Reprocessable Polyhydroxyurethane Network Composites: Effect of

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Cite This: ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

Reprocessable Polyhydroxyurethane Network Composites: Effect of Filler Surface Functionality on Cross-link Density Recovery and Stress Relaxation Xi Chen,† Lingqiao Li,† Tong Wei,† David C. Venerus,§ and John M. Torkelson*,†,‡ Department of Chemical and Biological Engineering and ‡Department of Materials Science and Engineering, Northwestern University, Evanston, Illinois 60208, United States § Department of Chemical and Biological Engineering, Illinois Institute of Technology, Chicago, Illinois 60616, United States Downloaded via NEW MEXICO STATE UNIV on January 5, 2019 at 08:36:53 (UTC). See https://pubs.acs.org/sharingguidelines for options on how to legitimately share published articles.

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ABSTRACT: Conventional polymer network composites cannot be recycled for high-value applications because of the presence of permanent covalent cross-links. We have developed reprocessable polyhydroxyurethane network nanocomposites using silica nanoparticles with different surface functionalities as reinforcing fillers. The property recovery after reprocessing is a function of the interaction between the filler surface and the network matrix during the network rearrangement process. When nonreactive silica nanoparticles lacking significant levels of surface functional groups are used at 4 wt % (2 vol %) loading, the resulting network composite exhibits substantial enhancement in mechanical properties relative to the neat network and based on values of rubbery plateau modulus is able to fully recover its cross-link density after a reprocessing step. When nanoparticles have surface functional groups that can participate in dynamic chemistries with the reprocessable network matrix, reprocessing leads to losses in mechanical properties associated with cross-link density at potential use temperatures, along with faster rates and lower apparent activation energies of stress relaxation at elevated temperature. This work reveals the importance of appropriate filler selection when polymer network composites are designed with dynamic covalent bonds to achieve both mechanical reinforcement and excellent reprocessability, which are needed for the development of recyclable polymer network composites for advanced applications. KEYWORDS: covalent adaptable network composites, polymer−filler interfaces, reprocessability, stress relaxation, surface functionality

1. INTRODUCTION Thermosets are used in a wide range of applications because of outstanding mechanical properties, heat stability, and solvent resistance. However, conventional thermosets cannot be reshaped or reprocessed in the melt state and thus cannot be recycled for high-value applications because of the existence of permanent, covalent cross-links. Over the past 15 years or so, research efforts have focused on incorporating dynamic covalent bonds into polymer networks, with the goal of making cross-linked polymers malleable or adaptable under certain conditions.1−7 The dynamic chemistries used in covalent adaptable networks8 (CANs) are divided into two types based on their mechanism and impacts on the cross-link density of the network. The first type is dissociative dynamic chemistry in which covalent, reversible bonds break under stimulus, hence leading to a reversion of cross-linked polymer to branched and linear chains that can be reprocessed and reshaped in the melt state, with cross-link bonds reforming upon stimulus removal. Examples of dissociative chemistry used in CANs include the Diels−Alder reactions1,9 and alkoxyamine chemistry.10,11 The second type is associative © XXXX American Chemical Society

dynamic chemistry in which dynamic bonds exchange with each other, maintaining the number of bonds constant during network rearrangement but leaving the network polymer, sometimes called a vitrimer, malleable at sufficient exchange rate. Examples of associative chemistry used in CANs include transesterification,2,12 transamination,13 and transcarbamoylation.14−16 Dynamic covalent bonds can endow networks with advanced characteristics including self-healing,17−20 reprocessability,21−28 degradability,29−31 and plasticity.32,33 Polymer properties and performance can be improved by incorporating reinforcing fillers to create polymer composites.34−38 This approach has been applied to CANs to enhance their mechanical properties. Many studies of covalent adaptable network composites (CANCs) have focused on selfhealing applications.39−44 By using reversible Diels−Alder chemistry, Li et al.41 designed a graphene oxide/polyurethane (PU) network composite with mechanical property enhanceReceived: October 31, 2018 Accepted: December 26, 2018 Published: December 26, 2018 A

DOI: 10.1021/acsami.8b19100 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

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Figure 1. Mechanism of synthesis and rearrangement of the PHU network composites.

ing.43,54,56,57 Compared to self-healing and plasticity, a higher extent of malleability is required to achieve reprocessability.15 Thus, there is a tradeoff between mechanical reinforcement and reprocessability. A few studies of reprocessable CANCs have shown excellent recovery of mechanical properties after reprocessing,52−54 but no study has focused on recovery of cross-link density, which is one of the most important characteristics of cross-linked polymers. For example, Guo and co-workers introduced exchangeable β-hydroxy ester bonds at the interfaces between epoxidized natural rubber and multiwalled carbon nanotubes.54 Their network composite exhibited improved mechanical performance and “an integrated recovery of tensile strength, tensile modulus, and ultimate elongation”54 after remolding. Polyhydroxyurethane (PHU), which can be synthesized by reaction of cyclic carbonate and amine functional groups, has received recent research focus as an isocyanate-free alternative to PU.14,16,58−63 As such, PHUs have sustainability benefits because they eliminate the use of toxic isocyanates involved in traditional PU production. Recently, several studies have investigated PHU networks with dynamic crosslinks.14,16,28,63,64 Dichtel, Hillmyer, and co-workers indicated that PHU thermosets derived from bis(six-membered cyclic carbonates) and amines can achieve network rearrangement through transcarbamoylation between carbamate groups and pendant hydroxyl groups.14 Their network recovered 75% of original room-temperature tensile properties after reprocessing

ment and thermally triggered self-healing characteristics; they estimated a 78% average healing efficiency from recovery of stress-at-break.41 Zhang et al. developed a self-healable carbon fiber-reinforced polymer (CFRP) composite based on an isocyanurate−oxazolidone network matrix, in which the isocyanurate and epoxide groups transform to oxazolidone rings under thermal stimulus.44 These CFRP composites have high glass transition temperature (Tg) up to 285 °C and show no thermal degradation below 300 °C.44 Other studies have demonstrated an array of stimuli-responsive applications of CANCs including recycling of carbon fibers in CFRP composites by depolymerization of the thermoset matrix,45−47 surface welding of thermoset composites,48 and shape reconfiguration of shape memory network composites.49,50 Yu et al.46 reported that an epoxy-based CFRP composite fully dissolves in ethylene glycol via transesterification exchange reactions; the reclaimed carbon fibers can be reused with the same properties as virgin fibers, and the solution reforms the polymer matrix upon heating. Despite many efforts to incorporate dynamic chemistries in polymer network composites, few studies have demonstrated reprocessability and remoldability of the materials.51−55 The limited number of reprocessable CANC studies may be attributed to the fact that although filler incorporation can reinforce the mechanical properties of the network matrix, fillers have adverse effects on the stress relaxation behavior and topological rearrangement associated with reprocessB


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product. Finally, the product was dried in a vacuum oven at 80 °C for 24 h. 2.3. Synthesis of Neat PHU Networks. Neat PHU network was synthesized using trifunctional cyclic carbonate THPMTC and longchain diamine ELASTAMINE RP-405, with 10 mol % DMAP catalyst with respect to the total number of amine functional groups. In a typical synthesis, 1.837 g of vacuum-dried THPMTC (3.10 mmol), 113.65 mg of DMAP (0.93 mmol), and 1.60 mL of anhydrous DMF were added into a Flacktek Max-10 mixing cup and heated at 80 °C until complete dissolution. Then, the mixture was cooled to room temperature, and 2.00 g of predried ELASTAMINE RP-405 (4.65 mmol) was added dropwise. The cup was sealed by a piece of parafilm, and reactants were mixed for 60 s in a FlackTek DAC 150.1 FVZ-K SpeedMixer at 2800 rpm. The homogeneous reactant solution was then poured into a 20 mL scintillation vial, and dry argon gas was flowed into the vial for 30 s to replace the air. The vial was sealed and placed in an oil bath at 80 °C to react for 12 h. Dry argon gas was flowed through the vial for 4 h to evaporate some solvent before the gel was removed and torn into small pieces and then dried in a vacuum oven at 80 °C for 48 h. 2.4. Synthesis of PHU Network Composites. Three different types of silica nanoparticles with same sizes were used to formulate PHU network composites: standard silica nanoparticles (OH-silica) with hydroxyl functional groups on the surface, amine-terminated silica nanoparticles (NH-silica) with amine functional groups on the surface, and superhydrophobic silica nanoparticles (SHP-silica) with some proprietary modifications of surface hydroxyl groups. With OHsilica and NH-silica, network composites were synthesized at a 4 wt % loading of nanofillers (OH-cPHU and NH-cPHU), which corresponds to ∼2 vol %. With the SHP-silica, network composites were synthesized with 4 and 2 wt % nanofillers (SHP-cPHU-4 and SHPcPHU-2). In a typical synthesis of OH-cPHU, 153.49 mg of OH-silica nanoparticle was added to a Flacktek Max-10 mixing cup and placed in a vacuum oven at 80 °C for 1 h to remove moisture. Then, 1.00 mL of anhydrous DMF was added, and the cup was sonicated for 90 min using a Branson 3510-DTH Ultrasonic Cleaner to disperse the nanoparticles. After sonication, 1.837 g of vacuum-dried THPMTC (3.10 mmol), 113.65 mg of DMAP (0.93 mmol), and 0.60 mL of anhydrous DMF were added to the cup, and the mixture was heated at 80 °C until complete dissolution. Predried ELASTAMINE RP-405 (2.00 g, 4.65 mmol) was added into the cup and mixed with other components for 60 s in a FlackTek DAC 150.1 FVZ-K SpeedMixer at 2800 rpm. The homogeneous reactant solution was poured into a 20 mL scintillation vial, and dry argon gas was flowed through the vial for 30 s. The reactants were placed in an oil bath at 80 °C to react for 12 h, followed by another 4 h flowing of argon gas. Then, the gel was removed, torn into small pieces, and dried in a vacuum oven at 80 °C for 48 h to remove DMF. 2.5. Reprocessing Procedure for the Neat PHU Network and Network Composites. Reprocessing was performed using a PHI press (model 0230C-X1). The as-synthesized materials in the form of small pieces were pressed into ∼1 mm thick sheets at 140 °C with a 7 ton ram force (generating a pressure of ∼11 MPa). The same 2 h reprocessing time was used for the neat PHU network and all network composites. After reprocessing the original materials, a uniform film was obtained, which was considered as the 1st molded sample. The 1st molded sample was then cut into small pieces and pressed again to give the 2nd molded sample. Similarly, the same procedure was performed to obtain the 3rd molded sample. 2.6. Characterization. Dynamic mechanical analysis (DMA) was performed with a TA Instruments Rheometrics Stress Analyzer III, characterizing storage modulus (E′), loss modulus (E″), and tan δ (E″/E′). Rectangular specimens (25 mm (L) × 4 mm (W) × 1 mm (T)) were mounted on the fixture and underwent temperature ramps from −30 to 110 °C. The measurement was carried out in tension mode at a frequency of 1 Hz, an oscillatory strain of 0.03%, and a 3 °C/min heating rate. At least three measurements were performed for each sample, and the E′ value at 100 °C was reported as the average rubbery plateau modulus with errors given by standard deviations.

at 160 °C for 8 h; the incomplete recovery was attributed to decomposition at high reprocessing temperature over long times.14 In a previous study, we showed that PHU networks synthesized with appropriate catalysis from five-membered cyclic carbonates and amines can be reprocessed multiples times at 140 °C for 1 h with full property recovery associated with cross-link density.16 We also discovered that both dissociative reversible cyclic carbonate aminolysis and associative transcarbamoylation exchange reactions participate in the catalyzed PHU network rearrangement.16 Here, we have developed the first reprocessable PHU network composites, focusing our study on the effect of silica nanoparticle surface functionality. As shown in Figure 1, the network formation and rearrangement process are similar to those determined in our previous study.16 Depending on surface functionality, during network rearrangement the nanofillers may participate in the reversible reaction or the exchange reaction, as with hydroxyl or amine groups, or in neither reaction in the absence of significant levels of surface functionality. With hydroxyl or amine group surface functionality, PHU network composites with 4 wt % nanofiller are reprocessable but with small ∼10−20% reductions in rubbery plateau modulus and thus cross-link density accompanying each reprocessing step, outside the range of experimental error. In contrast, when levels of nanofiller surface reactive groups are significantly reduced, within error the network composites with 2 and 4 wt % silica exhibit full recovery of rubbery plateau modulus and hence cross-link density after one reprocessing step. Additionally, Young’s modulus and ultimate tensile strength at potential use conditions are enhanced by ∼40% upon incorporating within the network composite 4 wt % silica nanofiller lacking significant surface functional groups. We also characterized the influence of filler surface functional groups on high-temperature stress relaxation of CANCs, an effect previously studied to only a limited extent.56,57

2. EXPERIMENTAL SECTION 2.1. Materials. Tris(4-hydroxyphenyl)methane triglycidyl ether, 4(dimethylamino)pyridine (DMAP, ReagentPlus, ≥99%), tetrabutylammonium iodide (TBAI, reagent grade, 98%), N,N-dimethylacetamide (ReagentPlus, 99%), propylene carbonate (ReagentPlus, 99%), and octylamine (99%) were purchased from Sigma-Aldrich and used as received. Anhydrous N,N-dimethylformamide (DMF, 99.8%) was purchased from Sigma-Aldrich and was kept on activated molecular sieves before use. Standard, superhydrophobic, and amine-terminated silica nanoparticles (size ∼80 nm) were purchased from NanoCym and used without further purification. ELASTAMINE RP-405 polyoxypropylenediamine (average Mn = ∼430 g/mol as reported by the supplier) was from The Huntsman Corporation and was kept on activated molecular sieve before use. 2.2. Synthesis of Tris(4-hydroxyphenyl)methane Tricyclocarbonate (THPMTC). THPMTC cross-linkers were synthesized, as described in ref 16. Tris(4-hydroxyphenyl)methane triglycidyl ether (20.00 g, 43.43 mmol) and TBAI (1.604 g, 4.343 mmol) were added to a test tube. The tube was placed into an oil bath set at 80 °C, and CO2 gas was bubbled through the mixture continuously until the reaction reached full conversion. N,N-Dimethylacetamide was added to the test tube when the mixture became too viscous for CO2 to flow through. The reaction progress was monitored using 1H NMR spectroscopy by examining the disappearance of peaks (∼2.72, 2.87, and 3.32 ppm) associated with epoxy moiety. After the reaction was completed, the mixture was dissolved in 25 mL acetone and poured into a beaker. Then, 75 mL of deionized water was added to precipitate the product. The aqueous phase was decanted, and the solid product was dried on a hot plate. The acetone−water wash was repeated several times until TBAI was fully removed from the C


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cases, consolidated and homogeneous films were obtained after reprocessing, which maintained their intact forms and swelled in swelling tests. Therefore, connections between broken pieces were successfully reformed during the reprocessing (remolding) step as a result of network rearrangement enabled by transcarbamoylation and reversible cyclic carbonate aminolysis. We first synthesized a neat PHU network and verified its reprocessability since a different catalyst load and reprocessing condition was used compared with our previous study.16 We also note that the diamine was switched from JEFFAMINE D400 in ref 16 to ELASTAMINE RP-405 in this study; both diamines are polypropylene-oxide-based materials. Figure 2

Uniaxial tensile testing was performed according to ASTM D1708 using an MTS Sintech 20/G tensile tester at 40 °C. Dog-bone-shaped samples (25 mm (L) × 5 mm (W) × 1 mm (T) and a gauge length of 18 mm) were cut from molded films and were kept on the fixture in an oven at 40 °C for 5 min to reach thermal equilibrium. Then, the samples underwent a uniaxial extension at a rate of 130 mm/min until break, with the data collected at a frequency of 350 Hz. The Young’s modulus, ultimate tensile strength, and strain-at-break results were reported as average values of at least five specimens with errors representing the standard deviations. Stress relaxation experiments were performed using a TA Instruments Rheometrics Stress Analyzer III. Rectangular specimens (25 mm (L) × 4 mm (W) × 1 mm (T)) were mounted on the fixture and allowed to equilibrate at the desired temperature for 10 min before being subjected to an instantaneous strain of 5%. The stress relaxation modulus at 5% strain was recorded until it decayed to ∼30% of its original value. Stress relaxation experiments were carried out at 120−160 °C at 10 °C increments, with at least three samples at each temperature. Attenuated total reflectance-Fourier transform infrared (ATRFTIR) spectroscopy was performed using a Bruker Tensor 37 FTIR spectrophotometer equipped with a diamond/ZnSe ATR attachment. All samples were scanned at a resolution of 2 cm−1, and 64 scans were collected in the range of 4000−600 cm−1. The cyclic carbonate group peaks at ∼1800 cm−1 were examined with the spectra normalized with respect to the aliphatic ether stretching peak at ∼1100 cm−1. 1H NMR spectra were obtained using a Bruker Avance III 500 MHz NMR spectrometer with a direct cryoprobe at room temperature and deuterated chloroform as solvent. Differential scanning calorimetry (DSC) was performed using a Mettler Toledo DSC 822e. Samples were annealed at 80 °C for 15 min and then cooled at −40 °C/min to −40 °C. Sample Tg’s (1/2 ΔCp method) were obtained from the following heating ramp at 10 °C/min. Scanning electron microscopy (SEM) samples were prepared by cryo-fracturing the as-synthesized and reprocessed PHU network composites. Each sample was mounted onto carbon tape and coated with a 15 nm thick layer of osmium using an OPC osmium coater to prevent charging. The morphologies of cryo-fractured section of the PHU network composites were obtained using a Hitachi S-8030 scanning electron microscope. Swelling tests were carried out for all molded samples to verify the effective network reconnection. Samples measuring ∼10 mm × ∼10 mm × ∼1 mm were immersed in chloroform or DMF, which are good solvents for linear PHU. Samples were left to swell for 1 day and were then subjected to external turbulence to determine whether they would break into small pieces.

Figure 2. Dynamic mechanical responses of the neat PHU network: E′ and tan δ as functions of temperature for 1st molded (squares), 2nd molded (circles), and 3rd molded (triangles) samples. Inset: zoomed-in rubbery plateau moduli as a function of temperature.

shows the DMA results of the 1st molded, 2nd molded, and 3rd molded neat PHU network. In all three molded samples, a rubbery plateau in the storage modulus E′ is observed at temperatures above Tg, confirming their cross-linked nature. The E′ rubbery plateau moduli determined at 100 °C are 2.27, 2.23, and 2.18 MPa, as shown in Table 1. These values are identical within experimental error, indicating a full recovery within error of cross-link density after multiple reprocessing steps. (According to ideal rubber elasticity theory, the rubbery plateau tensile modulus, E, is proportional to cross-link density.65 In the rubbery plateau regime, the values of loss moduli, E″, for our network and network composites are negligibly small compared with E′; thus, the rubbery plateau E′ values are essentially the same as E. Flory’s ideal rubber elasticity theory also predicts that E is proportional to temperature in the rubbery plateau regime, which is verified by the increase in storage modulus at temperatures well above Tg, as shown in the inset plot in Figure 2.) In addition to rubbery plateau modulus, the samples also show excellent consistency in Tg, as determined by both DMA and DSC (Table 1). Figure S1 shows the FTIR spectra of the three molded, neat PHU networks, normalized with respect to the aliphatic ether stretching peak at ∼1100 cm−1. Very small changes in cyclic carbonate groups (∼1800 cm−1), carbamate carbonyl groups (∼1700 cm−1), and urea carbonyl groups (∼1650 cm−1) are observed, which could be caused by low levels of side reactions occurring at high temperature for prolonged reprocessing time. For example, the increase in urea carbonyl groups after each reprocessing step may originate from the reaction of primary amine groups (liberated from reversible cyclic carbonate

3. RESULTS AND DISCUSSION 3.1. Effect of Filler Surface Functionality on Reprocessability of PHU Network Composites. Neat PHU networks and PHU network composites were synthesized using a trifunctional cyclic carbonate cross-linker THPMTC, a long-chain diamine ELASTAMINE RP-405, and DMAP catalyst at 10 mol % with respect to amine functional groups, as shown in Figure 1. Compared to our previous study of neat PHU networks,16 the amount of catalyst was increased because the incorporation of silica nanofillers restricted the mobility of the network matrix and led to a slower rate of network rearrangement; a higher catalyst load was needed to maintain the reprocessing time scale reasonably short as side reactions could possibly occur in PHU networks derived from fivemembered cyclic carbonates and diamines63 at the reprocessing temperature and cause adverse effects. In this study, the neat PHU network and network composites were all reprocessed under the same condition of 140 °C for 2 h for fair comparison of reprocessability and other properties. In all D


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ACS Applied Materials & Interfaces Table 1. E′ Rubbery Plateau moduli and Glass Transition Temperatures of the Neat PHU Network and PHU Network Composites sample neat PHU

OH-cPHU

NH-cPHU

SHP-cPHU-4

SHP-cPHU-2

1st molded 2nd molded 3rd molded 1st molded 2nd molded 3rd molded 1st molded 2nd molded 3rd molded 1st molded 2nd molded 3rd molded 1st molded 2nd molded 3rd molded

E′ (MPa)a

Tg (°C) by DMAb

Tg (°C) by DSCc

2.27 ± 0.05

23

20

2.23 ± 0.10

24

20

2.18 ± 0.04

23

20

2.93 ± 0.08

22

19

2.64 ± 0.10

21

19

2.16 ± 0.04

21

19

2.48 ± 0.07

22

18

1.91 ± 0.08

21

18

1.55 ± 0.06

21

17

2.75 ± 0.07

23

19

2.66 ± 0.11

22

20

2.37 ± 0.09

22

19

2.50 ± 0.05

21

19

2.44 ± 0.04

22

18

2.25 ± 0.06

21

18

the same conditions as with the neat PHU network. Figure 3a shows the DMA curves of 1st, 2nd, and 3rd molded OH-cPHU samples, and the E′ plateau moduli at 100 °C are summarized in Table 1. The 1st molded OH-cPHU sample shows an enhanced rubbery plateau modulus compared with the neat PHU network (increased by ∼29%), consistent with effective reinforcement by the incorporation of silica nanofiller. However, small but experimentally significant losses in OHcPHU rubbery plateau modulus (by ∼10% from 1st to 2nd molded samples and ∼26% from 1st to 3rd molded samples) are observed after reprocessing, with the 3rd molded sample dropping to the same E′ value as that of neat PHU samples. The loss in rubbery plateau modulus indicates a decreasing cross-link density as the network composite is increasingly reprocessed, consistent with exchange reactions occurring between hydroxyl groups on the nanoparticle surfaces and the network matrix. Because of their surface hydroxyl groups, the nanoparticles can participate in transcarbamoylation exchange reactions during network rearrangement, with the network matrix effectively losing part of its cross-link density because some links are not “exchanged back” as well as forming loops at the filler surface. We synthesized NH-cPHU using 4 wt % amine-functionalized silica nanoparticles with primary amine surface functionality. The E′ and tan δ profiles of molded NH-cPHU samples are shown in Figure 3b with rubbery plateau E′ values given in Table 1. Notably, the 1st molded NH-cPHU sample exhibits a rubbery plateau modulus only slightly enhanced (by ∼9%) relative to neat PHU and significantly below that (by ∼15%) of the 1st molded OH-cPHU sample. Several reasons may account for the lower rubbery plateau modulus in 1st molded NH-cPHU compared with 1st molded OH-cPHU. First, amine functional groups could perturb the stoichiometric balance in the reactants during network formation, leading to networks with more imperfections. Second, the 1st molded NH-cPHU sample may have already suffered from a higher extent of loss in cross-link density during the initial molding process, which occurred over a 2 h period. Compared to OH-cPHU, greater losses in rubbery plateau modulus are observed after reprocessing (by ∼23% from 1st to 2nd molded samples and ∼38% from 1st to 3rd molded samples). As with OH-cPHU, these losses can be attributed to interaction between the network and silica nanoparticles. With NH-cPHU, the primary amine functional groups on the filler surface can participate in reversible cyclic carbonate aminolysis

Determined at 100 °C by DMA. bDetermined as the temperature where the loss modulus E″ reaches its maximum value. cDetermined as the Tg,midpoint using the 1/2 ΔCp method.

a

aminolysis) with carbamate groups, which produces urea groups.63 The fact that the neat PHU network exhibits reprocessability with full property recovery despite these small changes indicates that low levels of side reactions do not significantly influence the bulk properties of the PHU network. We synthesized OH-cPHU network composites using 4 wt % (corresponding to ∼2 vol %) standard silica nanoparticles, which contain abundant hydroxyl groups on the surface. The network composites were synthesized and reprocessed using

Figure 3. Dynamic mechanical responses of (a) OH-cPHU and (b) NH-cPHU: E′ and tan δ as functions of temperature for 1st molded (squares), 2nd molded (circles), and 3rd molded (triangles) samples. Inset: zoomed-in rubbery plateau moduli as a function of temperature. E


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Figure 4. Dynamic mechanical responses of (a) SHP-cPHU-4 and (b) SHP-cPHU-2: E′ and tan δ as functions of temperature for 1st molded (squares), 2nd molded (circles), and 3rd molded (triangles) samples. Inset: zoomed-in rubbery plateau moduli as a function of temperature.

region, no comparison between the hydroxyl content in OHsilica and NH-silica could be drawn. Relative to OH-silica, a much smaller silanol group peak was observed in SHP-silica, consistent with a much lower yet finite level of residual hydroxyl groups on the surface of SHP-silica. We speculate that if all surface hydroxyl groups on silica nanoparticles were eliminated, then SHP-cPHU samples may exhibit full recovery of rubbery plateau modulus and cross-link density after multiple reprocessing steps, as with neat PHU networks. Future study is warranted. To verify the interfacial reaction between fillers and the polymer matrix, we performed a model study using monofunctional small-molecule model compounds. Propylene carbonate was reacted with octylamine in stoichiometric balance at 80 °C for 18 h until nearly full conversion of cyclic carbonate functional groups (Figure S3). DMAP catalyst was added at 10 mol % with respect to the cyclic carbonate functional groups. Then, the model compound was reacted in DMF at 140 °C for 8 h with silica nanoparticles at 10 wt % of nanoparticles relative to the model compound. (We note that this loading is ∼2.5 times the nanoparticle level used in our PHU composite studies and helps to ensure that reaction can be observed via FTIR peaks. However, because the nanoparticle loading is higher and the reaction occurs over longer time frames with the model compound, we caution that the model compound studies should not be used to infer anything about the quantitative levels of reaction in the PHU network composites.) The silica nanoparticles were collected using a high-speed centrifuge and washed seven times with methanol to remove any residual model compound. Figures S4−S6 show the FTIR spectra of silica nanoparticles before and after reaction. For all three types of particles, a peak associated with carbonyl groups at ∼1650 cm−1 can be observed, indicating the grafting of the model compound onto the filler surfaces. The reaction between fillers and the model compound can be transcarbamoylation exchange reactions or reversible cyclic carbonate aminolysis depending on the filler surface functionality. We examined the FTIR spectra of all molded PHU network composite samples; see Figures S7−S10. For each composite, the very small changes to FTIR spectra with reprocessing are no greater than that for the neat PHU network, indicating that the neat PHU network and PHU network composites undergo similar, very small levels of side reactions with reprocessing. Thus, the different effectiveness of reprocessing observed for

during network rearrangement, decross-linking part of the network. Relative to OH-cPHU, the worse property recovery in NH-cPHU after reprocessing indicates that the network loses more cross-links when reversible reactions are viable than when exchange reactions are viable between the network and fillers. To eliminate the filler−matrix interaction during network rearrangement, we synthesized SHP-cPHU-4 using 4 wt % superhydrophobic silica nanoparticles. In the absence of surface hydroxyl groups (or other functionality), this type of nanoparticle is not expected to undergo any significant reaction with the network matrix. Relative to neat PHU network, this PHU network composite exhibits an enhanced rubbery plateau modulus (by ∼21% for 1st molded samples). More importantly, within error the 2nd molded SHP-cPHU-4 sample fully recovers its original rubbery plateau modulus and thus cross-link density, as shown in Figure 4a and Table 1. This excellent recovery of cross-link density makes clear the importance of suppressing or eliminating reactions between a reprocessable network and filler surfaces during reprocessing. Despite the recovery of cross-link density within error from 1st to 2nd molded samples, a small loss in E′ is observed when the network composite is molded a 3rd time (∼14% reduction relative to the 1st molded sample). We prepared another network composite, SHP-cPHU-2, to determine if reducing the filler content to 2 wt % would aid the property recovery of the 3rd molded sample. Similar to SHP-cPHU-4, the 2nd molded SHP-cPHU-2 exhibits the same rubbery plateau modulus within error as the 1st molded sample, whereas a ∼10% loss in E′ was observed for the 3rd molded sample compared with the 1st molded sample. See Figure 4b and Table 1. Regarding the small losses in cross-link density in 3rd molded SHP-cPHU-4 and SHP-cPHU-2 samples relative to 1st molded samples, we hypothesized that the SHP-silica nanoparticles may have a very low but non-negligible surface density of hydroxyl groups, leading to a small loss in cross-link density caused by exchange reactions between the network matrix and the filler surface functional groups. To test this hypothesis, we did FTIR spectroscopy on the three types of silica nanoparticles to compare qualitatively the hydroxyl group content. The FTIR spectra were normalized using the siloxane group (Si−O−Si) stretching peaks (1000−1200 cm−1), and the silanol group (Si−OH) vibration peaks (3200−3600 cm−1) were examined, as shown in Figure S2. Because the N−H stretching peak of primary amine groups also appears in this F


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the PHU network composites cannot be attributed wholly or in part to qualitative differences in filler dispersion. 3.3. Effect of Filler Surface Functionality on the Rate and Apparent Activation Energy of Stress Relaxation at Elevated Temperature. We investigated the effect of filler loading and surface functionality on stress relaxation behaviors using 1st molded neat PHU network and PHU network composite samples. First, the rates of stress relaxation at 140 °C (reprocessing temperature) were evaluated; see Figure 5a.

each PHU network composite is attributable to the effect of filler surface functionality and not to different levels of deleterious side reactions. Also, all molded PHU networks and network composite samples have nearly identical Tg’s (Table 1), indicating that Tg is insensitive to small or modest reductions in cross-link density at low cross-link density. 3.2. Mechanical Property Enhancement in PHU Network Composites. We carried out uniaxial tensile testing on PHU network and network composites to evaluate the effect of nanofillers on mechanical reinforcement. As the Tg’s of these samples are close to room temperature, tensile tests were performed at 40 °C to ensure that the samples were in their rubbery states during the tests (tensile tests of networks taken in the glassy state or near Tg may more significantly reflect the (near) glassy nature than the cross-link density). Young’s modulus, ultimate tensile strength, and strain-at-break for 1st molded neat network and network composites are summarized in Table 2. Relative to neat PHU networks, Table 2. Tensile Properties (at 40 °C) of the 1st Molded neat PHU Networks and PHU Network Composites sample neat PHU OH-cPHU NH-cPHU SHP-cPHU-4 SHP-cPHU-2

Young’s modulus (MPa) 1.9 2.3 1.6 2.7 2.5

± ± ± ± ±

0.3 0.2 0.2 0.3 0.1

ultimate tensile strength (MPa) 1.3 1.4 1.4 1.8 1.7

± ± ± ± ±

0.2 0.1 0.1 0.1 0.1

Figure 5. (a) Stress relaxation curves of the neat PHU network and PHU network composites at 140 °C under a strain of 5%. (b) Arrhenius apparent activation energy of stress relaxation for the neat PHU network and PHU network composites.

strain-at-break (%) 190 224 202 190 217

± ± ± ± ±

31 12 13 8 8

Compared with the neat network, each network composite exhibits slower stress relaxation, with SHP-cPHU-4 being the slowest, followed by SHP-cPHU-2, OH-cPHU, and NHcPHU. When nanofillers are incorporated into the network matrix, the mobility of the network is reduced because of the fillers serve as nonflowable solids anchored in the polymer chains. At the polymer−filler interfaces, the polymer chains tend to adhere to or have suppressed mobility near the filler surface (relative to the mobility in the neat network). Also, the nanoparticles may block some functional groups on different polymer chains from contacting with each other, reducing the possibility of network rearrangement. This also explains why SHP-cPHU-2 shows faster stress relaxation than SHP-cPHU-4. Compared to network composites containing SHP-silica, OHcPHU and NH-cPHU have much higher stress relaxation rates; the time required to relax to 40% of initial stress at 140 °C decreases from 1710 s in SHP-cPHU-4 to 390 and 320 s in OH-cPHU and NH-cPHU, respectively. Because the hydroxyl and amine surface functional groups are capable of participating in the dynamic chemistries during network rearrangement, they can aid stress relaxation as compared with the nonreactive SHP-silica. These findings are in agreement with two previous studies, which also reported faster stress relaxation when the surface functionality on silica nanoparticles is able to participate in the dynamic chemistries.56,57 In addition, the faster relaxation rate of the NH-cPHU sample compared with the OH-cPHU sample indicates that the stress can be relaxed more quickly when a higher fraction of the reversible reaction is present than when a higher fraction of the exchange reaction is present. We also studied stress relaxation as a function of temperature from 120 to 160 °C to determine the apparent Arrhenius activation energy of stress relaxation. Most studies reporting the apparent activation energy in CANs13,14,20,24,27,28 have estimated the characteristic relaxation time (and also the average relaxation time), τ*, as the time

Young’s modulus increases ∼42% in PHU network composites containing 4 wt % SHP-silica nanoparticles. This enhancement is consistent with effective load transfer at filler−matrix interfaces under very small strain and correlates well with the improvement in the rubbery plateau modulus observed by DMA. The low Young’s modulus of NH-cPHU, lowest of the PHU network composites studied here, is consistent with a greater extent of property loss in the 1st molding process and is also in accordance with the DMA results and accompanying discussion given above. Notably, relative to neat PHU networks, there are also significant enhancements in ultimate tensile strength in PHU network composites containing SHPsilica nanoparticles (by ∼38 and ∼31% with 4 and 2 wt % nanofiller, respectively), demonstrating the effectiveness of using SHP-silica nanoparticles for mechanical reinforcement. Within error, all PHU network composites considered here exhibit strain-at-break values that are unchanged from that of the neat PHU network. It is well known that the quality of filler dispersion can play an important role in the reinforcement of polymer composites and that surface modification can influence filler dispersion.66−68 We took SEM images of all PHU network composites at 4 wt % loading and found similar dispersion levels. See Figures S11−S13. Both single silica nanoparticles and aggregations consisting of a few to hundreds of nanoparticles were observed. No difference in dispersion could be distinguished among as-synthesized and 1st, 2nd, and 3rd molded samples, indicating that the reprocessing step as done here did not significantly affect dispersion. Thus, for the filler loading and reprocessing steps employed here, neither filler surface functionality nor the network rearrangement process has a significant effect on the extent of filler dispersion. In turn, this means that differences in mechanical properties in G


Research Article


carbonate aminolysis reaction is smaller than that of the transcarbamoylation exchange reaction since the apparent activation energy represents a combination of the two, and OH-cPHU and NH-cPHU differ by the fraction of each chemistry involved in the network rearrangement process.

required for stress to relax to 1/e of its initial value, consistent with the assumption of a single-exponential decay response. We employed the Kohlrausch−Williams−Watts (KWW) stretched exponential decay function for fitting stress relaxation (in a recent study of a different reprocessable polymer network,12 we found that the KWW function provided a much better fit to stress relaxation data than a singleexponential decay). The KWW stretched exponential decay can be expressed as follows12,69 σ (t ) = exp{−(t /τ *)β } σ0

4. CONCLUSIONS Taking advantage of dynamic covalent bonds in hydroxyurethane linkages, we developed novel reprocessable PHU/silica network composites using silica nanoparticles with several different surface functional groups: hydroxyl groups (present with standard silica nanoparticles), amine groups, and the near absence of surface functional groups (using superhydrophobic (SHP) silica nanoparticles). Hydroxyl groups and amine groups can participate in transcarbamoylation exchange reactions and reversible cyclic carbonate aminolysis reactions, respectively, that constitute the dynamic chemistry associated with reprocessable PHU networks, providing routes for loss of cross-link density after reprocessing. Thus, when 4 wt % nanosilica with hydroxyl or amine surface groups is added to the PHU network, within error there is no enhancement of rubbery-state tensile properties relative to neat PHU. More importantly, in contrast to neat PHU which exhibits full recovery of rubbery plateau E′ values after two reprocessing steps, these PHU network composites exhibit significant reductions in rubbery plateau E′ values and thus cross-link density after one and two reprocessing steps. When 2 or 4 wt % SHP silica is added to the network, Young’s modulus and tensile strength are significantly enhanced, whereas strain-at-break is maintained constant relative to the neat network. More importantly, 2 and 4 wt % SHP-cPHU samples exhibit full recovery within error of the rubbery plateau E′ value (at 100 °C) after one reprocessing step, indicating full recovery of cross-link density in the network composite after this reprocessing step. After a second reprocessing step, these samples exhibit 10−14% reductions in rubbery plateau E′ values, outside experimental error, consistent with small reductions in cross-link density. Such reductions accompanying the second reprocessing step are consistent with small, residual levels of surface hydroxyl groups observed via FTIR spectroscopy on the SHP silica surfaces. High-temperature stress relaxation was also characterized. Relative to neat networks, the presence of filler increases the average stress relaxation time, to a small extent for silica with surface functional groups and to a large extent for SHP silica. Little or no change is observed in apparent activation energy of the stress relaxation when silica nanoparticles with surface functional groups are added to PHU networks. In contrast, addition of SHP silica leads to substantial increases in activation energy. These results indicate that if reprocessable networks with full recovery of cross-link density can be developed for a particular polymer species, then it is also possible to make similar robust, reprocessable polymer network composites. However, special care must be taken to ensure that the filler surface functionality does not participate deleteriously in the chemistries associated with the dynamic cross-links.

(1)

where σ(t)/σ0 is the normalized stress at time t, τ* is the characteristic relaxation time, and β is the exponent that controls the shape of the stretched exponential decay and reflects the breadth of the relaxation distribution. β takes values in the range 0 < β ≤ 1, with β = 1 being a single-exponential decay response and β ≪ 1 indicating an extremely broad distribution of relaxation times.69 For a KWW decay, the average relaxation time, ⟨τ⟩, is given by69 ⟨τ ⟩ =

τ * Γ(1/β) β

(2)

where Γ is the Gamma function. Table S1 shows the detailed fitting results for the neat PHU network, which were obtained from both KWW stretched exponential and single-exponential decay analyses. Although the stress relaxation data are not far from being well expressed by a single-exponential decay function, the stretched exponential decay fitting provides higher R2 values; thus, we did all further analyses using fits from KWW exponential decay. Figure 5b shows the Arrhenius fit of apparent activation energy of stress relaxation; the results are summarized in Table 3. The apparent activation energy of the neat PHU network Table 3. ⟨τ⟩ at 140 °C and Apparent Activation Energy of Stress Relaxation of the 1st Molded Neat PHU Network and PHU Network Composites sample neat PHU OH-cPHU NH-cPHU SHP-cPHU-4 SHP-cPHU-2

apparent activation energy (kJ/mol)b

⟨τ⟩ at 140 °C (s)a 281 491 369 2265 1362

± ± ± ± ±

17 26 19 115 61

a Determined from three samples. range of 120 ≤ T ≤ 160 °C.

135.2 142.2 132.8 159.1 155.0

± ± ± ± ±

3.1 1.6 3.6 6.5 2.4

adjusted R2 0.998 >0.999 0.997 0.993 0.999

b

Determined over temperature

was determined as 135.2 ± 3.1 kJ/mol, in reasonable agreement with values obtained in some other reprocessable PHU networks.28,63 When SHP-silica is added to the network, the apparent activation energy increases; the introduction of nonreactive fillers provides a significant barrier for dynamic chemistries and stress relaxation to occur. The OH-cPHU and NH-cPHU samples give apparent activation energies smaller than those of SHP-cPHU samples, with the NH-cPHU value being the same within error as that for the neat PHU network. These results reveal that the extra barrier to stress relaxation because of filler incorporation can be relieved if the filler surface functionality can participate in the dynamic chemistries associated with network rearrangement. Moreover, our results suggest that the activation energy of the reversible cyclic

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ASSOCIATED CONTENT

* Supporting Information S

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami.8b19100. H


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FTIR spectra; SEM images; stress relaxation fitting results; DSC results (PDF)

(15) Zheng, N.; Fang, Z.; Zou, W.; Zhao, Q.; Xie, T. Thermoset Shape-Memory Polyurethane with Intrinsic Plasticity Enabled by Transcarbamoylation. Angew. Chem., Int. Ed. 2016, 55, 11421−11425. (16) Chen, X.; Li, L.; Jin, K.; Torkelson, J. M. Reprocessable Polyhydroxyurethane Networks Exhibiting Full Property Recovery and Concurrent Associative and Dissociative Dynamic Chemistry via Transcarbamoylation and Reversible Cyclic Carbonate Aminolysis. Polym. Chem. 2017, 8, 6349−6355. (17) Lu, Y.-X.; Guan, Z. Olefin Metathesis for Effective Polymer Healing via Dynamic Exchange of Strong Carbon−Carbon Double Bonds. J. Am. Chem. Soc. 2012, 134, 14226−14231. (18) Zheng, P.; McCarthy, T. J. A Surprise from 1954: Siloxane Equilibration Is a Simple, Robust, and Obvious Polymer Self-Healing Mechanism. J. Am. Chem. Soc. 2012, 134, 2024−2027. (19) Gordon, M. B.; French, J. M.; Wagner, N. J.; Kloxin, C. J. Dynamic Bonds in Covalently Crosslinked Polymer Networks for Photoactivated Strengthening and Healing. Adv. Mater. 2015, 27, 8007−8010. (20) Han, J.; Liu, T.; Hao, C.; Zhang, S.; Guo, B.; Zhang, J. A Catalyst-Free Epoxy Vitrimer System Based on Multifunctional Hyperbranched Polymer. Macromolecules 2018, 51, 6789−6799. (21) Obadia, M. M.; Mudraboyina, B. P.; Serghei, A.; Montarnal, D.; Drockenmuller, E. Reprocessing and Recycling of Highly CrossLinked Ion-Conducting Networks through Transalkylation Exchanges of C−N Bonds. J. Am. Chem. Soc. 2015, 137, 6078−6083. (22) Fang, Z.; Zheng, N.; Zhao, Q.; Xie, T. Healable, Reconfigurable, Reprocessable Thermoset Shape Memory Polymer with Highly Tunable Topological Rearrangement Kinetics. ACS Appl. Mater. Interfaces 2017, 9, 22077−22082. (23) Röttger, M.; Domenech, T.; van der Weegen, R.; Breuillac, A.; Nicolaÿ, R.; Leibler, L. High-Performance Vitrimers from Commodity Thermoplastics through Dioxaborolane Metathesis. Science 2017, 356, 62−65. (24) Ishibashi, J. S. A.; Kalow, J. A. Vitrimeric Silicone Elastomers Enabled by Dynamic Meldrum’s Acid-Derived Cross-Links. ACS Macro Lett. 2018, 7, 482−486. (25) Chen, Y.; Tang, Z.; Zhang, X.; Liu, Y.; Wu, S.; Guo, B. Covalently Cross-Linked Elastomers with Self-Healing and Malleable Abilities Enabled by Boronic Ester Bonds. ACS Appl. Mater. Interfaces 2018, 10, 24224−24231. (26) Gao, W.; Bie, M.; Liu, F.; Chang, P.; Quan, Y. Self-Healable and Reprocessable Polysulfide Sealants Prepared from Liquid Polysulfide Oligomer and Epoxy Resin. ACS Appl. Mater. Interfaces 2017, 9, 15798−15808. (27) Zhang, B.; Kowsari, K.; Serjouei, A.; Dunn, M. L.; Ge, Q. Reprocessable Thermosets for Sustainable Three-Dimensional Printing. Nat. Commun. 2018, 9, No. 1831. (28) Fortman, D. J.; Snyder, R. L.; Sheppard, D. T.; Dichtel, W. R. Rapidly Reprocessable Cross-Linked Polyhydroxyurethanes Based on Disulfide Exchange. ACS Macro Lett. 2018, 7, 1226−1231. (29) Ma, S.; Webster, D. C. Degradable Thermosets Based on Labile Bonds or Linkages: A Review. Prog. Polym. Sci. 2017, 76, 65−110. (30) Liu, H.; Nelson, A. Z.; Ren, Y.; Yang, K.; Ewoldt, R. H.; Moore, J. S. Dynamic Remodeling of Covalent Networks via Ring-Opening Metathesis Polymerization. ACS Macro Lett. 2018, 7, 933−937. (31) Wang, C.; Goldman, T. M.; Worrell, B. T.; McBride, M. K.; Alim, M. D.; Bowman, C. N. Recyclable and Repolymerizable Thiol− X Photopolymers. Mater. Horiz. 2018, 5, 1042−1046. (32) Scott, T. F.; Schneider, A. D.; Cook, W. D.; Bowman, C. N. Photoinduced Plasticity in Cross-Linked Polymers. Science 2005, 308, 1615−1617. (33) Zheng, N.; Hou, J.; Xu, Y.; Fang, Z.; Zou, W.; Zhao, Q.; Xie, T. Catalyst-Free Thermoset Polyurethane with Permanent Shape Reconfigurability and Highly Tunable Triple-Shape Memory Performance. ACS Macro Lett. 2017, 6, 326−330. (34) Giannelis, E. P. Polymer Layered Silicate Nanocomposites. Adv. Mater. 1996, 8, 29−35.

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

John M. Torkelson: 0000-0002-4875-4827 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We acknowledge support of Northwestern University via discretionary funds associated with a Walter P. Murphy Professorship (J.M.T.), ISEN Fellowships (X.C.; L.L.), and a 3M Fellowship (X.C.). We thank Prof. L.C. Brinson for generously providing access to the DMA instrument at Northwestern University.

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REFERENCES

(1) Chen, X.; Dam, M. A.; Ono, K.; Mal, A.; Shen, H.; Nutt, S. R.; Sheran, K.; Wudl, F. A Thermally Re-Mendable Cross-Linked Polymeric Material. Science 2002, 295, 1698−1702. (2) Montarnal, D.; Capelot, M.; Tournilhac, F.; Leibler, L. SilicaLike Malleable Materials from Permanent Organic Networks. Science 2011, 334, 965−968. (3) Wojtecki, R. J.; Meador, M. A.; Rowan, S. J. Using the Dynamic Bond to Access Macroscopically Responsive Structurally Dynamic Polymers. Nat. Mater. 2011, 10, 14−27. (4) Yang, Y.; Ding, X.; Urban, M. W. Chemical and Physical Aspects of Self-Healing Materials. Prog. Polym. Sci. 2015, 49−50, 34−59. (5) Denissen, W.; Winne, J. M.; Du Prez, F. E. Vitrimers: Permanent Organic Networks with Glass-Like Fluidity. Chem. Sci. 2016, 7, 30− 38. (6) Zou, W.; Dong, J.; Luo, Y.; Zhao, Q.; Xie, T. Dynamic Covalent Polymer Networks: From Old Chemistry to Modern Day Innovations. Adv. Mater. 2017, 29, No. 1606100. (7) Zhang, Z. P.; Rong, M. Z.; Zhang, M. Q. Polymer Engineering Based on Reversible Covalent Chemistry: A Promising Innovative Pathway Towards New Materials and New Functionalities. Prog. Polym. Sci. 2018, 80, 39−93. (8) Kloxin, C. J.; Scott, T. F.; Adzima, B. J.; Bowman, C. N. Covalent Adaptable Networks (CANs): A Unique Paradigm in Cross-Linked Polymers. Macromolecules 2010, 43, 2643−2653. (9) Oehlenschlaeger, K. K.; Mueller, J. O.; Brandt, J.; Hilf, S.; Lederer, A.; Wilhelm, M.; Graf, R.; Coote, M. L.; Schmidt, F. G.; Barner-Kowollik, C. Adaptable Hetero Diels−Alder Networks for Fast Self-Healing under Mild Conditions. Adv. Mater. 2014, 26, 3561− 3566. (10) Higaki, Y.; Otsuka, H.; Takahara, A. A Thermodynamic Polymer Cross-Linking System Based on Radically Exchangeable Covalent Bonds. Macromolecules 2006, 39, 2121−2125. (11) Jin, K.; Li, L.; Torkelson, J. M. Recyclable Crosslinked Polymer Networks via One-Step Controlled Radical Polymerization. Adv. Mater. 2016, 28, 6746−6750. (12) Li, L.; Chen, X.; Jin, K.; Torkelson, J. M. Vitrimers Designed Both to Strongly Suppress Creep and to Recover Original Cross-Link Density after Reprocessing: Quantitative Theory and Experiments. Macromolecules 2018, 51, 5537−5546. (13) Denissen, W.; Rivero, G.; Nicolaÿ, R.; Leibler, L.; Winne, J. M.; Du Prez, F. E. Vinylogous Urethane Vitrimers. Adv. Funct. Mater. 2015, 25, 2451−2457. (14) Fortman, D. J.; Brutman, J. P.; Cramer, C. J.; Hillmyer, M. A.; Dichtel, W. R. Mechanically Activated, Catalyst-Free Polyhydroxyurethane Vitrimers. J. Am. Chem. Soc. 2015, 137, 14019−14022. I


Research Article


(54) Qiu, M.; Wu, S.; Tang, Z.; Guo, B. Exchangeable Interfacial Crosslinks Towards Mechanically Robust Elastomer/Carbon Nanotubes Vitrimers. Compos. Sci. Technol. 2018, 165, 24−30. (55) Xu, C.; Cui, R.; Fu, L.; Lin, B. Recyclable and Heat-Healable Epoxidized Natural Rubber/Bentonite Composites. Compos. Sci. Technol. 2018, 167, 421−430. (56) Legrand, A.; Soulié-Ziakovic, C. Silica−Epoxy Vitrimer Nanocomposites. Macromolecules 2016, 49, 5893−5902. (57) Huang, Z.; Wang, Y.; Zhu, J.; Yu, J.; Hu, Z. Surface Engineering of Nanosilica for Vitrimer Composites. Compos. Sci. Technol. 2018, 154, 18−27. (58) Guan, J.; Song, Y.; Lin, Y.; Yin, X.; Zuo, M.; Zhao, Y.; Tao, X.; Zheng, Q. Progress in Study of Non-Isocyanate Polyurethane. Ind. Eng. Chem. Res. 2011, 50, 6517−6527. (59) Blattmann, H.; Fleischer, M.; Bähr, M.; Mü lhaupt, R. Isocyanate- and Phosgene-Free Routes to Polyfunctional Cyclic Carbonates and Green Polyurethanes by Fixation of Carbon Dioxide. Macromol. Rapid Commun. 2014, 35, 1238−1254. (60) Maisonneuve, L.; Lamarzelle, O.; Rix, E.; Grau, E.; Cramail, H. Isocyanate-Free Routes to Polyurethanes and Poly(hydroxy Urethane)s. Chem. Rev. 2015, 115, 12407−12439. (61) Beniah, G.; Chen, X.; Uno, B. E.; Liu, K.; Leitsch, E. K.; Jeon, J.; Heath, W. H.; Scheidt, K. A.; Torkelson, J. M. Combined Effects of Carbonate and Soft-Segment Molecular Structures on the Nanophase Separation and Properties of Segmented Polyhydroxyurethane. Macromolecules 2017, 50, 3193−3203. (62) Beniah, G.; Heath, W. H.; Jeon, J.; Torkelson, J. M. Tuning the Properties of Segmented Polyhydroxyurethanes via Chain Extender Structure. J. Appl. Polym. Sci. 2017, 134, No. 44942. (63) Fortman, D. J.; Brutman, J. P.; Hillmyer, M. A.; Dichtel, W. R. Structural Effects on the Reprocessability and Stress Relaxation of Cross-Linked Polyhydroxyurethanes. J. Appl. Polym. Sci. 2017, 134, No. 44984. (64) Matsukizono, H.; Endo, T. Reworkable Polyhydroxyurethane Films with Reversible Acetal Networks Obtained from Multifunctional Six-Membered Cyclic Carbonates. J. Am. Chem. Soc. 2018, 140, 884−887. (65) Flory, P. J. Principles of Polymer Chemistry; Cornell University Press, 1953. (66) Zou, H.; Wu, S.; Shen, J. Polymer/Silica Nanocomposites: Preparation, Characterization, Properties, and Applications. Chem. Rev. 2008, 108, 3893−3957. (67) Ma, P.-C.; Siddiqui, N. A.; Marom, G.; Kim, J.-K. Dispersion and Functionalization of Carbon Nanotubes for Polymer-Based Nanocomposites: A Review. Composites, Part A 2010, 41, 1345−1367. (68) Askar, S.; Li, L.; Torkelson, J. M. Polystyrene-Grafted Silica Nanoparticles: Investigating the Molecular Weight Dependence of Glass Transition and Fragility Behavior. Macromolecules 2017, 50, 1589−1598. (69) Dhinojwala, A.; Hooker, J. C.; Torkelson, J. M. Retardation of Rotational Reorientation Dynamics in Polymers near the Glass Transition: A Novel Study over Eleven Decades in Time Using Second-Order Non-Linear Optics. J. Non-Cryst. Solids 1994, 172− 174, 286−296.

(35) Stankovich, S.; Dikin, D. A.; Dommett, G. H. B.; Kohlhaas, K. M.; Zimney, E. J.; Stach, E. A.; Piner, R. D.; Nguyen, S. T.; Ruoff, R. S. Graphene-Based Composite Materials. Nature 2006, 442, 282−286. (36) Paul, D. R.; Robeson, L. M. Polymer Nanotechnology: Nanocomposites. Polymer 2008, 49, 3187−3204. (37) Spitalsky, Z.; Tasis, D.; Papagelis, K.; Galiotis, C. Carbon Nanotube−Polymer Composites: Chemistry, Processing, Mechanical and Electrical Properties. Prog. Polym. Sci. 2010, 35, 357−401. (38) Iyer, K. A.; Torkelson, J. M. Importance of Superior Dispersion Versus Filler Surface Modification in Producing Robust Polymer Nanocomposites: The Example of Polypropylene/Nanosilica Hybrids. Polymer 2015, 68, 147−157. (39) Peterson, A. M.; Jensen, R. E.; Palmese, G. R. Thermoreversible and Remendable Glass−Polymer Interface for Fiber-Reinforced Composites. Compos. Sci. Technol. 2011, 71, 586−592. (40) Xu, Z.; Zhao, Y.; Wang, X.; Lin, T. A Thermally Healable Polyhedral Oligomeric Silsesquioxane (POSS) Nanocomposite Based on Diels−Alder Chemistry. Chem. Commun. 2013, 49, 6755−6757. (41) Li, J.; Zhang, G.; Deng, L.; Zhao, S.; Gao, Y.; Jiang, K.; Sun, R.; Wong, C. In Situ Polymerization of Mechanically Reinforced, Thermally Healable Graphene Oxide/Polyurethane Composites Based on Diels−Alder Chemistry. J. Mater. Chem. A 2014, 2, 20642−20649. (42) Lin, C.; Sheng, D.; Liu, X.; Xu, S.; Ji, F.; Dong, L.; Zhou, Y.; Yang, Y. A Self-Healable Nanocomposite Based on Dual-Crosslinked Graphene Oxide/Polyurethane. Polymer 2017, 127, 241−250. (43) Schäfer, S.; Kickelbick, G. Double Reversible Networks: Improvement of Self-Healing in Hybrid Materials via Combination of Diels−Alder Cross-Linking and Hydrogen Bonds. Macromolecules 2018, 51, 6099−6110. (44) Zhang, L.; Tian, X.; Malakooti, M. H.; Sodano, H. A. Novel Self-Healing CFRP Composites with High Glass Transition Temperatures. Compos. Sci. Technol. 2018, 168, 96−103. (45) Taynton, P.; Ni, H.; Zhu, C.; Yu, K.; Loob, S.; Jin, Y.; Qi, H. J.; Zhang, W. Repairable Woven Carbon Fiber Composites with Full Recyclability Enabled by Malleable Polyimine Networks. Adv. Mater. 2016, 28, 2904−2909. (46) Yu, K.; Shi, Q.; Dunn, M. L.; Wang, T.; Qi, H. J. Carbon Fiber Reinforced Thermoset Composite with near 100% Recyclability. Adv. Funct. Mater. 2016, 26, 6098−6106. (47) Wang, S.; Xing, X.; Zhang, X.; Wang, X.; Jing, X. RoomTemperature Fully Recyclable Carbon Fibre Reinforced Phenolic Composites through Dynamic Covalent Boronic Ester Bonds. J. Mater. Chem. A 2018, 6, 10868−10878. (48) Chabert, E.; Vial, J.; Cauchois, J.-P.; Mihaluta, M.; Tournilhac, F. Multiple Welding of Long Fiber Epoxy Vitrimer Composites. Soft Matter 2016, 12, 4838−4845. (49) Yang, Z.; Wang, Q.; Wang, T. Dual-Triggered and Thermally Reconfigurable Shape Memory Graphene-Vitrimer Composites. ACS Appl. Mater. Interfaces 2016, 8, 21691−21699. (50) Yan, P.; Zhao, W.; Jiang, L.; Wu, B.; Hu, K.; Yuan, Y.; Lei, J. Reconfiguration and Shape Memory Triggered by Heat and Light of Carbon Nanotube−Polyurethane Vitrimer Composites. J. Appl. Polym. Sci. 2018, 135, No. 45784. (51) de Luzuriaga, A. R.; Martin, R.; Markaide, N.; Rekondo, A.; Cabañ ero, G.; Rodríguez, J.; Odriozola, I. Epoxy Resin with Exchangeable Disulfide Crosslinks to Obtain Reprocessable, Repairable and Recyclable Fiber-Reinforced Thermoset Composites. Mater. Horiz. 2016, 3, 241−247. (52) Cai, C.; Zhang, Y.; Zou, X.; Zhang, R.; Wang, X.; Wu, Q.; Sun, P. Rapid Self-Healing and Recycling of Multiple-Responsive Mechanically Enhanced Epoxy Resin/Graphene Nanocomposites. RSC Adv. 2017, 7, 46336−46343. (53) Liu, Y.; Tang, Z.; Chen, Y.; Zhang, C.; Guo, B. Engineering of β-Hydroxyl Esters into Elastomer−Nanoparticle Interface toward Malleable, Robust, and Reprocessable Vitrimer Composites. ACS Appl. Mater. Interfaces 2018, 10, 2992−3001. J


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