Properties,morphology and structure of BPDA/PPD/ODA polyimide fibres

Materials

p-Phenylenediamine was obtained from Adamas Reagent Co., Ltd. 4,4′-Oxydianiline was obtained from Bengbu Zuguang Fine-Chemical Co., Ltd. 3,3′,4,4′-Biphenyltetracarboxylic dianhydride (s-BPDA) was obtained from Shijiazhuang Haili Fine-Chemical Co., Ltd. N, N′-Dimethylacetamide (DMAc) was obtained from Tianji Tiantai Reagent Co., Ltd and was distilled over CaH₂ under reduced pressure before use.

Preparation of PI fibres

All poly(amic acid) (PAA) spinning solutions were synthesised in DMAc, and the reaction process is shown in Scheme 1. First, distilled DMAc, PPD and ODA were charged into a three-necked flask equipped with a mechanical stirrer. After the diamine was dissolved, calculated amounts of dianhydride were added gradually; the solid concentration and inherent viscosity were controlled to 15 wt-% and 1·80–3·0 dL g⁻¹ respectively.

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Scheme 1

Synthetic route of PI fibre (m/n = 95∶5, 90∶10, 85∶15 and 80∶20)

The PAA solutions were filtered and degassed before use. The PAA fibres were spun by dry jet wet spinning. The PAA dopes were extruded through a spinneret (34 holes with 0·12 mm in diameter) into an air gap and then into a coagulation bath. The solidified filament entered into the second washing bath and then clustered at the third spool. The fibres were dried and then imidised and drawn in a furnace over 400°C.

Characterisation

Inherent viscosities η_inh were obtained on 0·5 g dL⁻¹ concentration with an Ubbelohde viscometer in DMAc at 30°C. The mechanical properties were measured with a Donghua XQ-1 instrument; the gauge length and crosshead speed for the tensile test were 20 mm and 10 mm min⁻¹ respectively. Thermogravimetric analyses (TGAs) were conducted with a PerkinElmer TGA7 at a heating rate of 10°C min⁻¹ in a flowing nitrogen. The linear coefficients of thermal expansion (CTEs) of fibres were performed on a TMA/SDTA841e/190 (Mettler Toledo) with a heating rate of 10°C min⁻¹ scanning between 25 and 520°C in atmosphere. A bundle of filaments (34 filaments) were clamped into the sample holder under different forces applied. T_g values were measured with a dynamic mechanical analyser (DMTA V; Rheometric Scientific) for samples of ∼10 mm length using the tensile mode at a constant frequency of 1 Hz and a heating rate of 10°C min⁻¹ scanning from 25 to 450°C in atmosphere. The fibres were embedded in Spurr's epoxide resin, and the morphology of the fibres was observed with an XL30 environmental scanning electron microscope field emission gun scanning electron microscope (FEI Company). The two-dimensional (2D) wide angle X-ray diffraction (WAXD) patterns were conducted with a Bruker Discovery general area detector diffraction system (40 kV and 30 mA) equipped with a Hi-star charge coupled device detector and diffracted beam graphite monochromator, and the diameter of pinhole collimator was 0·1 mm. The one-dimensional (1D) WAXD patterns were conducted with a Rigaku D/MAX 2550 installation (50 kV and 200 mA). Small angle X-ray scattering (SAXS) experiments were performed using a NanoSTAR-U (40 kV and 650 μA; Bruker AXS Inc.) with a Hi-star detector. The distance between the sample and the detector was L_SD = 1074 mm. The effective scattering vector q [q = (4πsin θ)/λ, where 2θ is the scattering angle] at this distance ranges from 0·044 to 2·0 nm⁻¹. The samples were adopted with the transmission mode using Cu K_α1 radiation (wavelength, 1·54 Å) in all X-ray experiments.

Mechanical properties of fibres

The mechanical properties of the co-polyamic acid fibres and thermal imidised (collectively as imidised) and heat drawn co-PI fibres are listed in Tables 1–3. As could be seen in the tables, the as spun PAA fibres had generally low strength and modulus, ranging from 0·20 to 0·31 and 5·8 to 8·1 GPa respectively, while the as imidised fibres had displayed a far more increased strength and modulus, ranging from 0·66 to 0·96 and 23·1 to 29·4 GPa respectively. For heat drawn fibres, the strength and modulus of the fibres ranged from 1·10 to 2·25 and 65·0 to 96·5 GPa respectively. The strength and modulus of the BPDA/PPD PI fibres produced with the same process were 0·75 and 70 GPa respectively, which were obviously lower than all co-PI fibres containing ODA. From Table 3, it could be seen that the tensile strength of PI3 fibre was approximately three times that of BPDA/PPD PI fibre, and the modulus was also higher. It indicated that the use of ODA improved efficiently the mechanical properties of BPDA/PPD PI fibres. The best strength and modulus were observed when the ratio of ODA was 15%, and a further increase in ODA component led to decreased mechanical properties. The introduction of ODA largely improved the flexibility of the polymer chain. As the content of ODA increased, the polymer chains could obtain bigger drawing ratio in the heat treatment stage, which reduced the influence of voids in the fibre and thus promoted the mechanical properties of PI fibres. However, the internal rotation of molecular chains was easier to happen with the increase in ODA content, which impaired the mechanical properties of fibres. Once the unfavourable factors were dominant, it was inevitable to reduce the mechanical properties of fibres.

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Table 1

Mechanical properties of as spun PAA3 fibres

As spun PAA fibres	PPD/ODA	Tenacity/GPa	Modulus/GPa	Elongation/%
PAA1	95∶5	0·20	5·8	22·0
PAA2	90∶10	0·23	6·0	19·9
PAA3	85∶15	0·30	8·1	28·9
PAA4	80∶20	0·31	7·0	20·6

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Table 2

Mechanical properties of as imidised PI3 fibres

As imidised PI fibres	PPD/ODA	Tenacity/GPa	Modulus/GPa	Elongation/%
PI1	95∶5	0·79	25·5	5·7
PI2	90∶10	0·65	23·1	4·9
PI3	85∶15	0·96	29·4	6·9
PI4	80∶20	0·66	23·3	4·5

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Table 3

Mechanical properties of heat drawn PI3 fibres

Heat drawn PI fibres	PPD/ODA	Tenacity/GPa	Modulus/GPa	Elongation/%
PI1	95∶5	1·10	78·8	2·9
PI2	90∶10	1·85	85·3	2·9
PI3	85∶15	2·25	96·5	4·0
PI4	80∶20	1·25	65·0	4·5

Thermal and dynamic mechanical properties of PI fibre

The thermal behaviours of PI fibres were characterised by TGA, thermomechanical analysis (TMA) and dynamic mechanical analysis (DMA). The TGA curves of PI fibres displayed excellent thermal stability (Fig. 1). All samples showed similar degradation patterns in N₂ atmosphere. There was almost no weight loss until 480°C, and a rapid weight loss occurred at >530°C because of the thermal decomposition of the polymer backbone. The T_5% of PI-1, PI-2, PI-3 and PI-4 were 540, 535, 530 and 522°C respectively, and the T_5% of BPDA/PPD PI fibres produced by us was 557°C. The 40–55% of the original mass was retained even after heating to 800°C, which was a reasonable carbon yield for this material. Incorporating 6,4′-diamino-2-phenylbenzimidazole in polymer chains significantly enhanced the mechanical properties without sacrificing the thermal properties.

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Figure 1

Thermogravimetric analysis curves of PI fibres

Figure 2a shows the CTE data for the as imidised and heat drawn PI3 fibres at 0·1 N force applied during the TMA measurements. The CTEs of two kinds of fibres exhibited slightly negative values (−7·3×10⁻⁸°C⁻¹) between 25 and 300°C, but large difference >300°C due to the internal stress frozen into fibre during the preparation process. When the force applied overcame the internal stress, the higher CTE was observed. Furthermore, the CTE data of fibres exhibited a drastic change at ∼290°C, which was an indication of the glass transition temperature T_g.

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Figure 2

a as imidised and heat drawn fibres at 0·1 N force applied; b heat drawn fibres at 0·1–0·6 N force applied

Figure 2b represents the CTE data for the heat drawn PI3 fibres at different force levels. Generally, a negative CTE could be found at a low force applied level and then gradually increased with force applied for overcoming the internal stress. As shown in Fig. 2, the change in absolute CTE values of heat drawn fibres was smaller than that of as imidised fibres, indicating that the dimensional stability of fibres was obviously enhanced by the heat drawn process.

The DMA spectra of as imidised and heat drawn PI3 fibres were investigated from room temperature up to 450°C and presented in Fig. 3. A clear transition of E′ curves was seen at ∼270°C, generally defined as the onset of glass transition in the polymer. The storage modulus E′ of the heat drawn fibres exhibited better retention than the as imidised fibres at >300°C due to the structural improvement of fibre at the heat drawn stage.

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Figure 3

Dynamic mechanical analysis curves of as imidised (AI) and heat drawn (HD) PI3 fibres

Two major transitions were easily seen in the loss modulus E″ curves, and the as imidised and heat drawn fibres show a similar trend all over the temperature ranges. The first transition at ∼150°C displayed β relaxation, which was related to the subglass transition.28^–31 At the higher temperature of ∼290°C, the second transitions (α relaxation) of two kinds of fibres were observed, which were generally regarded as the maximum of glass transition in the polymer and bigger than that of the E′ curve. The T_g value of the E″ curve was identical with the T_g value decided by TMA.

Below the glass transition temperature, tan δ values were very close for both as imidised and heat drawn fibres. Tan δ represents the energy consumption of the segmental movement, which is related to the structural homogeneity of the polymer. A smaller tan δ value above T_g and a lower α relaxation temperature in the heat drawn fibre indicated that the structural homogeneity was evidently improved by the heat drawn process.

Morphology of PAA and PI fibre

The morphology of the fibres was studied by scanning electron microscopy (SEM) and shown in Fig. 4. As for PAA3 fibre in Fig. 4a, the cross-section of the fibre was round and void free, and yet the ‘skin–core’ structure was not obvious. These meant that the coagulation conditions were very beneficial to form a homogeneous and dense fibrous structure. Generally, the soluble PAAs are spun into coagulation to prepare PAA fibres and then subsequently converted to PI fibres by thermal imidisation. However, voids in PI fibres can be caused due to the extrusion of volatile molecular such as water during thermal imidisation, leading to the difficulty to prepare PI fibres with dense interior structure in a two-step process.32 In our experiment, PI fibres were prepared by keeping PAA fibres in a furnace of >300°C and then were subsequently drawn in a temperature higher than T_g under N₂ atmosphere. The voids could be removed during thermal imidisation, and then the molecular chains could be further rearranged and oriented along the fibre axis to form a compact structure in the heat drawn stage. As shown in Fig. 4b–c, the cross-section of the as imidised and heat drawn PI3 fibre displayed lack of voids and distinct ‘skin–core’ structure.

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Figure 4

a–c cross-section of PAA3, as imidised and heat drawn PI3 fibre, which was embedded in Spurr's epoxide resin, and d directly fractured cross-section of heat drawn PI3 fibre

An alternative method of examining the structure of the heat drawn PI2 fibres was also investigated by fracturing fibres. Figure 4d shows the directly fractured cross-section of heat drawn fibres in liquid nitrogen, indicating the lack of a well defined fibrillar structure. It was a large difference from the two sorts of PI fibres, which were prepared by Goel et al.32 and Li et al., 33 and in the latter two PI fibres, there existed a highly fibrillated structure.

Wide angle X-ray diffraction of PI fibres

Fibre diagrams were recorded for parallel bundles of fibres (34 monofilaments in a bundle); both as imidised and heat drawn PI3 fibres are shown in Figs. 5 and 6. The structural parameters including the crystal size, crystal d-spacing and molecular chain orientation were measured with WAXD and can be calculated by equations (1)–(3).

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Figure 5

Two-dimensional WAXD for as imidised (top) and heat drawn (bottom) PI3 fibre: herein, a, e, c and g were diffraction patterns in meridional and equatorial directions, in which beamstop was 0° retrospectively, while b, f, d and h were diffraction patterns in meridional and equatorial directions, in which beamstop was 18° retrospectively

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Figure 6

One-dimensional WAXD for as imidised and heat drawn PI3 fibre in a meridional and b equatorial directions respectively

The apparent crystallite size (ACS) is estimated using the Scherrer equation (1) where λ is the X-ray wavelength, β is the full width at half maximum of the crystalline peak, θ is one-half of the diffraction 2θ angle and k is usually regarded as 0·89.

The crystal interplane d-spacing is calculated with the Bragg reflection equation (2) The crystalline orientation of the fibres is determined quantitatively by computation of the empirical formula34 (3) where H is the full width at half maximum of the strongest diffraction ring (Debye ring) in the equatorial direction. The f takes the value of 100% (H = 0°) and 0% (H = 180°) with complete orientation parallel and misorientation to the drawing direction respectively.

The 2D WAXD patterns are shown Fig. 5. The as imidised and heat drawn PI3 fibres exhibited clear diffraction streaks along the meridian, indicating that polymer chains formed highly ordered structures in the fibre axial direction. On the equator, as shown in Fig. 5d with h, the equatorial arcs originated from the lateral packing were obscure, especially to as imidised fibres, which were so called ‘amorphous halos’. It meant the low ordering degree of the lateral stack. Comparing Fig. 5b with f and Fig. 5d with h, the breadths of the meridional steaks and the equatorial arcs along the azimuthal direction became narrower by heat drawn process, meaning that the orientation of molecular chains was enhanced through the heat drawn stage. The azimuth curves of two kinds of fibres are shown in Fig. 7, and the molecular orientations were calculated with equation (3) as f_AI = 0·79 and f_HD = 0·90 respectively. The absence of off-axis peak was observed in Fig. 5; this indicated that a translationally disordered form of the as imidised and heat drawn fibre was present. A similar situation also existed in the poly(p-phenylene benzobisoxazole) structure.35, 36

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Figure 8

Two-dimensional SAXS for a PAA3, b as imidised and c heat drawn PI3 fibres respectively (horizontal direction is meridian of fibres)

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Figure 7

Azimuth curves of a as imidised and b heat drawn PI3 fibre respectively

The 1D WAXD patterns along the meridional and equatorial scans of two sorts of PI3 fibres are shown in Fig. 6. In Fig. 6a, the intensity of meridional diffraction became bigger through heat drawn stage, especially the first diffraction peak, indicating that the degree of crystalline perfection of the basic segments (BPDA-PPD) was evidently improved in the fibre axis direction. Moreover, there appeared a very weak crystal peak at the angle of ∼33·4° in the heat drawn fibre. In Fig. 6b, the equatorial diffraction of heat drawn fibres emerged two obvious crystallite peaks, meaning that the lateral stack of fibres was notably improved by the heat drawn process. The ACS and crystal d-spacings of two kinds of fibres were calculated with equations (1) and (2) respectively and listed in Table 4. It could be seen from Table 4 that the crystal d-spacings slightly increased, except for the first and third diffraction peaks, and the longitudinal crystal sizes were only of subtle difference except for the second diffraction peak after heat drawn process.

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Table 4

d-spacings and apparent crystal sizes of as imidised and heat drawn PI3 fibre

As imidised PI fibre			Heat drawn PI fibre
Diffraction angle (2θ)	Crystal d-spacing/Å	Apparent crystal size/Å	Diffraction angle (2θ)	Crystal d-spacing/Å	Apparent crystal size/Å
11·3	7·82	54	11·1	7·94	55
16·9	5·23	78	16·8	5·27	64
22·9	3·88	68	22·3	3·98	65
28·1	3·18	56	28·0	3·19	56
			33·4	2·68	88
39·8	2·26	63	39·6	2·27	60
45·7	1·98	56	45·1	2·01	59

Above all, the molecular orientation, the ordering degree of longitudinal and the lateral stack were markedly improved by the heat drawn process. It was easily understood that the mechanical properties could be enhanced with heat treatment stage.

Small angle X-ray scattering of PI fibres

The scattering from microvoids typically accounts for 97–99% of the overall scattering in semicrystalline polymers37, 38 because the SAXS intensity is proportional to the square of the electron density difference between adjacent domains. Therefore, SAXS is an ideal tool to characterise nanoscale microvoids in the polymer system.39 In this section, the PI3 fibres at each stage of the manufacturing process were adopted to investigate the evolution law of microvoids within fibres by SAXS due to the adverse influences on the mechanical properties of fibres.

As shown in Fig. 8, the 2D SAXS patterns of fibres from as spun to heat drawn processes were obtained. The common feature seen in these patterns was that there were intense scatterings nearby the beamstop along the equatorial direction yet no detection along the meridional direction. These indicated that fibres existed large size needle-like microvoids, which oriented with their long axes in the fibre axis direction. The microvoid systems with preferred orientation were also known to be present in cellulose fibres40, 41 and in carbon fibres.42^–44 From Fig. 8, the scattering width of fibres along the meridional direction obviously decreased (W_a>W_b, where W is the width of scattering peaks) in as spun and as imidised process, meaning that the orientation of fibres was significantly enhanced in the as imidised process due to the reciprocal law of SAXS. Comparing with Fig. 8b, there existed lots of new scattering dots along both sides of the equatorial direction in Fig. 8c, which came from the better lateral crystallites in the heat drawn fibre. This phenomenon could be confirmed with Fig. 6b. The results indicated that the higher orientation and the stack ordering of polymer chain were important factors that improved the mechanical properties of fibres.

Figure 9 shows the 1D SAXS patterns along the meridional and equatorial directions of fibres at each preparation stage. The common tendency in two patterns was that the scattering intensity decreased from as spun to as imidised process, indicating that the structure of the fibres became more homogeneous, mostly ascribing to the elimination of microvoids in the two processes. However, the scattering intensity of the heat drawn fibres remained unchanged and enlarged with respect to that of as imidised fibres in the meridional and equatorial directions respectively. The possible reason for the results was that the misorientation of microvoids reduced; meanwhile, the amount of microvoids slightly increased due to the excessive stretching in the heat drawn process.

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Figure 9

One-dimensional SAXS for PAA3, as imidised and heat drawn PI3 fibres along a meridional and b equatorial directions respectively

The fractal dimension D_s can be a measure of the internal surface roughness in a microvoid material, which follows a power law I(q)∝q^−α. For mass fractals, α = D_m and 1<α<3. In contrast, we have α = 6−D_s for surface fractals and 3<α≤4. If D_s = 2, we obtain the well known Porod's law I(q)∝q⁻⁴ for non-fractal structures with smooth interfaces.45^–47 The bigger the D_s value, the larger the surface roughness of the microvoid. As shown in Fig. 10b, the slopes of ln I(q) versus ln q curves in the linear region were calculated, with linear regression analysis, as −3·47, −3·86 and −3·81, corresponding to fractal dimensions of 2·53, 2·14 and 2·19 for PAA3, as imidised and heat drawn PI3 fibres respectively. The results indicated that the surface roughness of microvoids obviously decreased in as imidised process and then slightly increased in heat drawn process.

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Figure 10

a, b LnI(q) versus Lnq relationships of PAA3, as imidised and heat drawn PI3 fibres along equatorial direction respectively

Effects of heat treatment on mechanical properties of PI fibres

Methods for improving the strength of fibres usually involve increasing the degree of crystallinity and the molecular orientation, decreasing the content of voids and the structural homogeneity or some combination of these factors.48, 49 In this section, PI3 fibres were treated under variously thermal conditions; the effects of temperature, time and drawing ratios on the mechanical properties of fibres were discussed. There was only a variable that can be changed among three heat treatment conditions in one experiment.

As shown in Fig. 11a, the tenacity and modulus of PI3 fibres increased with drawing ratios until ∼1·8 (temperature and time remained 540°C and 11 s respectively) and then decreased. The orientation of molecular chains was first developed with the increase in drawing ratios to the critical point, yet meaning that the mechanical properties of fibres was enhanced at this stage. However, once exceeding the critical point, a few molecular chains with high orientation on the early stage began to be broken and led to the residual molecular chains in the fibres to be seriously disrupted. As a result, the mechanical properties of the fibres started to descend.

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Figure 11

a–c effects of heat drawn process parameters on mechanical properties of PI3 fibres

In Fig. 11b, it could be seen that the preferable mechanical properties were obtained at ∼540°C (drawing ratio and time remained 1·8 and 11 s respectively). Commonly, the flowability and rearrangement of molecular chains became easier as the temperature increased, which made fibres have better mechanical properties due to the higher orientation and less defects. Nevertheless, the degradation and misorientation of molecular chains started to play an important role when the temperature further increased. Once the adverse effect dominated the structure evolution, the mechanical properties inevitably declined.

As shown in Fig. 11c, the mechanical properties gradually increased with heat treatment time. The structural heterogeneity was the main restricted factor to the mechanical properties of fibres when the heat treatment time was short (temperature and drawing ratio remained 540°C and 1·8 respectively). As the heat treatment time prolonged, the mechanical properties could be enhanced by eliminating microvoids of fibres.