single-walled carbon nanotubes fibers

Composites Science and Technology 83 (2013) 47–53

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Stretching induced interfacial crystallization and property enhancement of poly(L-lactide)/single-walled carbon nanotubes fibers q Wei Zhang, Nanying Ning, Yao Gao, Fan Xu, Qiang Fu ⇑ College of Polymer Science and Engineering, State Key Laboratory of Polymer Materials Engineering, Sichuan University, Chengdu 610065, PR China

a r t i c l e

i n f o

Article history: Received 22 February 2013 Received in revised form 23 April 2013 Accepted 25 April 2013 Available online 3 May 2013 Keywords: A. Carbon nanotubes A. Nanocomposites B. Interface B. Mechanical properties Stretching induced

a b s t r a c t In this work, poly(L-lactide) (PLLA)/single-walled carbon nanotubes (SWNTs) fibers were first prepared at high drown ratio and low drown ratio, respectively. Then they subjected to tensile testing. It was found that the composite fibers obtained at low drown ratio showed no obvious enhancement of tensile strength and elongation at break, compared with the pure PLLA fibers obtained at the same low drown ratio. However, a significant property enhancement was observed for the composites fibers obtained at high drown ratio. Structure analysis of the as spun fibers before and after tensile testing suggests a possible stretching induced formation of brush-like hybrid structure in which the PLLA lamellae growing perpendicular to the SWNTs axis for fibers obtained at high drown ratio. This unique brush-like hybrid structure could largely enhance the interfacial interaction between PLLA and SWNTs, thus results in greatly improved tensile strength and elongation at break. Ó 2013 Elsevier Ltd. All rights reserved.

1. Introduction Filler reinforced polymer composites have gained many attentions for scientists and engineers in recent decades due to their good mechanical performance and a wide range of applications in industry. Since discovered by Iigima [1,2], the carbon nanotubes (CNTs) possessing superlative stiffness and strength are considered promising candidates for mechanical reinforcement of polymers. Many studies have focused on the mechanical properties of polymer/CNTs composites and attempted to maximize the reinforcement as much as possible [3–8]. However, the effect of the polymer reinforcement by adding CNTs is different in various systems, especially in the amorphous and semicrystalline polymers. Yuan et al. [3] fabricated the polystyrene and pristine multiwalled carbon nanotubes (MWNTs) composites by melt-mixing with uniform nanotube dispersion. However, they observed no mechanical enhancement with the addition of MWNTs and the microscopic morphology analysis showed the poor polymer–nanotube interfacial adhesion. In the study performed by Gorga and Cohen [4], the modulus and strength of polymethyl methacrylate (PMMA) were increased by 38% and 25% respectively with the addition of 10 wt% MWNTs. Meincke et al. [5] improved the modulus of poly-

q This is an open-access article distributed under the terms of the Creative Commons Attribution-NonCommercial-No Derivative Works License, which permits non-commercial use, distribution, and reproduction in any medium, provided the original author and source are credited. ⇑ Corresponding author. E-mail address: [email protected] (Q. Fu).

0266-3538/$ - see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.compscitech.2013.04.019

amide-6 (PA-6) from 2590 MPa to 4130 MPa (59% increment) by adding 10 wt% MWNTs. Ganß et al. [6] prepared the composites of polypropylene (PP) and MWNTs and they observed 13% and 10% increase in modulus and yield strength only with addition of 0.5 wt% MWNTs. It seems that the reinforcing efficiency of MWNTs in semicrystalline polymers is higher than that in amorphous ones. Cadek et al. [7] compared the mechanical enhancement of the CNTs-reinforced semicrystalline polymer (polyvinyl acetate (PVA)) and amorphous polymer (poly-9-vinyl carbazole (PVK)). They observed that the increase of the modulus and strength were more remarkable in PVA than that in PVK. They considered the crystalline layer of PVA nucleated by CNTs might lead to better stress transfer than the amorphous interface of PVK around CNTs. It is clear that the crystalline interface plays an important role in enhancing mechanical properties of CNT–semicrystalline polymer composites. Poly(L-lactide) (PLLA) as a biodegradable semicrystalline polymer has been wildly applied in biomedical use, food packaging, electronics and automobile field in recent years [9,10]. In order to enhance the properties and enlarge the application of PLLA, many nanofillers have been incorporated into PLLA matrix, such as CNTs [11–13]. To produce high performance PLLA/CNTs composites, the good interfacial adhesion between the PLLA and CNTs and the alignment of the CNTs in matrix could be the key factors, except for the well dispersion of CNTs. Many studies showed that the CNTs acting as nucleating agent for most semicrystalline polymers also significantly affected the crystallization kinetics and crystallization morphology of PLLA [14,15], as well as the mechanical properties [12]. Among the various methods of enhancing the

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interfacial adhesion between polymers and fillers, the interfacial crystallization is reported as an efficient and economical way [16], which means the good interaction between the PLLA and CNTs could be realized by PLLA crystallization on CNTs under certain condition. Unfortunately, the PLLA shows low crystallization kinetics [15,17,18], even if CNTs are incorporated [19,20]. Therefore, only amorphous PLLA matrix can be observed in PLLA/CNTs nanocomposites prepared by common processing methods (e.g., injection molding). On the other hand, shearing has been demonstrated to play an important role in controlling the structure and properties of polymer and its composites [21–23]. Very recently, many studies have confirmed the occurrence of a stretching-induced crystallization upon drawing amorphous PLLA [24–26]. Both of these interesting results drive us to investigate the effects of the stretching-induced crystallization on the interfacial interaction and the resultant mechanical properties of PLLA/CNTs nanocomposites. In this study, the single-walled carbon nanotubes (SWNTs) owning good nucleation ability for PLLA due to their rough surface [14,20,27–29] were selected as the reinforcing fillers for PLLA and the melt spinning method was used to prepare PLLA/SWNTs composite fibers. We attempt to obtain two kinds of fibers with low and high pre-orientation of polymer chains and SWNTs by adopting different draw ratios (DRs), and the evolution of structures especially the interfacial crystalline structure during the stretching process and the properties of the composite fibers are deeply investigated. The goal of this study is to obtain PLLA/CNTs composites with high performances and more importantly, to help us further understand the mechanism of filler-reinforced semicrystalline polymers. 2. Experimental 2.1. Materials All the materials used in this study are commercially available. PLLA (4032D, NatureWorks LLC) with high stereoregularity (1.2– 1.6% D-isomer lactide) was used as the matrix polymer. It exhibits a density of 1.25 g/cm3. The weight averaged molecular weight (Mw) and polydispersity are 207 kDa and 1.74, respectively. Raw SWNTs were purchased from Chengdu Organic Chemicals Co. Ltd., Chinese Academy of Sciences. Its diameter is about 1–2 nm and its length is about 5–30 lm, the purity is larger than 60 wt%. SWNTs were used as received without any chemical modification. 2.2. Preparation of PLLA/SWNTs composites Both SWNTs and PLLA were dried in vacuum at 60 °C for 15 h to remove moisture prior to composites preparation. The melt blending of PLLA with a weight percentage of 0.1 and 0.3 wt% of SWNTs was conducted using a TSSJ-2S co-rotating twins-crew extruder. The temperature was maintained at 160, 170, 170, 180, 185, 185, 185 and 180 °C from hopper to die and the screw speed was kept at 100 rpm. The composites were melt-spun as fibers at 185 °C using a set of melt-spun devices adapted to the piston-mode Rosand RH70 (Malvern, Bohlin Instruments) capillary rheometer. The capillary has a die diameter of 1.0 mm and a length-to-diameter ratio of 16. The extruded-composite fibers were air-cooled and picked up under tension by a windup spool. To prepare fibers with different draw ratio (DR) (defined as the section area of capillary die versus that of fiber), two take-up speeds were selected. The take-up speeds were about 10 and 50 m min1 for low and high take-up speed, respectively. In a succinct manner, the as-spun fibers were denoted as xSP-L/H, where x represents the percentage of SWNTs and L/H represents low/high draw ratio, respectively.

For example, 0.3SP-L represents the fiber containing 0.3 wt% SWNTs prepared upon low draw ratio. For comparison purposes, the fibers of pure PLLA have also been prepared using the same conditions described above. And these neat PLLA fibers are marked as PLLA-L/H. 2.3. Characterization The tensile experiments were carried out on an Instron 5567 universal testing machine with a 100 N-load cell. The fibers were fixed on a paper frame and tested with a crosshead speed of 5 and 100 mm min1 for modulus and strength measurements, respectively. The gauge length was 10 mm, and the measured temperature was around room temperature (23 °C). The reported values were calculated as averages over five specimens for each composite fiber. Microscopic morphology observations were conducted on an FEI Inspect F field emission scanning electron microscope (FESEM, USA) under an acceleration voltage of 5 kV. Two types of SEM measurement were adopted for identifying crystalline structure and interfacial interaction between SWNTs and matrix, respectively. For crystalline structure observation, the fibers were first chemically etched in sodium hydroxide solvent (pH = 13) for 5 h at 25 °C to remove the amorphous phase of PLLA and then the cylindrical surfaces of the fibers were coated with gold. In another method, the fibers were cut in liquid nitrogen perpendicular to the flow direction, then the fractured surfaces were gold-coated and observed. Polarized Raman spectroscopy was implemented using a microRaman spectrometer (Renishaw) equipped with a microscope. The laser with a wavelength of 514.5 nm was excited by a 136 M He+ resource. The excited laser with a power of 1.7 mW and a spot diameter of 2 lm was used. To estimate the orientation of SWNTs in composite fibers, the dichroism Raman spectra were achieved by recording Raman spectroscopy along two directions normal between each other, which were parallel and perpendicular to the long-axis (length) direction of fiber, respectively. All received spectra have been calibrated by baseline subtraction. As the D⁄ mode is the second-order overtone mode, and in principle its intensity does not depend on the presence of disorders (except for those effects that affect the first-or higher-order scatterings equally). Thus, it is a suitable reference for peak intensity normalization. In this study, all Raman resonances were normalized with respect to the D⁄ mode. The melting behaviors of the specimens were characterized with a Perkin–Elmer Pyris-1DSC (USA) under a dry nitrogen atmosphere. The mass of testing sample was around 5 mg. The sample was firstly held at 30 °C for 5 min and then heated from 30 to 200 °C at a rate of 10 °C min1. The degree of crystallinity (Xc) for each specimen was evaluated with following expression: Xc = (DHm  DHc)/DH0m, where DHm and DHc are the measured enthalpies of melting and cold crystallization, respectively, and DH0m is the fusion enthalpy of 100% crystalline PLLA. The value of DH0m was selected as 93.7 J g1 [30]. Two-dimensional wide-angle X-ray scattering (2D-WAXS) experiments were conducted on a Bruker Discover 8 diffractometer. The wavelength of the monochromated X-ray from Cu radiation was 0.154 nm, and the sample-to-detector distance was 273 mm. The background of all the 2D-WAXS patterns given in this article had been extracted and thus allows a qualitative comparison between various samples. The samples were placed with the orientation (drawing direction) perpendicular to the beams. Azimuthal scans (0–360°) 2D-WAXS were made for the (1 1 0) plane of PLLA at a step of 1°. The orientation degree of crystal planes could be calculated by the orientation parameter f from the following equation:

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W. Zhang et al. / Composites Science and Technology 83 (2013) 47–53

fH ðcos uÞ ¼

3hcos2 ui  1 2

hcos2 ui ¼

R p2 0

IðuÞ cos2 u sin udu R p2 0 IðuÞ sin udu

where u is the angel between the normal of a given (h k l) crystal plane and the drawing direction and I is the intensity. A perfectly perpendicular orientation gives f = 1.0. An un-oriented sample gives f = 0.

The elongation at break is sharply increased to 445% for 0.1SP-H and 590% for 0.3SP-H. Correspondingly, the strength at break is obviously enhanced from 65 MPa for PLLA-H to 95 MPa for 0.3SPH. Obviously, with the addition of SWNTs, the reinforcement of PLLA is more efficient in high DR samples than that in low DR samples. The superior strengthen and toughness PLLA/SWNTs composites fiber is obtained at high DR. 3.2. Orientation and interfacial interaction of PLLA molecular chains and SWNTs

3. Results and discussion 3.1. Tensile properties of the as-spun fibers The typical stress–strain curves of neat PLLA and PLLA/SWNTs composites prepared by melt spinning at different draw ratios are shown in Fig. 1. As for the samples prepared at a relatively low DR, both of the tensile strength and tensile modulus of PLLA are not significantly changed with the addition of SWNTs, whereas the elongation at break is increased slightly from 135% for pure PLLA-L to 180% for 0.1SP-L, and further increased to 205% for 0.3SP-L, as shown in Fig. 1a. This suggests that the mechanical properties of PLLA change a little with the addition of SWNTs for the composite fibers prepared at relatively low DR. However, as the DR increased, a much more remarkable increase in strength and elongation at break can be observed by adding SWNTs, as shown in Fig. 1b. For PLLA-H, the elongation at break is about 237%, which is larger than PLLA-L, which could be due to pre-orientation and disentanglement of PLLA chains in high DR fibers.

As we all know, the alignment of the polymer molecular chains and the fillers have a significant impact on the mechanical properties of polymer composites [31–33]. The 2D-WAXS technology was employed to get necessary information for the PLLA matrix in the melt-spun fibers. The fibers were annealed at 100 °C for 20 min because the as-spun fibers are almost amorphous that the orientation of PLLA lamellae cannot be detected through 2D-WAXS. It has been demonstrated that the molecular orientation in the crystalline phase of melt-spun PLLA fiber remains almost constant during the annealing [34]. As shown in Fig. 2, the scattering signals are designated to the (1 1 0) plane and (2 0 3) plane of orthorhombic PLLA lamellae. Corresponding azimuthal scans for (1 1 0) plane (Fig. 2e) transformed from 2D patterns are also presented for clarity. The orientation degree of all samples obtained from 2D-WAXS

(a)

70

(b)

(203)

(a) 60

(110)

Stress (Mpa)

50 40

0.3SP-L 30

0.1SP-L PLLA-L

20

(d)

(c)

10 0 0

30

60

90

120

150

180

210

Strain (%) 100

(b)

90

70

(e)

0.3SP-H

60

Relative intensity

Stress (Mpa)

80

0.1SP-H

50 40 30

PLLA-H

20 10

f

0 0

100

200

300

400

500

600

Strain (%) Fig. 1. The stress–strain curves of the PLLA/SWNTs fibers at different SWNTs loading prepared upon different draw ratios, (a) low draw ratio; (b) high draw ratio.

0.55

A-0.3SP-H

0

60

120

A-PLLA-H

0.58

A-0.3SP-L

0.23

A-PLLA-L

0.21

180

240

300

360

Azimuthal angle (degree) Fig. 2. 2D-WAXS patterns of fibers that annealed at 100 °C for 20 min (a) A-PLLA-L, (b) A-PLLA-H, (c) A-0.3SP-L, (d) A-0.3SP-H, and (e) azimuthal scans of these fibers.

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patterns is listed in Fig. 2e. For the fibers containing 0 and 0.3 wt% SWNTs, the orientation degree of the (1 1 0) plane is 0.21 and 0.23 for low DR and 0.58 and 0.55 for high DR, respectively. This indicates that the increase of DR certainly improves PLLA chains orientation and the addition of SWNTs cause little change of the orientation of PLLA matrix. Thus, the change of orientation of matrix after adding SWNTs is not the main reason for the mechanical improvement of high DR fibers. In a word, a higher orientation level of PLLA matrix were obtained though a simple melt-spun method by using high DR, which might have a significant impact on the mechanical properties of the composite fibers during the stretching and we will discuss it in the next section. In addition, the polarized Raman spectroscopy was conducted to detect the orientation level of SWNTs in PLLA/SWNTs composite fibers prepared at different draw ratios, as shown in Fig. 3a. For each sample, Raman spectra are recorded parallel to (0°) and perpendicular to (90°) the stretching direction of the fibers. The characteristic peak of SWNTs at about 1585 cm1, which is assigned to the G band and is associated with tangential C–C bond stretching motions, could be observed for both samples. The orientation degree of SWNTs is evaluated by the depolarization factor, R, which is defined as the ratio of the peak intensity for the G band in the parallel direction (I0) to that in the perpendicular direction (I90), (R = I0/I90). For the fibers prepared under low DR (0.3SP-L), the value of RL is approximate to 1.2, indicating the low orientation level of SWNTs in PLLA matrix. However, for the fibers prepared under high DR (0.3SP-H), the value of RH increases remarkably to 2.1, suggesting the preferential orientation of SWNTs in PLLA matrix along the draw direction. In brief, the SWNTs are preferentially aligned along the fiber axis for high DR samples. It is well known that extension could not only enhance the orientation of polymer matrix and fillers but also their interfacial interaction [35,36]. Raman spectroscopy was also used to explore the interfacial adhesion between PLLA and SWNTs. The high frequency Raman spectra for raw SWNTs, and the composite fibers with 0.3 wt% SWNTs prepared under low and high DR are shown in Fig. 3b. We can clearly observe that the characteristic peak of SWNTs (G band) shifts from 1584 cm1 for raw SWNTs to 1587 cm1 for 0.3SP-L. This could be due to the relatively weak CH–p interfacial interaction between PLLA molecular chains and SWNTs, as reported in many works [37,38]. For 0.3SP-H, the peak of G band further blue-shifts to 1591 cm1, indicating stronger interfacial interaction between PLLA and SWNTs as prepared under higher DR. 0.1SP-H get the same results. Similar result can be obtained from the Raman spectra of 0.1SP-L and 0.1SP-H. For briefness, only the spectra of 0.3SP-L and 0.3SP-H are presented in Fig. 3b. These results indicate that more PLLA molecular chains are absorbed onto SWNTs surface due to the orientation of extensive PLLA chains under a stronger extension field.

3.3. The structure changes during stretching To further explore the mechanism for the enhanced mechanical properties, the structure changes during stretching were investigated. As we know that PLLA has a relatively low crystallization rate under quiescent conditions, thus an amorphous state is usually obtained. The accelerated crystallization of PLLA could be observed during melt spinning due to the molecular extension under external force. However, the rapid solidification could result in a low crystallinity in the as-spun PLLA fibers. As shown in Fig. 4a, all the as-spun PLLA and PLLA/SWNTs composite fibers have very low crystallinity (0.3–2.6%) even for that prepared under a relatively high DR. However, for the broken fibers after mechanical test, the cold crystallization peak (Pc) is significantly decreased for both PLLA-L and 0.3SP-L after the stretching process, and the corresponding crystallinity is largely increased to about 18% for PLLA-L and 22% for 0.3SP-L. For PLLA-H and 0.3SP-H, Pc is even absent after stretching and the corresponding crystallinity is sharply increased to about 34% for both samples. Thus, the stretching-induced crystallization occurs in the fibers during the stretching process. Factually, for the shear induced crystallization of semicrystalline polymer filled with SWNTs, it has been widely demonstrated that SWNTs can dramatically alter the crystallization kinetics and the development of morphology under flow conditions [39,40]. The combined effects of shear and the presence of SWNTs may yield a synergistic increase in the number of active nuclei. To further clarify the structure changes during stretching process, the fibers were stretched to a designated elongation such as 50%, 100%, 150% and then the melting behavior of the stretched fibers were investigated. The crystallinity variation of 0.3SP-H during stretching is shown in Fig. 4c. Obviously, the area of Pc gradually decreases and then disappears, and the corresponding crystallinity versus strain is summarized in Fig. 4d. For 0.3SP-H, the evolution of the crystallinity during stretching could be described as two steps. Firstly, the crystallinity grows rapidly from 1.7% to 34.3% when the corresponding strain reaches 200%. Secondly, when the strain increases from 200% to 590%, the crystallinity remains almost constant around 35%, which means the stretching-induced crystallization mainly occurs in the first step. In addition, at the equivalent strain, the fibers prepared under high DR show higher values of crystallinity compared with these under low DR. In order to confirm the crystalline structure induced by stretching, the as spun PLLA/SWNTs composite fibers and the corresponding stretched fibers were all etched in sodium hydroxide solvent at 25 °C, and then were observed by SEM. For the as-spun composite fibers, only amorphous state can be observed for both fibers prepared under low and high DR, as shown in Fig. 5a and c, respec-

(a)

(b)

1480

Intensity

Indensity

0.3SP-L-0 0.3SP-L-90

0.3SP-H-0 0.3SP-H-90

1520

1560

1600

1640

Raman Shift (cm-1)

1680

SWNT 0.3SP-L 0.3SP-H

1584 1587

1591

1515 1530 1545 1560 1575 1590 1605 1620 1635

Raman Shift (cm-1)

Fig. 3. (a) Polarized Raman spectra of composite fibers (0.3SP-L and 0.3SP-H) parallel (0°) and perpendicular (90°) to fiber axis. (b) Raman spectrum (G band) of raw SWNTs and composite fibers with 0.3 wt% SWNTs prepared at different draw ratios.

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(a) As spun

Xc %

(b) After stretch

PLLA-L 0.3% 0.3SP-L

PLLA-L-S 17.7%

0.5%

0.3SP-L-S

endo 80

100

120

21.7%

PLLA-H-S 34.2%

PLLA-H 2.6% 0.3SP-H

Xc %

0.3SP-H-S

1.7%

140

160

80

180

100

O

120

34.8%

160

140

180

O

Temperature C

Temperature C 40

Xc %

0.3SP-H-0%

1.7%

0.3SP-H-50%

16.5%

0.3SP-H-100%

25.9%

0.3SP-H-150%

30.2%

0.3SP-H-200%

34.3%

0.3SP-H-400%

35.4%

35

Crystallinity (%)

endo

(c) During stretch

(d)

30 25 20

0.3SP-H 0.3SP-L PLLA-H PLLA-L

15 10 5 0

80

100

120

140

160

180

0

100

200

300

400

500

600

Strain (%)

O

Temperature C

Fig. 4. DSC melting curves of PLLA and PLLA/SWNTs composite fibers with different draw ratios, (a) the samples were as-spun ones without any treatment, (b) these samples were collected after stretching test, (c) the melting curves of 0.3SP-H fibers collected at different elongation state during stretching test, (d) the crystalline phase fraction versus plastic strain for these four kinds of fibers.

tively. However, for the corresponding stretched fibers, the obvious crystallization structure of PLLA can be clearly observed, as shown in Fig. 5b and d. Compared with 0.3SP-L, the PLLA lamellae of 0.3SP-H is much denser. This suggests the much higher nucleation density of PLLA chains on the SWNTs surface for 0.3SP-H than that for 0.3SP-L. In addition, the image (Fig. 5d) reveals the SWNTs are surrounded by large amounts of PLLA lamellae, showing a brush-

(a)

(b)

like hybrid structure. This special structure has been confirmed by large-scale observation of the stretched fiber 0.3SP-H through our experiment. This structure induced by stretching resembles the nanohybrid shish-kebab (NHSK) structure with lamellae growing perpendicular to the fiber axis, which could improve the interfacial bonding between PLLA and SWNTs. In contrast, for the sample produced by low DR, almost no such structure can be clearly observed in Fig. 5b. Based on the above results, this amazing brush-like hybrid structure caused by stretching-induced interfacial crystallization in the high DR composite fibers may play an important role in the improvement of the mechanical properties.

3.4. Mechanism for the enhanced mechanical properties

(c)

(d)

Brush-like hybrid

SWNTs 5 µm

Fig. 5. The SEM images of fibers etched by sodium hydroxide solvent, the surface morphology of the as-spun one (a) 0.3SP-L, (c) 0.3SP-H, and after stretching one (b) 0.3SP-L, (d) 0.3SP-H.

To deeply understand the mechanism for the enhanced mechanical properties, the morphology of the stretched composite fibers at an elongation of 200% and the fractured surfaces are observed by SEM, as shown in Fig. 6. For both composite fibers prepared at different DR, as shown in Fig. 6a and d, obvious necking occurs at an elongation of 200% due to the disentanglement and orientation of PLLA chains during stretching with accompanied by the crystallization of PLLA. For 0.3SP-L, a large number of cracks perpendicular to the tensile direction are induced at an elongation of 200%, and the break immediately occurs with further elongation (205%), as shown in Fig. 6a. The fracture surface is located at the region where necking behavior has not appear yet, as shown in the inset in Fig. 6a. This indicates that the fracture occurs at the amorphous region, not the crystallization region (necking area). Under higher magnification, we can see clearly the crack distance (about 3000 nm), as shown in Fig. 6b. For the fractured surface of 0.3SP-L shown in Fig. 6c, one can clearly see that the curved SWNTs bundles are indeed pulled out from the matrix, suggesting the relatively poor interfacial interaction. However, for 0.3SP-H, no crack

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W. Zhang et al. / Composites Science and Technology 83 (2013) 47–53

(b)

(a)

(c)

10 µm

200 µm

(f)

(e)

(d)

100 µm

5 µm

5 µm

3 µm

Fig. 6. SEM micrographs of fibers’ surface morphology, the fiber (a and b) 0.3SP-L, and (d and e) 0.3SP-H stretched to e = 200%, then etched by sodium hydroxide solvent, fractured surfaces of (c) 0.3SP-L, and (f) 0.3SP-H.

Fig. 7. Schematic of strain-induced interfacial crystallization and brush-like structural growth mechanism during stretching.

is induced under stretching, and the fracture occurs at the necking region at an elongation of 200%, as shown in Fig. 6d. In the necking region, a highly oriented crystallization structure is induced during stretching, as shown in Fig. 6e. Meanwhile, in the fractured surface (shown in Fig. 6f) the interface between PLLA and SWNTs is blurrier, and instead of pulled-out from the matrix as is occurred for 0.3SP-L, SWNTs bundles are broken under the high stretching force, which indicates the very good interfacial interaction between SWNTs and PLLA matrix in 0.3SP-H. To represent the mechanism for the enhanced mechanical properties, a schematic model is proposed based on the above discussions, as shown in Fig. 7. For the composite fibers prepared at a high DR (0.3SP-H), both PLLA chains and SWNTs are extended parallel to the drawing direction due to the strong shear stress imposed on the cooling melts during the melt spinning process. Under this condition, SWNTs not only act as nucleating sites as quiescent conditions but also retard the relaxation of the extended chains, thus forming large amounts of active nucleation sites on the SWNTs surface. Then, during subsequent stretching, PLLA chains crystallize gradually to form the brush-like hybrid superstructure in the end (200%). In the second strain domain (200–

590%), the interfacial adhesion between matrix and SWNTs is significantly enhanced due to the brush-like hybrid structure formed on the SWNTs surface in the first step, which is a the crucial factor for the mechanical enhancement of 0.3SP-H. In this case, the efficiency of load transfer from the matrix to SWNTs could be improved and the formation micro-cracks could be easily terminated, thus leading to the increase of the elongation until the break of SWNTs bundles. Finally, the elongation at break is nearly 600%, and the corresponding strength at break is obviously enhanced to 95 MPa, which are much larger than that of PLLA-H. In a word, the formation of stretching-induced brush-like hybrid structure enhances the interaction between the SWNTs and PLLA matrix and plays an important role in the largely enhanced mechanical properties of 0.3SP-H. For the composite fibers prepared under low DR (0.3SP-L), the low orientation degree of both PLLA chains and SWNTs are most likely existed in the initial state. In this situation, PLLA chains are hard to align onto the SWNTs surface during the subsequent stretching, which results in a weaker heterogeneous nucleation ability of SWNTs for PLLA. Therefore, it is unfavorable for the formation of interfacial crystallization during the stretching and the

W. Zhang et al. / Composites Science and Technology 83 (2013) 47–53

improvement of the interfacial adhesion is really limited. A large number of cracks are formed during stretching and the SWNTs are mainly pulled out from the matrix. In this case, the elongation at break (205%) is higher than that of PLLA-L (135%), but much lower than that of 0.3SP-H (600%). 4. Conclusion In summary, a mechanical reinforcement has been observed in the melt-spun PLLA/SWNTs fibers. The composite fibers prepared under high draw ratio show an order of improvement in toughness than the PLLA-H, meanwhile the corresponding break strength increases remarkably. With the increase of draw ratio, both of the PLLA chains and SWNTs can be oriented well along the draw direction. Combining with the observation of the structure changes during stretching, the mechanism for the enhanced mechanical properties could be understood as follows. Firstly, the as-spun fibers were obtained almost in an amorphous state, while the SWNTs in composite fibers obtained under high draw ratio (e.g. 0.3SP-H) shows better nucleation efficiency for PLLA matrix comparing with that prepared at low draw ratio during stretching. Secondly, the stretching-induced crystallization obviously happened in the high DR composite fibers and the PLLA chains nucleate and grow on the SWNTs surface to form the interfacial brush-like hybrid structure during the stretching. Therefore the tensile stress is increased in the second step sharply attributing to excellent interfacial adhesion between PLLA and SWNTs induced by the brush-like hybrid structure formed in the first step. In a word, the remarkable reinforcement effect in high DR fibers is mainly attributed to the interfacial crystallization formed during stretching process. Our work provides a good example that the interfacial adhesion between the fillers as well as the property enhancement could be achieved via interfacial crystallization. Acknowledgement We would like to express our sincere thanks to the National Natural Science Foundation of China for Financial Support (51121001 and 51210005). References [1] [2] [3] [4]

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