Enhanced Glassy State Mechanical Properties of Polymer

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Enhanced Glassy State Mechanical Properties of Polymer Nanocomposites via Supramolecular Interactions

Amir Hashemi a, Nicolas Jouault a,b*, Gregory A. Williams c, Dan Zhao a, Kevin J. Cheng c, Jeffrey W. Kysar d, Zhibin Guan c*, Sanat K. Kumar a*

a

Department of Chemical Engineering, Columbia University, 500 W. 120th St, New York, NY

10027, United States b

Sorbonne Universités, UPMC Univ Paris 06, CNRS, Laboratoire PHENIX, Case 51, 4 place

Jussieu, F-75005 Paris, France c

Department of Chemistry, University of California, Irvine, CA 92697-2025, United States

d

Department of Mechanical Engineering, Columbia University, 500 W. 120th St, New York, NY

10027, United States

KEYWORDS: Nanocomposites, supramolecular interactions, mechanical properties, hydrogen bonding, reinforcement

ACS Paragon Plus Environment

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ABSTRACT It is now well accepted that the addition of nanoparticles (NPs) can strongly affect the thermomechanical properties of the polymers into which they are incorporated. In the solid (glassy) state, previous work has implied that optimal mechanical properties are achieved when the NPs are well dispersed in the matrix and there is strong interfacial binding between the grafted NPs and the polymer matrix. Here we provide strong evidence supporting the importance of intermolecular interactions through the use of NPs grafted with polymers that can hydrogen bonding with the matrix which clearly show that their yield significant improvements in the measured mechanical properties. Our finding thus supports the previously implied central role of strong interfacial binding in optimizing the mechanical properties of polymer nanocomposites.


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Mixing inorganic nanoparticles (NPs) with an organic polymer is a commonly used strategy to create polymer nanocomposites (PNCs) with improved macroscopic properties. It’s now well accepted that the spatial organization of NPs controls such improvements and that the optimal NP structure may be different depending on the property of interest.1-3 For example, polystyrene (PS) grafted silica NPs maximize the mechanical properties in the rubbery state4 (i.e. above the glass transition temperature, Tg, of the material) when they assemble into fractal-like percolated structures. In contrast, well-dispersed NPs are mechanically most efficient in the glassy state (i.e. below Tg).5,

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Going further, this latter work showed that, strong interfacial

binding between the grafted NPs and the polymer matrix is another requirement to optimize the Young’s modulus (YM), yield stress (YS) and ductility of the PNCs. While the interfacial binding in this particular example was due to the entanglement of the grafts and matrix chains, we conjecture that the use of a grafted polymer capable of supramolecular binding to the matrix may be a powerful tool to not only control the NP dispersion state, but more importantly the binding between the NPs and the polymer matrix. This should therefore significantly improve the glassy-state mechanical properties of these hybrid materials, a hypothesis we test here. Supramolecular polymers are materials which are formed not through covalent bonds, but where monomers “bond” through non-covalent interactions,7 e.g., hydrogen bonds,8-11 metalligand coordination,12,

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π-π stacking14 or dynamic covalent bonds (such as imine groups) to

form chain-like structures. By definition, the bonds that make up this material are dynamic, but the lifetime of a bond can be strongly renormalized upward due to the fact that most groups that can bond are paired – thus, even though a single bond breaks fairly frequently, the unpaired groups cannot find other partners to bind with. Thus, they are forced to reconnect with their previous partners. Such repeated pairing causes a dramatic increase in the lifetime of the


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supramolecular bond. The use of such reversible bonds helps designing responsive and adaptive materials15, 16. In this paper, we focus on hydrogen bonding nanocomposite systems: in particular silica NPs grafted with polymers functionalized with 2-ureido-4-pyrimidinone (UPy) units, and then incorporated into a polymer matrix.17, 18 Excellent NP dispersion was achieved presumably due to the existence of multiple hydrogen bonds between the NPs and the polymer matrix, a concept that has been used in previous works. For example, Heo et al.19 used gold NPs grafted with 2.8 kg.mol-1 (noted K in the future use) P(S-r-2VP) chains (grafting density of 1.7 chains/nm2) in a 24.4 K P(S-r-4VPh) matrix. While the H-bonding between the OH group on the vinyl phenol moieties groups in the matrix and the nitrogen of the ligand grafted on the NPs ensures good miscibility, the consequences of this improved miscibility on mechanical properties has not been investigated. Here, we investigate the glassy state mechanical behavior in the glassy state of UPy grafted silica NPs incorporated in polymethylmethacrylate (PMMA) matrices. We find that the mechanical response in the glassy state surpass those obtained previously,5 where there was no specific interactions between the grafts and the matrix chains, except for transient chain entanglements. Our results clearly show that both good NP dispersion and strong adhesion between the NP and the matrix are critical to optimizing the glassy state response of PNCs.

NPs characterization in solution. For this study, AS-40 silica NPs (Grace Davison) were grafted with ligands end-functionalized with UPy units (see Figure 1).10 We used two types of UPy grafted silica NPs named UPy 1 and UPy 2. UPy 1 is composed of a UPy unit grafted to


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the NPs via a short alkane (10 carbons) linker; the grafting density is 0.34 chains/nm2 (961 chains/NP). UPy 2 is grafted to the NP via a poly(n-butylacrylate) (PBA) chain (n=60) linker with a grafting density of 1.2 chains/nm2 (3393 chains/NP). All details about the synthesis and the grafting process can be found in the Supplementary Information (SI) and the organic content grafted to the NPs is shown in Table 1. The UPy units strongly interact complimentarily through strong quadruple hydrogen bonds (2 N-donors and 2 H-acceptors10, 20) with other UPy groups.10, 21-23

In nonpolar solvents (such as chloroform), the UPy group has a high dimerization constant,10

implying that two adjacent groups from the same NP or from two different ones can strongly dimerize in this type of solvent. To prevent NP aggregation due to such interactions, we used a very polar solvent, dimethylformamide (DMF) to ensure good dispersion and stability of the system. To clearly show the effect of the UPy chains on the final mechanical properties we also studied samples with similar size ungrafted (bare) NPs (from Nissan Chemicals, initially dispersed in isopropanol). Note that these bare NPs are partially treated by the manufacturer to ensure better solution stability, but silanol groups are still available at the surface. In addition, negative surface charges ensure strong electrostatic repulsion between the NPs in DMF, adding to the solution stability of the NPs. Insert Figure 1 Dynamic and Static Light Scattering (DLS and SLS) and Transmission Electron Microscopy (TEM) were used to characterize the grafted NPs. From DLS measurements in DMF, the normalized time autocorrelation function g(2)(q,t) is measured as a function of the scattering vector q=(4π/λ)n0sin(θ/2) (θ is the scattering angle which has been varied from 30° to 140°). NPs solutions are characterized by a single relaxation mechanism, i.e. one population in size, with a characteristic time τ inversely proportional to q2, indicating a diffusive process (see


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Figure S1 in SI). Then, the extrapolation of the apparent diffusion coefficient, Dapp = 1/(τq2), to q=0 gives the mutual diffusion coefficient D0 from which the hydrodynamic diameter dh,i can be derived using the Stokes-Einstein relation. Note that the Dapp exhibits a linear dependence on q2 due to the polydispersity of the NPs. Table 1 presents the results for the different NPs. For comparison with the diameter determined from TEM images Table 1 also shows the hydrodynamic diameter obtained from the number size distribution, dh,n, at one angle. The discrepancy between dh,i and dh,n comes from the polydispersity of the NPs (since the scattering intensity I in DLS scales as dh6, larger objects contribute more to the scattering and thus shift the intensity size distribution to higher diameters). From the TEM images analysis the AS-40 NP diameter has been estimated to be 24 nm (confirmed by DLS and close to the size provided by the supplier). Then, in DMF, we measured a hydrodynamic diameter dh,n (silica core and polymer shell) of the grafted objects to be 28 nm for the UPy 1 and 48 nm for the UPy 2. By difference, we deduce the thickness e of the grafted layers: 2 nm and 12 nm for UPy 1 and UPy 2, respectively, in agreement with the length of the grafted chains. Complementary SLS measurements (see Figure S2 in SI) show no increase in scattering intensity at low scattering vector q and a slight correlation peak coming from the NP-NP repulsion, indicating that no aggregation occurs in DMF for UPy functionalized NPs. The grafted NPs are thus individually dispersed in solution prior to the preparation of the nanocomposites. Insert Table 1 Three different classes of PNCs were prepared by mixing the UPy-grafted NPs with PMMA matrices of three different molecular weights (Mw=98.4 K with PDI=1.16, Mw=193 K with PDI=1.17, Mw=387.9 K with PDI=1.07) in DMF. These solutions were spin coated onto a mica sheet to form thin PNC films. The PNC sample was then removed from the mica by


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floating it in water and directly placed (i) on a TEM grid for NP observation; (ii) silicon wafers for thickness measurement or (iii) bulge test window for mechanical measurements. The thicknesses, ≈100 nm (see Table S1 in SI for exact values for each sample), were determined by ellipsometry and it has been previously shown that mechanical properties measured by bulge test are thickness independent in this range.5,

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The mechanical setup will be described in detail

below. NP dispersion in PMMA matrices. Figure 2 presents TEM images of different NP/PMMA PNCs. Note that these samples were not annealed. We first consider the effect of grafted chains on NP dispersion in a 193K PMMA matrix (Figure 2a). Bare NPs (5.2 wt%) show uniform dispersion as expected since the PMMA can H-bond with the OH groups on the silica surface. UPy 1 (5 % by weight silica) arranges in small aggregates, due to the alkane chain and potential π-π stacking between UPy:UPy dimers while UPy 2 NPs are well dispersed. In that case, a bulky diisopropyl phenyl group and PBA linker were introduced to the UPy moiety and, as shown in a previous study21, such bulky substituents can effectively prevent such π-π stacking between UPy:UPy dimers. The dispersion state has been further quantified by the calculation of the radially averaged correlation function C(r), corresponding to the two-point correlation of the TEM image pixels as a function of distance (Figure 2)25. The C(r) gives information about the typical feature size in one image. The initial slope and first zero give the NPs size and distribution: around 23 nm for bare NPs and 30 nm for UPy NPs. Good NP dispersion manifests itself as a depletion hole in the C(r) coming from the repulsion between the NPs, as for the bare NPs. A deviation from this initial slope, as observed for the UPy 1 system, is due to the existence of small aggregates (of around 50 nm here). Insert Figure 2


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We now examine the role of PMMA Mw on the dispersion state of the UPy 2 grafted NPs. Figure 2b shows TEM images for three different molecular weights of the PMMA matrix: 98.4 K (left), 193 K (middle) and 387.9 K (right) PMMA filled with 6.3 wt% of the NPs. Phase separation occurs for low Mw 98.4 K (i.e., poor NP dispersion), while good dispersion is observed for 193 K (with few aggregates) and 387.9 K. The C(r) analyses confirm these observations. For 387.9 K, the negative values (depletion region) in the C(r) at intermediate r indicate repulsion between NPs, leading to the good dispersion observed on the TEM image. For 98.4 K multiple correlation peaks are clearly visible, indicating an apparently crystalline NP order, also confirmed by the image. The 193 K is intermediate, illustrating the “coexistence” between individual NPs and aggregates. Before annealing, the NP dispersion state may originate from the role of the solvent from which the PNCs are processed (here DMF), as well as the polymer matrix Mw, which modulate the polymer/NP interaction and thus the final spatial NP dispersion26. Note that annealing the UPy 2/98.4 K samples for 3 days at 150 °C, where hydrogen bonds are very dynamic, leads to good NP dispersion (Figure 2c, right and also confirmed by the C(r) calculation). At equilibrium, the UPy NPs are well dispersed in PMMA matrix, independent of the Mw. We conjecture that this good NP dispersion is due to the multiple hydrogen bonds between the UPy units and the PMMA. Indeed, Meijer and coworkers27 showed on oligo(ethylene glycol) (OEG) substituted UPy that the UPy dimerization constant is reduced by three order of magnitude caused by intramolecular interaction between OEG and the UPy. Similarly, in our case, the UPy units are embedded in an acrylate environment (from PBA linker and PMMA matrix) whose ester groups are significantly more Lewis basic than the ether groups in OEG, ensuring stronger interaction with the UPy. The dominant mode of H-bonding in PNCs


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is thus not between the UPy groups, but rather between the UPy moieties and the acrylate monomers on the matrix chains (see Scheme 1, more on this below). The H-bond strength between one N-H of the UPy unit and a carbonyl group of PMMA is around 11 kJ/mol (2.6 kcal/mol)28 and multiple H-bonds can be formed between the UPy NPs and the polymer matrix at room temperature, making the total NP/polymer binding energy very high. Moreover, regarding the grafting density (√ = 9.3 with N=60 for UPy 2) and the high matrix to brush ratio P/N (P and N denote the degrees of polymerization of the matrix and the grafted chains respectively: P/N=16.4, 32.2 and 64.6 for 98.4 K, 193 K and 387.9 K respectively), autophobic dewetting should occur, leading to NP aggregation6. However, we conjecture that the strongly favorable UPy/PMMA interactions overcome the entropic penalty of high polymer matrix Mw and insure good NP dispersion over a wide range of PMMA Mw. Insert Scheme 1 Glassy state mechanical behavior. To investigate the mechanical properties of the PMMA PNCs below Tg (i.e., at room temperature) we used the bulge test.5,

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Briefly, after spin

coating solutions of NP/PMMA, the non-annealed PNCs are transferred from mica sheets onto a silicon wafer with a rectangular hole in the middle (called bulge window) by immersing it in water (PMMA is hydrophobic). A pressure (using atmospheric air) is then applied to the deposited film through this hole, modifying the film shape (pressure increased steadily at 2 kPa/second until approximately zero stress/strain slope). A scanning laser confocal microscope is used to measure the profile and radius of curvature of the deformed film from which we determine the stress and strain30 through a parabolic modeling of the profile (see Scheme S4 in SI).


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The stress versus strain raw data show a residual stress (i.e. non-zero value of the stress) before deformation, possibly due to the tension induced in the film during the transfer onto the bulge window. A vertical offset has been applied to obtain a zero stress at zero strain (see Figure S6). Then, we deduced the Young’s Modulus YM (initial slope of the stress-strain curve) and the yield stress (σ yield) and strain (γ yield), which correspond to end of the linear regime, and compare it with the pure PMMA matrix. Table 2 reports all the values for the different systems with the percentage increase relative to the neat polymer. In pure glassy polymers the YM is associated with chain resistance to low deformation and the σ

yield

can be correlated to the thermally

activated β-relaxation, i.e. to the local motion of ester side groups allowing chain arrangements and thus energy dissipation. The post-yield behavior, at higher deformation, describes the nonlinear response of the material to larger deformation. YM, σ yield, γ yield and post-yield will be discussed below. Insert Table 2 Insert Figure 3 Figure 3 shows the mechanical response (stress versus strain) for different pure PMMA Mw: (a) 98 K, (b) 193 K and (c) 387.9 K. Pure PMMA samples (i.e. without NPs) exhibit similar behavior for all Mws with an elastic regime at low strain characterized by a linear increase of the stress followed by a maximum (YS) and a stress decrease, indicating softening and failure of PMMA. The YM is similar for all the Mw’s with a value of 2.2+/-0.4 GPa, in good agreement with literature values,24, 31 but large variations in the yield (both σ yield and γ yield, see Figure S7), are attributed to the brittle fracture by the nucleation and breakdown of crazes at room


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temperature (note that the temperature Tβ associated to the β–relaxation is around 25°C for PMMA)32. Since the UPy 2 systems present different dispersion state with changes in PMMA Mw we will focus on these systems below. The addition of UPy 2 grafted NPs (6.3% by weight silica) leads to different mechanical behavior depending on the PMMA Mw. For 98.4 K there is no change in stress/strain response compared to the pure polymer. For 193 K and 387.9 K clear increases in YM and YS are observed. Insert Figure 4 Young’s modulus (YM). In Figure 4 we compare the YM of the UPy 2 PNCs with different PMMA Mw. For PNCs consisting of PMMA matrices with Mw of 387.9 K, 193 K and 98.4 K we measured a 114% increase, a 64% increase and nearly no change (within the error bar), respectively, for the YM relative to the pure polymer (Figure 4c). As shown by the TEM images (Figure 4a), the large increase in YM is related to the perfect individual NPs dispersion in the polymer matrix. The more dispersed the NPs, the larger the increase in YM. Our maximum increases (+114%) are roughly 2 times larger than increases reported for non H-bonded systems in the literature (+52% for PS grafted silica NPs in a PS matrix).5 Similarly, in Figure 5, we compare the nature of the NP/PMMA interaction on YM, i.e., by comparing UPy1, UPy 2 and ungrafted bare NPs in 193 K PMMA. Insert Figure 5 The YM only increases for the UPy 2 system (3.6 GPa, i.e. + 64% increase), for which the dispersion is similar to the ungrafted one (YM = 1.9 GPa, similar to the pure PMMA) and better than the UPy 1 system (YM = 1.7 GPa, slightly lower than the pure one but still in the error bar).


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It’s now commonly accepted that the PNC YM or stiffness increases upon the addition of NPs. Its evolution has been first described with a simple free volume argument, i.e. NPs occupy voids and decrease polymer free volume, hindering the polymer segment motions and a predictive expression has been proposed for the YM determination in the case of well dispersed spherical NPs: YM=YM0.(1+2.5Φ+14.1Φ2), where YM0 is the YM of the pure PMMA and Φ is the NPs volume fraction. One obtains for a predicted YM of 2.4 GPa for 6.3 %wt UPy 2/387.9 K PMMA, which is below the measured YM of 4.7 GPa. In the same vein, basic continuum models, in which only two phases (filler + matrix) exist, also fail to reproduce our YM increase, indicating that the supramolecular interfacial interaction has to be taken into account33. Below Tg, the UPy unit strengthens the bonding between the NPs and the polymer and plays a major role in the mechanical response. Finally, note that similar YM were found after annealing (see Figures S9 in SI), suggesting that the residual solvent (if there is any) in the thin films does not affect its glassy mechanical behavior, as similarly observed in our previous work34. To obtain a more quantitative estimate of the role of H-bonds on mechanical behavior, we postulate that each bond contributes ⁄ , to the YM, i.e. that each bond is worth one unit of thermal energy, and that this is distributed over the diameter of the NP35, 36. If one then uses the fact that the PNC have 6.3% by weight of the NP, that the Si NP cores constitute 36% by weight of the grafted NP (the other 64% represent the UPy functionalized chains), and that the silica core density is 2200 kg/m3, we obtain that there are 7.5 × 10 NPs/m3 in these PNC. A single H-bond between a NP and the matrix, on average, then results in a modulus increase of ~ 0.2 GPa. Given that the YM increases by ~2.5 GPa for the 387.9 K PMMA compared to the pure PMMA (Figure 4), it follows that there are ~13 H-bonds/NP on average between the UPy and the PMMA chains. This number is about 10% of the maximum number of H-bonds (~120)


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that is predicted to occur between UPy and PMMA chains, given that the free energy of association between two UPys and between a UPy and a PMMA are ~11 and 4 , respectively. While we do not understand why only such a small number of H-bonds form between the NPs and the PMMA, these results nevertheless highlight the important role played by these relatively small number of bonding events on NP/polymer miscibility. This final fact is in good agreement with recent ideas of Hooper and Schweizer37 who suggested, based on integral equation theory, that polymer-NP miscibility can only be achieved for intermediate values of NP-polymer attraction. Too little attraction causes direct enthalpically driven phase separation, while too strong an interaction leads to particle agglomeration driven by polymer bridging. Our results, which show minimal NP clustering in many cases (especially after annealing), are thus consistent with this relatively modest amount of NP-matrix interaction. For comparable NP dispersion state, our mechanical results can thus be correlated to the number of H-bonds between NPs and PMMA, which is larger for the UPy 2 than the ungrafted NPs. This is the main result of this paper. The strong reinforcement is correlated to the H-bonding induced connectivity between the UPy unit and the PMMA. A higher grafting density, the presence of the PBA linker (compared to alkane for UPy 1) and the lack of π-π stacking between UPy:UPy dimers in UPy 2 NPs provide better miscibility and enhance this mechanical phenomenon. Yielding and post-yield behavior. The yield is defined as the maximum of the stress-strain curve, i.e. the end of the linear regime and can be seen as the highest stress and strain the sample can sustain before irreversible structural reorganization of polymer chains and/or NPs38. For unfilled samples, the brittle nature of PMMA makes its determination inaccurate since large variations are observed. However we can qualitatively compare the PNCs σ yield with its unfilled equivalent. No change is observed for 98.4 K; while for 193K increases of 49% and 188% for


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bare NPs and UPy 2 respectively are reported and 94% for UPy 2 in 387.9K. As previously observed5 changes in σ yield can be correlated to the nature of the NP/polymer interface. The σ yield increase is only observed for good NP dispersion, i.e., when the surface area to volume ratio is optimized and thus interfacial properties dominate. Then, the interaction between ester groups of PMMA and the UPy unit or silanol groups for bare NPs reduces the chain local motions and thus increases the σ

yield.

Note that from the σ

yield

and γ

yield

one obtains an elastic energy density

E=1/2.(σ yield γ yield) or toughness, corresponding to the area under the curve in the linear regime, of 90, 84.4 and 62.5 kJ.m-3 for 98.4 K, 193K and 387.9 K UPy 2/PMMA, respectively. The UPy 2 systems are stiffer but less tough when improving the NPs dispersion, commonly observed in PNCs. For the post-yield behavior let’s focus on the well-dispersed case of the UPy 2 in 387.9 K PMMA (Figure 3c). At larger deformations, after the YS, pure PMMA shows a stress decrease due to the nucleation, propagation and breakdown of crazes, reducing the local modulus and leading to the final fracture of the sample. On the contrary the UPy 2 shows a constant stress value characteristic of shear banding5 (zone of high strain). Here again, uniform individual NP dispersion and strong interaction with the polymer matrix may explain such behavior by preventing craze nucleation. However, for large deformation, such a hypothesis needs to be confirmed by supplementary data and complementary experiments (uniaxial stretching experiments for instance). In summary a significant increase in Young’s modulus and yield stress in the glassy state has been measured in PMMA polymers containing NPs grafted with supramolecular UPy units. Through the multiple hydrogen bonds between those units and the PMMA, NP/polymer miscibility is achieved and a 114% increase in YM has been obtained for uniform NP dispersion.


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This work provides clear evidence that mechanical reinforcement of the solid-state properties of a glassy polymer requires strong “adhesion” between NPs and polymers in addition to good NP dispersion and brings fundamental and practical evidence of the promising features of supramolecular grafted NPs as fillers in PNCs.


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Figure Legends Figure 1. Chemical structures of the different NPs used in this study. Figure 2. TEM images for different PMMA PNCs: (a) Influence of NPs nature in 193K PMMA matrix: bare silica (left), UPy 1 (middle) and UPy 2 (right). (b) Influence of PMMA Mw in UPy 2/PMMA PNCs: 98.4 K (left), 193 K (middle) and 387.9 K (right). (c) Influence of annealing on 98.4 K PMMA/UPy 2 PNCs: no annealing (left), 4h annealing at 150 °C (middle) and 3 days annealing at 150 °C (right). For all TEM images the autocorrelation functions of pixels, C(r), have been calculated as a function of distance and are shown on the right of each line. Figure 3. Stress as a function of strain for (a) 98.4 K, (b) 193 K and (c) 387.9 K PMMA PNCs containing UPy 2 NPs (red squares). The data are compared to the pure PMMA, i.e. without NPs (black circle). For 193 K (b) we also plotted the results for UPy 1 NPs (green cross). The mechanical measurements have been performed in the glassy state at room temperature using a bulge test (see main text). Figure 4. (a) TEM images of UPy 2/PMMA PNCs for 387.9 K (top), 193 K (middle) and 98.4 K (bottom). (b) Corresponding stress versus strain curves. (c) Young’s modulus obtained from the stress/strain curves (initial slope) of the different systems. Figure 5. (a) TEM images of UPy 2 (top), UPy 1 (middle) and bare silica (bottom) in 193 K PMMA matrix. (b) Young’s modulus obtained from the stress/strain curves (initial slope).


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Scheme Legends Scheme 1. Schematic representation of supramolecular crosslinks between UPy 2 grafted NPs and PMMA matrix.

Table Legends Table 1. Hydrodynamic diameters of the different NPs: dh,i and dh,n are the average diameter by intensity and by number respectively. The discrepancy between dh,i and dh,n comes from the polydispersity of the NPs. Thicknesses e of the grafted chains UPy 1 and UPy 2 are also reported. We also report the organic content grafted to the NPs, corresponding to the mass loss at 800°C using TGA. Table 2. Young’s modulus (YM), Yield stress and strain (σ yield (MPa) / γyield (%)) values of the pure PMMA and PNCs for different polymer Mw’s.


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Figures

Bare NP O-H UPy unit

20 nm UPy 1

~ 2 nm 24 nm UPy unit

UPy 2

24 nm ~ 12 nm

Figure 1.


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a

Bare – 5.2 %wt – 193 K

UPy 1 – 5 %wt – 193 K

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UPy 2 – 6.3 %wt – 193 K

Bare NPs UPy 1 UPy 2

0,4 0,3 0,2 0,1 0 -0,1

0

50

100

150

200

250

Distance (nm) 0,5

UPy 2 - 98.4 K PMMA

UPy 2 - 193 K PMMA

UPy 2 - 387.9 K PMMA

Correlation function C(r)

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98.4 K 193 K 387.9 K

0,4 0,3 0,2 0,1 0 -0,1

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UPy 2 - 98.4 K – No annealing

UPy 2 - 98.4 K – 4 h – 150 °C

UPy 2 - 98.4 K – 3 days 150 °C


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98.4 K - Before annealing 98.4 K - After 3 days annealing

0,4 0,3 0,2 0,1 0 -0,1

0

50

100

150

Distance (nm)

Figure 2.


19

200

250

Nano Letters

30

a

M = 98.4 K w

Stress (MPa)

25 20 15 10 5 0

Pure PMMA UPy 2 0

0,5

1

1,5

Strain (%) 30

b

Pure PMMA Bare NPs UPy 2

M = 193 K w

Stress (MPa)

25 20 15 10 5 0

0

0,5

1

1,5

Strain (%) 30

c

M = 387.9 K w

25

Stress (MPa)

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20 15 10 5 0

Pure PMMA UPy 2 0

0,5

1

1,5

Strain (%)

Figure 3.


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a b

c

30

UPy 2

387.9 K

387.9K

UPy 2 387.9 K

25

+114%

193K

193 K

Stress (MPa)

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Nano Letters

98.4K 98K

20

UPy 2 193 K

+64%

15

UPy 2 98.4 K

10 5 0

≈

Pure matrix 0

0,2

0,4

0,6

0,8

1

1,2

1,4

0

Strain (%)

1

2

3

4

5

6

Young Modulus (GPa)

98.4 K

Figure 4.


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Nano Letters

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UPy 2

UPy 2

+64%

UPy 1 UPy 1

≈

Bare NPs

≈

Pure Matrix (193 K)

Bare 0

1

2

3

4

5

Young Modulus (GPa)

Figure 5.


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Scheme

PMMA

Supramolecular crosslink

---- -----

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Scheme 1.


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Table 1 dh,i (nm)

dh,n (nm)

Graft

Organic

thickness e

content

(nm) (c)

grafted to the NPs (wt%) (d)

20.4

-

20

-

25(a)

-

24(b)

-

UPy 1

44

28

2

4.8

UPy 2

58

48

12

64.1

Nissan

Bare NPs

Chemicals AS-40

(a) Obtained at one angle using the Zetazier Nano ZS instrument (b) From TEM analysis (c) Obtained by difference using the dh,n (d) Obtained by TGA


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Nano Letters

Table 2 Mw (K)

NPs

YM (GPa)

% change in

σ yield (MPa) /

% change in

YM

γyield (%)

σ yield

98.4

-

2.2

-

18 / 0.9

-

98.4

UPy 2

2.0

≈

18 / 1

≈

193

-

2.2

≈

7.8 / 0.65

-

193

Bare

1.9

≈

11.6 / 0.65

+49%

193

UPy 2

3.6

+64%

22.5 / 0.75

+188%

387.9

-

2.2

-

12.9 / 0.65

-

387.9

UPy 2

4.7

+114%

25 / 0.5

+94%


25

Nano Letters

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ASSOCIATED CONTENT Supporting Information. Materials and Methods, details of the grafted NPs synthesis, DLS and SLS details, bulge test description, thickness values from ellipsometry, annealing effect. This material is available free of charge via the Internet at http://pubs.acs.org.

AUTHOR INFORMATION Corresponding Author * (S.K.K) Email: [email protected] * (Z.G) Email: [email protected] * (N.J) Email: [email protected] Notes The authors declare no competing financial interest.

ACKNOWLEDGMENT The authors thank Dr Alamgir Karim from University of Akron (College of Polymer Science and Polymer Engineering) for providing the PSS base layer. Z.G. acknowledges the financial support from the US Department of Energy, Division of Materials Sciences (DE-FG02-04ER46162).


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Table of Contents Graphic.

PMMA

Supramolecular crosslink

114% Young Modulus

---------

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Enhanced Glassy State Mechanical Properties of Polymer

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