Research Article Cite This: ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX
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Restricting Molecular Mobility in Polymer Nanocomposites with Self-Assembling Low-Molecular-Weight Gel Additives Symone L. M. Alexander†,§ and LaShanda T. J. Korley*,†,‡ †
Department of Materials Science and Engineering, University of Delaware, Newark, Delaware 19716, United States Department of Chemical and Biomolecular Engineering, University of Delaware, Newark, Delaware 19716, United States
‡
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ABSTRACT: Multiscale investigation of molecular gel additives in polymer matrices guides understanding of how solution-state assemblies result in mechanically enhanced, solid-state nanocomposites. Model polymers, poly(ethylene oxide-co-epichlorohydrin) (EO−EPI) and poly(vinyl acetate) (PVAc), were utilized as matrices and reinforced by cholesterol−pyridine (CP) nanofiber networks. The CP nanofillers suppress ethylene oxide segment melting for EO−EPI composites, whereas for PVAc nanocomposites, cause a polymer−gel dissociation transition. Incorporation of crystalline CP fiber networks led to an order of magnitude increase in tensile storage modulus due to restrictions on polymer chain mobility. This decrease in molecular mobility was confirmed by decreased loss moduli for both EO−EPI and PVAc composites. Excitingly, PVAc nanocomposites display an additional relaxation mode caused by release of PVAc chains from the transient molecular gel assembly. For both EO−EPI and PVAc composites, bulk flow can be suppressed to temperatures up to 100 °C by simply increasing the CP concentration. KEYWORDS: polymer nanocomposite, low-molecular-weight gels, self-assembly, nanofibers, molecular gel
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absorbing onto the nucleating nanofiber network.2 However, the polymer additives must fit certain parameters to guide nucleation in this way. The polymer chains are required to have rigid structures, which necessitates the use of lowmolecular-weight species, as well as include a strong interaction between the additive and substrate.2 Yet, these guiding principles do not match the polymer properties necessary to fabricate mechanically robust composites. For instance, polymers used to fabricate composites need to be above their entanglement molecular weight, which usually implies high-molecular-weight species.6 Additionally, the polymer would be present in higher concentrations than the molecular gel. Under those conditions, the polymer solution behavior would also influence viscosity and the degree of interaction between the polymer chains and the MG network. Therefore, it was necessary to investigate how these new requirements for polymer additives would influence the selfassembly of the molecular gel system. In our recent work, we examined the effects of highmolecular-weight polymer additives on the nucleation of lowmolecular-weight gels.7 Poly(ethylene oxide-co-epichlorohydrin) (EO−EPI) and poly(vinyl acetate) (PVAc) were utilized as model polymers for a collapsed chain conformation (poor solubility in gelation solvent) and an extended chain conformation (good solubility in gelation solvent), respectively. A cholesterol−pyridine (CP) molecular gel in a single solvent (anisole) served as the control molecular gel network.
INTRODUCTION Self-assembled, one-dimensional nanostructures can be utilized to achieve orders of magnitude mechanical enhancements in polymer matrices. In our previous work, Stone et al. designed a polymerizable, diacetylene-based, low-molecular-weight gel (LMWG) capable of forming fiber networks across multiple length scales in a poly(ethylene oxide-co-epichlorohydrin) (EO−EPI) matrix.1 An up to 2 orders of magnitude increase in the tensile storage modulus was achieved, which outperformed a similar filler incapable of forming self-assembled nanostructures. Interestingly, polymerizing the gel network did not lead to significant enhancement in properties. Thus, the gel network structure is the key factor in achieving superior mechanics. This realization prompted two fundamental questions about LMWGs as additives: (1) how does the presence of polymers affect nucleation of the self-assembled fiber network in solution, and (2) how does the molecular gel (MG) structure lead to such drastic mechanical enhancement? The fabrication process of molecular gel-reinforced composites begins with gelation of the polymer solution, so the logical first step is to understand how polymers affect gel nucleation. Interestingly, although we were using molecular gels (MGs) as additives in composites, the molecular gel field was utilizing polymers as additives to tune the structure of MG self-assembled fiber networks.2−5 In a review by Li et al., MG structures are classified as having either permanent or transient nodes.2 Architectures with permanent nodes have fiber networks that trap solvent by random or Cayley tree branching patterns, whereas those containing transient nodes utilize fiber wrapping and bundling to achieve gelation. Furthermore, polymers can be employed to tune these assemblies by © XXXX American Chemical Society
Received: September 5, 2018 Accepted: November 14, 2018
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DOI: 10.1021/acsami.8b15112 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX
Research Article
ACS Applied Materials & Interfaces
Figure 1. Schematic representation of fabrication of CP-reinforced polymer composites.
relaxation modes should be present: one for free matrix polymer, and another for the MG-incorporated chains. In this study, our goal is to provide fundamental understanding of how structure and interaction between the fiber network and polymer matrix influence the function of the MG as an additive in polymer nanocomposites, with the hope that this knowledge will inform and inspire future composite design.
EO−EPI, because of its collapsed chain conformation, resulted in a viscous solution comprised of large polymer aggregates. When included in a CP gel, the EO−EPI aggregates served as physical barriers to gelation and resulted in the highly branched CP fiber network characteristic of gels with permanent node junctions. At low CP concentration, the aggregates caused a confinement effect that resulted in higher dissociation temperatures compared to pure CP gels, but the viscosity of the polymer solution diminished the elasticity of the gel network. In contrast, the freely extended PVAc chains in anisole allowed for co-assembly and incorporation into the self-assembling fiber network. At all concentrations, the CP gels with high-molecular-weight PVAc additives exhibited higher dissociation temperatures than pure CP gels. Instead of branching, the PVAc−CP gels exhibited transient node structures, characterized by fiber wrapping and bundling, which resulted in gels that were more elastic than their EO− EPI−CP counterparts. Thus, this fundamental study provided the knowledge that polymer chain conformation dictates whether a MG forms a permanent or transient network structure, which directly influences the dissociation and mechanics of the molecular gel. Now armed with an in-depth understanding of the influence of high-molecular-weight polymers on MG self-assembly, we seek to address the second fundamental question of how the solution-state assembly influences solid-state mechanical behavior. In this investigation, solution-cast polymer nanocomposites with molecular gel fillers were fabricated using a cholesterol−pyridine molecular gel with EO−EPI or PVAc as the polymer matrix. The effects on MG network structure caused by changing the fabrication mode to solution casting are investigated via microscopy. Additionally, differential scanning calorimetry (DSC) was utilized to elucidate any variation in thermal transitions of the polymer matrices due to the CP additive, such as the glass transition temperature and melting transitions. Also of interest is whether the CP fiber network, which could be considered crystalline in the absence of solvent, would display any thermal transitions related to gel dissociation. Finally, the MG additive’s influence on polymer chain mobility and the material’s ability to dissipate energy were investigated utilizing dynamic mechanical analysis (DMA). Our hypothesis being that if the polymer chains are incorporated into the self-assembled network, then two
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EXPERIMENTAL SECTION
Materials. All chemicals, except for EO−EPI (67% epichlorohydrin), were purchased and used as received from Sigma-Aldrich. EO− EPI was obtained from Scientific Polymer Products. Synthesis of Cholesterol−Pyridine Gelator. A cholesterol− pyridine (CP) gelator was synthesized using a previously described procedure.8 Briefly, dichloromethane (300 mL) and cholesterol chloroformate (5 g, 0.011 mol) were added to a 500 mL roundbottom flask. Then, 4-aminopyridine (2.09 g, 0.022 mol) was added to the flask, and the reaction was stirred at room temperature overnight. The product was then filtered and dried under rotary evaporation to yield an off-white solid. The solid filtrate was rinsed first with acetone and then with hexanes. The remaining product was filtered and dried using vacuum filtration to yield a white solid (4.118 g, 74% yield). The product was characterized using a 600 MHz 1H nuclear magnetic resonance spectrometer (NMR). 1H NMR (600 MHz, CDCl3) δ = 0.68 (s, 3H), 0.86−0.87 (m, 6H), 0.92−2.44 (m, 34H), 4.60−4.66 (m, 1H), 5.41−5.42 (m, 1H), 7.64 (d, 1H), 8.02 (s, J = 5.8 Hz, 2H), 8.46 (d, J = 5.8 Hz, 2H). CP-Reinforced Polymer Composite Fabrication. The desired amount of gelator filler was added to an empty vial. Then, 10 mL of a polymer solution (50 mg mL−1 EO−EPI or PVAc in anisole) was added to the vial. The mixture was heated at 120 °C and stirred until the gelator was fully dissolved, indicated by a transition from an opaque to a transparent solution. The solution was poured into a room-temperature (25 °C) Teflon mold, and loss of solvent flow was observed within ∼1 min. The composite was covered and allowed to dry (∼3−5 days). To remove the composite from the Teflon substrate without causing mechanical deformation, the composite was cooled with liquid nitrogen and then peeled off the substrate. Characterization. A Leica (polarized) optical microscope was utilized for visualization of the crystalline self-assembled nanofiber network at different concentrations in the polymer matrices. Scanning electron microscopy (SEM) was conducted on a JSM-7400F highresolution scanning electron microscope. Samples were mounted on carbon tape and then sputter-coated with an Au/Pd mixture for 15 s before imaging. B
DOI: 10.1021/acsami.8b15112 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX
Research Article
ACS Applied Materials & Interfaces Differential scanning calorimetry (DSC) was conducted on a TA Instruments Discovery Series DSC equipped with an autosampler. Neat EO−EPI and EO−EPI-based composites were subjected to a heat-cool-heat cycle from −70 to 100 °C at a rate of 10 °C min−1. Neat PVAc and PVAc-based composites were subjected to a heatcool-heat cycle from −20 to 100 °C at a rate of 10 °C min−1. Dynamic mechanical analysis (DMA) was performed on a TA Instruments DMA Q800 operating in controlled force mode. Rectangular films (∼2 mm width and 20 mm length) were cut from a bulk film. Neat EO−EPI and EO−EPI-based composites were heated from −70 to 110 °C at a rate of 5 °C min−1. Neat PVAc and PVAc-based composites were heated from −20 to 110 °C at a rate of 5 °C min−1.
in SEM. Upon introduction of CP to the composite, crystalline, branched fiber structures are observed in both POM and SEM (Figure 2). A transition from isolated, branched structures to a more interconnected, branched fiber network with increasing CP concentration is also observed, as shown in Figure 2. Thus, the confinement effect observed in prior work is still present in composite systems and is characteristic of the E-CP20 composition not the fabrication mode. Additionally, the branched fiber structures confirm that the permanent network structure is retained for CP-reinforced EO−EPI composites. P-CP composites also exhibited identical structural characteristics to their xerogel counterparts. As with neat EO−EPI, neat PVAc does not display any crystalline structures in POM or fibrous structures in SEM (Figure 3). Upon introduction of CP, crystalline, fibrous structures are observed in POM and SEM (Figure 3). In this case, the fiber structures are transient in nature, characterized by fiber wrapping or bundling rather than branching. As the CP concentration increases, the selfassembled fiber network becomes denser in nature, with the PCP60 structure exhibiting long-range directionality in both POM and SEM (Figure 3). The observed directionality could be due to the rate of self-assembly of the fiber network, which becomes faster as the CP concentration is increased. Additionally, the hot solution is casted into a roomtemperature Teflon mold. This fabrication procedure may cause gel assembly to align along the path of the spreading solution, which applies a shear force as the fiber network forms.9 However, this behavior was observed in samples taken from gels that were formed in vials,7 so this feature is more likely a characteristic of the mode and/or rate of assembly. Thus, it is concluded that the CP fiber network is templated along the high MW PVAc chains in solution via non-covalent interactions, which has been demonstrated in other MG− polymer systems.5,10 In all cases, the transient network characteristics are preserved in the P-CP composites, independent of fabrication conditions. Influence of CP Self-Assembled Fiber Networks on Glass Transition Temperatures of EO−EPI and PVAc. The next step toward understanding the influence of molecular gels in polymer composites was an examination of the changes in the thermal behavior compared to the neat polymers. Differential scanning calorimetry and dynamic mechanical analysis were utilized as tools to investigate the thermal response induced by coupling molecular gels with polymers. Cholesterol−pyridine xerogels and solution-cast polymer films were used as controls for comparison with the E-CP and P-CP composites. Interestingly, CP xerogels did not display any significant thermal transitions throughout the experimental temperature range (Figure S1). Thus, any network dissociation observed in the composites is expected to be a result of polymer−MG interactions. Likewise, any variations in the glass transition temperature (Tg) compared to the neat polymer matrices are due to the introduction of the CP self-assembled fiber network as a filler. E-CP Composites. In polymer composite materials, variations in Tg from the neat polymer matrix to the composite reflect the ability of the matrix to relax into a rubbery state from its glassy state.6 Utilizing DSC, E-CP composites display glass transition temperatures within 1 °C of each other (Figure 4a). This result is expected, as EO−EPI and CP do not interact during the gelation progress. Thus, the energy required to
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RESULTS AND DISCUSSION Fabrication and Structural Characteristics of CPReinforced Composites. Our goal in introducing highmolecular-weight polymers as majority components in molecular gel systems was to understand how the polymer solution behavior influences molecular gel self-assembly. With this approach, it is possible to elucidate the structural characteristics guiding mechanics of low-molecular-weight gel-reinforced composites. Thus, composites were fabricated utilizing the same compositions chosen from our previous work,7 since an in-depth understanding of the network structure is available as a framework. The fabrication procedure for the molecular gel-reinforced composites is described in the Experimental Section and is shown schematically in Figure 1. An important distinction must be made regarding the concentration of molecular gel in the initial solution-based system vs in the final dry composite. The initial weight percentages are calculated as weight/weight ratios of component to anisole. For example, EO−EPI and PVAc are 5 w/w % in anisole, whereas CP is calculated as 1, 2, or 3 w/w % in anisole. However, in the final composite, anisole has been removed via evaporation. Therefore, the CP to polymer weight ratio is either 1/5, 2/5, or 3/5. Thus, in the final MG−polymer composites, what were denoted as polymer-CP1, polymerCP2, and polymer-CP3 in our previous work are equivalent to 20, 40, and 60 wt % of CP in either EO−EPI or PVAc (Table 1). For the sake of clarity, the nomenclature in this study is in Table 1. Weight Ratios and Nomenclatures of E-CP and PCP Composites composite nomenclature
MG concentration in solution (wt %) (w/w solvent)
in composite (wt %) (w/w polymer)
EO−EPI
PVAc
1 2 3
20 40 60
E-CP20 E-CP40 E-CP60
P-CP20 P-CP40 P-CP60
reference to the CP to polymer weight ratio and is labeled as polymer-CP20, polymer-CP40, and polymer-CP60 (Table 1). At these compositions, the CP network is well above its percolation threshold. Post-fabrication, the composites were examined via polarized optical microscopy (POM) and scanning electron microscopy (SEM) to ensure that the fabrication process did not alter the previously observed permanent or transient fiber network structures. POM of a solution cast, neat EO−EPI film shows no crystalline structures, as evidenced by a completely dark field (Figure 2). Additionally, no fiber networks are observed C
DOI: 10.1021/acsami.8b15112 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX
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Figure 2. Polarized optical microscopy (POM) and scanning electron microscopy (SEM) of E-CP composites. Polarized optical microscopy depicts incorporation of crystalline self-assembled nanofiber networks in the EO−EPI matrix. Structural evolution of the fiber network is exhibited via both POM and SEM.
Figure 3. Polarized optical microscopy (POM) and scanning electron microscopy (SEM) of P-CP composites. Polarized optical microscopy depicts incorporation of crystalline self-assembled nanofiber networks in the PVAc matrix. Structural evolution of the fiber network is exhibited via both POM and SEM.
modulus. With the inclusion of 20 wt % CP, the Tg increases over 10 °C compared to neat EO−EPI. In our previous work utilizing in situ small-angle X-ray scattering with temperature variation, a confinement effect was observed for E-CP20 gels.7 At this concentration, CP fiber networks are concentrated into discrete, confined assemblies by EO−EPI aggregates, resulting in much higher dissociation temperatures than neat CP gels as well as the higher wt % E-CP40 and E-CP60 gels. These discrete, concentrated structures are also present in E-CP20 composites (Figure 2). In the E-CP20 solid-state systems,
induce the glass transition of EO−EPI chains does not vary significantly due to inclusion of CP. However, DSC only provides information regarding the energy barrier of the composites and does not take into account the viscoelastic behavior and overall thermal response.11 Therefore, dynamic mechanical analysis (DMA) was employed to further examine the ability of the EO−EPI matrix to transition to a rubbery state in the presence of the CP nanofiller. Figure 4b displays the Tg of the E-CP composites as a function of temperature determined from the peak of the loss D
DOI: 10.1021/acsami.8b15112 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX
Research Article
ACS Applied Materials & Interfaces
Figure 4. Thermal and thermomechanical analyses of composites using DSC and DMA. (a) DSC first heating cycle (exotherm up) for E-CP composites shows a reorganization endotherm for all systems except E-CP60, which causes the greatest restrictions on molecular mobility. (b) DMA of E-CP composites shows an increase in the rubbery plateau storage modulus, with the E-CP60 system restricting bulk flow for T > 100 °C. (c) Loss modulus decreases due to introduction of MG fiber network; the most substantial increase in Tg is in the E-CP20 system due to dense, isolated fiber structures. (d) DSC first heating cycle (exotherm up) for P-CP composites shows a disassembly endotherm for E-CP40 and E-CP60 composites. (e) Disassembly for E-CP20 composite is observable via DMA as well as the evolution of a rubbery plateau for E-40 and E-CP60 composites. (f) Loss modulus decreases with increasing CP concentration, whereas Tg increases with increasing CP concentration.
these structures are restricting bulk flow of EO−EPI, resulting in a significant increase in the glass transition temperature. ECP40 and E-CP60 composites also exhibit higher Tgs, however, without the added benefit of confinement, their influence only results in increases of 2 and 4 °C, respectively. The combination of DSC and DMA results implies that if there are no direct interactions between the polymer and molecular gel, the effects of the MG on Tg are dependent more so on the MG network’s ability to physically restrict relaxation of the polymer matrix to its rubbery state, rather than on concentration of the MG network in the composite. P-CP Composites. In contrast to E-CP composites, P-CP systems incorporate non-covalent interactions between PVAc and the nucleating CP fiber network. The interactions resulted in changes in the glass transition temperature compared to neat PVAc exhibited in both DSC and DMA results (Figure 4d,e). The interactions between the PVAc polymer chains and the self-assembled cholesterol−pyridine networks are expected to restrict the chain mobility of PVAc, causing an increase in the glass transition temperature. In turn, this causes an increase in the energy barrier for the glass transition as the CP concentration increases, which is observed via DSC. Although no significant increase in Tg was observed for P-CP20, increases of ∼4 and 6 °C were observable via DSC for PCP40 and P-CP60 composites, respectively. DMA results depicting the overall thermal response further reinforced this hypothesis. Again, the Tg of the P-CP20 composite fell within error of neat PVAc, whereas P-CP40 and P-CP60 composites displayed increases of over 10 °C. These observations imply that because the molecular gels are directly interacting with the
polymer chains via non-covalent interactions, they are able to significantly influence the Tg with changes in concentration. Secondary Relaxation and Mechanical Enhancement in CP-Polymer Composites. Examination of Tg for E-CP and P-CP composites provided insight into the mechanism guiding the influence of CP on the glassy to rubbery transition, which was found to be dependent on direct interaction with the matrix. However, the glass transition is just the beginning of the thermal response exhibited by these systems. Secondary thermal transitions were also observed for both E-CP and PCP composites via DSC (Figure 4). Therefore, we sought to understand the cause of these transitions utilizing a heat-coolheat cycle in DSC as well as DMA analysis. E-CP Composites. Neat EO−EPI, E-CP20, and E-CP40 composites all display an endothermic transition at ∼51 °C (Figure 4a). Initially, this was thought to be an ethylene oxide segment melting peak. However, in the second heating cycle, the endotherm was no longer present in the neat EO−EPI composite (Figure S2). It is well known that the fabrication history or thermal history of a polymeric material can result in additional thermal transitions.11 Thus, the second heating and cooling cycles are typically what is reported in the literature. In this work, we are interested in the effects of the assembly or fabrication process on the polymer matrix, so the first heating cycle provides key information on the relaxation behavior of our composites. In our previous work, we showed that anisole is a poor solvent for the EO−EPI matrix, resulting in collapsed chain conformations in solution.7 Due to the rapid gelation of CP molecules, EO−EPI chains are trapped in this energetically unfavorable conformation in the composite during fabrication. Upon heating in DSC, the polymer chains relax into a more E
DOI: 10.1021/acsami.8b15112 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX
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not display a secondary transition peak. However, the P-CP20 system should require the least amount of energy to dissociate, so this observation could be due to the detection limits of DSC at the chosen experimental conditions. Thus, DMA was utilized to further examine the cause of this phenomenon. Above the glass transition temperature, the elastic behavior indicated by the storage modulus varies greatly between PVAc composite systems. Though the P-CP20 system flows at nearly the same temperature as neat PVAc, an additional plateau is observed at approximately 70 °C (Figure 4e). Therefore, although an additional endotherm peak was not observed in DSC, a secondary relaxation occurs in the P-CP20 composite, observable via dynamic mechanical analysis. The CP xerogels do not exhibit any gel dissociation endotherms, so it is more likely that flow of incorporated PVAc chains induces network dissociation. Therefore, the secondary transitions and this plateau behavior are attributed to the polymer chains incorporated within the gel network, which then disassemble as the PVAc matrix begins to flow. The existence of this additional relaxation mode further supports the conclusion that the PVAc chains and CP nanofibers are interacting intimately with one another, and that PVAc is incorporated within the self-assembled structure of the CP fiber network. Similar to the effects on Tg, the magnitude of incorporation and subsequent dissociation of PVAc from the MG network should be tunable based on gel concentration, as was observed in P-CP40 and P-CP60 systems. Indeed, a small rubbery plateau is observed for the P-CP40 composite (Figure 4e). However, upon dissociation of the fiber network, the composite fails. Neither DSC of CP xerogels nor neat PVAc exhibit these transitions; therefore, this solid-state dissociation of the fiber network is representative of the point at which the PVAc chains free themselves from the CP assembly. In contrast to the P-CP20 system, the fiber network assembly in the P-CP40 system incorporates a significant portion of the polymer matrix. Consequently, disengagement of PVAc chains from the fiber network leads to bulk failure of the entire composite. Furthermore, concentration can be utilized as a handle to control this effect. The P-CP60 composite also displayed a gel dissociation peak in DSC, but in this case a well-defined rubbery plateau is observed in DMA until temperatures above 100 °C (Figure 4e). This extension of the rubbery plateau is due to the density of the fiber network, which undoubtedly incorporates more PVAc into its selfassembled structure and hinders bulk material flow of PVAc chains to temperatures of up to 105 °C, compared to the ∼87 °C flow of the neat PVAc matrix. In contrast to the Tg of P-CP composites, the loss modulus of P-CP composites decreases as a function of CP concentration (Figure 4f). This trend is attributed to the incorporation of more PVAc chains into the self-assembled fiber network as CP concentration increases. This inherently reduces chain mobility of the system, requiring more energy to induce molecular motion and disrupt the inter- and intramolecular forces present in the nanocomposites. Thus, not only do we observe increases in glass transition temperature, we also observe thermal transitions in both DSC and DMA related to these secondary relaxations from these MGreinforced PVAc composites. Composite Aging and Future Investigations. Molecular gels are known to exist in metastable state or in a condition between solubilization and crystallization. Weiss and Terech describe the kinetics of reaching a stable state post-
favorable conformation in neat EO−EPI, E-CP20, and E-CP40 composites. Interestingly, the E-CP60 composite does not exhibit this transition. The hypothesis is that the strength of the E-CP60 gel network is able to prevent the transition to a more favorable conformation. Dynamic mechanical analysis was utilized to confirm this hypothesis. The storage modulus as a function of temperature for E-CP composites is shown in Figure 4b. After the glass transition temperature, the neat EO−EPI composite begins to flow at ∼35 °C. All E-CP composites exhibit storage modulus plateaus at least an order of magnitude higher than that of neat EO−EPI. Though the E-CP20 composite has the highest Tg caused by the confinement effect, its rubbery plateau storage modulus is less than those of the other E-CP composites. This phenomenon is believed to be due to the segregation of the self-assembled structures, which diminishes the overall mechanical enhancement.12,13 E-CP40 and E-CP60 have similar Tgs; however, above Tg the difference in viscoelastic behavior is evident. E-CP60 has a higher storage modulus, but more importantly it does not show any evidence of the bulk flow observed in the other three systems. The ability to maintain this mechanical enhancement nearly 70 °C above bulk flow of the neat EO−EPI matrix and over 20 °C above the E-CP40 system displays the resilience of the E-CP60 composite. This observation strongly supports the hypothesis that the CP60 molecular gel network is able to suppress molecular motion of the EO−EPI matrix. Additionally, unlike many fiber-reinforced polymer nanocomposites, a reduction in the loss modulus of E-CP nanocomposites is observed, rather than the increase observed with the storage modulus (Figure 4c).14−18 The peak of the loss modulus represents an increase in molecular motion of a material at a given temperature.19 The magnitude of the loss modulus is lower for all of the composites containing the CP molecular gel (Figure 4c). Furthermore, the temperatures at which these transitions occur are higher than for neat EO− EPI, meaning that more energy is required to induce molecular motion. It is concluded that though the molecular gels enhance EO−EPI’s ability to store energy and resist deformation, the molecular assembly of the CP nanofiber networks restricts overall chain mobility in the composites. Since the loss modulus did not increase at the same magnitude as the storage modulus, the dampening capabilities of the composite systems also decrease. Thus, the MG additive fiber network functions by restricting the flow of the polymer matrix, which enhances the mechanics of the bulk material. P-CP Composites. For the P-CP composites, the CP nanofiber network incorporates some of the polymer chains into its self-assembled structure, leading to a transient nanofiber network. Secondary transitions not present in neat PVAc were observed for P-CP40 and P-CP60 composites in DSC results, at approximately 41 and 43 °C, respectively (Figure 4d). CP xerogels did not display any thermal transitions in the experimental temperature range, so it is likely that the secondary endotherms were caused by interactions between PVAc and the CP fiber assembly. The increase in transition temperature from CP40 to CP60 is due to the higher concentration of CP in the composite, which was also observed in our previous work.7 The concentration of CP in the composite is proportional to the number of selfassembled fiber networks that are formed, meaning that higher concentrations require more energy for dissociation. The PCP20 system, which has the lowest fiber network content, does F
DOI: 10.1021/acsami.8b15112 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX
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Figure 5. Aging effects in P-CP60 composites. (a) SEM image of a new P-CP60 composite, (b) SEM image of a P-CP60 composite after 1 month, (c) thermomechanical behavior of new versus aged composites, (d) high-magnification SEM image of aged P-CP60 composite exhibiting fiber bundles leading to needle-like points.
gelation as secondary aggregation, coarsening, or aging.20 These transitions are often restricted to solution or melt-state materials; however, in the course of this investigation on solidstate systems, aging effects were also observed. Figure 5 displays SEM and DMA of the P-CP60 system immediately after fabrication (1 day) and after 1 month (aged). Dramatic changes in morphology are observed for the aged P-CP60 composite, where large, fibrillar structures protrude from the composite (Figure 5b). Upon closer examination, these fibrillar features are comprised of the nanofiber bundles observed in PCP samples, with some ending in a needle-like point (Figure 5d). This effect is not due to rapid drying rates, as the composites were dried slowly beneath a large crystallization dish over several days. To examine the impact of these morphological differences, the thermomechanical behaviors of the fresh and aged samples were compared. The storage modulus increased, whereas the glass transition temperature decreased for the aged sample (Figure 5c). This is potentially caused by separation of these needle-like fiber assemblies from the polymer matrix, which indicates a lack of interfacial adhesion with the PVAc matrix. Therefore, though the storage modulus increases due to these crystalline assemblies, their separation from the PVAc matrix frees the polymer chains to flow at lower temperatures. Long-term effects of aging in MG− polymer nanocomposites will be investigated and discussed in future work.
polymers on a nucleating molecular gel network in solutionstate systems, we inversely sought to understand the impact of the CP nanofiber network on thermal and thermomechanical transitions of the EO−EPI and PVAc matrices in solid-state composites. By solution casting these polymer−MG mixtures, it was possible to fabricate molecular gel-reinforced polymer nanocomposites. Even with a change in fabrication technique, the CP molecular gel retains its permanent node structure in EO−EPI and transient node structure in PVAc. Thus, this approach facilitated a connection between our knowledge of MG filler architecture and solid-state nanocomposite mechanics. Differential scanning calorimetry was coupled with dynamic mechanical analysis to examine the changes in thermal transitions induced by the CP gel network in each matrix. Changes in Tg for E-CP composites were most evident in DMA results, revealing that physical restriction of molecular motion is the primary mode of mechanical enhancement for an MG filler without non-covalent interactions with the matrix. This idea was confirmed by lack of an endotherm in the first and second heating cycles of the highest wt % E-CP60 system that was otherwise present in the first heating cycle of all other E-CP systems. Furthermore, the drop in loss modulus also supports the conclusion that the molecular gel filler reduces overall molecular motion in these composites. The thermal and thermomechanical behaviors of P-CP composites proved to be tunable with concentration, due to the non-covalent interactions between CP and PVAc. Increases in Tg were observed in both DSC and DMA and followed increases in concentration. Secondary relaxations were apparent in the P-CP40 and P-CP60 systems in DSC, but further investigation using DMA highlighted an additional thermal transition in the P-CP20 system as well, along with the
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CONCLUSIONS AND FUTURE DIRECTIONS In this study, we addressed fundamental questions regarding use of molecular gel fiber networks to influence the mechanics of their host matrices. Given lessons learned in our previous work regarding the influence of high-molecular-weight G
DOI: 10.1021/acsami.8b15112 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX
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copy. Both facilities are a part of University of Delaware’s shared user facilities.
evolution of rubbery plateaus in P-CP40 and P-CP60 composites. Because the CP xerogels did not display dissociation or melting temperatures in the experimental temperature range, this secondary transition is believed to be the disassembly of PVAc chains from the gel network. The persistence of a rubbery plateau in the P-CP60 system up to ∼105 °C shows that increasing MG filler concentration can provide mechanical enhancement that restricts disassembly from the matrix and resists bulk flow. As a final observation, aging effects were observed in the PCP60 system, resulting in very interesting, needle-like fiber structures, which protruded from the polymer matrix. This change in morphology also influenced thermomechanical behavior, resulting in an increase in storage modulus and a decrease in glass transition temperature. Molecular gels have proven to be a stimulating addition to the class of nanofillers in polymer nanocomposites due to their molecular assembly and hierarchical fiber structures. Along with polymer chain conformation, it would be interesting to investigate the effects of molecular gels on block co-polymers assemblies. Additionally, solid-state switchability of the gel network could find application in shape-memory and/or selfhealing materials, along with multifunctional responsive materials. It is our hope that this work is a step toward pushing the huge library of molecular gels toward solid-state materials and removing the restrictions on their understanding and usefulness outside of solution-state assemblies.
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ASSOCIATED CONTENT
S Supporting Information *
The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami.8b15112.
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REFERENCES
DSC of CP xerogels; first cooling and second heating cycles of differential scanning calorimetry (DSC); loss modulus and Tan Delta data from dynamic mechanical analysis (DMA) (PDF)
AUTHOR INFORMATION
Corresponding Author
*E-mail:
[email protected]. ORCID
LaShanda T. J. Korley: 0000-0002-8266-5000 Present Address §
Department of Chemical and Biomolecular Engineering, Georgia Institute of Technology, Atlanta, Georgia 30332, United States (S.L.M.A.) Author Contributions
The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS S.L.M.A. would like to thank the NSF Graduate Research Fellowship for financial support. This research utilized the Advanced Materials Characterization Laboratory (AMCL) for differential scanning calorimetry and dynamic mechanical analysis and the Keck Center for Advanced Microscopy and Microanalysis (Keck CAMM) for scanning electron microsH
DOI: 10.1021/acsami.8b15112 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX
Research Article
ACS Applied Materials & Interfaces (20) Weiss, R. G.; Terech, P. Molecular Gels: Materials with SelfAssembled Fibrillar Networks; Springer, 2006.
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DOI: 10.1021/acsami.8b15112 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX
