Pre-drawing induced evolution of phase, microstructure and property in para-aramid fibres containing benzimidazole moiety

Abstract

Copoly(p-phenylene-benzimidazole-terephthalamide) (PBIA) fibre was spun by wet-spinning and drawn in a coagulating bath with different pre-drawing ratios (R). The evolution of phase, microstructure structure and conformation induced by pre-drawing were studied by 2D-wide angle X-ray diffraction (2D-WAXD), attenuated total reflection-Fourier transform infrared (ATR-FTIR) and scanning electron microscope (SEM). WAXD results indicate that mesomorphic/amorphous coexisting phases transform to crystal/mesomorphic/amorphous coexisting phases when R is higher than 75%. Even when R reaches an ultimate value of 100%, the content of crystalline is only 10.7% but the mesomorphic phase content is higher than 40%, which indicates that the 3D well defined crystalline structure is disturbed by the introduction of asymmetric benzimidazole units. The ATR-FTIR results indicate that the increase of R leads to conformational order, contributing to phase transition. The results of SEM show that PBIA fibres have a typical skin–core morphology. The sizes of micro-fibrils in the core part become more homogeneous with the increase of R. Moreover, with the increasing R from 0% to 100%, the tensile strength and initial modulus of the PBIA fibre significantly increase from 7.8 cN dt⁻¹ to 21.5 cN dt⁻¹ and from 238 cN dt⁻¹ to 884 cN dt⁻¹, respectively. Dynamic mechanical analysis (DMA) of PBIA fibres shows two transition peaks, and some amorphous phase is confined by the order phases.

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1. Introduction

Para-aramid fibres are considered a high-performance material owing to their superior mechanical and thermal properties, which make them useful for advanced technologies.^1,2 The best-known commercial aramid is poly(p-phenylene terephthalamide) (PPTA). In the late 1960s, the lyotropic liquid-crystalline behavior of rigid-rod PPTA solution was discovered by Du Pont, which led to the development of Kevlar® fibres.^3,4 The phase structure, microstructure and structure–property relation in Kevlar® fibres were then studied in detail by many researchers.^3,5–10 The measurements of Farris et al. indicated the values of crystallinity of Kevlar®119, Kevlar®29, Kevlar®49 and Kevlar®149 fibres were 75–77%.¹⁰ Kevlar®49 and Kevlar®149 fibres showed higher crystal orientation and lower paracrystalline parameters than those of Kevlar®119 and Kevlar®29, which led to the higher initial modulus of Kevlar®49 and Kevlar®149 fibres.¹⁰ The previous experimental findings indicated that the fibre structure included a mesophase between crystalline and amorphous fractions in many polymer fibres.^11,12 On the basis of this, the crystal, mesomorphic and amorphous fractions of Kevlar 49 fibres were firstly studied and calculated via on-line synchrotron WAXD techniques by Ran et al.⁵ Their results showed that crystal fraction was 50%, mesophase fraction was 20% and 30% was amorphous in the Kevlar 49 fibre.⁵ The chains of the mesophase have only 1D or 2D order and pack along the fibre axis, which may result in higher tenacity and higher elongation than the crystal fraction. However, the effect of the mesophase on the mechanical properties of Kevlar 49 fibre may be limited owing to the low total mass fraction of the mesophase. Thus, the authors considered that it is probably desirable to increase the fraction of the mesophase for the future development of the polymer fibres with higher tenacity.⁵

Another important para-aramid fibre is copoly(p-phenylene-benzimidazole-terephthalamide) (PBIA), introducing 5-(6)-amino-2-(4-aminobenzene)benzimidazole (PABZ) into the PPTA main chain.¹³ PBIA fibre was firstly developed by researchers from the former Soviet Union, which showed higher tensile strength and interfacial properties with epoxy resin than Kevlar®.^14–16 The introduction of PABZ leads to the random occurrence of head-to-head and head-to-tail isomerism, increasing the solubility and reducing the order of chain packing. To our best knowledge, reports on the phase structure and structure–property relationships of PBIA fibres are rare.^14,17–19 Perepelkin et al.^16,17 reported that the structure of PBIA fibre (Armos® in Russia) would transform to liquid-crystalline state during heat treatment. They thought that the less regular molecular chain structure in PBIA fibre would result in a higher contribution of stress holding and the maximal level of mechanical properties among all para-aramid fibres. However, no quantitative data have been provided yet about this kind of liquid-crystal form. The evolution of the structure and properties of PBIA fibres during heat treatment has been studied by Li et al.²⁰ The enhancement of crystallite size and spontaneous orientation led to the increase of tensile strength after the decomplexation of hydrogen chloride.²⁰ However, the analysis is on the basis of crystalline–amorphous two phase structure and the mesophase was not taken into account in the study.

Since the PBIA solution is an isotropic state, there is a possibility to control the phase structure during fibre-forming and thermal treatment process. Hot-drawing is a common method to control the phase structure and obtain high macromolecular orientation in polymer fibres. However, hot-drawing is difficult for rigid-rod para-aramid fibres with poor stretchability owing to strong hydrogen bonding interactions.^5,21 In this study, the PBIA fibres were spun by wet-spinning and pre-drawn with various drawing ratios in the coagulating bath, and were then heat treated without strain. The effects of pre-drawing on the phase structure and microstructure were studied in detail. The pre-drawing induced phase transition is firstly reported. Unlike Kevlar® fibres, a high fraction of mesomorphic phase is found in PBIA fibres, which would be in favour of the improvement of mechanical properties.

2. Experiment

2.1 Material

Dimethylacetamide (DMAC) was obtained from Shanghai Qunli Chemical Company and distilled over CaH₂ under reduced pressure before use. Paraphthaloyl chloride (TPC), p-phenylenediamine (PPD) and LiCl were obtained from ChengDu Kelong Chemical Co., Ltd., and used as received. 5-(6)-Amino-2-(4-aminobenzene)benzimidazole (PABZ) was obtained from Changzhou Sunlight Medical Raw Material Co. Ltd.

2.2 Synthesis of PBIA

LiCl was heated at 400 °C in a muffle furnace for 5 hours and then grinded into powder. 280 g of DMAC and 8.4 g of dry LiCl were added into a 100 ml three-necked bottle and then started to stir. When the LiCl was dissolved, 2.734 g (25.28 mmol) of PDA and 5.6668 g (25.28 mmol) of PABZ were added and dissolved. The solution was cooled in an ice/acetone bath to −5 °C in a nitrogen atmosphere, and then 10.26 g (50.56 mmol) of TPC was added accompanied by rapid stirring for 1 hour. Continuous stirring was carried out at 30 °C for 2 hours and a yellowish-brown organo-soluble PBIA solution was obtained. Fig. 1 shows the synthetic process and the chemical structure of the copolyamide.

Fig. 1 The synthetic process and the chemical structure of the copolyamide (m : n = 1 : 1).

2.3 Preparation and pre-drawing of as-spun PBIA fibres

The PBIA solutions were filtrated and degassed at reduced pressure for 24 h prior to use. The PBIA fibres were prepared by wet-spinning. The detailed spinning process can be seen in our previous paper.²⁰ The as-spun PBIA fibres were drawn at 0%, 25%, 50%, 75% and 100% in the coagulating bath, respectively, which is denoted as PBIA-0, PBIA-25, PBIA-50, PBIA-75 and PBIA-100, respectively.

It is worth noting that the as-spun PBIA fibre would be broken if the pre-drawing ratio is higher than 100%. Immediately after drawing, the fibres were washed with deionized water and then dried under even vacuum at 100 °C for 0.5 h. The pre-drawn PBIA fibres were further heated at 360 °C for 0.5 hour with fixed ends.

2.4 Characterization

The two-dimensional wide angle X-ray diffraction (2D-WAXD) patterns of the fibres were collected on Rigaku R-AXIS RAPID (MM007HF) diffractometer with VariMax HR optics at 40 kV and 30 mA. The background scattering was recorded and subtracted from the sample patterns and the d-spacing of each diffraction peak was calculated using the Bragg equation. The degree of molecular orientation can be calculated by integrating the corresponding intensity of azimuthal scans along the isolated and preferred crystalline plane. The degree of molecular orientation of the fibres is calculated based on the Hermans equation:^5,22where f₂ is the degree of molecular orientation along the fibre axis direction and ψ represents the angle between the fibre axis and c-axis crystal unit cell. The numerical values of the mean-square cosines in the above equation are determined by corrected intensity distribution I(ψ) diffracted from the crystalline plane by Gaussian fitting following the equation:

Once all the values of (cos² ψ) have been obtained, the molecular orientation in the direction of (hkl) can be evaluated.

The PBIA fibres were peeled and skin and core layer morphologies were observed on a JEOL JSM-5900LV scanning electron microscope (SEM).

Attenuated total reflection-Fourier transform infrared (ATR-FTIR) spectra of PBIA fibres were measured at a Nicolet Magna 650 spectroscope in the range 4000–600 cm⁻¹. The frequency scale was internally calibrated with a reference helium–neon laser to an accuracy of 0.2 cm⁻¹.

Small angle X-ray scattering (SAXS) experiments were performed using a NanoSTAR-U (BRUKER AXS INC.) with Cu Kα radiation (λ = 0.154 nm). The generator was operated at 40 kV and 650 μA. Two-dimensional SAXS patterns were obtained using a HI-STAR detector. The sample to detector distance was 1074 mm. The effective scattering vector q (, where 2θ is the scattering angle) at this distance ranges from 0.044 to 2.0 nm⁻¹.

Mechanical properties including the initial modulus, tensile strength and elongation at break of the PBIA fibres were measured on a KOL KD III-5000 with a strain rate of 25 mm min⁻¹. The fixture span was 215 mm. Each sample was measured ten times and average values of tensile strength, initial modulus and elongation were calculated.

Dynamic mechanical analysis (DMA) was performed on a TA Q800 instrument at a heating rate of 10 °C min⁻¹ under N₂ atmosphere in the range of 50 to 450 °C and the load frequency was 0.5 Hz.

3. Results and discussion

3.1 Pre-drawing induced the evolution of phase

The 2D-WAXD and 1D-WAXD curves of as-spun PBIA fibres with different pre-drawing ratios (R) are shown in Fig. S1 and S2 in ESI,† respectively. Five as-spun PBIA fibres show a broad peak at about 25°, indicating the amorphous structure. The as-spun fibres were heated at 380 °C for 0.5 hour with fixed-ends. In order to study the effect of pre-drawing on the phase structure and properties, no strain was applied to the fibres during heat treatment. The 2D-WAXD patterns of five PBIA fibres with different pre-drawing ratios (R) are shown in Fig. 2. The diffuse halo of the WAXD pattern at R = 0% exhibits a slight reinforcement on the equator, indicating an ordered intermolecular packing in the transverse fibre direction. There are weak and broad diffraction spots along the meridian, demonstrating a relatively poor order and orientation along the fibre direction. With the increase of R from 0% to 50%, the diffraction spots on the equator become stronger and narrower, which indicates the order and orientation is enhanced with the increase of R. Moreover, these WAXD patterns are similar with the results observed previously for nematic liquid crystalline.^23–26 The results of Levchenko et al. demonstrated that PBIA fibre showed smectic liquid crystalline structure in addition to some part of the nematic mesophase.¹⁸ It is in accordance with our observation. However, it is difficult to separate smectic liquid crystalline and nematic liquid crystalline from 2D-WAXD, so they are called a mesomorphic phase. The three-dimensional crystalline order is absent for PBIA-0, PBIA-25 and PBIA-50 fibres as evidenced by the absence of off-equatorial diffraction arcs.^23,27,28 It is an indicator of mesomorphic/amorphous coexisting phases for PBIA-0, PBIA-25 and PBIA-50 fibres. Actually, a mesomorphic form at room temperature is not specific to PBIA fibre. A number of aromatic polyesters, such as PET,^29–32 poly(ethylene-naphthalate),^33,34 polylactide^23,35 and poly-(esteramides),³⁶ have been reported to develop a mesomorphic order upon drawing or other treatments. As R reaches 75%, some weak off-equatorial diffraction arcs are observed, manifesting in the formation of a well-defined 3D crystalline structure for PBIA-75 fibre. The off-equatorial diffraction arcs become more obvious and the intensity is stronger for PBIA-100 fibre at R = 100%, indicating the enhancement of the degree of 3D well defined crystalline.

Fig. 2 The 2D-WAXD patterns of PBIA fibres: (a) R = 0%, (b) R = 25, (c) R = 50, (d) R = 75 and (e) R = 100.

In addition, the corresponding 1D-WAXD on the equatorial direction and peak fitting curves of PBIA fibres with different R are shown in Fig. 3. The main peak is observed at 2θ ≈ 20.4°, giving clear evidence of an interchain spacing d ≈ 0.43 nm. This peak becomes narrower and narrower with the increase of R, indicating the enhancement of interchain order induced by the pre-drawing. As the R is higher than 50%, another weak peak at 22.5° is obviously observed for PBIA-75 and PBIA-100 fibres, suggesting a well-defined ordered structure. Moreover, those two peaks in WAXD profile are similar to the crystalline peaks of (110) and (200) in PPTA fibres.^3,10 Thus, we consider that the copolyamide PBIA fibres show a similar crystalline structure with that of PPTA.²¹ However, PBIA-75 and PBIA-100 fibres do not show two separate diffraction spots in the 2D-WAXD pattern. This is different from that of Kevlar®,^3,10 which is owing to the overlap of the coexistence of crystalline and mesomorphic phases.

Fig. 3 The WAXD curves and peak fitting curves for five PBIA fibres: (a) PBIA-0, (b) PBIA-25, (c) PBIA-50, (d) PBIA-75 and (e) PBIA-100; (f) evolution with the PBIA amorphous, mesomorphic and crystalline fractions with the increase of R.

In order to further investigate phase contents and pre-drawing induced phase transition, the peaks were fitted assuming Gaussian type, as shown in Fig. 3(a)–(e). Peaks at about 20.4° and 22.5° are referred to as crystalline peaks. Peaks at about 21.5° and 28° are referred to as mesomorphic peaks. Evolution of the amorphous, mesomorphic and crystalline fractions with the increase of pre-drawing ratios (R) is shown in Fig. 3(f). When R is lower than 75%, the fraction of crystalline is 0% and only mesomorphic and amorphous phases coexist in PBIA-0, PBIA-25 and PBIA-50 fibres.

With the increase of R from 0% to 50%, the fraction of amorphous phase decreases from 80.4% to 64.5% while the fraction of mesomorphic phase gradually increases from 19.6% to 35.5%. Mesomorphic phase transition to crystalline phase takes place as R increases from 50% to 75%. The fractions of mesomorphic phase of PBIA-75 and PBIA-100 fibres are 40.6% and 40.3%, respectively. The fraction of crystalline phase increases from 4.3% to 10.7%. The results indicate that further increasing R from 75% to 100% is conducive to the enhancement of crystalline phase. Moreover, Ran and co-workers studied crystal, mesomorphic and amorphous phases fractions of poly(p-phenylene terephthalamide) fibre (Kevlar 49) via synchrotron 2D-WAXD.⁵ Their results indicated that about 20% of the fraction was mesomorphic phase, 50% was crystalline and 30% was amorphous phases in the Kevlar 49 fibre. Obviously, the copolyamide PBIA fibre shows lower crystalline fraction compared to Kevlar®. The reason is that the 3D well defined crystalline structure is disturbed by the introduction of asymmetric benzimidazole units.¹⁶ Moreover, they thought that the mesomorphic phase fraction would result in higher tenacity and higher elongation than the crystal fraction.⁵ Thus, the high mesomorphic phase fraction in PBIA fibre is in favour of the improvement of tensile strength.

In addition, the azimuthal intensity I(ϕ) profile of the reflection at 2θ = 20.4° is shown in Fig. 4(a). The peak width at half height of the I(ϕ) profile gradually reduces with the increase of R, indicating the improvement of orientation along the fibre direction.

Fig. 4 (a) The azimuthal intensity I(ϕ) profile of the reflection at 2θ = 20.5°; (b) the increase of f₂ with the increase of pre-drawing ratios.

The degree of molecular orientation (f₂) of five PBIA fibres is calculated based on the Hermans equation and is shown in Fig. 4(b). The degree of molecular orientation (f₂) increases from 0.71 to 0.87 with the increase of R from 0% to 100%. Even though the pre-drawing ratio is 0% in the coagulating bath, PBIA fibre shows an orientational packing along the fibre direction, which is owing to the formation of molecular orientation from the spinning jet and spontaneous orientation during the heat treatment.²⁰

3.2 Pre-drawing induced the evolution of microstructure

The PBIA fibres were peeled and skin and core layer morphologies were observed on a scanning electron microscope (SEM), as shown in Fig. 5. All of the PBIA fibres show an obvious skin–core structure, which is similar to that of Kevlar® fibres.^37–39 The skin morphology shows a smooth ribbon-like structure without apparent microfibrils. The thickness of the skin layer was measured roughly and was about 750 nm, 500 nm, 240 nm, 200 nm and 180 nm for PBIA-0, PBIA-25, PBIA-50, PBIA-75 and PBIA-100, respectively. The thickness of the skin layer decreases with the increase of pre-drawing ratios, since the low drawing ratios at the coagulating bath would lead to bad solidification and apparent skin–core structure. Higher magnification (160 000×) SEM images of the core layer of five PBIA fibres are shown in Fig. 5(A-2)–(E-2). The core morphology exhibits an apparent microfibril structure. The diameter of the microfibrils and microfibril numbers (n) in an area of 1 μm × 1 μm was measured from Fig. 5, as shown in Table 1. The diameter of the microfibrils is about tens of nanometers and the size distribution decreases with the increase of the pre-drawing ratio. The values of n in the area of 1 μm × 1 μm increase from 18 to 27 with the increase of pre-drawing ratios, which may result from the increase of crystallinity and orientation. In addition, SAXS measurement was used to characterize the micro-structure of PBIA fibres. The 2D-SAXS curves are shown in Fig. 6. The common feature seen in these patterns is that the SAXS intensity has a shape of a streak along the equator and there is no detectable scattering along the meridian direction, which is similar with that of Kevlar® fibre.⁶ It is evidence of the absence of a lamellar structure and a long period. The interpretation of the equatorial scattering in SAXS is complicated since it may contain two kinds of contributions, including the microfibrillar structure and voids morphology.⁵ Grubb et al. pointed out that the scattering objects in the Kevlar 49 were mainly associated with the microfibrillar structure, not the void morphology.⁴⁰ From the results of SEM, we are in complete agreement with them. The scattering at the low q region from long rods should follow the Guinier approximation,⁴¹ given bywhere r is the radius of gyration of the cross-section and G is a scaling constant. Fig. 7 shows Guinier plots of the scattered intensities along the layer line in the horizontal (equator) direction. The values of r are calculated and tabulated in Table 1. This data shows that the average radius of gyration for the PBAI fibres is in the range of 13–14 nm. The average radius of the microfibrils was almost unchanged from the results of SAXS while the distribution of microfibril size is more uniform and microfibril numbers increase with the increase of pre-drawing ratios. Thus, from the above results of WAXD and SAXS, schematic diagrams of pre-drawing induced phase transition and structural evolution pathways are proposed, as shown in Fig. 8. When the pre-drawing ratio (R) is lower than 75%, PBIA fibres show mesomorphic/amorphous coexisting phases. Further increasing R leads to phase transition from two coexisting phases to crystal/mesomorphic/amorphous coexisting phases.

Fig. 5 SEM images of peeled HT-PBIA fibres (general view) of (A-1) PBIA-0, (B-1) PBIA-25, (C-1) PBIA-50, (D-1) PBIA-75 and (E-1) PBIA-100; high magnification (160 000×) SEM images of the core layer of (A-2) PBIA-0, (B-2) PBIA-25, (C-2) PBIA-50, (D-2) PBIA-75 and (E-2) PBIA-100.

Table 1 Diameter of microfibril (D) observed from SEM, microfibril numbers (n) in the area of 1 μm × 1 μm observed from SEM and average radius of microfibril (r) calculated from SAXS

Samples	Diameter of microfibril from SEM (nm)	Microfibril numbers in 1 μm × 1 μm	Radius of microfibril from SAXS (nm)
PBIA-0	18–55	18	13.2
PBIA-25	18–45	22	13.5
PBIA-50	18–35	23	13.6
PBIA-75	15–28	25	13.0
PBIA-100	15–30	27	14.0

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Fig. 6 2D SAXS patterns of PBIA fibres at five different pre-drawing ratios: (a) 0%, (b) 25%, (c) 50%, (d) 75% and (e) 100%.

Fig. 7 Guinier plots of the scattered intensities along the layer line in the horizontal (equator) direction.

Fig. 8 Schematic diagrams of estimated phase transition and structure evolution for PBIA fibres.

3.3 The evolution of conformation and H-bonding interaction

ATR-FTIR is used to study the evolution of conformation and hydrogen bonding (H-bonding) interactions. The ATR-FTIR curves of five PBIA fibres are shown in Fig. 9. The absorption band at 1300–1320 cm⁻¹ could be assigned to Ph–N vibrations.⁴² The peak at 1300–1330 cm⁻¹ is assigned to Ph–N stretching vibrations and the wavenumber gradually shifts from 1308 cm⁻¹ to 1320 cm⁻¹ with the increase of R.

Fig. 9 (a) ATR-FTIR curves of five PBIA fibres at 1300–1450 cm⁻¹; (b) ATR-FTIR curves of five PBIA fibres at 600–900 cm⁻¹; (c) the wavenumber of C–H out of plane bending and I₁₄₀₅/I₁₄₂₄ with the increase of draw ratio; (d) linear fitting curve of I₁₄₀₅/I₁₄₂₄ and order phase content.

The absorption band at about 840 cm⁻¹ could be assigned to C–H out of plane bending of the benzene ring,^42,43 as shown in Fig. 9(b). As the R increases from 50% to 75%, C–H out of plane bending of benzene peak immediately shifts from 837 cm⁻¹ to 827 cm⁻¹, as shown in Fig. 9(c). Ishida and Huang studied the crystallization of a thermoplastic polyimide using infrared spectra.⁴⁴ Crystallization of the polyimide produces spectral shifts from 1380 to 1401 cm⁻¹ for C–N–C stretching vibration and from 828 to 822 cm⁻¹ for C–H out-of-plane bending of the benzene ring, respectively.⁴⁴ It is owing to the increase of conjugation upon crystallization.⁴⁴ Thus, in our copolyamide system, the shifts of Ph–N stretching vibration and C–H out of plane bending of the benzene ring are also explained by the enhancement of conjugation of the benzene ring and the heterocyclic ring, which leads to the increase of macromolecular order and the transformation of 3-D well defined crystalline induced by pre-drawing. In addition, from Fig. 9(a), 1424 cm⁻¹ and 1405 cm⁻¹ are assigned to the in-plane vibration of the benzimidazole ring and the C–C skeleton vibration of the poly(p-phenylene terephthalamide) (PPTA) units, respectively.^45,46

When R values are 0% and 25%, the intensity of 1424 cm⁻¹ (I₁₄₂₄) is higher than that of 1405 cm⁻¹ (I₁₄₀₅). When R is higher than 50%, I₁₄₂₄ starts to exceed I₁₄₀₅. Fig. 9(c) shows the values of I₁₄₀₅/I₁₄₂₄. This value gradually increases from 0.87 to 1.45 with the increase of R. Moreover, it is found that the variation of I₁₄₀₅/I₁₄₂₄ is similar with the ordered phases (crystalline and mesomorphic phases) content. Linear fitting between I₁₄₀₅/I₁₄₂₄ and ordered phases content is done, as shown in Fig. 9(d). The linear correlation value (R₂) is higher than 0.97. It indicates that I₁₄₀₅/I₁₄₂₄ in ATR-FTIR can reflect the order of the macromolecular chain. Thus, we can conclude that extended and conjugated macromolecular conformation is induced by pre-drawing, resulting in the transitions from mesomorphic phase to crystalline phase.

Moreover, in order to evaluate the evolution of H-bonding interactions of C=O, the FTIR spectra in the range of 1620–1700 cm⁻¹ belonging to V (C=O) of amide I are carefully investigated. Fig. 10 shows curve-resolved FTIR spectra in the range of 1580–1700 cm⁻¹ for PBIA-0, PBIA-25, PBIA-50, PBIA-75 and PBIA-100. Two peaks are separately identified for C=O band, among which band I (center: 1665 cm⁻¹) and band II (center: 1641 cm⁻¹) correspond to “free” C=O stretching hydrogen-bonded C=O and of amides, respectively.

Fig. 10 Curve-resolved FTIR spectra in the range of 1580–1700 cm⁻¹ for (a) PBIA-0, (b) PBIA-25, (c) PBIA-50, (d) PBIA-75 and (e) PBIA-100. Band I belongs to H-bonding “free” C=O; band II is H-bonding associated C=O; band III belongs to C–C of the benzene ring. (f) The variation of PWH and H-bonding C=O W (%) with the increase of pre-drawing ratio (R).

For a quantitative evaluation, the percentage of the hydrogen bonded C=O (W) can be calculated by the equation

where A_I and A_II are the relative areas of band I and band II, respectively. Hydrogen bonded C=O (W) and the peak width at half height (PWH) are shown in Fig. 10(f). The values of W for five fibres are in the range of 86–88%, which indicates that most of the C=O groups are involved in the formation of hydrogen bonds and the percentage of W stays unchanged.

Moreover, the peak width at half height (PWH) decreases from 46.7 cm⁻¹ to 39.6 cm⁻¹ with the increase of R. It shows that the intensity of hydrogen bonding interactions is more uniform with the increase of R, which is in accordance with the enhancement of the order of the macromolecular chain.

3.4 Mechanical properties and glass transition

The tensile strength and initial modulus of five PBIA fibres with different pre-drawing ratios are shown in Table 2. With the increasing R, the tensile strength and initial modulus of the PBIA fibres significantly increase from 7.8 cN dt⁻¹ to 21.5 cN dt⁻¹ and from 238 cN dt⁻¹ to 884 cN dt⁻¹, respectively, while the elongation at break decreases from 4.6% to 2.5%. Obviously, the enhancement of tensile strength and initial modulus with the increase of R results from the improvement of ordered phases content and orientation. When the R is higher than 75%, the tensile strength of PBIA-75 and PBIA-100 fibre is higher than 20 cN dt⁻¹ and the initial modulus is higher than 800 cN dt⁻¹, indicating outstanding mechanical properties. Actually, the tensile strength of industrial PBIA fibre is higher than 28 cN dt⁻¹ and the initial modulus is up to that of Kevlar® fibre.⁴⁷ The extremely high tensile strength may be owing to the high mesomorphic phase fraction in PBIA fibre.⁵

Table 2 The mechanical properties of PBIA fibres

Samples	Tensile strength (cN dt^−1)	Initial modulus (cN dt^−1)	Elongation at break (%)
PBIA-0	7.8 ± 0.3	238 ± 17	4.6 ± 0.3
PBIA-25	10.6 ± 0.6	324 ± 22	4.1 ± 0.3
PBIA-50	15.6 ± 0.6	652 ± 38	3.7 ± 0.2
PBIA-75	20.7 ± 0.4	838 ± 24	2.6 ± 0.1
PBIA-100	21.5 ± 0.5	884 ± 22	2.5 ± 0.1

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The glass-transition temperatures (T_gs) of PBIA fibres were characterized by dynamic mechanical analysis (DMA), as shown in Fig. 11. The storage modulus increases with the increase of pre-drawing ratio, which is in accordance with the results of mechanical properties. The storage modulus of PBIA fibres substantially decreases at 200–300 °C, representing glass transition behavior. In Fig. 11(b), an obvious α transition peak corresponding to segment motion in the amorphous domains is observed at about 250 °C for five PBIA fibres. The magnitude in tan δ gradually reduces from 0.1 to 0.06 with the increase of R. It indicates that segment motion is suppressed due to the improvement of macromolecular order. Moreover, it is interesting to find that there is another small peak at about 400 °C, indicating the appearance of a new segmental process. This phenomenon, previously reported for semicrystalline poly(ethylene terephthalate),^48,49 poly (trimethylene terephthalate),⁵⁰ poly(L-lactic acid)^51,52 and polyimide,⁵³ can be described by an additional α transition process, labelled α′, appearing at higher temperature. It is attributed to the restricted motions of the amorphous phase confined by the crystalline structures, which is called a rigid-amorphous phase⁵⁴ in flexible and semi-rigid polymers. Moreover, the comparable intensities of the α′ and α peaks (I_α′/I_α) increase with the increase of R. It implies that the content of confined amorphous phase is enhanced, resulting from the increase of order phases content.

Fig. 11 DMA curves of five PBIA fibres: (a) storage modulus; (b) tan δ.

4. Conclusions

Copoly(p-phenylene-benzimidazole-terephthalamide) (PBIA) fibre was spun by wet-spinning and drawn in a coagulating bath with different pre-drawing ratios (R). The evolution of phase, microstructure and mechanical properties induced by pre-drawing were studied. When R is lower than 75%, PBIA fibres show mesomorphic/amorphous coexisting phases. Further increasing R leads to phase transition from two coexisting phases to crystal/mesomorphic/amorphous coexisting phases. Even though R reaches an ultimate value of 100%, the crystalline content is only 10.7% but a high content of mesomorphic phase (40.3%) is found, which is different with Kevlar®. It indicates that the 3D well defined crystalline structure is disturbed by the introduction of asymmetric benzimidazole units. The increase of R leads to more extended and conjugated conformation of PBIA chains from the results of ATR-FTIR, which contributes to phase transition. The PBIA fibres show an obvious skin–core structure. The skin morphology is a smooth ribbon-like structure and the core morphology exhibits an apparent microfibril structure. Microfibril numbers per unit area observed from SEM increase with the increase of R. Moreover, with the increasing R, the tensile strength and initial modulus of PBIA fibres significantly increase from 7.8 cN dt⁻¹ to 21.5 cN dt⁻¹ and from 238 cN dt⁻¹ to 884 cN dt⁻¹, respectively, which results from the improvement of the mesomorphic and crystal phases content and orientation induced by pre-drawing. An α transition peak at about 250 °C and an α′ transition peak at about 400 °C are observed in the DMA curve, which implies that some mobile amorphous phase is confined by the ordered phases. The content of confined amorphous phase increases owing to the improvement of ordered phase.

Acknowledgements

This work was financially supported by the National Natural Science Foundation of China (Grant No. 51573105) and State Key Laboratory of Polymer Materials Engineering (Grant No. sklpme 2014-2-04). The authors acknowledge the Analytical & Testing Centre of Sichuan University, People's Republic of China for characterization.

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Footnote

† Electronic supplementary information (ESI) available. See DOI: 10.1039/c6ra10184d

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