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Microstructure and Self-Assembly of Supramolecular Polymers Center-Functionalized with Strong Stickers Cyril Véchambre,†,‡ Xavier Callies,§ Cécile Fonteneau,∥,⊥ Guylaine Ducouret,§ Sandrine Pensec,∥,⊥ Laurent Bouteiller,∥,⊥ Costantino Creton,§ Jean-Marc Chenal,†,‡ and Laurent Chazeau*,†,‡ †

Université de Lyon, CNRS, F-69621, Lyon, France MATEIS, INSA-Lyon, CNRS UMR5510, F-69621, Lyon, France § ESPCI-Paristech, SIMM, CNRS, UPMC Université Paris 06, 10 rue Vauquelin, Paris Cédex 05, 75231, France ∥ IPCM, Chimie des Polymères, Sorbonne Université, UPMC Université Paris 06, UMR 8232, F-75005 Paris, France ⊥ IPCM, Chimie des Polymères, CNRS, UMR 8232, F-75005 Paris, France ‡

S Supporting Information *

ABSTRACT: This manuscript describes the microstructure of a series of nearly monodisperse poly(n-butyl) acrylate (PnBA) chains center-functionalized with a triurea interacting moiety, able to self-associate by six hydrogen bonds. Different molecular weights have been investigated, from 5000 g·mol−1 up to 80 000 g·mol−1. For molecular weights (Mn) below 40 000 g·mol−1, Xray scattering experiments and atomic force microscopy at ambient temperature clearly show that the systems organize as nanofibers hexagonally packed in oriented domains. This supramolecular structure explains the solid-like gel behavior of these polymers, which is suppressed at high temperature (at an order−disorder transition temperature). For higher molecular weights, nanofibers still form at ambient temperature but their concentration is too low to self-assemble in oriented domains. This is consistent with the reported viscoelastic behavior of these systems described in the companion paper.1

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adhesives.14 Using short (Mn = 3 kg/mol) polyisobutylene chains center-functionalized with a bis-urea group, Courtois et al. demonstrated that interesting adhesive properties could be obtained.15,16 Yet, for such systems, structure and rheological properties should depend crucially both on the strength of the sticker group and on the molecular weight of the polymer chains. In order to systematically study this dependence, we have examined the influence of the molecular weight of the side chains on the self-assembly of PnBA polymers center functionalized with functional groups able to form multiple hydrogen bonds. A first study focused on the linear rheology of PnBA center functionalized with bis-urea xylene groups.17Although the rheological properties were significantly affected by the presence of this associating group, no longrange structure was observed. In the present paper, we focus on the linear rheology and structure of the same system functionalized with a more strongly associating sticker, the triurea moiety. A companion paper analyzes its rheological properties1 while this paper focuses on its long-range structure

INTRODUCTION Since the pioneering work of Lehn, Meijer, and others, supramolecular chemistry has become a common name to describe structures that self-associate through weak bonds such as hydrogen bonds in a nonpolar solvent.2−4 In polymer science, this method can lead to very large structures and/or networks.5−8 In recent years, the study of self-assembly of supramolecular structures has extended from solution to bulk, with a growing interest in tuning polymer properties in the bulk by using molecules that are able to interact.9,10 The nonpolar solvent is then replaced with a nonpolar or weakly polar polymer which still plays the role of the solvent but is also functionalized with interacting groups, so-called stickers by physicists. Depending on the molecular weight of the polymer, and the strength, number, and position of those stickers along the chain, several types of material properties can result, going from high molecular weight polymers reversibly cross-linked by dilute and weak stickers to association of nanofibers or networks formed by strong and concentrated stickers attached to short unentangled chains.11−13 While the majority of the studies on these systems has focused on producing elastic and fully soluble supramolecular networks, some of us have discovered that supramolecular structures could also be used to design supramolecular © XXXX American Chemical Society

Received: July 16, 2015 Revised: October 22, 2015

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DOI: 10.1021/acs.macromol.5b01584 Macromolecules XXXX, XXX, XXX−XXX

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that the AFM observations are independent of the substrate and would not be modified by a modification of its surface chemistry. Rheology. Shear oscillatory tests were carried out on a stresscontrolled rheometer Anton Paar Physica MCR 501 in the linear viscoelastic regime, over a large temperature range. The sample was set up between the two parallel plates of the rheometer (diameter 8 mm, gap 0.8 mm) at room temperature. As the samples flow under stress, the final gap was reached by compressing the sample with the upper plate. Before the temperature ramp, the sample was annealed for 2 h at 75 °C in the rheometer. The viscoelastic moduli G′ and G″ were measured at 1 Hz during a temperature ramp from 25 to 180 °C at 5 °C·min−1. The fast heating rate was chosen to minimize the degradation of PnBA3Us at high temperature. SAXS and USAXS. Small angle X-ray scattering experiments were carried out on the D2AM beamline of the European Synchrotron Radiation Facility (ESRF), France. The energy was set at 8 keV. The two-dimensional (2D) patterns were recorded by a CCD Camera (Princeton Instrument). The distance between sample and detector was set at 880 mm. The exposure time was set at 10 s. Ultrasmall-angle X-ray scattering experiments were carried out on the SWING beamline of the SOLEIL Synchrotron, France. The energy was set at 10 keV. The 2D patterns were recorded by a CCD Camera (AVIEX). The distance between sample and detector was set at 6650 mm. The exposure time was 20 s. 2D spectra for unstretched samples did not show any preferential orientation. Thus, scattered intensity was plotted as a function of the scattering vector q (equal to (4π/λ) sin θ) from the integration of the scattered intensity over the azimuthal angle, from 0 to 360°. For static measurements, all samples were first diluted in toluene and deposited on a thin kapton film; the solvent was slowly evaporated overnight and the sample was then placed in partial vacuum during a week. Since no attempt was made to ensure a fixed mass of the samples, absolute intensity comparisons of the patterns for different samples is not quantitative. For all patterns the kapton signal was subtracted. SAXS spectra of PnBA3U5 sample were recorded under elongation with a homemade tensile testing machine at ambient temperature. The sample was deposited, in its gel state, over a thin uncross-linked SBR film, on which it perfectly sticks. This system was then stretched at a fixed strain rate 4.1 × 10−3 s−1. SAXS measurements were also carried out on static samples at different temperatures in a homemade chamber in which samples were placed between two kapton thin films. Spectra were recorded from 25 °C up to the temperature of the order−disorder transition, with a temperature ramp of 5 °C·min−1. FTIR. Infrared spectra were recorded on a Nicolet iS10 spectrometer between KBr windows. The temperature was controlled with a heating device (P/N21525) from Specac.

characterized by AFM and small-angle X-ray scattering (SAXS). The molecular weight of the polymer was varied from 5000 to 85 000 g·mol−1 with the same tri urea group in the center of the chain, resulting in both an increase in molecular weight of the polymer and the dilution of the stickers. Using atom transfer radical polymerization (ATRP) from the center moiety results in a precise control of the polymer molecular weight in both arms of the sticker.

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EXPERIMENTAL SECTION

Materials. The synthesis of the supramolecular polymers was previously described by Fonteneau et al.18 Briefly, supramolecular polymers were synthesized by a functional initiator approach. First a self-assembling compound bearing three urea groups and two carbon− halogen moieties was prepared. In a second step this compound was used as initiator for the ATRP polymerization of n-butyl acrylate. A schematic structure is shown on Figure 1. The ATRP procedure allows

Figure 1. Schematic PnBA−triurea−PnBA chemical structure. to control the molecular weight and dispersity of the poly(n-butyl acrylate) chains. All the resulting supramolecular polymers have a narrow distribution of molecular weights (Đ < 1.3). The molecular characteristics of the supramolecular polymers studied in this paper, which we will refer to as PnBA3UM with M the molecular weight of the polymer in kg/mol, are reported in Table 1.

Table 1. Macromolecular Characteristics of the Studied Supramolecular Polymers samples

Mn (g/mol)

Đ

DP

sticker wt fraction Φ (%)

PnBA3U5 PnBA3U8 PnBA3U12 PnBA3U18 PnBA3U23 PnBA3U40 PnBA3U85

5200 8000 12000 18000 22600 40000 85000

1.3 1.2 1.2 1.2 1.3 1.3 1.3

36 58 89 140 172 306 660

6.4 4.3 2.7 1.9 1.5 0.9 0.4

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RESULTS The frequency and temperature dependent rheological properties of the different systems in the range −30/+50 °C have been extensively studied in the companion paper of Callies et al.1 and will only be highlighted here. At room temperature, the PnBA3U supramolecular polymers behave as a gel at low frequency when their molecular weight is below 40 kg mol−1 while they flow for higher molecular weights. To extend the temperature range of the rheological characterization, the shear modulus of PnBA3U5, PnBA3U8, PnBA3U12, and PnBA3U18, was monitored by dynamic mechanical analysis using a temperature ramp of 5 °C.min−1, at a frequency of 1 Hz and with a low shear amplitude γ = 2% . Results are plotted on Figure 2 as a function of the temperature. For the whole series, G′ shows a pseudoplateau at room temperature, the height of which increases with decreasing Mn, i.e., with increasing sticker weight fraction Φ. In the pseudo plateau domain, the G′ value is several decades larger than the values measured in the same conditions with an unfunctionalized PnBA with a molecular weight of 18 kg mol−1, as shown

Atomic Force Microscopy (AFM). AFM experiments were performed with a Di3100 AFM apparatus connected to a Nanoscope V scanning probe controller (VEECO Instruments, Plainview, NY). All images were obtained at ambient temperature in tapping mode using silicon cantilevers (SSS-NCH tips from Nanosensor) with a nominal spring constant of 42 N·m−1 and a high resonance frequency of 330 kHz. The average scanning speed was less than 20 μm·s−1. The sample was diluted in toluene, a droplet was then placed on a hard surface (indifferently a mica or a silica wafer) and the solvent was slowly evaporated overnight; then the sample was submitted to a partial vacuum during a week to allow complete solvent evaporation. The vertical z resolution was about 0.5 nm, while the x−y resolution was about 5 nm or less. For all samples, at least three different areas were analyzed to ensure image repeatability. Height images (not shown) indicate a flat surface regardless of molecular weight with a maximum difference in height of 5 nm for a square image of 1 μm. The observations were found to be the same on silica or mica wafer substrates. Moreover, given the sample thickness which is above 100 μm, i.e., much larger than the characteristic length of the microstructure discussed in this manuscript, one can reasonably conclude B


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Figure 4. left) 2D image of the SAXS spectrum of sample PBA3U-18, right) Scattered intensity as a function of q for samples PnBA3U5 (blue □), PnBA3U8 (red ○), PnBA3U12 (green △), PnBA3U18 (orange ◇), and PnBA3U23 (gray ×). The arrow indicates the presence of a second Bragg peak (the second bragg peak of PnBA3U8, PnBA3U12, and PnBA3U18 is more clearly visible in Figure S1).

Figure 2. Shear modulus as a function of the temperature for different supramolecular polymers: (blue ■) PnBA3U5, (red ●) PnBA3U8, (green ▲) PnBA3U12, and (orange ◆) PnBA3U18 (filled symbols represent G′ while empty symbol represent G″). For comparison, data calculated for a linear PnBA18 sample from ref 19 (× symbols, Mn = 18 kg mol−1) and data for the PnBAX5 sample extracted from ref 17 (+ symbols, Mn = 5400 g·mol−1) are also reported (full line for G′, dashed line for G″).

as a function of the scattering vector, q, in Figure 4 (right). An intense Bragg peak is observed for samples up to a molecular weight of 22 600 g·mol−1 and reveals a long-range order with a well-defined characteristic distance. For the PnBA3U40, an unresolved peak can be observed (Figure S1) but is so small that it will be disregarded in the following analysis. No peak can be observed for PnBA3U85, i.e., for Mn > 40 000 g·mol−1. The scattering vector at the Bragg peak q* is plotted on Figure 5 as a function of sticker volume fraction, Φ. These data

in Figure 2 on which the data extracted from ref 19 for this linear PnBA are also reported. This difference indicates the presence of the supramolecular structure made of strong selfassociations of the triurea cores. A rapid drop of G′ is however observed at a temperature which increases with Φ, and reflects the weakening or disappearance of this structure. Note that for PnBA3U8, the drop in G′ occurs in two steps suggesting some first level of disorganization before the long-range order is lost. The temperature at which G′ and G′′ cross, usually taken as the transition temperature, Tc, between a viscous fluid behavior and a gel like behavior, is reported on Figure 3 as a function of the

Figure 5. Log−log plot of the Bragg peak (q*) as a function of the sticker volume fraction (Φ). The line is the curve fitting with a powerlaw (R2 the square root correlation coefficient).

can be fitted by a power law relationship with a power 0.5, which is generally found for scattering of aligned 1D objects when their concentration increases.21,22 Figure 6 shows the scattered intensity obtained from USAXS and SAXS experiments for PnBA3U5, PnBA3U18 and PnBA3U85 as a function of the scattering vector. A correction factor has been applied to the SAXS data to superimpose them to the USAXS data in their common q domain. The scattered intensity curve of PnBA3U85 shows different regimes: a q−2.3 dependence is found for q below 0.2 nm−1, and a q−1 dependence above, followed at large q, by a rapid drop in scattered intensity. In the case of PnBA3U5 and PnBA3U18, the presence of the Bragg peak at large q is still present, but there is no q domain with a q−1 dependence. In addition, a q−3 dependence is observed for q below 0.3 nm−1. The difference between the exponent(−2.3) at small Φ and the one at large sticker concentrations (−3) may be related to the presence of the Bragg peak, i.e., of ordered domains. At large q, the q−1 dependence observed for PnBA3U85 indicates the presence of rod-like scattering entities with characteristic dimensions of the

Figure 3. Order−disorder transition temperature as a function of the sticker volume fraction ϕ. (◇) Tc, determined from the G′ − G″ crossover, and (red □) Tx, determined from the Bragg peak disappearance.

sticker volume fraction. Note that using strictly the criterium of the G′-G″ crossing, Tc deduced for PnBA3U8 is equal to 168 °C (i.e., very close to Tc of PnBA3U5); this temperature is in this case not related to the first modulus drop, which occurred at a lower temperature, around 130 °C. Tc increases nonlinearly with ϕ, as is the case for low molecular weight gelators (LMWG)20 (see Figure S5 in Supporting Information). 2D isotropic spectra were obtained from SAXS and USAXS for all the samples (see for instance Figure 4 left), indicating that the supramolecular structure is isotropic, at the scale of the X-ray beam (0.3 × 0.3 mm2), in the explored length scale range (2−100 nm). The corrected scattered intensity measured by SAXS, averaged over all the azimuthal angles (360°) is plotted C


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as for the other samples (Figure S2), probably because of their very different mechanical properties. Interestingly, an increase in the average distance between rods with increasing molecular weight of the polymer is directly measured on the AFM images. This characteristic distance is almost the same as the distance extracted from the Bragg peak in SAXS (Table 2), suggesting indeed a packing of cylindrical objects. To investigate the abrupt decrease in shear modulus observed on Figure 2, SAXS experiments were performed as a function of temperature with a temperature ramp of 5 °C.min−1. For samples up to a molecular weight of 20 000 g· mol−1, the Bragg peak remains visible up to a critical temperature Tx above which it rapidly disappears (as illustrated in Figure 9 for PnBA3U5). The peak position does not shift, and its amplitude does not change much up to its disappearance. The critical temperature Tx where the peak vanishes can be compared with the transition temperature, Tc, previously deduced from the rheology measurement. As shown on Figure 3 both are remarkably superimposed, suggesting that the temperature at which the material loses its gel-like behavior is closely related to a loss of long-range order in the supramolecular structure (note that for PnBA3U8, the agreement between Tx and Tc is even better if Tc is taken as the temperature of the first modulus drop, i.e., 130 °C,). This temperature also corresponds to a decrease (but not a disappearance) of hydrogen bonds as detected by FTIR spectroscopy (see Figures S3 and S4). SAXS experiments have also been performed in situ during tensile tests with sample PnBA3U5 (cf. details in the Experimental Section). Regardless of the azimuthal angle, the position of the Bragg peak does not change with stretching ratio λ = l/l0 where l0 is the initial gauge length. Nevertheless, its intensity is anisotropically changed as shown on Figure 10a, and becomes more intense in the direction perpendicular to the stretching direction, suggesting the orientation of the packed nanofibers domain in the tensile direction. Herman’s orientation function, ⟨ P2 cos(φ)⟩, generally referred to as the order parameter S for liquid crystals, was calculated to quantify this phenomenon, following

Figure 6. Log−log plot of the scattered intensity versus the scattering vector for PnBA3U5 (□), PnBA3U18 (×), and PnBA3U85 (Δ). The dotted line is a fit of the PnBA3U85 data with the scattering equation for rod-like entities.24

order of several nanometers; in the low q domain, the power law dependence between −2 and −3 suggests that these entities are in concentrations large enough to interact with each other, and/or form a percolating network and/or a hierarchical structure.23 The absence of a q−1 domain for the largest concentration of triurea cores might be due to an overlapping of the scattering of the individual entities with the scattering of the larger objects they form. At this point, the interpretations of these (U)SAXS data need confirmation by the more direct but not averaging over a large volumeobservation by AFM. 1D objects that could indeed be surface images of nanofibers, are visible on the AFM images of all the samples for which a Bragg peak is observed (cf. Figure 7 and Figure S2). In PnBAU5, PnBAU8, and PnBAU12 samples, nanofibers are organized in bundles in which nanofibers are parallel to each other; the higher the concentration of self-associative units, the better defined the bundles of nanofibers are for the largest sticker weight fraction, these bundles even seem to form long entangled ropes at larger scale. This multiscale supramolecular structure is oriented locally (in the direction of the bundle) at the scale of the nanofibers, but is isotropic at scales above a few micrometers; this is therefore consistent with the isotropic SAXS and USAXS patterns previously presented. For Mn > 20 000 g·mol−1 (Figure 8), rod-like objects are still present (in PnBA3U40 and PnBA3U85) but do not show a parallel packing and, in PnBA3U85, seem randomly dispersed. Note also that it was technically impossible to obtain an AFM image of those two samples (Mn > 20 000 g·mol−1) at a magnification as high

90

3⟨cos2(θ )⟩ ⟨P2cos(φ)⟩ =

∫0 I(φ) cos2(φ) sin(φ) dφ 90

∫0 I(φ) sin(φ) dφ

2

−1 (1)

where φ is the azimuthal angle from the direction of the maximum intensity, and θ is the scattering angle. Here ⟨cos2 θ⟩

Figure 7. AFM phase images of the PnBA3U5 at two different magnifications. D


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Figure 8. AFM phase images of sample PnBA3U40 (left) and sample PnBA3U85 (right).

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DISCUSSION AFM images (for all the samples) as well as SAXS results dependence of q* with Φ0.5 and for PnBA3U85 the q−1 dependence of the intensityshow that the triurea-functionalized polymers organize as nanofibers in all the samples. In addition to the Bragg peak previously discussed, one can also note on the SAXS scattering curves of PnBA3U8, PnBA3U12 and PnBA3U18 (see arrows in Figure 4, and Figure S1) a second Bragg peak at a q value equal to √3q*; this additional peak is characteristic of an hexagonal packing of these nanofibers, the first Bragg peak being the signature of the (100) distance and the second of the (110) distance. Assuming that all the samples showing a Bragg peak contain nanofibers hexagonally packed, i.e., for Mn < 23 000 g·mol−1, one can deduce the average distance between nanofibers DSAXS from eq 2:

Table 2. Average Distance between the Nanofibers Obtained by AFM and SAXS sample designation

PnBA3U5

PnBA3U8

PnBA3U12

PnBA3U18

DAFM, nm DSAXS, nm

6±1 7±0.3

9±1 8.5±0.3

10±1 9.7±0.3

12±1 12±0.3

DSAXS =

4 2π × 3 q*

(2)

The calculated values are reported in Table 2 and are in excellent agreement with DAFM, in spite of the large experimental error inherent to the AFM technique. The dependence of this distance on the core weight fraction to the power −0.5 suggests that most of the triurea moieties participate to the nanofibers formation. One can therefore estimate the radius of these nanofibers hexagonally packed from eq 3:

Figure 9. SAXS of PnBA3U5 at different temperatures.

can be approximated as 1. The orientation parameter, ⟨P2 cos φ⟩, should take a value of 1 for scattering entities perfectly parallel to the elongation direction. Its variation with stretch λ is plotted on Figure 10b and shows the progressive orientation of the nanofibers domain.

Figure 10. (a) Intensity at the Bragg peak as a function of the azimuthal angle for different values of λ (1.35, 1.7, and 2.1) for PnBA3U5. (b) Orientation parameter ⟨P2 cos φ⟩ as a function of λ for the PnBA3U5. Dotted line represents the predicted affine orientation of rigid nanofibers. E


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Macromolecules ⎛ 3 ⎞1/2 × (Φ)⎟ Rc = D⎜ ⎝ 2π ⎠

The generation of nanofibers by supramolecular chemistry has received a great deal of attention in the past decades as it is a promising way to obtain peculiar properties in materials (conductivity, rheology, mechanics, optics).25 In particular, many studies have been devoted to self-associative small molecules (low molecular weight gelators (LMWG)) where an organization as nanofibers was expected in solution.26−28 In these systems, the self-assembly is the result of phase separation (solvophobic effect) and is stabilized by physical interactions pistacking and/or H-bonds. After demonstrating the possibility for these small molecules to self-assemble and form anisotropic objects, many authors were interested to increase the size of these molecules, obtaining in this way supramolecular polymers.29 For example, supramolecular polymers forming nanofibers involve conjugated polymers, which assemble by π−π interactions, and are potentially interesting for their transport properties.30 Reports of similar supramolecular assembly resulting from hydrogen bond interactions is also relevant for polymer gels or elastomers: indeed, these systems can combine self-healing properties due to the reversibility of the bonds creating the gel, and specific mechanical properties originating from the formation and organization of the supramolecular nanofibers.31−34 In our systems, a strong correlation is indeed found between the gel-like behavior and the presence or absence of relatively large domains of oriented nanofibers. At high molecular weight, i.e., above 20 000 g·mol−1, at ambient temperature or above, the nanofibers are present but diluted ; their concentration, and the interactions between them are too weak to lead to correlated motions of these objects.35 As a result, the material behavior is characteristic of that of an entangled polymer melt. Its elastic modulus value is however much higher than the modulus of the corresponding polymer filled with rods with the same aspect ratio (around 50).36 This is likely due to the fact that the system cannot be seen as rod-like fillers in a polymer matrix as all the polymer chains of the matrix are grafted on these fillers. For such systems, from the author’s knowledge, there is no straightforward model to predict the measured modulus. Below 20 000 g·mol−1, a gel-like behavior is observed at low frequency despite the fact that the behavior of the unfunctionalized polymer should be that of an unentangled polymeric fluid (Mc = 2Me, Me ≈ 23 000 g·mol−1); this gel behavior arises from the dense and regular packing of the nanofibers in large oriented domains. This packing, usually observed for concentrated rods suspensions, is mainly controlled by the different interaction energies (mainly the rod−rod interactions) and the balance between the orientational entropy of the nanofibers and the configurational entropy of the side chains.37−40 When deformation is applied, the nanofibers inside these domains must move in a correlated fashion, so that the rheology of the system becomes that of large colloidal objects with a mechanical response resulting from the difficulty for these domains to orient in the flow direction.35 More than the deformation of these clusters, the macroscopic deformation is mainly due to the movements of these clusters in relation the one to another, which should explain the weak orientation of the rods with the deformation. Such behavior is sensitive to temperature, as its increase modifies the different thermodynamic contributions involved in the nanofibers packing, and leads to the weakening of the H-bonds, which in turn reduces the average nanofiber length and concentration. At still higher temperatures the long-range order vanishes and is accompanied by a loss of the gel-like behavior.

(3)

The slope deduced from the linear plot of DSAXS (more precise than DAFM) as a function of ((√3/2π) × Φ)−0,5 can be used to calculate a radius of 0.92 nm (cf. Figure 11). Such value is remarkably consistent with the length of a triurea unit, the extended length of which should be about 1.9 nm (estimated by molecular modeling, cf. Figure 1).

Figure 11. DSAXS as a function of (√3/2π) × Φ)−0,5, with the slope from which the radius of the triurea rods is deduced.

The absence of Bragg peaks obviously makes such analysis impossible for PnBA3U40 and PnBA3U85; however, AFM images (cf. Figure 8) indicate that nanofibers are also formed in these samples. Assuming that the radius of the fibers does not depend on Φ, it is tempting to fit the USAXS data with the theoretical scattering function of randomly oriented and diluted nanofibers of radius Rc (see Figure 6).24 The fitting parameters are the intensity factor, and the rod length. Fitting is only relevant for PnBA3U85 for which a q−1 dependence is clearly visible and no Bragg peak is present in the relevant q domain. The rod length obtained from the best fit is around 50 nm, which is actually consistent with the corresponding AFM images. This length seems however largely distributed, and the nanofibers show also some flexibility on the AFM images, which is not taken into account in our fit. The nanofibers length of the other samples (i.e., with higher sticker volume fraction) is very difficult to estimate from AFM images but seems to slightly increase with sticker volume fraction. When hexagonally packed and stretched uniaxially, the nanofibers orient in the tensile direction (Figure 10a). The evolution of the orientation function with the stretching ratio of PnBA3U5 is compared with the prediction for the progressive orientation of rigid rods having no extensibility and whose rotation around their barycenter follows the affine deformation of an incompressible matrix. Such calculation corresponds to an upper bound of the orientation function and as shown on Figure 10b, greatly overestimates the results. To sum up, the microstructure of the different systems is clearly multiscale, isotropic at the sample scale but oriented at the more local scale, and made of semiflexible nanofibers. When the molecular weight is below 20 000 g·mol−1 these nanofibers are packed hexagonally, in domains of several micrometers, which slightly orient with stretching. F


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An interesting comparison can be made on Figure 2 between PnBA3U5 and another PnBA center functionalized supramolecular polymer (PnBAX5) made of the same molecular weight with a weaker sticker containing only two urea functions (cf. ref 17 from which data have been extracted). The order disorder transition of PnBAX5 occurs at ambient temperature, and SAXS and AFM performed at this temperature did not show any evidence of structure similar to that of PnBA3U5. Thus, for a molecular weight below 20 000 g·mol−1, the strength,of the stickers, by controlling the nanofibers formation, their length and concentration, may control the temperature below which the system organizes as oriented domains of nanofibers and adopts a gel like behavior. For this reason, rheological and mechanical characterization of the PnBAX samples described in ref 17 in a temperature domain well below the ambient temperature should lead to the same observations as those made at and above the ambient temperature for PnBA3U samples.

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was financially supported by the ANR SUPRADHESION program (Project ANR-10-BLAN-0801). The authors are indebted to the CLYM (Centre Lyonnais de Microscopie http://www.clym.fr) for the access to the Di3100 atomic force microscope, particularly to D. Albertini for his very helpful technical support; they are indebted to the SOLEIL synchrotron for providing beam time on the SWING beamline, to G. Stocklet for his technical support, to the ESRF synchrotron for providing beam time on the D2AM beamline, and to R. Seguela for his technical support.

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CONCLUSION This manuscript reports the microstructure of triurea center functionalized poly(butyl acrylate) (PnBA) chains with a precisely controlled molecular weight. Various molecular weights have been investigated, from 5000 g·mol−1 up to 80 000 g·mol−1. For Mn below 20 000 g·mol−1, X-ray-scattering experiments and AFM images clearly show that at room temperature, the systems organize as nanofibers hexagonally packed in micrometer-size oriented domains. The existence of this supramolecular structure of parallel nanofibers explains the gel behavior of these polymers, resulting from the strongly correlated movements of the nanofibers inside the domains. Such behavior is suppressed at high temperature, when the balance between the different thermodynamic contributions at the origin of the ordering of the nanofibers (orientational entropy of the nanofibers, configurational entropy of the side chain, and the interactions energies) is modified, leading to the disappearance of the long-range order. This rheological behavior is also consistent with the very weak level of orientation of the nanofibers in the flow direction during stretching, as this one depends on the weak interaction between the side chains. For systems with molecular weight above 20 000 g·mol−1, nanofibers still form but their concentration is too low to promote their organization in oriented domains. This is then consistent with the viscoelastic fluid like behavior of these systems. These conclusions made at ambient temperature with PnBA polymers center functionalized triurea units, might be generalized to other center functionalized PnBA of same molecular weight at another temperature, since the strength of the stickers indirectly controls the temperature at which they can assemble as oriented domains of nanofibers.1

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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/acs.macromol.5b01584. SAXS plots, AFM images, FTIR spectra, and thermodynamics (PDF)

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REFERENCES

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