Highly Flexible Hybrid Polymer Aerogels and Xerogels Based on

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Highly Flexible Hybrid Polymer Aerogels and Xerogels Based on Resorcinol-Formaldehyde with Enhanced Elastic Stiffness and Recoverability: Insights into the Origin of Their Mechanical Properties George Hasegawa,*,† Taiyo Shimizu,‡ Kazuyoshi Kanamori,‡ Ayaka Maeno,§ Hironori Kaji,§ and Kazuki Nakanishi‡ †

Department of Applied Chemistry, Graduate School of Engineering, Kyushu University, Motooka, Nishi-ku, Fukuoka, 819-0395, Japan ‡ Department of Chemistry, Graduate School of Science, Kyoto University, Kitashirakawa, Sakyo-ku, Kyoto 606-8502, Japan § Institute for Chemical Research, Kyoto University, Gokasho, Uji, 611-0011, Japan S Supporting Information *

ABSTRACT: Flexible low-density materials, such as aerogels and polymer foams, have received increasing attention as energy absorbers and cushions that protect artificial products and human bodies. Microscopic geometry is a crucial factor determining their mechanical functions, i.e. strength and toughness (flexibility). However, it is a formidable challenge to combine these two properties because they are mutually elusive in general; stiff materials are brittle, while flexible ones are soft. Here, we demonstrate lightweight porous polymeric materials based on a common phenolic resin, resorcinolformaldehyde (RF) gels, with salient combinatorial properties of high stiffness (up to 100 MPa) and good recoverable compressibility (against 80−90% strain), which can deliver remarkable energy absorption and dissipation performances repetitively. The detailed investigation reveals that the unique mechanical features originate from the synergetic effect of interdigitated hard and soft components in polymer matrices as well as exquisitely designed highly branched microstructures both generated through the spontaneous supramolecular self-assembly of the nonionic block copolymer (F127) and RF oligomer, which is essentially analogous to how natural organisms create biological structural materials, e.g. nacre and bone. scaling as a function of the bulk density (ρ) by E ∼ ρn,8,9 the reinforced aerogels inevitably sacrifice their lightweight feature. The trade-off relationship between stiffness and bulk density drove researchers to focus on another mechanical function, viz. flexibility. As with the pioneering works pertaining to aerogels,3 the studies on flexible aerogels started with polyorganosiloxane materials. Kanamori et al. reported that the polymethylsilsesquioxane (PMSQ) aerogels prepared solely from methyltrimethoxysilane show recoverable deformation against uniaxial compression up to 80% in addition to high optical transparency.10 This unique flexible trait is derived from the moderately reduced cross-linking density of the gel networks as well as the repulsive interaction between hydrophobic moieties, allowing for the ambient pressure drying by evaporation toward PMSQ xerogels with similar pore properties to those of the correspondent aerogels owing to the spring-back effect.11

1. INTRODUCTION Highly porous monolithic materials, commonly referred to as aerogels, have intrigued scientists and engineers due to their unique characteristics, such as low density, low thermal conductivity and low sound velocity, which originate from fine internal void spaces and open-pore geometry. Despite a large number of potential applications developed by many researchers since the invention of silica aerogels in 1931,1 aerogels inherently have serious limitations for widespread practical use except for in the specific applications as cosmic dust capturers and Cherenkov radiation detectors2−5 because of their friable nature. To cope with the mechanical friability of aerogels, early works attempted to consolidate the gel networks by various approaches, e.g. aging, modification of surface functionality, cross-linking with polymers and hybridization with reinforcing agents.3−7 Some of these strategies can improve the mechanical strength of aerogels, which, however, in turn incur a laborious and costly synthesis process. In addition, since the Young’s modulus (E) obeys the power-law © 2017 American Chemical Society

Received: November 3, 2016 Revised: February 21, 2017 Published: February 22, 2017 2122

DOI: 10.1021/acs.chemmater.6b04706 Chem. Mater. 2017, 29, 2122−2134

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Chemistry of Materials

Very recently, we have reported the facile one-pot synthesis of hierarchically porous RF gels comprising mesoporous nanorods with 2-D hexagonal ordering by integrating the micellar templating and phase separation techniques.42 The supramolecular self-assembly of amphiphilic triblock copolymer/RF oligomer yields columnar nanobuilding blocks affected by the anisotropy of ordered cylindrical mesopores, which are organized to form the isotropic macroporous structure. Since the microscopic architecture is deemed a kind of highly branched truss structure where several struts connect with each other at a node, the carbon monoliths that inherit the exquisite porous morphology are imbued with high mechanical strength.42 In this study, we have synthesized a series of RF aerogels and xerogels with multifold morphologies including truss-like microstructures by regulating the polymerization-induced phase separation with the aid of an amphiphilic block copolymer, Pluronic F127 (EO106PO70EO106). The obtained RF gels display diverse mechanical properties depending on their morphologies, all of which involve notable flexibilities that have not been reported for the conventional RF aerogels. The origin of the unique mechanical characteristics is discussed from the aspects of the microscopic geometry and chemical composition of the polymer backbones. The optimization of the two factors provides RF gels with an unprecedented mechanical trait combining high stiffness and salient recoverable compressibility against up to 90% deformation. The energy absorption capabilities of the stiff and flexible polymer monoliths have also been investigated in detail.

Incorporation of bifunctional silicon alkoxides to this system gives rise to the marshmallow-like xerogels with improved flexibility both for compression and for bending over a wide temperature range.12−14 The notable advantage of such polyorganosiloxanes lies in the facile functionalization like superamphiphobicity either by employing different organoalkoxysilane precursors or by a postreaction with surface silanol groups.14 Lately, a variety of flexible aerogels other than polyorganosiloxanes have been developed. Aerogels composed of onedimensional (1-D) nanofibers create a new category called “nanofibrous” aerogels. The fibrous components range from biopolymers15−17 and carbon nanotubes18−20 to inorganic combounds.21,22 Taking this concept a step further, aerogels consisting of two-dimensional (2-D) graphene sheets have been intensively studied as well.23−27 The graphene aerogels are characterized by extraordinarily low density (as low as 180 μg cm−3),24 high electric conductivity, and excellent reversible recovery from large strains with negligible Poisson’s ratio.27 More practical flexible lightweight materials are a family of polymer foams and sponges,28−33 some of which can be routinely found in our daily life. The above-mentioned flexible aerogels have a wide range of applications, such as oil absorbents for environment remediation, vibration damping, and mechanical energy dissipation.18−20,25−33 With respect to the mechanical energy absorption, two important factors determine the performance of flexible aerogels: stiffness and compressibility, which can be represented by the combination of mechanical strength, elastic modulus, and resilience from large compressive strains.34,35 However, the discrete modulation of the two key properties faces a great obstacle because they are mutually exclusive in general.36 Materials that can recover from large strains are usually very soft (extremely low Young’s modulus), while stiff materials are often hard and brittle and hardly revert after a large deformation. A film comprising vertically aligned carbon nanotubes reported by Cao and co-workers seems to be an only material showing reversible deformation under large strain (∼85%) together with high elastic modulus (>50 MPa).18 This striking mechanical response is, however, anisotropic due to the structural anisotropy, and the production and cost issues also need to be addressed for commercial viability. The mechanical properties of aerogels are dictated by the chemical configuration of solid phase, the pore volume fraction (porosity), and the porous morphology. Hence, the refined architectural design of porous geometry in aerogels is of great importance for the individual control over mechanical strength and reversible compressibility. Resorcinol-formaldehyde (RF) aerogels are mundane yet promising materials in terms of controlling pore characteristics because the polycondensation of RF networks in aqueous media is highly tunable by varying the synthesis parameters, such as pH and temperature, in analogy to the alkoxysilane-derived sol−gel systems. In the past quarter century since the first report by Pekala,37 a myriad of papers on RF aerogels has been published regarding the morphological control.38−41 Nevertheless, only a few flexible RF aerogels can be found in the literature published so far40,41 and are still far from the application to energy absorbers. This is because the previous studies lack the adequate modulation of phase separation during polymerization, resulting in polymer scaffolds composed of nanoparticles whose interparticle necks are responsible for the fragility of those RF aerogels.

2. EXPERIMENTAL SECTION 2.1. Synthesis of RF Aerogels and Xerogels. The RF gels were prepared by the one-pot sol−gel process according to the procedure reported previously42 with minor modifications. In a typical synthesis, 3.0 g of F127, 3.0 mL of 1,3,5-trimethylbenzene (TMB), and 3.0 mL of benzyl alcohol (BzOH) were added to a given amount of a triethylene glycol (TEG)/1 M HCl aq. mixed solvent in a glass tube followed by stirring at room temperature to obtain a homogeneous solution. Then, 2.2 g of resorcinol was dissolved in the resultant solution, and 6.0 mL of formaldehyde solution (37 wt % in H2O containing 10−15% methanol) was subsequently added. After stirring for 30 min at room temperature, the sol was kept at 60 °C for 48 h for gelation and aging. The obtained wet gels were subjected to solvent exchange with 2-propanol followed by drying from supercritical CO2 at 80 °C and 14 MPa or evaporative drying under ambient pressure. The supercritically dried sample notations are defined as TEGx-Hy, where x and y specify the volume (in mL) of TEG and 1 M HCl aq., respectively. The xerogels obtained by evaporative drying are denoted as TEGx-Hy-X. Some samples were heat-treated at different temperatures for 30 min under nitrogen atmosphere. The heat-treated samples are designated as TEGx-Hy-T and TEGx-Hy-X-T, where T represents the calcination temperature (°C). 2.2. Characterization. The porous morphologies of the samples in the micrometer range were observed with a scanning electron microscope (SEM, JSM-6060S, JEOL). Nitrogen adsorption− desorption measurements (Belsorp Max, Bel Japan Inc.) were performed at 77 K to characterize the micro- and mesoporous properties. Helium pycnometry (Pentapyc 5200e, Quantachrome Instruments) was employed to determine the skeletal densities. Smallangle X-ray scattering (SAXS) measurements of the monolithic specimens were carried out with a RINT system (RINT Ultima III, Rigaku Corp.) equipped with a Cu Kα X-ray generator (λ = 0.154 nm). The parameters, such as d-spacing d(10) and lattice constant a0, were calculated as d(10) = λ/(2sin θ) and a0 = 2d(10)/√3 for a 2-D hexagonal lattice. FT-IR spectra were recorded on an FT-IR spectrometer (IR Affinity-1, Shimadzu Corp.) using the KBr pellet 2123


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Figure 1. (a-h) SEM images of the RF aerogels prepared with varied starting compositions. (i) Appearance of the RF aerogels (TEG60-Hy). technique, while13C solid-state dipolar decoupling/magic angle spinning (DD/MAS) nuclear magnetic resonance (NMR) measurements were performed with an NMR spectrometer Advance III 800 (Bruker Corp.) under a static magnetic field of 18.8 T. Thermogravimety (TG) and differential thermal analysis (DTA) were carried out by Thermo plus TG 8120 (Rigaku Corp.) at a heating rate of 5 °C min−1 with continuously supplying air or argon at a rate of 100 mL min−1. Mechanical characteristics of the specimens were measured by a material tester (EZGraph, Shimadzu Corp., Japan) under a quasi-static condition at a crosshead speed of 0.5 mm min−1 unless otherwise stated. Contact angles were measured with Drop Master DM-561Hi (Kyowa Interface Science Co., Ltd.) with fixing the volume of water droplet as 5 μL. The detailed experimental procedures are described in the Supporting Information.

1 M HCl, the relatively coarse macroporous morphology (TEG60-H8) becomes finer associated with the overall structure turning into the assembly of thin cylinders interconnected with each other (TEG60-H12 and TEG60H16). The further increase of 1 M HCl coarsens the macroporous structure again accompanied by thickening the macropore frameworks (TEG60-H20), ending up with a stringof-beads-like morphology consisting of a few micron-size globules fused mutually (TEG60-H28). In the systems with different amounts of TEG (TEG30-Hy and TEG90-Hy), the similar trends were observed in the sequential morphological change with respect to the amount of 1 M HCl, as depicted in Figures S1 and S2. The observed geometrical transition on the length scale from submicrometers to micrometers reflects the variation of the gelation timing relative to the spinodal phase separation between comparatively hydrophobic polymeric phase and polar solvent phase; transient structures during the phase separation are fixed as a microscopic gel morphology by the sol−gel transition.42 An increase of 1 M HCl, which acts as an acid catalyst, facilitates the polymerization of RF networks, thereby shortening the gelation time, yet simultaneously stimulates the phase separation tendency by enhancing the polarity of solvent. The former leads to the sol−gel transition at the earlier stage of phase separation, whereas the latter sets off the opposite trend. Hence, when the amount of 1 M HCl

3. RESULTS AND DISCUSSION 3.1. Synthesis and Characterization of RF Aerogels. The synthesis strategy to fabricate porous RF gels in this study is based on the sol−gel method accompanied by phase separation.42,43 Various microstructures can be tailored by changing the synthesis parameters such as starting composition. For simplicity, the system described herein has only two valuables, i.e. the amounts of TEG (x) and 1 M HCl (y). Morphological evolution of the RF aerogels with varied amounts of 1 M HCl for the TEG60-Hy system is demonstrated in Figure 1 (a)-(f). In increasing the amount of 2124


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would readily disturb the periodical arrangement. Consequently, the ordered state is more disintegrated as the time lag between the onset of phase separation and the subsequent sol− gel transition is prolonged, rendering the coarser macroporous structure owing to the longer duration allowed for coarsening. The macroporous geometries consisting of columnar frameworks reminiscent of TEG60-H12 and TEG60-H16 can be tailored with a varied amount of TEG solely by adjusting the amount of 1 M HCl, as shown in Figure 1 (g) and (h). The increase of TEG leads to the lower concentration of polymer phase, thereby resulting in the lower density of the aerogels and vice versa. The finer morphology comprising shorter nanorods was observed in TEG30-H5, while the connecting struts became thicker with a lotus root-like appearance in TEG90H25. The SAXS profiles in Figure 2 (b) testify the formation of well-defined 2-D hexagonal symmetry, and the nitrogen sorption isotherms (Figure 3 (b)) showing a H1-type hysteresis loop of type IV isotherm classified by IUPAC also support the presence of cylindrical mesopores with a narrow pore size distribution in both specimens. The FT-IR measurement (Figure S3 (a)) disclosed that, whether the aerogels involved ordered mesoporosity or not, a lot of F127 still persisted in the dried specimens even after careful washing and subsequent supercritical drying (supercritical fluids generally have high extraction ability), because the incorporated F127 strongly interacts with the RF polymer backbone via hydrogen bonding.45,46 The TG curve for the aerogel (Figure S3 (b)) consists of two stages of weight loss, in which the first drop at 230−320 °C is mainly ascribed to the pyrolysis of F127. Although the decomposition temperature of F127 is lower than 180 °C in air, the remaining F127 in the RF scaffolds appears to be stable even at >200 °C, which supports the strong interaction between F127 and RF. As it is difficult to determine the amount of remaining F127 precisely because the RF gel is gradually pyrolyzed over the wide temperature range up to 500 °C in air, we roughly estimated the RF/F127 ratio as 1/1 by weight from the TG curve in Figure S3 (b). It is noteworthy that all the samples with different starting compositions (but fixed amounts of resorcinol, formaldehyde, and F127) gave almost the same TG curves, implying that the compositions (proportions of F127 to RF) of the dried samples were almost constant irrespective of their morphology. This result suggests that the identical RF/F127 oligomer units formed in the incipient stage of polymerization are organized to become a gel. The rest of F127 would be distributed in the fluid phase, which is rinsed off through the washing process, but the amount could not be very large considering the RF/F127 ratio in the starting composition. The detailed pore properties in the nanometer scale of the RF aerogels are compiled in Table 1. Comparing TEG60-Hy, it is found that the bulk density decreased as the amount of 1 M HCl was increased. First, this is because the higher amount of 1 M HCl, which is distributed to the solvent (fluid) phase, reduces the polymer concentration in the gel, thereby decreasing the bulk density. Besides, it should be noted that the samples with finer morphology resembling a truss structure (TEG60-H12 and TEG60-H16) shrank to a greater extent due to the spontaneous syneresis during gelation and aging, as shown in Figure S4, while the shrinkage during supercritical drying was almost the same (about 10%) for all the specimens. The skeletal density of the RF aerogel was measured as ca. 1.25 g cm−3 by helium pycnometry, resulting in the porosities of 90−92% for TEG60-Hy.

increases, the coarse phase-separated domains (TEG60-H8) first become finer with decreasing the macropore diameter from >5 μm to 2−3 μm (TEG60-H12 and TEG60-H16) and then resume coarsening with enlarging macropores and thickening macropore skeletons (TEG60-H20), as a result of the competition between the shortened gelation time and the accelerated phase separation. As reported previously,42,43 the columnar constituents in TEG60-H12 and TEG60-H16 incorporate highly ordered 2-D hexagonal arrays (space group of p6mm) of mesopores, as determined by SAXS measurements (Figure 2) and nitrogen

Figure 2. SAXS profiles of the RF aerogels: (a) RF aerogels with different amounts of 1 M HCl (TEG60-Hy) and (b) RF aerogels consisting of columnar macroframeworks with ordered mesoporosity. The peaks can be indexed to a 2-D hexagonal symmetry.

physisorption (Figure 3). The crystallographic orientation of straight cylindrical mesopores in a parallel manner gives rise to the anisotropic nanorods43,44 that are homogeneously distributed to form isotropic macroframeworks. As shown in Figure 2 (a), the periodic ordering was waned with coarsening macroporous morphology, resulting in the complete loss of mesopores in TEG60-H24. As losing the periodic architecture, the mesopore volume and surface area decrease, resulting in negligible open meso- and microporosity in TEG60-H24. In this sol−gel system, the ordered mesostructure arises from the supramolecular self-assemblies of RF/F127 associated with the auxiliary agents, TMB and BzOH.42−44 Since the lyotropic phase formed in a fine balance of composition is fairly sensitive to the change in an ambient medium, the phase separation 2125


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Figure 3. (a,b) Nitrogen adsorption−desorption isotherms and (c,d) mesopore size distributions obtained by the BJH method for the RF aerogels: (a,c) TEG60-Hy and (b,d) RF aerogels with ordered mesopores.

Table 1. Specific Surface Area Obtained by the BET Method

TEG30-H5 TEG60-H12 TEG60-H16 TEG60-H20 TEG60-H24 TEG90-H25

SBETa /m2 g−1

Vmicromesob /cm3 g−1

Dpc /nm

d(10)d /nm

a0d /nm

TWalle /nm

110 83 96 24 3.7 71

0.283 0.189 0.261 0.135 0.027 0.174

8.2 6.3 7.2 N.A. N.A. 6.3

16.6 14.6 16.5 N.A. N.A. 15.1

19.2 16.9 19.1 N.A. N.A. 17.4

10.9 10.5 11.9 N.A. N.A. 11.1

a Specific surface area obtained by the Brunauer−Emmett−Teller (BET) method. bMicro- and mesopore volume obtained by nitrogen adsorption isotherms at p/p0 = 0.99. cMean mesopore size calculated from the adsorption branches by the BJH method. dCalculated from SAXS measurements. eWall thickness calculated as TWall = a0 − Dp.

3.2. Mechanical Properties of RF Aerogels. The microstructural diversity of the RF aerogels presented in Figure 1 delivers a variety of mechanical properties. Typical stress− strain curves for the RF aerogels during uniaxial compression/ decompression with a maximum strain of 50% are exemplified in Figure 4, and the details are summarized in Table 2. All the samples displayed good flexibility and recoverability to >90% of the original sizes after 50% deformation. Note that additional restitution took place continuously with time after the stress was totally removed. The softer aerogels showed the superior resilience, while the slower recovery rate was observed for the stiffer ones. Each loading curve exhibits a linear elastic region at small strains (100 °C) for the RF gel. The weight loss ahead of the carbonization is attributable to the partial dehydration condensation of −OH groups in the RF networks. When the heat-treatment temperature increases, the loss of F127 and the decrease of hydrophilic functional groups on the RF network render the polymer surface more hydrophobic with increasing water contact angle beyond 145° (see Figure S13). The physical properties of the representative calcined samples are compiled in Table 4. Since the samples featured by the truss-like morphologies exhibited similar mechanical properties regardless of the drying procedure, we examined the effects of thermal treatment employing the xerogels, which would be favorable from a practical perspective. It was confirmed that the thermal treatment at 250−450 °C under inert atmosphere does not impair the tailored macroporous morphology (Figure S14) and that the change in bulk density is limited, as tabulated in Table 4. The mechanical behaviors against compressive stress for the soft aerogels (TEG60-H20 and TEG60-H24) in relation to the calcination temperature are shown in Figure S15. In both cases, an increase of the calcination temperature enhances the elastic modulus and improves the resilience after unloading. However, TEG60-H20-350 was always cracked at ca. 15% strain, while

maximum temporary volume shrinkage (usually observed at the end of the constant-rate period),47 when all the liquid in the pores evaporates and hence the capillary force is no longer loaded on the networks.10 As detailed in the preceding section, the obtained aerogels displayed good resilience from at least 25% strain, which indicates that they can be dried under ambient conditions with small shrinkage owing to the springback effect. Figure S10 shows the macroporous morphologies of the xerogels, which manifest that the microstructures of the specimens were not significantly disturbed by the stress during the solvent evaporation under an ambient condition. The shrinkage ratios through evaporative drying are listed in Table 3. It is found that TEG60-H20-X and TEG60-H24-X show larger irreversible shrinkage than that of the corresponding aerogels, whereas the shrinkage during evaporative drying was relatively small for TEG60-H12-X and TEG60-H16-X (see also Figure S4). Accordingly, the bulk densities of TEG60-H20-X and TEG60-H24-X are higher than those of TEG60-H12-X and TEG60-H16-X. This disparity can be attributed to the difference in elasticity arising from the different macroporous morphologies. It signifies that the gels with more elastic network are advantageous in terms of the alleviation of irreversible shrinkage during drying. Figure 6 depicts the effect of drying process on the mechanical properties. Because of the larger shrinkage during

Figure 6. Compressive stress−strain curves for the RF xerogels (TEG60-Hy). The curves for the correspondent aerogels are overlaid as broken lines.

drying resulting in the higher bulk densities, the xerogels became more robust than their supercritically dried counterparts. In particular, the Young’s modulus considerably increased by a factor of 2−3 for TEG60-H20-X and TEG60-H24-X due to the large difference in shrinkage ratio between a pair of the correspondent aerogel and xerogel as mentioned above. In these cases, readily deformable parts of the macroframeworks (e.g., branching nodes and interparticle necks) were deduced to be densified by the compressive stress derived from the 2128


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Chemistry of Materials Table 4. Various Properties of the Heat-Treated Samples TEG60-H20-250 TEG60-H20-350 TEG60-H24-250 TEG60-H24-350 TEG60-H24-450 TEG30-H5-X-350 TEG60-H12-X-350 TEG60-H16-X-350 TEG60-H16-X-450 TEG60-H16-X-280*g TEG60-H20-X-350 TEG60-H24-X-350 TEG90-H25-X-350

linear shrinkagea /%

weight lossb /%

ρbulkc /mg cm−3

porosityd /%

Young’s modulus /MPa

yield strengthe /MPa

7 21 8 20 28 15 14 14 24 20 14 12 14

4 42 4 42 57 40 41 43 58 48 40 43 42

110 114 104 105 105 239 132 117 116 115 130 134 82.6

91.3 91.0 91.7 91.7 91.9 81.2 89.6 90.3 91.1 91.1 89.8 89.4 93.5

2.7 4.6 0.76 1.1 2.6 102 28 26 31 28 8.8 1.9 5.1

0.058 0.090 0.016 0.044 0.062 2.1 0.57 0.51 0.73f 0.65 0.24 0.070 0.11

Linear shrinkage during heating calculated as 100 × (1−[calcined gel diameter]/[aerogel diameter]). bWeight loss during heating, most of which is due to the decomposition of F127. cBulk density calculated as [weight]/[bulk volume]. dCalculated as 100 × (1−[bulk density]/[skeletal density]). e Yield stress at 0.2% offset strain assessed from each stress−strain curve. fFractured stress was ca. 1.9 MPa at ∼9% strain. gCalcined at 280 °C for 2 h under 10% air/nitrogen atmosphere. a

Figure 7. (a,b) Stress−strain curves of the heat-treated samples with truss-like morphologies. (c) Sequential photographs of the uniaxial compression test on TEG30-H5-X-350 illustrating the high recoverability from 80% compression.

TEG60-H24-350 endured up to ∼60% compression yet was broken into small pieces by the further deformation (see also Figure S16). The heat treatment at 450 °C significantly compromised the flexibility for TEG60-H24. Figure S17 exhibits the stress−strain curves (50% compression) for the xerogels treated at 350 °C in comparison with those of the as-dried specimens. Both TEG60-H20-X-350 and TEG60-H24-X-350 were cracked before the strain reached 50%, indicating the heat treatment enhances the stiffness at the expense of compressibility. In the case of the samples with the truss-like morphologies, however, the calcination at 350 °C has a positive influence on the overall viscoelasticity, i.e. both mechanical strength and flexibility. Their Young’s moduli and yield strengths were increased by a factor of 1.5−2, and they were restored almost completely maintaining the structural integrity after 50% compression. Remarkably, the improved resilience was also observed when compressing to 80% strain, as shown in Figures 7 and S18 (see also Movie S1). Besides, it was also disclosed that TEG60-H16-X-350 can be compressed to 90% without serious damage (Figure S19). After the removal

of loading, the remaining deformation continued to be gradually relaxed with time. As such, TEG30-H5-X-350 left the permanent strain of ca. 7% after 80% compression (Figure 7 (c)) in stark contrast to that of more than 60% for TEG30-H5X (Figure S6). In this respect, TEG30-H5-X-350 combines considerably high elastic modulus of approximately 100 MPa and good resilience from large deformation, which is hitherto unreported for such lightweight materials. 3.5. Insights into Origin of the Unique Mechanical Properties. Most of the natural biomaterials renowned for their mechanical strength and toughness, such as bone, wood, nacre, and shells, comprise a combination of contrasting constitutive elements: hard and brittle components (frequently minerals) and soft and tough components (usually organic fibrils).36,48 Exquisite arrangements of the hard and soft domains in such biomaterials offer remarkable stiffness and toughness that surpass the properties of the discrete components, which has thus far inspired a myriad of artificial materials: for example, hybrid aerogels,3−7 polymer−clay nanocomposites,49 and double-network (DN) gels.50 Hereafter, 2129


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degradation of flexibility in TEG60-H16-X-280* was observed (Figure S20), indicating that the deteriorated flexibility is not caused by the increase of cross-linking density in the RF networks. In light of the above-mentioned results, it is obvious that the soft component of F127 is essential for imparting high fracture resistance to the RF/F127 hybrid polymer monoliths. In addition, there seems to exist an optimal RF/F127 ratio for the highest restoring force; TEG60-H16-X-350 obviously exhibited the superior recoverability to the as-dried specimen, and the same holds for the other specimens with similar microscopic geometries. Hence, it is assumed that excessive F127 adversely affect the resilient capability. The content of F127 in TEG60H16-X-350 was roughly estimated as 7−10 wt % from the TG results. It means that the optimum amount of F127 is likely to be small relative to RF, as is the case with most of the structural materials created by nature, which are hard/soft composites with a low proportion of soft element; for instance, nacre involves only 5 vol % of soft biomacromolecules.36,48 As discussed above, the incorporation of F127 in an appropriate manner and the highly ramified macroframeworks are the key for the unique mechanical properties of the RF/ F127 hybrid polymer monoliths. However, there still remains a question whether RF polymer matrix is also pivotal, in other words, whether the present material design strategy is valid for any kinds of hard components. Fortunately, by substituting silicon alkoxides for RF source, the same synthesis route based on the confluence of phase separation and supramolecular self-assembly can produce the SiO2/F127 and ethylene-bridged polysilsesquioxane/F127 composites bearing the truss-like microstructures akin to TEG60-H16, which are labeled as TMOS-F127 and BTMEF127, respectively (Figure S23).53 The columnar skeletons in both inorganic−organic hybrid gels also possessed cylindrical mesopores akin to the RF gels obtained in this study, as depicted in Figure S23 (c) (see also ref 51). Although the detailed pore characteristics such as pore size, volume, and porosity were different as listed in Table S1, the ordered mesoporous structures in these samples were undoubtedly generated by the cooperative self-assembly of polymeric species and a surfactant. Hence, it can be postulated that F127 molecules are integrated in silica or in polysilsesquioxane matrix in a manner analogous to that in the RF/ F127 counterparts. According to the TG measurement, TMOSF127 and BTME-F127 contained ca. 18 wt % and ca. 8 wt % of F127, respectively. In spite of the higher or comparable F127 content relative to that in the RF/F127 composites treated at 350 °C, both inorganic−organic hybrid gels were found to be rather fragile and disrupted by a compressive strain less than 5%, as shown in Figure S23 (d). We therefore conclude that the nature of the hard component is also an important factor governing the mechanical properties, especially the flexibility of composite materials. It appears that, if the truss structure embraces considerably cross-linked hard matrix as a major component, the incorporation of F127 hardly endows a flexible feature. In summary, the unique mechanical features consolidating high stiffness and good flexibility are traced to (i) the truss-like architecture built up of columnar units interconnected with each other, (ii) the integration of the soft polymer chains (F127) into the hard RF networks in a proper fashion in terms of disposition and proportion of soft and hard components (the details were not thoroughly elucidated in this work), and (iii) a

we discuss the function of hard (RF) and soft (F127) segments in the composite gels on their mechanical features. In an effort to unravel the contribution of F127 to the flexible trait, the mechanical properties of TEG60-H16-X-T were compared with respect to the content of F127. The stress− strain relationship as a function of calcination temperature for TEG60-H16-X is displayed in Figure S20. It seems evident that, as opposed to the thermal treatment at 350 °C which effectively enhances the compressibility, stiffness, and resilience, the treatment at 450 °C is detrimental to the fracture resistance; TEG60-H16-X-450 shattered into fragments prior to plastic deformation. Although the decomposition of F127 occurs above ∼300 °C under inert atmosphere as shown earlier in Figure S12, a part of F127 still remained in TEG60-H16-X-350, whereas almost all the F127 were excluded by the calcination at 450 °C, as evidenced by the 13C NMR spectra51,52 (Figure 8) and TG curves and FT-IR spectra (Figure S21) for TEG60-H16-X-T. It follows that the fragility of TEG60-H16-X-450 is predominantly attributed to the absence of F127.

Figure 8. 13C solid-state NMR spectra for the samples treated at different temperatures (TEG60-H16-X-T). The peak assignment is based on refs 49 and 50. The sharp strong peaks at 18 ppm (marked by open arrows) and at 70−76 ppm (marked by solid arrows) are assigned to the methyl carbons and the carbons of ethylene oxide chains in F127, respectively. The small band at 50−80 ppm in TEG60H16-X-450 and TEG60-H16-X-280* is probably attributed to the spinning sideband.

As the heat-treatment at elevated temperatures somewhat increases the cross-linking density of RF networks through dehydration condensation (see Figure S22), which could also impair the flexibility of polymer scaffolds, we also inspected an F127-free sample (TEG60-H16-X-280*) obtained by the calcination at 280 °C under slightly oxidative atmosphere with less thermal effects on the RF network. The thorough removal of F127 in TEG60-H16-X-280* and its similar crosslinking density to TEG60-H16-X-350 are proved by the 13C NMR, TG, and FT-IR measurements (Figures 8, S21, and S22). As for the mechanical property, the substantial 2130


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Table 5. Energy Absorption Capabilities of the Flexible Heat-Treated RF Gels for the First Loading−Unloading with a Maximum Strain of 50% or 80%

TEG60-H20-250 TEG60-H24-350 TEG30-H5-X-350 TEG60-H12-X-350 TEG60-H16-X-350 TEG90-H25-X-350 TEG30-H5-X-350 TEG60-H16-X-350 TEG90-H25-X-350

maximum strain /%

energy absorptiona /J g−1

energy absorptiona /J cm−3

energy dissipationb /J g−1

energy dissipationb /J cm−3

energy loss coeffc

50 50 50 50 50 50 80 80 80

0.67 0.63 8.9 4.5 4.4 1.7 31 12 5.7

0.073 0.066 2.1 0.60 0.51 0.14 7.5 1.4 0.47

0.52 0.44 6.6 3.0 2.7 0.74 24 9.2 4.5

0.057 0.046 1.6 0.39 0.31 0.099 5.7 1.1 0.37

0.77 0.70 0.74 0.65 0.61 0.68 0.76 0.75 0.80

a Total work done by compression obtained from the integrated area under the stress−strain curves during loading stress. bObtained from the integrated area in the hysteresis loop of each stress−strain curve. cCalculated as [energy dissipation]/[energy absorption].

Figure 9. (a,b) Cyclic stress−strain curves of TEG60-H16-X-350 under a quasi-static condition (0.5 mm min−1) at a maximum strain of (a) 50% and (b) 80%. (c) Effect of the curing process (wetting and drying) on the fatigued sample (TEG60-H16-X-350) after the successive compression to 80% for 10 cycles. (d) Mechanical response of the heat-treated samples to a dynamic compression (1000 mm min−1). (e) Energy absorption capabilities of the heat-treated samples during the cycling compression to 70% with a fast straining rate of 0.5 s−1. After the 10th and 20th cycles, the worn gels were cured by soaking in 2-propanol followed by vacuum drying at 60 °C.

microlattices (0.33 J g−1 and 0.0046 J cm−3 at 50%),34 boehmite nanofiber (BNF)/PMSQ gels (1.4−2.0 J g−1 and 0.22−0.29 J cm−3 at 50%),22 graphene aerogels (1 J g−1 and 0.005 J cm−3 at 80%),27 marshmallow-like gels (0.2 J g−1 and 0.02 J cm−3 at 80%),12 and PMSQ aerogels (7.9−9.1 J g−1 and 1.7−1.9 J cm−3 at 83%).10 The stress−strain behavior of TEG60-H16-X-350 during the cyclic loading/unloading test is represented in Figure 9 (a) and (b). As with most of the flexible elastomers like rubbers and polyurethanes, TEG60-H16-X-350 showed a typical softening behavior.28 In the second cycle, the monolith became more compliant than that in the first cycle, and the stress−strain curves in the following consecutive compression cycles were prone to converge to a softened mechanical profile. Meanwhile, a subtle change in the stress level at the maximum strain was found in the successive cycles, indicating little deterioration in mechanical stiffness. We also observed the so-called Mullins effect54 upon reloading with the higher maximum strain (Figure S24); the stress−strain curve moves along the softened curve

moderate cross-linking density of the three-dimensional (3-D) RF networks. 3.6. Energy Absorption Capabilities of Thermally Modified Polymer Monoliths. Flexible materials characteristic of low density and high mechanical strength are potent energy absorbers that can be used repetitively.35 In this sense, the high stiffness in conjunction with the large hysteresis loops in the stress−strain curves observed in Figure 7 (a) and (b) corroborates a great potential of the heat-treated RF gels in the art. The energy absorption capabilities of the selected samples for the first loading/unloading cycle are tabulated in Table 5. In compressing to 80%, they showed an energy loss coefficient, which is the ratio of dissipated energy over total work done by compression, of 0.75−0.80, reflecting splendid energy dissipation ability. It is notable that the stiffest sample, TEG30H5-X-350, delivered fairly high specific energy absorption of 8.9 J g−1 (2.1 J cm−3) and 31 J g−1 (7.5 J cm−3) during the compression to 50% and 80%, respectively, in contrast to the values for flexible low-density materials reported previously: Ni 2131


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isotropic cocontinuous macroframeworks, which are identified as highly branched truss-like structures. The polymer scaffolds featured by such robust microstructures acquire high compressive elasticity as well as good flexibility, which allow the gels to be dried under ambient condition with mitigated shrinkage compared to the samples with the other morphologies. In addition, the gels with similar truss-like morphologies yet varied bulk densities can be fabricated simply by altering the amounts of TEG and 1 M HCl in the starting composition. It follows that the mechanical strength of the gels can be controlled to some extent, while the elastic modulus is still dictated by the bulk density according to the power law (Figure S27). It is put forward that, like natural structural materials, the hybridization of rigid 3-D cross-linked RF networks and linear polymer chains of F127 overcomes the inherent fragility of RF aerogels and gives rise to the unique flexible features. The detailed investigation on the correlation between mechanical properties and chemical composition has unveiled that the simple thermal treatment under inert atmosphere can further enhance both mechanical strength and restorable compressibility of the composite gels with the truss-like morphologies. The adequate reduction of the [soft component]/[hard component] proportion is found to be key for improving the mechanical traits especially regarding the recoverability from a largely compressed state. Modulation of the amount of F127 through the thermal post-treatment evoked the unprecedented polymeric materials having the combination of good compressibility, high stiffness, and noticeable recoverability. By virtue of such unique mechanical properties, the heat-treated specimens delivered remarkable energy absorption and dissipation performances repetitively, which paves the way for a new category of flexible materials. The monolithic gels fatigued with cyclic compression to a large strain can be facilely restored to their original shape through wetting and drying. Further efforts to develop flexible materials affording quick recovery and high fatigue resistance are currently ongoing.

until the strain reaches the maximum strain in the previous cycles followed by overlapping with the curve in the first cycle. In the serial compression cycles, the residual strain becomes rather substantial due to the accumulated strain imposed by a relatively slow recovery rate, albeit the deformed sample gently resumes with time, regaining the elastic stiffness, as shown in Figure S25. Specifically, when it comes to the cyclic 80% compression, the unrecoverable strain seems to be an acute issue from an application viewpoint. Concerning TEG30-H5-X350 and TEG90-H25-X-350, the accumulated strain reached around 20−30% after 10 cycles of 80% compression (see Figure S26). For the purpose of addressing this drawback, we propose a facile and feasible method to restore a sample strained by iterative compression to its original shape, that is, wetting and drying. By the simple treatment in which the sample is immersed in alcohols like ethanol and 2-propanol followed by drying, the fatigued gel can revert almost completely with retrieving elasticity to some extent, as demonstrated in Figure 9 (c). On wetting, a swollen gel becomes softer, which allows the deformed skeletons to be relaxed. As discussed above, the highly elastic and flexible truss-like polymer framework affords the stress arising from capillary pressure both during wetting and drying, enabling the stressed gel skeletons to achieve almost full recovery in shape. A part of local damage in the macroframeworks, however, cannot be cured by this treatment, thereby resulting in more or less compromised mechanical properties (elastic modulus and yield stress) compared to those of the corresponding fresh specimens. Finally, we demonstrate the energy absorption and dissipation capabilities of the heat-treated samples under a dynamic condition, which is closer to that in the practical uses of cushioning and shock damping. The compressive mechanical profiles at a strain rate of 1000 mm min−1 (corresponding to ca. 1.0−1.3 s−1), which is the fastest limit of the instrument, are shown in Figure 9 (d). All the samples with truss-like structures can be deformed and recovered as following the high-speed compression platen without any cracks (see Movie S2), implying their high durability against axial impact. Remarkably, in spite of the high stiffness that is generally inextricably linked to brittleness,36 TEG30-H5-X-350 also withstood the highspeed compression to 80%, as shown in Figure 9 (d). The cyclic energy absorption performances of the heattreated samples with the truss-like morphologies are summarized in Figure 9 (e). Under the dynamic condition (0.5 s−1), each specimen delivered the energy absorption, energy dissipation, and energy loss coefficient comparable to those under quasi-static conditions in the first compression cycle. In the subsequent cycles, these values gradually decreased due to the softening behavior28 with the energy loss coefficient hovering at a relatively high value of 0.5−0.6. It is also found that the straightforward curing process did work out well.

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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/acs.chemmater.6b04706. Detailed experimental procedure, SEM images, shrinkage during aging and drying, compressive and flexural mechanical properties, FT-IR spectra, TG-DTA curves, water contact angles, 13C solid-state NMR spectra, N2 sorption isotherms for the RF samples and polysilsesquioxane monoliths (PDF) Movie S1 of the mechanical tests (AVI) Movie S2 of the mechanical tests (AVI)

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4. CONCLUSIONS The morphological evolution through the course of phase separation in the RF sol−gel system assisted by F127 provides flexible hybrid polymer gels with diverse porous architectures, demonstrating a variety of mechanical characteristics depending on their morphologies. Within the specific starting compositions where several requirements (such as micelle formation through the sol−gel transition) are fulfilled, the columnar building blocks with 2-D hexagonally ordered mesopores generated by the supramolecular self-assembly constitute

AUTHOR INFORMATION

Corresponding Author

*Phone/Fax: +81 92 802 2862. E-mail: [email protected]. ORCID

George Hasegawa: 0000-0003-4546-5197 Kazuyoshi Kanamori: 0000-0001-5087-9808 Notes

The authors declare no competing financial interest. 2132


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ACKNOWLEDGMENTS The authors express their sincere gratitude to Dr. G. Hayase for providing data related to the BNF/MSQ and PDMS/MSQ gels. The present study has been performed with financial support from the Advanced Low Carbon Technology Research and Development Program (ALCA, JST Japan). Financial support by Grant-in-Aid for Scientific Research (No. 16K05935 for G.H. and No. JP26288106 for K.N.) from the Japan Society for the Promotion of Science (JSPS) is also gratefully acknowledged.

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Highly Flexible Hybrid Polymer Aerogels and Xerogels Based on

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