Advertisement

Characterizations of Nanocomposites of Liquid Crystalline Polymers

  • Tae Young Ha
  • Yong-Ho Ahn
  • Bo-Soo Seo
  • Donghwan Cho
  • Jin-Hae ChangEmail author
Living reference work entry

Latest version View entry history

  • 15 Downloads
Part of the Polymers and Polymeric Composites: A Reference Series book series (POPOC)

Abstract

Nanocomposites of three thermotropic liquid crystalline polymers (TLCPs) with organoclay were prepared. The first TLCP, poly(2-ethoxyhydroquinone-2-bromoterephthaloyl), EHBT, consists of wholly aromatic ester type mesogenic units containing an ethoxy side group, and the second poly(oxybiphenyleneoxy-2,5-dihexyloxyterephthaloyl) (OBDT) is an aromatic polyester TLCP having alkoxy side groups on the terephthaloyl moiety. The last TLCP polyazomethine (PAM) consists of diad aromatic azomethine type mesogenic units. An EHBT with an alkoxy side-group was synthesized from 2-ethoxyhydroquinone and 2-bromoterephthalic acid. Nanocomposites of EHBT with Cloisite 25A (C25A) as an organoclay were prepared by the melting intercalation method above the melt temperature (Tm) of the TLCP. Liquid crystallinity, morphology, and thermo-mechanical behaviors were examined with increasing organoclay content from 0 to 6 wt%. Liquid crystallinity of the C25A/EHBT hybrids was observed when organoclay content was up to 6 wt%. Regardless of the clay content in the hybrids, the C25A in EHBT was highly dispersed in a nanometer scale. The hybrids (0–6 wt% C25A/EHBT) were processed for fiber spinning to examine their tensile properties. Ultimate strength and initial modulus of the EHBT hybrids increased with increasing clay content and the maximum values of the mechanical properties were obtained from the hybrid containing 6 wt% of the organoclay. A TLCP (OBDT)/organoclay nanocomposite was synthesized via in-situ intercalation polycondensation of diethyl-2,5-dihexyloxyterephthalic acid and 4,4′-biphenol in the presence of organically modified montmorillonite (MMT). The organoclay, C18–MMT, was prepared by the ion exchange of Na+–MMT with octadecylamine chloride (C18–Cl). OBDT/C18–MMT nanocomposites were prepared to examine the variations of the thermal properties, morphology, and liquid crystalline phases of the nanocomposites with clay content in the range 0–7 wt%. It was found that the addition of only a small amount of organoclay was sufficient to improve the thermal behavior of the OBDT hybrids, with maximum enhancement being observed at 1 wt% C18–MMT. Nanocomposites of PAM with the organoclay C12-MMT were also synthesized by using the in-situ interlayer polymerization method. The variations with organoclay content of the thermal properties, morphology, and liquid crystalline mesophases of the hybrids were determined for concentrations from 0 to 9 wt% C12-MMT. The wide-angle X-ray diffraction (XRD) analysis and transmission electron microscope (TEM) micrographs show that the levels of nanosize dispersion can be controlled by varying the C12-MMT content. The clay particles are better dispersed in the matrix polymer at low clay contents than at high clay contents. With the exception of the glass transition temperature (Tg), the maximum enhancement in the thermal properties was found to arise at an organoclay content of 1 wt%. Further, the PAM hybrids were shown to exhibit a nematic liquid crystalline phase for organoclay contents in the range 0–9 wt%.

Keywords

Thermotropic liquid crystalline polymers Nanocomposite Organoclay Montmorillonite Intercalation method 

Definition

TLCPs having an alkyl substituent with a base on a nematic liquid crystalline phase were reviewed. We also examined the correlation between the thermal properties and clay contents of TLCP nanocomposites and the dispersed morphology of the clay particles. Morphological studies showed that the levels of nanosize dispersion can be controlled by varying the organoclay content. The addition of only a small amount of organoclay was found to be sufficient to improve the thermal behavior of the TLCP hybrids.

Introduction

Thermotropic liquid crystalline polymers (TLCPs) have already been established as high-performance commercial engineering polymers owing to their specific chemical structures, high strengths, high moduli, low viscosities, and other good mechanical properties. The structure–property relationships of TLCPs have been the subject of much research. Despite their inferior physical strength when compared with lyotropic liquid crystalline polyamides, TLCPs are attracting a great deal of interest based on their melt processability (Baird and Sun 1990; Lusignea 2001).

Despite the increasing interest and numerous papers on composites with TLCPs, most studies have treated only wholly aromatic rigid-rod-type polymers; less attention has been paid to the processing and properties of TLCPs containing substituents, side group structures, or a flexible alkyl group in the main chains. Although wholly aromatic rigid-rod-type TLCPs exhibit very attractive mechanical properties, they generally have high melting points, which can give rise to difficulties in processing. The presence of substituents or side groups in main chain structures in otherwise aromatic polyesters lowers their melting points, thus widening the processing window. Thus, despite the expected loss of thermomechanical properties that results from their use (compared to the properties of wholly aromatic TLCPs), they may have considerable advantages in particular applications (McArdle 1989).

The recent development of nanocomposites was stimulated by work with nylon 6/clay by Toyota (Kojima et al. 1994; Usuki et al. 1995). Progress, especially with respect to nanocomposite formation, has been reviewed by many researchers (Giannelis 1996; Yang et al. 1998). Nanocomposites are one of the most important classes of synthetic engineering materials. Their makeup is such that they can be transformed into new materials possessing the advantages of both organic materials, such as light weight, flexibility, and good moldability, and inorganic materials, such as high strength, heat stability, and chemical resistance. The incorporation of organic/inorganic hybrids can yield materials possessing excellent stiffness, strength, and gas barrier properties with far less inorganic content than is used in conventionally filled polymer composites: the higher the degree of delamination in polymer/clay nanocomposites, the greater the enhancement of these properties. Nanocomposites such as these can also be employed as scratch- and abrasive-resistant hard coatings, nonlinear optical materials, and reinforcements for elastomers and plastics (LeBaron et al. 1999; Vaia 2000).

Nanofillers can be categorized on the basis of their dimensions, such as one-dimensional fillers (nanotubes and nanowires), two-dimensional fillers (clays and graphene), and three-dimensional fillers (spherical and cubic nanoparticles). Among them, clays have sandwich structures with one octahedral Al sheet and two tetrahedral Si sheets, the so-called phyllosilicates. There are many types of phyllosilicates, including kaolinite, montmorillonite (MMT), hectorite, saponite, and synthetic mica. These clays consist of stacked silicate sheets with lengths of about 46 nm for hectorite, 170 nm for saponite, 218 nm for MMT, 1230 nm for synthetic mica, and so on. They have the same sheet thickness of 1 nm. Therefore, the only difference among them is the length of the silicate sheets (Yano et al. 1997; Garcia et al. 2000).

To overcome problems of macro- and micro-phase separation between organic polymers and inorganic clays, organic/inorganic polymer hybrids have mostly been synthesized using three methods: solution intercalation, melt intercalation, and in-situ intercalation polymerization. Additionally, other approaches, such as the sol–gel process and monomer/polymer grafting to clay layers, have resulted in organic/inorganic polymer hybrids (Ishida et al. 2000; Shen et al. 2002).

Solution intercalation is based on a solvent system in which the polymer is soluble and the clay layers swell. The layered clay is first swollen in a solvent, such as N,N′-dimethylacetamide (DMAc). When the polymer and clay solutions are mixed, the polymer chains intercalate and displace the solvent from between the layers of clay. Upon solvent removal, the intercalated structure remains, resulting in hybrids with nanoscale morphology. In the process of melt intercalation, the layered silicate is mixed with a molten polymer matrix. If the silicate surfaces are sufficiently compatible with the chosen polymer, the polymer can enter the interlayer space and form an intercalated or an exfoliated nanocomposite. Finally, in-situ intercalation polymerization is based on the use of one or more monomers that may be in-situ linearly polymerized or crosslinked and was the first method used to synthesize polymer-layered silicate nanocomposites based on nylon 6. In situ intercalation relies on swelling of the organoclay due to the monomer, followed by in-situ polymerization initiated thermally or by the addition of a suitable compound. Chain growth in the clay galleries triggers clay exfoliation and nanocomposite formation. Thus, an advantage of the in-situ method is the preparation of polymer hybrids without physical or chemical interactions between the organic polymer and the inorganic material (Min and Chang 2012).

Organic modification of the clay surface introduces reactive moieties that disrupt the bundle structure and can potentially make it possible to obtain individual sheets. Organic modification involves attachment of functional moieties to the open ends and sidewalls of the clay, primarily to improve the solubility and dispersibility of the clay sheets. Accordingly, one of the best methods of achieving a homogeneous dispersion of clays in a polymer matrix is the use of organoclays. This involves organically modifying the clays with polymers that are structurally similar to the matrix polymer to ensure that the dispersed clays are compatible with the polymer matrix and to limit any microscopic phase separation in the nanocomposites. Many papers reported large improvements in the thermal stabilities of TLCP nanocomposites by using organoclay. This enhancement of the thermal stabilities explains reasonably well the dispersed structure of clay in the nanocomposites caused by the formation of a large aspect ratio of the clay particles (Chang 2014).

The aim in this chapter was to investigate the effectiveness and influence of two different processes (melt and solution intercalation) on the morphology and thermal properties of TLCP/organoclay nanocomposites in order to obtain a material that combines the excellent thermal behavior provided by addition of inorganic particles with the versatility and easy processing characteristics of TLCP composites. The general goal of this work was to use a minimum amount of clay in the hybrids and still obtain thermal properties significantly superior to those of the matrix polymer. The properties of these nanocomposites were studied as a function of the organoclay content of the TLCP matrix.

Experimental

Materials

The source clay, Kunipia-F (Na+-MMT), was obtained from Kunimine Co. (Tokyo, Japan). By screening this Na+-MMT clay with a 325-mesh sieve to remove impurities, we obtained a clay with a cationic exchange capacity of 119 meq/100 g. Cloisite 25A (organically modified MMT; C25A) was obtained from Southern Clay Product Co. (Gonzales, LA, USA). All reagents were purchased from Aldrich Chemical Co. (Yongin, Korea). Commercially available solvents were purified by distillation.

Syntheses of Organoclays (C12- and C18-MMT)

A dispersion of Na+-MMT was added to solutions of the ammonium salts of dodecylamine (C12) and octadecylamine (C18). These organophilic MMTs were obtained through a multi-step route (Park and Chang 2000) and have been termed C12-MMT and C18-MMT, respectively.

Syntheses of TLCPs and Their Nanocomposites

Preparation of Poly(2-ethoxyhydroquinone-2-bromoterephthaloyl) (EHBT) Nanocomposites

The compound 2-ethoxyhydroquinone was synthesized via a multistep route (Lenz et al. 1991), and 2-bromoterephthalic acid was purchased from Aldrich Chemicals. Poly(2-ethoxyhydroquinone-2-bromoterephthaloyl) (EHBT) was prepared by direct polycondensation of equivalent weights of the appropriate 2-ethoxyhydroquinone and 2-bromoterephthalic acid in the presence of thionyl chloride (SOCl2) and pyridine. The detailed procedure was also described earlier by us (Chang et al. 1994). The ethoxy substituent and -Br on the TLCP not only lower the melting point but also improve the suitability for some applications. The polymer formed was thoroughly washed with methanol, with dilute HCl, and then with water before drying at 60 °C in a vacuum oven. The chemical structure of EHBT is shown in Scheme 1. Figure 1 shows the thread nematic textures of pure EHBT at 202 °C and 215 °C.
Scheme 1

Chemical structures of EHBT, OBDT, and PAM TLCPs

Fig. 1

Optical micrographs of EHBT taken at (a) 202 °C and (b) 215 °C (×250)

The inherent viscosity of the EHBT was 0.64 dL/g, which was measured at 30 °C in solutions with a concentration of 0.2 g/dL in a phenol/1,1,2,2-tetrachloroethane (TCE) = 50/50 (v/v) mixture. The solution viscosity values (see Table 1) ranged from 0.64 to 0.78. Considering that these viscosity values were obtained from pure EHBT in EHBT hybrids from which the clay content had been removed; these numbers can be regarded as constant.
Table 1

General properties of EHBT hybrids with various organoclay contents

C25A

wt%

I.V.a

Tg

°C

Tm

°C

Ti

°C

TDib

°C

wtR600c

%

LC phase

0 (pure EHBT)

0.64

92

143

225

330

37

Nematic

2

0.72

95

150

226

352

42

Nematic

4

0.68

98

150

225

352

44

Nematic

6

0.78

98

149

225

353

47

Nematic

aInherent viscosities were measured at 30 °C using 0.2 g/100 mL solutions in a phenol/1,1,2,2-tetrachloroethane (w/w = 50/50) mixture

bInitial weight reduction onset temperature

cWeight percent of residue at 600 °C

For simplicity, the hybrids are referred to in this paper as 0 wt% organoclay/TLCP, 3 wt% organoclay/TLCP, and so on, in which 3 wt% organoclay and TLCP are the organoclay and polymer components used to prepare the hybrids, respectively, and the number denotes the organoclay weight percent in the hybrid.

Because the synthetic procedures for C25A/EHBT nanocomposites with different wt% organoclay are very similar, only a representative example for the preparation of C25A/EHBT (2 wt%) is given. EHBT and 1 g of C25A were dry-mixed and melt-blended at 190 °C, which is within the nematic region of the polymer, for 30 min using a mechanical mixer. The synthetic route for EHBT nanocomposite is shown in Scheme 2.
Scheme 2

Synthetic routes for the EHBT hybrid

To identify chemical reactions such as transesterification and thermal degradation at the processing temperature, the 4 wt% C25A/TLCP hybrid was annealed at 190 °C. Differential scanning calorimetry (DSC) thermograms of the heat-treated hybrids are shown in Fig. 2. The DSC scans did not change significantly as the heat treatment time increased from 10 to 60 min at 190 °C. Chemical changes thus do not take place to any appreciable extent at the extrusion processing temperature of 190 °C. It was also confirmed by 1H- and 13C-NMR spectroscopy that no detectable transesterification reaction occurred in TLCP under the processing conditions (not shown here).
Fig. 2

DSC thermograms of 4 wt% C25A in EHBT hybrid annealed at 190 °C for different times

Preparation of Poly(Oxybiphenyleneoxy-2,5-Dihexyloxyterephthaloyl) (OBDT) Nanocomposites

Diethyl-2,5-dihexyloxyterephthalic acid was synthesized via a multistep route, and 4,4′-biphenol was purchased from Aldrich Chemical Co. The polyester-based poly(oxybiphenyleneoxy-2,5-dihexyloxyterephthaloyl) (OBDT) nanocomposite was prepared by in situ polycondensation of equivalent weights of diethyl-2,5-dihexyloxyterephthaloyl chloride and 4,4′-biphenol in the presence of C18-MMT. Because the synthetic procedures for the C18-MMT/OBDT nanocomposites with various organoclay contents are very similar, only the preparation of the 3 wt% C18-MMT/OBDT is described here. TCE/pyridine (v/v = 50/50) (100 mL) mixed solvent and C18-MMT (0.45 g) were placed in a polymerization tube, and the mixture was stirred for 24 h at room temperature. Diethyl-2,5-dihexyloxyterephthaloyl chloride was obtained from diethyl-2,5-dihexyloxyterephthalic acid and SOCl2 following a method in the literature. Diethyl-2,5-dihexyloxyterephthalic acid (20.0 g, 5.4 × 10−1 mol) was dissolved in SOCl2, and the solution was refluxed for 12 h, after which the excess SOCl2 was removed by distillation at room temperature under vacuum. In a separate tube, 10.85 g (27 mmol) of diethyl-2,5-dihexyloxyterephthaloyl chloride was added to TCE (100 mL), and this mixture was added to the organoclay/pyridine solution and 5.1 g (27 mmol) of 4,4′-biphenol/pyridine solution, with vigorous stirring to obtain a homogeneously dispersed system. This mixture was heated for 10 h at 100 °C under a steady stream of N2 gas. The product was cooled to room temperature and poured into methanol. The precipitated product was repeatedly washed with methanol and water, and dried under vacuum at 80 °C for 12 h to obtain the polyester-based OBDT nanocomposites. The chemical structure of OBDT is shown in Scheme 1. The synthetic route for OBDT nanocomposite is shown in Scheme 3.
Scheme 3

Synthetic routes for the OBDT hybrid

The polymers are soluble in mixed solvents; therefore, the solvent mixture p-chlorophenol/TCE (w/w = 50/50) was used in the solution viscosity measurement. The solution viscosity values (see Table 2) were found to range from 0.55 to 0.82. These viscosity numbers can be regarded as constant except for that for the 7 wt% C18-MMT/OBDT hybrid. In this hybrid, excess organoclay seems to prohibit production of the higher-molecular-weight polymer.
Table 2

General properties of OBDT hybrids with various organoclay contents

C18-MMT

wt%

I.V.a

Tg

°C

Tm

°C

Ti

°C

LC phase

0 (pure OBDT)

0.82

110

283

316

Nematic

1

0.79

133

289

324

Nematic

3

0.71

113

283

313

Nematic

5

0.8

113

282

309

Nematic

7

0.55

112

297

n.o.b

aInherent viscosities were measured at 30 °C using 0.1 g/100 mL solutions in a p-chlorophenol/1,1,2,2-tetrachloroethane (w/w = 50/50) mixture

bNot observed

Preparation of Polyazomethine (PAM) Nanocomposites

PAM was prepared by direct polycondensation of equivalent weights of 2,5-diaminotoluene and terephthaldicarboxyaldehyde in ethanol. The chemical structure of PAM is shown in Scheme 1. Because we used the same synthetic procedure for all the C12-MMT/PAM nanocomposites, we describe here only the method employed to prepare 1 wt% C12-MMT/PAM. 2,5-Diaminotoluene (3.1 g, 0.016 mol), and C12-MMT (0.053 g) in ethanol (300 mL) were placed in a 500 mL three-necked flask. This mixture was stirred at 25 °C for 5 h under a nitrogen atmosphere. A solution of terephthaldicarboxyaldehyde (2.2 g, 0.016 mol) in 100 mL of ethanol was then added to this mixture, which was stirred vigorously at 80 °C for 12 h. The resulting polymer was thoroughly washed with water and then with ethanol prior to drying at 60 °C in a vacuum oven. The chemical structures relevant to this synthetic route are shown in Scheme 4. As shown in Table 3, the inherent viscosity of PAM was found to be 0.77–0.87, as measured at 30 °C at a concentration of 0.1 g/dL in solution in sulfuric acid.
Scheme 4

Synthetic routes for the PAM hybrid

Table 3

General properties of PAM hybrids with various organoclay contents

C12-MMT

wt%

I.V.a

Tg

°C

Tm

°C

Ti

°C

TDib

°C

wtR650c

%

LC phase

0 (pure PAM)

0.85

113

209

236

445

68

Nematic

1

0.82

116

226

243

453

71

Nematic

3

0.77

117

218

242

451

70

Nematic

5

0.87

124

219

241

451

74

Nematic

9

0.81

126

217

239

381

75

Nematic

aInherent viscosities were measured at 30 °C using 0.1 g/100 mL solution in a sulfuric acid

bAt a 2% initial weight-loss temperature

cWeight percent of residue at 650 °C

Extrusion

The EHBT hybrids were processed for fiber spinning to examine their tensile properties. The dried blends were pressed at 160 °C and 2500 kg/cm2 for a few minutes on a hot press. The film-type blends were dried in a vacuum oven for 24 h before being extruded through the die of a capillary rheometer. From the capillary rheometer, the hot extrudates were immediately drawn at constant take-up speed to form extended extrudates having the same diameters. The cylinder temperature of the extruder was 190 °C, and the mean residence time in the capillary rheometer was about 2–3 min.

Characterizations

Thermal and thermogravimetric analyses of the hybrids were conducted under N2 atmosphere using DuPont 910 equipment (New Castle, DE, USA). The samples were heated and cooled at a rate of 20 °C/min. Wide-angle X-ray diffraction (XRD) measurements were performed at room temperature on a Rigaku (D/Max-IIIB) X-ray diffractometer (Tokyo, Japan) using Ni-filtered Cu-Kα radiation. The scanning rate was 2°/min over a range of 2θ = 2–14°.

The tensile properties of the extrudate were determined using an Instron Mechanical Tester(Model 5564) (Norfolk County, USA) at a crosshead speed of 2 mm/min. The specimens were prepared by cutting strips 5 × 70 mm2 in size. An average of at least eight individual determinations was obtained. The experimental uncertainties in the tensile strength and modulus were ±1 MPa and ±0.05 GPa, respectively.

A polarizing microscope (Leitz, Ortholux) (Lahn-Dill-Kreis, Germany) equipped with a Mettler FP-5 hot stage was used to examine the liquid crystalline behavior. The morphology of the fractured surfaces of the extrusion samples was investigated using a Hitachi S-2400 scanning electron microscope (SEM) (San Jose, CA, USA). The fractured surfaces were sputter-coated with gold for enhanced conductivity using an SPI Sputter Coater. Transmission electron microscope (TEM) photographs of ultrathin-section polymer/organoclay hybrid samples were taken on an EM 912 OMEGA (Carl Zeiss) TEM (Tokyo, Japan) using an acceleration voltage of 120 kV.

Results and Discussion

Dispersibility of Organoclay in TLCP

Figure 3 shows the XRD patterns of C25A organoclay, pure EHBT, and their EHBT hybrids with 2–6 wt% C25A for 2θ = 2–14°. The observed interlayer spacing was 2θ = 5.64° (d = 15.67 Å) for C25A. A peak was observed at 2θ = 4.69° (d = 18.84 Å) for pure EHBT. As the amount of organoclay increased from 2 to 6 wt%, C25A/EHBT hybrids showed the same peak at the same position (2θ = 4.69°). No obvious clay peaks appeared in the XRD curves of the EHBT hybrids with 2–6 wt% C25A. This indicated that these clay layers were exfoliated and dispersed homogeneously in the TLCP matrix. This was also direct evidence that the C25A/EHBT hybrids formed nanocomposites.
Fig. 3

XRD patterns of C25A and EHBT hybrids with various organoclay contents

The XRD results for the pristine clay, organoclay, and C18-MMT/OBDT nanocomposites are shown in Fig. 4. The d001 reflection for Na+-MMT was observed at 2θ = 7.38°, which corresponds to an interlayer distance of 11.98 Å. The XRD peak for the organically modified clay, C18–MMT, appeared at 2θ = 4.14°, corresponding to an interlayer distance of 21.35 Å. As expected, ion exchange between Na+-MMT and octadecylamine chloride (C18-Cl) increases its basal interlayer spacing over that of pristine Na+-MMT and shifts the diffraction peak greatly toward lower values of 2θ. Surface modification by alkylamine produced substantially larger interlayer distances for MMT; thus, it is clear that the XRD data shift to smaller angles for the modified clay. A larger interlayer spacing should generally be advantageous in intercalation of polymer chains. It should also lead to easier dissociation of the clay and thus result in hybrids with better clay dispersion (Lagaly 1999).
Fig. 4

XRD patterns of Na+-MMT, C18-MMT, and OBDT hybrids with various organoclay contents

Small peaks at d = 14.30 Å (2θ = 6.18°), 11.92 Å (2θ = 7.42°), and 9.74 Å (2θ = 9.08°) appeared in the XRD results for pure OBDT. Similar XRD peaks appeared for the OBDT hybrid with 1 wt% organoclay content. The intensity of the XRD peak at d = 14.63 Å (2θ = 6.04°), however, increased as the clay loading was increased from 3 to 7 wt%, suggesting that dispersion is more effective at lower clay loadings than at higher clay loadings. Higher clay loadings are expected to result in increased agglomeration of some portion of the clay within the TLCP matrix; however, the presence of the organoclay was found to have no effect on the location of the peak, which indicates that perfect exfoliation of the clay layer structure of the organoclay does not occur in the TLCP matrix (Chang et al. 2006).

For the PAM hybrids, Fig. 5 illustrates the XRD patterns of the pristine clay and organoclay. For Na+-MMT reacted with alkylamine (C12-MMT), this peak is broadened and shifted to 2θ = 5.28° (d = 16.74 Å), which suggests that the clay is swollen to the range of the d spacing. Figure 5 also shows the XRD curves for pure PAM and PAM hybrids with organoclay loadings of 1–9 wt%. The characteristic XRD peaks of pure PAM synthesized with an MMT interlayer were observed. For the PAM hybrids, however, a peak is present at 2θ = 6.26° (d = 14.12 Å). The intensities of the XRD peaks increased substantially as the clay loading was increased from 1 to 9 wt%, which suggests that dispersion is more effective at lower clay loadings than at higher clay loadings. Higher clay loadings are expected to result in increased agglomeration of some portion of the clay within the PAM matrix (Khonakdar et al. 2008). In addition to the main diffraction peak, an additional small peak is present at 2θ = 8.52° (d = 10.38 Å). This secondary peak is possibly related to the XRD spectrum of the PAM matrix.
Fig. 5

XRD patterns of C12-MMT and PAM hybrids with various organoclay contents

Morphology

XRD is the conventional method of determining the interlayer spacing of clay layers in the original clay and in intercalated polymer/clay nanocomposites. Unfortunately, XRD cannot detect regular stacking exceeding a layer spacing of 88 Å. Note that the commonly used definition of an exfoliated nanocomposite is based on layer spacing larger than this value. It was the electron-microscopic analyses that provided evidence for the formation of nanoscale hybrids.

Fractured surfaces of the films were viewed under SEM. A comparative analysis of SEM photographs of EHBT hybrids with different clay contents reveals the fibrous morphology and platelet orientation distribution, including the overall projection, as shown in Fig. 6.
Fig. 6

SEM photomicrographs of (a) 0 (pure EHBT), (b) 2, (c) 4, and (d) 6 wt% C25A in EHBT hybrids

The structure of polymer nanocomposites is typically elucidated using XRD and TEM. XRD offers a convenient method of determining the interlayer spacing of the periodic arrangement of the clay layers in virgin clay and polymer nanocomposites, whereas TEM provides a qualitative understanding of the microstructures through direct visualization. More direct evidence of the formation of a true nanocomposite is provided by TEM of an ultramicrotomed section. TEM micrographs of EHBT with C25A contents of 2, 4, and 6 wt% are shown in Fig. 7a–c, respectively. The dark lines are intersections of the 1-nm-thick clay layers, and the spaces between the dark lines are interlayer spaces. This TEM photograph proves that most of the clay layers of the organoclay were exfoliated and dispersed homogeneously into the TLCP matrix. This is consistent with the XRD observations shown in Fig. 3. In conclusion, we were able to successfully synthesize EHBT nanocomposites using C25A via a melting intercalation method. The preceding results indicate that the existing state of the clay particles affects the thermal behaviors and tensile mechanical properties of each organoclay/polymer hybrid.
Fig. 7

TEM photomicrographs of (a) 2, (b) 4, and (c) 6 wt% C25A in EHBT hybrids

Figure 8 shows SEM images of the fractured surfaces of the OBDT hybrids containing 0–7 wt% organoclay. Figure 8b shows that clay phases form in the OBDT hybrids: the hybrid with 1 wt% C18-MMT has morphologies consisting of clay domains that are well dispersed in a continuous TLCP phase. In contrast, the micrographs of the 5 and 7 wt% C18-MMT/OBDT hybrid (Fig. 8d, e, respectively) show voids and deformed regions that can be attributed to the coarseness of the fractured surface. The fractured surfaces were more deformed when the hybrids contained more organoclay, likely as a consequence of agglomeration of clay particles. Figure 9 shows TEM photographs of the OBDT hybrid containing 1 wt% C18-MMT. Some of the clay is dispersed well in the OBDT matrix, and some is agglomerated to a size of approximately 10–20 nm. As in the case of 5 wt% C18-MMT (see Fig. 10), these clays are for the most part agglomerated in the polymer matrix. The peaks in the XRD patterns of these samples should be attributed to these agglomerated layers (see Fig. 4). Unlike the hybrids containing 1 wt% C18-MMT, the clay layers of the 5 wt% hybrid are not intercalated into the matrix polymer, consistent with the XRD and SEM data in Figs. 4 and 8, respectively. Agglomeration of the dispersed clay phase clearly increases with increasing organoclay content. It is conceivable that the agglomerated clay domain may decrease the thermal properties of the hybrids, probably because of poor interfacial adhesion between the organoclay and the TLCP hybrid matrix.
Fig. 8

SEM micrographs of (a) 0 (pure OBDT), (b) 1, (c) 3, (d) 5, and (e) 7 wt% C18-MMT in OBDT hybrids

Fig. 9

TEM micrographs of 1 wt% C18-MMT/OBDT hybrid at (a) lower and (b) higher magnification

Fig. 10

TEM micrographs of 5 wt% C18-MMT/OBDT hybrid at (a) lower and (b) higher magnification

The TEM micrographs in Figs. 11 and 12 show the morphologies of PAM hybrids at different magnifications. Figure 11 (1 wt% C12-MMT/PAM) shows that C12-MMT is well dispersed in the polymer matrix at all magnification levels, although some agglomerated particles have formed. For the 5 wt% hybrid (Fig. 12), some of the clay is well dispersed within the PAM matrix, and the remainder is present in agglomerations. The average diameter and length of the clays in Figs. 11 and 12 are 5–10 nm and 10–20 nm, respectively. As previously mentioned, the TEM results indicate that at low organoclay contents, the clays are well dispersed throughout the PAM matrix and that agglomerated structures are evident at higher clay contents. The peaks in the XRD patterns of these samples are attributed to these agglomerated layers (see Fig. 5).
Fig. 11

TEM micrographs of 1 wt% C12-MMT/PAM hybrid at (a) lower and (b) higher magnification

Fig. 12

TEM micrographs of 5 wt% C12-MMT/PAM hybrid at (a) lower and (b) higher magnification

Thermal Behaviors

The thermal properties of EHBT hybrids with different organoclay contents are listed in Table 1. The glass transition temperatures (Tg) of EHBT hybrids increased linearly from 92 to 98 °C with clay loading from 0 to 4 wt% and leveled off at more than 4 wt% of organoclay. The increase in the Tg value of these hybrids could be the result of two factors (Agag et al. 2001). First, the effect of small amounts of dispersed clay layers on the free volume of TLCP is significant and does influence the glass transition temperature of TLCP hybrids. The second factor is ascribed to confinement of the intercalated polymer chains within the clay galleries, which prevents segmental motions of the polymer chains.

DSC traces of pure TLCP and the hybrids are shown in Fig. 13. The endothermic peak of pure TLCP appears at 143 °C and corresponds with the melt transition temperature (Tm). The maximum transition peaks of the TLCP hybrids containing different clay contents in the DSC thermograms are slightly increased to 150 °C (see Table 1). This increase in the thermal behavior of the hybrids may result from the heat insulation effect of the clay layer structure, as well as the strong interaction between the organoclay and TLCP molecular chains (Chang et al. 2014). The isotropic transition temperatures (Ti) of pure EHBT was virtually unchanged in the EHBT hybrids regardless of the organoclay loading.
Fig. 13

DSC thermograms of C25A and EHBT hybrids with various C25A contents

In addition to the higher melting point, the thermal degradation properties of the TLCP hybrids also show improvement. A comparative thermal gravimetric analysis (TGA) of pure EHBT and three hybrids with 2–6 wt% C25A is shown in Table 1 and Fig. 14. The TGA curves do not show a weight loss below 100 °C, as shown in Fig. 14, indicating that no water remained in the samples. The weight loss due to decomposition of EHBT and its hybrids was nearly the same up to a temperature of about 300 °C. After this point, the initial thermal degradation temperature (TDi) was influenced by the organoclay loading in the hybrids. Table 1 shows that the TDi values of the C25A/EHBT hybrids (at 2% weight loss) increased with increasing organoclay content. TDi was observed to be 352–353 °C depending on the composition of the clay from 2 to 6 wt% in the EHBT hybrids, with a maximum increase of 23 °C for the 6 wt% C25A/EHBT as compared to that of the pure EHBT. The weight of the residue at 600 °C increased from 37% to 47% as the clay loading increased from 0 to 6 wt%. This enhancement of the char formation is ascribed to the high heat resistance exerted by the clay itself.
Fig. 14

TGA thermograms of C25A and EHBT hybrids with various C25A contents

The thermal properties of OBDT hybrids with various organoclay contents are listed in Table 2. The Tg values of the OBDT hybrids increase from 110 to 133 °C as the clay loading is increased from 0 to 1 wt%. This increase in Tg is ascribed to confinement of the intercalated polymer chains within the clay galleries, which prevents segmental motion of the polymer chains. However, the Tg values of the TLCP hybrids with 3–7 wt% organoclay content were found to be virtually unchanged at 112–113 °C (see Fig. 15). This implies that varying the clay content does not improve the confinement of the TLCP chains. The endothermic peaks of pure TLCP appear at 283 and 316 °C, which correspond to Tm and Ti, respectively (Table 2). The maximum transition peaks of the OBDT hybrids increase with the addition of clay up to a critical content and then decrease above that critical loading. For example, the Tm and Ti values of the 1 wt% C18-MMT/OBDT hybrid are 289 and 324 °C, respectively. When the organoclay content in OBDT reaches 5 wt%, Tm is 282 °C, and Ti is 309 °C (see Fig. 16).
Fig. 15

DSC thermograms of Tg values of OBDT hybrids with various C18-MMT contents

Fig. 16

DSC thermograms of Tm and Ti values of OBDT hybrids with various C18-MMT contents

These decreases seem to be the result of clay agglomeration, which occurs upon the addition of clay to the polymer matrix above a critical clay loading. Agglomeration also reduces the heat insulation effect of the clay layers in the polymer matrix. However, agglomerated structures form and become denser in the TLCP matrix above a critical clay content (1 wt%). The presence of organoclay agglomeration in TLCP was confirmed using the XRD and electron microscopy results (see Figs. 4, 8, 9, and 10).

The thermal properties of PAM hybrids with various C12-MMT contents are listed in Table 3. The Tg values of the PAM hybrids increase linearly from 113 to 126 °C as the clay loading increases from 0 to 9 wt%. These increases in the Tg values of these hybrids might be the result of two factors, as described above. In contrast to the Tg values, the Tm, Ti, and TDi values of the hybrids increase with increasing C12-MMT content up to 1 wt% and then decrease with further increases in the organoclay loading up to 9 wt%. For example, the Tm, Ti, and TDi values of the C12-MMT/PAM hybrid with 1 wt% clay loading are higher by 17, 7, and 8 °C, respectively, than those of pure PAM.

As mentioned above, the increase in Tm upon the addition of the organoclay might result from the heat-insulating effects of the clay layer structure, as well as from interactions between the organoclay and the PAM molecular chains. The presence of the clay also enhances the initial decomposition temperatures by acting as an insulator and a mass-transport barrier to the volatile products generated during decomposition. This increase in the thermal stability can also be attributed to the high thermal stability of the clay and to interactions between the clay particles and the polymer matrix (Gilman 1999).

In contrast to the behavior observed for C12-MMT contents of 0 to 1 wt%, the Tm, Ti, and TDi values of the hybrids decreased with increasing organo-MMT content from 3 to 9 wt%. This decrease in Tm, Ti, and TDi seems to be the result of clay agglomeration, which occurs when the clay content in the polymer matrix exceeds some critical value (see Figs. 5 and 12). The DSC and TGA curves for the PAM hybrids with various organoclay contents are shown in Figs. 17 and 18, respectively.
Fig. 17

DSC thermograms of PAM hybrids with various C12-MMT contents

Fig. 18

TGA thermograms of PAM hybrids with various C12-MMT contents

We conclude that the introduction of an inorganic clay component into an organic polymer can improve the polymer’s thermal properties because of the good thermal stability of the clay. However, in this hybrid system, the maximum effect on the thermal properties, with the exception of Tg, was obtained at an organoclay loading of 1 wt%. The weight of the residue at 650 °C ranges from 68% to 75% and increases as the clay loading increases from 0 to 9 wt%.

Considering the above results, it is consistently believable that the introduction of inorganic components into organic polymers can improve their thermal stability on the basis of the fact that clays have good thermal stability.

Tensile Properties

The pure EHBT and EHBT hybrids were extruded through a capillary die with a draw ratio (DR) of 1 to examine the tensile strength and modulus of the extrudates. The DR was calculated from the ratio of the diameter of the drawn extrudate to that of the extruder die.

The tensile mechanical properties of pure EHBT and its hybrid fibers are shown in Fig. 19. The tensile strength and initial modulus of the C25A/EHBT hybrids increased with increasing organoclay content. The ultimate tensile strength of EHBT hybrid fibers increased as the organoclay content increased. When the C25A was increased from 0 to 6 wt% in the hybrids, the strength improved linearly from 11.03 to 17.28 MPa. The ultimate strength of 6 wt% C25A/EHBT was 1.6 times higher than that of pure EHBT. The same type of behavior was observed for the initial moduli. For example, the initial tensile modulus of the 2 wt% C25A hybrid was 4.03 GPa, which was about 140% higher than that of pure TLCP. When the organoclay in EHBT reaches 6 wt%, the modulus (5.76 GPa) is about 2.0 times that of the pure EHBT TLCP.
Fig. 19

Effects of C25A contents on the ultimate strength and initial modulus of EHBT hybrid fibers

This large increase in the tensile properties of hybrids owing to the presence of organoclay can be explained as follows: the amount of the increase due to the clay layers depends on interactions between rigid, rod-shaped TLCP molecules and layered organoclays, as well as on the rigid nature of the clay layers (Fornes et al. 2002). Moreover, the clay was much more rigid than the TLCP molecules and did not deform or relax as the TLCP molecules did. This improvement was possible because organoclay layers could be highly dispersed and exfoliated in the TLCP matrix. This is consistent with the general observation that the introduction of organoclay into a matrix polymer increases its strength and modulus. The percent elongation at break of all samples, however, decreases from 2% to 1% and then remains constant with further clay addition.

Liquid Crystalline Mesophase

Figure 20 shows the thread nematic textures for 2 and 6 wt% C25A/EHBT hybrids. Regardless of their clay content, the C25A/EHBT hybrids exhibited liquid crystallinity as the organoclay content increased to 6 wt%.
Fig. 20

Polarizing optical micrographs of (a) 2 and (b) 6 wt% C25A in EHBT hybrids taken at 200 °C (×250)

Figure 21 shows the threaded nematic textures of OBDT hybrids with various organoclay contents from 0 to 7 wt% C18-MMT. A substantial decrease in the liquid crystallinity of the OBDT hybrids was observed for clay loadings between 0 and 7 wt%, which suggests that the dispersion is better at a lower clay loading than at a higher clay loading. This also suggests that liquid crystallinity is maintained by a small amount of organoclay content. However, no Ti could be detected in the DSC thermogram of the OBDT hybrid with 7 wt% C18-MMT (see Table 2 and Fig. 16). This hybrid does not show any birefringence at any temperature up to decomposition, suggesting that the higher clay domains become agglomerated in the TLCP matrix more easily than those at lower organoclay contents (see Fig. 21e). This agglomerated clay could also more easily destroy the rigid mesogens. Similar to our findings, previous studies have found that nanocomposites obtained by intercalation between an organoclay and TLCP in a nematic state resulted in a loss of liquid crystallinity for the polymer.
Fig. 21

Polarizing optical micrographs of (a) 0 (pure OBDT), (b) 1, (c) 3, and (d) 5 wt% C18-MMT in OBDT hybrids taken at 295 °C and (e) 7 wt% C18-MMT in OBDT hybrid taken at 311 °C (×250)

Typical polarizing microscope photomicrographs of the optical textures of the liquid crystalline structures of PAM hybrids with various organoclay contents are shown in Fig. 22. All the liquid crystalline compositions with C12-MMT contents between 0 and 9 wt% have a threaded Schlieren texture typical of nematogens.
Fig. 22

Polarizing optical micrographs of (a) 0 (pure PAM), (b) 1, (c) 3, and (d) 5 wt% C12-MMT in PAM hybrids taken at 235 °C (×250)

Conclusions

In this chapter, TLCP hybrids with different organoclay contents were prepared using different intercalation methods and monomers containing different substituents to investigate their thermomechanical properties, morphology, and liquid crystallinity. We found that the addition of a small amount of organoclay was sufficient to change the thermal properties of the TLCP matrix polymer. The TLCP hybrid used in this study showed considerably higher thermal properties than neat TLCP. In the OBDT and PAM hybrids, the thermal properties (Tm, Ti, and TDi) of the TLCP hybrids increase with the addition of the organoclay up to a critical content and then decrease with further organoclay loading. The XRD, SEM, and TEM results for the TLCP hybrids show that the clay particles are better dispersed in the matrix TLCP at low clay contents than at high clay contents. Higher clay loadings are expected to result in increased agglomeration of some portion of the clay within the TLCP matrix. EHBT hybrids of different C25A contents were extruded with a draw ratio of 1 from a capillary rheometer to investigate their mechanical properties. The ultimate strength and initial modulus of the hybrids increased with increasing C25A content. When the amount of organoclay in EHBT reached 6 wt%, a 1.6-fold increase in the ultimate strength and a 2.0-fold increase in the initial modulus were obtained, as compared to the strength and modulus of the pure polymer matrix.

Future Directions

Nanocomposites are a class of composites derived from ultrafine inorganic particles, such as clays (with sizes in the nanometer range), that are homogeneously dispersed in a polymer matrix. Because of the nanometer sizes of the particles, these materials possess properties that are superior to those of conventional composites. In particular, the high interfacial adhesion in the nanocomposites resulting from their nanometer particle dimensions improves the physical material properties. Many works demonstrated a potential fabrication technique for nanocomposites of a TLCP with an organoclay that yields high-performance materials. Nanostructured materials often possess a combination of physical and mechanical properties that are not present in conventional composites. Even at low clay concentrations (<10 wt%), the thermomechanical properties can be substantially improved.

The various structural factors influencing the thermal and thermotropic properties of main-chain, aromatic polyesters with and without substituents or side-groups based on nanocomposites have not been systematically studied. Thus, essential studies of the structure and length of mesogenic units and substituents or side-groups, the nature of lateral substituents in the mesogenic parts, inclusion of a kink unit in the mesogenic structure, the molecular weight and thermal history of the polymer, and the copolymers’ structures and nanocompositions are the most important factors investigated. Further investigation of the thermal properties should account for the chemical structures in side groups, given that the morphology and liquid crystallinity depend strongly on the side group of the TLCP.

Cross-References

Notes

Acknowledgments

This research was supported by the Ministry of Trade, Industry & Energy (MOTIE, Korea) under the Industrial Technology Innovation Program (No. 10063420, Development of high strengthen thermotropic liquid crystal polyester fiber).

References

  1. Agag T, Koga T, Takeichi T (2001) Stidies on thermal and mechanical properties of polyimide-clay nanocomposites. Polymer 42:3399–3408CrossRefGoogle Scholar
  2. Baird DG, Sun T (1990) Novel composites from blends of amorphous and semicrystalline engineering thermoplastics with liquid-crystalline polymers. In: Weiss RA and Ober CK (ed) Liquid crystalline polymers, 1st edn. ACS Symposium Series 435, Washington, p 416–438Google Scholar
  3. Chang J-H (2014) Preparation and characterization of poly(trimethylene terephthalate) nanocomposites. In: Pandey et al (eds) Handbook of polymer nanocomposites. Processing, performance and application, Volume a: layered silicates, 1st edn. Springer, New York, pp 267–292CrossRefGoogle Scholar
  4. Chang J-H, Jo B-W, Jin J-I (1994) In situ composites of a new thermotropic LCP and PBT. Korean Polym J 2:140–147Google Scholar
  5. Chang J-H, Ju CH, Kim SH (2006) Synthesis and characterization of a series of thermotropic liquid crystalline copolyester nanocomposites. J Polym Sci B Polym Phys 44:387–397CrossRefGoogle Scholar
  6. Chang J-H, Ham M, Kim J-C (2014) Comparison of properties of poly(vinyl alcohol) nanocomposites containing two different clays. J Nanosci Nanotechnol 14:8783–8791CrossRefGoogle Scholar
  7. Fornes TD, Yoon PJ, Hunter DL, Keskkula H, Paul DR (2002) Effect of organoclay structure on nylon 6 nanocomposite morphology and properties. Polymer 43:5915–5933CrossRefGoogle Scholar
  8. Garcia-Martinez JM, Laguna O, Areso S, Collar EP (2000) A thermal and mechanical study under dynamical conditions of polypropylene/mica composites containing atactic polypropylene with succinil-fluoresceine grafted groups as interfacial modifier from the matrix side. J Polym Sci B Polym Phys 38:1564–1574CrossRefGoogle Scholar
  9. Giannelis EP (1996) Polymer layered silicate nanocomposites. Adv Mater 8:29–35CrossRefGoogle Scholar
  10. Gilman JW (1999) Flammability and thermal stability studies of polymer layered-silicate (clay) nanocomposites. Appl Clay Sci 12:31–49CrossRefGoogle Scholar
  11. Ishida H, Campbell S, Blackwell J (2000) General approach to nanocomposite preparation. Chem Mater 12:1260–1267CrossRefGoogle Scholar
  12. Khonakdar HA, Jafari SH, Asadinezhad A (2008) A review on homopolymer, blends, and nanocomposites of poly(trimethylene terephthalate) as a new addition to the aromatic polyesters class. Iran Polym J 17:19–38Google Scholar
  13. Kojima Y, Usuki A, Kawasumi M, Okada A, Kurauchi T, Kamigaito O (1994) Fine structure of nylon-6-clay hybrid. J Polym Sci B Polym Phys 32:625–630CrossRefGoogle Scholar
  14. Lagaly G (1999) Introduction: from clay mineral-polymer interactions to clay mineral-polymer nanocomposites. Appl Clay Sci 15:1–9CrossRefGoogle Scholar
  15. LeBaron PC, Wang Z, Pinnavaia TJ (1999) Polymer-layered silicate nanocomposites: an overview. Appl Clay Sci 15:11–29CrossRefGoogle Scholar
  16. Lenz RW, Furukawa A, Bhowmik P, Garay RO, Majnusz J (1991) Synthesis and characterization of extended rod thermotropic polyesters with polyoxyethylene pendant substituents. Polymer 32:1703–1712CrossRefGoogle Scholar
  17. Lusignea RW (2001) LCP extrusion and applications. In: Chung TS (ed) Thermotropic liquid crystal polymers, 1st edn. Technomic Publishing, Lancaster, pp 303–350Google Scholar
  18. McArdle CB (1989) The application of side chain liquid crystal polymers in optical data storage. In: McArdle CB (ed) Side chain liquid crystal polymers, 1st edn. Blackie, New York, pp 357–394Google Scholar
  19. Min U, Chang J-H (2012) Colorless and transparent polyimide nanocomposite films: thermo-optical properties, morphology, and oxygen permeability. In: Panzini MI (ed) Thick films: properties, technology and applications, 1st edn. Nova Science, New York, pp 261–282Google Scholar
  20. Park D-K, Chang J-H (2000) Nanocomposites based on montmorillonite and thermotropic liquid crystalline polyester. Polymer (Korea) 24:399–406Google Scholar
  21. Shen Z, Simon GP, Cheng Y-B (2002) Comparison of solution intercalation and melt intercalation of polymer-clay nanocomposites. Polymer 43:4251–4260CrossRefGoogle Scholar
  22. Usuki A, Koiwai A, Kojima Y, Kawasumi M, Okada A, Kurauchi T, Kamigaito O (1995) Interaction of nylon 6-clay surface and mechanical properties of nylon 6-clay hybrid. J Appl Polym Sci 55:119–123CrossRefGoogle Scholar
  23. Vaia RA (2000) Structural characterization of polymer-layered silicate nanaocomposites. In: Pinnavaia TJ, Beal GW (eds) Polymer-clay nanocomposites, 1st edn. John Wiley & Sons, Chichester, pp 229–266Google Scholar
  24. Yang F, Qu Y, Yu Z (1998) Polyamide 6/silica nanocomposites prepared by in situ polymerization. J Appl Polym Sci 69:355–361CrossRefGoogle Scholar
  25. Yano K, Usuki A, Okada A (1997) Synthesis and properties of polyimide-clay hybrid films. J Polym Sci Part: A Polym Chem 35:2289–2294CrossRefGoogle Scholar

Copyright information

© This is a U.S. Government work and not under copyright protection in the U.S.; foreign copyright protection may apply 2020

Authors and Affiliations

  • Tae Young Ha
    • 1
  • Yong-Ho Ahn
    • 1
  • Bo-Soo Seo
    • 1
  • Donghwan Cho
    • 1
  • Jin-Hae Chang
    • 1
    Email author
  1. 1.Department of Polymer Science and EngineeringKumoh National Institute of TechnologyGumiSouth Korea

Personalised recommendations