Understanding the nano- and macromechanical behaviour, the failure and fatigue mechanisms of advanced and natural polymer fibres by Raman/IR microspectrometry^*

Synthetic and natural fibres are stored under controlled water partial pressure at room temperature. Some fibres are annealed at various temperatures or encapsulated with water (H₂O) or heavy water (D₂O) in long Pyrex tubes [29]. Fresh hand-spun, cocoon extracted and aged silkworm/spider fibres were studied before (the bave consists of two fibres embedded in a common sericin sheath) and after degumming (see [8, 16–18] for details). Regenerated silk was studied either in the gel or in the dried film state (thickness 10–50 μm). The regenerated silk was also used to prepare silk fibre-reinforced regenerated silk-matrix composites [19, 30].

The fibres (diameter between ∼5 (spider silk) and ∼30 (PA66) μm) are cut either in pieces a few tens of micron long for study under hydrostatic pressure (figure 4(d)) or a few tens of mm for tensile studies (figure 4(a)). Fibre tips are taped onto gold-coated glass plates, or mounted on paper to be assembled in a universal tester as previously described [7, 16].

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The mechanical tensile properties of the samples were obtained using (computer controlled) Universal Fibre Testers (UFT) previously described [16]. The gauge length was 30 mm. The Tester was set under the ×50 or ×100 long working distance objective (total magnification ×500 or ×1000) of high resolution Raman spectrometers (figure 4, [7, 8]). The orientation of the fibre/tester versus the laser polarization was chosen to obtain the stronger signal.

The Raman spectra were recorded at controlled temperature with different spectrometers equipped with lasers delivering (514.5 or 532 nm) green or (785 nm) red lines. A motorized XY table makes it possible to carry out mapping with a > 0.1 μm step. The laser power is controlled, typically <2 mW at the sample. A neon lamp is added to control the absolute wavenumber position [8].

A Lorentzian (L) shape for the 'crystalline/molecular' modes and a Gaussian (G) shape for the 'amorphous' signal are expected according to the Raman theory and the configuration distribution. Lab-developed software was used for the data extraction from the linescan/mapping hyperspectrum.

Measurements under hydrostatic pressure are taken through a diamond anvil cell, the pressure being calibrated with the fluorescence lines of chromium ions in ruby crystals [29]. The infrared microscope/ATR is an Equinox 55 Fourier transform Michelson interferometer with an ATR 'golden gate' input (Bruker, France). It is designed to examine the surface (a few microns in depth) of a material brought into 'perfect' contact with the ATR crystal. An Irscope II microscope with Cassegrain objectives makes it also possible to work in transmission for samples with thickness less than 30 μm. The spot diameter of ∼200 μm is reduced to 20 μm with a diaphragm, to separately analyse the core and the skin of a fibre.

X–H vibrators (X=O or N) are considered as isolated vibrators because of the large mass difference between X and H atoms. Furthermore, the very anharmonic character of the X–H bond potential due to the H-bonding makes the stretching mode wavenumber and bandwidth very sensitive to the local structure [17]. Reliable curves between the wavenumber and the X–H...Y distance have long been established [32, 33]. However, the intrinsic mode width (resulting from the structural disorder alone, without mechanical coupling) is obtained only in the case of an isotopic dilution [H]/[D] <10%: very light vibrators normally couple in a complex way (mechanical, electric, quantum) with the other phonons and only with isotopic dilution, by removing the essence of these couplings, it becomes possible to obtain a spectrum giving details on the vibrators' environment [11, 34].

As an example, figure 6(a) compares the infrared signature of the PA66 fibre in its hydrogenated and deuterated forms: some modes shift according to the D/H mass effect [5] but the prominent change is the decrease of the bandwidth. When the H/D mole ratio is less than 10%, mechanical couplings are almost suppressed and the bandwidth of the X–H vibrator reflects the static and dynamic disorder [11]. In the case of PA66 fibre the N–H stretching mode appears made of two bands, a narrow one assigned to N–H vibrators in 'crystalline' environment and a broader one assigned to those in amorphous matrix [5]. Similar features are observed for Raman N–H fingerprint. The main component is the narrowest Raman peak at 3303 cm⁻¹ ('Cr' bandwidth at half height ∼30 cm⁻¹, reduced to L = 20 cm⁻¹ in diluted H/D isotopic polyamide fibre). The IR 'Cr' counterpart peaks at 3304 cm⁻¹ (bandwidth ∼100 cm⁻¹, reduced to L ∼ 40 cm⁻¹ in diluted H/D isotopic polyamide fibre). These narrow peaks are well described with a Lorentzian function and can be assigned to 'crystalline' macromolecular chains. The N–H wavenumber is insensitive to isotopic dilution and the peak corresponds well to a N–H vibrator almost free from any hydrogen bonds (dN–H···O > 0.293 nm, see [33]). The much larger IR bandwidth results from the intrinsic larger broadness of IR modes due to the larger volume probed by the electric IR probe. In orthogonal polarization the intensity of this Lorentzian Raman component strongly decreases to the benefit of a broader Gaussian component (L, ∼80 cm⁻¹, figure 6(b)), ascribable to the 'amorphous' macromolecules. The higher wavenumber (G, 3310 cm⁻¹) indicates an absence of hydrogen bond, in agreement with a less compact structure. Two other weak bands are observed but it is difficult to know whether they both correspond to other types of N–H vibrators, or more probably to combinations, or result from a bad description of the line-shapes. The conservation of an amorphous component in V polarization confirms the isotropic character of the amorphous connections. Note that the re-hydrogenation of the deuterated fibres is slow and no remarkable difference between spectra carried out a few minutes after exiting the heavy water or 1 h afterwards is observed. The whole of these results confirms the existence of 'crystalline' and 'amorphous' or, rather, ordered and disordered conformation/structures of the macromolecular chains. The R/IR splitting being almost inexistent in the isotopically diluted state, it may be concluded that the local 'crystalline' and 'amorphous' structures in PA66 fibre are very similar and thus in agreement with a model of continuous angular disorder (para-crystal).

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Actually, because of the very large intensity of the low wavenumber Raman external/lattice modes Raman analysis should be preferred (figures 6(a) and (b)). Figure 6(b) compares the signatures of silk raw material (spectra directly recorded on the Bombyx mori silkworm fibre, dry gland and hand-spun fibre from the gland). The spectral signatures of the amide group are systematically present, as observed for the PA66 fibre [7]. Assignments can be found in references [7, 8]. Figure 6(c) shows the low-frequency collective modes, the narrow one for 'ordered' polyamide chains and the broader for amorphous one. A large Rayleigh wing is also present due to the density heterogeneity at the subnanometric scale. The polarized Raman signatures have been normalized versus the C–H massif. The analysis agrees with the results of x-ray diffraction highlighting amorphous and more ordered 'crystalline' conformations and validates the choice to apply this description to the collective mode around 100 cm⁻¹. The fact that the νN–H wavenumber of the amorphous 'phase' is a little higher is in agreement with the lower energy of the corresponding chain collective mode, as observed in figure 6. The disappearance of the 'crystalline' component signature with polarization is nearly complete. Therefore, the analysis of the PA66 chain modes at ∼100 (crystalline)/80 (amorphous) cm⁻¹ confirms the alignment of the crystalline macromolecular chains as well as the isotropic character of the amorphous phase. Similar features are observed for PET fibres, but a higher number of components can be distinguished in accordance with the higher crystallinity and lower symmetry of this compound [12].

For silks, this character is less marked and variable, according to the type of silk considered and the 'crystalline' collective mode measured at ∼80 cm⁻¹ instead 100 cm⁻¹ in PA66 fibre.

Raman shift as a probe of the (residual/applied) stress/strain was first developed for carbon fibre and C-fibre reinforced composite material because of the strong Raman signal of carbon and its resonance character. However, it is possible to easily measure the second harmonic spectrum (1350/1600 × 3 cm⁻¹ ∼ 4000/4800 cm⁻¹) that multiplies by three times the sensitivity of the methods [21, 35–37]. The stress-induced Raman shift does not only depend on the stress but also—and mainly—on the bond anharmonicity [3, 36]. Anharmonicity is great for aromatic and hydrogen-bonded bonds. Note the Raman shift originating from temperature change is much larger than the stress-induced shift, thus a careful control of the temperature—i.e., use of low laser illumination power—is mandatory [36–38].

We will now address the collective low wavenumber signature of 'amorphous' and 'crystalline' polymer chains to understand the mechanical tensile and compressive behaviour with the examples of PA66 (90 and 100 cm⁻¹), PET (60 and 75 cm⁻¹) and i-PP (50 and 110 cm⁻¹) [12–18]. For keratin fibres, it was not possible to extract the low wavenumber components and we had to use the Amide I (νC = O), the N–H stretching and the S–S bridge modes [14].

We will first consider the stronger narrow (crystalline) and broad (amorphous) component of PET fibre spectra (figure 7(a)). This shows an initially viscoplastic behaviour until 450–500 MPa (∼10–12% of lengthening), slightly varying according to the fibre producer. The amorphous bonds first and the crystalline ones later, are put into an increasingly rigid elastic tension (figure 7 [12]). The thresholds measured on the scale of the molecular bonds correspond to the transitions on the macroscopic stress–strain curve measured on single fibre between an initial softening, a pseudo-elastic regime and a nonlinear pre-rupture. This behaviour is very similar to that of a composite material with initial loading of the matrix alone (the amorphous matter), then that of the reinforcement (continuous ordered macromolecular chains) and the matrix simultaneously and, eventually, the degradation of the matrix followed by that of the reinforcement. The slight upward Raman shift measured for the crystalline signature indicates a Poisson's effect; the squeezing of the fibre diameter involved by the tensile stress compresses the crystalline islands, as sketched in the Oudet's model (figure 6(c) [39]). A very similar behaviour is measured for polypropylene but the downward shift starts at a low stress level in accordance with the high viscoelasticity of the material.

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A very different behaviour is observed for PA66 fibre (figure 8): the Raman shift measured for crystalline and amorphous signatures is rather similar. This is consistent with the Prevorsek's model [40], in which a macromolecular chain belongs both to amorphous and to 'crystalline' regions (figure 6(b)). We observe a plateau up to ∼500 MPa, i.e. up to the inflexion point of the macroscopic tensile stress–strain curve (figure 1(a)) and then a down slope, more or less pronounced as a function of the PA66 fibre producer. Changes are more significant for the amorphous component, as already observed for other fibres (figure 7). Actually, figure 8 compares data of two types of PA66 fibres, with very similar ultimate fracture strengths but produced by two different companies (their behaviour under fatigue is different). No significant difference is measured for the crystalline component but the slopes of the 'amorphous' chain signal are different above the ∼500 MPa threshold, indicating different behaviours in amorphous environment.

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Figure 9(a) compares the macroscopic tensile stress–strain curve and the Raman wavenumber shift of N–H (∼3277 cm⁻¹) and Amide I (∼1650 cm⁻¹) bands of Caucasian white hair fibre. The relative intensity of the S–S bond (∼510 cm⁻¹) versus strain is given in figure 9(a). A perfect relationship is observed between the macroscopic stress–strain behaviour and the nanoscopic Raman shift one: a linear elastic behaviour up to 4% of strain and then a plateau assigned to the α-helix opening transition allowing the possibility of a β-sheet arrangement. This transition is prepared by the breaking of S–S bonds (figure 9(a)).

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The evolution of the Raman shift induced by controlled strain is shown in figure 9(b) for fibres exhibiting the characteristic types presented in figure 5(a). The behaviour measured at the nanometric scale by Raman scattering—the probe is the chemical bond polarizability—fits well with the macroscopic stress–strain behaviour.

The recording of a series of spectra across the fibre diameter (figure 3) with high magnification microscope objectives (spot diameter 5–1 μm) makes it possible to analyse the texture anisotropy of the various fibres. An example is given with PA66 fibre (figure 10) but measurements have also been conducted for PBO and PET [12] fibres and for silks [16–18]. A core–skin effect is obvious [12] from the wavenumber shift across the fibre diameter. These variations indicate a weak tension of the crystalline chains and a strong compression of the amorphous bonds in the fibre core.

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The strong width reduction of the 'amorphous' macromolecular chain component in the fibre core undoubtedly indicates a much better organization, at the local scale, of the amorphous 'phase'. This characteristic can be related to the slower cooling of the fibre core.

Such anisotropy is consistent with a radial temperature gradient experimented by the fibre during the processing steps. A comparison of the data recorded for PA66 fibres thermally treated at various temperatures below and above the glass transition (T_g) temperature shows that the core/skin anisotropy disappears by annealing above T_g (figure 6), according to the above-proposed origin. The prominent role of a quenched skin layer in silicate glass to maximize the mechanical properties is well known in the glass industry. Different core/skin textures are observed as a function of fibre composition and producer.

Fatigue resistance studies, i.e. the solicitation of a material during a series of mechanical cycles between two load levels, one very low with respect to the ultimate strength (e.g. 5–10%) and the other not very far away from the ultimate value (e.g. 60–70%) are the most representative tests to design a material. The fatigue fracture mechanism is specific. This is illustrated by considering the habit of the fibre fracture, as shown in figure 11. A fracture performed by increasing the load up to the ultimate fracture gives a habit as visible in figure 3(a), a mirror-like surface in the area where the crack starts and then a more corrugated one. In contrast, a fatigue fracture shows an elongated tongue (figure 11(a)). The Raman analysis of fibres broken in fatigue along the line A–B–C highlighted a state of compressive stress of the amorphous phase close to the point of initiation of the rupture and its progressive decrease over 200–300 μm beyond this point [41]. As shown in figure 11(b) the upward shift measured in B–C points is located at the intersection between the y-axis and the slope of the 'amorphous' signal at a high level of strain. All occurs as if the rupture in fatigue resulted from the loss of viscoelasticity—and its characteristic plateau—at certain points, of the amorphous phase.

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Figure 8(c) compares the wavenumber shift of the low wavenumber lattice, Amide I and νN–H stretching modes of PA66 fibre in the diamond anvil cell (figure 4(d)). A threshold is observed at 4 GPa. Under hydrostatic pressure the coming together of the chain segments (reduction in νN–H wavenumber for both crystalline and amorphous chain conformation) can be clearly seen together with the existence of a threshold after which the geometry of the fibres is modified: the Amide I wavenumber, constant below 4 GPa, starts to shift and a slope change is observed for the low wavenumber mode. This confirms the higher sensitivity of N–H stretching and collective lattice mode with respect to the Amide I (or Amide III) modes, the bands used by other spectroscopy teams [42].