Orientation of Well-Dispersed Multiwalled Carbon Nanotubes in Melt

Article pubs.acs.org/Macromolecules

Orientation of Well-Dispersed Multiwalled Carbon Nanotubes in Melt-Spun Polymer Fibers and Its Impact on the Formation of the Semicrystalline Polymer Structure: A Combined Wide-Angle X‑ray Scattering and Electron Tomography Study Thomas Vad,*,† Johannes Wulfhorst,† Tian-Tian Pan,† Wilhelm Steinmann,† Sarah Dabringhaus,† Markus Beckers,† Gunnar Seide,† Thomas Gries,† Wiebke F. C. Sager,‡ Markus Heidelmann,§ and Thomas E. Weirich§ †

Institut für Textiltechnik (ITA), RWTH-Aachen University, D-52074 Aachen, Germany Soft Matter, Institute of complex systems, Forschungszentrum Jülich, D-52425 Jülich, Germany § Central Facility for Electron Microscopy (GFE), RWTH-Aachen University, D-52074 Aachen, Germany ‡

ABSTRACT: The orientation behavior of well-dispersed multiwalled carbon nanotubes (CNTs) within high-speed melt-spun semicrystalline polymer fibers has for the first time been studied using three-dimensional reconstructions from bright-field transmission electron microscopy (TEM) tomography. The local investigation technique allows separating contributions stemming from additionally present CNT aggregates. Over a relatively narrow draw ratio range applied during the fiber production process, a transition region is found, in which the CNTs change their orientation from being aligned perpendicular to being aligned parallel to the fiber axis. Complementary performed wide-angle X-ray scattering measurements and mechanical analysis of the polymer/CNT nanocomposite fibers reveal a strong correlation between the CNT orientation and the structural and mechanical properties of the fibers. Characteristic quantities such as crystallinity, crystal size, and correlation length parameters of crystalline and amorphous polymer chains undergo significant changes within the CNT orientation transition region indicative of a cooperative process.

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state drawing of the fibers by a certain draw ratio λD enhances these effects, thus leading to fibers with higher tensile strength and, correspondingly, to a loss of flexibility. Since it is known that unidirectional polymer nanocomposites show the largest reinforcement effects, the axial orientation of the CNTs with respect to the fiber axis plays an important role for the improvement of the mechanical fiber properties.13 To date, no comprehensive or systematic study has been reported in the literature that focuses on the dependence of CNT orientation, fiber processing, and mechanical properties under industrially relevant processing conditions such as takeup velocities in the range 1000 m/min ≤ vt ≤ 10000 m/min and MWCNTs as filler material. The published results do, however, not even provide a uniform picture. While Pötschke et al.13 report a significant increase in the tensile strength of the fiber by adding MWCNTs to polycarbonate, Bhattacharyya and co-workers14 find that the mechanical properties of polypropylene fibers remain more or less unchanged when reinforced with SWCNTs.

INTRODUCTION The properties of polymer fibers can generally be adjusted by incorporating different types of nanoparticles in order to meet specific requirements, such as mechanical, electrical, or optical properties. Carbon nanotubes (CNTs) have in this context intensively been investigated as reinforcement fillers, since they exhibit a high mechanical strength modulus, high aspect ratio, and high electrical conductivity, enhancing thus the mechanical,1−4 thermal,5,6 and electronic7,8 properties of polymer fibers. While single-walled CNTs (SWCNTs) display due to their well-defined rod-like shape higher tensile strength and a better mechanical load transfer, multiwalled CNTs (MWCNTs) are, because of their low costs, preferred in the industrial manufacturing of CNT reinforced polymer fibers. Besides solution mixing9 and in-situ polymerization,10 melt mixing11 is the preferred method for the fabrication of polymer−CNT nanocomposite fibers especially on industrially relevant scales. In general, the mechanical properties of melt-spun polymer fibers are determined by the spinning process parameters applied.12 The orientation of the polymer chains as well as the formation of the semicrystalline polymer structure can be controlled by the take-up velocity vt, whereby additional solid© 2013 American Chemical Society

Received: January 18, 2013 Revised: June 22, 2013 Published: July 5, 2013 5604

dx.doi.org/10.1021/ma4001226 | Macromolecules 2013, 46, 5604−5613

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Table 1. Processing Parameters for the Melt Spinning of PET/CNT and PA6/CNT Composites parameter capillary geometry no. of filaments cap. diameter [mm] cap. length [mm] extrusion temp [°C] mass flow rate [g/min] shear rate (in capillary) [s−1] extrusion velocity [m/min] winding speed [m/min] inline draw ratio take-up speed [m/min]

PET + 0.5 wt % CNT

PA6 + 0.1 wt % CNT

17 0.35 0.70 295 54 1.40 × 104 36.7 3500 1.00/1.51/2.14/2.22/2.66 3500/2318/1636/1577/1316

34 0.25 0.50 265 43 1.21 × 104 22.6 4500 1.00/1.15/1.25/1.35 4500/3930/3600/3333

changes in the CNT orientation are complementary retrieved by performing wide-angle X-ray scattering experiments and compared with measurements of the mechanical properties of the fibers. In this way the complex interplay between CNTs and the surrounding polymer structure can be revealed and correlated to the properties of the resulting fibers, allowing thus for a more directed tuning of the process conditions.

In order to investigate the effect of the CNT orientation with respect to the fiber axis, so far, conventional transmission electron microscopy (TEM),13,15,16 small- and wide-angle X-ray scattering (SAXS, WAXS),11,16−18 and polarized Raman spectroscopy14,16,19−21 have been employed. Since the majority of melt-spun polymer fibers form a periodic semicrystalline structure12,22 that gives rise to pronounced scattering contributions in the SAXS regime (i.e., form factor contributions from individual crystallites depending on their size and shape and contributions from structural correlations between individual crystallites) and the WAXS regime (i.e., crystal structure, size, and orientation) that superimpose with the comparatively weak scattering signal from the CNT carbon layers, determination of the CNT orientation is restricted to a limited number of cases, such as SWCNTs in a mainly amorphous polymer matrix.18 In addition, the shape of MWCNTs is irregular corresponding to worm-like objects23 rather than to stiff rods. Therefore, results from routinely performed azimuthal WAXS intensity scans around characteristic CNT scattering peaks may not exhibit significant differences even if the CNT orientation may in fact change considerably due to changes in the fiber production process. While the presence of a semicrystalline polymer structure does not affect the applicability of TEM and Raman spectroscopy, the irregular shape of MWCNTs can cause biases for a reliable determination of their orientation. A further complication arises from the strong attractive van der Waals forces between single carbon nanotubes that lead to the formation of CNT aggregates and thus a reduced homogeneous distribution of the CNTs within the polymer matrix. Within the aggregates, the CNTs are on average randomly oriented and do not interact with the surrounding polymer chains. Therefore, only the amount of well-dispersed CNTs is relevant in correlating the CNT orientation with structural, mechanical, and/or electroconductive properties of the reinforced fibers. Since Raman spectroscopy cannot distinguish between dispersed and aggregated CNTs, the separation of these contributions can only be achieved by visual inspection of the spatial distribution of the CNTs inside the fiber, i.e., by TEM investigations. In order to obtain the full 3D information on the shape and spatial configuration of the CNTs in the fiber, especially for determining the orientation of irregularly shaped MWCNTs, 3D acquisition of TEM micrographs and their tomographic reconstruction is required. In this work emphasis is laid on determining MWCNT orientation in melt-spun fibers produced at high take-up velocities and varying draw ratios using electron tomography. Structural changes in the polymer matrix that are induced by

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EXPERIMENTAL SECTION

Materials. Polyamid-6 (PA6) B24 N03 (melting temperature Tm = 220 °C, density d20 = 1.14 g/cm3)24 and polyethylene terephthalate PET 4048 (Tm = 250 °C, d20 = 0.9 g/cm3)25 were purchased from BASF AG, Germany, and Invista Resins & Fibers GmbH, Germany, respectively, and used as standard polymers for high-speed melt spinning of the semicrystalline polymer fibers. Multiwalled carbon nanotubes Nanocyl 7000 were received from Nanocyl S.A., Belgium, with an average diameter of 9.5 nm and a length of 1500 nm, yielding thus an aspect ratio of roughly 1:150.26 The two, for the melt-spinning process necessary, MWCNT-modified polymer compounds PET/ CNT and PA6/CNT were produced via melt-mixing by Nanocyl S.A., Belgium. Both custom-made polymer/CNT master batches (labeled as Plasticyl CUSTOM 10% NC 7000) contained MWCNTs at a total mass fraction of 10 wt %. Prior to fiber spinning, the corresponding polymers were added to obtain final CNT concentrations of 0.5 wt % for PET/CNT and 0.1 wt % for PA6/CNT. At CNT concentrations of 0.1 wt % for PA6/CNT and 0.5 wt % for PET/CNT stable high-speed melt-spinning of the respective composite fibers was possible. With increasing CNT concentration, the number and size of the CNT aggregates increase, leading to an unstable spinning process due to frequent rupture of the fibers at the positions of larger CNT aggregates. Melt Spinning. Melt spinning was carried out at the spinning plant ITA Plus, which allows a broad spinning parameter variation at industrial scale. Details on the working principle of the spinning facility are given by Steinmann and co-workers.27 Depending on the filament application, a wide variety of drawing rates can be realized before winding up the filaments on bobbins. The solid-state drawing of the fibers is performed between two godets, where the consecutive godet revolves faster than the first. In this study melt spinning was performed with inline drawing. The winding speed vw is kept constant while the take-up velocity vt of the first godet is varied; i.e., with increasing draw ratio λD = lw/lt (lt is the fiber length at the first godet and lw is the final fiber length at the second godet), the take-up velocity vt is decreased. For the stable production of the polymer/CNT composite fibers, draw ratios between 1.0 ≤ λD ≤ 1.35 at a winding speed of vw = 4500 m/min were used for the PA6/CNT compound, and for PET/CNT, the draw ratio range was extended to 1.0 ≤ λD ≤ 2.66, however, at a somewhat smaller winding speed of vw = 3500 m/min. All process parameters involved are summarized in Table 1. WAXS. X-ray scattering experiments on the fiber samples were carried out at the Institute of Crystallography (RWTH Aachen 5605


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Figure 1. Principle of the WAXS analysis shown for a PA6/CNT composite fiber sample. Left: 2D WAXS intensity distribution. For the analysis, radial 2Θ scans and azimuthal φ scans around the (200) Bragg reflection of γ-PA6 are performed. The γ-phase of crystalline PA6 is monoclinic with unit cell parameters a = 9.33 Å, c = 4.78 Å, b = 16.88 Å, and β = 121°.28 Arrows indicate the scan directions; broken lines mark the intensity data considered for the different scans. Center: 2Θ scan. Right: φ scan with intensity contributions from the crystalline and the amorphous part of the polymer. All intensity distributions are modeled by sets of Gaussians. and B0) are obtained from a least-squares fit of the model function (eqs 5 and 6) to the azimuthal intensity distribution (eq 4). The average squares of the cosines for the respective contributions are obtained from

University) using a single-crystal diffractometer STOE & Cie. IPDS II equipped with an image plate for digital readout. Molybdenum Kα radiation with an X-ray wavelength of λ = 0.710 73 Å was chosen for the experiments. The experimental setup as well as a detailed description of the specimen holder used for the experiments is given elsewhere.27 In the present study, WAXS was used to detect changes in the crystallinity, the crystallite dimensions, the correlation length of the amorphous polymer material, and the orientation distribution of crystalline and amorphous polymer. The principle of the data analysis is depicted in Figure 1. From the radial 2Θ scan, the position 2Θc of a crystalline Bragg reflection and the position of the amorphous correlation peak 2Θa are obtained. The crystallite dimension dc and the amorphous correlation length da are calculated via the Scherrer formula dc/a

λ = Δ2Θc/a cos(Θc/a)

π

2

⟨cos (φ)⟩c/a =

Ac Ac + A a

⎡3 1⎤ fc/a = − 2⎢ ⟨cos2(φ)⟩c/a − ⎥ ⎣2 2⎦

(2)

3 1 ⟨cos2(φ)⟩c/a − 2 2

(3)

requires the analysis of the azimuthal intensity distribution I(φ) =

∫2Θ I(2Θ, φ) sin(2Θ) d2Θ/∫2Θ sin(2Θ) d2Θ

(4)

which is obtained by averaging the 2D intensity distribution I(2Θ,φ) over a 2Θ range around the crystalline and amorphous peak positions. The integration range is chosen such that both crystalline and amorphous scattering contributions are correctly taken into account. The azimuthal intensity distribution is modeled by

I(φ) = Ic(φ) + Ia(φ) + B0

(5)

where Ic/a(φ) =

Ac/a 2π σc/a

exp[− (φ − φc/a)2 /(2σc/a 2)]

(7)

(8)

It should be noted that the method used for the analysis of the orientation is only an approximation, which can be applied for the case of a narrow orientation distribution. The full treatment for preferred orientation in polymer fibers is thoroughly discussed by Burger, Hsiao, and Chu.30 Electron Tomography. In order to retrieve 3D nanostructural information on the distribution of MWCNTs within the semicrystalline polymer fibers, bright-field transmission electron microscopy tomography was performed. Inspection by TEM requires electron transparent samples of a few hundred nanometers in thickness. For this purpose microtome cuts perpendicular and parallel to the fiber axis were produced using an EM UC6 ultramicrotome from Leica Microsystems, Germany. Tomographic tilt series were recorded with a 200 kV FEI Tecnai F20 field emission gun (FEG)-TEM equipped with a Gatan image filter. Acquisition of the tilt series was performed with a Gatan 916 single-tilt low profile sample holder allowing a maximum tilt angle of ±80°. Data acquisition and reconstruction was carried out using the Gatan Digital Micrograph tomography suite. To obtain the 3D data, a series of images in conventional bright-field TEM-mode were taken, while the sample was rotated around the main tilt axis within an angular range of ±60° at a step size of 1°. During the acquisition, the specimen drift was carefully monitored and corrected semiautomatically by opposing beam shifts or stage movements. Residual drift was removed by cross-correlating subsequent images prior to the reconstruction by weighted back-projection.31−36 The advantage of electron tomography compared to a conventional TEM analysis is shown in Figure 2. The projections from a rotation series around the fiber axis clearly demonstrates that from one single projection, incorrect CNT orientations are derived. The full 3D information on the spatial configuration of the CNTs is absolutely necessary to unambiguously determine their orientation relative to the fiber axis.

(1)

The determination of the Hermans orientation factor fc/a =

π

∫0 Ic/a(φ) sin(φ) dφ

Since an equatorial reflection is used for the determination of the orientation behavior, the Hermans orientation factors are given by29

where Δ2Θc and Δ2Θa are the full widths at half-maximum of the crystalline and amorphous peak, respectively. The crystallinity νc is computed from the ratio of the area of the crystalline diffraction peak Ac and the total area (Ac + Aa) in the 2Θ scan range υc =

∫0 cos2(φ)Ic/a(φ) sin(φ) dφ

(6)

describes the azimuthal contributions of the crystalline and amorphous polymer material, respectively, and B0 is a constant background scattering contribution. The seven model parameters (Ac/a, φc/a, σc/a, 5606


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Figure 3. Top: two 2D projections of a CNT in the x−y plane of the reconstructed sample volume at different depths z. Only the marked sections of the CNT lie in the respective x−y planes. Bottom left: 3D model of a set of six neighboring CNTs extracted from the investigated sample volume for a PET/CNT composite at a draw ratio of λD = 1.51 (the arrow indicates the orientation of the fiber axis f). Bottom right: definition of the simplified CNT shape using the end-to-end vector and the vector components parallel and perpendicular to the fiber axis for the determination of the orientation factor. Figure 2. Different projections of three MWCNTs (in a PA6/CNT composite, λD = 1.25) from a tilting series around the fiber axis after reconstruction and modeling of the nanotubes using the IMOD software package (the arrow indicates the direction of the fiber axis). For each projection, the apparent alignment of the CNTs relative to the fiber axis is different and demonstrates the importance of the 3D reconstruction for an unbiased determination of the CNT orientation.

Mechanical Properties. A Statimat 4U from Textechno GmbH, Germany, was used to determine the stress−strain curves of the PET fiber material with and without MWCNTs with a clamping length of 100 mm and an extension velocity of 100 mm/min under standard conditions. These curves are used for the calculation of the tensile strength and the elongation at break. The mechanical properties are compared with the structural data determined by the other experimental techniques and thus serve as an additional cross-check of the results obtained by WAXS and electron tomography.

Determination of the CNT Orientation from Electron Tomograms. Electron tomography on the fiber samples permits to trace each MWCNT inside the reconstructed volume (see Figure 3). A flexible cylinder-model of the CNT is constructed making use of the IMOD software package. Finally, the orientation with respect to the fiber axis is determined by defining a simplified cylindrical shape using the end-to-end vector of the CNT and by calculating the square of the cosine of the angle between the end-to-end vector and the fiber axis from the vector components df parallel and dr perpendicular to the fiber axis

cos2(φ) =

df 2 df 2 + dr 2

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RESULTS AND DISCUSSION The present study has been performed with the two polymer/ CNT nanocomposite fibers poly(ethylene terephthalate) (PET) containing 0.5 wt % CNTs and polyamide 6 (PA6) with a CNT content of 0.1 wt % (see Experimental Section) that have been produced by high-speed melt spinning at different draw ratios λD. In addition, unmodified PET fibers were spun under processing conditions similar to those used for the PET/CNT composites in order to identify the changes induced by the incorporation of the CNTs into the polymer melt and to estimate the nanotubes’ impact on the structural and mechanical fiber properties in the investigated processing window. For the analysis of the polymer structure of the PA6/CNT composite fibers using WAXS, radial 2Θ scans and azimuthal φ scans were carried out around the (2 0 0) Bragg reflection of monoclinic γ-PA6 (see Experimental Section, Figure 1). For the PET fibers, the region around the (0 1 0), (1 −1 0), and (1 0 0) reflections was examined by WAXS. Crystalline PET is triclinic with unit cell parameters a = 4.56 Å, b = 5.95 Å, c = 10.76 Å, α = 99.8°, β = 118.2°, and γ = 111.4°.38 Selected results from the WAXS experiments are shown in Figure 4.

(9)

The CNT orientation factor is computed according to eq 3 by averaging over the individual cos2(φ) values of all CNTs in the investigated sample volume Vsa = 1 × 1 × 0.5 μm3. Since the number of CNTs that can be found in the sample volume is quite small (in this study, the number of CNTs ranges between NCNT = 5 and NCNT = 16), the resulting average orientation factor is prone to be strongly biased by just one or two outliers. Therefore, a robust analysis37 is performed where the average Hermans factor is given by the median of the set of individual orientation factors, and the estimated standard deviation corresponds to the average deviation of the single values from the median. Thus, a reliable trend of the CNT orientation in dependence of the draw ratio is obtained without introducing a systematic bias arising from the orientation behavior of aggregated CNTs. 5607


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Figure 4. Selected two-dimensional WAXS intensity distributions in dependence of the draw ratio. Top row: unmodified PET. The region of maximum intensity corresponds to the (0 1 0), (1 −1 0), and (1 0 0) Bragg reflections of triclinic PET. Center row: PET/CNT composites containing 0.5 wt % CNTs. Bottom row: PA6/CNT composites containing 0.1 wt % CNTs. Figure 5. Part of the reconstructed sample volume of a PA6/CNT composite fiber (λD = 1.25) close to the cut surface of the sample in different projections. The grooves in the background originate from the ultramicrotome sectioning. The two nanotubes in the center of projection (a) are entangled and almost parallel aligned to the cut surface. The tube in the lower left of projection (a) penetrates the cut surface (see projections d−f) and extends itself into the fiber forming an S-shaped tube which can be seen in projections (b−e). Tubes penetrating the cut surface demonstrate their extraordinary mechanical properties, which enables CNTs to persist even ultramicrotome cutting with diamond knives.

The 2D intensity distributions of the unmodified PET and the PET/CNT composite fiber samples exhibit significant changes with increasing draw ratio and indicate that the properties of the polymer structure respond quite sensitive to changes in the processing conditions. Especially the amorphous rings become less pronounced with increasing draw ratio. Therefore, an overall increase in the orientation of the polymer material is expected. In contrast, the intensity distributions of the PA6/CNT composites remain unchanged; i.e., the PA6 polymer structure appears to be nearly independent of the draw ratio. The results of the WAXS analysis are discussed in the following two sections. A typical reconstructed fiber sample volume from the electron tomography investigations is depicted in Figure 5 that stems from the same tilting series as the three CNTs shown in Figure 2. The regular pattern of grooves in the background of the projections arises from the ultramicrotome cutting of the fiber sample. The CNTs in the reconstructed sample volume can be clearly separated from the polymer matrix due to their higher TEM-image contrast in comparison to the surrounding polymer material. It is, therefore, straightforward to model the respective 3D CNT pixel distributions by worm-like objects, to define the CNT cylinder axis, and to finally determine the CNT orientation factors relative to the fiber axis as described in the Experimental Section (Figure 3). PET Fibers Containing 0.5 wt % CNTs. The CNT orientation obtained from the analysis of the TEM tomograms (see Figure 6) exhibits anat first sight quite unusualSshaped behavior in dependence of the draw ratio λD. At low draw ratios the CNTs are found to be perpendicular aligned

with respect to the fiber axis (f CNT = −0.5). In the range 2.1 ≤ λD ≤ 2.3, a transition from perpendicular to parallel alignment is observed and appears to reach a constant value of f CNT = 0.87 at λD ≥ 2.66. While the orientation of the crystalline polymer (obtained from the analysis of the WAXS experiments) remains more or less unchanged ( fc = 0.93), the amorphous polymer chains become increasingly aligned parallel to the fiber axis with increasing draw ratio. Similarly, the crystallite dimension (Figure 7) increases almost monotonically with the draw ratio. In the transition zone, where the CNTs change their orientation, the crystallite sizes decrease slightly and grow again after the CNTs are aligned more parallel to the fiber axis. The effect of the changes in the CNT orientation appears to be much stronger for the amorphous polymer material. Here, a significant loss from da = 1.2 nm to da = 0.9 nm in the spatial correlation is observed when the CNTs start to change orientation. Once the CNTs are broadly aligned parallel to the fiber axis ( f CNT = 0.56 at λD = 2.22), the correlation length increases very 5608


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polymer material (bottom part of Figure 7). At a CNT orientation perpendicular to the fiber axis, the crystallinity is very low (around νc = 2%) and shows a slight increase with increasing draw ratio. In the transition zone, a “jump” in the crystallinity from νc = 3% at λD = 2.14 to νc = 7% at λD = 2.22 is observed. At first sight, a reasonable explanation of these results may be that at low draw ratios the shear forces in the spinning capillaries are larger than at high draw ratios. While the CNTs are aligned parallel to the fiber axis at low shear rates, their orientation changes to perpendicular alignment with increasing shear rates. The tendency to a so-called vorticity alignment of CNTs above a critical shear stress was already discovered by Hobbie et al.,39 when performing rheological investigations on semidilute polymer/CNT mixtures, and is consistent with previous experimental and theoretical studies on the orientation of nanoscopic fibers (MWCNTs, SWCNTs) in weakly elastic fluids.40−42 Since the Reynolds numbers for the PET and PA6 polymer melts are very small (Re ≈ 0.01 for PA6 and Re ≈ 0.07 for PET), the flow conditions for the extrusion are laminar. Therefore, the vorticity alignment of the CNTs under turbulent flow conditions, which would require much higher Reynolds numbers (Re ≥ 2000) is not a convincing explanation for the perpendicular alignment of the CNTs with respect to the fiber axis. However, a recently published study on the orientation behavior of wormlike micelles in solution passing a narrowed section when flowing through a capillary shows that, even under laminar flow conditions, the cylindrical particles align parallel to the flow direction when entering a narrowed capillary section (as expected) and change to perpendicular alignment after leaving the narrowed section (in contrast to common expectations).43 The change in orientation is explained by the formation of an extensional flow component perpendicular to the flow direction in the expansion zone behind the narrowed capillary section, arising from the change in the capillary geometry. This result suggests that the CNTs will a priori be aligned perpendicular to the fiber axis after passing the spinning nozzles and that a change of the CNT orientation requires drawing of the fiber material. The perpendicular alignment of the CNTs suppresses thus the formation of the semicrystalline polymer structure in fiber direction at low draw ratios. The decrease in the correlation length of the amorphous polymer during the transition in the CNT orientation from perpendicular to parallel alignment also indicates that the CNTs build up chemical bonds with the surrounding (amorphous) polymer matrix. Although the orientation of the amorphous material is overall increasing with increasing draw ratio, the polymer chains attached to the CNTs align parallel to the CNT axis rather than to the fiber axis,44 thus causing a loss in the spatial correlation. With increasing draw ratio, the uniaxial forces become sufficiently large to align the amorphous polymer chains and, hence, the CNTs parallel to the fiber axis. This behavior explains the “jump” in the amorphous correlation length as well as in the crystallinity once a threshold draw ratio is exceeded. The significance of these findings becomes more pronounced by comparing them to the results obtained for unmodified PET fibers produced with the same parameters at draw ratios λD = 1.31, 2.14, and 2.66 and a winding speed of vw = 3500 m/min (see Figure 8). Without CNTs, the PET crystallinity is a priori much higheralmost by about 1 order of agnitude at low draw ratio

Figure 6. Hermans orientation factors f of CNTs (derived from electron tomography) and crystalline as well as amorphous PET orientation (determined by WAXS) in melt-spun PET/CNT composite fibers containing 0.5 wt % CNTs processed with different draw ratios λD. A sharp transition in the CNT orientation from an alignment perpendicular to the fiber axis to parallel alignment occurs at draw ratios 2.1 ≤ λD ≤ 2.3. The lines through the experimental data points are guides to the eye.

Figure 7. Polymer crystallite size and amorphous correlation length determined by WAXS experiments (top) and crystallinity of the polymer material (bottom) in melt-spun PET/CNT composite fibers vs draw ratio λD. The region where the CNTs change orientation is marked by the vertical dashed lines. The lines through the experimental data points are guides to the eye.

rapidly to da = 1.8 nm and does not change significantly at an even higher draw ratio of λD = 2.66. Even more impressive is the effect of the CNT orientation on the crystallinity of the 5609


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Figure 8. Top: Hermans orientation factors for crystalline as well as amorphous polymer orientation in unmodified PET fibers vs draw ratio λD. Center: corresponding polymer crystallite size and amorphous correlation length. Bottom: crystallinity of the unmodified polymer material. The region where the CNTs change orientation in the PET/CNT composite fiber is marked by the vertical dashed lines. The lines through the experimental data points are guides to the eye.

Figure 9. Results from the mechanical analysis of melt-spun PET/ CNT composite and unmodified PET fibers vs draw ratio λD. Top: tensile strength at break. Center: elongation at break. Bottom: tensile energy absorption at break. The lines through the experimental data points are guides to the eye.

(νcPET/νcPET/CNT ≈ 9.5 at λD = 1.31). In comparison to the PET/CNT composite, the relative change in the crystallinity of roughly 16% over the entire explored λD range is small. While the results for the crystallite and amorphous orientation factors are well comparable to those of the PET/CNT composite fibers, the crystallite sizes as well as the amorphous correlation lengths in the unmodified PET fibers are significantly larger (dcPET/dcPET/CNT ≈ 3.8 and daPET/daPET/CNT ≈ 4.5 at λD = 1.31). A loss in spatial correlation of the amorphous polymer material isdue to the absence of CNTsnot observed. The mechanical analysis of the PET fibers with and without CNTs is depicted in Figure 9 and confirms the reliability of the results obtained by electron tomography and WAXS. At low draw ratios, the tensile strength at break F is about 1.5 times higher for the unmodified PET fibers than for the PET/CNT composites, which is explained by their higher crystallinity. In

contrast, the elongation at break ε is larger for the CNT modified PET fibers, which is caused by the low crystallinity and the perpendicular alignment of the CNTs at low draw ratios. With increasing λD, the CNTs become aligned and the elongation at break decreases steeper than for the unmodifed PET fibers. Although the mechanical properties of the PET/ CNT composites are not improved in comparison to the unmodified PET fibers, the reinforcement effect of the CNTS is clearly visible. Since the crystalline fraction in the umodified PET fibers is 10 times larger, a similar difference in the tensile strength in comparison to the PET/CNT fibers should be expected, which is, however, not observed here. Concerning the 5610


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tensile strength of the fibers, the much lower crystallinity of the PET/CNT composites is almost compensated by the relatively small amount of 0.5 wt % CNTs. The combination of the two mechanical parameters W = F × ε, which is proportional to the tensile energy absorption at break, reflects the dependence of the mechanical properties of the PET/CNT fibers on the CNT orientation. For the CNT modified PET fibers, a steep decrease in the W parameter is observed, which changes its behavior in the λD region where the CNTs change their orientation from perpendicular to parallel alignment, and finally remains constant if the CNTs are oriented parallel to the fiber axis. In comparison, the tensile energy absorption at break of the unmodified PET fibers decays exponentially. PA6 Fibers Containing 0.1 wt % CNTs. Qualitatively, the CNT orientation in the melt-spun PA6/CNT composite fibers shows the same behavior as in the PET/CNT fibers (Figure 10). The change in the CNT orientation, however, starts at a

transition zone from perpendicular to parallel CNT alignment around f CNT = 0 (Figure 11).

Figure 10. Hermans orientation factors f of CNTs, and crystalline as well as amorphous PA6 orientation in melt-spun PA6/CNT composite fibers containing 0.1 wt % CNTs at different draw ratios λD. The transition in the CNT orientation from perpendicular to parallel alignment relative to the fiber axis occurs at draw ratios 1.1 ≤ λD ≤ 1.3. The lines through the experimental data points are guides to the eye.

Figure 11. Structural behavior of the PA6 polymer in PA6/CNT composite fibers vs draw ratio λD. Top: polymer crystallite size and amorphous correlation length determined by WAXS experiments. Bottom: crystallinity of the polymer material. The region where the CNTs change orientation is marked by the vertical dashed lines. The lines through the experimental data points are guides to the eye.

comparatively low draw ratio of λD = 1.15, which may be attributed to the polymer material, the different processing conditions (such as the extrusion temperature, extrusion velocity, and winding speed), and/or the smaller amount of only 0.1 wt % CNTs in the PA6/CNT composite. The orientation factors of the amorphous polymer as well as the crystalline PA6 fraction are higher than in the CNT modified PET fibers. These results can be expected due to the lower CNT content in the PA6/CNT composite. Moreover, a slight decrease in the orientation of the amorphous PA6 from fa = 0.83 to fa = 0.78 is observed when the CNTs start to change orientation. This result supports the hypothesis that the CNTs build up chemical bonds with the amorphous polymer. For the PET/CNT composites (Figure 6), this effect is probably blurred by the strong dependence of the amorphous polymer orientation on the draw ratio. As for the PET/CNT composites, the correlation length of the amorphous PA6 polymer material is decreasing when the CNTs change from perpendicular to parallel alignment which is in full agreement with the orientation behavior of the amorphous PA6 polymer. Surprisingly, the PA6 crystallinity appears to be nearly independent of the draw ratio in the

This result is different to the behavior of the PET/CNT composites and shows that the formation of the crystalline polymer structure is not only dependent on the CNT orientation but is substantially influenced by the total amount of CNTs incorporated into the polymer melt. Therefore, the complex interplay between fiber processing conditions, CNT content, CNT orientation, and polymer structure may serve as an explanation for the various results reported in the literature concerning the reinforcement effect of CNTs in polymer fibers.

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CONCLUSIONS In this study the orientation behavior of well-dispersed multiwalled carbon nanotubes within polymer/CNT composite fibers produced by high-speed melt spinning has for the first time been addressed using bright-field electron tomography. The performed 3D reconstructions from the tomographic tilt series allow tracking the trajectories of individual CNTs within the fiber volume, thus determining their orientation parameters independent from additionally present CNT aggregates. For both polymer/CNT composite fibers investigated, a so far unprecedented transition region in the CNT orientation could be assigned as a function of the draw ratio applied during the 5611


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fiber production process, in which the nanotubes change from perpendicular to parallel orientation with respect to the fiber axis. Complementary performed wide-angle X-ray scattering experiments and mechanical analysis of the polymer/CNT composite fibers disclose that the change in the CNT orientation is moreover apparently reflected in the structural as well as the mechanical properties of the polymer composite fibers. While the semicrystalline polymer structure in the polymer/CNT composite fibers is less developed compared to the unmodified polymer fibers, the crystallinity of the CNT loaded fibers changes nonmonotonically within the assigned CNT orientation transition region, which is indicative for a cooperative crystallization process of the polymer chains. During fiber processing, folded polymer chains are generally pulled out in fiber direction and crystallites form from elongated and aligned unfolded chains.12 Therefore, the mere presence of the CNTs in the polymer melt seems to perturb the growth of the crystalline polymer chain domains rather than enhancing the crystallization process. This finding may explain the varying results reported in the literature concerning the reinforcement effect of CNTs in polymer materials. At low draw ratios, when the dispersed CNTs are found to orient perpendicular to the fiber axis, the formation of the semicrystalline polymer structure is restrained. At high draw ratios, CNT alignment parallel to the fiber axis is observed, thereby leading to an increase in the crystallinity of the polymer material. A strong correlation with the CNT orientation could concomitantly be shown for the correlation length of oriented amorphous polymer chains that first decreases and then increases, while the CNTs align parallel to the fiber axis. In addition, mechanical analysis of the fibers reveals that the CNT modified polymer fibers exhibit a lower tensile strength and larger elongations at break. The minimum in the elongation, which is even more pronounced in the tensile energy absorption, cannot be explained by a suppressed crystallinity due the presence of the CNTs. The fact that the measured structural and mechanical properties of the polymer matrix undergo significant changes in the transition region from perpendicular to parallel CNT alignment hints at the existence of chemical bonds between the CNTs and the surrounding (amorphous) polymer matrix. A deeper insight into the conditions defining the CNT orientation transition region and the trade-off between CNT content and polymer crystallization that would permit a directed fine-tuning of the structural and mechanical fiber properties asks not only for more experimental but also for simulation studies.45

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected] (T.V.). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS

The authors thank Mr. Christopher Holub for assistance with the WAXS measurements. Financial support by the German Research Foundation (DFG), project Fiber Sage (Grants GR1311/34-1 and WE 2579/3-1), is gratefully acknowledged. 5612


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Orientation of Well-Dispersed Multiwalled Carbon Nanotubes in Melt

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