POLYMER STRUCTURE AND DYNAMICS UNDER CONFINEMENT

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1 POLYMER STRUCTURE AND DYNAMICS UNDER CONFINEMENT K. Chrissopoulou, 1 H. Papananou, 1,2 K. Androulaki, 1,2 and S. H. Anastasiadis 1,2 1 Institute of Electronic Structure and Laser, Foundation for Research and Technology-Hellas, Heraklion Crete, Greece 2 Department of Chemistry, University of Crete, Heraklion Crete, Greece ABSTRACT Polymer materials are often filled with inorganics to improve their properties. Nevertheless, their behaviour when they are close to surfaces or when they are restricted in space can be very different from that in the bulk. In this work, we investigate the influence of nano-confinement on the polymer structure and dynamics for hydrophilic hybrids of PEO with sodium montmorillonite, Na + -MMT or silica nanoparticles as well as of hyperbranched polymers with multiple functional groups with Na + -MMT. In the case of the layered silicate nanocomposites, X-ray diffraction (XRD) shows that intercalated structures are formed for all systems whereas Transmission Electron Microscopy (TEM) reveals the dispersion of the silica nanoparticles within the polymer matrix. In the case of semi-crystalline PEO, the morphology and crystallization behaviour of the hybrids were investigated with XRD, Fourier Transform Infrared Spectroscopy (FTIR), Differential Scanning Calorimetry (DSC) and Polarised Optical Microscopy (POM). Confinement is shown to modify the polymer structure and crystallinity with the effect being qualitatively different for the different types of confinement. The dynamics was investigated by Quasielastic Neutron Scattering (QENS) and Dielectric Relaxation Spectroscopy (DRS). The very local rotation of side groups seems to be unaffected by the confinement unless there are restrictions that modify the motion even in the bulk polymers. Moreover, the segmental dynamics of the confined polymers depends very strongly on the polymer / inorganic intercations. INTRODUCTION Inorganic additives are frequently utilized as fillers in polymer materials in order to improve their properties, [1] especially when they exist as a fine dispersion within the polymeric matrix, thus producing a nanocomposite.[2,3] In these cases the final properties of the hybrids are determined mainly by the existence of many interfaces.[4] Among nanocomposites, the ones comprising of polymer and inorganic layered silicates have been considered as a novel generation of composite materials due to their anticipated unique properties.[5] In these systems, three different types of structure can be identified depending on the interactions between the polymer and the inorganic nanomaterial: the phase separated, where the polymer and the inorganic are mutually immiscible, the intercalated,[6] where the polymer chains intercalate between the layers of the inorganic material forming thin polymer films with nm thickness, and the exfoliated,[7,8] where the periodic structure of the inorganic material is destroyed and the inorganic platelets are dispersed within the polymeric matrix. The intercalated nanohybrids are of particular interest, since they constitute model systems for the investigation of the static and dynamic properties of macromolecules in nano-confinement using, however, macroscopic samples and conventional analytical techniques. The behavior of polymers in thin films or close to interfaces can be very different from that in the bulk, especially when the molecules are confined to dimensions comparable to their size. Moreover, the investigation of polymer dynamics, which includes vibrational motions, rotations of side groups, the segmental -process as well as the overall chain dynamics, has attracted the scientific interest because of the complexity it exhibits over many length- and time-scales which affects greatly many of their macroscopic properties. [9] The dynamics of polymers confined in thin films or within porous media has been attracting the interest of the scientific community for more than a decade investigating the influence of both the confinement and the polymerinorganic interactions on the dynamics.[10,11] EXPERIMENTAL PART Materials: Poly(ethylene oxide) homopolymer, PEO, was purchased from Aldrich. Its molecular weight is,000 g/mol and its polydispersity index is M w/m n = 2.4, as determined by size exclusion chromatography utilizing polystyrene standards. The polymer possesses hydroxyl chain ends. It exhibits a glass transition temperature T g=-67c and a melting temperature T m=65c. Hybrane S 1200, a hyperbranched polyesteramide with a number-average molecular weight M n=1200 g/mol, was kindly supplied by DSM. Hybrane is amorphous with a glass transition temperature, T g, of K. A series of the hyperbranched polyester polyol Boltorn, H20, H30 and H40, was

2 supplied by Polymer Factory, Sweden AB. They are amorphous polymers; their branches consist mainly of ester groups and possess a number of hydroxyl end-groups. The inorganic material used was a multilayer sodium montmorillonite, Na + -MMT, Cloisite Na +, supplied by Southern Clay. It is a layered silicate clay, which is hydrophilic due to the hydrated cations that exist within its galleries. Due to this hydrophilicity, Na + -MMT can be mixed with hydrophilic polymers such as the ones used in this work without the need of any surface modification. The period of the multilayer structure is ~1 nm (when completely dry) and exhibits a cation exchange capacity, CEC, of 92.6 mmol/g. Alternatively, silica nanoparticles (Ludox LS) were purchased from Aldrich in an aqueous solution. They possess hydroxyl groups and their surface area is given by the provider to be ~220 m 2 /gr. Experimental Techniques: Structural characterization of the pure materials and of the nanocomposites was performed with X-ray diffraction, XRD whereas the investigation of their thermal properties was performed with Differential Scanning Calorimetry, DSC. The crystallinity of the materials was estimated by XRD, DSC as well as Raman Spectroscopy, RS, and the chain conformations with RS and Fourier Transform Infrared Spectroscopy, FTIR. Polymer dynamics in the bulk and under confinement was studied by Quasielastic Neutron Scattering, QENS, and Dielectric Relaxation Spectroscopy, DRS. RESULTS AND DISCUSSION Figure 1 shows the XRD measurements of all the hybrids synthesized utilizing H20 and Na + -MMT, after solution mixing and solvent evaporation, together with the diffractograms of the pure inorganic material. H20 is amorphous and exhibits just a weak amorphous halo with full width at half maximum of ~4. The main peak (001) of pure Na + - MMT is found at 2θ = 8.9, corresponding to an interlayer spacing of d 001 = 0.99 nm for the empty galleries. For intercalated nanocomposites to be realized, this main peak should be found shifted to lower angles, indicating that the polymer has entered the inorganic galleries causing an increase of the interlayer distance. At low polymer content there is a peak at 2θ = 5.7 ±0.1 corresponding to an interlayer spacing of d 001 = 1.55±0.05 nm whereas above wt% in H20 the peak jumps at 2θ = 4.5 ±0.1 corresponding to an interlayer spacing of d 001 = 1.95±0.05 nm, indicating existence of monolayers and bi-layers of polymer chains inside the galleries. [12] Intensity (a.u.) H20 weight % ( o ) Figure 1: X-ray diffractograms of H20/Na + -MMT nanohybrids. The curves have been shifted vertically for clarity. 5 0 Figure 2: TEM image of a hybrid with 80wt% PEO and 20wt% silica nanoparticles Similar is the structure of all nanocomposites synthesized from hydrophilic polymers and Na + -MMT. Figure 2 shows a TEM of another nanohybrid system consisted of PEO and silica nanoparticles. The composition of the hybrid is 80wt% polymer and 20wt% SiO 2 and the good dispersion of single nanoparticles within the polymeric matrix is obvious. The thermal properties of all systems have been investigated by DSC. Figure 3 shows the DSC thermograms of H20, H30 and H40 and of the respective hybrids with, and wt% in polymer. DSC confirms the amorphous structure of the macromolecules since only a step in the heat capacity is observed and no melting transition is present. The glass transition temperature of the pure polymers H20, H30 and H40 is obtained at 14±1 o C, 35±1 o C and 46±1 o C, respectively, i.e., a significant dependence of the T g on the generation is observed as a consequence of the increased molecular weight and the higher degree of branching. Despite the dependence of the T g on the generation, the heat capacity step at the transition seems to be constant for the three polymers having the value of ΔC p = 0.11±0.01 cal. g - 1.o C -1. The DSC thermograms of the three polymers are compared with the respective ones of the intercalated

3 nanocomposites with high polymer content. The respective Boltorn/Na + -MMT nanocomposites with wt% and wt% polymer content exhibit no identifiable glass transition, (the same holds for all hybrids with low polymer content) and only the hybrids with wt% polymer content exhibit a clear glass transition. In all cases, the transition is broader than the respective one of the neat polymers and it is found at higher temperatures. Such thermal behavior, i.e. suppressed glass transition for intercalated polymers has been observed before in similar systems of nanohybrids of both linear and hyperbranched polymers. C p (cal g -1 o C -1 ) H20/Na + -MMT (a) H30/Na + -MMT (b) H40/Na + -MMT (c) 14 o C 35 o C 46 o C 0.2 cal g -1 o C -1 Heat Flow/mass (cal g -1 o C -1 ) 20% 30% 32% 35% 38% 40% % 60% % 80% % 95% % PEO content - 0 Temperature ( o C) Temperature ( o C) Figure 3: DSC heating curves expressed as specific heat, Cp, for the bulk polymers H20, H30 and H40 and for the nanocomposites Boltorn/Na + -MMT with wt%, wt% and wt% polymer content. Figure 4: DSC heating curves for PEO and PEO / SiO 2 nanohybrids of different polymer content. The non-isothermal melting endotherms for the series of PEO/SiO 2 nanohybrids with different polymer content are shown in Figure 4 where it is obvious that DSC is able to discern two different populations with different crystallization and melting temperatures.[13] For the pure PEO as well as for hybrids with high polymer content one melting is observed at T m=67 C. Nevertheless, for the nanocomposite with % polymer, a second transition peak is first observed. The new melting transition shows a composition dependent T m and as more nanoparticles are being added, this peak grows larger (while the main peak becomes smaller), until it finally substitutes the main peak for the 30% PEO hybrid. The existence of the second peak can be attributed to polymer crystallizing under confinement and close to the silica surface. The polymer gets restricted in the space between the nanoparticles which gets smaller as the silica content increases, leaving no space for the polymer spherulites to grow. Thus the polymer crystallizes differently than it does in the bulk with the smaller spherulites giving a shifted melting transition as observed by DSC. Further than the investigation of the structure and thermal properties of these systems, their dynamic behavior has been examined as well utilizing Dielectric Relaxation Spectroscopy.[14] Figure 5 shows the relaxation times of the different processes for the H30 Boltorn hyperbranched polymer in the bulk (solid symbols) and in a nanohybrid (empty symbols) with wt% polymer i.e. one where almost all of the chains are confined. The relaxation times of the - and the -processes coincide and follow an Arrhenius temperature dependence, = 0 exp[e/rt]; this behavior is similar for the other two H20 and H40 hyperbranched polymers. The activation energies of the -process are very similar E,H20 = 65.0±1.5 kj/mol, E,H30 = 69.5±1.0 kj/mol and E,H40 = 66.5±1.5 kj/mol and 0 = O(10-18 s). This process is due to local fluctuations and rotations of the hydroxyl groups. As far as the -process is concerned, which is attributed to ester reorientation, there are few temperatures from which we can derive a relaxation time since, as temperature increases, it is in close proximity to the other two processes. The temperature dependence of its relaxation times seems very similar to the one of the -relaxation. The activation energies derived for the three polymers are E,H20 = ±3 kj/mol, E,H30 = 81±9 kj/mol and E,H40 = 86±2 kj/mol, respectively. Around the T g of H30, as determined

4 by DSC, the segmental process appears. For this process a different behavior is observed between the three polymers, as expected, due to the differences of the calorimetric glass transition of the three generations. Moreover, the temperature dependence of this process appears to follow the Vogel-Fulcher-Tammann (VFT) equation. -log( max /sec) H30 & H30 / Na + -MMT slow ' ' ' /T (K -1 ) Figure 5: Relaxation times of the different processes appearing in H30 and H30/Na + -MMT with wt% polymer content. Solid symbols corresponds to the bulk polymer whereas empty symbols to the confined polymer dynamics On the other hand, the nanocomposites exhibit a different scenario for both the sub-t g and the segmental dynamics. The relaxation times of the different processes present for all nanohybrids are shown in Figure 5, as well. Two sub-t g processes exist in this case as well; nevertheless, their relaxation times are much faster than the respective ones of the neat polymer and with much weaker Arrhenius temperature dependence. On the other hand, both processes have very similar relaxation times when the three confined polymers are compared. For the fastest process, called ', the activation energies that result from the temperature dependence of the relaxation times have values E',H20 = 32.5±2.5 kj/mol, E',H30 = 31±2 kj/mol and E',H40 = 23±1 kj/mol. This process is identified as the motion of the polar hydroxyl groups of the confined chains. It is noted that these values are very similar to the activation energy obtained for the sub-t g process for a linear poly(ethylene oxide), the dynamics of which was measured with DRS.[15] The second sub- T g process, called ', shows similar temperature dependence to the ' process. The activation energies for this process are E',H20 = 23.5±1.0 kj/mol, E',H30 = 32.5±2.5 kj/mol and E',H40 = 26±1 kj/mol. The easier carbonyl reorientation is attributed to the less restricted motion due to the fewer hydrogen bonds in this case as well. At higher temperatures, even more significant deviations in the behavior between the pure polymer and the nanohybrid appear. Firstly, there is an intermediate process that emerges at temperatures much below the bulk polymer T g, with an Arrhenius temperature dependence. This behavior resembles the '-process, observed in previous studies on PMPS and PEO under confinement. In the case of the Boltorns, the '-process is much faster than the one of the neat polymers for lower temperatures or in the proximity of the glass transition temperature but, at higher temperatures, it appears to cross the -process of the bulk polymer becoming much slower. Finally, at even higher temperatures, a slow process appears, for all confined systems, probably due to interfacial polarization due to the large number of interfaces constituting the nanocomposites and the ions trapped in their proximity. This process shows an Arrhenius temperature dependence and a significant effect on the generation as well. CONCLUSIONS The structure and dynamics of different hydrophilic polymers in the bulk and under confinement or in proximity to silica surfaces was investigated as a function of the polymer and the additive architecture. In all cases where the additive was the hydrophilic sodium montmorillonite, intercalated structure was found whereas in the case of silica nanoparticles good dispersion was obtained. The thermal behavior, the morphology and the crystallinity were found to be strongly dependent on the interactions with the surfaces and on the severe confinement. As far as the dynamics is concerned, the very local rotation of side groups seems to be unaffected by the confinement unless there are restrictions that modify the motion even in the bulk polymers whereas the segmental dynamics depends very strongly on the polymer / inorganic interactions.

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