Thin-film field-effect transistors are essential elements of stretchable electronic devices for wearable electronics1,2. All of the materials and components of such transistors need to be stretchable and mechanically robust3,4. Although there has been recent progress towards stretchable conductors5,6,7,8, the realization of stretchable semiconductors has focused mainly on strain-accommodating engineering of materials, or blending of nanofibres or nanowires into elastomers9,10,11. An alternative approach relies on using semiconductors that are intrinsically stretchable, so that they can be fabricated using standard processing methods12. Molecular stretchability can be enhanced when conjugated polymers, containing modified side-chains and segmented backbones, are infused with more flexible molecular building blocks13,14. Here we present a design concept for stretchable semiconducting polymers, which involves introducing chemical moieties to promote dynamic non-covalent crosslinking of the conjugated polymers. These non-covalent crosslinking moieties are able to undergo an energy dissipation mechanism through breakage of bonds when strain is applied, while retaining high charge transport abilities. As a result, our polymer is able to recover its high field-effect mobility performance (more than 1 square centimetre per volt per second) even after a hundred cycles at 100 per cent applied strain. Organic thin-film field-effect transistors fabricated from these materials exhibited mobility as high as 1.3 square centimetres per volt per second and a high on/off current ratio exceeding a million. The field-effect mobility remained as high as 1.12 square centimetres per volt per second at 100 per cent strain along the direction perpendicular to the strain. The field-effect mobility of damaged devices can be almost fully recovered after a solvent and thermal healing treatment. Finally, we successfully fabricated a skin-inspired stretchable organic transistor operating under deformations that might be expected in a wearable device.
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This work was supported by Samsung Electronics and the Air Force Office of Scientific Research (grant number FA9550-15-1-0106). S.R.-G. acknowledges the Fonds de Recherche Québécois, Nature et Technologie (FRQNT) for a postdoctoral fellowship. Y.-C.C. acknowledges the Ministry of Science and Technology, Taiwan, for partial financial support (project 104-2923-E-002-003-MY3). F.L. thanks the Swiss National Science Foundation for an Early Mobility Postdoc grant. B.C.S. acknowledges the National Research Fund of Luxembourg for financial support (project 6932623). J.L. acknowledges support from the National Science Foundation Graduate Research Fellowship Program under grant DGE-114747. T. Kurosawa acknowledges support from the Office of Naval Research (N00014-14-1-0142). X.G. acknowledges support from the Bridging Research Interactions through the collaborative Development Grants in Energy (BRIDGE) programme under the SunShot initiative of the Department of Energy (contract DE-FOA-0000654-1588). Use of the Stanford Synchrotron Radiation Lightsource, SLAC National Accelerator Laboratory, is supported by the US Department of Energy, Office of Science, Office of Basic Energy Sciences under contract DE-AC02-76SF00515. X-ray diffraction studies were performed at the Stanford Nano Shared Facilities.
The authors declare no competing financial interests.
Extended data figures and tables
a, Polymer composition by 1H NMR. The ratio of PDCA moieties incorporated in the polymer backbone is determined by the integration of protons (1) versus the alkyl-chain terminal protons (2). b, General characterization of P1–P4. aDetermined from thermogravimetric analysis. bThe HOMO (highest occupied molecular orbital) energy level was calculated from cyclic voltammetry. Potentials versus Ag/AgCl using 0.1 M TBAPF6 (tetrabutylammonium hexafluorophosphate) in CH3CN (acetonitrile) as the electrolyte solution. cCalculated using the following equation: gap = 1,240/λonset of polymer film. (PDI, polydispersity; Td, degradation temperature; Mw, weight average molecular weight; λ, absorption wavelength in ultraviolet/visible spectroscopy).
a, Chemical structure of model compound 1 and 1H NMR at various concentrations of compound 1 in CDCl3. Upon increasing concentration from 0.05 M to 0.8 M, a distinct shift of the amide proton (black arrow) towards low fields is observed. This indicates hydrogen bonding between the PDCA moieties, as previously observed31. A dimerization constant of 0.18 M−1 was determined by plotting concentration versus chemical shift and fitting using a dimer association model32. b, Molecular structure of 1 showing intermolecular hydrogen bonds determined by single-crystal X-ray diffraction. Ellipsoids are set at the 30% probability level. Selected hydrogen atoms are omitted for clarity. c, Chemical structure of model oligomer M1 and 1H NMR of M1 at various temperatures in 1,1,2,2-tetrachloroethane-d2 showing the amide NH protons peak. The chemical shift upon temperature increase indicates a breaking of the hydrogen bonds formed between the polymer chains (shown as insets). Oligomer M1 was used for this study because the solubility of the polymer was not sufficiently high to perform a similar study under high concentrations. (a.u., arbitrary units; ppm, parts per million.)
a, Optical microscope images of P1, P3 and P5 as function of applied strain (0%–100%). b, Height and phase AFM images of P3 under 100% strain showing no crack formation. c, Table of grazing incidence X-ray diffraction data for P1 and P3 films as a function of strain (0%–100%). The samples are annealed at 150 °C for 10 min. A reduction in the mean size of crystallites is observed for P1 to P3. The semiconductor is 35 nm thick. (FWHM, full-width at half-maximum.)
a, b, Atomic force microscopy images of damaged (a) and healed (b) thin film of P3 after solvent and thermal annealing. We note that all the previously observed nanocracks were absent after the healing process. c, Dichroic ratio of P3 healed thin film as determined by polarized ultraviolet–visible spectroscopy. d, e, Atomic force microscopy images of damaged (d) and healed (e) thin film of P5 after solvent and thermal annealing. We observed that a small number of nanocracks remained in the film. f, Dichroic ratio of P5 healed thin film as determined by polarized ultraviolet–visible spectroscopy. We observed that the dichroic ratio of healed film of P3 fully recovered to a value similar to that of the pristine film without damage. On the other hand, when P5 is subjected to the same treatment, the dichroic ratio was not restored, indicating that the movement of the polymer chains was insufficient to restore the film’s mechanical properties.
Extended Data Figure 5 Fabrication and electronic properties of a fully stretchable 5 × 5 transistor array.
a, Fabrication process of fully stretchable OTFTs. (1) Transfer printing of the carbon nanotube gate electrode as prepared by spray-coating a carbon nanotube solution (10 mg ml−1 in CHCl3) on SEBS substrate (thickness 200 μm). (2) Contact transfer printing of PDMS dielectric layer as prepared by spin-coating a diluted PDMS (220 mg ml−1 in CHCl3) on OTS-treated SiO2 substrate on a carbon nanotube gate electrode. (3) Contact transfer printing of semiconducting polymer layer (prepared by spin coating onto OTS-treated SiO2 substrate) on PDMS dielectric layer. (4) Spray coating of carbon nanotube (70 nm)/PEDOT:PSS (30 nm) source and drain electrodes on the semiconducting layer. b, c, On/off current mapping (b) and statistical distribution (c) of 5 × 5 fully stretchable OTFT arrays (width 1,000 μm; length 150 μm). Number of devices from a single substrate, n = 25.
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Oh, J., Rondeau-Gagné, S., Chiu, Y. et al. Intrinsically stretchable and healable semiconducting polymer for organic transistors. Nature 539, 411–415 (2016). https://doi.org/10.1038/nature20102
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