The monomers, 4-aminophenyl sulfone (APS) 97%, sebacoyl chloride (SCC) 97%, 4-4′-oxydianiline (ODA) ≥98%, isophthaloyl chloride (IPC) ≥98% purchased from Aldrich were used as received. Triethylamine (TEA) ≥99.5%, dimethylsulfoxide (DMSO) ≥99.9%, methanol (99.8%), and hydrochloric acid >99% procured from Fluka were used as such. Montmorillonite K-10 (cation exchange capacity of 119 meq/100 g), silver nitrate (99.9%), andN,N-dimethyl acetamide (DMAc) >99% (dried over molecular sieves before use) obtained from Aldrich were used.

Amide oligomer was synthesized by reacting ODA (2 mol) and IPC (1 mol) in DMAc under anhydrous conditions. Both the monomers were dissolved in DMAc separately and then mixed by drop wise addition of ODA into IPC solution with constant stirring. The reaction mixture was placed in the ice bath to avoid any side reactions. A stoichiometric amount of TEA was added to the contents of the flask with high speed stirring for 3 h in order to quench HCl produced during the reaction. Oligomerization reaction is shown in Scheme 1. The oligomer solution was precipitated in excess methanol, filtered, and then dried under vacuum.

For the synthesis of nanocomposites, nature of the clay was first changed from hydrophilic to organophilic through an ion exchange reaction using oligomeric species as a modifier. Since oligomer was soluble in DMSO, the intercalation was carried out in the non-aqueous medium (Scheme 1). Solid oligomer (25.23 g) was dissolved in DMSO (100 mL) followed by slow addition of concentrated hydrochloric acid (4.8 mL) with constant stirring and heating at 80 °C. Montmorillonite was dispersed in another beaker in DMSO at 80 °C. This suspended clay was added to the cationic oligomer solution with stirring at 60 °C for 3 h. The precipitates of organoclay were collected by filtration and washed repeatedly with DMSO to remove the residual ammonium salt of oligomer until no AgCl precipitates identified with AgNO₃solution. These precipitates were dried in a vacuum oven at 60 °C for 24 h. The dried cake was ground and screened with a 325-mesh sieve. The powder obtained was termed as oligomer-MMT and used for the preparation of nanocomposites.

Aromatic–aliphatic polyamide matrix was synthesized by condensing 0.05 mol of 4-aminophenylsulfone with 0.05 mol of sebacoyl chloride in DMAc at low temperature and under anhydrous conditions. The reaction mixture was cooled to 0 °C in order to avoid any side reactions because the reaction was highly exothermic. After 1 h, the reaction mixture was allowed to come to ambient temperature and stirring was continued for 24 h to ensure the accomplishment of the reaction. To the reaction contents, 1% of sebacoyl chloride was added in order to generate carbonyl chloride terminal ends. The polyamide formed was viscous and golden yellow in color. To this polyamide solution, stoichiometric amount of TEA was added to quench HCl produced during the reaction. Centrifugation was carried out to separate the precipitates from the pristine polyamide resin. The above synthesized polyamide resin serve as a stock solution for nanocomposite formation. Scheme 2illustrates the formation of aromatic–aliphatic polyamide chains.

Appropriate amounts of polyamide solution were mixed with oligomer-MMT to yield various concentrations ranging from 2 to 20 wt% of nanocomposite films. The mixture was stirred vigorously for 24 h at 25 °C in order to achieve uniform dispersion of organoclay in the polyamide matrix. Nanocomposite films were prepared by pouring the solutions into petri dishes, followed by solvent evaporation at 70 °C for 12 h. The nanocomposite films were further dried in vacuum oven at 80 °C to a constant weight. Scheme 2represents the formation of aromatic–aliphatic polyamide/oligomer-MMT nanocomposites.

FT-IR data for amide oligomer and thin polyamide film were recorded using Excalibur series FT-IR spectrometer, Model No. FTSW 3000MX (BIO-RAD). Weight-average (M _w) and number-average (M _n) molecular weights of polyamide was determined using a GPC equipped with Waters 515 pump. AbsoluteN,N-dimethylformamide (DMF) was used as an eluent monitored through a UV detector (UV S3702 at 270 nm) with a flow rate of 1.0 mL/min at 60 °C. XRD analysis was performed by a Philips PW 1820 diffractometer which uses Cu Kα as a radiation source. SEM micrographs were taken on a LEO Gemini 1530 scanning electron microscope at an accelerating voltage of 5.80 kV. The samples were fractured in liquid nitrogen prior to imaging. TEM images were obtained at 200 kV with FEI Tecnai F20 transmission electron microscope. The nanocomposite films were first microtomed into 60 nm ultra thin sections with a diamond knife using Leica Ultracut UCT ultramicrotome. Tensile properties of the composite films (rectangular strips) were measured according to DIN procedure 53455 at 25 °C using Testometric Universal Testing Machine M350/500. Thermal stability of nanocomposites was determined using a METTLER TOLEDO TGA/SDTA 851^ethermogravimetric analyzer at a heating rate of 10 °C/min under nitrogen.T _gof nanocomposites was recorded using a METTLER TOLEDO DSC 822^edifferential scanning calorimeter at a ramp rate of 10 °C/min in nitrogen atmosphere. The water uptake measurements of nanocomposites were performed under ASTM D570-81 procedure at 25 °C.

XRD was exploited to characterize the microstructure of Na-MMT, a layered silicate with an interlayer spacing around 1.006 nm (2θ = 8.78°). The organophilic MMT has a characteristic peak at low 2θ equal to 4.68° corresponding to a basal spacing of 1.886 nm. Data indicate that stiff and long chain structure of oligomer leads to the greaterd-spacing of montmorillonite helping for the intercalation of polyamide into interlayers of clay. The XRD pattern for Na-MMT, neat polyamide, oligomer-MMT-based nanocomposites is shown in Fig. 1. Absence of diffraction peaks in XRD pattern of composites containing up to 14 wt% oligomer-MMT is indicative of the disruption of ordered platelets to a delaminated dispersion. An exfoliated dispersion was observed at low organoclay concentration. Increase in clay concentration from 16 to 20 wt% increases the basal spacing but the order is retained that appeared in the form of small peaks (Fig. 1) resulting in intercalated nanocomposites. At low clay concentration, polyamide clay interactions overcame the van der Waals forces between silicate interlayers resulting in complete disruption of clay structure. Due to an increase in clay concentration, van der Waals interactions dominated polymer clay interactions resulting in a finite expansion of silicate interlayers and retention of clay structure.

SEM micrographs of fractured surface of the nanocomposites are presented in Fig. 2. These images did not exhibit inorganic domains at the maximum possible magnification, which means nanolayers are distributed well in the polyamide matrix. The absence of MMT particles indicates that the agglomerate is broken down to a size (submicron) that cannot be seen at this magnification. The thickness measured from the cross-sectional view of the micrograph (Fig. 2a) is found to be 0.28 mm.

The state of delamination and intercalation inferred from XRD studies was further analyzed by TEM. Transmission electron micrographs of various polyamide-based oligomer-MMT nanocomposites are demonstrated in Fig. 3. Individual crystallites of the silicate are visible as regions of alternating narrow, dark, and light bands showing a strip distribution of silicate layers. Figure 3a shows a disruption of ordered platelet with an average platelet separation of 20 nm for polyamide/oligomer-MMT composites containing 6 wt% clay content. This is an indication of dominating delaminated dispersion. TEM photographs of 10 and 20 wt% nanocomposites are represented in Fig. 3b and c, respectively. These composites showed separation from 9 to 13 nm indicating an intercalated dispersion. The silicate dark lines have variable thickness due to stack of platelets one above each other and even high level of stacking occurred in the 20 wt% clay content. The trend in platelet spacing indicated by TEM matched with the XRD results.

Tensile behavior of the system is shown in Table 1 and Fig. 4. The tensile strength of hybrid material increased up to 6 wt% oligomer-MMT (32.12 MPa) relative to the neat polyamide (18.86 MPa) and then decreased with further incorporation of organoclay. The tensile modulus increased up to 6 wt% oligomer-MMT, and then decreased with further addition of clay content. Both elongation at break point and toughness showed a decreasing behavior as compared to the pure polyamide. Mechanical data revealed improvements in the tensile strength of the hybrid materials because the stress is more efficiently transferred from the polymer matrix to the inorganic filler. Many polymeric matrices have been reinforced with MMT having no interphase interactions among the phases 47 48 49. Polyimide-clay nanocomposites derived from poly(amic acid) and modified MMT with 12-aminododecanic and dodecylamine exhibited lower thermal expansion and gas permeation properties of composite films 8 50. These modifier developed no interaction with the poly(amic acid) and remained as low molecular weight compounds after imidization thus deteriorating the thermal and mechanical properties of resulting nanocomposites. However, when a modifier containing two amine functional groups were employed where one cationic end of modifier replaced with the negatively charged silicate layers while the other group of the swelling agent reacted with poly(amic acid) molecules diffused into space between the nanolayers of MMT. In this way, modifier attached chemically to the organoclay yielding mechanically stronger nanocomposites. Similarly, chemically bonded and unbonded nanocomposites based on polyamides have also been documented by the present authors using both sol-gel and solution intercalation techniques 56 7232426 27 28. Enhancement in modulus results due to strong interactions through chemical and hydrogen bonding between the polyamide matrix and layered silicate. Nevertheless upon high loading of oligomer-MMT, silicate layers may stack together in the form of crystallites and interlayer spaces do not expand much, limiting the diffusion of the polymer chains and deteriorating the mechanical properties.

Thermal stability of the polyamide/oligomer-MMT composites determined under inert atmosphere is shown in Fig. 5. Thermal decomposition temperatures of the nanocomposites were found in the range 400–450 °C. However, the pure polyamide shows initial weight loss between 100 and 200 °C, which may be due to the removal of moisture and/or some volatiles. Thermograms indicated that nanocomposites are thermally stable, which increased with the addition of oligomer-MMT in the polyamide. Nanocomposites prepared from polyamides and different ceramic phases showed enhanced thermal stability upon the addition of these inorganic materials 23 24 27 28. The weight retained at 800 °C is roughly proportional to the amount of organoclay in the nanocomposites. Inclusion of the inorganic filler into the organic phase was found to increase the thermal stability presumably due to superior insulating features of the layered silicate which also acts as mass transport barrier to the volatile products generated during decomposition.

The glass transition temperatures of nanocomposites were recorded using DSC technique that increased with augmenting organoclay contents (Table 1). These results described a systematic increase in theT _gvalues as a function of organoclay showing greater interaction between the two disparate phases. The maximumT _gvalue (91.87 °C) was obtained with 16 wt% addition of organoclay relative to pristine polyamide (72.34 °C). Further inclusion of the oligomer-MMT decreased theT _gbecause the entire clay may not interact with the polymer matrix resulting in poor interfacial interactions. Introduction of modified clay impeded the segmental motion of the polymer chains and increased amount of organoclay shifted the baseline of DSC curve toward higher temperature. This also suggested that polyamide chains developed interactions with organophilic silicate layers. As a result, the motions of polymer chains were restricted, thereby, increasing theT _gvalues of the composite materials. Glass transition temperatures of nanocomposites increased for all the compositions studied. The change of glass transition temperature of the polymer composites relative to pure polyamide is attributed to the interaction between the filler and matrix at interfacial zones.

The presence of silicate layers may be expected to decrease the water uptake due to a more tortuous path for the diffusing molecules that must bypass impenetrable platelets. The improved barrier characteristics, chemical resistance, reduced solvent uptake, and flame retardance of clay–polymer nanocomposites take advantage from the hindered diffusion pathways through the nanocomposite. The water uptake of composite materials measured under the saturation conditions (168 h) are shown in Table 1 .The results showed maximum water absorption for the neat polyamide film 16.1% due to exposure of amide and sulfonyl polar groups to the surface of polymer where water molecules developed secondary bond forces with these polar groups. The increase in weight of the hybrid films due to uptake of water gradually decreased as the organoclay content in nanocomposites increased. This decrease is apparently due to the mutual interaction between the organic and inorganic phases. This interaction resulted in lesser availability of amide and sulfonyl groups to interact with water.