Article pubs.acs.org/Macromolecules
From Molecular Structure to Macromolecular Organization: Keys to Design Supramolecular Biomaterials Marie Hutin, Ewelina Burakowska-Meise, Wilco P. J. Appel, Patricia Y. W. Dankers, and E. W. Meijer* Institute for Complex Molecular Systems & Laboratory of Macromolecular and Organic Chemistry, Eindhoven University of Technology, 5600 MB Eindhoven, The Netherlands S Supporting Information *
ABSTRACT: In the past decade, significant progress has been made in the field of biomaterials, for potential applications in tissue engineering or drug delivery. We have recently developed a new class of thermoplastic elastomers, based on ureidopyrimidinone (UPy) quadruple hydrogen bonding motifs. These supramolecular polymers form nanofiber-like aggregates initially via the dimerization of the UPy units followed by lateral urea-hydrogen bonding. Combined kinetic and thermodynamic studies unravel the pathway complexity in the formation of these polymorphic nanofibers and the subtlety of the polymer’s design, while these morphologies are so critically important when these materials are used in combination with cells. We also show that the cell behavior directly depends on the length and shape of the nanofibers, illustrating the key importance of macromolecular and supramolecular organization of biomaterials. This study leads to new design rules that determine what factors are decisive for a polymer to be a good candidate as biomaterial.
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INTRODUCTION In contrast to traditional synthetic polymers, supramolecular polymers rely on the use of potentially strong but reversible interactions such as coordination, electrostatic, dipole−dipole or hydrogen bonds to link the monomeric units.1−7 Due to their unusual and highly adaptable physical and mechanical properties, many applications are foreseen for such compounds particularly as biomaterials in the field of tissue engineering and drug delivery.8−11 The connections between molecular structure and macromolecular organization confer to these materials an exceptional but extremely complex dynamic character.12,13 A slight modification of the intermolecular interactions can have dramatic effects at the macromolecular level, which will directly affect the efficiency of the polymer as biomaterial.14,15 Therefore it is very important to gain a deep understanding of the molecular and supramolecular organization of these assemblies. This can be performed through the study of their fundamental characteristics (i.e., the melting and crystallization behavior), and some of these characteristicssuch as the presence of several melting transitions and morphologiesare challenging to explain. Unlike traditional polymers, very few studies have been performed to understand the origin of multimelting transitions of supramolecular polymers, both from a kinetic and thermodynamic point of view.16−19 These polymers, generally classified as thermoplastic elastomers, are built upon hydrogen-bonding motifs that crosslink block copolymers. These materials are known for their dynamic behavior and their ability to phase-separate, forming fiber-like structures. Thermodynamic aspects of the formation of polymer aggregates with different morphologies have recently © 2013 American Chemical Society
been reported. To cite a few examples, in 2001 it has been shown by ten Brinke et al. that some of the cylindrical aggregates based on diblock copolymer have a predictable shape.20 Later in 2005 Wilkes et al. studied the time-dependent change in the morphology of the hard-phase domain from a polyurethane polymer linked to 1,4-phenylene diisocyanate units.21 In the past decade, we have introduced and developed the use of the ureidopyrimidinone (UPy) units to build supramolecular thermoplastic elastomers (Figure 1). Their unique physical properties are the result of the dimerization of the enol or keto tautomers of the UPy group through self-complementary
Figure 1. Supramolecular polymer 1 with a poly(ethylene-co-butylene) spacer equipped with UPy-unit end groups and urea functionalities for lateral interactions. Received: July 23, 2013 Revised: October 13, 2013 Published: November 1, 2013 8528
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consequently also for the DSC, FTIR, and AFM experiments in order to have comparable results. Molecules substituted at both ends by UPy units are known to dimerize intra- or intermolecularly in solution.47−50 On the basis of the theoretical study of the competition between macrocyclization and linearization of Jacobsen and Stockmayer51 and of Flory,52 our group has previously reported that at low concentration the intramolecular cyclization tends to be favored, while at high concentration polymeric chains are formed. We therefore expect that the way we process the polymer here would lead to different molecular organization,23 a very important factor for cell adhesion. Indeed Figure 2 reveals the absence of the C−O
quadruple hydrogen bonding. The high dimerization constant (Kdim = 6 × 107 M−1) of the UPy in the keto form shows that this association is extremely stable in solution.22 The urea groups allow the polymer to self-organize in one-dimensional rods; the UPy layers being held on top of each other via π-stacking and dipole−dipole interactions.23,24 On a macromolecular level, these one-dimensional rods aggregate into nanofibers. Recently we reported that the shape of the nanofibers can be influenced by variation in the substituents at the 6-position of the UPy-unit and by the solvent used to self-assemble the polymer.8,23,25,26 In the light of our recent results on competing pathways for hydrogenbonded based aggregates,24 we present here detailed studies of the temperature effect on the crystallization and melting behavior of one of our polymers using DSC, AFM, and VT-FTIR analytical techniques.27−29 We show that the shape of the nanofibers depends strongly on the annealing temperature of the films and use pathway complexity to explain our results. These polymers are investigated in order to finally process them into membranes that can be used in a bioartificial kidney setup.30,31 In such a bioartificial kidney these membranes function as support for kidney epithelial cells ex vivo.32−34 In this way the kidney epithelial cells can retain their function ex vivo, and ultimately could ameliorate dialysis by removal of uremic toxins from the blood.30,31,35 In order for the epithelial cells to retain their epithelial phenotype and to function well they have to receive signals from the scaffold material. Both the surface morphology (fiber-like structure) and bioactivity are important.30,35 Great progress has been achieved since the 1990s in understanding the role of surface morphology on the behavior of cells.36−39 Many studies have reported the effect and importance of topology, morphology and nanostructures on the cells.40−46 Here we show that the history of the supramolecular materials has an effect on the shape of the nanofibers, and subsequently that these various shaped nanofibers influence cell behavior. Primary tubular epithelial cells (PTEC) are used in order to in future apply these materials for the development of membranes for a bioartificial kidney as described above. Results of these cell studies are reported here to illustrate the importance of mesoscopic morphology in the design of new biomaterials.
Figure 2. FT-IR spectra of 1 isolated by solvent evaporation and precipitation before and 12 h after the melt, emphasizing how critically important the choice of the polymer processing technique can be. The spectra have been recorded at 25 °C, and the samples melted and cooled down at a rate of 10 °C/min.
vibrations between 1000 and 1200 cm−1, as well as a considerable shift of the band at 800 cm−1. Moreover, the way of isolation influences the polymer behavior, as melting the polymer does not change its IR spectrum. Surprisingly no differences are noticed on the DSC spectra. It is clear from these data that the polymer studied here has different macromolecular organizations based on the way it is processed and might not form only 1D aggregates, therefore any comparison with our previous studies23,24 will be done cautiously. Characterization of the Supramolecular Elastomer with DSC. The typical run of a DSC experiment in this study is shown in Figure 3. The DSC sample, prepared as described above, is heated from the initial temperature (Ti = 25 °C) to the final temperature (Tf = 115 °C) at a rate of 10 °C/min, kept at this temperature for 1 min, then cooled down to the annealing temperature Ta. After the first time of crystallization t1, the same sample is heated again to Tf, allowed to sit for 1 min, and cooled down to the same Ta, where it is left for a longer crystallization time t2. This cycle is repeated n times, until the heating curve is measured after the crystallization time ti = 24 h. When not specified, tn‑1 < tn. A fresh sample has been used for each set of curves. Due to the time spent in the heating and cooling processes, the values and results obtained from the shorter crystallization times (ti < 60 min) are comparable between two set of experiments only if the annealing temperatures are the same.
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RESULTS AND DISCUSSION Material and Methods. The supramolecular polymer studied here is depicted in Figure 1. Since the bulkiness of the substituent at the 6-position of the UPy-unit is an important factor on the crystallization kinetics and the melting enthalpies of the supramolecular elastomers made,23 we chose to study polymer 1, substituted by a (R)-2,7-dimethylheptyl group at the 6-position of the UPy-unit and the amorphous poly(ethylene-cobutylene) spacer with Mn = 3500 g mol−1 and a DPI of 1.08. It has been synthesized according to the procedures published before.23 Its crystallization is nearly complete in 24 h after the melt and can therefore be conveniently followed in time with DSC, AFM and VT-FTIR. The equilibrium melting enthalpy of 1 is rather low but still high enough to be measured with an acceptable accuracy. The poly(ethylene-co-butylene) (PEB) spacer used does not show any melting transition; only a single glass transition temperature Tg of 1 is observed at −56.6 °C. Although having a relatively low melting point, it would be difficult to melt 1 to make films, but instead it needs to be solubilized (for instance in chloroform) and the film obtained by drop-casting and solvent evaporation need to be annealed and dried; this has been done for the cell studies described later but 8529
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Figure 3. Left: Successive heating DSC curves of 1 followed by several crystallization times, Ta = 25 °C. The arrows highlight the evolution of the melting transitions. Right: Melting enthalpy of the two transitions in function of the time of crystallization, with calculated melting enthalpy at the equilibrium.
Figure 4. Left: Successive heating DSC curves of 1 followed by several crystallization times (in this case tn‑1 > tn), Ta = 25 °C. The arrows highlight the evolution of the melting transitions. Right: Melting enthalpy of the two transitions in function of the time of crystallization, with calculated melting enthalpy at the equilibrium.
shortened (tn‑1 > tn), the same tendency is highlighted: the melting transitions are shifted to a higher melting temperature (Figure 4). This shows that the evolution of the transitions is not due to decomposition. More interestingly, it seems that the state of one transition after a short crystallization time is the most stable (higher melting temperature) and the one reached after a long crystallization time is the less stable (lower melting temperature). It could appear contradictory, however it is important to consider that only a small fraction of polymer 1 can be measured by DSC (the PEB part being amorphous is silent by DSC), and that by DSC only the properties related to the crystalline part are measured. Therefore the system in its whole probably becomes more stable when the fraction of less stable crystallite is increased. Multiple melting transitions have been reported for several semicrystalline polymers and can have many different origins. For example, Bas et al. attributed the double melting peaks of the poly(ether ether ketone) (PEEK) to the presence of two different crystallites with different thicknesses.54 In their case the more stable lamellae (e.g., with the highest melting temperature) are the first formed. Moreover, when the time of crystallization is increased, the first transition is shifted to the higher temperatures. Sasaki et al. as well concluded a difference of stability of the crystallites concerning the double melting transitions of the syndiotactic 1,2-polybutadiene.55 Paukkeri and Lehtinen interpreted the double melting transition of the polypropylene to a reorganization of the material during the melting scan.56 Supaphol studied the multi endotherms of the syndiotactic polypropylene. He suggested a melting−recrystallization−
Melting several times the same sample might not be judicious because of possible decomposition. However, it has been shown that a dimeric-UPy polymer is thermally degraded only when heated at 225 °C.53 Moreover, the total melting enthalpy measured on a sample of 1 that has undergone the described cycle eleven folds (the total time spent in the melt was twelve minutes, Tf = 115 °C, Ta = 25 °C, ti = 24 h) was the same as a fresh sample that has spent only one minute in the melt at the same Tf (same Ta and ti) (see in the Supporting Information). Consequently the decomposition due to the multi melting of the polymer can be discarded. As can be seen in Figure 3, one single, weak, and broad transition was observed at 85 °C 1 h after the melt. When the crystallization time was raised, it resolved into two distinct transitions, one at a higher and the other one at a lower temperature (24 h after the melt, these temperatures are 82.6 and 95 °C). These two melting transitions succeed to a broad cold crystallization transition between 65 and 80 °C that reaches its maximum between 1 and 3 h after the melt. No crystallization transition was observed in the cooling part, and modulated DSC did not reveal any duplication of the mentioned transitions (see the Supporting Information). An interesting feature on these heating DSC curves is that when the crystallization time is increased the intensity of the two melting transitions is increasing while the melting temperatures are shifted to lower temperatures. This indicates that the crystallites become less stable over time (e.g., for example due to geometry change, molecular reorganization). However, when the same experiment is reversed, i.e., when the crystallization time is 8530
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Figure 5. Left: DSC heating curves obtained for different times of crystallization at an annealing temperature of 70 °C. Right: Melting enthalpy of the transition in function of the time of crystallization, with calculated melting enthalpy at the equilibrium.
Figure 6. Heating DSC curves of 1 annealed at different temperatures 1 (a), 4 (b), and 24 (c) h after the melt.
remelting process.57 Bassett and Patel reported the lowering of the stability of the crystallites of isotactic poly(4-methylpentene1).58 They explained this phenomenon as constraints encountered during the in-filling process. Strobl et al. showed a similar trend involving the syndiotactic poly(propene-co-octene).59 They attributed this shift of the melting temperature to a higher proportion of noncrystalline units in the polymer. However in our present study the crystals that are developed when ti is short have a higher melting point than with a longer ti and therefore they possess a higher stability. This would suggest that Ostwald’s law of stage is not followed.60 In the light of the studies of the origin of several melting transitions reported so far, we can mainly consider the following possibilities: (i) the polymer is a thermotropic liquid crystal, (ii) it exhibits only one polymorph, in which case we have a phenomenon melting−recrystallization− remelting, (iii) it exhibits two polymorphs melting at different temperatures, or (iv) the polymer is present in one polymorph but phase separates when increasing the temperature; e.g., the softer part melts first, leading to a phase separation, followed by the melting of the most rigid part. First, polymer 1 is not expected to be a thermotropic liquid crystalline polymer because its rigid component is very small compared to the amorphous component. This was confirmed by pictures taken with the polarized optical microscope which did
not show any crystallization process under polarized light.61 Second, it could be the melting−recrystallization−remelting phenomenon. This is generally checked by changing the heating and cooling rate. The intensity of the second peak should decrease while increasing the rate, because the thicker crystallites (corresponding to the higher melting temperature) do not have the time to form. Figure S3 (in the Supporting Information) highlights the influence of the heating and cooling rates on the heating curves. The second melting transition decreases in intensity while the other one increases. However, a third transition appears at a lower temperature, indicating that the two transitions have another origin. In order to assess the other possibilities, more DSC experiments have been performed. Interestingly, if the same experiment than the one described in Figure 3 was run at an annealing temperature of 70 °C, only one transition was observed (Figure 5). This one was very weak after 1 h of crystallization, but became more intense after 24 h. Interestingly the temperature of the melting transition (Tm = 99.5 °C) remained in this case constant. Consequently, this suggests that the shift of the melting transitions and the appearance of the double melting peaks should be related. Therefore, several annealing temperatures were tested (Ta = 25, 30, 35, 40, 50, 55, 60, 70, 85, 100 °C) and 8531
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40 °C. By curve fitting, the values of the melting enthalpies at the equilibrium were estimated to be 2.2 J g−1 when Ta = 30 °C and only 0.9 J g−1 when Ta is 40 °C. The first transition behaves in the reverse way. The value of the associated melting enthalpy at the equilibrium goes from 2.3 J g−1 when Ta = 30 °C to 3.1 J g−1 when Ta = 40 °C, showing that the intensity of the second transition is transferred at least partially to the first one. In summary, two separate phenomena can be distinguished from the study of the effects of the annealing temperature on the DSC heating curves. When Ta is raised from 25 to 40 °C, the main crystallite is the less stable one. However, the difference of intensity between the two bands tends to be minimal by increasing the annealing time, thus suggesting that the first band corresponds to a kinetically favored product. Above 55 °C, the crystallite corresponding to the first transition is melted and the more stable crystallite is favored. This may suggest that the more stable crystallite does not necessarily correspond to the more stable molecular organization. Instead, in the crystallization temperature regime 25−40 °C, the more stable molecular organization corresponds to the less stable crystallite. The DSC studies show that it is very likely that polymer 1 aggregates into two main polymorphs, explaining the origin of the two transitions. Although polymorphs have the same chemical composition, they crystallize in different arrangements; this is why only one glass transition is seen. Because of their structural similarity, they should possess complementary hydrogen bonding motifs, and therefore can mix, forming a blend.24 When the annealing time is short after the melt, they are mixed and only one DSC transition is seen. When the annealing time is increased, the two main polymorphs self-organize exhibiting two separate transitions, both shifted to lower temperatures. It is indeed known that polymer blends can have synergistic properties compared to their separate constituents,63,64 and therefore this can explain why the temperatures of the melting transitions become lower when the annealing time increases. When the annealing temperature is augmented, the first polymorph is melted, and therefore the polymer is reorganized into the second polymorph. This is why only one transition is observed when Ta = 70 °C, with a fixed melting temperature, independent of the crystallization time. Annealing-Dependent AFM Studies. In order to assess these hypotheses and to tentatively elucidate the differences between the two polymorphs, additional experiments are performed with atomic force microscopy (AFM); a widely used technique to study the macromolecular organization of polymers.65−68 The formation of nanofibers from UPy-ureabased polymers observed by this technique has already been the objective of several publications. Before investigating the temperature effects on polymer 1 by AFM, it appears essential to understand, retrospectively, the subunits in this class of polymer allowing the formation of fibers. Sijbesma et al. has previously monitored the formation of fibers from aliphatic polymer incorporating two urea groups by AFM.69−71 In addition, we have also shown the necessity of incorporating urea groups in bis-ureidopyrimidinone-based polymers to form fibers.72 Other studies showed a strong influence of the substituent at the 6-position of the UPy unit in UPy-urea based polymers in this class of polymers on the width and length of the fibers.23 In conclusion, the fibers are formed in UPy-based polymers only when urea (or urethane) groups are present, the length and shape of the fibers depending on the group at the 6position of the UPy. The DSC measurements presented above
the heating curves were measured after several crystallization times (ti = 60, 90, 120, 240, 360, 720, 1440 min) (see Figure 6 for ti = 60, 240, and 1440 min). No transitions are observed when Ta = 100 °C (full data are provided in the Supporting Information). Several trends can be seen. First, the polymer behaves similarly for all the crystallization times when Ta = 25, 30, 35, and 40 °C. Thus, 1 h after the melt, the second transition can hardly be distinguished, after 4 h the first one is slightly more intense than the second one. Both transitions have the same intensity 24 h after the melt when Ta = 25, 30, and 35 °C. At 40 °C, the first transition is still predominant. Second, when Ta = 50 °C the two transitions exhibit closer temperatures, and coalesce at Ta = 55 °C. This temperature appears to be critical, because 1 h after the melt, the first transition is the most intense, 4 h after the melt they have the same intensity and can hardly be distinguished and 24 h after the melt the second one has the highest intensity. Finally, above this Ta, the second transition remains the main one for all the crystallization times and at 70 °C only one transition is seen. When Ta = 85 °C, the melting temperature is highly shifted and the sample has to be heated at 140 °C to see a weak transition at 125 °C. This may be attributed to an increase in the barrier height and in the crystallite size.29,62 It appears so far that this polymer exists in two different complex forms that can be interconverted below and above a critical point. Although challenging to understand, this behavior is representative to the fact that the molecular organization and supramolecular assembly are tightly bound via pathway complexity.24 It can clearly be seen from these experiments that the intensity of the first transition is transferred to the second one when Ta is increased. This is also shown by the total melting enthalpies measured 24 h after the melt for Ta = 70 °C and for Ta = 25 °C, which are in the same order of value (respectively 5.4 and 5.0 J g−1). A cold crystallization can be seen between 65 and 80 °C. When Ta becomes closer to this temperature, the crystallite corresponding to the first melting transition cannot be formed, and the material is crystallized in the form associated with the second melting transition. While the crystallite corresponding to the first transition is melted at 55 °C, it appears that its intensity increases between 30 and 40 °C, being preferred over the second one. Figure 7 shows that the intensity of the second transition decreases when the annealing temperature increases from 30 to
Figure 7. Evolution of the melting enthalpy measured from the integration of the second transition. The data have been fitted (see details in the Supporting Information) to obtain the melting enthalpy at the equilibrium. 8532
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Figure 8. Tapping mode AFM phase images taken at room temperature of samples of 1, drop cast from a solution 1 mg/mL in chloroform on a glass substrate; scan sizes 500 nm × 500 nm: (a) before the melt, 12 h at 40 °C; 24 h after the melt at an annealing temperature of (b) 25 °C, (c) 40 °C, (d) 55 °C, (e) 70 °C, (f) 85 °C, (g) 100 °C, (h) 130 °C; scan sizes 1 μm × 1 μm, 24 h after the melt at an annealing temperature of (i) 70 °C, (j) 85 °C, (k and l) 100 °C. Parts k and l have been measured from the same sample (heating and cooling rate 10 °C/min). The samples have been quenched at room temperature before taking the pictures.
evaporation of the solvent and picture a of Figure 8 is representative of all of them. As can be seen, the picture highlights very short and not well-defined nanofibers. Pictures b−g highlight a clear correlation between the annealing temperatures and the shape of the fibers. When Ta = 25 °C, the fibers are shorter than before the melt. When Ta = 40 and 55 °C, the fibers become longer, but seem still randomly organized in a three-dimensional space. At Ta = 70 °C, the fibers appear much longer, more parallel and more organized in a surface rather than in a volume.73 When Ta = 85 °C, the same tendency is observed: they are longer, more parallel, and lay onto the same surface, but they also look much thinner and bent. When Ta = 100 °C, they are organized in a completely parallel, straight
should now show the effect of annealing on the nanofibers’ formation. With this understanding, a series of AFM phase pictures in the tapping-mode have been taken at room temperature 24 h after the melt from different annealing temperatures (25, 40, 55, 70, 85, 100, and 130 °C) (Figure 8). For all samples, a solution of 1 in chloroform (1 mg/mL) was drop cast on a glass substrate. After the solvent evaporated under air, the sample was dried overnight under vacuum at 40 °C. The polymer film was afterward melted at a rate of 10 °C/min to 130 °C before being cooled down to the annealing temperature at the same rate. It was then kept at this temperature for 24 h. As reference, two pictures of three different samples have been taken before the melt after overnight 8533
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Figure 9. Variable temperature infrared experiment: 1 has been melted and cooled down at a rate of 10 °C/min, and left for 12 h at 25 °C before the measurement. This figure highlights the presence of the different temperature ranges effects.
compared with the DSC and AFM data.74,75 A sample of 1 isolated by solvent evaporation has been melted at a rate of 10 °C/min, then cooled down to 25 °C at the same rate and left for a crystallization time of 12 h. The IR spectra have been recorded by increasing the temperature stepwise, in the way that the polymer is melted with an average rate of 10 °C/min. The resulting spectra are shown in Figure 9. Three different temperature regimes can be reported from Figure 9, similar to those noticed on the DSC experiments. Between room temperature and 80 °C, all the vibrations decrease in intensity. Between 80 and 90 °C, all the peaks increase again in intensity. This is consistent with a rearrangement in the melt. Between 90 and 100 °C, two different behaviors are noticed. The vibration between 2800 and 3000 cm−1, corresponding to the polymer and the peaks corresponding to C−H vibrations between 650 - 800 cm−1, increase in intensity, suggesting that the polymer does not return in the solid state after the first melting transition−confirming that there is no re-entering phenomenon of melting−recrystallization−remelting here. A “regularity” band75 at 740 cm−1 is present in all these spectra except at 115 °C in the melt. The UPy group in the keto form is known to exhibit vibrations at 1697, 1660, 1579, and 1526 cm−1 by solid-state infrared spectroscopy.76 The first conclusion concerning the analysis of this part is that the UPy does not adopt the enol tautomer even at the highest temperature, in contrast to what has been shown in some of our previous studies.77 These vibrations decrease in intensity between 25 and 80 °C, then increase between 80 and 90
fashion on the surface, much thinner and all aligning within the same direction. This becomes more noticeable when considering the 1 μm scale pictures. Pictures k and l show a very dense network of long and thin fibers. No fiber formation can be seen from the melt at 130 °C. The pictures taken 1 week after the melt did not highlight any remarkable change in the shape of the fibers (see the Supporting Information). In general, it seems that the fibers are not split nor nucleate from another, but extend at both ends, forming a dense twodimensional network of parallel and long fibers. This is consistent with the cooperative growth of a one-dimensional stack, which phenomenon has already been reported by our group.72 These pictures are all in agreement with the DSC studies. At room temperature, where the two DSC transitions coexist, the fibers are randomly organized. When the temperature is increased at 55 °C, it is seen by DSC that the two transitions coalesce, and by AFM that the fibers start becoming longer and more parallel. When the annealing temperature is 70 °C, only one transition is seen by DSC, and the fibers are well organized and aligned. However, when Ta = 85 and 100 °C only a very weak and broad transition or no transition at all is seen by DSC, when the AFM pictures highlight even longer and aligned fibers, but all organized in a surface rather than in a volume. This suggests that the DSC transitions might not be correlated to the formation of the fibers themselves, but rather to an interfiber interaction. Infrared Spectroscopy Studies. To help elucidating the molecular arrangement and conformation of these aggregates, VT-FTIR studies have been performed and the outcome is 8534
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°C than on the polymers that were annealed at 85 or 100 °C (Figure 11a−c). This was also reflected in the metabolic activity of the cells which was measured after 3, 6, and 9 days of culturing (Figure 11d). Noticeably, after 3 days the cells on the polymer annealed at 40 °C showed lower activity than the cells cultured on the polymers annealed at 85 and 100 °C. These results were also found at day 6, but became less pronounced after 9 days of culturing. Furthermore, no differences in the activity of cells grown on the polymer annealed at 85 and 100 °C were found. Additionally, the metabolic activity of the cells became higher in time, which indicates cell proliferation on the materials studied. AFM studies showed that under various conditions the polymer undergoes morphological changes from disorganized short fibers for Ta = 40 °C to longer and more organized fibers when Ta > 40 °C (Figure 8). Although the initial AFM pictures have been taken under air at 21 °C, in contrast to the aqueous environment of the cell study at 37 °C, these results suggest that cell activity is strongly related to the mesoscopic morphology of the polymer. Many studies have shown a correlation between cell behavior and surface morphology at the mesoscopic level,40−46 but to our knowledge this is the first example showing differences in cell behavior on a supramolecular polymeric material that only varies in aggregation morphology via phase separation, without changing the chemical nature. We show that the cells perform better on the surfaces with long fibers than on the surfaces with short fibers. This might be due to a higher contact surface with the longer fibers when the film has been annealed at higher temperatures. In future work the relationship between surface morphology and cell adhesion and activity will be investigated.
°C and, on the contrary to the aliphatic vibrations, decrease in intensity until 100 °C. At 115 °C, these bands broaden and increase in intensity, suggesting a loss of the hydrogen bonding. These spectra exhibit also two C−O vibrations at 1020 and 1096 cm−1. As seen previously, these bands decrease in intensity between 25 and 80 °C, consistent with a weakening of the hydrogen bonding, then increase in intensity between 80 and 90 °C, consistent with a rearrangement and after 90 °C their intensity lower again. In contrast with the PEB and UPy vibrations, these C−O vibrations do not broaden or increase in intensity in the melt. The crystallization processes were also followed in time at different annealing temperatures: 25, 40, and 70 °C. The spectra taken 24 h after the melt and in the melt are shown in Figure 10.
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Figure 10. Infrared spectra of 1 after 24 h of crystallization at different temperatures (25, 40, and 70 °C). The spectrum in the melt is shown for comparison.
CONCLUSIONS In the past few years, our group has made tremendous progresses in the design and understanding of biologically relevant supramolecular polymers that assemble into one-dimensional stacks. We have for example noticed that the incorporation of peptides within the stack allowed for better cell recognition.26 However when thermal analyses of some of these polymers were performed, two melting transitions were observed and unexplained so far.23 Because a good understanding of its molecular and macromolecular organization in the solid state is of vital importance for our future projects and can be particularly relevant to the entire supramolecular and polymers chemistry community, we performed a detailed study of the melting and crystallization behavior of one of these polymers by DSC, AFM, and FT-IR. These results were confronted with cells study. DSC experiments led to the conclusion that polymer 1 exists in two main polymorphs forming a blend. When the annealing time was higher, the two polymorphs separated by self-recognition. AFM pictures showed a clear correlation between the shape of the fibers and the annealing temperatures. It also seemed that, at an annealing temperature higher than the second melting transition seen by DSC, the fibers aligned on a surface, suggesting that the melting transitions result from a three-dimensional organization. While VT-FTIR experiments showed the same trend as DSC and AFM, it gave several clear structural indications of the polymer: first only one tautomer was observed even at the highest temperature, the keto one. Finally, different cell behavior was seen on the same supramolecular material which displayed, depending on the annealing procedure, different fiber lengths at the surface. These results show the importance of the mesoscale structure of a biomaterial on cell behavior.
For further details see the Supporting Information. The C−O vibrations found at 1096 and 1020 cm−1 lose half of their intensity with annealing temperatures of 40 and 70 °C and they completely disappeared in the melt, similarly to what has been seen in Figure 2. The same observation is made concerning the aliphatic vibration at 800 cm−1. The regularity band at 740 cm−1 is absent in the spectrum of the melt. In a general way, the spectra of the polymer annealed at 40 °C resemble very much to the spectra taken at Ta = 70 °C. We have demonstrated that the molecular and macromolecular organizations of polymer 1 depend on the annealing conditions employed. While DSC and VT-FTIR techniques highlighted a similar temperature-dependent trend of the polymer on the bulk, AFM provided direct evidence of the temperature effects on the fiber length, shape and dimensional organization observed from the polymer on a glass surface. Cell Behavior. Finally, in order to investigate the influence of the nanostructures on cell behavior we cultured primary tubular epithelial cells (PTEC)30,31 on films of polymer 1 annealed for 24 h at three different temperatures (40, 85, and 100 °C). The morphology of the cells was studied with optical microscopy after 24 h of culturing. After 3, 6, and 9 days the cell activity was investigated with the resazurin assay, in which the metabolic activity of living cells is measured through conversion of the nonfluorescent resazurin dye (7-hydroxy-3H-phenoxazin-3-one10-oxide) into the highly red fluorescent dye resorufin (Figure 11).31,78,79 Already, after 24 h of culturing, it was clearly visible that fewer cells adhered and spread on the sample which was annealed at 40 8535
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Figure 11. Optical micrographs of primary tubular epithelial cells (PTEC) after three days of culturing on a film of polymer 1 annealed for 24 h at (a) 40, (b) 85, and (c) 100 °C. Scale bars represent 100 μm. (d) Metabolic activity of the cells on polymer 1 annealed at 40, 85, and 100 °C determined by the resazurin assay.
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ASSOCIATED CONTENT
ketone); DPI, polydispersity index; PTEC, primary tubular epithelial cells
S Supporting Information *
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Experimental section, DSC and MDSC spectra, DSC analyses of the polymer isolated by precipitation, determination of the growth dimensionality, AFM phase images, AFM height images, FT-IR studies, and details of the cell experiments. This material is available free of charge via the Internet at http://pubs.acs.org.
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REFERENCES
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AUTHOR INFORMATION
Corresponding Author
*(E.W.M.) E-mail:
[email protected]; Telephone: +31 40 247 3101. Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS The authors thank the financial support from the Council for Chemical Sciences of The Netherlands Organization for Scientific Support (CW-NWO). This research forms part of the Project P1.01 iValve of the research program of the BioMedical Materials institute, cofunded by the Dutch Ministry of Economic Affairs, Agriculture and Innovation. The financial contribution of the Nederlandse Hartstichting is gratefully acknowledged.
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ABBREVIATIONS UPy, ureidopyrimidinone; AFM, atomic force microscopy; DSC, differential scanning calorimetry; VT-FTIR, variable temperature Fourier transform infrared spectroscopy; PEEK, poly(ether ether 8536
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Macromolecules
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