Metallopolymers Containing Excess Metal–Ligand Complex for

Article pubs.acs.org/Macromolecules

Metallopolymers Containing Excess Metal−Ligand Complex for Improved Mechanical Properties Aaron C. Jackson,* Scott D. Walck, Kenneth E. Strawhecker, Brady G. Butler, Robert H. Lambeth, and Frederick L. Beyer U.S. Army Research Laboratory, Weapons & Materials Research Directorate, Aberdeen Proving Ground, Aberdeen, Maryland 21005-5069, United States S Supporting Information *

ABSTRACT: This work incorporates ML complexes as unbound entities that interact with ML complexes bound to the backbone of the polymer. The π−π interactions and Coulombic forces between bound and unbound ML complexes hold the ML-rich phase together and result in improved mechanical properties over polymers containing only the bound ML complexes. The ML-rich phase formed ordered, cylindrical domains. The storage modulus, surface elastic modulus, and high temperature stability of these metallopolymers increased with increasing concentration of ML complex in the polymer while an optimal concentration and morphology are necessary to improve the strength and creep resistance of the polymer. Ultimately, the successful addition and patterning of unbound ML complexes as a hard phase in a polymer matrix provides an important template for the design of a new type of supramolecular nanocomposite.

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INTRODUCTION Supramolecular polymers are polymers held together by reversible bonds capable of interacting with their surrounding environment.1−5 Research demonstrates that this reversibility is useful for interesting properties such as self-healing capabilities6−9 and low temperature processing.10 Supramolecular bonds remain intact under ambient conditions but break with mechanical, thermal, or radiative stimuli. When this happens, the local connectivity breaks down and the effective chain length decreases, enhancing the molecular mobility. Upon removal of the stimulus, the bonds re-form, and the polymer regains its original properties. Types of supramolecular interactions include hydrogen bonding,6,7,11−14 π−π stacking,9,15,16 and metal−ligand (ML) bonding.8,10,17−23 Metallopolymers are particularly versatile because the metal ion choice affects the bond strength of the metal−ligand bond. The metal−ligand complex is also highly aromatic and charged. This leads to a high potential for phase separation within solid polymer films that has important implications for mechanical and transport properties in these materials.8,19,21,22 To understand the dynamic mechanical properties that arise from adjustable bond strength, the polymer configuration, glass © XXXX American Chemical Society

transition temperature of the polymer, and strength of the metal−ligand bond are important variables. Metallopolymers studied to date have metal−ligand bonds in the main chain of the polymer,8,17,18,24,25 as cross-links,22,26 and as pendant groups27,28 off of the polymer backbone. Much of the work on mechanical properties in mechanical films has focused on main-chain and cross-linking metal−ligand bonds since those bonds contribute directly to the overall molecular weight and connectivity within the polymer. Within the research concerning the mechanical properties of metallopolymer films, the glass transition temperature (Tg) of the backbone of the polymer is also important. A polymer with a Tg lower than the working temperature has a minimal effect on the mechanical properties. Finally, the metal−ligand bond strength must be sufficient to produce polymers with useful molecular weights or cross-linking. For supramolecular polymers, the molecular weight is proportional to the square root of the bond strength bond strength.5,29,30 Indeed, many of the examples for Received: March 10, 2014 Revised: June 10, 2014

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the result of changes to the polymer itself. This approach also provides the framework for developing future supramolecular nanocomposites. This is particularly important in applications where the use of preformed nanoparticles requires higher temperatures to set a workable viscosity for the resin system.32 By using supramolecular interactions to form these nanocomposites, the energy required for manufacture can be reduced by breaking the reinforcement down during manufacture. The reinforcement can then re-form after the material shape has been determined. While this is the first attempt to demonstrate nanoscale phase separation of ML complexes at the nanoscale, other work describes similar approaches. Examples where favorable interactions lead to nanoscale or microscale phase separation include polymerdispersed liquid crystals33 and fullerenes in poly(3-hexylthiophene).34 This work is unique because the intermolecular reactions are strong enough to act as reinforcement. To develop these supramolecular nanocomposites, this work builds on our previous work using poly(butyl acrylate) crosslinked with copper(II)−MeBIP bonds.22 Based on that work, the cross-linked polymer contains ca. 4 nm in diameter clusters of ML complexes that have a well-defined small-angle X-ray scattering (SAXS) fingerprint. The low-Tg backbone and weakly bonding copper ion are chosen here to maximize rearrangement within the polymer during manufacture. While many variables will affect the properties of this system, this work focuses on manufacturing metallopolymers with excess unbound Cu−MeBIP complexes and the new morphologies associated with those polymers. The dynamic modulus and surface elastic modulus provide a measure of the stiffness of this new polymer and are compared to the stiffness of aggregates of the ML complex to understand how the rule of mixtures applies to this system. The creep resistance provides a qualitative measure of the strength of the new polymer and the stability of the aggregates of ML complex contained in the polymer.

polymers with tunable mechanical properties utilize polymers with low Tg and multidentate ligands that allow strong ML bonds. Ultimately, the choice of a polymer system depends on the type of metal-binding ligand used. Ligands such as pyridine and carboxylic acid interact well with metal ions, and indeed, this type of interaction is a viable and useful interaction in commercial polymers such as Surlyn. However, the resulting metal−ligand bonds are also labile and require a polymer backbone with inherently robust mechanical properties to form robust films. For research that focuses on the effect of adjustable metal−ligand bond strength on the mechanical properties, stronger ML bonds are required. Work by several groups have found success using multidentate ligands based on pyridine such as terpyridine 1 8 , 1 9 and 2,6-bis(1′methylbenzimidazolyl)pyridine (MeBIP).8,10,20−22 Other ligands such as a palladium/platinum pincer ligand17 and pyridine-2,6-dicarboxylic acid25 can act as ligands for metallopolymers, but to date, research on tunable mechanical properties has focused primarily on the viscosity of solutions of those polymers. The effects of variable ML bond strength and strong phase separation on the mechanical properties of metallopolymers have been discussed in previous literature. Stronger ML bonds improve the stiffness, strength, and mechanical stability of a metallopolymer at higher temperatures and strains.20−22 By using combinations of metal ions, these properties can be tuned. The interactions between adjacent metal−ligand complexes are important at low temperature and strains. As an example, polymers made by Rowan et al.20,21 require interactions between metal−ligand complexes to maintain mechanical properties above the melting temperature of the metallopolymer’s backbone. The high aromaticity of the ligand and Coulombic forces between ML complexes promote phase separation between the ML complexes and the polymer matrix. These physical cross-links are also responsible for the high modulus observed in metallopolymers at ambient temperatures.20−22 In particular, the modulus for a polymer crosslinked with ML bonds is 10 times higher than the theory of rubber elasticity predicts.21 The strong interactions between ML complexes suggest that the proportions of the ML-rich phase and polymer matrix phase can also be changed to tune mechanical properties. The use of hard segments and soft segments to modify mechanical properties is common in thermoplastic elastomers. It is well documented that increasing the fraction of hard phase, up to approximately 60−70%, improves the hardness of the polymer.31 Beck et al. indirectly considered the concentration of ML complex in metallopolymers by modifying the molecular weight of the backbone between ML complexes in a linear metallopolymer. The two concentrations of ML complex in a metallopolymer tested were 80% w/w ML complex and 30% w/w ML complex. However, the metallopolymer with 80% w/ w ML complex was mechanically unstable. No studies exist where the concentration of ML complex in a metallopolymer is systematically varied. This work incorporates unbound ML complex into a typical metallopolymer to increase the volume fraction of the hard phase in the polymer providing reinforcement. Here, unbound ML complexes, as opposed to bound ML complexes, are not covalently linked to the primary polymer backbone. This ensures that any changes in the mechanical properties are the result of an increase in the ML complex concentration and not

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EXPERIMENTAL SECTION

Materials. All solvents and reagents were purchased from commercial sources and used as received unless otherwise noted. Polymer 1 was recycled from polymer studied in previous work by Jackson et al.22 Based on that work, the concentration of ML complex in the backbone of polymer 1 is 0.4 mmol per gram polymer. MeBIP (2) was synthesized using previous methods outlined by Rowan et al.35 Film Formation. Films of 1-Cu were formed by dissolving 1 and 0.2 mmol of Cu(ClO4)26H2O per gram of polymer. The concentration of polymer in solution was kept below 3% w/v. The solvent was a mixture containing 50% v/v chloroform and 50% v/v acetonitrile. The polymer was cast in a Teflon drying pan at room temperature for 24 h and left uncovered. The resulting film was cut into sections less than 2 cm × 2 cm in size and pressed between 2 steel plates backed with release ply. Two Kapton spacers, 400 μm thick, 2 cm wide, and 10 cm long, were used to set the thickness of the films. The press cycle for these polymers was as follows: (1) 200 °C for 10 min, no pressure; (2) 200 °C and 2 tons of force for 30 min; (3) cooldown to 60 °C while maintaining 2 tons of force on the sample. The metallopolymers used in this work contained different concentrations of ML complex as outlined in Table 1. To make polymer with unbound ML complex, a stock solution was made containing Cu(ClO4)2·6H2O and 2 with a 1:2 mole ratio in a chloroform/acetonitrile (1:1 by volume) mixture. The solids content of this solution was less than 3% w/v. Separately, metallopolymer was dissolved in chloroform/acetonitrile (1:1 by volume) at less than 1% w/v. A portion of the stock solution of ML complex, corresponding to the appropriate weight fraction of unbound ML complex in the polymer, was added to the solution of metallopolymer. In addition to samples prepared using 1-Cu as the starting polymer, one sample of B

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electron loss energies before and after the nitrogen K loss peak at 401 eV. X-ray energy dispersive spectroscopy (XEDS) was performed using an Oxford Incax-Sight system. Electron counts from the Kα1 electron shell for oxygen, nitrogen, copper, and chlorine were collected per pixel and in parallel with HAADF STEM imaging. Samples for TEM were microtomed on a Leica Ultra cut UCT with the Leica EMFCS cryostage using a Microstar diamond knife with a 6° cutting angle. Most samples were microtomed at −60 °C to a thickness of 90 nm. Samples that curled or could not be transferred to a TEM grid easily were microtomed at −80 °C to a thickness of 90 nm. Atomic force microscopy (AFM) was performed using a MultiMode8 AFM with a Nanoscope V controller using PeakForce Quantitative Nanomechanical Mapping mode. Commercial AFM tips, AC200TS (Asylum Research), were used as received. The bulk of each metallopolymer film was exposed for AFM by shaving off part of the surface using a razor blade. Surface elastic modulus map images were obtained using the relative method described by PF QNM40 using the Derjaguin−Muller−Toporov (DMT) modulus model.41 The reference sample employed for this method was a spin-cast polystyrene (PS) film sample from Bruker-Nano, which has a reported surface elastic modulus of 2.7 GPa. A similar approach to modulus mapping has been described elsewhere.42−44 DMT derived surface elastic modulus values for each polymer were obtained by averaging the modulus found across each unmodified image. Images shown in this work have been contrast-adjusted for viewing.

Table 1. Polymer Components sample

Wf ML complex (% w/w)

Wf polymer (% w/w)

Wf unbound ML complex (% w/w)

1-Cu MP-04 MP-16* MP-36

18.7 19.9 21.0 47.9

100 96.5 84.2 64.1

0 3.5 15.8 35.9

MP-16 and one sample of MP-36 were prepared sequentially by adding 2-Cu to MP-04 and MP-16, respectively. Data were averaged over both types of samples. The polymer films were cast in a Teflon Petri dish overnight and pressed using the process described for making 1-Cu films. 2-Cu was cast onto glass slides for light microscopy imaging and AFM of crystals of the ML complex. First, Cu(ClO4)2·6H2O and 2 were dissolved in a solution of chloroform and acetonitrile. 0.3 mL of this solution was dropped onto a glass slide. For faster evaporation times, the slide was left uncovered and kept at 30 °C. For slower evaporation times, the slide was covered and kept at 22 °C. Instruments. Polymer film samples were cut with dimensions of approximately 5 mm wide, 25 mm long, and 0.5 mm thick and were tested in tensile mode using a DMA Q800 from TA Instruments. In creep experiments, the sample was equilibrated at 35 °C for 10 min, stressed at 3 MPa for 20 min, and allowed to rest for 30 min. Temperature sweeps were taken at 1 Hz. The stress used in creep tests, 3 MPa, provided a point of reference for obtaining a qualitative measure of strength for the polymers. Small-angle X-ray scattering (SAXS) data were collected as described elsewhere.22 Briefly, radiation (λ = 1.542 Å) was produced at a operating voltage and current of 45 kV and 100 mA. Data were collected for scattering vector magnitudes ranging from 0.007 to 0.45 Å−1. Using Igor Pro v6.3A and a package of data analysis procedures provide by Argonne National Laboratory, SAXS data were analyzed for structural information by fitting the data using a combination of different functions.36,37 Scattering at low q was fit using a power law of the form I(q) = BqP, where q is the magnitude of the scattering vector magnitude given by q = 4π sin(θ)/λ and where 2θ is the scattering angle. At intermediate angles, around q = 0.1 Å−1, the data were fit using a constant background, the form factor for a cylinder, and a lognormal distribution for cylinder radius.38 Interparticle scattering was included by using interferences calculated following Beaucage, using a modified Borne−Green approximation and quantified by the centerto-center distance and the number of particles in the nearest-neighbor sphere.39 High angle annular dark field scanning transmission electron microscopy (HAADF STEM) and energy filtered transmission electron microscopy (EFTEM) were performed on a JEOL 2100F operated at 200 kV with a Gatan energy filtering system. HAADF STEM was performed using the HAADF 5 camera length. EFTEM imaging was performed using an Ultrasoft 1000 camera. The zero-loss peak imaging used filtered electron loss energies centered around 0 eV with a 10 eV slit width. The nitrogen jump ratio imaging used filtered

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RESULTS AND DISCUSSION Scheme 1 describes the method used for making polymers containing unbound ML complex. The metallopolymer mixtures MP-04, MP-16, and MP-36 are prepared by combining the parent polymer 1-Cu with different amounts of the unbound complex 2-Cu. Based on analysis of the storage modulus of 1-Cu and visual cues, sequential reprocessing steps did not degrade the polymer. After five reprocessing steps, the storage modulus of 1-Cu decreased by 5 MPa. As a result, any changes to the polymer mechanical properties were attributed to the addition of ML complex and not polymer degradation. The new metallopolymers contained ML complexes at various concentrations as outlined in Table 1. As a point of reference, MP-16 contains equal concentrations of unbound and bound ligands. As we demonstrated in our previous work, thermal processing of the metallopolymers was critical.22 Cast films of MP-04 and MP-16 had a uniform green color, but the film thickness was not uniform. In addition to a nonuniform thickness, cast films of MP-36 were unique since the films had a lighter green layer on the surface in contact with the Teflon dish and a darker green layer on its exposed surface. This suggested that larger scale phase separation had occurred during the casting phase. Once pressed, all films had a uniform green color and a uniform thickness. Pressed films were

Scheme 1. Process for Making Polymers Cross-Linked with Metal−Ligand Bonds and Containing Unbound Metal−Ligand Complexesa

a

The process involves dissolving 1-Cu and 2-Cu separately. Both solutions are then mixed, cast, and pressed to form MP-x. C

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to clusters of ML complex approximately 4 nm in diameter.22 Under no mechanical load, the scattering data do not change with temperature, suggesting that the clusters are stable. Based on SAXS and HAADF STEM, poly(butyl acrylate) has no similar morphology confirming that the morphology observed in 1-Cu is specific to the aggregation of ML complexes. As the concentration of unbound ML complex increased, the peak shifted to lower q and the scattering intensity at low q increased. Scattering from these aggregates of unbound and bound ML complex in MP-16 and MP-36 resulted in a shoulder in the intensity profile as opposed to the peak observed in 1-Cu and MP-04. In addition, the shift in the shoulder location for MP-36 suggested that larger size aggregates were present in that polymer. At low q, the increase in the intensity is attributed to scattering from features larger than approximately 100 nm, the resolution of the SAXS data. This increase in low-q intensity corresponds with the appearance of cylindrical domains and is related to either the increasing length of the cylinders or the aggregation of those cylinders. Based on HAADF STEM images (Figure 2), these larger size-scale features are primarily cylindrical structures. HAADF STEM of cast MP-36 suggested that the cylindrical structures were inherent to the casting process and were not affected by the pressing process (see Supporting Information). In all HAADF STEM images, dark regions corresponded to material with low electron density while the light regions corresponded to material with high average atomic number. XEDS and EFTEM (see Supporting Information) confirmed that the bright regions in HAADF-STEM corresponded to material containing high concentrations of nitrogen, chlorine, and copper. These elements are characteristic of the ML complex. Geometrical analysis of the microscopy data matched well with fits of the scattering data based on a cylinder form factor.

sufficiently robust to be handled. As a control, poly(butyl acrylate) when cast with unbound metal−ligand complex was a viscous liquid. This demonstrates that the interactions between the phase separated ML complexes and the polymer matrix was important for robust films. SAXS and HAADF STEM confirmed that unbound ML complex dispersed well into microphase-separated regions within each polymer. SAXS data of 1-Cu had a peak at q = 0.1 Å−1 (Figure 1), and our previous work attributed that peak

Figure 1. SAXS of metallopolymer films cast with varying ML complex content.

Figure 2. HAADF STEM of metallopolymer films cast with varying ML complex content. 1-Cu contains aggregates of the ML complex while MP04, MP-16, and MP-36 contain a cylindrical morphology with high aspect ratio. D

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diffraction peaks corresponding to molecular order in the MLrich phase. For MP-04 and MP-16, the most prominent spacing occurred at 2θ = 13° corresponding to a lattice spacing of 6.8 Å. The spacing was not consistent across each polymer and suggested that the polymer played a role in determining extent of order in the ML-rich phase. It is also important to note that the peaks in MP-36 were not as sharp as the peaks for MP-04 and MP-16. The large scale phase separation observed in the cast polymer, based on differences in color, was the most likely cause for this. While the final polymer looked uniform in color, the opaque color of the sample made it difficult to observe phase separation throughout the polymer. The phase separation, if it is the culprit, decreased the intensity of the observed peaks in the XRD spectrum by decreasing the overall concentration of the nanoscale ML complex aggregates. Light microscopy of as-cast 2-Cu suggested that there was a direction of growth inherent to the ML complex that gives rise to the formation of 1D structures in these polymers. Similar to other crystallization processes, the concentration of the ML complex in solution and the time for solvent evaporation during the casting process affected by the shape of ML aggregates formed. Samples required low concentrations of ML complex and long solvent evaporation times to form large crystals of the ML complex. For solutions drop-cast onto glass coverslips, the best sample had a concentration lower than 0.3 mg/mL and the solvent evaporation occurred over 5 min. The crystals had a preferred growth direction such that most of them were needleshaped. In comparison, the method for casting metallopolymers utilized a higher concentration of ML complex but occurred over 24−48 h. This suggested that the ML complexes in the polymer had sufficient time to stack in a preferred direction in the polymer. The diameter of the cylinders in MP-04, MP-16, and MP-36 was much larger than the expected width of a single ML complex. This suggested that the nanoscale phase separation contains several stacks of ML complexes. The storage modulus, mechanical stability, creep resistance, and strength all improved with the addition of unbound ML

Of the samples, only MP-16 contained a sufficiently high concentration of uniform structures to compare geometrical calculations and SAXS fits. For geometrical calculations, the features analyzed by microscopy corresponded to a 2D projection of the 3D cylinders. The cylinder length for the 2D projections observed in MP-16 was 70 ± 26 nm, and the cylinder diameter was 5 ± 1 nm. Based on calculations for an ideal, monodisperse system of cylinders, this size distribution corresponded to randomly oriented cylinders 105 nm long and 5 nm in diameter. The geometrical analysis of MP-16 matched well with fits of the scattering data. The best fit model to the SAXS data predicted cylinders 150 nm long and 3 nm in diameter (see Supporting Information). To test the possibility that the morphology might be lamellar, a fit was performed using the form factor for a disk. The resulting disk-shaped parameters were not physically reasonable and did not agree with the object dimensions obtained from microscopy. Between the different polymers, there were slight differences in the diameter of the cylindrical phase separation observed in HAADF STEM (Figure 2). MP-04, MP-16, and MP-36 all contained thin cylinders with a diameter of ca. 5 nm. MP-04 and MP-36 also contained thicker cylinders with a diameter of ca. 13 nm. Qualitatively, the cylinders observed in MP-36 were longer than those observed in MP-04 or MP-16. Other noncylindrical shapes existed in the samples. Closer inspection of polymer-rich phase in MP-04 also showed clusters of the ML complex similar to the morphology observed in 1-Cu. Because of the low concentration of unbound ML complex, the clusters are associated with bound ML complexes that do not interact with unbound ML complexes. Particles were observed in MP16 and were ca. 25 nm in diameter. Based on this value and the expected thickness of the sample, 100 nm, the particles in MP16 made up less than 0.5% of the sample. These particles form where a high density of cylindrical structures were located and suggested that a particularly high local density of ML complexes was responsible for local breakdown of the structures. Despite these differences, there was no macroscale phase separation. Sharp peaks in XRD data suggested that the ML complexes stack within the microphase-separated regions (Figure 3).

Figure 3. XRD of metallopolymer films cast with varying ML complex content.

While previous research defined the geometry of similar crystals of ML complexes,45 common scattering peaks were not found. This suggests that the exact crystalline structure is highly dependent on the ML complex environment. XRD of 1-Cu contained two broad peaks that were similar to previous descriptions of ML complex aggregates.21 As a result, the broad peaks were assigned to the spacing between adjacent ML complexes (2θ = 8.6°) and the amorphous halo (2θ = 20°). Polymers containing unbound ML complex also had sharp

Figure 4. Light microscopy of ML complex cast from a solution of 2Cu (0.3 mg/mL) over 5 min at room temperature. E

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polyurethanes where a phase transition to a more brittle polymer occurs as the concentration of the hard phase approaches 60−70%.31 AFM mapping of the surface elastic modulus modulus provided a route to calculating approximate bounds for the stiffness of these metallopolymers as a function of concentration of unbound ML complex. The modulus maps of the polymer showed a morphology that correlated well to the morphology found by microscopy and X-ray scattering (Figure 6) and demonstrate that there is a difference in stiffness associated with the two phases. Although the radius of the AFM tip limited resolution, it was clear that the modulus of the MLrich phase increased in the samples with unbound ML complex. The surface elastic modulus of cast ML complex, 3.79 ± 0.33 GPa, provided a more precise measure of the expected stiffness of the ML-rich phase. Using the surface elastic modulus of the ML complex and the average surface elastic modulus for 1-Cu, 130 ± 20 MPa, the rules of mixtures and inverse mixtures were calculated. These calculations assumed a uniform density of 1 g/cm3. The surface elastic modulus for MP-16, 210 ± 20 MPa, correlated well with the inverse rule of mixtures. This was consistent with a polymer containing particulate reinforcement and confirmed that the unbound ML complex was uniformly distributed in the polymer and interacted well with the polymer. The surface elastic modulus found for MP-36 had a wider distribution of values but remained within the expected range calculated. Variation in the modulus is not surprising since both the HAADF STEM and modulus maps (see Supporting Information) showed nonuniform distributions of the cylindrical morphology.

complex (Figure 5). The storage modulus increased from 23.7 ± 1.1 MPa for 1-Cu to 72 ± 0.1 MPa for MP-36 (Figure 5a).

Figure 5. Storage modulus as a function of temperature (a) and creep tests (b) for polymers containing unbound ML complex. In creep tests, MP-04 broke prior to any data recording.

This confirmed that an increase in the ML-rich phase, the hard phase, improved the polymer stiffness. The improvement in mechanical stability at high temperatures suggested that the new interactions between ML complexes constrained and stabilized the ML bond. During creep tests (Figure 5b), MP-16 remained intact and retained most of its original shape after the application of 3 MPa of stress. In comparison, 1-Cu, MP-04, and MP-36 all broke prior to the completion of the test. As a result, MP-16 is qualitatively stronger than the other polymers since the polymer consistently withstood 3 MPa of stress while the other samples did not. The strain rate prior to break for 1Cu was larger than the strain rate observed at the same time for MP-16 and demonstrated that MP-16 had better creep resistance. However, an optimal morphology or optimal concentration is necessary to improve the creep resistance and strength since the improvement in mechanical properties did not hold at the highest concentration, MP-36. This decline in creep resistance and strength is similar to behavior in

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CONCLUSIONS The addition of unbound copper−MeBIP complexes improved the modulus, mechanical stability, creep resistance, and strength of a poly(butyl acrylate) polymer cross-linked with copper− MeBIP complexes. Unbound ML complexes interacted with the bound ML complexes in the material to form a cylindrical MLrich microphase that acts as reinforcement in the polymer. The storage modulus, surface elastic modulus, and mechanical stability of the polymers at high temperature improved with higher concentrations of unbound ML complex. The creep resistance and strength also improved, but there was an optimal concentration and morphology required to access those improvements. AFM measurements of the surface elastic modulus for these polymers and cast ML complex suggested that the modulus of these polymers could approach 1 GPa if

Figure 6. AFM surface elastic modulus maps of 1-Cu (a), MP-16 (b), and the average surface elastic modulus found by AFM as a function of the ML complex volume fraction (c). The density of the polymer is assumed to be 1 g/cm3 and uniform for the calculations. F

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(16) Colquhoun, H. M.; Williams, D. J.; Zhu, Z. J. Am. Chem. Soc. 2002, 124 (45), 13346−13347. (17) Yount, W. C.; Juwarker, H.; Craig, S. L. J. Am. Chem. Soc. 2003, 125 (50), 15302−15303. (18) Schmatloch, S.; Gonzalez, M. F.; Schubert, U. S. Macromol. Rapid Commun. 2002, 23 (16), 957−961. (19) Schmatloch, S.; van den Berg, A. M. J.; Alexeev, A. S.; Hofmeier, H.; Schubert, U. S. Macromolecules 2003, 36 (26), 9943−9949. (20) Kumpfer, J. R.; Jin, J.; Rowan, S. J. J. Mater. Chem. 2010, 20 (1), 145−151. (21) Beck, J. B.; Ineman, J. M.; Rowan, S. J. Macromolecules 2005, 38 (12), 5060−5068. (22) Jackson, A. C.; Beyer, F. L.; Price, S. C.; Rinderspacher, B. C.; Lambeth, R. H. Macromolecules 2013, 46 (14), 5416−5422. (23) Xu, D.; Liu, C.; Craig, S. L. Macromolecules 2011, 44 (7), 2343− 2353. (24) Knapp, R.; Schott, A.; Rehahn, M. Macromolecules 1996, 29 (1), 478−480. (25) Vermonden, T.; van der Gucht, J.; de Waard, P.; Marcelis, A. T. M.; Besseling, N. A. M.; Sudhölter, E. J. R.; Fleer, G. J.; Stuart, M. A. C. Macromolecules 2003, 36 (19), 7035−7044. (26) El-ghayoury, A.; Hofmeier, H.; de Reuter, B.; Schubert, U. S. Macromolecules 2003, 36 (17), 3955−3959. (27) Aamer, K. A.; Tew, G. N. Macromolecules 2007, 40 (8), 2737− 2744. (28) Shunmugam, R.; Tew, G. N. J. Am. Chem. Soc. 2005, 127 (39), 13567−13572. (29) Ciferri, A. Macromol. Rapid Commun. 2002, 23 (9), 511−529. (30) Würthner, F.; Thalacker, C.; Sautter, A.; Schärtl, W.; Ibach, W.; Hollricher, O. Chem.Eur. J. 2000, 6 (21), 3871−3886. (31) Thermoplastic Elastomers, 3rd ed.; Holden, G., Kricheldorf, H. R., Quirk, R. P., Eds.; Hanser Publishers: Munich, Germany, 2004; pp 1− 21. (32) Verber, R.; de Schepper, I. M. Phys. Rev. E 1997, 55 (3), 3143− 3158. (33) Doane, J. D.; West, J. L.; Zumer, S. US Patent 4,685,771, August 11, 1987. (34) Liao, M.; Tsai, C.; Lai, Y.; Cao, F.; Wu, J.; Wang, C.; Hsu, C.; Liau, I.; Cheng, Y. Adv. Funct. Mater. 2013, DOI: 10.1002/ adfm.201300437. (35) Rowan, S. J.; Beck, J. B. Faraday Discuss. 2005, 128, 43−53. (36) Ilavsky, J.; Jemian, P. R. J. Appl. Crystallogr. 2009, 42, 347−353. (37) Beaucage, G. J. Appl. Crystallogr. 1995, 28, 717−728. (38) Guinier, A.; Fournet, G. Small-Angle Scattering of X-Rays; John Wiley and Sons: New York, 1955. (39) Beaucage, G.; Ulibarri, T. A.; Black, E. P.; Schaefer, D. W. In ACS Sympositum Series 585: Hybrid Organic-Inorganic Composites; Mark, J. E., Lee, C. Y.-C., Bianconi, P. A., Eds.; American Chemical Society: Washington, DC, 1995; pp 97−111. (40) Peakforce QNM User Guide 004-1036-000 Rev. E; Bruker, 2011; pp 1−84. (41) Derjaguin, B. V.; Muller, V. M.; Toporov, Y. P. J. Colloid Interface Sci. 1975, 53, 314−326. (42) Schön, P.; Dutta, S.; Shirazi, M.; Noordermeer, J.; Vancso, J. G. J. Mater. Sci. 2011, 46, 3507−3516. (43) Dokukin, M. E.; Sokolov, I. Langmuir 2012, 28, 16060−16071. (44) Strawhecker, K. E.; Hsieh, A. J.; Chantawansri, T. L.; Kalcioglu, Z. I.; Van Vliet, K. J. Polymer 2013, 54, 901−908. (45) Zhu, Y.; Xia, C.; Meng, S.; Chen, J.; Chen, J.; Xie, J. Polyhedron 2013, 61, 181−187.

the morphology could be maintained at high concentrations of ML complex. While breakdown of the ideal morphology occurred at high concentrations of unbound ML complex and affected the mechanical properties, we anticipate that the future choice of components and manufacturing methods will provide a route toward eliminating or minimizing large scale phase separation. Ultimately, this work provides an important example of a supramolecular approach to building nanocomposites where the reinforcing phase is held together by reversible metal−ligand, aromatic, and charge interactions.

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ASSOCIATED CONTENT

S Supporting Information *

Additional electron microscopy, AFM, and SAXS fits. This material is available free of charge via the Internet at http:// pubs.acs.org.

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AUTHOR INFORMATION

Corresponding Author

*E-mail [email protected] (A.C.J.). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS A.C.J. is supported by the Postgraduate Research Participation Program at the US Army Research Laboratory, administered by the Oak Ridge Institute of Science and Education through an interagency agreement between the US Department of Energy and Army Research Laboratory (Contract ORISE 1120-112099).

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REFERENCES

(1) Brunsveld, L.; Folmer, B. J. B.; Meijer, E. W.; Sijbesma, R. P. Chem. Rev. 2001, 101 (12), 4071−4097. (2) Fox, J. D.; Rowan, S. J. Macromolecules 2009, 42 (18), 6823− 6835. (3) de Greef, T. F. A.; Meijer, E. W. Nature 2008, 453 (7192), 171− 173. (4) Rowan, S. J. Angew. Chem., Int. Ed. 2005, 44 (31), 4830−4832. (5) Dobrawa, R.; Würthner, F. J. Polym. Sci., Part A: Polym. Chem. 2005, 43 (21), 4981−4995. (6) Cordier, P.; Tournilhac, F.; Soulie-Ziakovic, C.; Leibler, L. Nature 2008, 451 (7181), 977−980. (7) Colquhoun, H. M. Nat. Chem. 2012, 4 (6), 435−436. (8) Burnworth, M.; Tang, L.; Kumpfer, J. R.; Duncan, A. J.; Beyer, F. L.; Fiore, G. L.; Rowan, S. J.; Weder, C. Nature 2011, 472 (7343), 334−U230. (9) Burattini, S.; Greenland, B. W.; Merino, D. H.; Weng, W.; Seppala, J.; Colquhoun, H. M.; Hayes, W.; Mackay, M. E.; Hamley, I. W.; Rowan, S. J. J. Am. Chem. Soc. 2010, 132 (34), 12051−12058. (10) Burnworth, M.; Mendez, J. D.; Schroetert, M.; Rowan, S. J.; Weder, C. Macromolecules 2008, 41 (6), 2157−2163. (11) Sivakova, S.; Bohnsack, D. A.; Mackay, M. E.; Suwanmala, P.; Rowan, S. J. J. Am. Chem. Soc. 2005, 127 (51), 18202−18211. (12) van Beek, D. J. M.; Spiering, A. J. H.; Peters, G. W. M.; te Nijenhuis, K.; Sijbesma, R. P. Macromolecules 2007, 40 (23), 8464− 8475. (13) Hirschberg, J.; Beijer, F. H.; van Aert, H. A.; Magusin, P.; Sijbesma, R. P.; Meijer, E. W. Macromolecules 1999, 32 (8), 2696− 2705. (14) Rowan, S. J.; Suwanmala, P.; Sivakova, S. J. Polym. Sci., Part A: Polym. Chem. 2003, 41 (22), 3589−3596. (15) Burattini, S.; Greenland, B. W.; Hayes, W.; Mackay, M. E.; Rowan, S. J.; Colquhoun, H. M. Chem. Mater. 2011, 23 (1), 6−8. G

dx.doi.org/10.1021/ma500516p | Macromolecules XXXX, XXX, XXX−XXX