Sri Harish Kumar
Paleti†
a,
Youngseok
Kim†
a,
Joost
Kimpel
a,
Mariavittoria
Craighero
a,
Shuichi
Haraguchi
a and
Christian
Müller
*ab
aDepartment of Chemistry and Chemical Engineering, Chalmers University of Technology, 41296 Göteborg, Sweden
bWallenberg Wood Science Center, Chalmers University of Technology, 41296 Göteborg, Sweden. E-mail: christian.muller@chalmers.se
First published on 24th January 2024
Conjugated polymers exhibit a unique portfolio of electrical and electrochemical behavior, which – paired with the mechanical properties that are typical for macromolecules – make them intriguing candidates for a wide range of application areas from wearable electronics to bioelectronics. However, the degree of oxidation or reduction of the polymer can strongly impact the mechanical response and thus must be considered when designing flexible or stretchable devices. This tutorial review first explores how the chain architecture, processing as well as the resulting nano- and microstructure impact the rheological and mechanical properties. In addition, different methods for the mechanical characterization of thin films and bulk materials such as fibers are summarized. Then, the review discusses how chemical and electrochemical doping alter the mechanical properties in terms of stiffness and ductility. Finally, the mechanical response of (doped) conjugated polymers is discussed in the context of (1) organic photovoltaics, representing thin-film devices with a relatively low charge-carrier density, (2) organic thermoelectrics, where chemical doping is used to realize thin films or bulk materials with a high doping level, and (3) organic electrochemical transistors, where electrochemical doping allows high charge-carrier densities to be reached, albeit accompanied by significant swelling. In the future, chemical and electrochemical doping may not only allow modulation and optimization of the electrical and electrochemical behavior of conjugated polymers, but also facilitate the design of materials with a tunable mechanical response.
Key learning points– Conjugated polymer-based materials cover the full spectrum of mechanical behavior from stretchable polymers and elastic blends to stiff composites.– The mechanical properties of polymers depend on the processing history and the resulting nano- and microstructure, and should only be compared if measured with the same bulk or thin-film measurement technique. – Chemical and electrochemical doping can strongly alter the rheological and mechanical properties of initially soft conjugated polymers, while materials with a high elastic modulus are less affected. – The electrical and mechanical properties of conjugated polymers such as the electrical conductivity and elastic modulus tend to correlate but can be partially decoupled through the use of multi-component systems and the addition of suitable dopants. – In the case of devices that operate at low charge-carrier densities, the mechanical properties of the undoped semiconductor should be considered, while the properties typical for doped polymers govern the behavior of highly charged devices. |
The mechanical properties of conjugated polymers change when charge carriers are introduced because of a change in the rigidity of the backbone upon oxidation or reduction, but also because of the counterions that are often introduced to balance the charge on the polymer. The function of many types of thin-film as well as bulk electronic devices involves the modulation of the charge-carrier density, which can also alter the mechanical properties during operation. Depending on the dimensions of the polymer film, tape or fiber that is used to construct various devices, different types of measurements can be used to elucidate the thin-film or bulk mechanical response (Section 3).
Thin-film devices such as organic light emitting diodes (OLEDs), organic solar cells and organic field-effect transistors (OFETs) employ conjugated polymers in their semiconducting state and then modulate their charge-carrier density through the application of an electric potential or through the exposure to an external stimulus such as light. In addition, modest chemical doping of thin-film devices can be used to improve charge transport through trap filling and to reduce contact resistance effects.1,2 The charge-carrier density in OLEDs and organic photovoltaic (OPV) devices typically reaches values of 1021 to 1023 m−3 (Fig. 1),3 which implies that only few sites are charged (total number of sites about 1027 m−3). In the case of OFETs higher values are reached but charges accumulate in a nanometer thin region at the interface with the gate dielectric,1 meaning that most of the material remains uncharged. Hence, to understand the mechanical response of conjugated polymer films in thin-film devices, their semiconducting (or weakly doped) state should be considered (Section 4).
Instead, bulk devices such as organic thermoelectric (OTE) generators employ conjugated polymers in a strongly oxidized or reduced state with charge-carrier densities in the range of 1026 to 1027 m−3 (Fig. 1), often brought about via chemical doping (Section 5).2 Other types of thin-film devices such as organic electrochemical transistors (OECTs)4 or bulk devices such as polymer actuators5 modulate the electrical properties of the polymer via electrochemical doping, with the conducting state often exhibiting a very high charge-carrier density of up to 1027 m−3 (Fig. 1), which is only limited by the high degree of swelling that the material experiences (cf. Section 6). For such highly charged devices the mechanical properties of the doped (and swollen) polymer are decisive.
This review will first explore the configuration and conformation of conjugated polymers from a synthesis and structure–processing–property perspective (Section 2). Then, various techniques are discussed that can be used to study the mechanical properties of thin-films or bulk materials (Section 3), which are composed of either the neat semiconductor or a multi-component system of the conjugated polymer together with an acceptor or dopant molecule, another conjugated or an insulating polymer, or a reinforcing agent such as a carbon allotrope or a nanocellulose particle (Section 4). Emphasis is put on the impact of chemical and electrochemical doping on the mechanical properties of conjugated polymers (Sections 5 and 6). Finally, the mechanical properties are discussed in the context of different applications, with focus on organic solar cells as an example of a device where the undoped semiconductor is more relevant, as well as OTE devices and OECTs, which rely on highly chemically or electrochemically doped materials, respectively (Section 7).
Alternatively, unsubstituted conjugated polymers can be prepared by chemical oxidative polymerization and are processable in their oxidized form provided suitable counterions are chosen that impart solubility. The most prominent example is oxidative polymerization of poly(3,4-ethylenedioxythiophene) (PEDOT) in the presence of polystyrene sulfonate (PSS).6 The resulting polymer:polyanion PEDOT:PSS complex can be processed as a dispersion from water, yielding p-type conducting films or fibers with an electrical conductivity of up to 3500 S cm−1 and a Young's modulus E = 22 GPa (Section 5).7 Recently, a similar approach has been reported for poly(benzodifurandione) (PBFDO), which involved oxidative polymerization and in situ reductive n-doping, resulting in a polymer:proton complex that can be processed from dimethyl sulfoxide (DMSO),8 which can be wet-spun into n-type conducting fibers with an electrical conductivity of up to 1600 S cm−1 and a Young's modulus E = 19.5 GPa.9
The introduction of side chains greatly enhances the solubility of growing polymer chains, which has led to the development of a plethora of polymerization techniques for conjugated polymers.10,11 Polymers with a simple repeat unit such as poly(3-alkylthiophene)s (P3ATs) can be prepared by chemical oxidative polymerization, which yields a regio-random and thus disordered polymer with a low elastic modulus (Section 4). Instead, Kumada (McCullough method) or Negishi coupling (Rieke method) can be used, involving monosubstituted organomagnesium halide (Grignard) or organozinc halide reagents (Fig. 3a), which permit the synthesis of regio-regular P3ATs that feature an elastic modulus of several 100 MPa at room temperature (Section 4). Kumada coupling can result in high-molecular weight polymers, e.g. P3HT with a number-average molecular weight Mn >300 kg mol−1 has been reported,12 which is significantly larger than the entanglement molecular weight (see Section 2.3).13
Polymers with more complex repeat units can be realized by combining two monomers, usually done via polycondensation methods using precious metal catalysts.10,11 A dibrominated aromatic unit undergoes polymerization with difunctionalized aromatic comonomers through Stille coupling (diorganotin), Suzuki coupling (diorganoboronate) or Kumada coupling (diorganomagnesium) (Fig. 3b). Some monomers feature sufficiently active aromatic C–H bonds, which enable direct arylation polymerization through reaction with a dibrominated aromatic monomer (Fig. 3b).14
The solubility of a polymer generally decreases with the degree of polymerization, and many conjugated polymers already become insoluble once even modest molecular weights are reached, which complicates synthesis, workup, characterization and, finally, processing. Molecular weight control in polycondensation reactions is – according to Carother's equation – linked to the stoichiometric ratio of the two monomers involved in the coupling reaction. Any stoichiometric imbalance means that at high monomer conversion the majority of chain ends are populated by the reactive groups associated with the excess monomer, which hinders any further reaction, leading to shorter chains. Hence, a stoichiometric imbalance can be deliberately selected to limit the molecular weight of the polymer, ensuring that the prepared material remains soluble. Conversely, high molecular weights are difficult to achieve with polycondensation reactions because (1) a 1:1 stoichiometric ratio is difficult to ensure and (2) the solubility of high-molecular-weight conjugated polymers is limited, which explains why only few conjugated polymers feature a high level of ductility (Section 4.2).
Conjugated repeat units can be incorporated into the full spectrum of copolymer architectures, either with other conjugated repeat units and/or with non-conjugated parts. Random copolymerization involving two or more conjugated comonomers, again usually through polycondensation reactions, is widely used to combine the optical and electrical properties of the respective homopolymers. A variety of regular copolymer architectures have been described such as random or alternating multiblock copolymers comprising conjugated segments separated by flexible, non-conjugated spacers, which is a widely used approach for adjusting the mechanical properties of conjugated polymers (Fig. 4).15,16 Moreover, conjugated and non-conjugated polymers can be combined in different ways to create copolymers with different types of mechanical response. For instance, P3HT blocks have been combined with a polyethylene block resulting in AB type block copolymers that display greatly enhanced ductility17 but retain their electronic properties even if the conjugated block only comprises as little as 10 wt% of the copolymer.18 Another intriguing example are ABA type copolymers comprising two P3HT blocks connected via a flexible polymethacrylate (PMA) block, resulting in a material with a nanostructure that is typical for thermoplastic elastomers.19
A rigid backbone implies that the entropy change upon dissolution or melting is low leading to high dissolution and melting temperatures, which can limit polymerization as well as processing. Some unsubstituted polymers can be processed when oxidized with suitable dopants. For instance, polyaniline (PANI) becomes soluble in organic solvents and can be melt-processed upon blending with a variety of commodity polymers when protonated with, e.g., dodecylbenzenesulfonic acid (DBSA),21 where the alkyl chain increases the overall conformational entropy of the polymer:counterion complex.
A more common approach to facilitate polymerization (cf. Section 2.1) and impart processability is the decoration of the conjugated backbone with flexible alkyl or oligoether side chains (Fig. 4), again leading to a higher overall conformational entropy even though in some cases long side chains can actual increase lp.22,23 As a result, neutral polymers with an appreciable degree of polymerization (molecular weight) remain soluble in the reaction medium as well as the solvent(s) chosen for workup, characterization and, finally, processing.
The optimal side-chain length represents a compromise between solubility, which benefits from longer side chains, and maximizing the fraction of the (opto)electronically active conjugated part, which implies that side chains should be short. In case of P3ATs, for instance, hexyl side chains have an optimal length, resulting in a soluble material with good charge transport characteristics.24 However, if mechanical properties are also considered slightly longer heptyl side chains may be preferable because the glass transition temperature Tg is lowered to below room temperature, which results in a significant reduction in elastic modulus (Section 4).25
Most conjugated polymers don alkyl, alkoxy, thioalkyl or oligoether side chains that are either linear or branched. Side chains can be functionalized with, e.g., sulfonic or carboxylic acid groups, amines, urethane and ester groups, which can introduce intermolecular interactions such as ionic or hydrogen bonds. Moreover, crosslinkable moieties can be added, which facilitate the formation of covalent network points (see Section 4.2).
The nanostructure of relatively flexible polymers such as regio-regular P3HT, which has a persistence length of lp ≈ 3 nm at room temperature in dichlorobenzene,26 in many ways resembles that of polyethylene (modelling of the chain conformation yielded a similar value of lp ≈ 4 nm27). P3HT chains can fold but when crystallized from the melt tend to organize in a fringed-micelle type nanostructure with crystalline domains that are separated by rigid amorphous as well as amorphous regions.28 Crystalline domains are connected via tie chains, provided the molecular weight is sufficiently high, i.e. number-average molecular weight Mn ≥ 25 kg mol−1 in case of regio-regular P3HT, leading to a ductile material with a tensile elastic modulus E of about 0.2 MPa and charge-carrier mobility μ < 0.1 cm2 V−1 s−1 (Section 4).13,29 The molecular-weight distribution strongly impacts the probability of tie-chain formation and concomitantly charge transport in thin films.30 The regio-regularity dictates to which extent the polymer can order. The strong tendency to π-stack can result in the growth of highly elongated, fibril-like crystallites, which can form individual “whiskers” or “nanofibrils” when solidified from dilute solution.31
More rigid conjugated polymers such as PBTTT, which has a persistence length of lp ≈ 4–5 nm in chlorobenzene at room temperature,32 display a liquid–crystalline phase above the melting temperature Tm, which facilitates the development of extended ordered domains, resulting in a higher μ > 0.3 cm2 V−1 s−1 and E ≈ 1.8 GPa (buckling method) compared with P3HT.29 Many conjugated polymers have a large persistence length that does not favor chain folding. Some rigid conjugated polymers do not feature any long-range order when studied with X-ray diffraction and should instead be thought of as comprising somewhat ordered regions with varying size and degrees of para-crystallinity, that are embedded in a disordered matrix.33
One class of rigid materials are diketopyrrolopyrrole (DPP) based copolymers such as PDPP-3T (see Fig. 2 for chemical structure), which has a persistence length of lp ≈ 16 nm in o-dichlorobenzene at room temperature23 and thus is only able to form a fringed-micelle type nanostructure. PDPP-3T with an intermediate molecular weight of about Mn ≈ 90 kg mol−1 features a nanostructure comprising ordered regions that are connected by tie chains, resulting in a peak in E of 460 MPa (buckling method; 700 MPa if measured with force microscopy; cf. Section 3) and μ ≈ 4 cm2 V−1 s−1.34 Instead, lower molecular-weight material features chain-extended crystals but no connectivity, while entanglements of higher molecular weight PDPP-3T hinder the formation of ordered regions, both resulting in a decrease in E and μ. Attempts to improve the poor ductility of DPP-based copolymers by introducing, e.g., hydrogen-bonding moieties have been met with limited success. The introduction of 10% amide- or urea-containing side chains does not strongly influence E but can alter the strain at break εbreak, measured with the film-on-water method (see Section 3.2), which appeared to occur because of the concomitant change in the degree of order.35
Support | d | T g | E′ or G′ | ε break | ε crack | Swelling | |
---|---|---|---|---|---|---|---|
Tensile testing | — | ≫μm | (✓) | ✓ | ✓ | — | — |
DMA | — | ≫μm | ✓ | ✓ | (✓) | — | — |
DMA, fiber mesh | ✓ | <μm | ✓ | — | — | — | — |
DMA, elastic support | ✓ | <μm | ✓ | ✓ | (✓) | — | — |
Shear rheometry | ✓ | ≫μm | ✓ | ✓ | — | — | — |
QCM-D | ✓ | <μm | — | (✓) | — | — | ✓ |
Nanoindentation | ✓ | ≈μm | (✓) | ✓ | — | — | ✓ |
AFM | ✓ | <μm | — | ✓ | — | — | ✓ |
Buckling | ✓ | ≪μm | — | ✓ | — | (✓) | — |
Film-on-water tensile testing | ✓ | ≪μm | — | ✓ | ✓ | ✓ | — |
The stress σ = F(ε)/A, i.e. the applied force per cross sectional area of the sample, can be expressed as engineering stress where the area is the initial area, A = A0, and true stress if the area corresponds to the actual area A = A(ε). Initially, the sample experiences elastic deformation and the material would return to its original shape if the stress was removed at this point. Stress increases linearly with strain ε, and the slope is referred to as the Young's modulus E = σ(ε)/ε (note that for static deformation the symbol E is written without an apostrophe; cf. Section 3.1.2). The tensile stiffness of the sample depends on E as well as the geometry of the sample according to:
(1) |
Tensile deformation beyond the yield stress σyield results in plastic (permanent) deformation, i.e. the material would not recover its original shape if the stress was removed. Hence, stretching can be used to produce polymer tapes and fibers with a high degree of permanent orientation, which has been exploited for aligning a wide range of conjugated polymers including PA, P3ATs, poly(p-phenylene vinylene)s (PPVs) as well as various blends comprising conjugated polymers and insulating polymers, including polyethylene and polyaramide.37
Ultimately, the sample breaks, yielding the strain and stress at break, εbreak and σbreak. A material with a low εbreak is referred to as brittle, while ductile materials feature an εbreak ≫ 100%. The area under the stress–strain curve is the toughness, i.e. the energy per volume absorbed by a material during tensile deformation, meaning that an initially stiff and then ductile material has a high toughness.
Definition of mechanical propertiesStiffness (stiff): the ability of a material to resist deformation. Flexibility (flexible): inverse of stiffness. The ability of a material to deform. Elasticity (elastic): the ability of a material to reversibly deform under application of a mechanical stress, including the return to its original shape once the stress is removed. Ductility (ductile): the ability of a material to irreversibly deform under application of a mechanical stress without fracture. Yield and tensile strength (strong): the maximum stress that a material can experience without plastic deformation (yield) and fracture, respectively. Stretchability (stretchable): the ability of a material to reversibly or irreversibly deform through tensile drawing. Toughness (tough): the amount of energy per volume that a material can absorb during deformation prior to fracture. |
Fig. 6 Selected modes of deformation available via (a) tensile testing, (b) rheometry and (c) dynamic mechanical analysis (DMA); the moving part is shown in yellow. |
Alternatively, a thin film of the conjugated polymer can be deposited on a glass fiber mesh,41 an elastic substrate made of, e.g., polydimethylsiloxane (PDMS)42 or a kirigami-cut polyimide film43 but only the latter two can provide absolute measurements of the elastic modulus. In case of fiber-mesh supported samples the cross-sectional area is ill-defined since the mesh and sample are entwined with each other. Hence, the glass fiber mesh method only allows to determine relative changes in storage and loss modulus, which is however widely used for the determination of the Tg of conjugated polymers and organic photovoltaic blends (Table 1).41,44–46
Oscillatory deformation is carried out within the linear viscoelastic regime (typically, the strain amplitude is much less than 1%) so that the sample experiences no permanent (plastic) deformation at each temperature and frequency f = ω/2π where a measurement is carried out. The oscillating strain and stress are given by ε(t) = ε0sin(ωt) and σ(t) = σ0sin(ωt + δ), respectively, where δ is the phase lag. For a perfectly elastic material, e.g. a glassy or rubbery material, the phase lag δ = 0, i.e. the material responds instantaneously to a change in strain or stress. Conversely, any viscous component, which is most prominent when a material transitions from the glassy to the rubbery regime, and vice versa, leads to the dissipation of energy and a phase lag δ ≠ 0. The storage (elastic) and loss modulus are given by G′ = σ0/ε0× cosδ and G′′ = σ0/ε0× sinδ. Note that the symbols G′ and E′ are typically used for dynamic shear or tensile deformation, respectively (see Section 3.1.3). The loss tangent is defined as:
tanδ = G′′/G′ | (2) |
Fig. 7 Typical DMA thermograms of the storage and loss moduli, G′ (solid) and G′′ (dashed), vs. temperature T of an amorphous polymer (red), semi-crystalline polymer (blue) and crosslinked, amorphous polymer (turquoise) with rubber plateau modulus G0N, featuring a β-relaxation temperature Tβ, glass transition temperature Tg (also referred to as α-relaxation temperature Tα) and melting temperature Tm; the inset shows G′ and G′′ of a regio-random (red; Mn = 42 kg mol−1, polydispersity index PDI = 2.4) and regio-regular P3HT (blue; Mn = 24 kg mol−1, PDI = 1.2, regio-regularity = 95%) measured with oscillatory shear rheometry at a frequency of 10 rad s−1; data reproduced from ref. 47. |
Entangled polymer melts and solutions feature a rubber plateau modulus G0N and a rheometer can hence be used to study the impact of entanglements on the mechanical properties of conjugated polymer melts and solutions (see Section 4). For example, static shear rheometry has been used to record the viscosity of P3HT solutions as a function of solvent type and polymer molecular weight, which allowed to investigate the impact of chain entanglement on pre-aggregation.48 Oscillatory shear rheometry can be used to study gel formation in conjugated polymer solutions by identifying the crossover point where the shear elastic (storage) and loss modulus are of equal magnitude, G′ = G′′; for a viscous liquid G′ < G′′ while for a gel the elastic response dominates, i.e. G′ > G′′.49,50
Oscillatory shear rheometry also facilitates DMA type measurements where G′ and G′′ are recorded as a function of temperature or deformation rate (frequency), which allows to determine, e.g., the Tg of a material from the peak in G′′ or the loss tangent tanδ = G′′/G′ (Fig. 7).47,51,52 For example, Gomez et al. have used oscillatory shear rheometry to determine the Tg of a wide range of conjugated polymers, which allowed to develop an empirical model for predicting the Tg based on the makeup of the repeat unit.51
(3) |
E = 2G(1 + ν) | (4) |
In case of nanoindentation, an indentation tip with a conical, triangular pyramidal (Berkovich) or cylindrical shape applies a variable load and as a result penetrates the sample causing deformation. The force F is recorded as a function of depth h during loading-unloading cycles (strain- or force-controlled mode), whose frequency can be varied enabling dynamic measurements similar to DMA (see Section 3.1.2). For relatively soft materials such as polymers, it is important that the maximum indentation depth does not exceed approximately one tenth of the total film thickness (typically d > 1 μm) so that the stiffness of the underlying substrate does not influence the measurement. The stiffness S = dF/dh of an elastic material can be extracted by the Oliver–Pharr method from the first derivative of the part of the force–displacement curve that is recorded at the beginning of the unloading cycle (Fig. 8) and is related to Er according to:55
(5) |
Unlike nanoindentation, mechanical measurements with AFM only allow to access the linear (elastic) deformation regime since the interaction between the AFM tip and sample surface occurs through adhesive/repulsive forces rather than prolonged contact. Hence, only Er can be extracted. Typically, the repulsive force between sample and tip is recorded as a function of the cantilever displacement, and Er is obtained by evaluating the force–displacement curve with the Derjaguin–Muller–Toporov (DMT) model (Fig. 8):61
(6) |
The high surface tension and low viscosity of the water surface results in a pseudo free standing specimen, which helps to minimize any influence from the substrate. However, the sample can be affected by the uptake of water. A comparison of tensile deformation of (1) free-standing films of regio-regular P3HT and (2) samples supported by water (d ≈ 80 nm) revealed a slight reduction in elastic modulus but increase in εbreak from less than 50% to more than 100% in case of the latter, which could be explained with plasticization by water despite the hydrophobic nature of the polymer.53
(7) |
A variation of the buckling method involves cyclic tensile deformation of a film on an unstrained substrate with the maximum strain increasing for each cycle. Buckling occurs during the relaxation step once the film starts to undergo plastic deformation because the maximum strain exceeds εyield, thus providing information about the latter.65 Instead, if tensile deformation is allowed to proceed, fracture of the polymer ultimately occurs, which can be used to determine εcrack. A disadvantage of the buckling method is the limited control over the deformation rate. A comparison of the film-on-water and buckling method yielded a higher elastic modulus for P3HT in case of the latter, which was explained with an at least 100 times higher strain rate in case of the buckling method as well as differences in the mode of deformation (tension vs. compression).66 Moreover, buckling experiments are typically carried out at room temperature.
(8) |
Instead, to describe the response of a sensor coated with a viscoelastic material, for which energy dissipation must be considered, a Kelvin–Voigt model can be used with a complex shear modulus given by:68
G* = G′ + iG′′ = ζ + i2πfη | (9) |
Molecular dynamics simulations can be used to estimate the Tg. For instance, the density of the simulation box as a function of temperature can provide information about the Tg at which the linear expansion coefficient changes, leading to a more rapid decrease in density upon further heating.72 Moreover, molecular dynamics simulations can be used to study the temperature dependence of thermal vibrations of different building blocks such as aromatic rings and side chains, which provides information about local relaxation dynamics and thus the Tg.73
In the glassy regime, polymer chains are unable to relax on a global scale. The conformation of the polymer is “frozen in” on the experimental timescale and the material is brittle, characterized by a large G′ (Fig. 7) but low εbreak. However, local relaxation processes can occur such as relaxation of the side chains of conjugated polymers, which are only frozen in at low temperatures (e.g. below a β-relaxation temperature Tβ; see Fig. 7). Local relaxation allows glassy polymers to absorb impact energy even below Tg resulting in some degree of impact toughness.46
The transition from the glassy to the rubbery regime (or to the viscous regime if the polymer is not entangled) occurs around the Tg, marked by a strong drop in G′ and a peak in G′′ as well as tanδ = G′′/G′ (note that the Tg increases with heating/cooling rate). The Tg of conjugated polymers can be predicted despite strong variations in the makeup of the conjugated (aromatic) backbone as well as the side-chain length and density by considering the difference in mobility of conjugated and non-conjugated atoms.51 A more refined prediction of the Tg could be obtained with a machine-learning model that considered the side-chain fraction, the number of isolated, fused, and bridged aromatic rings, the number of halogen atoms, and the number of double and triple bonds that are not part of aromatic rings.73 Generally, a decrease in side-chain fraction correlates with an increase in Tg.52,74 For instance, a decrease in the side-chain fraction through the incorporation of unsubstituted aromatic units leads to an increase in Tg, as reported for, e.g., DPP-based copolymers whose backbone comprised varying numbers of thiophene rings or fused thiophene units such as thienothiophene.75 Instead, an increase in the side-chain length without altering the makeup of the backbone tends to result in a reduction in Tg, as observed for, e.g., P3ATs.76 The Tg of a (conjugated) polymer determines its mechanical properties around room temperature, i.e. the temperature range where most conjugated polymers are used (cf. Section 7). Hence, a change in side-chain fraction can be used to adjust the elastic modulus.77 Polymers with a low Tg < 0 °C such as polythiophenes with oligoether or long alkyl side chains are soft at room temperature with a tensile elastic modulus E′ as little as 1–10 MPa. Instead, semi-flexible conjugated polymers with a Tg > 50 °C feature E′ ≈ 1 GPa (Fig. 10), while rigid ladder type polymers feature even higher values of, e.g., 8 GPa in case of poly(benzimidazobenzophenanthroline) (BBL) measured with nanoindentation,78 likely because of the absence of side chains.
Fig. 10 Tensile elastic modulus close to room temperature vs. the Tg of conjugated polymers with oligoether (open diamonds) and alkyl side chains (filled circles). Figure adapted with permission from ref. 79; Copyright 2023 (CC-BY), American Chemical Society, with values from tensile testing or DMA (grey),38,80–83 oscillatory shear rheometry (blue; converted using eqn (4) and assuming υ = 0.5)47,51 and one datapoint added from ref. 84. |
In the rubbery regime longer sections of polymer chains are able to relax, i.e. they can adopt a new conformation and thus dissipate stress. Polymer chains are held in place by entanglements provided they are sufficiently long, which is the case if Mn is larger than the entanglement molecular weight Me, i.e. the minimum molecular weight required for polymer chains to entangle (Me ≈ 25 kg mol−1 for P3HT; see Fig. 11a for schematic of an entangled melt).13 Entanglements only persist at timescales less than the time required by chains to disentangle.
Fig. 11 Illustration of different types of network points. (a) Entanglements; (b) crystallites; (c) ionic interactions; and (d) covalent crosslinks. |
In addition to entanglements, polymer chains can form a network via chemical crosslinks, which can be covalent bonds or strong secondary interactions, such as ionic crosslinks. The plateau modulus of a polymer network melt or solution (see Fig. 7) is given by:57
(10) |
In the viscous regime, polymer chains are able to relax freely, unimpaired by entanglements since polymer chains have sufficient time to disentangle and adopt a new conformation. The cross-over time between the rubbery and viscous regime is given by the disentanglement time, meaning that (1) low molecular weight polymers, which do not entangle, only show a glassy and viscous regime, and (2) chemically crosslinked materials cannot disentangle, i.e. the “rubber plateau” continues to exist also at higher temperatures (Fig. 7) and longer times/lower frequencies.
Physical network points are mostly dynamic since they tend to disappear when certain stimuli are applied, such as, chiefly, temperature. For example, crystallites (or more generally ordered domains) disappear upon heating above Tm, hydrogen bonds dissociate above a certain critical temperature, and entanglements of a non-crosslinked polymer melt can disappear when the material experiences elongational deformation.
Since many conjugated polymers π-stack at least to some extent, ordered domains are typically present, which reinforce the material between Tg and Tm, resulting in a storage modulus of 100 MPa to 1 GPa depending on the degree of order. For example, regio-random P3HT, which cannot order, features a low G′ of not more than 10 MPa at room temperature, while crystallites in case of regio-regular P3HT lead to a much higher G′ of about 100 MPa despite a similar Tg.47 Semi-crystalline polymers such as regio-regular P3HT with a low Tg tend to be ductile at room temperature with εbreak ≫ 100% provided the molecular weight is high enough (and the right processing technique is selected; cf. spin-coating vs. spray-coating92) so that tie chains and trapped entanglements can connect nearby crystallites (see Fig. 11b).13 Instead, conjugated polymers with a high Tg are usually reported to be brittle since most measurements are carried out at room temperature76 and would only become ductile if deformed at elevated temperatures (provided the molecular weight is sufficiently high).
A common type of multi-component system within the field of organic electronics are bulk-heterojunction blends that form the active layer of OPV devices. A donor polymer is mixed with a small-molecular or polymeric acceptor, which – once the polymer absorbs light – undergo excited-state electron transfer (Fig. 12). The acceptor can cause embrittlement of the bulk-heterojunction blend,93 for example because of the presence of acceptor-rich domains with poor mechanical connectivity. Besides the distribution and size of domains, the connectivity in donor:acceptor bulk-heterojunction blends depends on the molecular weight of the polymer located in donor-rich as well as mixed domains, which influences both the electrical properties (solar cell efficiency) as well as the mechanical properties (fracture energy).45
Small-molecular additives can be added that at low concentrations function as, e.g., a plasticizer that decreases Tg, at larger concentrations swell the polymer and at high concentrations function as a solvent. Additives can interact with the polymer. A common way to plasticize polymers that form hydrogen bonds is the addition of polar molecules such as water, which decrease the interactions between polymer chains by forming hydrogen bonds with the polymer instead. Another common type of additive is a small-molecular dopant that undergoes ground-state electron transfer with a polymer, resulting in a counterion that balances the charge (polaron) that has been created on the polymer (Fig. 12; see Section 5). Alternatively, the polymer is electrochemically oxidized and takes up counterions as well as solvent molecules from an electrolyte solution (Fig. 12; see Section 6).
Other types of multicomponent systems are blends (or AB copolymers) of two or more polymers; a conjugated polymer can be combined with another conjugated polymer, an insulating polymer or a polyelectrolyte. Blends of two polymers can – depending on their energy levels – undergo ground-state or excited-state electron transfer (Fig. 12), resulting in a chemically doped material or a bulk-heterojunction blend for OPV devices, respectively. The second polymer can also be an insulator, e.g. polyethylene or polystyrene, whose purpose is to adjust the rheological and mechanical properties of the conjugated polymer.94 Moreover, the second polymer can be a polyelectrolyte, e.g. PSS, which then provides the counterions for electrical charges (polarons) on the conjugated polymer. The most widely studied conjugated polymer:polyelectrolyte complex is PEDOT:PSS,6 and numerous ways to modify its mechanical properties have been explored, including plasticization and blending with other polymers.95
The conjugated polymer can also form a composite with a reinforcing agent such as cellulose nanocrystals (CNC) or cellulose nanofibrils (CNF), a carbon allotrope such as graphene or carbon nanotubes (CNTs) as well as other 2D materials such as MXenes. Carbon allotropes are usually added to modify the electrical properties of conjugated polymers,96 but it can be anticipated that they also act as a reinforcing agent, which is commonly observed for nanocomposites with commodity polymers.97 Nanocellulose, instead, is more widely used to modify the rheological and mechanical properties of conjugated polymers,98e.g. for wet-spinning of fibers composed of CNF and PEDOT:PSS,99 but can also enhance the ionic mobility (see Section 7.4).79
Thin films of conjugated polymers can be uniaxially aligned through tensile drawing on a stretchable substrate and by shear, e.g. through rubbing of solid films on a rigid substrate.100,101 Similarly, orientation of free-standing bulk samples can be achieved by solid-state tensile drawing that is terminated prior to fracture or through compression molding of a solid material.102 Fiber spinning of conjugated polymers paired with solid-state drawing, for example, tends to yield filaments with a high degree of uniaxial alignment along the fiber axis, which is essential for achieving a high stiffness and stress at break σbreak (tensile strength).37 How the properties along the two directions perpendicular to the direction of chain alignment are affected depends on the secondary interactions between chains. Strong interactions such as hydrogen bonding or π-stacking tend to strengthen a material perpendicular to the direction of chain alignment and may even enhance charge transport, as observed for solid-state pressed P3HT, which features the highest charge-carrier mobility along the π-stacking direction.103
(11) |
Redox doping entails the addition of an oxidizing or reducing agent (the dopant) that exchanges one or several electrons with the conjugated polymer resulting in either p- or n-doping. The polymer donates or accepts electrons resulting in polarons on the polymer backbone, whose charge is balanced by the ionized dopant molecules that remain as counterions. In case of acid–base doping, the dopant (e.g. a protonic acid or N-DMBI; see Fig. 2 for chemical structure) and the semiconductor exchange a proton (H+)105,106 or hydride (H−)105 resulting in a p- or n-doped material, respectively.
Doped films can be prepared by a variety of different processing methods, such as co- and sequential processing. Co-processing involves mixing of the conjugated polymer and dopant, e.g. via dissolution in a common solvent, followed by spin coating, drop casting, fiber spinning, etc. Sequential processing, instead, entails the preparation of neat polymer films followed by exposure of the solidified material to a dopant solution or vapor. Another variant is ion-exchange doping, where a semiconductor film is exposed to a strong oxidizing (or reducing) agent that is dissolved in an electrolyte solution. Oxidation of the film is followed by exchange of the dopant counterions with alternative ions provided by the electrolyte solution.107,108 For an in-depth discussion of various doping mechanisms and methods, we refer the reader to two recent reviews.2,109
If a high degree of control over the microstructure of a sample is required, sequential doping is typically the preferred method. However, sequential processing of thick films or bulk samples, which are needed for certain mechanical measurements (see Section 3) as well as applications that require free-standing architectures (see Section 7), can result in inhomogeneous doping throughout the material because dopant molecules must diffuse into the sample.110 To achieve a homogeneous distribution of the dopant, porous structures such as foams can be used that ease the ingression of dopant molecules into the material.111 Alternatively, some combinations of polymer and dopant can be co-processed into bulk samples, such as PANI and DBSA, where the dodecyl chains of the dopant impart melt-processability.21
Doping can strongly alter the nano- and microstructure of conjugated polymers, which can significantly influence electronic charge transport but also alter the mechanical properties of the semiconductor through a range of effects including plasticization, ionic crosslinks, planarization of the conjugated backbone and a change in the degree of order (π-stacking). A further intriguing avenue is the use of doping reactions or counterion exchange for the preparation of organogels, hydrogels or coacervates. For example, conjugated polymers (including conjugated polyelectrolytes) can form networks in solution through ground-state electron transfer112 or via ionic crosslinks that involve multivalent acids such as phytic acid (see Fig. 2 for chemical structure),113 polyelectrolytes (cf. PEDOT:PSS), ionic liquids114 or salts that comprise multivalent ions.115 A recent example involved the formation of a complex composed of an anionic conjugated polyelectrolyte – a polythiophene with both hexyl and hexyl sulfonate side chains – and a cationic bottlebrush polyelectrolyte. Ionic interactions resulted in a coacervate that after drying yielded a soft material with a low E = 0.7 MPa but εbreak = 430%, which upon doping with H-TFSI turned into an elastic conductor with a lower E = 0.2 MPa and εbreak = 94%.116
Fig. 13 Stress–strain curves measured by tensile deformation of free-standing p(g42T-T) films co-processed with 3 mol% F4TCNQ or 6 mol% F2TCNQ, resulting in a similar oxidation level of 5.7 and 6.4%, respectively; Figure adapted with permission from ref. 39; Copyright 2022 (CC-BY), The Royal Society of Chemistry. |
Fig. 14 (a) A tensile-drawn regio-regular P3HT film clamped in a dynamic mechanical analyzer, and DMA thermograms of isotropic P3HT films (b) prior to doping and (c) after sequential doping with Mo(tfd-COCF3)3. Figure adapted with permission from ref. 38; Copyright 2019, American Chemical Society. |
In some cases, co-processing of polymer and dopant, or doped polymer and counterion, can result in mechanically robust materials. For example, PANI can be co-dissolved with camphor or aryl sulfonic acids as well as, optionally, insulating polymers such as poly(methyl methacrylate) (PMMA), polystyrene or polypropylene.21 Direct ink writing of PANI and DBSA123 or dinonylnaphthalene sulfonic acid, polystyrene and fused silica124 has been demonstrated, in the latter case resulting in printed filaments with a storage modulus of about 0.2 GPa at room temperature. PANI and DBSA can also be wet-spun into fibers with a diameter below 5 μm that feature an E = 30 GPa and σbreak = 1080 MPa but εbreak of only 4%.125 Similarly, aqueous dispersions of PEDOT:PSS can be used to wet-spin fibers with a diameter of 5–10 μm, an E of up to 22 GPa, σbreak = 550 MPa, εbreak = 7.5% and an electrical conductivity of 3500 S cm−1.7 Similar values of E = 19.5 GPa and 1600 S cm−1 have recently been reported for wet-spun and drawn PBFDO fibers.9
In the high doping regime materials tend to feature brittle failure with a low εbreak of typically less than 10%. Chemical doping can both increase or decrease the elastic modulus of conjugated polymers depending on how a particular dopant affects the Tg and degree of crystalline order. In case of conjugated polymers with E > 0.1 GPa, chemical doping only results in a relatively minor change in stiffness, which can both decrease or increase depending on the polymer:dopant pair (Fig. 15). For example, oriented PA tapes and poly(2,5-dimethoxy-p-phenylenevinylene) (PDMPV) fibers show a decrease in elastic modulus upon doping, as measured with tensile drawing.126,127 Likewise, ion exchange doping slightly reduces the stiffness of PBTTT films, as measured with AFM and nanoindentation,128 and regio-regular P3HT, measured with DMA, shows a decrease in elastic modulus when doped with, e.g., EBSA or Mo(tfd-COCF3)3, likely due to a plasticization effect as evidenced by a concomitant decrease in Tg.38,121 Instead, doping of P3HT with F4TCNQ or FeCl3 tends to result in an increase in elastic modulus, as measured with AFM and tensile drawing,129,130 respectively, possibly because of an increase in π-stacking.
Fig. 15 Impact of chemical doping on the tensile storage modulus of PDMPV, PA and various polythiophenes with alkyl or oligoether side chains. Blue circles represent the semiconductor in its neat state, while orange circles are doped polymers; data from ref. 38, 39, 84, 118, 121, 126, 127, 130 and 131. |
A significant increase in Tg as well as π-stacking is observed upon chemical doping of conjugated polymers with a lower stiffness. For instance, p(g42T-T) with E = 8 MPa at room temperature can experience a 29-fold increase in Young's modulus to 232 MPa, along with a change in Tg from −43 to 3 °C, when doped with 30 mol% F4TCNQ (measured with a dynamic mechanical analyzer using static tensile deformation; see Section 3.1.2).39 Similarly, the modulus of p(g42T-T) doped with 18 mol% H-TFSI increases 20 times, reaching 164 MPa (Fig. 15).118 Shorter triethylene glycol side chains give rise to a stiffer polymer, p(g32T-T), with E = 76 MPa, which upon doping with 20 mol% F4TCNQ increased to 826 MPa.84
Similar correlations are observed for chemically doped polymers. For example, both the electrical conductivity and Young's modulus of p(g42T-T) increase in tandem upon doping with F4TCNQ.39 Doping increases the charge-carrier density but also induces π-stacking and thus enhances the charge-carrier mobility, both of which lead to an increase in the electrical conductivity (see eqn (11)). At the same time, the increase in π-stacking as well as the introduction of Coulomb interactions between polarons and dopant counterions lead to stronger interactions between molecules. As a result, E is enhanced, which is the energy per volume stored in a material upon elastic deformation (small strains) and can be understood as the resistance to rearrangement of nearby molecules. Other dopants such as H-TFSI induce π-stacking in the low doping regime, while in the high doping regime π-stacking is disrupted, resulting in a breakdown of the correlation between electrical conductivity and Young's modulus,118 which may prove useful for the design of conducting but not overly stiff materials.
Uniaxial orientation tends to result in a strong enhancement of both the electrical conductivity and Young's modulus, which both scale with the degree of orientation. For example, a linear correlation has been reported for oriented P3AT fibers doped with FeCl3,130 PDMPV fibers doped with I2127 and PEDOT:PSS fibers.7 Wet-spun poly(3-octylthiophene) (P3OT) fibers doped with FeCl3 feature values of E ≈ 0.5 GPa and an electrical conductivity of 25 S cm−1, which increase to 2.2 GPa and 180 S cm−1 for a draw down ratio of 5.5.131 The approximately linear correlation between the electrical conductivity and Young's modulus in case of doped fibers is also evident when comparing champion values reported for different conjugated polymers.37
The number of transferred electrons depends on the potential that is applied at the electrode and the speed of electrochemical doping is governed by the drift of ions into the semiconductor. The degree of electrochemical doping can be readily altered, or even reversed, by changing the potential that is applied at the working electrode. The accompanied uptake/expulsion of ions and solvent molecules leads to switchable changes in the volume of the polymer, which is widely used for the design of actuators5 and can be anticipated to lead to electrochemically mutable mechanical properties, which could be exploited for the design of new types of mechatronic devices.
More hydrophilic materials such as PEDOT:PSS and polythiophenes with oligoether side chains are able to take up not only ions but also solvent molecules. Hence, these materials undergo passive swelling, i.e. the uptake of solvent molecules from the electrolyte in the absence of any applied electric field, leading to an increase in volume. As a result, ion-conduction pathways are present, i.e. the electrolyte is in contact with the conjugated polymer throughout its entire volume, which facilitates the ingression of ions and hence oxidation/reduction of the whole film once an electrochemical potential is applied.
In case of PEDOT:PSS, which is initially electrically conducting, the application of a negative potential at the working electrode reduces the conjugated polymer and cations ingress into the material to compensate the charge of PSS counterions (depletion mode). Instead, a positive (negative) working electrode potential leads to oxidation (reduction) of initially neutral polymers, accompanied by ingression of anions (cations) to balance the generated charges (accumulation mode). The ions that enter the material are accompanied by solvent molecules, and the amount of solvent that is taken up depends on, e.g., the anion size and electrolyte133 as well as the ionic strength of the electrolyte.69 This so-called active swelling results in an additional increase in volume, which can exceed 100% in case of polythiophenes with oligoether side chains in combination with an aqueous electrolyte134,135 and 20 to 60% in case of PEDOT:PSS.136,137 The ingression of ions tends to be faster than for apolar polymers with a diffusion coefficient of up to 10−9 cm2 s−1 in case of Cl− in p(g32T-TT).70
Electrochemical doping can significantly alter the nanostructure of the material due to swelling of amorphous domains, as well as the oxidation/reduction of the backbone, leading to, e.g., enhanced (or reduced) π-stacking, expansion of the lamellar stacking distance and changes in texture. For example, the backbone of p(g32T) planarizes upon electrochemical oxidation using an aqueous NaCl electrolyte, which leads to a significant increase in π-stacking that results in a high charge-carrier mobility provided that ordered domains are well connected.138 Some changes are irreversible, such as the transformation of initially densely packed films into an open network structure upon pronounced active swelling.69,139,140
Similar to chemical doping, the elastic modulus of a conjugated polymer is affected by a number of processes, including (1) plasticization by counterions and in particular solvent molecules, (2) stiffening of polymer chains due to oxidation (possibly accompanied by a change in ordering), (3) ionic crosslinks between oxidized polymer chains and counterions, and (4) swelling through the uptake of counterions and solvent molecules (Fig. 16).141 The relative importance of these in part counteracting effects determines how the elastic modulus of a material changes upon electrochemical doping. For example, a polymer that takes up counterions but repels solvent molecules may show an invariant or even enhanced elastic modulus because ionic crosslinking outweighs plasticization. Instead, a polymer that experiences considerable swelling due to the uptake of solvent molecules that accompany counterions is likely to display a decrease in stiffness.
PPy, widely studied as an actuator material, shows a decrease in E from 1 to 0.8 GPa upon oxidation, using aqueous NaPF6 as the electrolyte, which was explained with a plasticization effect due to the ingression of PF6− anions accompanied by solvent (water) molecules.143 Furthermore, the ductility can significantly change with the oxidation level. For instance, PPy electropolymerized in aqueous p-toluenesulfonic acid is brittle in its oxidized state, which was explained with ionic crosslinking between charged polymer chains and counterions, but becomes more ductile with εbreak increasing from 5 to 21% upon electrochemical reduction using aqueous electrolytes such as NaCl with monovalent cations due to plasticization as a result of the ingression of Na+.144 Instead, the material remained brittle upon reduction using an aqueous MgCl2 electrolyte, during which divalent Mg2+ cations enter the material. Evidently, the type of counterion can influence the mechanical properties (cf.Fig. 13; chemical doping with F2TCNQ or F4TCNQ). Plasticization due to active swelling has also been inferred in case of electropolymerized P3HT films, where electrochemical oxidation let to a 15% thickness increase due to the ingression of PF6− anions and a further 48% increase due to solvent swelling (propylene carbonate), overall resulting in a decrease in elastic modulus.145
Other materials such as polythiophenes with oligoether side chains such as p(g32T), p(g42T-T) and p(g32T-TT) tend to experience a more significant volume change ΔV upon oxidation.134,135,146,147 For example, p(g32T) with triethylene glycol side chains turns from a solid material into a gel accompanied by ΔV > 1000% during the first oxidation cycle using aqueous KCl, an increase that is not completely reversible because of the partial retention of counterions and solvent molecules during reduction,134 as well as permanent structural changes. During subsequent oxidation/reduction cycles, p(g32T) undergoes reversible active swelling by ΔV > 200%, which is much larger than passive swelling in the same electrolyte (Fig. 17). A similar polymer with diethylene glycol side chains, p(g22T), only shows a ΔV = 27%, which highlights the importance of sufficiently long side chains.146 In addition to swelling, p(g32T) experiences an increase in π-stacking upon electrochemical oxidation.138 Nevertheless, it can be anticipated that polar polymers such as p(g32T) exhibit a significant reduction in elastic modulus upon electrochemical doping because the large degree of swelling-induced plasticization likely outweighs changes in ordering, consistent with the observed solid-to-gel transition.134 PEDOT:PSS experiences considerable passive swelling in aqueous electrolytes, with the degree of swelling depending on the PSS content,148,149 as well as in non-aqueous solvents such as acetonitrile and methanol.137 The passive swelling ratio of PEDOT:PSS tends to be higher than that of accumulation mode materials (e.g., conjugated polymers without polyelectrolyte) because the protonation/deprotonation of the polyelectrolyte (i.e., PSS) facilitates additional water uptake.150 However, PEDOT:PSS undergoes limited active swelling,150 which suggests that the material will soften once it is brought in contact with an electrolyte, but exhibits only limited further changes in mechanical properties during reduction/oxidation cycles (see Fig. 17).
A high degree of passive and/or active swelling can limit the stability of electrochemical devices because of irreversible structural changes due to the retention of counterions and solvent as well as gradual delamination of device active layers. Hence, materials that experience minimal swelling upon electrochemical doping are highly sought after. One example is the polythiophene poly(3-(6-hydroxy)hexylthiophene) (P3HHT), which undergoes minimal and hence reversible passive and active swelling with ΔV < 10% and therefore recovers its stiffness after each oxidation/reduction cycle.147
The substrate material, which can be rigid, flexible (bendable) or even stretchable, determines the mechanical properties of the device. Thin-film devices on a rigid substrate do not require organic semiconductor layers with any specific mechanical properties, which arguably is the reason why this type of architecture is often selected for screening of new materials in research laboratories. However, the overly use of rigid substrates tends to divert attention from the mechanical properties that are ultimately required for the design of materials for flexible and/or stretchable electronics.
In contrast, if the substrate is non-rigid, the thin-film layer stack must be able to accommodate the same type of deformation that the substrate (surface) experiences upon bending or stretching, without fracture of any of its constituent layers or loss of adhesion between layers. A substrate with a low stiffness can be realized by selecting a material with a low elastic modulus such as a thermoplastic elastomer or other type of rubber, which will also exhibit a high degree of reversible stretchability.
Alternatively, the thickness of the substrate can be reduced to achieve a low stiffness (see eqn (1), Section 3.1). Plastic foils with a thickness of 0.3 to 3 μm have been used for the design of imperceptible electronics composed of conformable and low-weight devices from OFETs to OLEDs and OPV devices, which can be placed on skin or implanted.151 A conjugated polymer layer deposited on top of a low-stiffness substrate will experience compressive and/or tensile stresses upon bending or stretching. The resulting deformation should not exceed εcrack (see Section 3.2), which will result in the loss of integrity of the semiconductor layer, especially upon repeated bending or stretching. Most conjugated polymers feature a low yield strain εyield < 10%,152 which means that they will undergo plastic deformation when deformed beyond this limit, leading to irreversible changes that may negatively affect the device performance. Hence, depending on the application a polymer with a sufficiently high εyield and εcrack must be selected, which may decrease upon blending with acceptor (cf. bulk-heterojunction blends; Section 4.3) or dopant molecules (cf. chemical doping; Section 5), but can also be enhanced through suitable additives such as polymeric binders.
Bulk devices where the organic semiconductor provides both the electrical as well as mechanical performance have been explored in the context of wearable electronics. For instance, conducting polymer tapes and fibers can function as electrical conductors, as actuators, as strain or electrochemical sensors (see Section 7.4), and they can be used as components in energy harvesting (e.g. thermoelectric generators; see Section 7.3) and storage devices (e.g. batteries, supercapacitors).37 Bulk materials must have a thickness of at least several μm so that they can handle the mechanical load without the support of a substrate. At first sight, bulk processing is straightforward because the electrical and mechanical properties of conjugated polymers tend to correlate (cf. Section 5.4). A wide range of conventional polymer processing methods can be readily utilized. Melt processing is often not feasible because of the prohibitively high melting temperatures of many conjugated polymers. Instead, solution processing methods such as wet spinning of fibers37 and 3D printing of gels153 are widely explored. However, processing of bulk materials as well as the operation of thick devices is limited by the rate of mass transport of auxiliary species such as solvent molecules, dopants, counterions, etc. For example, solution processing requires the removal of the processing solvent, which takes time if thick materials are to be created. The impact of the drying kinetics on nanostructure formation is well understood in case of thin films but is more difficult to control when bulk materials are prepared. As a result, there is a tendency for thin films to exhibit superior electrical properties compared with bulk materials.
Deformation of devices on flexible substrates can cause mechanical degradation via adhesive failure between layers and cohesive failure of the active layer,25 and hence the selection of robust materials is critical for ensuring a stable performance.
During the last decade, the synthesis of new types of conjugated polymers has to a significant extent been fueled by the demand for new donor materials for organic photovoltaics. Most donor polymers have a high Tg,76 which is thought to arrest (or at least slow down) phase separation of donor:acceptor bulk-heterojunction blends, resulting in a brittle material with a low εcrack. The addition of the acceptor – a fullerene derivative or a so-called non-fullerene acceptor (NFA) – tends to lead to further embrittlement because acceptors also tend to exhibit a high Tg.82,157 Below the blend Tg (s), decohesion of the bulk-heterojunction active layer occurs via brittle failure, which can be mitigated by selecting a high-molecular weight polymer, as observed for P3HT:[6,6]-phenyl-C61-butyric acid methyl ester (PC61BM) devices.158
A number of approaches have been explored to improve the ductility of bulk-heterojunction blends including plasticizers,25 the addition of small-molecular additives that form an internal network159 and binder materials such as a polystyrene-b-poly(ethylene-ran-butylene)-b-polystyrene (SEBS) block copolymer (see Section 4).160,161 Moreover, the constitution of the conjugated polymer itself can be modified, e.g. through the introduction of a flexible spacer (Fig. 4), which increases the flexibility of the backbone and in the context of all-polymer OPV devices has yielded an εcrack > 20%.162
To create bulk materials with conducting polymers, the processing method must be carefully selected. Co-processing of P3HT and F4TCNQ results in aggregation of the polymer in solution, and thus a brittle solid.121 Instead, millimeter-thick architectures of P3HT can be solid-state pressed, followed by sequential doping with F4TCNQ, which is however ineffective because of diffusion of the dopant is prohibitively low.110 Polymers such as p(g42T-T) that show better compatibility with dopants such as F4TCNQ and H-TFSI can instead be shaped into bulk materials via co-processing from solution.166 One of the most promising p-type materials is PEDOT:PSS, which can be readily processed as an aqueous dispersion and has been utilized for the fabrication of free-standing films167 and fibers7,168 with a very promising thermoelectric performance.
Bulk materials such as silk and cellulose yarns coated with PEDOT:PSS164,169 have been used to fabricate thick textile devices by embroidering the conducting yarn into a wool fabric to create devices with a thickness of about 1 cm. Fused filament fabrication (FFF) 3D printing is another method to create thick out-of-plane devices. For example, a device with 100 leg pairs could be realized by first printing 1.6 mm thick, porous legs of a Nafion precursor on a textile substrate, which were then used as a template for the oxidative polymerization of PEDOT.170 Inkjet printing is being explored as a technique to combine solution processing with patterning of smaller leg pairs. For example devices with a leg thickness of 25 μm have been printed comprising PEDOT:PSS as the p-type material and a doped fullerene derivative as the n-type material.171 Finally, binder materials can be used to enhance the mechanical properties of thermoelectric materials, including semicrystalline polymers such as poly(ethylene oxide) PEO to increase the robustness of F4TCNQ doped P3HT172 and polyurethane to impart stretchability to PEDOT:PSS173,174 or p(g42T-T).175
The active layer of an OECT is composed of an OMIEC, which can be a conjugated polymer that is in contact with a gate electrode via an electrolyte, e.g. a salt dissolved in water or acetonitrile (Fig. 18). Application of a suitable electrical potential Vg at the gate electrode triggers a redox reaction in the polymer. Electronic charge is exchanged with the source or drain electrode and the generated charge is compensated through the exchange of ions with the electrolyte. The OMIEC can initially be a semiconductor or a conductor such as PEDOT:PSS, and the channel conductance is increased (accumulation mode) or decreased (depletion mode) during device operation. In both cases, ions accompanied by solvent molecules enter so that overall charge neutrality is maintained (see Section 6).
Thin-film devices can be deposited on a planar or curved substrate, including filaments and yarns, which makes OECTs ideally suited for textile-based logic circuits.179 Thin active layers (d < 100 nm) are preferred because the conductance of an OECT is altered upon changing Vg. The device switching speed depends on the rate of ion exchange with the electrolyte, which is determined by both the ionic mobility and the thickness of the polymer layer. Hence, very thin layers maximize the switching speed for a given material, an important criterion for circuit design. The substrate can be rigid, flexible or elastic, and the polymer layer must be able to accommodate any imposed deformation. For instance, OECTs on a PDMS substrate based on p(g32T-T) with an Mn = 68 kg mol−1 and E ≈ 50 MPa could be stretched to ε = 100% at least 5000 times without a significant change in device performance, which enabled positioning on skin for real-time recording of electrocardiogram (ECG) signals.180 Instead, PEDOT:PSS devices on a thermoplastic polyurethane substrate could only be stretched to ε = 50%, likely because of the higher stiffness of the active layer.181
OECTs can be deposited on a sacrificial substrate that can be used to position devices directly on skin or tissue, which takes over the role of the substrate.176 A significant part of the overall volume of an OECT is occupied by the electrolyte. Hence, to fabricate fully functional devices, it is important to use a mechanically robust electrolyte, which can be either a liquid (e.g. an aqueous electrolyte) or a solid electrolyte with suitable mechanical properties. The electrolyte can be used as the substrate, as demonstrated for devices on elastic gelatin hydrogel films with E < 1 MPa, on top of which meander-shaped PEDOT:PSS patterns were deposited.182 OECTs where a bulk material functions as both the channel and provides the mechanical integrity must have a thickness of at least a few micrometers. As a result, the rate of ion exchange and hence the switching speed is lower compared to thin-film devices. Ion ingression can be aided by maximizing the contact area between the channel material and the electrolyte through the use of a porous material. For instance, OECTs with a channel thickness of 1 mm could be created by dyeing the internal cell walls of balsa wood with PEDOT:PSS, which enabled ingression of the electrolyte and hence relatively fast switching on the order of a few seconds.183 Another approach is the use of non-planar architectures such as filaments. For example, oriented PEDOT:PSS microfibers with an E of up to 4 GPa have been used as the channel, yielding OECTs with a record device performance thanks to a very high μ of up to 13 cm2 V−1 s−1.149 Alternatively, a hydrophilic reinforcing agent such as CNF can be added to the conjugated polymer, which allows to modulate the mechanical properties and at the same time increases the ionic mobility.79
Many doped organic semiconductors are characterized by poor stability due to degradation reactions with, e.g., water and oxygen. Moreover, unreacted dopants and counterions can diffuse and aggregate within thin-film layer stacks, leach out or sublime from devices and drift in an electric field, resulting in a change in not only the electrical but likely also mechanical properties of doped conjugated polymers. Further studies that explore the stability of doped materials are needed, as well as strategies that mitigate degradation reactions and hinder diffusion.
It can be anticipated that an in-depth understanding of the mechanical properties of (doped) conjugated polymers will enable the design of truly robust and/or elastic (semi)conducting materials, which promises to advance the fields of wearable electronics and bioelectronics. Moreover, the increased use of fatigue testing of conjugated polymers upon not only repeated electrical stress but also deformation is needed to realize materials that exhibit the long-term mechanical properties that are typical for engineering polymers. It can be anticipated that doping will be utilized as a tool to optimize the stiffness and ductility of conjugated polymers and may even allow the introduction of reversible behavior. One further opportunity is the more widespread use of the chemical toolbox developed in the context of thermoplastic elastomers and dynamic networks, which may facilitate the development of materials that not only feature attractive electrical and mechanical properties but can also be reused at the end of their lifetime.
Footnote |
† These authors contributed equally. |
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