Trade-off of mechanical and electrical properties in stretchable P3HT/PDMS blending films driven by interpenetrating double networks formation Open Access

In recent years, flexible electronics is booming as a hot and challenging research field due to the rapid development of health and intelligent automation industries, such as the wearable medical apparatus, sensors, and soft robots.^1–4 In these applications, constituent electronic devices may need to adapt intimately with many surfaces such as the skins of human body and involve repeatedly in various dynamical motions in order to realize specific biocompatible functions. Thus, the mechanical properties including deformability and mechanic strength of plastic electronic devices are strongly desirable together with improvements in optoelectronic and chemical performance. Polymeric semiconductors are tremendously promising as active materials in applications of flexible electronics owing to their intrinsic flexibility, large-area processability, and good optoelectronic properties.^5,6 π-electron delocalization and relatively strong π–π couplings of conjugated backbones give rise to the origin of semiconducting nature of conjugated polymers.⁷ Thus, the electronic properties of conjugated polymers are highly relevant with the conjugation length of chains and the crystallization degree of the film. With production of optoelectronic performance, this strong π–π packing or crystallization yet tends to introduce rigidity and fragility rather than stretchability and deformability to polymer films due to the difficulty in strain release at the rigid grain boundaries of the crystalline domains.^8,9 In this way, for flexible electronics applications, compromised excellent mechanical properties and optoelectronic performance are essential. Novel chain synthesis strategies and film processing technologies for flexible electronics are highly demanded and are still challenging.

Currently, there are mainly two categories of strategies to improve the deformability of flexible electronic devices. First is the integration of special structural designs by making non-stretchable materials into specific configurations that can absorb the loaded strain without fracture.^10–13 In this strategy, intricate manufacturing techniques are always involved and might be not applicable for low-cost, large-scale, and high-density device manufacture. Second is the materials innovation by synthesizing novel materials that are stretchable in single or aggregated forms.^14–18 From the materials point of view, this may be a straightforward and intrinsic approach to deal with the issues during the development of soft electronics. One of the successful examples is the incorporation of flexible organic insulating units into organic semiconductor chains by fine molecular design and chemical synthesis, which is capable of imparting various additional functions to the organic thin films with improved electrical properties.¹⁷ 

Instead of the aforementioned complicated chemical synthesis, mechanical blending of polymeric semiconductors with flexible elastomers or plastomers may be an alternative facile way to combine both merits, that is, the mechanical properties of the flexible elastomers and the electrical performance of the organic semiconductors.^6,19–21 Intuitively, the blended elastomer matrix can undertake partially the external strain loadings due to phase separation occurrence and possible interpenetrating network formation and thus enhance the deformability of the blending films. Although the presence of insulating elastomers tends to dilute the conducting networks of the conjugated polymers,^22,23 charge transport properties of the blending films may be well maintained on a condition of good connectivity and percolation of the conjugated chains in the blends.²⁴ 

Common elastomers such as polydimethylsiloxane (PDMS),²⁵ polystyrene-block-poly(ethylene-co-butylene)-block-polystyrene (SEBS),²⁶ and polyurethane (PU)²⁷ are often employed for applications in the stretchable electronics because of their excellent mechanical elasticity and easy processing. Shin et al.²⁶ prepared a highly stretchable channel layer on the PDMS substrate by mixing P3HT with SEBS due to the occurrence of vertical phase separation of P3HT fibers toward the top of the blended films. Zhang et al.²⁸ utilized a phase-separated interpenetrating polymer network (IPN) strategy in the P3HT/PDMS chloroform system. The semiconducting P3HT chains were found to embed in the rubbery matrix, and the blended films demonstrated an improved elasticity, transparency, and hole mobility. In a previous work of Song et al.,¹⁹ they further proposed a slip and rotating mechanism underlying the superior stretching stability of the P3HT/PDMS blend layer at high strains (20%–100%), which might originate from the less developed tight junctions among P3HT nanowires with the supporting effect of the PDMS matrix.

In this work, we blend P3HT nanowires and the PDMS elastomer in order to investigate the synergistic effect of the elastomer on the balance of the mechanical and electrical properties of the blending films. The P3HT/PDMS blending films were prepared via spin-coating on two kinds of substrates: SiO₂/Si and PDMS. Organic field-effect transistors (OFETs), X-ray photoelectron spectroscopy (XPS), atomic force microscopy (AFM), and optical microscopy are employed to characterize the electronic properties, phase separation structures, surface morphologies, and the deformation of films under stretching, respectively. It is found that an interpenetrating double networks structure is formed in P3HT/PDMS blending films that coated on PDMS in contrast to a stratified layer structure on the SiO₂/Si substrate. The mobility of the blending films on both substrates can be reserved with the P3HT content being as low as 5 wt. % due to the well-connected P3HT network. Simultaneously, the interpenetrating double networks structure of the P3HT/PDMS blending film effectively improves its stretching performance.

P3HT (Mw: 20–45 kDa, regioregularity: ≥90%, Polymer dispersity index (PDI): ∼2.1) was purchased from Sigma Aldrich. PDMS (Sylgard 184) and a silicone cleaning solvent (DS-2025) were purchased from Dow Corning. The toluene solvent was purchased from Sinopharm Chemical Reagent Co., Ltd. All materials were used without further purification.

10 mg/ml P3HT/toluene solution was prepared in a water bath at 60 °C. Then, it was cooled down naturally to 23 °C, followed by a 48 h aging at 23 °C to gain P3HT nanowires. The color of the solution changed from orange to dark purple, indicating the formation of nanowires.²⁹ After that, the mother solution (10 mg/ml) was diluted to 1 mg/ml. PDMS was dissolved in toluene [200 mg/ml, base:curing agent = 10:1 (weight ratio)] and mixed with the diluted 1 mg/ml P3HT/toluene solution by the designed weight ratio.

The OFET devices with bottom-gate, bottom-contact configurations were fabricated to test the mobility of the films. The heavily doped n-type Si wafer with 300 nm thick SiO₂ layer was used as the substrate. The substrate was cleaned sequentially with acetone, isopropanol, and deionized water for 15 min via ultrasonic cleaning. With shadow masks, source and drain electrodes (Au/Ti = 50 nm/5 nm) were deposited by using an E-beam evaporator (Lesker LAB18). The channel length (L) and width (W) were 90 µm and 1500 µm, respectively. OFET devices were prepared by spin-coating the P3HT/PDMS solution onto the SiO₂/Si substrate at a spin speed of 1000 rpm for 60 s. All the devices were annealed and cured at 140 °C for 1 h. The prepared OFET devices were characterized using Keithley 2612A units under nitrogen atmosphere. The transfer curves (I_D–V_G) were obtained in the saturation regime. The mobility (μ) of OFETs was calculated in the saturation regime (drain voltage, V_D = −60 V) by plotting the square root of the drain current (I_D) vs the gate voltage (V_G) using the following equation:

where W (1500 µm) and L (90 µm) are the transistor channel width and length, respectively. V_T is the threshold voltage, and C_i is the capacitance per unit area of the silicon dioxide gate dielectric.

For strain-dependent field-effect mobility measurement, all films were spin-coated on a stripped PDMS sheet [base:curing agent = 20:1 (weight ratio)] from toluene solution. The PDMS supported films were stretched to different strains by using a home-made micro-stretcher and then flipped over with strain reserved to attach to the SiO₂/Si substrate onto which an octadecyltrichlorosilane (ODTS) self-assembled monolayer had been grown and source and drain electrodes (Au) had been deposited. The electrical properties of the films parallel (Parallel) and perpendicular (Perpendicular) to the stretching direction were both measured, as schematically shown in Fig. 1. In the repeat-strain tests, the device performance was measured after releasing the films back to 0% strain.

XPS analyses were performed on ESCALAB 250 apparatus equipped with a monochromated Al Kα radiation source and a flood gun for removing the charging of insulating samples. Both sides of the blending films were measured after peeling off from the SiO₂/Si substrates via immersing the sample in a 5 wt. % hydrofluoric acid solution. The depth profiles of the film on the PDMS substrate were obtained using argon ion sputtering.

The surface morphologies of the films on the SiO₂/Si substrate and PDMS substrate were measured by atomic force microscopy (AFM) (Dimension Icon, Bruker) in the tapping mode and peak force Quantitative Nanomechanical Mapping (QNM) mode, respectively.

Stretching test of P3HT and P3HT/PDMS films was performed on a PDMS substrate. The film/PDMS structure was prepared by spin-coating of the solution onto the PDMS substrate. The stretching behavior of the thin films under increasing strain was observed using an optical microscope (OLYMPUS BX51).

First, we investigate the electrical properties of the P3HT/PDMS blending films with respect to those of a pure P3HT film, both spin-coated on SiO₂/Si substrates. The representative transfer characteristics are shown in Fig. 2(a). The average mobilities of the pure P3HT film and P3HT/PDMS blending films are shown in Fig. 2(b). For P3HT/PDMS blending films, the mobilities maintain almost equivalent to those of the pure P3HT film (100 wt. %: 2.62 × 10⁻³ cm² V⁻¹ s⁻¹, 5 wt. %: 2.93 × 10⁻³ cm² V⁻¹ s⁻¹) and have higher on–off current ratios until the content of P3HT is reduced lower than 5 wt. %. Further reducing the content of P3HT to 1 wt. %, the mobility decreases sharply with nearly one order of magnitude. The mobility maintenance of the blending films in a wide range of content pronouncedly suggests that the introduction of PDMS is not certainly detrimental to the electrical performance of the blending films, depending on the microstructures strongly. It also means that the manufacturing cost of flexible OFETs may be decreased substantially by using blending films, instead of pure organic semiconductors.

In order to reveal the microstructures underlying the electronic properties of the P3HT/PDMS blending films, XPS and AFM were employed to observe the phase separation and the nanowire networks in the blending films. Figures 3(a) and 3(b) typically show S 2p and Si 2p signals from both the top surface and bottom surface of the 5 wt. % blending film. On the top surface of the film that contacts with air, no S 2p but only Si 2p is detected, suggesting the separation of PDMS toward the outer surface. On the bottom surface that contacts with the SiO₂/Si substrate, S 2p and Si 2p both appear, indicating the attachment of P3HT to the polarized surface. It is known that the surface energy of PDMS (∼7.3 mJ/m²) is much smaller than that of P3HT (∼19.3 mJ/m²).²⁸ Thus, the PDMS component tends to migrate to the film/air interface and P3HT migrates to the hydrophilic SiO₂/Si substrate during film preparation in order to minimize the interfacial free energy.^19,30 In addition, the P3HT has lower solubility in toluene than that of PDMS,²⁸ leading to P3HT crystallization ahead of PDMS in the spin-coating process. This separated and confined P3HT layer at the interface with SiO₂ can act as the charge transport channel, leading to an equivalent electronic performance as the P3HT net film. Moreover, with the decrease in P3HT content in the blending film, the P3HT separated layer turns thinner and thus may give rise to the larger on–off ratio,³¹ as shown in Fig. 2(a). While at content lower than 5 wt. %, the dramatically declined charge mobility of the blending film can be ascribed to worse network connectivity that will be addressed in the next paragraph.

Figures 4(a) and 4(b) are AFM images for the P3HT/PDMS blending films in a content of 100 wt. % and 5 wt. %, respectively. For the pure P3HT film, nanowires networks are clearly observed since the marginal solvent toluene was employed. For the 5 wt. % blending films, no obvious morphological and phase characteristics can be discerned, supporting the phase separated amorphous PDMS layer at the outer surface. In order to shed light on the morphology of the buried P3HT component, we immersed the blending films (5 wt. %, and 1 wt. %) in a silicone cleaning agent (DS-2025) for 24 h to remove the PDMS phase,²⁸ and the morphologies of the remaining P3HT networks are displayed in Figs. 4(c) and 4(d). Figure 4(c) shows an interconnected P3HT nanowires network within the 5 wt. % blending film; thus, an equivalent charge transport property to the pure P3HT film can be expected. In contrast, the P3HT nanowires appear sporadic and worse connected in the 1 wt. % blending film, interpreting well the sharp drop in mobility. The formation and percolation of the P3HT network that concentrated at the interface with SiO₂ can explain why the P3HT/PDMS blending films demonstrate equivalent charge transport properties as net P3HT in a wide range of P3HT content.

A pure P3HT film has been reported to be brittle and easily broken under strains.³² In order to investigate the deformability of the P3HT/PDMS film, we prepared the blending film on the PDMS sheet so that it can be stretched to different strains by stretching. Considering the completely different surface polarity of PDMS with respect to the SiO₂/Si substrate, the vertical phase separation structure of the blending film on PDMS may be different. Therefore, we further investigated the elemental distribution along the normal direction of the film using XPS depth profiles, and here, a typical result of the 10 wt. % P3HT/PDMS blending film on the PDMS sheet is shown in Fig. 5. Different from the stratified layered structure on the SiO₂/Si substrate (Fig. 3), it is seen that S and Si both can be detected from the top surface to the depth of the 10 wt. % blending film on the PDMS sheet, suggesting that P3HT nanowires penetrate through the PDMS matrix. This result is in agreement with the previous reports that ascribed the absence of stratification between the P3HT and PDMS matrix on PDMS to the low surface energy and substrate surface selectivity of PDMS.¹⁹ 

For mobility measurement of the blending film on the PDMS sheet, we peeled the PDMS sheet off and flip it over onto a SiO₂/Si substrate with the source and drain electrodes fabricated. Figure 6 shows the transfer curves and the calculated mobilities with the variation of P3HT content. Interestingly, the mobilities of the blending films above 5 wt. % are similar (∼1.78 × 10⁻³ cm² V⁻¹ s⁻¹) and equivalent to that of the pure P3HT film (∼2.29 × 10⁻³ cm² V⁻¹ s⁻¹). At lower content than 5 wt. %, the mobility decays sharply by about two magnitudes. This mobility variation with P3HT content of the blending films on the PDMS sheet is quite similar to that on the SiO₂/Si substrate (Fig. 2). It is also noted that the on–off current ratios of the films on the PDMS sheet appear less dependent on the P3HT content, owing to the formation of interpenetrating P3HT and PDMS networks in contrast to a layered structure on the SiO₂/Si substrate. The inset in Fig. 6(b) is AFM height images of the 5 wt. % and 1 wt. % blending films on the PDMS sheet. The 5 wt. % blending film exhibits a well-connected P3HT network on the surface, underlying the observed good electrical property as the pure P3HT film. In contrast, a worse P3HT nanowires network can be seen on the surface of the 1 wt. % blending film, thus leading to a rapid drop of the mobility.

Overall, the field-effect mobility of the P3HT/PDMS blending film demonstrates a similar content-dependence relationship no matter spin-coated on the SiO₂/Si or PDMS substrate although the phase separation structures are completely different on these two substrates (a stratified layer structure on SiO₂/Si and an interpenetrating network on PDMS). The determinant factor on the mobility of the film is the connectivity of the P3HT networks. The electronic properties of the P3HT/PDMS blending films exhibit a rapid drop after an approximate critical content of 5 wt. % where the nanowires interconnection becomes much worse.

Furthermore, the strain-dependent electrical properties of the P3HT/PDMS blending film (5 wt. %) are investigated. Figure 7 shows the stretching performance of the pure P3HT film and 5 wt. % P3HT/PDMS blending film both supported on PDMS sheets through an optical microscope. In Fig. 7(a), the pure P3HT film reveals many short curves and twisted cracks of several to ∼10 μm width immediately at 25% strain, and the cracks propagate wider when stretching to 50% strain. The cracks also demonstrate roughly perpendicularly to the strain direction. In Fig. 7(b), the blending film does not show any macroscopic cracks even under 100% strain, suggesting effective energy dissipation in the blending film. It should be pointed out that we cannot discern the breakage of P3HT nanowires in the blending film with the optical microscope. It has been reported that the elongation at break of self-supported P3HT was molecular weight dependent and was less than 10% for the P3HT film of Mn ∼ 20 kDa (equivalent to the P3HT used in this work).³³ In our case, the pure P3HT film supported on the PDMS sheet still maintains despite of cracks formation at 25% strain. This could be due to the nanowire network formation as well as a relative slippage between the P3HT layer and PDMS support. The latter could also be reflected from the observed curve twisted cracks, which otherwise would be long and straight as revealed in the literature.³⁴ 

The electrical properties of the pure P3HT film and 5 wt. % P3HT/PDMS blending film were examined both along and perpendicular to the strain direction under 0%–100% strains. Figures 8(a) and 8(b) show the transfer characteristics curves of the OFETs and the calculated mobilities according to Eq. (1), respectively. Upon stretching, the mobility of the pure P3HT film drops rapidly from ∼2.3 × 10⁻³ cm²V⁻¹s⁻¹ to ∼1.4 × 10⁻⁴ cm²V⁻¹s⁻¹ at the strain of 25% along the strain direction. It drops further with nearly two orders of magnitude at the strain of 100%. However, the P3HT/PDMS blending film demonstrates enhanced field-effect mobility under strain, displaying higher tolerance to the mechanical stretching with respect to the pure P3HT film. Its mobility drops by about only one order of magnitude at the strain of 100%, much less than the P3HT film. In the repeat-strain tests as shown in Fig. 8(c), the measured mobility of the blending films is always higher than the pure P3HT film, suggesting that the cracking of the P3HT nanowires is less in the blending film. The improved mechanical and electrical properties of the P3HT/PDMS blending film can be explained by the interpenetrating double networks formation that helps the energy dissipation on both networks. In previous work,¹⁹ a slip and rotating mechanism was proposed to underpin the enhanced stretchability of the P3HT/PDMS blend film. This nanowires slippery and rotation of the P3HT network readily leads to a partial release of the external strain and thus an improved tolerance to stretching.

In addition, the field-effect mobilities measured perpendicular to the strain direction are also present. The P3HT pure film displays an isotropic transport behavior, while the P3HT/PDMS blending film shows an anisotropic behavior. The mobility during stretching on the perpendicular direction is a bit superior to the parallel direction at a large strain (ε > 50%). This anisotropy of mobility can be correlated with two factors, nanowire alignment and crack formation.³⁵ The difference of mobility at large strain may be dominated by crack formation effect. Because cracks are usually formed perpendicular to the strain direction, it usually has a greater negative effect on the charge transport in the parallel direction. These results demonstrate that the stretching stability of the blending film with an elastomeric matrix is superior to that of the pure film. The blending of PDMS with a small amount of P3HT nanowires in a content of 5 wt. % presently appears to compromise well the mechanics and the electrical character of the semiconducting film.

In summary, we blended one-dimensional P3HT nanowires with elastomeric insulator PDMS to prepare stretchable semiconducting blending films on two kinds of substrates with different polarity, SiO₂/Si and PDMS sheet. The surface morphology, phase separation, and film deformability strain were investigated with multiple techniques such as AFM, XPS, and optical microscopy. On the SiO₂/Si substrate, a stratified layer structure was formed with P3HT nanowires separated mainly toward the interface with SiO₂. On the PDMS substrate, an interpenetrating double networks structure was obtained in the P3HT/PDMS blending films. On both substrates, the mobilities of the P3HT/PDMS blending films are comparable to the pure P3HT film in a wide range of P3HT content (≥5 wt. %) due to the good connectivity of P3HT networks in the films. Furthermore, the blending films supported on PDMS display much improved electrical and stretching performance with respect to the pure P3HT film. This balanced mechanical and electrical performance of the blending films is determined by the well-connected P3HT nanowires network dispersed in the PDMS elastomer matrix.

This work was supported by the National Natural Science Foundation of China (Grant No. 51673182) and the Fundamental Research Funds for the Central Universities (Grant No. WK2060190053).