Silica–Epoxy Vitrimer Nanocomposites - Macromolecules (ACS

Article pubs.acs.org/Macromolecules

Silica−Epoxy Vitrimer Nanocomposites Aurélie Legrand and Corinne Soulié-Ziakovic* Laboratoire Matière Molle et Chimie, CNRS, ESPCI Paris, PSL Research University, 10 rue Vauquelin, Paris, France S Supporting Information *

ABSTRACT: Reinforced silica−epoxy vitrimer nanocomposites have been made in a solvent-free, easily processable, and economical way with filler contents up to 40 wt %. Increasing silica content leads to higher modulus materials in both glassy and rubbery regions. These nanocomposites are insoluble like permanent cross-linked networks but can completely relax stresses by thermoactivated exchange reactions that rearrange the network topology. Furthermore, the surface functionalization of particles improves the dispersion of fillers and accelerates the relaxation process. Similar results were obtained with industrial precipitated silica, which would allow vitrimer nanocomposites to be produced on an industrial scale.

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INTRODUCTION Vitrimers, presented by Leibler and co-workers in 2011, are organic permanent networks which rearrange their topology via thermally triggered exchange reactions without changing the cross-link density.1 A few vitrimer materials have been reported in the literature and are gathered in a recent review of Denissen et al.2 Currently, vitrimers utilize exchange reactions based on catalyzed transesterifications,1,3−5 transamination of vinylogous urethanes,6 transcarbamoylation of urethanes,7 olefin metathesis,8,9 and transalkylation of triazolium salts.10 Vitrimers are amazing covalent adaptable networks;11,12 the exchange mechanism is associative so no depolymerization occurs upon heating and the cross-link density is preserved at all temperatures. Therefore, these materials can swell in chemically inert solvents and like thermosets remain insoluble even at high temperatures. Two transition temperatures are characteristic of the viscoelastic behavior of vitrimers: the glass-transition temperature Tg between the glassy and the rubbery state of polymer networks and a second transition temperature Tv related to the exchange reactions which marks the transition between the viscoelastic solid state to the viscoelastic liquid state of vitrimer materials.3 At temperatures below Tg and Tv, the topology of the network is quenched and the material behaves like a permanently cross-linked network. Above Tg and Tv, the viscosity decreases following an Arrhenius law, behavior which is like that of silica. The exchange reactions become fast enough to allow the network to flow, thus conferring to these materials the ability to be reshaped, welded, and healed. To improve the mechanical and thermal properties of vitrimers, we explored the addition of fillers. Commonly, mechanical and thermal properties of cross-linked elastic materials are improved by adding large amount of inorganic fillers such as silica nanoparticles.13,14 Not only is the cost of final materials decreased but, depending on the kind of fillers used, the modulus, hardness, strength, thermal and rheological © XXXX American Chemical Society

properties, and fracture toughness of the nanocomposites can be tuned.15−17 However, adapting these strategies to vitrimers poses a challenge because adsorption of polymer chains at the surface of the fillers could hinder topology rearrangements and reduce the malleability. Fiber-reinforced thermoset composites made of a dynamic epoxy resin have been presented.18,19 These materials can be repaired and reshaped, but the extent of the deformation is limited due to the fibers. CNT epoxy vitrimer composites have been recently achieved.20 Yet, the filler content was too low to alter significantly the mechanical and/or vitrimer properties of the matrix. Our study focuses on the impact of adding large amount of particles on vitrimer properties. Especially, we studied to what extent fillers could impact the topology rearrangements depending on the surface chemistry of the particles. The mechanical properties of composites depend on several factors including type,21−25 size,16,26−30 aspect ratio,31,32 filler content, quality of the particles dispersion within the polymer matrix, and the chemical nature of the interface between fillers and polymer chains.33−36 Indeed, nanocomposites properties depend on the interactions between the fillers and the matrix referred as the interfacial adhesion.33 The strength strongly depends on the state of dispersion of the particles and even more importantly on the effectiveness of stress transfer between the matrix and the fillers. Large aggregates can create local defects that weaken materials,37 and generally, for equivalent loadings, smaller particles create stronger materials.16,28,29,33 Poor interfacial adhesion between nanofillers and the matrix leads to a dewetting along the phase boundary that causes extra stress concentration.38,39 Thus, limiting the aggregation of the fillers and improving their interfacial adhesion with the matrix Received: April 20, 2016 Revised: July 27, 2016

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Macromolecules Scheme 1. Adapting Surface Chemistry of Fillers to the Epoxy−Vitrimer Networka

a Epoxide-functionalized silica surface is linked to the matrix through β-hydroxy ester functions. Exchange reactions can happen both at fillers−matrix interface and in the network.

increases the strength of the composites.33,40−42 A common way to improve both the dispersion state and the interfacial adhesion is functionalizing the surface of the fillers.14,43−47 Steric stabilization induced by grafting polymer chains on the particles limits fillers aggregation. Fillers can also be covalently attached to the matrix using linkers or coupling agents leading to both good dispersion state and high interfacial adhesion values, but such linkages could prevent topology rearrangements in vitrimer composites. Adapting the functionalization of the fillers surface to the vitrimer matrix chemistry is key to achieving malleable vitrimer composites. We studied composites made from an epoxy-based vitrimer. Epoxy-based vitrimers are permanent polyester/polyol networks that contain a transesterification catalyst; the topology rearrangements result in this particular case from thermoactivated and catalyzed transesterification reactions. Depending on the chosen chemistry, epoxy-based vitrimers behave like classical soft or hard thermosets.1,3 In this article, we demonstrate that the elastic modulus of soft epoxy-based vitrimers can be reinforced with monodisperse silica nanoparticles up to 40 wt % without preventing the topology rearrangements. We show that dispersion states and rearrangement rates are improved when silica nanoparticles are linked to the vitrimer matrix through exchangeable bonds that participate in the relaxation of the composite (Scheme 1). Additionally, we achieved reshapable and reinforced epoxy−vitrimer composites with industrial precipitated silica in a solvent-free process, opening the way to large scale production of vitrimer nanocomposites.

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was kindly provided by Solvay. Water was obtained from a Millipore Milli-Q water purification system. Analytical Methods. Dynamic Light Scattering (DLS). DLS measurements were performed on monodispersed silica NPs to determine particles size with an ALV/CGS-3 compact goniometer system equipped with a 22 mW He−Ne laser. Dispersions were diluted with deionized water to a solid volumic fraction of 10−5 before analysis. The detection angle was varied between 30° and 150° with a 10° step. Transmission Electronic Microscopy (TEM). TEM images were obtained with a Zeiss CEM 902 microscope with an accelerating voltage of 90 kV and at a magnification of 3000×. 60−80 nm thick slices of composite were cut by ultracryotomy with a Leica Ultracut UCT equipped with a diamond knife (Diatome Cryo 35°, 2.0 mm) at −100 °C and a cutting speed of 1 mm/s. N2 Physisorption Analyses. These analyses were performed on a Micromeritics ASAP 2020. The samples were outgassed at 160 °C overnight before analysis. Specific surface areas were determined using the Brunauer−Emmet−Teller (BET) method. Thermogravimetric Analyses (TGA). TGA analyses were conducted on a Netzsch TG 209 F1 Libra. Samples were submitted to a 60 min isotherm at 120 °C before increasing temperature to 590 °C at a speed of 5 °C/min under a 20 mL/min O2 flow. Grafting density was determined evaluating the mass change between 130 and 600 °C of functionalized silica in comparison with bare silica (see TGA experiments in the Supporting Information for details). Tensile Tests. These tests were performed at 25 ± 1 and 100 °C on dog-bone samples (10 mm × 2 mm × 1.5 mm) with a strain rate of 10 mm/min using an Instron 5564 tensile machine. A video extensometer was used to determine the deformation. Each sample was kept at least 24 h at 25 °C in a climatic chamber before experiment. Young’s modulus was determined by linear fitting of stress−strain curves between 1% and 3% strain. Dynamic Mechanical Analyses (DMA). DMA analyses were conducted on a TA Q800 apparatus in the film tension geometry. Heating ramps were applied at 3 °C/min from −25 to 200 °C. Rectangular samples of 5 mm × 1.5 mm cross section and about 10 mm length were tested at 1 Hz and 15 μm amplitude. Relaxation. These experiments were also made with the DMA apparatus in the rubbery region of nanocomposites (190 °C). A 3% constant deformation was applied during the test. The stability of samples at this temperature was checked by TGA experiments (Figure SI-3). Differential Scanning Calorimetry (DSC). DCS was performed on a TA Q1000 apparatus. Two heating cycles, from −25 to 150 °C, were recorded at 10 °C/min.

MATERIALS AND METHODS

Materials. All the chemical products are commercially available and were used as received without further purification. Pripol 1040, a mixture of C18 fatty acids derivatives, containing about 23 wt % dimers and 77 wt % trimers (296 g/mol COOH), was provided by Croda. Bisphenol A diglycidyl ether (DGEBA, DER 332), tetraethyl orthosilicate (TEOS, 98%), and absolute ethanol were purchased from Sigma-Aldrich. Zinc acetate dihydrate (ACS, 98.0−101.0%) and (3glycidyloxypropyl)trimethoxysilane (97%) were purchased from Alfa Aesar. Ammonium hydroxide solution (0.88 S.G. 35% NH3) was purchased from Fisher Chemicals. Precipitated silica Zeosil 1165 MP B

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Figure 1. Synthetic procedures. (a) Synthesis and functionalization of monodisperse silica NPs. (b) Composites composition and formulation. evaporation of the solvent at 55 °C under vacuum, the silica was dried by increasing the temperature to 90 °C under vacuum and let for 1 h to enable the condensation of the organosilane moieties at the silica surface. Grafting density was estimated by TGA at 0.4 mmol/g or 1.5 coupling agents/nm2 of silica surface (see Figure SI-2 for details). The specific surface area measured by BET before functionalization is 160 m2/g. GLYMO−Zeosil silica and BARE−Zeosil silica refer respectively to epoxide-functionalized and nonfunctionalized precipitated silica. Nanocomposite Formulation. All composites were prepared with the same epoxy vitrimer matrix. Solvent-free dispersion process was developed and optimized for each filler type, namely, monodispersed silica NPs and precipitated silica nanofillers. For both types of fillers, powders were dried at 130 °C under vacuum for 30 min to activate the particles surface before mixing with the organic matrix. Preparation of the Catalyst/Fatty Acid Mixture (Pripol 1040). In a 1000 mL round-bottom flask under vacuum, 14.8 g of zinc acetate dihydrate (10 mol % to the acid functions) was dissolved in 200 g of Pripol 1040, by heating gradually from 100 to 180 °C. The mixture was maintained at 180 °C under vacuum until no more gas evolved (acetic acid resulting from the ligand exchange reaction) and until the complete dissolution of the zinc catalyst particles (24 h). Formulation of Monodisperse Silica Nanoparticles Composites (25 and 40 wt %). Silica NPs (BARE- or GLYMO-silica NPs) and the diepoxide reactant DGEBA were ground in a mortar until a homogeneous mixture was obtained and then added in a PTFE beaker with an adequate amount of the catalyst/fatty acid mixture (equimolar ratio COOH/epoxide functions). The blend was manually stirred at 130 °C until complete miscibility and then quickly poured into a brass mold (10 cm × 10 cm × 0.15 cm) sandwiched with two antiadhesive silicone papers. The mold was placed in a heating press and left under 5 tons for 15 h at 130 °C. Vitrimer composites were prepared with 25 and 40 wt % of silica nanoparticles, corresponding to 15% and 25% v/v, respectively. Formulation of Industrial Precipitated Silica Composites (25 wt %). Precipitated silica (22.5 g, BARE- or GLYMO- Zeosil silica), DGEBA (24.8 g), and the catalyst/acid mixture (42.6 g) were mixed together and poured in a Haake Polydrive mixer equipped with Banbury rotors. The blend was mixed 1 h at 25 °C and 100 rpm. The viscous mixture was then spread on a silicone paper, heated 2 min at

Welding Experiments. These experiments were conducted on rectangular samples of material (30 mm × 5 mm × 1.5 mm). The two ribbons were surimposed on a 15 mm length and held together under pressure for welding times ranging from 5 to 30 min at 190 °C. A good contact between the two pieces was ensured by applying a ∼ 30% strain during the treatment (assembly compressed from 3 to 2.1 mm). The welding efficiency was then evaluated by carrying out lap shear tests on the assembly at room temperature with a cross-head speed of 5 mm/min and comparing the forces at break. Synthesis and Characterizations. Monodisperse Silica Particles. Silica nanoparticles (NPs) of 100 nm in diameter were prepared using Stöber et al.’s method.48 Briefly, 800 mL of absolute ethanol and 48 mL of ammonium hydroxide solution were added in a roundbottom flask and stirred for 15 min at room temperature. 24 mL of TEOS was then quickly poured, and the resulting solution was stirred overnight at room temperature. Silica NPs were collected by centrifugation (8500 rpm, 45 min), washed with absolute ethanol by four cycles of centrifugation−dispersion, air-dried over 3 h at 100 °C, and then ground in a mortar. The particles size determined by DLS was 120 nm with a polydispersity index of 4%. The size determined from TEM images analysis was about 100 nm. The specific surface area (BET) is 41 m2/g. Functionalization of Monodisperse Silica Particles. Functionalization of particles was performed on the silica nanoparticles washed and dried as described above and dispersed in dry toluene (1 g of particles in 4 mL of solvent). The required quantity of organosilane was calculated to have an excess of about 100 organosilane molecules per nm2 of silica surface. This corresponds to approximately 1 mL of (3glycidyloxypropyl)trimethoxysilane per gram of 100 nm silica nanoparticles synthesized as described above. The mixture was stirred at 90 °C overnight. Silica NPs were collected and washed as previously described for nonfunctionalized NPs. Grafting density was estimated by TGA at 0.13 mmol/g corresponding to 2 coupling agents/nm2 of silica surface (see Figure SI-1 for details). GLYMO−silica NPs and BARE−silica NPs refer respectively to epoxide-functionalized and nonfunctionalized silica particles. Functionalization of Commercial Precipitated Silica Nanofillers. Precipitated silica Zeosil 1165MP (80 g) was mixed with 320 mL of dry cyclohexane and 10 mL of (3-glycidyloxypropyl)trimethoxysilane. The dispersion was homogenized under stirring for 30 min allowing the impregnation of silica micropearls with the organosilane. After the C

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Macromolecules 130 °C under vacuum to remove air bubbles, and poured into a brass mold (10 cm × 10 cm × 0.15 cm) sandwiched with two antiadhesive silicone papers.. The cure step was conducted in a heating press for 15 h under 5 tons at 130 °C.

Nanocomposites were insoluble when swelled in trichlorobenzene at 160 °C for 16 h, which confirmed that they are cross-linked materials and that exchange reactions do not induce depolymerization of the network (Figure 2b and Figure SI-5). The gel fraction was 80% for both the matrix and the 25 wt % composites. Because of transesterification equilibrium and topology reorganization, a few small molecules or clusters are dissolved during the experiment. Flory−Stockmayer gelation theory49,50 predicts a conversion of 77% at the gel point of the epoxy network. At ambient temperature, no topology rearrangements are observed, and the gel fraction determined by swelling experiments in chloroform at 25 °C is 90% (swelling experiments in Supporting Information). As also observed in the literature for epoxy resins/silica particles composites,14 TEM analyses revealed strikingly mesoscopic structures of composites depending on the surface chemistry. BARE−silica NPs form micrometer-size aggregates within filler-free regions of matrix (25 wt % composite is presented in Figure 2c and 40 wt % composite in Figure SI-4b). In contrast, functionalization of surfaces by epoxide groups remarkably improves the state of dispersion (Figure 2d and Figure SI-4d). Although some micrometer-size aggregates are still present, the vast majority of nanoparticles are not aggregated. This suggests that for epoxy-functionalized particles covalent links between functional groups and the matrix chains are formed and steric stabilization during mixing stage combined with good interfacial adhesion occurs to create an efficient dispersion within the matrix. Dynamic mechanical analysis (DMA) was performed in the tension film geometry to determine the effect of adding silica NPs on the mechanical properties of the epoxy−vitrimer matrix (Figure 3 and Table 1). Tα was taken at the E″/E′ maximum, the temperature location of the main relaxation related to the glass transition of the materials. For all composites, Tα is slightly higher than the value measured for the matrix in the absence of nanoparticles (Table 1). Similar increases (less than 10 °C) have also been reported for other polymer−silica nanocomposites.17,51−54 Over the full investigated range of temperatures, E′ of all composites is higher than the matrix modulus Em ′ and depends on the silica content but not on the type of interactions between the silica NPs and the matrix (Figure 3). Other reports33,40 have found interfacial adhesion has little effect on the modulus of composites since this value is determined at small loads or displacements for which debonding cannot be observed. Epoxy−vitrimer-based nanocomposites are no exception: whatever the surface treatment of the fillers, materials behave the same way with identical moduli at all temperatures (Figure 3). The Young’s modulus of particle−polymer composites may be predicted using empirical or semiempirical equations.55,56 In the case of spherical particles, the modulus depends on the modulus of components, particle loading, and particle size. Guth’s equation is based on the assumption of rigid particles in a dilute regime corrected with a particle−particle interaction term56 (the power term in eq 1).

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RESULTS AND DISCUSSION Monodisperse Silica NPs−Epoxy Vitrimer Composites. The silica−epoxy vitrimer nanocomposites were

Figure 2. The 25 wt % silica NPs epoxy vitrimer nanocomposites. (a) Picture of materials. (b) Swelling of the GLYMO−silica composite in trichlorobenzene. TEM analysis of the fillers state of dispersion in (c) the BARE−silica composite and (d) the GLYMO−silica composite.

Table 1. Tα Determined by DMA and Young’s Modulus at Room Temperature Determined by Tensile Tests; Experimental (25 °C) and Predicted Young’s modulus Ratio Ec/Em for the Vitrimer Matrix, the BARE−Silica Composites, and the GLYMO−Silica Composites (25 and 40 wt % Filler Contents) Tα (°C)

E (MPa)

material

DMA

25 °C (±2 MPa)

epoxy−vitrimer matrix 25 wt % BARE− silica 25 wt % GLYMO− silica 40 wt % BARE− silica 40 wt % GLYMO− silica

30

12

34

Ec/Em tensile tests (25 °C)

Guth’s prediction

22

1.8

1.7

34

20

1.7

1.7

33

35

2.9

2.5

37

44

3.7

2.5

obtained by dispersing BARE− or GLYMO−silica monodisperse NPs with diglycidyl ether of bisphenol A (DGEBA), a mixture of tricarboxylic and dicarboxylic fatty acids, and zinc acetate as a catalyst (Figure 1b). High silica contents of 25 and 40 wt % were chosen to reach a significant level of reinforcement and thus discriminate the impact of the chemical nature of the silica surface on properties. The latter concentration is close to the percolation threshold. Silanols and silanolates at the surface of the BARE−silica NPs can interact with the matrix trough noncovalent bonds (H-bonds and electrostatic interactions, respectively). In contrast, surface epoxide functions of GLYMO−silica NPs react with the fatty acid mixture and covalently link silica NPs to the network (Scheme 1).

Ec = 1 + 2.5Vp + 14.1Vp2 Em

(1)

Vp is the particle volume fraction, and Ec and Em are the Young modulus of the composite and the matrix, respectively. The Ec/ Em ratio indicates the amount of rigidity imparted to the D

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Figure 3. E′ and E″/E′ of (a) 25 wt % BARE− and GLYMO−silica composites compared to vitrimer matrix and (b) 40 wt % BARE− and GLYMO−silica composites (dashed lines) compared to 25 wt % GLYMO−silica composite (solid line). Blue: vitrimer matrix; black: BARE−silica composites; red: GLYMO−silica composites. Measurements made by DMA in tension film geometry, 1 Hz, 15 μm.

material by the fillers. Predicted and experimental Ec/Em ratios of BARE−silica and GLYMO−silica are of the same order of magnitude at 25% of fillers. At 40%, Ec/Em is underestimated probably due to aggregation (Table 1). The effect of the nature of silica/matrix interactions (expected to be mostly non covalent for BARE−silica NPs and covalent for GLYMO−silica NPs) on the strength and elastic properties of the nanocomposites can be highlighted with tensile test experiments. Stress−strain curves of experiments of the silica−epoxy vitrimer nanocomposites are presented in Figure 4. At 25 °C, experiments were conducted very close to the alpha transition of materials (Tα is 30 °C for the matrix and around 35 °C for composites, Table 1). The stress−strain behavior of the matrix is elastomeric at this temperature, whereas composites behave like solids in the glassy state (Figure 4a). The modulus, the tensile strength, and the elongation at break are deeply modified during the alpha transition. As this transition region is different for all materials, we have to be careful with the interpretation of these results: ultimate properties can only be compared if E′ versus T curves are superimposed. Thus, 25 wt % composites (BARE and GLYMO) can only be compared to each other and not to the matrix (the same for the 40 wt % composites). The strength value of the 25 wt % composites is 11 MPa with BARE−silica and 15 MPa with GLYMO−silica with a variation of 5% and 7%, respectively (see Figure SI-8). The elongation at break is around 100% for both composites (error on the value under 5%). The 40 wt % filler content leads to a diminished ultimate strain while the strength remains approximately the same (see Figure SI-9). However, due to a higher amount of aggregates in these composites, the error on the value of ultimate properties (strength and elongation at break) is

around 25%. The decrease of ultimate properties with increasing filler content has been observed for composites with particles larger than approximately 80 nm.33 Yet, the strength of GLYMO−silica composites is always around 40% higher than the strength of the BARE−silica composites (Figure 4a), which indicates a stronger interfacial adhesion with the vitrimer matrix as a result of covalent bonding. At 100 °C all materials, raw matrix and composites, are in the rubbery plateau so well above the glass transition. The 25 wt % composites show higher strength and modulus than the matrix (Figure 4b). At this temperature, the GLYMO−silica composite exhibits comparable mechanical properties than the BARE−silica composite (see Figure SI-8 for reproducibility details). As previously mentioned, the defining property of vitrimer materials is their ability to rearrange their topology by catalyzed exchange reactions at high temperature. Relaxation experiments were conducted on the nanocomposites at 190 °C, and a 3% constant deformation was applied (Figure 5). The corresponding relaxation times were determined by exponential fitting of the relaxation curves (Table 2). Several observations can be made from the data. First, higher silica content corresponds with longer relaxation time, which indicates that fillers slow down the network relaxation (Figure 5a). Yet, even with high filler contents, the nanocomposites are able to completely relax stresses at high temperature (190 °C) and to flow just as the epoxy−vitrimer matrix. Second, GLYMO−silica composites relax faster than the BARE−silica composites. In the latter, the matrix is locally adsorbed through noncovalent bonds at the particles surface so local reorganization can only be achieved through adsorption/desorption rearrangements. These interfacial movements, combined with exchange reactions within the E

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Figure 4. Stress−strain curves obtained at 25 °C (a) and 100 °C (b) for the epoxy vitrimer matrix (blue) and the BARE−silica (black) and GLYMO−silica (red) composites (25 wt %: solid line; 40 wt %: dashed line). Measurements made at a strain rate of 10 mm/min.

Figure 5. Relaxation curves obtained at 190 °C and under a 3% constant deformation: (a) materials with 10 mol % of Zn and (b) materials with 0.1 mol % of Zn. Blue: vitrimer matrix; black: BARE− silica composite; red: GLYMO−silica composite. Solid lines: filler content is 25 wt %. Dashed lines: filler content is 40 wt %.

matrix, allow the BARE−silica composites to entirely relax stresses, but more slowly than the GLYMO−silica composites. Indeed, for these composites, silica particles are linked to the vitrimer matrix through exchangeable β-hydroxy ester bonds that can also participate in the relaxation process of the whole material (Scheme 1). To evidence the adsorption/desorption process at the particles surface, similar experiments were performed on composites containing a reduced amount of catalyst (0.1% of zinc). In this case, the epoxy−acid network is formed but is not able to reorganize its topology by transesterifications. The small relaxation observed is attributed to the rearrangements of flexible parts of the network (Figure 5b). The relaxation extent is higher for the silica composites than for the raw vitrimer matrix (respectively 15% and 8% stress relaxation for the 25 wt % BARE and GLYMO composites compared to 5% for the raw vitrimer matrix, after 10 min). These results demonstrate that the relaxation of silica−epoxy vitrimer composites depends on at least two phenomena: exchange reactions (transterifications) and adsorption/desorption rearrangements at the particles/ matrix interface. Interestingly, these surface rearrangements are also observed for the GLYMO−silica composite. Zhuravlev showed that the surface silanol density, when the surface is

hydroxylated to the maximum degree, is a physicochemical constant which varies as a function of heat treatment. The total maximum silanol density is 4.6−4.9 OH per nm2.57 The grafting density at the GLYMO−silica NPs surface was estimated at 2 molecules per nm2. It is unlikely that all surface epoxide groups react with a fatty acid to form a β-hydroxy ester link due to limited accessibility. Consequently, some epoxide functions could have been converted to diols forming non covalent interactions with the matrix. Toward Industrially Relevant Nanofillers. We tried next to make vitrimer composites containing fillers commonly used in industry, such as precipitated silica nanofillers. 100 g blends of composites containing 25 wt % of bare or functionalized Zeosil 1165MP nanofillers were made solvent-free in a mixer. However, 40 wt % composites blends could not be properly processed due to their high viscosity as the silica content being beyond the percolation threshold. The functionalization of the fillers with organosilanes was performed prior their introduction in the catalyzed epoxy/acid mixture. TEM analysis revealed that the dispersion of GLYMO− Zeosil is only slightly improved compared to BARE−Zeosil. F

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Figure 6. TEM analysis of the state of dispersion of (a) the BARE− Zeosil and (b) the GLYMO−Zeosil 25 wt % composites (scale bars are 1 μm).

The aggregates are polydisperse in both composites (approximately 100 nm to 1 μm in diameter) (Figure 6). The viscosity of the mixtures before curing is much higher than when using monodisperse 100 nm silica particles. The specific surface area of precipitated silica fillers is 4 times higher than that of the silica NPs. The quantity of interfaces at similar filler contents is consequently higher. Dynamic mechanical analysis and tensile test results (Figure 7) were similar to the ones obtained with the 25 wt % monodisperse silica composites (Figures 3 and 4). First, storage moduli E′ are similar over the entire range of temperatures regardless of the surface chemistry of the particles (Figure 7a). The modulus of composites containing precipitated silica is roughly twice the modulus of composites with monodisperse silica NPs (25 °C, 25 wt % silica content, DMA additional results in Supporting Information). Second, we also observe that the strength of composites is improved by using epoxide functionalized silica. The reached strength is twice higher with GLYMO−Zeosil silica than with BARE−Zeosil silica at 25 °C (Figure 7b) and 100 °C (Figure SI-10). Finally, relaxation experiments show that both composites relax almost completely a 3% deformation (Figure 8). Composites relax stresses 3 times faster with functionalized precipitated silica than with bare silica. This result confirms the necessity to adapt the fillers surface chemistry to the exchangeable nature of the matrix in order to preserve the vitrimer properties, especially when dealing with fillers with high specific surface area. The ability to be reshaped has been tested with GLYMO-functionalized epoxy−vitrimer composites

Figure 7. (a) Modulus E′ of BARE−Zeosil (black) and GLYMO− Zeosil (red) composites compared to vitrimer matrix (blue) measured by DMA in tension film geometry (1 Hz, 15 μm). (b) Stress−strain curves obtained for these composites at 25 °C and a deformation rate of 10 mm/min.

Figure 8. Relaxation curves and relaxation times of BARE−Zeosil (black) and GLYMO−Zeosil (red) vitrimer composites compared to vitrimer matrix (blue).

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Macromolecules Table 2. Experimental Tensile Strength, Elongation at Break and Relaxation Time for the Epoxy−Vitrimer Matrix and Silica Composites Obtained from Tensile Tests (25 °C, 10 mm/min) and DMA Relaxation Experiments (190 °C, 3% Deformation) material vitrimer matrix 25 wt % BARE− silica 25 wt % GLYMO−silica 40 wt % BARE− silica 40 wt % GLYMO−silica

tensile strength (MPa)

elongation at break (%)

relaxation time τ (min)

24 (±2) 11 (±1)

220 100

1.6 4.5

14 (±1)

90

3.1

11 (±3)

55

7.4

14 (±3)

50

5.4

Figure 10. Stress−strain curves of lap-shear experiments performed on samples welded at 190 °C for various welding times. Blue: vitrimer matrix; black: BARE−Zeosil composite; red: GLYMO−Zeosil composite. Solid line: peeling. Dashed line: bulk rupture (sample does not peel and breaks in the bulk).

(Figure 9). A ribbon of GLYMO−silica or GLYMO−Zeosil composite can be durably twisted by local heating at 200 °C for a few minutes with no residual internal stresses. Welding experiments have also been performed on Zeosil composites. As vitrimer materials, two pieces of composites compressed together under heat should show a cohesive adhesion at the interface. The obtained welded assembly should behave like a new bulk material with the same mechanical properties as the two previous ones. Welding times were chosen according to the characteristic times determined from the relaxation experiments. As can be seen in Figure 10a, the adhesion increases with increasing the welding time for all materials. This means that exchanges take place at the interface and allow the establishment of chemical links between the two surfaces. After 30 min at 190 °C, no more peeling was observed for composites, and the rupture occurred directly in the bulk material, above the welded interface. A picture of such a broken assembly, welded for 30

Figure 9. GLYMO−silica composites can be reshaped in a few minutes by local heating (200 °C).

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min, is shown in Figure 10b. These experiments also evidence that the Zeosil composites are vitrimer materials and are (re)processable. However, the composites, and specially the BARE one, cannot handle large deformation and break quickly at 25 °C. It was then difficult with these experiments to show the impact of the epoxide functionnalization of the fillers on the welding efficiency of materials.

CONCLUSION Vitrimer silica nanocomposites, with filler contents up to 40 wt %, can easily be processed in relatively large scales without any solvent. Increased loadings lead to higher modulus materials in both glassy and rubbery regions. When functionalized with epoxide functions that can react during the vitrimer preparation, silica fillers are covalently linked to the network through the same exchangeable β-hydroxy ester bonds formed within the vitrimer matrix. Thanks to these bonds, their interfacial adhesion with the vitrimer network and their dispersion state are improved. Furthermore, nanocomposites can completely relax stresses and flow at high temperature. Even though nanoparticles slow down the stress relaxation of the vitrimer network, surface exchangeable bonds speed up the relaxation of composites compared to nonfunctionalized fillers. In conclusion, by adapting the chemistry of the filler surface to the vitrimer one, it was possible to reinforce the network by enhancing interfacial adhesion and dispersion state, without drastically altering the stress relaxation of the vitrimer matrix. The process can be performed in a mixer to produce large batches of composites containing industrial precipitated silica (up to 25 wt %) that can be easily reshaped. This work can be readily extended to other vitrimer chemistries by adapting the functionalization of fillers to the chemical nature of the exchange reactions. ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.6b00826.

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REFERENCES

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Article

TGA experiments, swelling experiments, additional DMA results and DSC experiments (PDF)

AUTHOR INFORMATION

Corresponding Author

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

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We deeply thank Professor Ludwik Leibler for the numerous discussions which inspired this work and brought decisive enlightenment in the understanding of these systems. We are indebted to Szilvia Karpati for her help with TEM experiments. We thank François Tournilhac for useful discussions and Tyler Stukenbroeker for critical reading of the manuscript. We are grateful to SOLVAY for providing Zeosil 1165MP and CRODA for providing Pripol 1040. The authors acknowledge funding from ESPCI and CNRS. I

DOI: 10.1021/acs.macromol.6b00826 Macromolecules XXXX, XXX, XXX−XXX

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DOI: 10.1021/acs.macromol.6b00826 Macromolecules XXXX, XXX, XXX−XXX