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Cite this: Chem. Commun., 2018, 54, 10092
Enabling shape memory and healable effects in a conjugated polymer by incorporating siloxane via dynamic imine bond†
Received 5th July 2018, Accepted 10th August 2018
Yaling Zhang,
Chunhui Dai,
Shiwei Zhou and Bin Liu
*
DOI: 10.1039/c8cc05410j rsc.li/chemcomm
A conjugated polymer of poly(fluorene-co-benzothiadiazole) (PFBT) was modified by a soft segment of poly(dimethylsiloxane) (PDMS) to yield a copolymer via dynamic imine bonds. The copolymer was casted as a free-standing film, which was intrinsically flexible, stretchable, and showed tunable shape memory, healable and degradable effects.
How to manipulate elasticity and flexibility of conjugated polymers becomes important with burgeoning flexible and wearable electronics.1 However, the rigid and planar building blocks of conjugated polymers diminish the polymer flexibility. A straightforward approach is to incorporate soft segments to conjugated polymers, either in the form of physical blends or chemical modifications, and meanwhile to keep a balance between electronic property and flexibility of the materials.2 Physical blends are simple to be prepared and could gain polymer flexibility directly,3 whereas it’s hard to achieve homogenous mixtures due to miscibility problems between conjugated polymers and most soft polymers.4 In comparison, chemical modifications could connect soft segments to conjugated polymers from the molecular level and increase flexibility effectively. Common methods, such as side-chain engineering and block copolymer synthesis, have been widely applied in the synthesis of flexible conjugated polymers.5 So far, most of those chemical modifications were carried out with covalent bonds for achieving the long-term stability. Fewer examples used non-covalent interactions,6 which may be due to their vulnerable features. Recently, it’s discovered that conjugated polymers could show increased flexibility, stretchability and even healability when they were properly incorporated with non-covalent interactions like ionic, metal–ligand and hydrogen bonds among polymer chains.7 This effect is closely
Department of Chemical and Biomolecular Engineering, National University of Singapore, 4 Engineering Drive 4, Singapore 117585, Singapore. E-mail:
[email protected] † Electronic supplementary information (ESI) available: Polymer synthesis and characterizations. See DOI: 10.1039/c8cc05410j
10092 | Chem. Commun., 2018, 54, 10092--10095
related to the quick association and dissociation process of those non-covalent interactions under mild conditions, which could provide reversible interactions among polymer chains. However, non-covalent interactions are known to be much weaker than covalent bonds, which may limit the performances and applications of those polymers at a high temperature or for a long-term usage.8 Noteworthy, dynamic covalent bonds, which have the comparable stability as covalent bonds, are capable to be broken and reformed reversibly in an equilibrium control under stimuli conditions.9 Namely, they could provide reversible interactions just as non-covalent interactions under a higher energy input. Thus, they have been used to prepare polymers to fulfil functions like self-healable, degradable and various kinds of stimuliresponsive properties.10 So far, there are only a few reports on conjugated polymers containing dynamic covalent bonds in the polymer backbones to achieve polymer degradability.11 For example, Lei et al. fabricated polymer transistors based on a conjugated copolymer with dynamic imine bonds in the polymer backbone to render the material degradable under acidic stimulus.11a Nevertheless, conjugated polymers with soft segments linked via dynamic covalent bonds have not been reported so far and potential new properties in such polymers remain to be explored. This work presents a primitive example in which dynamic imine bond was used to incorporate a soft segment of poly(dimethylsiloxane) (PDMS) into a conjugated polymer of poly(fluorene-cobenzothiadiazole) (PFBT) to construct a dynamic imine bonded copolymer (Fig. 1). Dynamic imine bond is one of the most widely used dynamic covalent bonds,12 and it was used to induce dynamic connections among polymer chains. PDMS and PFBT were chosen as polymer models for their respectively well-known flexibility and wide applications.13 A crosslinked polymer network was formed since it could increase chain entanglements and material toughness, which enabled bulk material characterizations.14 Thus, a dynamic imine bonded copolymer of PFBT–PDMS was synthesized and casted as a free-standing film, and its potential features including tunable shape memory, healable and degradable effects were explored.
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Fig. 1 Synthetic route to the PFBT–PDMS and the 1H NMR spectra of (a) isophthalaldehyde end-capped PFBT and (b) PFBT–PDMS (CDCl3). Inset pictures: The copolymer film under natural light (left) and under 365 nm UV illumination (right).
Firstly, an isophthalaldehyde end-capped PFBT was synthesized via a one-pot Suzuki polymerization by using the 5-bromoisophthalaldehyde as the end-capping unit (detailed synthesis in ESI†). In 1H NMR spectrum of the polymer (Fig. 1), peaks of isophthalaldehyde (aldehyde –CHQO, 10.24 ppm; benzene ring –CH–, 8.48 and 8.37 ppm) were clearly observed. By calculating the integral ratios between isophthalaldehyde and aromatic backbone (7.80–8.20 ppm), the polymer end-capping ratio was obtained and referred as the aldehyde content, which was specified below. Afterwards, imine reaction was carried out between the isophthalaldehyde end-capped PFBT (aldehyde content B1/7 K) and amino-terminated PDMS (amino content B2/5 K) in toluene under reflux. After the reaction was completed, the mixture showed reversible sol–gel transitions when it was heated above 40 1C and cooled to room temperature (r.t. B 20 1C) (Fig. S1, ESI†). This phenomenon suggested the formation of copolymer PFBT–PDMS by imine bonds, which formed a lightly crosslinked polymer network. Besides, p–p stacking in the conjugated polymers could also contribute to reversible interactions which may together cause the sol–gel transition.15 A free-standing film was obtained by solution casting on a Teflon mould (Fig. 1 inset picture). For comparison, reaction mixtures using PFBT polymers without isophthalaldehyde end-capping units could neither form sol–gel transition nor free-standing film, which proved that the aldehyde content was crucial to induce the polymer crosslinking and film formation. In 1H NMR spectrum of the PFBT–PDMS, the isophthalaldehyde peak disappeared
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and new peaks belonging to imine (–CHQ, 8.42 ppm) and methylene proton (–CH2–, 3.68 ppm) next to imine were observed (Fig. 1). The isophthalaldehyde disappearance and imine formation were also confirmed by FT IR spectra (Fig. S2, ESI†). The copolymer dissolved in tetrahydrofuran (THF) solution showed absorption peaks at 320 nm and 450 nm, which was typical for PFBT. Its emission is similar to that of PFBT as well (Fig. S3, ESI†). The absorption peak of spin-coated polymer film was B10 nm red-shifted as compared to solutions. Afterwards, thermogravimetric analysis (TGA) showed that the polymer only had B5% weight loss up to 430 1C (Fig. S4, ESI†). Differential scanning calorimetry (DSC) results did not show obvious signals within the range of 20 to 200 1C, and one possible reason may be due to a very wide transition. As PFBT has a glass transition temperature (Tg) B 110 1C and PDMS has a Tg B 120 1C, the combination of both may explain the wide transition in PFBT–PDMS. The copolymer film is intrinsically flexible, foldable and pliable, which enabled bulk material characterizations to be performed. The uniaxial tensile test showed that the film was stretchable with B6% elongation at break, a tensile strength of B19.3 MPa and a Young’s modulus of B0.6 GPa at room temperature (Fig. S5, ESI†). When the film was heated to B140 1C, it became highly stretchable and showed B300% elongation at break (Fig. S6 and Movie S1, ESI†). On dynamic mechanical analysis (DMA), the film showed a storage modulus of B1.3 GPa at 30 1C (Fig. 2). When temperature was increased, the storage modulus generally decreased. A short plateau was reached at B140 1C, and the film began to flow after 160 1C, revealing the high temperature induced plasticity in the film. The loss factor tan d showed a wide transition range from 100 to 200 1C with a peak at B140 1C, which was corresponded to Tg of the film. This wide transition may be due to Tg differences in each polymer component as stated above. Besides, there would
Fig. 2 Dynamic mechanical analysis of the copolymer film, (a) storage modulus and loss factor change within a temperature range, (b) creep recovery, (c) normalized stress relaxation behaviours and (d) stress relaxation activation energy (Ea) calculated by fitting characteristic relaxation time t* versus 1000/T by Arrhenius analysis.
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be dangling and free chains considering incomplete endcapping of isophthalaldehyde units in PFBT, which could also cause the wide transition. The viscoelastic property of the copolymer film was tested by creep recovery and stress relaxation experiments. Creep recovery was performed by stress loading and unloading under different temperatures and stresses. The film didn’t show creep effects until the temperature reached 140 1C and it showed more obvious creep effects under higher temperatures and larger loading stresses (Fig. 2 and Fig. S7, ESI†). In the stress relaxation analysis, the film began to show stress relaxation behaviours at temperature around 100 1C. When temperature was increased to 160 1C, the stress relaxation modulus quickly relaxed to zero in several minutes. This stress relaxation behaviour revealed that the film had a crosslinked network based on covalent dynamic bonds. Meanwhile, un-crosslinked free chains and dangling chains should also be concerned since they relax in a shorter time scale. By fitting characteristic relaxation time t* in accordance with temperature based on the Arrhenius analysis, the activation energy for stress relaxation was calculated to be 80.0 0.2 kJ mol 1. According to references of vitrimer,16 on assumption that when the viscosity reached 1012 Pa s, the crosslinked network topology freezing transition temperature (Tv) was obtained to be B43 1C, which was lower than its Tg. Therefore, once the temperature entered glass transition range of the polymer, the film began to show stress relaxation behaviours as observed around 100 1C. Dynamic imine bond is reported to be efficient to construct healable and malleable crosslinked polymers due to the exchangeable feature of the imine crosslinked polymer chains.8 The polymer healability was tested on a bulk film with a surface scratch. After heating at 160 1C for 10 min, the scratch disappeared without any melt of the bulk film, revealing healable effects of the film at microscale (Fig. 3). Besides, concerning the high temperature induced plasticity of the film, malleability of the film was tested by heating two overlapped films under a clamp pressure. After heating the clamped sample at 160 1C for 10 min, the films were able to form a bulk film with slight melts at the overlapped area (Fig. S8, ESI†). However, the recovered film lost a lot of materials toughness, showing limited malleability at macroscale. This may be due to the restricted chain mobility of the PFBT polymer and a relatively low content of dynamic imine bonds as exchangeable and recoverable crosslink points, which overall limited the polymer malleability. Tunable shape memory effects of the polymer film were further tested since that there are potentials in the dynamic
Fig. 3 The surface scratch on a bulk film under microscope (10) (a) before and (b) after healing treatment. The ink serves as a reference to identify the location.
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Fig. 4 The copolymer film shows (a) dual, (b) reconfigurable dual and (c) triple shape memory effects. (i) Deform at 100 1C and fix at 20 1C, (ii) recover at 100 1C and (iii) deform at 140 1C for 20 min and fix at 20 1C; (I) deform at 140 1C and fix at 100 1C, (II) further deform at 100 1C and fix at 20 1C, (III) recover at 100 1C and (IV) recover at 140 1C.
bonded hard–soft copolymer structures.17 As shown in Fig. 4a, a flat piece of film was deformed at 100 1C and fixed at 20 1C to obtain a temporary shape V. When it was heated to 100 1C again, it recovered to the original shape quickly in B20 s. The shape recovery ratio was larger than 90% based on twisting angle changes. Another temporary shape S again demonstrated the dual shape memory effect, which was due to the hard–soft copolymer structures. In Fig. 4b, a film was deformed at 140 1C for 20 min to induce network change, and then cooled to 20 1C to obtain a fixed V shape. Then the film was deformed at 100 1C and fixed at 20 1C to obtain a flat shape. When it was heated to 100 1C, it recovered back to V shape, which served as a new permanent shape. However, not fully recovered shapes were observed, which may be due to stress residues. Another shape S repeated the process, revealing the reconfigurable dual shape memory effects. This was due to the dynamic imine bond serving as reconfigurable crosslink points in the polymer. In Fig. 4c, a film was firstly deformed at 140 1C and fixed at 100 1C to obtain a temporary shape S1. Then it was further deformed at 100 1C and fixed at 20 1C to obtain the temporary shape S2. When the film was subjected to heating, it recovered from S2 to S1 shape at 100 1C, then recovered to its initial shape at 140 1C, revealing the triple shape memory effect of the film, which was mainly due to the wide glass transition of the polymer. In addition to temperature, another typical stimulus that dynamic imine bonds act responsively to is acid. Under acidic conditions, imine equilibrium would be influenced and shifted to dissociation of the bond,11b which would cause degradation of the copolymer. Thus the acid induced degradation effect of the polymer was explored. As shown in Fig. 5, a piece of film was emerged in THF which was added with acetic acid (pH B 1), and it swelled and dissolved quickly within 3 hours. Meanwhile, the one emerged in pure THF still existed as a swollen film (Fig. S9, ESI†), which didn’t dissolve even after 20 hours and showed a swelling ratio of B10 fold. The mixture was subjected to heating at 50 1C for another 6 hours to induce further degradation of the polymer. On GPC analysis results, a decrease of polymer molecular weight and an increase of smaller fragments were clearly observed, revealing the acid-induced degradation of
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References
Fig. 5 The copolymer film degradation under acidic conditions. Photos of films emerged in (a) pure THF, (b) THF added with acetic acid (pH B 1) for different times at B20 1C and (c) GPC analysis of the acid-induced polymer degradation.
the polymer. Similar degradation effect was also observed when the copolymer was treated with a large amount of benzaldehyde (Fig. S10, ESI†). The stimulus as benzaldehyde would participate in the imine equilibrium and compete with the steric hindered isophthalaldehyde in PFBT to react with amine in PDMS, thus also caused degradation of the copolymer. In conclusion, an isophthalaldehyde end-capped PFBT was synthesized first and then reacted with amine-terminated PDMS to yield a copolymer of PFBT–PDMS via dynamic imine bonds. The copolymer was casted as a free-standing film, which was intrinsically flexible and bendable. The film showed stability and pliability with a high mechanical strength at room temperature. While at elevated temperatures, it became highly stretchable and showed obvious viscoelastic properties as stress relaxation and creep. Along with a wide glass transition, the film showed potential dual, triple and reconfigurable shape memory effects. The film was healable at microscale and degradable under stimuli conditions. This work presents a novel exploration of applying dynamic covalent bonds to integrate soft segments into conjugated polymers, which provides new potentials such as shape memory, healable and degradable functions to empower conjugated polymers for applications in the next generation of flexible electronics. We thank the Singapore National Research Foundation (R279-000-444-281 and R279-000-483-281) and the National University of Singapore (R279-000-482-133) for financial support. We also thank Prof. Chaobin He and Dr Jiaotong Sun of National University of Singapore for their help on DMA analysis.
Conflicts of interest There are no conflicts to declare.
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