Molecularly Engineered Intrinsically Healable and Stretchable

Article Cite This: Chem. Mater. 2017, 29, 8850-8858

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Molecularly Engineered Intrinsically Healable and Stretchable Conducting Polymers Paul Baek,†,‡ Nihan Aydemir,†,‡ Yiran An,†,‡ Eddie Wai Chi Chan,†,‡ Anna Sokolova,§ Andrew Nelson,§ Jitendra P. Mata,§ Duncan McGillivray,†,‡ David Barker,† and Jadranka Travas-Sejdic*,†,‡ †

Polymer Electronics Research Centre, School of Chemical Sciences, University of Auckland, Auckland, New Zealand The MacDiarmid Institute of Advanced Materials and Nanotechnology, Wellington, New Zealand § Australian Centre for Neutron Scattering, Australian Nuclear Science and Technology Organization, Kirrawee, NSW 2234, Australia ‡

S Supporting Information *

ABSTRACT: Advances in stretchable electronics concern engineering of materials with strain-accommodating architectures and fabrication of nanocomposites by embedding a conductive component into an elastomer. The development of organic conductors that can intrinsically stretch and repair themselves after mechanical damage is only in the early stages yet opens unprecedented opportunities for stretchable electronics. Such functional materials would allow extended lifetimes of electronics as well as simpler processing methods for fabricating stretchable electronics. Herein, we present a unique molecular approach to intrinsically stretchable and healable conjugated polymers. The simple yet versatile synthetic procedure enables one to fine-tune the electrical and mechanical properties without disrupting the electronic properties of the conjugated polymer. The designed material is comprised of a hydrogen-bonding graft copolymer with a conjugated backbone. The morphological changes, which are affected by the composition of functional side chains, and the solvent quality of the casting solution play a crucial role in the synthesis of highly stretchable and room-temperature healable conductive electronic materials.

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cross-linking of the conjugated segments.15 Other dynamic interactions that afford healing exist,8,20 but hydrogen bonds are desirable for spontaneous interaction, which enables materials to heal at room temperature.21,22 Nonetheless, it is imminent that the disruption of conjugation in the presence of nonconjugated segments will affect the electrical properties of the CP.14,23 Thus, advancement in the molecular design of intrinsically stretchable and healable CPs will require a careful selection of hydrogen-bonding components and unique approaches to incorporate such moieties without disturbing the inherent electrical properties of the CP. Herein, we present a novel multiphase design for highly stretchable, room-temperature healable, and electrically conductive CPs based on a graft copolymer concept. The key feature of this design is the simple but versatile synthetic procedure that enables one to easily tailor the composition of both the CP backbone and the polymeric side chains that possess dynamic healing motifs through hydrogen bonding. Our concept is analogous to the self-healing thermoplastic elastomer design22 that incorporates both stretchability and healing, properties afforded by a graft copolymer structure comprised of “soft” amorphous and “hard” crystalline phases.

INTRODUCTION Stretchable and healable electronic components and devices are new frontiers in organic electronics.1−4 Healing, which is intrinsic to biological systems, is no longer limited to biomaterials5,6 as it has been well-researched in applications of polymers7−10 and skin-inspired electronics.1,11 Current approaches to stretchable electronics entail engineering of strain-compliant conductive or semiconductive materials onto an elastomer or fabricating hybrid materials comprised of noncompliant conductive materials in a stretchable matrix.4,12,13 The realization of innovative design and synthesis of intrinsically stretchable and healable materials presents an ideal strategy over the currently established approaches,13 which enables the key mechanical and electronic properties required for stretchable electronics to be incorporated into a single molecular entity. To date, there have been only a few reports on intrinsically stretchable and healable materials based on conjugated polymers (CPs).14−17 For instance, the mechanical stretchability of CPs can be enhanced through incorporation of flexible polymer segments within the conjugated backbone.13,16,18,19 Similarly, the only design concept for stretchable as well as healable conjugated polymers that is currently presented involves insertion of molecules within the conjugated backbone that enhances the molecular stretchability and reversible healing through dynamic hydrogen bonding-induced © 2017 American Chemical Society

Received: August 3, 2017 Revised: October 1, 2017 Published: October 3, 2017 8850

DOI: 10.1021/acs.chemmater.7b03291 Chem. Mater. 2017, 29, 8850−8858

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Chemistry of Materials

Figure 1. (a) Synthesis of the acrylate urethane monomer. The inset shows the 1H NMR spectra of acrylate urethane at various concentrations in CDCl3 (focused on the N−H peak), which indicate the presence of hydrogen bonding. (b) Synthetic scheme of three series of graft copolymers, PTh-g-PAU, from the macroinitiator PMI via ATRP of acrylate urethane. ATRP was performed in tetrahydrofuran using Cu(I)Br and PMDETA. (c) General mechanism for stretchability and healing afforded by the noncovalent cross-linking among PAU side chains.

Figure 2. (a) 1H NMR spectra of the PMI and graft copolymers in CDCl3, which show the gradual increase in the intensity of peaks associated with the PAU in comparison to that of the P3HT peaks with increasing graft lengths. (b) Increase in solution QY as a function of PAU DP. (c) Parameters of the flexible cylinder model. (d) Comparison of persistence length lp (left axis) and radius r (right axis) of the polymers, which were obtained from fitting the SANS data with the flexible cylinder model. The graft copolymers with a side chain DP of >5 have a lower rigidity (indicated by a lower lp), which implies a change in the overall behavior of the graft copolymer with respect to graft length, where the grafted PAUs begin to dominate the physical properties of the graft copolymer.

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Chemistry of Materials Rather than a flexible polymer backbone as described by Chen et al.,22 we demonstrated our concept by utilizing a CP as the backbone. Importantly, we reasoned that the electrical and mechanical properties of the CP graft copolymers can be altered in a desired fashion using graft lengths or a solvent mixture, from which the polymers are cast. First, both theoretical and experimental studies of conventional graft copolymers have emphasized the importance on the effect of graft lengths on their structural properties.24−26 For instance, persistence length, which reflects the rigidity of the polymer chain,27,28 increases with side chain length.24 It is well established in the field of CPs that understanding the conformation in solution is crucial for optimizing the electronic and mechanical properties of CP films.27−32 Hence, small angle neutron scattering33−35 is employed here for the first time to probe the structure of CP graft copolymers. Lastly, CP graft copolymers processed from a solvent mixture, consisting of a poor solvent for the CP backbone but a good solvent for the polymeric side chains, exhibited high stretchability up to 380% and room-temperature healing. The resulting films are highly amorphous (glass transition temperature below room temperature) yet contain ordered domains of collapsed CPs that can be attributed to a significant improvement in their electrical properties.36 Such properties in addition to the solution processability of the graft copolymers were exemplified by fabricating a simple stretchable device.

close to the feed value of 20%, which confirms the copolymerization efficiency between two functional thiophene monomers when the monomers are incorporated at similar reaction rates. A number-average molecular weight of 41 kDa with a dispersity of 5.9 was obtained, which is typical for the oxidative polymerization with FeCl3.41 From the PMI, a series of three graft copolymers were successfully synthesized by ATRP, as confirmed by the first-order kinetics of the polymerization reaction37 (Figure S3), 1H NMR (Figure 2a), Fourier transform infrared spectroscopy (FTIR) (Figure S4), and gel permeation chromatography (GPC) (Figure S5). The DPs of the grafted PAU obtained from the kinetic study are consistent with the DPs obtained from 1H NMR (Table S2). To further elaborate on the structural characterization of the graft copolymers, we estimated the PTh content in the graft copolymers from the ultraviolet−visible (UV−vis) measurements in tetrahydrofuran (THF) at various concentrations42 (Figure S7). The results show that PTh-g-PAU5, -17, and -48 consisted of 33, 26, and 20 wt % PTh, respectively. The effect of graft length on the optoelectronic and physical properties of the copolymers was initially evaluated by measuring the quantum yields of the graft copolymers in a THF solution (Figure S8 and Table S5). Interestingly, all graft copolymers showed a broad absorption maximum at approximately 436 nm, which corresponds to the π−π* transition43 of the P3HT backbone and is slightly red-shifted compared to that of the PMI (λmax = 432 nm). This spectroscopic shift initially suggests that there is sufficient spacing between the grafts because a densely grafted copolymer typically shows a blueshift, which results from the twisting of the backbone by the steric hindrance of the side chain polymers.44 Even in dilute solutions, P3HT is prone to aggregation.34 This behavior was evident in the solution quantum yields (QYs) for PMI, which showed a notably low QY of 2.0% compared to the literature value of 33% for pristine P3HT (in chlorobenzene).45 However, the graft copolymers showed increased QY values with increasing graft lengths, and PTh-g-PAU48 had a QY of 8.5% (Figure 2b), which we attribute to the weakened interchain interactions within the conjugated backbone that causes the photoluminescence quenching.31 To investigate the conformation of the PMI and dissolved graft copolymers in THF, small angle neutron scattering (SANS) was measured with the dilute solutions (0.5 wt %). We report, for the first time, the use of the scattering techniques to probe the structure of CP-based graft copolymers in solution. In THF, PMI and PTh-g-PAU5 showed a Porod scaling of q−1.5, whereas PTh-g-PAU17 and PTh-g-PAU48 scaled with q−1.7 (Figure S9). The Porod scaling between the values of 1.5 and 1.7 indicates that THF is a good solvent for all the polymers and that they are well-solvated in THF.46 To analyze the effect of the graft lengths on the intrinsic flexibility (reflected in the persistence length, lp) of the PTh backbone, a flexible cylinder model (Figure 2c) for CPs and graft copolymers24,34 was fitted to the mid-q region (0.005 Å−1 < q < 0.2 Å−1). The lp of 63.9 ± 2.6 Å for unfunctionalized PMI appeared to be slightly larger than the reported values (20−30 Å) in the literature.34,47 The higher lp of PMI compared to that of pristine P3HT can be attributed to the presence of large ATRP initiator groups.28 There is a general increase in lp up to a PAU DP of 5 (Figure 2d). Assuming that lp reflects the rigidity of the polymer, we can attribute the increase in lp to the increased level of repulsion among the PAU side chains with an increasing DP48 and the better solubility of the CP backbone, which is aided by the

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RESULTS AND DISCUSSION Molecular Design of Conducting Polymer Graft Copolymers. Our polymer design consists of grafting soft poly(acrylate urethane) (PAU) from a conjugated poly(3hexylthiophene) (P3HT) backbone via atom transfer radical polymerization (ATRP)37 (Figure 1). The P3HT backbone is a random copolymer of nonfunctionalized 3-hexylthiophene (3HT) and ATRP initiator-bearing thiophene [2-(thiophen-3yl)ethyl 2-bromo-2-methylpropanoate]. Previously, there have been discussions in the literature about the compromise between the grafting density of insulating polymer side chains and the electrical properties of CP-based graft copolymers.38 Thus, we selected the P3HT copolymer macroinitiator (termed PMI), which would afford a low grafting density (17% ATRP functionality) for the PAU side chains, in our efforts to preserve the inherent electrical properties of P3HT after grafting. PAU was selected as the functional side chain because of its healing properties, which originate from its low glass transition temperature [Tg = 9.5 °C (Figure S12)] and dynamic hydrogen bonding between the urethane groups.39 The presence of H bonding was confirmed using 1H nuclear magnetic resonance (NMR) experiments (Figure S1). We synthesized a series of three PTh-graf t-PAU copolymers with different graft lengths under identical ATRP conditions, where the lengths of the PAU grafts were controlled by the ATRP time. Here, PTh denotes the copolymer of 3HT and ATRP-bearing thiophene. For the sake of simplicity, the graft copolymers with PAU degrees of polymerization (DP) of 5, 17, and 48 are here termed PTh-g-PAU5, PTh-g-PAU17, and PThg-PAU48, respectively. Characterization of the Material. The chemical structures of the macroinitiator PMI and graft copolymers were extensively characterized. The 1H NMR spectra in Figure 2a indicate that the PMI is a regiorandom copolymer40 (head:tail ratio of 78%) with a grafting density of 17% (see the Supporting Information for further discussion). This value is 8852


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Figure 3. (a) UV−vis−NIR absorption spectra of thin films of PMI and PTh-g-PAU48 doped in the 0.05 M FeCl3/acetonitrile solution, which had a higher degree of doping for graft copolymers than for unfunctionalized CP. Because of the difference in thickness, the spectra of PTh-g-PAU48 were normalized against PMI for visual comparison. (b) DSC curves of PTh-g-PAU48 as a function of the solvent quality (first heating scans; heating rate of 10 °C min−1). The first heating scans are shown here to evaluate the pristine structures of the as cast polymer films as a function of solvent quality. The films cast from a mixed solvent have Tg values and melting points lower than the values of those cast from THF, which suggests the presence of an aggregated P3HT backbone. See the Supporting Information for further discussion. (c) Conductivity of the drop cast polymer films (∼5 μm) on a glass substrate, where a higher conductivity is achieved for films cast from a mixed solvent casting solution. (d) Young’s modulus and strain at break of the free-standing PMI and graft copolymer parallel specimens (crosshead speed of 5 mm min−1). The graft copolymer films cast from the mixed solvent casting solution have are exceptionally stretchable.

absorbance ratio between the polaron band (typically λmax ∼ 800 nm for P3HT52) and the bipolaron band (typical λmax ∼ 1500 nm52) for PTh-g-PAU48 than for the doped PMI. This result suggests a greater and more efficient conversion of polarons to bipolarons for PTh-g-PAU48. However, the thin films of the doped graft copolymers showed a significant decrease in electrical conductivity with an increase in graft length (Figure 3c). As expected, the insulating PAU chains disrupted the charge transport in the films, although the copolymer was fully doped. We hypothesize that the electrical performance of the graft copolymers can be controlled using the quality of the solvent from which the polymers are cast. For example, we anticipate that via the introduction of methanol, which is a nonsolvent for P3HT but a good solvent for PAU, the formation of the aggregate of the graft copolymers in the solution will improve the electrical performance of the cast films. This anticipation is based on an assumption that the crystalline PTh domains will be preserved in the films and facilitate better charge transport in the graft copolymer films.36 First, we performed UV−vis measurements to follow the formation of aggregated segments of the PTh backbone in the solvent mixtures with a 0−50% volume fraction of methanol (Figure S11). We found distinctive absorption peaks at approximately 550 and 600 nm, which typically correspond

grafted side chains.49 Additionally, the contour length of the polymer also increased with side chain length. This result is consistent with the observed red-shift in the UV−vis absorbance and the increase in solution QYs with increasing graft lengths. Thus, we suggest that the polymeric side chains isolate the PTh segments and render the graft copolymer in an extended conformation.44 Interestingly, the graft copolymers with longer grafts (PAU DP > 17) appear to be less rigid (lower lp). This result implies a change in the overall behavior of the graft copolymer with respect to the graft length, where the grafted PAUs begin to dominate the physical properties of the graft copolymer, thereby changing the conformation of the graft copolymer (PTh-g-PAU17 has ∼74 wt % PAU, whereas PTh-gPAU48 has ∼80 wt % PAU).24 Thus, the overall lp of graft copolymers with long PAU side chains (DP > 17) is estimated as the average lp of the grafted side chains and CP backbone.50 After characterizing the graft copolymers in solution, we investigated the effect of graft length on the electrical properties of the graft copolymer thin films. Despite the common belief that the doping of CPs would be hindered by the presence of insulating grafted polymers,38,51 all graft copolymer thin films showed successful doping, which is indicated by the formation of new charge carrier bands (polarons and bipolarons)52 in the visible−near-infrared (vis−NIR) region (λ = 500−1700 nm) upon doping (Figure S10). In fact, Figure 3a shows a greater 8853


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Figure 4. (a) Schematic diagram of the healing experiment with (undoped) PTh-g-PAU48. To induce healing, free-standing polymer strips were superimposed and pressed (2.6 kPa) between two silanized glass substrates for various amounts of time. (b) The healing of the PTh-g-PAU48 cast from THF or a mixed solvent (30 vol % methanol in THF) at room temperature or 60 °C was qualitatively analyzed on the basis of the SEM observation of the junction after various healing times. Healing is significantly improved at both room temperature and 60 °C for the films cast from the mixed solvent casting solution.

to the aggregation-induced interchain π−π stacking of P3HT,36 for all graft copolymer solutions that contained 30 vol % methanol. The differential scanning calorimetry (DSC) curves (Figure 3b) showed a melting point (133 °C) for PTh-gPAU48 cast from the 30 vol % methanol/THF mixture, which indicates the presence of a crystalline aggregate structure of PTh because of the addition of methanol. Thus, we propose that 30 vol % methanol is sufficient to induce aggregates without the precipitation of the graft copolymer from solutions. In terms of conductivity, the solvent mixture significantly affects the conductivity of the PTh-g-PAU48 thin films (Figure 3c): the conductivity of the film cast from the 30 vol % methanol mixture was 10 ± 0.9 mS cm−1, whereas that of the film cast from THF alone was 1.6 ± 0.7 mS cm−1. A stronger effect of the solvent mixture was observed when the mechanical properties of the graft copolymers were considered. First, when PAU side chains were introduced, the structures of the graft copolymer films were completely changed, as indicated by the disappearance of the melting and recrystallization peaks of PMI at 170 and 89 °C, respectively, in the DSC curves (Figure S12). The atomic force microscopy (AFM) images confirm this observation, where the distinctive crystalline domains of PMI are completely absent for the graft copolymers (Figure S16). A well-defined Tg was observed below room temperature (21 °C) for all grafted copolymers, whose value was 7.6−15 °C with an increasing PAU DP of 5−48 (Tg was measured from the second heating scans in Figure S12). To evaluate the pristine structures of the as cast PTh-g-PAU48 films as an effect of the solvent mixture, the first heating scans were analyzed (Figure 3b). When cast from the 30 vol % methanol/THF solvent mixture, the graft copolymers showed a Tg of −3.9 °C that was significantly lower than the Tg of those cast from THF (13.6 °C). The significant decrease in Tg, which is an effect of the solvent mixture, suggests the increased entanglement53 of grafted PAUs in the graft copolymer films. The mechanical properties (Young modulus and maximum strain) of free-standing graft copolymer films were investigated using strain−stress measurements as a function of side chain

length and the quality of the casting solvent (Figure 3d). As expected, the unfunctionalized PMI showed a Young’s modulus that was higher than those of the graft copolymers as CPs are well-known for their stiffness compared to more elastic nonconjugated polymers such as PAU.33,34 As a function of side chain length, graft copolymers showed a lower modulus with a longer PAU, where a significant decrease in the modulus observed for a PAU DP of >5 is in strong agreement with the significant decrease in lp. Considering the significant weight fraction of PAU (80%) of PTh-g-PAU48, we assumed that the H-bonded cross-linking among the PTh chains would increase the modulus of the polymer film. However, the highly amorphous nature of the PAU matrix of PTh-g-PAU48 appeared to have a stronger effect on the modulus and reduced it.54 The graft copolymers showed greater strain at break than the unfunctionalized CPs did. We attribute this observation to the increased level of (noncovalent) cross-linking of the CP chains via H-bonding interactions between the grafted PAUs, as well as the increased level of local reinforcement4,55,56 of the entangled PAU matrix by the presence of stiffer CP, especially for the graft copolymers with sufficiently long side chains (PAU DP > 5). When cast from a 30 vol % methanol/THF solvent mixture, graft copolymers PTh-g-PAU17 and PTh-g-PAU48 showed maximum strain values greater than those of the films cast from only THF. The PTh-g-PAU48 had exceptional stretchability (≤380%) and a low elastic modulus (1.5 ± 0.1 MPa). The significant increase in stretchability and a lower modulus, as a result of the addition of methanol (a nonsolvent for the CP), can be explained by the change in the size and conformation of the CP backbone, which acts as the “filler” or reinforcing element within the PAU matrix. As demonstrated by Chen et al.,22 we assume that in the presence of a nonsolvent, the CP backbone of the graft copolymer partially collapses, producing a hard−soft core−shell type structure, different from an extended random coil conformation in THF alone. The significant improvement in the stretchability then can be attributed to the entanglement of flexible PAU side chains, as a result of the 8854


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Figure 5. (a) Optical images of the stretchable PMI- or PTh-g-PAU48-coated cotton gauze strain sensor. (b) SEM images of the sensors at different magnifications. (c) AFM phase images of (undoped) polymer thin films spin-coated onto mica. The time-dependent normalized resistance changes (ΔR/R0) of (d) PMI, (e) PTh-g-PAU48 (cast from the THF solution), and (f) PTh-g-PAU48 (cast from the 30 vol % methanol mixture)-coated gauze sensors in response to repetitive stretching from 0 to 100% strain.

The optical images in Figure 5a show the successful fabrication of the polymer-coated textile strain sensors (see the Experimental Section for details). The structure of the polymer coating on the cotton fibers was evaluated via SEM imaging (Figure 5b). The images of the pristine cotton fibers (Figure S15) and PMI- or PTh-g-PAU48-coated fibers were compared, which reveals a uniform coating with weblike structures between adjacent fiber strands. These structures are likely polymer films that formed during the rapid drying process in vacuum. However, the PMI-coated fibers showed a surface (clearly shown on its weblike films) that was much rougher than that of the PTh-g-PAU48-coated fibers. Because coated fibers were not suitable for AFM imaging, and we wished to gain further insight into the surface morphology of the polymers’ thin films, the films were cast onto mica and imaged using AFM (Figure 5C). The phase image of the PMI thin film exhibits a notably distinctive phase contrast between the PTh crystalline domains and the amorphous phase. The PTh-gPAU48 thin films that were spin-coated from either THF or a 30 vol % methanol/THF mixture showed phase images much more featureless than those of PMI. Hence, there is a more homogeneous film of the graft copolymer, which is physically cross-linked via hydrogen bonding. Strains of ≤100% were applied with repetitive stretching while measuring the time-dependent normalized resistance change (ΔR/R0) for the graft copolymer-coated textile strain sensors, as shown in Figure 5d−f. The strain sensor coated with PMI showed a slightly more sensitive ΔR/R0 response (gauge factor of 0.40 ± 0.03) with a baseline signal significantly more stable than the signal of the sensor coated with PTh-g-PAU48 cast from THF (gauge factor of 0.13 ± 0.02). However, when the strain sensor was fabricated with PTh-g-PAU48 cast from the mixed solvent [30:70 (v/v) methanol/THF], the baseline signal became notably stable, and the strain signal was reproducible (gauge factor of 0.19 ± 0.01). The PTh-g-

partially collapsed CP core. This description is also consistent with the known effect of filler sizes on the reinforcement of elastomeric nanocomposites.57 Considering the given hydrogen-bonding interactions of PAU segments (see the discussion of FTIR in the Supporting Information), the low modulus, and the Tg of the PTh-gPAU48 films cast from the solvent mixture, we anticipated that the films would have healing abilities at room temperature (21 °C). To qualitatively evaluate the healing process, scanning electron microscopy (SEM) was used to view the fusion of two completely separate, superimposed graft copolymer strips after a gentle pressure of 2.6 kPa11 was applied to enforce entanglement at the polymer−polymer interface (Figure 4a). Compared to the PTh-g-PAU48 films cast from THF, SEM revealed successful room-temperature healing of the films cast from the solvent mixture when a small pressure of 2.6 kPa was applied for 6 h (Figure 4b). The upper and lower polymer strips diffused well into each other at the junction. More efficient healing of films cast from the solvent mixture can be explained by unrestricted diffusion of the graft copolymer chains, as well as the greater level of exposure of the hydrogenbonding side chains to each other, when the graft copolymer assumes the conformation discussed above. Regardless of the effect of the solvent mixture, PTh-g-PAU48 films cast from THF or the solvent mixture showed complete diffusion at the films’ junction when exposed to an elevated temperature of 60 °C for 1 h because of the greater mobility of both PTh and PAU segments at 60 °C. Our current studies of self-healing grafted conjugated copolymers focus on a range of conjugated polymers and a more quantitative evaluation of such properties. Stretchable Textile Strain Sensor. To demonstrate the use of the stretchability and electrical properties of the graft copolymers, stretchable textile strain sensors were fabricated using PTh-g-PAU48 and PMI for comparison. 8855


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SANS experiment was carried out at Australian Nuclear Science and Technology Organization under the proposal #5601.

PAU48 cast from the mixed solvent is more electrically stable than those cast from THF, which we attribute to the methanolinduced change in morphology that better supports the stretchability of the films.

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(1) Chortos, A.; Liu, J.; Bao, Z. Pursuing prosthetic electronic skin. Nat. Mater. 2016, 15 (9), 937−950. (2) Chen, D.; Liang, J.; Pei, Q. Flexible and stretchable electrodes for next generation polymer electronics: a review. Sci. China: Chem. 2016, 59, 659−671. (3) Root, S. E.; Savagatrup, S.; Printz, A. D.; Rodriquez, D.; Lipomi, D. J. Mechanical Properties of Organic Semiconductors for Stretchable, Highly Flexible, and Mechanically Robust Electronics. Chem. Rev. 2017, 117 (9), 6467−6499. (4) Onorato, J.; Pakhnyuk, V.; Luscombe, C. K. Structure and design of polymers for durable, stretchable organic electronics. Polym. J. 2017, 49 (1), 41−60. (5) Brochu, A. B. W.; Craig, S. L.; Reichert, W. M. Self-healing biomaterials. J. Biomed. Mater. Res., Part A 2011, 96A (2), 492−506. (6) Casuso, P.; Odriozola, I.; Pérez-San Vicente, A.; Loinaz, I.; Cabañero, G.; Grande, H.-J.; Dupin, D. Injectable and Self-Healing Dynamic Hydrogels Based on Metal(I)-Thiolate/Disulfide Exchange as Biomaterials with Tunable Mechanical Properties. Biomacromolecules 2015, 16 (11), 3552−3561. (7) Wilson, G. O.; Andersson, H. M.; White, S. R.; Sottos, N. R.; Moore, J. S.; Braun, P. V. Self-Healing Polymers. In Encyclopedia of Polymer Science and Technology; John Wiley & Sons, Inc.: Hoboken, NJ, 2002. (8) Wu, D. Y.; Meure, S.; Solomon, D. Self-healing polymeric materials: A review of recent developments. Prog. Polym. Sci. 2008, 33 (5), 479−522. (9) Cui, J.; Daniel, D.; Grinthal, A.; Lin, K.; Aizenberg, J. Dynamic polymer systems with self-regulated secretion for the control of surface properties and material healing. Nat. Mater. 2015, 14 (8), 790−795. (10) Shi, Y.; Wang, M.; Ma, C.; Wang, Y.; Li, X.; Yu, G. A Conductive Self-Healing Hybrid Gel Enabled by Metal−Ligand Supramolecule and Nanostructured Conductive Polymer. Nano Lett. 2015, 15 (9), 6276−6281. (11) Benight, S. J.; Wang, C.; Tok, J. B. H.; Bao, Z. Stretchable and self-healing polymers and devices for electronic skin. Prog. Polym. Sci. 2013, 38 (12), 1961−1977. (12) Rogers, J.; Someya, T.; Huang, Y. Materials and mechanics for stretchable electronics. Science 2010, 327 (5973), 1603−7. (13) Savagatrup, S.; Printz, A. D.; O’Connor, T. F.; Zaretski, A. V.; Lipomi, D. J. Molecularly Stretchable Electronics. Chem. Mater. 2014, 26 (10), 3028−3041. (14) Savagatrup, S.; Zhao, X.; Chan, E.; Mei, J.; Lipomi, D. J. Effect of Broken Conjugation on the Stretchability of Semiconducting Polymers. Macromol. Rapid Commun. 2016, 37 (19), 1623−1628. (15) Oh, J. Y.; Rondeau-Gagné, S.; Chiu, Y.-C.; Chortos, A.; Lissel, F.; Wang, G.-J. N.; Schroeder, B. C.; Kurosawa, T.; Lopez, J.; Katsumata, T.; Xu, J.; Zhu, C.; Gu, X.; Bae, W.-G.; Kim, Y.; Jin, L.; Chung, J. W.; Tok, J. B. H.; Bao, Z. Intrinsically stretchable and healable semiconducting polymer for organic transistors. Nature 2016, 539 (7629), 411−415. (16) Savagatrup, S.; Printz, A. D.; Rodriquez, D.; Lipomi, D. J. Best of Both Worlds: Conjugated Polymers Exhibiting Good Photovoltaic Behavior and High Tensile Elasticity. Macromolecules 2014, 47 (6), 1981−1992. (17) Peng, R.; Pang, B.; Hu, D.; Chen, M.; Zhang, G.; Wang, X.; Lu, H.; Cho, K.; Qiu, L. An ABA triblock copolymer strategy for intrinsically stretchable semiconductors. J. Mater. Chem. C 2015, 3 (15), 3599−3606. (18) Müller, C.; Goffri, S.; Breiby, D. W.; Andreasen, J. W.; Chanzy, H. D.; Janssen, R. A. J.; Nielsen, M. M.; Radano, C. P.; Sirringhaus, H.; Smith, P.; Stingelin-Stutzmann, N. Tough, Semiconducting Polyethylene-poly(3-hexylthiophene) Diblock Copolymers. Adv. Funct. Mater. 2007, 17 (15), 2674−2679. (19) Printz, A. D.; Savagatrup, S.; Burke, D. J.; Purdy, T. N.; Lipomi, D. J. Increased elasticity of a low-bandgap conjugated copolymer by

CONCLUSION We report a novel approach and exemplify the versatile synthetic procedure for fabricating intrinsically stretchable and room-temperature healable CPs. In comparison to molecularly stretchable and healable segmented CPs,15,17 our graft copolymer concept enables one to preserve the inherent electronic structure of the CP backbone and introduce new mechanical properties such as a low modulus, a high stretchability, and structural healing by grafting nonconjugated polymers with hydrogen-bonding functionalities. In this study, we demonstrate that the molecular parameters (such as the graft lengths and solvent type of the casting solution) can be altered to obtain a fundamental understanding of the mechanical, electrical, and structural properties of CP-based graft copolymers. The length of the side chain plays a crucial role in providing the essential noncovalent cross-linking that can compensate for the stress during stretching. Furthermore, the solvent type of the casting solution dictates the graft copolymer structure that can enhance the electrical conductivity, stretchability, and healing abilities of the cast films. Our developed technique is highly suitable for fabricating a broad range of stretchable organic electronics, and our current studies focus on the realization of such devices.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.7b03291. Materials, synthetic procedures for ATRP initiatorbearing thiophene [2-(thiophen-3-yl)ethyl 2-bromo-2methylpropanoate], acrylate urethane, macroinitiator, and graft copolymers via ATRP, experimental procedures, and further discussions about the results for the characterization of materials, including 1H NMR, FTIR, GPC, UV−vis, FL, SANS, DSC, tensile experiment, SEM, and AFM (PDF)

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REFERENCES

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Andrew Nelson: 0000-0002-4548-3558 David Barker: 0000-0002-3425-6552 Jadranka Travas-Sejdic: 0000-0002-1205-3770 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors thank the University of Auckland, New Zealand, and the Australian Institute of Nuclear Science and Engineering, Australia, for the Ph.D. scholarships for P.B. The authors acknowledge all technical staff members, particularly Radesh Singh and Roger van Ryn, at The University of Auckland, New Zealand, for their kind assistance and technical support. The 8856


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