in both cases, can be seen to be in the range of 70 °C to 90 °C (similar to the values obtained by DMA). It can further be noted from these figures, that when loading at 90°, the thermal strain of the composite is approximately the same as the resin thermal strain. At fibre orientation angles, the composite thermal strain is very low and a change in slope at T
g becomes quite indistinct. Interestingly, when the flax composites are at 0° fibre orientation, the thermal strain turns negative, above the matrix T
g temperature. From the results, it can be concluded that the composite thermal strain decreases as the fibre orientation decreases from 90° to 0°.
In a similar manner to the analysis of the composite and fibre modulus, the coefficient of thermal expansion of the composites determined by TMA can be used to establish the coefficient of thermal expansion of the fibre using micromechanical models. The longitudinal coefficient of thermal expansion of the fibre (α
f1) can be obtained using a rearrangement of the well-established Schapery [
23] equation:
where α
C1 and α
m are the longitudinal coefficient of thermal expansion (CTE) of the composite and the linear CTE of matrix [
17]. There is not, as yet, a universally applicable micromechanical model or equation for predicting the transverse CTE of composites; and so, we have compared the results from the published methods of both Chamberlain [
24] and Chamis [
25]. The rearrangement of the equation from Chamis gives the following expression for the transverse CTE (α
f2) of the fibres:
Whereas, the equation from Chamberlain results in the following expression:
where α
C2 is the transverse CTE of the composite and F is the fibre packing factor with a value of 0.785 for the square packing [
24].
The longitudinal and transverse CTEs of the flax and sisal fibres obtained from this analysis are compared with the epoxy matrix CTE in
Figure 8 respectively. It can be seen that the longitudinal CTEs for both natural fibres are not only an order of magnitude smaller than their transverse CTEs, but also negative in both cases. In other words, flax and sisal fibres shrink along their length when heated up. This is in agreement with the results of Cichocki and Thomason who found that jute fibre also has a negative longitudinal CTE [
12].
The two equations for the fibre transverse CTE give similar values in the case of both flax and sisal, although the Chamberlain equation results in systematically higher values for α
f2 in comparison to the Chamis equation. The sisal fibres appear to have a slightly smaller transverse CTE compared to flax fibres over the −50 °C to +50 °C temperature range. In both cases, the fibre transverse CTE is of the same order of magnitude as the epoxy matrix CTE, which is probably a reflection of the thermal response of the amorphous polymeric components of these fibres. An overall summary of the temperature dependence of the various thermomechanical parameters obtained for sisal and flax fibres in this study is shown in
Table 1. It should be noted that the values of α
f2 in
Table 1 are an average of the values obtained using Equations (6) and (7).
3.4. Implications for Natural Fibre Performance as a Composite Reinforcement
The results presented in
Table 1 illustrate the high levels of thermomechanical anisotropy present in the natural fibres. Although this anisotropy does not, in itself, present insurmountable challenges to the further application of such fibres as reinforcements for engineering composites, the implications do bear some further discussion.
Table 2 compares values of the thermomechanical properties [
22] of some other well-known reinforcement fibres at 25 °C, with those obtained in this work for flax and sisal.
It can be observed from the values present in
Table 2 that other reinforcements, such as carbon fibre and aramid fibre, used in engineering composite applications also have high levels of thermomechanical anisotropy. However, the challenge for those who wish to use natural fibres is not so much the anisotropy itself; rather it is the very low values of fibre transverse and shear modulus. These values for natural fibre are much lower than most polymer matrices; consequently, natural fibres perform poorly as reinforcement for the composite material in any off-axis loading scenario. In fact, natural fibres are not seriously considered by many as a replacement for carbon or aramid fibres, but they are regularly promoted as a possible replacement for glass fibres. In this case, the applications which are being considered are most often low performance, low cost, moulded composites; where a large proportion of the reinforcement fibres experience off-axis loading. In this situation, the isotropic nature of glass fibres still results in a significant level of reinforcement; whereas, the natural fibres will perform poorly.
There are many articles to be found in the literature which strongly identify the fibre–matrix interface as a challenge for the further development of natural fibre composites. Many authors have seen this as an issue of chemical compatibility and adhesion. However, it has recently been shown that the residual compressive radial stresses at the fibre–matrix interface caused by the fibre–matrix mismatch in CTE can account for a large proportion of the interfacial stress transfer capability in many composites [
15,
26,
27,
28,
29,
30]. Thomason has published an analysis for jute–PP composites which indicates that this hypothesis can be used to explain the very low levels of apparent interfacial shear strength in these composites [
13,
15]. Clearly, the data presented above where the value of α
f2 for flax and sisal is very close to that of the polymer matrix, implies a similar lack of any interfacial radial compressive stress to the interfacial stress transfer capability in composites using these fibres. In fact, if, as appears to be the case at some points in
Figure 8, α
f2 is greater than the CTE of the polymer matrix, this could result in an interface under radial tension, which could be very weak.
In further discussion of the interface in fibre reinforced composites, it is worth observing that the accepted stress transfer mechanism is shear along the interface. This has mainly been considered in terms of its effects on the interface and surrounding matrix, but there has been little attention paid to the effects of this interfacial shear stress on the reinforcing fibres. However, considering the very low values obtained for the natural fibre shear modulus in comparison to glass fibres, one must assume that the levels of shear strain on the fibre side of the interface must be at least an order of magnitude higher for natural fibres than for glass fibres. It is interesting to ask what effects such high levels of shear deformation might have on the internal structure of these natural fibres, how that might manifest itself in the apparent interfacial stress transfer capability, and the overall performance of the natural fibre composite. Charlet et al. [
31] have published the results of a unique study of the properties on the internal interphase region between elementary fibres within technical flax fibres. They gave a value for the average shear strength of these interphases of 2.9 MPa and a value for the shear modulus of 0.02 MPa. It seems clear that there is a much higher probability of failure of these internal interphase regions due to high shear strains, rather than the failure of the fibre–matrix interface. Such a result would go a long way to explaining why the chemical modification of the interface has such little effect on improving the composite performance with natural fibres in comparison to glass fibres. It also suggests that the way to move forward to improve natural fibre composite performance should be more focused on the internal structure of the technical fibres.