Investigation of crystalline and tensile properties of carbon nanotube-filled polyamide-12 fibers melt-spun by industry-related processes

Uppsala University

This is an accepted version of a paper published in Journal of Engineered Fibers and Fabrics. This paper has been peer-reviewed but does not include the final publisher proof-corrections or journal pagination.

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Chatterjee, S., Reifler, F., Chu, B., Hufenus, R. (2012)

"Investigation of crystalline and tensile properties of carbon nanotube-filled polyamide-12 fibers melt-spun by industry-related processes"

Journal of Engineered Fibers and Fabrics, 7(3)

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Investigation of crystalline and tensile properties of carbon nanotube-filled polyamide-12 fibers melt-spun by

industry-related processes

S. Chatterjee

, F A Reifler

, B T T Chu

, R Hufenus

a

Laboratory for Functional Polymers, Swiss Federal Laboratories for Materials Science and Technology (Empa), Überlandstrasse 129, 8600 Dübendorf, Switzerland.

b

Laboratory for Advanced Fibers, Swiss Federal Laboratories for Materials Science and Technology (Empa), Lerchebfeldstrasse 5, 9014 St Gallen, Switzerland

* sanjukta.chatterjee@empa.ch

ABSTRACT

The paper addresses the influence of carbon nanotubes (CNT) on the structure and mechanical properties of high tensile strength thermoplastic polymer fibers. Polyamide (PA) fibers with different draw ratios, with and without CNTs as fillers, and having mechanical properties close to industrial standards were spun in a pilot melt spinning plant.

The morphology of the fibers was investigated using optical microscopy, nuclear magnetic resonance (NMR) and 2-D wide angle x-ray diffraction (WAXD). Differential scanning calorimetry (DSC) was carried out to get an estimation of the crystallinity. For a concise interpretation of the results of tensile measurements performed on the fibers, a parameter was developed to account for the detrimental influence of polymer extrusion on their mechanical properties. CNTs seem to act as sites for the growth of un-oriented crystalline domains converted from oriented regions, without yielding a mechanical reinforcing effect.

INTRODUCTION

Owing to their exceptional strength as well as their electrical, thermal and electronic transport properties, carbon nanotubes (CNTs) have attracted wide attention for the last decade

. The influence of CNTs as fillers in polymer matrices is of great interest

. Several methods of polymer-CNT composite manufacture have been used, for example melt mixing

, solution mixing

and in situ polymerization

. For industrially relevant scales it is advantageous to use melt compounding, a very effective method for thermoplastic polymers.

Previous studies on mechanical properties of polymer

carbon nanotube composite fibers have shown varied results. With the inclusion of CNTs, Pötschke et al.

reported an increase in Young’s modulus for polycarbonate, whereas Bhattacharyya et al.

found that the mechanical properties of polypropylene fibers are mostly unaffected with the presence of single wall CNTs.

The tensile strength of CNT filled PA fibers melt-

spun at laboratory scale is usually much lower than

the tensile strength of industrial grade PA fibers

.

Spinning parameters are key factors determining the

quality of a fiber

. Hence, we used our fiber pilot

melt-spinning plant developed for spinning fibers in

industrially relevant scales with properties

comparable to commercial fibers. However, strong

van der Waals interactions amongst CNTs and their

high aspect ratio make it extremely difficult to

uniformly disperse CNTs in polymer matrices. CNT

agglomeration has proven to be a detrimental factor

for the mechanical properties of composites

.

Polyamide 12 (PA12) can exhibit domains consisting

of one or both of two polymorphic phases, namely α

and γ, as well as amorphous regions

. In the α phase,

the polymer chains are oriented parallel to each other

and the H-bond is in plane, whereas in the γ phase the

chains run anti-parallel resulting in the H-bond to be

twisted out of the plane

. The polymorph obtained

by quenching from the melt and subsequent

crystallization is denoted as γ´. It has a structure

similar to that of the γ phase, giving rise to similar

WAXD patterns. Hence, it can only be differentiated

from it by NMR

. Degree of crystallinity and

crystalline orientation act as governing factors to

determine the tensile strength of the fibers. Inclusion

of CNTs can bring about changes in the crystalline

structure of the polycrystalline matrix, e.g. they can

act as nucleation sites for the growth of new crystallites

. The relationship between the structure and the physical properties of these fibers is of primary interest

. In this paper we investigate how these important factors are influenced by changes in draw ratio (DR) and by the incorporation of CNTs into the fibers.

EXPERIMENTAL DETAILS

For fiber spinning, PA12 pellets (Grilamid L16 from EMS-GRILTECH, Switzerland, Mn 15,000-19,000) were used. An industrial grade masterbatch (Plasticyl PA1502 from Nanocyl SA, Belgium) comprising 15% multi walled CNTs of 90% purity in PA12 (Grilamid L16) was used for the melt-spinning of the CNT-filled fibers.

Unfilled PA12 fibers with DR 3, 4, and 4.5 were melt-spun from virgin PA12 with the in-house pilot melt-spinning plant

. PA12 together with the CNT masterbatch was extruded in a twin-screw extruder to obtain compounds with different CNT concentrations. These compounds were used to spin a set of PA12 fibers with varying CNT concentrations:

0.003 wt%, 0.0075 wt%, 0.015 wt%, 0.03 wt%, 0.075 wt% and 0.15 wt%, each with DR 3, 4 and 4.5. Some of the as-spun fibers were annealed for 3 hours at 160°C and allowed to cool slowly.

Tensile testing was carried out on the Tensorapid 3 (UTR3) tensile tester (Uster Technologies, Uster, Switzerland) with 500 N load cell, using testing standard ISO 2062:2009

. Ten single filaments of each fiber type were measured with a test length of 250 mm at a constant rate of extension of 250 mm/min and a preload of 0.5 cN/tex. Optical microscopy was performed using an Olympus SZX16 microscope to visualize the CNT dispersions in the fibers. WAXD measurements were done at the beamline I 711 of MAX-lab, Sweden, with an X-ray wavelength of 1.1Å. For WAXD measurements, a bundle containing 20 filaments was used and each diffractogram was recorded twice. Solid state

C NMR was carried out using a Bruker AVANCE-400 MHz NMR Spectrometer at 100.61 MHz with a 7 mm CP-MAS probe at MAS rates of 3500 Hz. For solid state NMR a bunch of fibers was chopped into very short pieces and measured. For the DSC measurements, the fibers were cut into small pieces.

DSC was performed using samples with an average sample weight of 7 mg on a Mettler DSC822 instrument in the temperature range of 25°C to 250°C with a heating rate of 20°C/min in nitrogen atmosphere.

FIGURE 1: Optical micrographs of fibers (DR 3) with CNT content of a) 0.0075 wt% b) 0.003 wt% c) 0.015 wt% d) 0.030 wt%.

RESULTS AND DISCUSSIONS Optical Microscopy

Figure 1 shows optical micrographs of some CNT filled fibers with sites of CNT agglomeration. Such agglomerations have negative effects on the mechanical properties of the fibers

.

Nuclear Magnetic Resonance (NMR)

The chemical shifts and line widths of individual resonances were determined by non-linear least- square fits of a sum of Gaussian/Lorentzian curves using the software of Massiot et al.

. The chemical shifts of the unfilled as-spun fibers (Table I) correspond to the chemical shifts reported in the literature

for the γ´ polymorphic phase; the same result was also found for the CNT filled fibers. For the annealed fibers, the chemical shift refers to the γ form. This is in accordance with the literature, where it can be seen that PA12 quenched from the melt at atmospheric pressure is prone to crystallize in the γ´

phase instead of α or γ phase, and that on annealing above 110°C at atmospheric pressure, a γ´ to γ transformation can take place

.

TABLE I:

C chemical shifts (in ppm) for the unfilled PA12 fibers compared to literature values

. Sampl

e

CH

gauc he

CH

all trans

C

C

C

C=O

PA12 fiber

30.9 33.1 37.9 27.3 40.4 173.6 Annea

led PA12

30.5 33.5 37.0 26.9 39.9 173.0

γ

30.6 33.5 37.1 ≈ 28 40.0 173.7 γ´

30.8 33.2 37.0 ≈ 27 40.6 173.7 α

31.2 34.3 38.7 27.3 42.4 172.8

Crystallinity and Orientation

WAXD is an extremely effective tool to deduce information about the morphology and crystalline structure, e.g., the crystalline polymorphic phases present and the percentage crystallinity in the samples. The 2D WAXD pattern of the fibers showed meridional and equatorial reflections as seen in Figure 2. These peaks can be attributed to the γ

and the γ

planes, respectively

.

FIGURE 2: 2-D WAXD images: a) undrawn unfilled fiber showing less orientation b) fiber with 0.03 wt% CNT and DR 3.

The WAXD data were processed using Version 4.1.

of the XRD2DScan displaying and analyzing Software (A. Rodriguez Navarro; Universidad de Granada, Granada, Spain). For the equatorial γ200 peak, the 1D intensity vs. 2Θ plot was deduced and de-convoluted using MATLAB programming to separate the crystalline and the amorphous parts. The crystalline peak was fitted with Lorentzian line shape as seen in Figure 3. For the amorphous region the original line shape was taken from that part of the spectrum which has no contribution to the equatorial crystalline peak. The equatorial crystallinity index (ECI) was calculated as:

(1)

and the Herman’s orientation factor (HOF) for the equatorial γ200 peak was calculated using the equation

(2)

where

(3)

ϕ is the azimuthal angle between the reference direction and the crystallographic axis. I

( ϕ) ^{is the} scattered intensity of the hkl plane in the direction ϕ .

FIGURE 3: De-convolution of the equatorial peak with fitted Lorentzian curves, shown for the fiber with 0.03 wt% CNT and DR 3. The peaks are assigned as (1) original curve (2) fitted peak 1 (unassigned; necessary to improve the fitting) (3) fitted peak 2 (2Θ=15.1°, d=4.19 Å, γ200) (4) fitted un-oriented crystalline peak (2Θ=15.3°, d=4.13 Å, γ200) (5) amorphous peak (from original line shape) (6) fitted background (7) total fit which overlaps with the original curve.

For all unfilled and CNT-filled PA12 fibers, the position of the fitted equatorial Lorentzian peaks showed a variation in the 2Θ peak position from 14.9° to 15.3°. This corresponds to interplanar d spacings ranging from 4.11 Å to 4.24 Å. Thus, these peaks are attributed to γ

.

The HOF values for the equatorial γ200 peak of the unfilled fibers are 0.19, 0.21 and 0.23 for DR 3, 4 and 4.5, respectively. The HOF value of the same peak in the CNT filled fibers varies from 0.10 to 0.21, 0.11 to 0.18 and 0.17 to 0.25 for the fibers with DR 3, 4 and 4.5, respectively. Hence, as expected, there is a tendency of higher orientation for higher DR.

The ECI of the unfilled fibers vary from 58% to 67%,

whereas the ECI of the CNT filled fibers vary in the

range of 60% to 68%. The ECI is largely influenced

by the DR, but for fibers with the same draw ratio,

irrespective of their CNT content, the ECI is very

close. However, the ECI values do not represent the total crystallinity of the polymer as only the equatorial crystalline peak is taken into account during the calculation. This value is influenced by the orientation of the crystallites

and therefore should not be directly correlated with the percentage crystallinity calculated from DSC measurements. An interesting phenomenon observed is the occurrence of an un-oriented crystalline peak for fibers filled with 0.03 wt% and 0.075 wt% CNT at DR 3 and 4 (Figure 3). This may be due to the fact that CNTs act as sites for the formation of un-oriented crystallites, as previously reported for PA6

. The CNTs are curved and randomly oriented in the matrix, thus the crystalline domains centered on them are un-oriented.

The fact that we do not see an increase in the crystallinity for CNT filled fibers suggests an interchange of crystalline phases (from oriented to un-oriented) instead of the formation of new crystallites. A possible explanation would be that the CNTs inhibit the orienting process of the polymer lamellae. However, at high draw ratios (DR 4.5), the polymer chains get oriented, no longer exhibiting the un-oriented crystalline peak.

Differential Scanning Calorimetry (DSC)

From the DSC data, the percentage crystallinity can be calculated as:

(4)

where ∆H is the enthalpy of fusion (area under the endotherm) of the sample, ∆H

is the enthalpy of fusion for a 100% crystalline PA 12 which was taken to be 209.34 J/g

. The factor F denotes the fraction of polymer present in the composite.

For the unfilled PA12 fibers the percentage crystallinity increased with draw ratio from 27.5%

for the undrawn fiber to 31.8% for DR 4. For the CNT filled fibers the percentage crystallinity varied from 27.6 % to 31.7 % depending rather more on the DR than on the CNT content. Fibers with the same DR manifest similar melting curves and the values of crystallinity do not depend on the CNT concentration. As also seen with WAXD analysis (ECI) we do not see an increase in crystallinity for samples with un-oriented crystallites, thus making it more probable that instead of the creation of new crystalline domains there are changes in the existing structure occurring.

Mechanical Properties

It has to be pointed out that the following analysis applies to fibers melt-spun with DR 4.5. We observe

that extrusion has a crucial influence on the mechanical properties of melt-spun fibers, as shearing forces during extrusion cause damage in the polymer chains, diminishing molecular weight. To dilute and mix the CNT masterbatch the polymer had to be twin-screw extruded multiple times, and a combination of extrusion and CNT incorporation seems to be detrimental to the tensile strength. We have defined a factor called "extrusion factor" (Ext), to quantify this effect. Virgin PA12 was extruded once and twice and the specific tensile strengths of the resulting fibers were measured as 55.4 cN/tex, and 41.7 cN/tex, respectively. The ratio of these two values (1.33) gives us an approximation of the damage as reflected in the tensile strength. Thus we assume that Ext for the polymer extruded once (Ext1) is 1 and Ext of the polymer extruded twice (Ext2) is 1.33.

To achieve a CNT concentration of 0.15 wt% in the as-spun fibers, a compound (Comp15) comprising 99 wt% virgin PA12 and 1 wt% CNT masterbatch was produced. This once extruded compound was melt- spun as such, but it was also used to prepare compounds with CNT concentrations of 0.075, 0.030 and 0.015 wt%, mixing it with virgin PA12. The respective extrusion factors were calculated as: Ext = fraction of virgin polymer × Ext1 + fraction of Comp15 × Ext2. The compound Comp015 with 0.015 wt% CNT consisted of 90 wt% virgin PA12 and 10%

Comp15; the corresponding extrusion factor can be calculated as Ext3 = 0.9 × Ext1 + 0.1 × Ext2 = 1.03.

This compound was also melt-spun as such and used to prepare compounds with CNT concentrations of 0.0075 and 0.003 wt%, mixing it with virgin PA12.

The respective extrusion factors were calculated as:

Ext = fraction of virgin polymer × Ext1 + fraction of Comp015 × Ext3. Table II summarizes the extrusion factors.

In order to characterize the combined effect of

extrusion and CNT incorporation ("strain factor"),

several mathematical combinations of the effects of

extrusion and CNT concentration have been

evaluated. The best fit (coefficient of determination

factor R² = 0.95) for the specific tensile strength

plotted against this strain factor (Str) could be

achieved with the linear combination Str = Ext + f ×

CNT concentration (Figure 4), with the weighting

factor for the CNT concentration in wt% being f =

0.37. Table II summarizes strain factors and specific

tensile strengths of the fibers with various CNT

concentrations. It clearly shows the influence of the

CNT incorporation for fibers with a low extrusion

factor; for higher extrusion factors, the influence of

the extrusion prevails. The tensile strength is seen to

be minimal for the fiber with 0.075 wt% CNT, which has both an elevated value in terms of extrusion factor and CNT concentration. As molecular orientation is crucial for the mechanical strength of fibers, even though new crystallites are created in CNT filled fibers, their un-oriented nature does not contribute to fiber reinforcement.

TABLE II: Summary of extrusion factors, strain factors and specific tensile strengths (with standard deviations) for the fibers with DR 4.5.

CNT (%)

Extrusion factor (Ext)

Strain factor (Str)

Specific tensile strength (cN/tex)

0 1.000 1.000 55.3±1.6

0 1.329 1.329 41.6±2.5

0.150 1.000 1.056 45.0±3.4 0.075 1.164 1.192 35.5±3.4 0.030 1.066 1.077 44.3±1.9 0.015 1.033 1.038 45.1±2.9 0.003 1.007 1.008 51.0±3.4 0.0075 1.016 1.019 49.5±1.7

FIGURE 4: Fit showing specific tensile strength in relation to the strain factor Str for fibers melt-spun at draw ratio 4.5 with different CNT concentrations.

CONCLUSIONS

The PA12 fibers with and without CNTs crystallize in the γ and γ' phase depending on the pretreatment, without any trace of the α phase. Crystallinity studies show that the DR has a more pronounced effect than

the inclusion of the CNTs. However, CNTs seem to act as sites for the development of un-oriented crystallites converted from the oriented phases, as detected by WAXD analysis. Virgin PA12 is a thermoplastic polymer that renders very strong synthetic fibers, but treatments like multiple extrusions or inclusion of nanofillers seem to have detrimental effects on the tensile strength of such fibers.

ACKNOWLEDGEMENTS

Authors would like to sincerely thank Daniel Rentsch for NMR, Pierluigi Barbadoro for extrusion, Laura Gottardo and Benno Wüst for melt-spinning, and Tomás Plivelic for his scientific inputs. Financial support from the Swiss National Science Foundation (SNSF) is thankfully acknowledged.

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