Elastomeric Conducting Polyaniline Formed Through Topological Control of Molecular Templates Hangjun Ding,†,‡,# Mingjiang Zhong,†,∇,# Haosheng Wu,‡ Sangwoo Park,† Jacob W. Mohin,† Luke Klosterman,‡ Zhou Yang,∥ Huai Yang,*,⊥ Krzysztof Matyjaszewski,*,† and Christopher John Bettinger*,‡,§ †
Department of Chemistry, ‡Department of Materials Science and Engineering, and §Department of Biomedical Engineering, Carnegie Mellon University, 5000 Forbes Avenue, Pittsburgh, Pennsylvania 15213, United States ∥ School of Materials Science and Engineering, University of Science & Technology Beijing, 30 Xueyuan Road, Beijing 100083, People’s Republic of China ⊥ School of Engineering, Peking University, Beijing 100871, People’s Republic of China S Supporting Information *
ABSTRACT: A strategy for creating elastomeric conducting polyaniline networks is described. Simultaneous elastomeric mechanical properties (E < 10 MPa) and electronic conductivities (σ > 10 S cm−1) are achieved via molecular templating of conjugated polymer networks. Diblock copolymers with star topologies processed into self-assembled elastomeric thin films reduce the percolation threshold of polyaniline synthesized via in situ polymerization. Block copolymer templates with star topologies produce elastomeric conjugated polymer composites with Young’s moduli ranging from 4 to 12 MPa, maximum elongations up to 90 ± 10%, and electrical conductivities of 30 ± 10 S cm−1. Templated polyaniline films exhibit Young’s moduli up to 3 orders of magnitude smaller compared to bulk polyaniline films while preserving comparable bulk electronic conductivity. Flexible conducting polymers have prospective applications in devices for energy storage and conversion, consumer electronics, and bioelectronics. KEYWORDS: conducting polymer, flexible electronics, nanostructure, block copolymer
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serve as stretchable electrode materials in next-generation photovoltaics, energy storage, and biosensors.14−16 Stoyanov17 et al. recently explored elastomeric conducting polymers for use in voltage controlled artificial muscles. Conjugated polymers exhibit electrical conductivities ranging from 10−12 to 103 S cm−1 depending on the polymer composition, doping level, and processing-dependent microstructure.18,19 The rigid backbones of conducting polymers such as polyaniline (PANI) form dense crystalline domains through π−π stacking20 that simultaneously increase both the electronic conductivity (σPANI = 0.1−100 S cm−1) and the Young’s modulus (EPANI = 1−10 GPa).21 Decoupling these properties to create flexible and elastomeric conducting polymers could enable stretchable energy storage materials and conformal bioelectronics interfaces.22−25 However, increasing mechanical compliance and electronic conductivity in concert is challenging because these properties are
echanically compliant electronic conductors are important materials for next-generation electronic devices that are stretchable, flexible, and conformable.1 Intrinsically elastomeric conductors have applications in consumer electronics,2 systems for energy storage and conversion,3 wearable and implantable devices,4 multielectrode arrays,5 and electronic skin6 and components in soft robotics.7 While many existing strategies for stretchable conductors utilize deterministic composites,8 large-scale lithography-free fabrication of flexible conductors using nondeterministic composites is advantageous. Combining orthogonal properties of elastomeric mechanical properties and electronic conductivity in randomly ordered conducting phases is challenging.9 Elastomeric matrices impregnated with carbon nanotubes,10 graphene,11 and inorganic nanoparticles12 produce materials that exhibit mechanical compliance and preserve electrical conductivity during mechanical deformation. Elastomeric conductors can also utilize conjugated polymers as the conducting phase. Conjugated polymers exhibit tunable nanostructures, hybrid ionic-electronic conduction, redox activity, and controlled protein adsorption.13 Elastomeric conducting polymers can © 2016 American Chemical Society
Received: March 1, 2016 Accepted: May 13, 2016 Published: May 13, 2016 5991
DOI: 10.1021/acsnano.6b01520 ACS Nano 2016, 10, 5991−5998
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Scheme 1. (a) Synthetic Route for Star-AAx-b-BAy Block Copolymers and (b) Processing Star-AAx-b-BAy Copolymers into ThinFilm Templates for in Situ ANI Polymerization and Elastomeric Conducting Star-AAx-b-BAy/PANI Composite Formation
orthogonal and competitive.26 Several strategies have emerged to create elastomeric conducting polymers including side-group modification27 and nondeterministic composites with elastomeric matrices and conjugated polymer conducting phases.28,29 However, large volume fractions of rigid conjugated polymers with isotropic microstructures are required to create percolating conducting phases.30,31 Intrinsic limitations in composite design lead to predictable trade-offs between electrical conductivity and mechanical compliance. An elastomeric PANI composite that preserves bulk electronic conductivity of σPANI > 10 S cm−1 is suitable for flexible energy storage systems and conformal bioelectronic interfaces. Bicontinuous structures formed from block copolymer compositions can reduce the percolation threshold compared to bulk materials ϕCP,block,crit≪ ϕCP,bulk,crit31−33 Monomers are confined within one phase of a bicontinuous block copolymer and polymerized into a continuous secondary conducting polymer network.34,35 Linear block copolymer composites have desirable macroscopic properties, but stable bicontinuous morphologies are challenging to prepare in large quantities (gram scale) because of the narrow window in a phase diagram.29 It is also difficult to stabilize bicontinuous morphologies after introducing additional components.36 Here we present a polymer template that confers elastomeric properties to networks composed of PANI while achieving electrical conductivities that are comparable to bulk PANI. Nanostructured block copolymer templates that employ star
topologies can achieve this unique combination of mechanical and electrical properties.
RESULTS AND DISCUSSION Synthesis, Preparation, and Characterization of Primary Star Block Copolymer Networks. The ideal polymer template for flexible and stretchable conducting polymers should preserve elastomeric properties by promoting percolation in the conducting phase. The template should also maximize the intrinsic conductivity of the conducting polymer through doping and control of oxidation states. These design criteria define the materials processing strategy for elastomeric conducting polymers (Scheme 1). Star copolymers are composed of linear block copolymers radially oriented around a central core. Star block copolymers form robust nanostructured molecular templates from solution, thereby avoiding complex processing conditions that are required for phase separation of linear block copolymers. Macroscopic films composed of star block copolymers with outer poly(acrylic acid) (PAA) hard blocks as shells and inner poly(n-butyl acrylate) (PBA) soft blocks as cores were prepared in a scalable manner by drop casting. Templates composed of star block polymers prepared in this manner form continuous phases of PAA coronas at relatively small volume fractions compared to mechanically compliant PBA cores. Aniline (ANI) monomers are then selectively partitioned to the continuous PAA phase through Coulombic interactions and polymerized into percolat5992
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Figure 1. (a) Stress−strain curves of thin films composed of star-AAx-b-BAy block copolymer templates with different molecular weights of each block. Cyclic uniaxial tensile stress−strain response curves of (b) star-AA20-b-BA80 block copolymers and (c) star-AA20-b-BA80/PANI composites (8 wt % PANI). (d) Experimental and predicted values of the electrical conductivity of star-AA20-b-BA80/PANI composites and linear-AA20-b-BA80/PANI composites as a function of PANI content. (e) The Young’s modulus (E) and maximum elongation at break (εmax) are plotted as a function of PANI loading within star-AA20-b-BA80/PANI composites.
processing of blank templates could produce loosely packed networks of star diblock copolymers that are condensed during uniaxial strain. The degree of plastic deformation is consistent in both star-AA20-b-BA80 templates and star-AA20-b-BA80/PANI composites, which suggests that the plastic deformation behavior is governed by the topology afforded by star block copolymers. In situ Formation of Polyaniline Networks. Star-AA20-bBA80 was chosen as the template material for in situ oxidative polymerization of ANI with ammonium persulfate (APS) as the oxidizing agent. The electrical conductivity of star-AA20-bBA80/PANI composites is governed by the quantity and oxidation state of the PANI. The oxidation state of templated PANI was controlled by choosing the ratio of [ANI]0:[APS]0 during in situ ANI polymerization. The electrical conductivity of star-AA20-b-BA80/PANI composites varied nonmonotonically with the ratio of [ANI]0:[APS]0 (Supporting Information, Figure S3). The maximum electrical conductivity was observed when [ANI]0:[APS]0 ratios produce the emeraldine salt (ES) form of PANI, demonstrated by FTIR and UV−vis spectra (Supporting Information, Figures S4 and S5). The amount of PANI in star-AA20-b-BA80/PANI composites was controlled by prescribing the ANI incubation time. The electrical conductivity of star-AA20-b-BA80/PANI composites increases sharply to 30 ± 10 S cm−1 for PANI loading >8 wt %, which is comparable to that of the pure bulk PANI (Figure 1d).44 Conversely, the electrical conductivity of linear-AA20-b-BA80/ PANI composites increases gradually until PANI loading exceeds 20 wt % to yield electrical conductivities of 0.010 ± 0.005 S cm−1. Mechanical and Electrical Characterization of Star Block Copolymer/Polyaniline Networks. Elastomeric starAA20-b-BA80/PANI composites exhibit more robust mechanical properties compared with viscoelastic linear-AA20-b-BA80 templates (Supporting Information). They are elastomeric yet
ing PANI conducting networks. PAA blocks serve as anionic dopants to promote conducting PANI while PBA phases exhibit low Tg to preserve elastomeric properties of the composite and minimize the percolation threshold through excluded volume effects.37 The materials synthesis strategy was evaluated by synthesizing a series of poly(acrylic acid)-block-poly(n-butyl acrylate) copolymers (PAA-b-PBA) with (linear-PAA-b-PBA) or star (star-PAA-b-PBA) topologies using atom-transfer radical polymerization (ATRP) (see Supporting Information for detailed synthesis).38−40 Star block copolymers were synthesized using a three-step process (Scheme 1a). First, linear block copolymers were prepared using Activator ReGenerated by Electron Transfer (ARGET) ATRP at low ppm of Cu catalyst.41 Next, the arm-first approach was implemented to form star topologies by cross-linking active chain ends with divinylbenzene.42 Finally, t-butyl groups were deprotected to form star polymers with pendant acrylic acid blocks.43 Pristine star-PAA-b-PBA films are mechanically robust and elastomeric (Figure 1a). The Young’s moduli (E) of thin films composed of star-AA20-b-BA80 and star-AA80-b-BA20 are 2.5 ± 0.4 MPa and 115.6 ± 5.1 MPa, respectively. This trend is attributed to increased PAA content from 12.3 to 69.2 wt % since Tg,PAA ≫ Tg,PBA (Tg,PAA ∼ 100 °C; Tg,PBA ∼ −50 °C). The ultimate tensile strength (UTS) and maximum elongation at break (εmax) increase proportionally with PBA content (Supporting Information, Table S1). Films composed of star-AA20-b-BA80 with an outer corona of low molecular weight PAA segments (12.3 wt %) exhibit significant hysteresis loops in the first cycle (n = 1), followed by smaller hysteresis loops in subsequent cycles (n = 2−5) of uniaxial tensile strain measurements (Figure 1b). Unrecovered strains for cycles n > 2 are 1 is attributed to mechanically induced reorganization of star diblock copolymers. Solvent 5993
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Figure 2. AFM images of films composed of (a) star-AA20-b-BA80 block copolymers and (b) linear-AA20-b-BA80 copolymers (400 × 400 nm2 negative phase). The values of the z-axis correspond to degrees out of phase in both (a) and (b). The RMS roughness across all samples was RRMS = 1.1 nm. SAXS patterns and azimuthally averaged data (insets) for thin films with the following compositions: (c) star-tBA20-b-BA80 copolymers; (d) star-AA20-b-BA80 molecular templates ;and (e) star-AA20-b-BA80/PANI composites.
linear block copolymer templates are shown in Supporting Information, Table S3. Values of t typically range from 1.1 to 3.1, depending on the anisotropy of the filler and the dimensionality of the matrix (e.g., 2D or 3D).48 Composites with isometric conducting particles in a 2D matrix have values of t = 1.0−1.4, while highly anisotropic particles in a 3D matrix exhibit values of t = 3 due to increased sensitivity to the volume fraction of the conducting phase. The value of t = (d − 1)ν, where d is the dimensionality of the system and ν is a parameter related to the geometry of the conductors. The value of tstar = 0.66 extracted from conductivity data for star-AA20-b-BA80/ PANI composites suggests that the effective dimensionality of PANI networks in star-AA20-b-BA80 templates is dstar < 2. This value is different than extracted values for PANI networks formed in linear-AA20-b-BA80 templates, where tlinear = 2.61. The value of tlinear suggests that PANI networks formed using linear-AA20-b-BA80 templates adopt a higher order dimensionality (d = 3) coupled with a characteristic anisotropy that is given by ν > 1.49 Taken together, these data suggest that starAA20-b-BA80 templates promote PANI percolation by reducing the critical exponent. The different values for percolation thresholds and critical exponents in composites that utilize star-AA20-b-BA80 templates versus linear-AA20-b-BA80 templates are attributed to the distinct nanoscale architecture of each polymer networks as determined by atomic force microscopy (AFM) (Figure 2a,b). Phase imaging of star-AA20-b-BA80 films suggests that hard PAA blocks form continuous networks that surround spherical PBA cores. These structural components are largely absent in films
exhibit electrical conductivities comparable to bulk PANI. Bulk PANI films are stiff and brittle with Young’s moduli EPANI = 1− 10 GPa and maximum elongations of εmax,PANI = 5−10%.45,46 The tensile mechanical properties of star-AA20-b-BA80/PANI composites are shown as a function of PANI content (Figure 1e; Supporting Information, Table S2). PANI networks formed within elastomeric star-AA20-b-BA80 templates exhibit a reduced Young’s modulus and increased elongation at break compared to bulk PANI. For example, star-AA20-b-BA80/PANI8 (composites with 8 wt % PANI) exhibit Young’s moduli of Estar‑AA20‑b‑BA80/PANI‑8 = 8.0 ± 1.4 MPa and maximum elongations at break of εmax,star‑AA20‑b‑BA80/PANI‑8 = 94.8 ± 4.8%. PANI-bearing composites exhibit mechanical properties that are comparable to elastomeric pristine templates. Both starAA20-b-BA80 templates and star-AA20-b-BA80/PANI8 composites exhibit ∼5% unrecoverable deformation during cyclic uniaxial strain for cycles n > 2 (Figure 1c). These results suggest that the intermolecular interactions that confer elastomeric properties to pristine star-AA20-b-BA80 templates are preserved in star-AA20-b-BA80/PANI composites. The percolation threshold of PANI networks created within block copolymer templates can be calculated using the following relationship:47 σ = σ0(ϕCP − ϕCP,crit)t(ν , d)
(1)
The critical exponent t is related to the anisotropy ν and dimensionality d of the conducting phase. Extracted values of ϕCP,crit, σ0, and t for composites formed from both star and 5994
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Figure 3. (a) Circuit diagram using bulk star-AA20-b-BA80/PANI composites to connect LEDs to a voltage source. (b) LED circuits completed with twisted star-AA20-b-BA80/PANI composites: off (left) and on (right). (c) LED circuits completed using bulk conductive star-AA20-bBA80/PANI composites in the following states: no deformation (left), stretching (middle), and bending (right). (d) Normalized changes in resistance values in two-point resistance measurements are plotted for star-AA20-b-BA80/PANI composites as a function of bending curvature in a test sequence with five bending cycles. (e) Normalized changes in two-point resistance measurements are plotted for star-AA20-b-BA80/ PANI composites as a function of uniaxial cyclical strain. The test sequence consists of five cycles of uniaxial strain for a given value of ε. Briefly, each cycle consists of straining the sample from 0 to ε while measuring the relative increase in resistance at the two strain states. The overall test sequence consists of repeating this process for five times for a given value of ε (see Methods section).
PANI) is much smaller than that observed during the formation of composites that utilize linear topologies (linear-AA30-b-BA50/ PANI) (See Supporting Information). These observations suggest that ANI polymerization occurs within the continuous polyanionic PAA domains thereby reducing the percolation threshold for PANI networks. The overall structures of the star block copolymer templates are likely only partially ordered. The most critical component of creating percolating networks is the molecular arrangement of the PAA domains on the exterior of the star block copolymer templates. We cautiously speculate that this molecular configuration promotes mechanically robust networks through strong interaction between nitrogen atoms in PANI and PAA group while simultaneously promoting efficient percolation of PANI formed in situ.
composed of linear-AA20-b-BA80. The morphology is corroborated by small-angle X-ray scattering (SAXS) (Figure 2c,d). Pristine films composed of star-AA20-b-BA80 exhibit short-range order, as indicated from a single peak in the SAXS spectra that corresponds to a d-spacing of 11 nm (Figure 2d). Scattering features are absent in films composed of star-tBA20-b-BA80 (Figure 2c). Star-AA20-b-BA80/PANI composites exhibit a peak that is slightly shifted (d-spacing of 13 nm) (Figure 2e) compared to PANI-free pristine star-AA20-b-BA80. These data suggest that in situ ANI polymerization largely preserves the ordering of star block copolymers templates. The increase in dspacing is attributed to the swelling of the PAA corona phase upon infusion and subsequent polymerization of ANI. The subtle shifts in d-spacing during the formation of composites using block templates with star topologies (star-AA30-b-BA80/ 5995
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METHODS
The combination of the continuous PAA network and the resulting partially ordered network produces templates that both reduce the percolation threshold while preserving flexible mechanical properties. Conversely, it is likely that linear diblock copolymer templates adopt a lamellar or hexagonal morphology based on the relative DP (N), anticipated values for the χ parameter, and large phase space for these morphologies. Both lamellar and hexagonal phases would likely prohibit the formation of continuous PANI structures within the PAA phase, which is consistent with the relatively high percolation thresholds and low intrinsic conductivities. Finally, linear diblock copolymers limit extended intermolecular H-bonding between polymers, which compromises the mechanical properties of the templates. Conversely, block copolymers with star morphologies and PAA coronas utilize acrylic acid groups that can serve as both H-bond donors and acceptors. Coalescence of the PAA phases can produce robust H-bonding that can also deform due to the soft PBA phase. We cautiously speculate that the combination of compliant PBA phases with intermolecular cross-links produces the unique elastomeric properties of the template. In situ polymerization of PANI within the PAA can also serve as an additional component to support H-bonding and electrical conductivity while preserving the underlying nanostructure of the star diblock copolymer template. LED Made of Star Block Copolymer/Polyaniline Networks. The strain-dependent conductivity of star-AA20-bBA80/PANI composites was calculated in multiple deformation modes from two-point resistance measurements (Figure 3). Star-AA20-b-BA80/PANI composites preserved electrical conductivity during mechanical strain. The brightness of lightemitting diodes (LEDs) wired in series with macroscopic films composed of star-AA20-b-BA80/PANI composites (L × W × t = 20 × 5 × 0.2 mm3) was maintained during twisting and uniaxial tensile strain for 100 cycles (Supporting Information, Video S1). The electrical conductivity of star-AA20-b-BA80/PANI composites was maintained during bending as inferred by twopoint resistance measurements, as the radius of curvature was decreased from Rcurve = ∞ (flat) to Rcurve = 1 mm (Figure 3). The bulk electrical conductivity of star-AA20-b-BA80/PANI composites was also preserved in macroscopic films during uniaxial tensile strain. Normalized strain-dependent resistance values ΔRi,tens = ΔRi,tens = ΔRi(ϵ)/R(ϵ = 0) are shown in Figure 3e. The value of ΔRi,tens was constant for a given value of ε and for the ith cycle, where i > 1.
Materials and Detailed Synthesis of Star-AAx-b-BAy. See Supporting Information. Preparation of Star-AAx-b-BAy/PANI Composites. Films were formed by drop casting of star block copolymer templates solutions (100 mg mL−1 in THF) on glass substrates, followed by room temperature solvent removal under vacuum. Star-AAx-b-BAy/PANI composites were formed by incubating star-AAx-b-BAy films in bulk ANI solution and acidic solutions of APS ([APS]0 = 0.0125−2 M). The total amount of PANI in the final composite was controlled by choosing ANI incubation times between 2 and 30 min. Mechanical and Microstructural Characterization. Tensile tests (n = 4) were conducted (Instron 5943 equipped with Bluehill 3 software, Norwood, MA) using a 10 N load cell. Samples between 0.2−0.5 mm in thickness were strained uniaxially at 5 mm min−1 until failure. The Young’s modulus was calculated using the slope of stress− strain curves for ε < 5%. SAXS measurements were conducted using a rotating anode (Rigaku RAMicro 7) X-ray beam with a pinhole collimation and a 2D detector (Bruker Highstar; 512 × 512 pixels). A double graphite monochromator for the Cu Kα radiation λ = 0.154 nm was used with a beam diameter of ∼0.8 mm and a sample-to-detector distance of 1.8 m. Intensity distributions were integrated over the azimuthal angle and are presented as functions of the scattering vector q = 4πsin 2θ/λ, where 2θ is the scattering angle. Phase images of thinfilm morphology were acquired using AFM (NTegra AFM, NT-MDT, Tempe, AZ) at ambient conditions using silicon cantilevers (Budget Sensors, Sofia, Bulgaria; k = 5 N m−1; nominal resonance frequency of f = 150 kHz). Scanning electron microscopy (SEM) analysis was conducted using a Hitachi 2460N scanning electron microscope. The specimens were attached to SEM stubs and coated with gold using a Pelco SC-6 sputter coater prior to imaging. Electrical Characterization of Elastomeric Conducting Polymers. The sheet resistance values of elastomeric conducting polymers in the unstrained state were characterized in ambient conditions using four-probe measurement on a S-1160A probe station (Signatone Corporation, Gilroy, CA, U.S.A.) equipped with two source measuring units (2400 SMU, Keithley Instruments, Cleveland, OH, U.S.A.). Contact electrodes in a square van der Pauw configuration were patterned on the surface of the conducting polymers using silver conductive paste. Both bending- and stretchingdependent resistances were recorded in strained samples using twopoint resistance measurements and a custom-made strain device. Resistance values were measured for composites under bending for five cycles. Two-point resistance values were measured for elastomeric conductors under uniaxial tensile strain for five cycles in succession for each of the following sequences: 0−15%, 0−20%, and 0−25%. Video of cyclic strain was recorded as samples were subjected to uniaxial tensile testing for 100 cycles at nominal strains of 25% (Supporting Information, Video S1). The estimated actual strain in each cycle for n > 1 is 18−20%.
CONCLUSIONS Macromolecular engineering of star block copolymers using ATRP yields templates that confer elastomeric properties to PANI while preserving electrical conductivity. Elastomeric PANI composites exhibit electrical conductivities similar to bulk PANI while reducing the bulk tensile Young’s modulus by almost 3 orders of magnitude from 1 to 10 GPa to
