On the Glass Transition of Polymer Semiconductors and Its Impact on

Review pubs.acs.org/cm

On the Glass Transition of Polymer Semiconductors and Its Impact on Polymer Solar Cell Stability Christian Müller* Department of Chemistry and Chemical Engineering, Chalmers University of Technology, 41296 Göteborg, Sweden ABSTRACT: The glass transition temperature is a critical processing parameter that governs the kinetics of molecular organization of polymer semiconductors during solidification. Yet, little attention is paid to the resulting structureprocessing-property relationships that lead to optimal optoelectronic performance, which is usually obtained with nonequilibrium nanostructures. This review elucidates the interplay of molecular design and glass transition phenomena that are common to the most well-studied families of conjugated polymers, including polyfluorenes, polythiophenes, and poly(p-phenylenevinylene)s. The influence of key structural factorsknown from classical polymer science such as molecular weight, chain rigidity, side-chain architecture, and intermolecular π−π interactions, is explored in order to provide rationales that can guide the synthesis of new polymer semiconductors with tailored glass transition temperatures. Moreover, the discussion is anchored in an overview of the main measurement techniques with emphasis on rate-dependency and sub-glass transition phenomena as well as differences between bulk and thin films. The second half of the review focuses on the glass transition temperature(s) of polymer:fullerene bulk-heterojunction blends, which represent the most promising active layer architecture for organic solar cells, and highlights the relevance of fullerene diffusion. A challenging perspective is provided with regard to the thermal stability of the blend nanostructure vs the mechanical robustness and ductility of the active layer material. Conflicting demands on the blend glass transition temperature, i.e., higher vs lower than the processing and operating temperature, require a satisfactory compromise that must be achieved before truly flexible polymer solar cells with a high lightharvesting efficiency can be realized.

1. INTRODUCTION Polymer semiconductors attract considerable attention for a wealth of applications ranging from thin-film electronic circuitry to organic light-emitting diodes (OLEDs), solar cells and thermoelectric generators. Solution-processing of these materials permits deposition of thin films on (flexible) substrates by rapid printing or coating techniques, which has particular appeal for cheap, large-area electronics. Typically, the nanostructure of the as-deposited semiconductor material is not in thermodynamic equilibrium and thus tends to evolve with time, which can strongly affect the optoelectronic properties. In order for a polymer to adopt an equilibrium conformation, main-chain structural relaxation must take place, which requires cooperative reorganization of chain segments that is facilitated by translational and in particular rotational modes of motion.1 Such conformational changes are temperature activated, and hence a characteristic temperature exists that marks the onset of main-chain relaxation processes, i.e., the so-called glass transition temperature Tg. At temperatures far below Tg segmental relaxation is prohibited on an experimental time scale and, therefore, the existing chain conformation is “f rozen in”. A prime example is rapid cooling of liquid-crystalline poly(9,9-dioctylfluorene) (F8) thin films on alignment layers, which permits quenching of highly oriented monodomains into © XXXX American Chemical Society

a glassy state for polarized electroluminescence applications (see Figure 1 for the chemical structure of some of the semiconductors discussed in this review).2−4 In contrast, at temperatures above Tg segmental relaxation can take place on an experimental time scale. The time scale that is required in order for this process to occur decreases exponentially with increasing temperature. The associated decrease in viscosity above Tg permits reorganization of polymer chains and diffusion of additives, which in a polymeric material that was initially out-of-equilibrium leads to viscoelastic stress relaxation, phase separation, crystallization, etc. Thus, an intimate understanding of glass transition phenomena is required in order to gain control over the precise nanostructure and hence optoelectronic properties of polymer semiconductors. The remaining part of the introduction will discuss why knowledge of the glass transition temperature is critical for (1) optimization of postdeposition processing protocols, (2) the thermal stability of the semiconductor nanostructure, (3) dewetting and delamination of thin films, and (4) the mechanical properties of polymer Received: January 4, 2015 Revised: March 9, 2015

A

DOI: 10.1021/acs.chemmater.5b00024 Chem. Mater. XXXX, XXX, XXX−XXX

Review

Chemistry of Materials

Figure 1. Chemical structures of some of the organic semiconductors discussed in this review.

semiconductors, which must be considered if flexible optoelectronic applications are to be targeted. 1.1. Optimization of Postdeposition Processing Protocols. Postdeposition thermal annealing is a widely used tool that permits the modification of the nanostructure of polymer semiconductors in a controlled way. Most thermal annealing protocols that are disclosed in the literature require heating above the glass transition temperature of the semiconductor material. As already discussed above, heat treatment is an important step during the fabrication of uniaxially oriented thin films on alignment layers.2−5 Knaapila et al. studied alignment of the liquid-crystalline polyfluorene F2/6 on rubbed polyimide substrates and found that a heating step above Tg ∼ 80 °C is required and that the optimal annealing conditions strongly depend on the polymer molecular weight.6 Similarly,

Zheng et al. were able to achieve a high degree of in-plane alignment through nanoimprinting of a low molecular-weight batch of the fluorene−benzothiadiazole copolymer F8BT above its Tg ∼ 90 °C (number-average molecular weight Mn ∼ 9 kg mol−1).7 Moreover, thermal annealing above Tg can induce the development of ordered phases that are not readily accessible via solution processing. For instance, Lu et al. observed a unique bilayer crystalline structure in the polycarbazole derivative PCDTBT,8 which only developed upon annealing above its Tg ∼ 130 °C.9 In some cases thermal annealing above Tg can be employed to improve the performance of optoelectronic devices such as the electroluminescence from polymer-based OLEDs10 and the charge-carrier mobility in polymer field-effect transistors (OFETs).11 Furthermore, thermal annealing can be used to B


Review


and inferred local changes in nanostructure, which improved the performance of corresponding photovoltaic devices.20 Furthermore, for conjugated polymers sub-Tg relaxation can be light activated. Botiz et al. have shown that limited segmental relaxation of the poly(p-phenylenevinylene) derivative MEHPPV in the form of main-chain planarization can be induced by light illumination below Tg.22 This “light-induced plasticization”, which was first observed for azobenzene functionalized polymers,23,24 can readily occur at temperatures below the Tg that is measured in the absence of light. 1.3. Dewetting and Delamination of Thin Films. Most organic optoelectronic devices comprise a submicrometer thin semiconductor film. Hence, besides evolution of the nanostructure, dewetting of the active layer from the underlying substrate can limit the operational window. This phenomenon is well documented for insulating polymers such as atactic polystyrene (PS), for which dewetting occurs when thin films are heated above Tg ∼ 100 °C.25−28 The tendency for dewetting can be reduced by increasing the polymer molecular weight26 and, in the case of conjugated polymers, by light exposure, which according to Botiz et al. leads to stiffening of polymer chains due to exciton formation.29 Instead, multistacks of thin films on flexible substrates may experience delamination of adjacent layers due to thermomechanical stresses. In the case of P3HT:PCBM solar cells the active layer has been found to display poor adhesion to the often used PEDOT:PSS30 hole-transport layer, which however improves upon thermal annealing.31,32 Though annealing above can also lead to the development of interfacial strain due Tblend g to evolution of the semiconductor nanostructure. For instance, Dupuis et al. have shown that P3HT:PCBM solar cells can degrade during operation at 60 °C as a result of crack formation at the active layer/electrode interfacecaused by P3HT crystallizationthat leads to complete loss in photovoltaic performance.33 1.4. Mechanical Properties of Polymer Semiconductors. Naturally, polymer semiconductors exhibit classical viscoelastic behavior, meaning that they are brittle below Tg but plastic, elastic, or liquid-like at more elevated temperatures, depending on, e.g., molecular weight, entanglement density, crystallinity, and the degree of cross-linking. This strong variation in mechanical response close to Tg is exemplified by regioregular P3HT that displays a glass transition just below room temperature, which however increases with strain rate and may vary for thin films due to confinement effects and/or interactions with the substrate (cf. sections 2.2 and 2.4 below). Thus, at room temperature P3HT can display both ductile as well as brittle behavior. For instance, Koch et al. have demonstrated plastic deformation of free-standing P3HT films by tensile drawing provided that the polymer molecular weight is larger than Mn ∼ 25 kg mol−1, which is needed to ensure elastic percolation, i.e., connectivity of crystalline entities through tie molecules.34,35 Instead, tensile deformation of submicrometer thin P3HT films on an elastic polydimethylsiloxane (PDMS) support can be more ambiguous. The tensile strain that is needed to induce crack formation canlikely due to differences in Tgrange from values as low as 9% up to 150% despite a similar molecular weight, which indicates brittle and more ductile behavior, respectively.36,37 In contrast, for thin films of P3HT:PCBM blends the onset of crack formation during tensile deformation is consistently is situated observed at strains as low as 2% because the Tblend g above room temperature.37,38 Evidently, addition of PCBM to

enhance the power conversion efficiency of polymer solar cells. The most prominent example are solar cells based on bulkheterojunction blends of poly(3-hexylthiophene) (P3HT) and phenyl-C61-butyric acid methyl ester (PCBM), which display a strong tendency to vitrify during solution processing, meaning that a largely disordered nanostructure is frozen in below the blend glass transition temperature Tblend ∼ 40 °C for a 1:1 g stoichiometry.12 However, thermal annealing of P3HT:PCBM above Tblend leads to a significantly enhanced photovoltaic g performance due to crystallization of in particular P3HT, which results in a red-shifted absorption spectrum and improved extraction of photogenerated charges from the solar cell active layer.12−15 1.2. Thermal Stability of the Semiconductor Nanostructure. The thermal stability of a polymer semiconductor material is critical for its long-term use. Gradual evolution of the nanostructure through phase separation and/or crystallization during operation is likely to affect the optoelectronic performance. Polymer:fullerene bulk-heterojunction blends, which typically display a single Tblend due to the fine blend g nanostructure, can undergo rapid coarsening as well as the growth of micrometer-sized fullerene crystals when annealed above Tblend as the system progresses toward thermodynamic g equilibrium (Figure 2).16−20 Moreover, vertical phase separa-

Figure 2. Transmission electron micrographs and electron diffraction patterns (insets) of 1:1 TQ1:PCBM films: (a) pristine, spin-coated, (b and c) annealed at 130 °C, and (d) annealed at 150 °C. Note that ∼ 110 °C. Adapted with permission from ref 19. Copyright 2013 Tblend g The Royal Society of Chemistry.

tion of the polymer can lead to the formation of an electron blocking layer at the electron extracting electrode.21 Phase separation and fullerene crystallization inevitably lead to a detrimental reduction in photovoltaic performance and thus Tblend defines the maximum operating temperature for polymer g solar cells. However, it should be noted that local conformational changes can even occur below Tblend . Recently, Bergqvist g et al. have explored sub-Tblend annealing of a polymer:fullerene g blend based on the thiophene−quinoxaline copolymer TQ1 C


Review


2. EXPERIMENTAL TECHNIQUES FOR THE MEASUREMENT OF RELAXATION PHENOMENA AND THE TG The glass transition of a polymeric material is a second order phase transition, i.e., the volume and enthalpy continuously change with temperature, and can be treated as a kinetic phenomenon.1 Main-chain segmental motion allows a polymer chain to adopt an unstrained conformation. This process is typically referred to as α-relaxation and can be described by a temperature-dependent relaxation time τα, i.e., the time needed for segmental relaxation, which strongly increases with decreasing temperature (cf. eq 1). Then, Tg describes the αtransition temperature, at which τα becomes comparable to the experimental time scale that is available for molecular rearrangements. As a result, the Tg will shift to a higher temperature with, e.g., increasing cooling rate or measurement frequency and must not be treated as a physical constant. Instead, it is important to consider the technique and experimental parameters that were used to determine a particular Tg value as well as the processing history of the sample. It should also be noted that only the amorphous, disordered fraction of a polymeric material displays a glass transition and, hence, the “signal strength” of any T g measurement depends on the polymer crystallinity. 2.1. Broadband Dielectric Spectroscopy (BDS). The perhaps most comprehensive method for the study of relaxation phenomena in polymeric materials is broadband dielectric spectroscopy (BDS), which measures the complex relative dielectric permittivity ε* = ε′ − iε″ (where ε′ is the dielectic constant and ε″ the dielectric loss) as a function of temperature and frequency of the applied electric field, ranging from mHz up to GHz or more.1 BDS can be employed to all insulating and semiconducting polymers and is sensitive to the strength, interactions, and relaxation kinetics of dipolar moieties. Besides the primary α-relaxation of polymer chain segments BDS also permits the study of secondary or sub-Tg processes, i.e., so-called β, γ, etc. relaxation processes, which involve local intramolecular motion. However, the scarce availability of most experimental polymer semiconductor batches complicates the use of BDS, which typically requires samples with a thickness of at least tens of micrometers. Another limitation is oxygen-doping of many conjugated polymers such as P3HT at ambient conditions, which increases the electrical conductivity and complicates BDS analysis. Therefore, only few BDS studies concerning the glass transition of polymer semiconductors exist. Papadopoulos et al. used BDS to study a series of defect-free oligofluorenes as well as a higher molecular-weight polyfluorene with ethyl-hexyl side chains (F2/6).49 Dielectric loss spectra were recorded over a wide range of temperatures and frequencies, ranging from −50 to 180 °C and 3 × 10−3 to 106 Hz, respectively. Then, fitting of the dielectric loss spectra yielded the α-relaxation time τα at maximum loss, which could be described with the empirical Vogel−Fulcher−Tammann (VFT) equation:50

P3HT turns a largely ductile polymer into a brittle material and hence the P3HT:PCBM active layer typically suffers from poor cohesion in polymer solar cells.39,40 A recent review by Savagatrup et al. discusses mechanical degradation of polymer solar cells in detail and argues that device operation below Tblend can lead to debonding, cohesive failure, and cracking g when exposed to large mechanical strains.41 Bruner et al. have carried out a detailed study concerning the temperaturedependent decohesion kinetics of P3HT:PCBM solar cells based on four different P3HT molecular weights ranging from Mn ∼ 16 to 42 kg mol−1.42 Below Tblend ∼ 40 °C decohesion g occurred through brittle failure with higher-Mn P3HT:PCBM devices displaying considerably improved resistance toward decohesion. Instead, above Tblend plastic deformation was g inferred, which aids dissipation of mechanical energy and therefore resists decohesion. Notably, significant plastic deformation was only observed for blends comprising P3HT with a Mn > 25 kg mol−1, which coincides with the molecular weight that is required for chain entanglement and the formation of tie molecules.35 Mechanical degradation during uniaxial and biaxial tensile deformation can be mitigated by increasing the side-chain length, e.g., by selecting poly(3-dodecylthiophene) (P3DDT) as the donor material (Figure 3), which reduces the Tg of the

Figure 3. Photographs of (a) P3HT:PCBM and (b) P3DDT:PCBM photovoltaic devices on an elastic PDMS substrate under a strain of ∼10%. Adapted with permission from ref 37. Copyright 2014 Wiley.

polymer and results in a more ductile bulk-heterojunction blend even at room temperature.37,43,44 Alternative strategies to improve the robustness are (1) deliberate buckling, which occurs when compressive strain is applied to the brittle active layer by relaxation of a prestrained elastomeric substrate,45 and (2) the use of ternary systems of the polymer:fullerene blend and a tough insulating polymer such as polyethylene.46 The remainder of this review serves two purposes. Following a brief overview of the main measurement techniques, the influence of molecular structure on the Tg will be discussed in order to provide rationales that can guide the synthesis of new polymer semiconductors. The second half of the review focuses on the glass transition temperature(s) of bulk-heterojunctions that comprise a donor polymer and a fullerene acceptor. These binary blends represent the most promising active layer architecture for organic solar cells47 and today give rise to a light-harvesting efficiency of more than 10%.48 In particular the thermal stability of the blend nanostructure will be discussed, which is a prerequisite for widespread use of this emerging technology.

⎛ B ⎞ τα = τ0·exp⎜ ⎟ ⎝ T − T0 ⎠

(1)

where τ0, B, and T0 are constants. For F2/6 the authors found that τα decreases from 100 s to 10−4 s when increasing the temperature from 70 to 140 °C (Figure 4). According to common practice an operational glass transition temperature, D


Review


by the large decrease in elastic modulus by about 1 to 2 orders of magnitude from more than 1 GPa to 0.1−0.01 GPa depending on the regioregularity and hence degree of crystalline order. Again, the relaxation processes are frequency dependent: van de Leur et al. observed for P3HT that the peak temperature of the loss tangent shifts from 10 to 30 °C as the frequency of the applied harmonic tensile strain is increased from 0.1 to 10 Hz.55 Clearly, a change in deformation rate can alter the mechanical response of P3HT at room temperature from rubbery to glassy and hence the α-transition value measured at the lowest strain rate should be used as the operational Tg. 2.3. Differential Scanning Calorimetry (DSC). The most straightforward technique to measure the bulk glass transition temperature of polymers is differential scanning calorimetry (DSC). The difference in heat flow to/from a sample as compared to a reference, which arises due to absorption or release of heat during phase transitions such as melting or crystallization, is recorded as a function of time or temperature. Moreover, the change in heat capacity that occurs during the glass transition can be determined from the change in heat flow and serves as a reliable indicator for Tg. In addition, modulatedtemperature DSC (MTDSC) can be employed to more readily detect subtle glass transitions by separating overlying enthalpic relaxation or recrystallization phenomena (cf. ref 42 for an example concerning P3HT:PCBM blends). DSC only requires a few milligrams of material, and, therefore, can be readily carried out with most available polymer semiconductor batches. Thus, most bulk Tg values that are available for conjugated polymers have been measured with DSC using relatively low heating/cooling rates of 5−20 °C min−1. Here, it is interesting to note that the measurement sensitivity increases with the scan rate, which can be exploited to detect weak thermal transitions (cf. Figure 5b and ref 60 for an example of a rapid heat−cool calorimetry study of several cyclopentadithiophene-benzothiadiazole (PCPDTBT) derivatives using a scan rate of 500 °C min−1). A typical DSC measurement is shown in Figure 5, using scans of the liquid-crystalline polyfluorene derivative APFO-3 with a Mn ∼ 5 kg mol−1 as an example. A distinct kink indicates the glass transition and the inflection pointhere at 111 and 105 °C for a scan rate of 10 °C min−1 in the heating and cooling thermogram, respectivelyis typically quoted as the Tg. Note that the DSC signal improves with scan rate. In addition, it can be seen that the glass transition occurs over a broad range of temperatures with an onset significantly below the nominal Tg. 2.4. Variable-Temperature Ellipsometry: Tg of Thin Polymer Films. The above-discussed techniques only provide information about the bulk glass transition temperature. However, most applications of polymer semiconductors concern thin films. For instance, a typical film thickness of an optimized polymer solar cell is on the order of only 100 nm. Variable-temperature ellipsometry, first demonstrated by Beaucage et al. and then others for atactic PS,61−65 is a powerful technique that permits directly probing the thermal behavior of thin polymer films. Ellipsometry provides information about the evolution of the film thickness d with temperature and thus, similar to classical dilatometry, can be used to probe thermal transitions of thin polymer films. Phase changes manifest themselves as a change in volumetric expansion coefficient α, which is directly proportional to variations in the rate of change of the density ρ and hence d with temperature T according to α

Figure 4. α-Relaxation time of F2/6 measured with broadband dielectric spectroscopy (BDS). The data are taken from ref 49: open circles are data measured with BDS; the solid line represents a best fit using the VFT equation and the reported fit parameters −log(τ0/s) = 9.2 ± 0.2, B = 1470 ± 230 K, and T0 = 272 ± 7 K. The arrow indicates the extrapolated operational glass transition temperature Tg = 56 °C, which corresponds to a τα = 100 s.

which corresponds to a τα = 100 s, can be defined, yielding a value of Tg = 56 °C. For all investigated oligomers ranging from 2 to 7 repeat units the operational Tg agreed well with differential scanning calorimetry (DSC) measurements carried out with a heating rate of 10 °C min−1. Moreover, all oligomers displayed a distinct β-relaxation with a typical Arrhenius temperature dependence of relaxation times. A similar BDS study has been carried out for “hairy-rod” poly(p-phenylene)s with a variety of different side chains.51,52 2.2. Dynamic Mechanical Analysis (DMA). Structural relaxation of polymers can also be investigated with dynamic mechanical analysis (DMA), which measures the complex dynamic modulus, i.e., the ratio of an oscillating tensile, compressive, flexural, or shear stress and the resulting strain in a periodically deformed sample, as a function of temperature and frequency.53 The strain must be kept small ( 10 kg mol−1 will be compared in order to avoid any major influence from variations in chain length. 3.2. Constitution of Chain Segments: Influence of Chain Rigidity on Tg. A further important structural factor that raises the Tg is incorporation of rigid units such as cyclic structures into the polymer backbone that resist internal rotation of main chain bonds. The rigidity of a polymer can be expressed in terms of its persistence length lp. McCulloch et al. used neutron scattering to measure the persistence length of a range of P3ATs and found a value of lp ∼ 2.9 ± 0.1 nm for P3HT with a regioregularity >97%.79 In contrast, regiorandom P3HT with lp ∼ 0.9 ± 0.1 nm was found to be significantly more flexible. Available literature data for P3ATs suggest that for a given alkyl side-chain length regiorandom P3AT features an approximately 10 to 20 °C lower Tg as compared to regioregular P3AT (Figure 8), which can be rationalized with an increase in chain flexibility. For instance, several studies have measured a Tg ∼ 12−14 °C for P3HT with a regioregularity of at least 92%42,57,80 and a slightly lower Tg ∼ −3 to 9 °C for regiorandom P3HT.56,57,81,82 The three widely studied conjugated polymers MEH-PPV, F2/6, and F8 have a larger

Figure 7. Molecular-weight dependence of Tg of F2/6 (top) and Tg as well as liquid-crystalline to isotropic transition temperature Tlc→i of APFO-3 (bottom), measured with BDS + DSC and DSC, respectively. Figures adapted with permission from ref 49 (copyright 2004 The American Institute of Physics (top)) and from ref 77 (copyright 2010, Wiley (bottom)).

Flory−Fox equation provides an empirical description for the increase in Tg with chain length and is given by:76 K Tg(M n) = Tg(M∞) − Mn (2)

Figure 8. Transition temperatures of poly(3-alkylthiophene)s (P3AT)s as a function of alkyl side-chain length: α-transition from DSC or TMDSC ( ● regioregularity ≥92%; ○ regiorandom);42,54,56,80−82,93,94 α-transition from DMA (◆ regioregularity >97%; ◇ regiorandom) and β-transition from DMA (▲ regioregularity >97%; △ regiorandom).56,57 DMA data correspond to loss modulus peak values measured at ∼1 Hz; for PT (m = 0) the loss tangent tan δ peak value is used. Note that several reports that are not included in this analysis quote widely different values for P3HT; e.g., Tg ∼ −16 to 24 °C (DMA peak loss modulus at 1 Hz; regioregularity ∼82%)95 and Tg ∼ 35 °C (broad DMA tan δ peak value at 1 Hz; regioregularity >98%).59

where Tg(M∞) is the glass transition temperature of a polymer with infinite molecular weight and K is a polymer-specific constant. More recently, the same trend was observed for APFO-3 (Figure 7).77 In both cases, the Tg rapidly saturates with increasing molecular weight and, accordingly, the Tg of polymers with a Mn > 10 kg mol−1 can be considered less dependent on chain length. Nevertheless, a gradual increase still G


Review


Table 1. Glass Transition Temperature Tg of Fluorene-Based Polymers as Well as the Corresponding Number-Average Molecular Weight Mn and Polydispersity Index PDI, for Which These Values Have Been Reporteda

a

For references that study several molecular weights, the Tg of the highest Mn is quoted.

persistence length than P3HT with values of lp ∼ 6 nm, lp ∼ 7 nm, and lp ∼ 8.5 nm, respectively.83−85 As a result, these polymers feature a higher glass transition temperature than P3HT with values of Tg ∼ 65−66 °C,10,82,86 Tg ∼ 56−80 °C,5,49,87 and Tg ∼ 65−72 °C.88,89 Another example is given by the two polyfluorene derivatives F8T2 and F8BT (Table 1). F8T2 contains, besides the fluorene unit, two relatively flexible thiophene units and features a glass transition temperature of Tg ∼ 73−109 °C.90−92 Instead, inclusion of the more rigid benzothiadiazole unit, which also has a larger rotational volume

(cf. section 3.2 below), results in a much higher Tg ∼ 135 °C for F8BT. 3.3. Constitution of a Single Repeat Unit: Influence of the Length and Bulkiness of Pendant Groups on Tg. Conjugated polymers typically carry alkyl side chains, which provide the good solubility in organic solvents that is required for solution processing. In accordance with insulating polymers such as poly(n-alkyl acrylate)s and poly(n-alkyl methacrylate)s the side-chain length and grafting density strongly influence the Tg. Here, P3ATs represent the most thoroughly studied system H


Review


Table 2. Glass Transition Temperature Tg of PPV Derivatives as Well as the Corresponding Number-Average Molecular Weight Mn and Polydispersity Index PDI, for Which These Values Have Been Reported

relatively high Tg ∼ 100 °C despite a low lp ∼ 1 nm (note that replacement of the phenyl with much smaller methyl groups results in polypropylene with a much lower Tg ∼ −10 °C).53 This trend is exemplified by a series of three fluorene copolymers comprising APFO-3, APFO-18, and APFOGreen9 (Table 1). APFO-3 features a high Tg ∼ 123 °C.77,98 Replacement of the benzothiadiazole with a diphenyl-quinoxaline unit yields APFO-18, which has a higher Tg ∼ 142 °C due to the two pendant phenyl rings.99 APFO-Green9, which carries four pendant phenyl rings, displays the highest Tg ∼ 177−192 °C.72,99 A further structural motif that increases the rotational volume is a kinked polymer backbone.4 For instance, the kinked polymer backbone of TFB is consistent with a high Tg ∼ 140 °C.71 For a TFB derivative that instead carries a straight butyl side chain on the triphenylamine unit an even higher Tg ∼ 156 °C has been reported.100 Segmental mobility can also be restricted through addition of a bulky pendant group onto the end of a flexible alkyl side chain. This behavior is illustrated by the PPV derivative MPEPPV, which carries a phenyl-ethyloxy side chain and thus features a high Tg ∼ 111 °C (Table 2).97,101 This value can be directly compared to the much lower Tg ∼ 65−66 °C of MEHPPV, which features a less bulky ethyl-hexyloxy side chain with a similar molar mass.10,82,86 3.4. Influence of π−π Interactions between Polymer Chain Segments on Tg. Interchain interactions can strongly influence the Tg since they restrict the molecular mobility of polymer chain segments. Wang et al. used variable-temperature ellipsometry to measure the glass transition temperature of PCDTBT and PCDTBT:fullerene thin films and observed a reduction by about 10−15 °C upon thermal annealing.70 PCDTBT was found to be largely amorphous but exhibited π−π stacking between the conjugated backbones with an average interchain distance of 4 Å. A decrease in coherence length from 10 to 8.5 Å indicated a decrease in the degree of interchain interaction, which appeared to correlate with the observed reduction in glass transition temperature. Overall, the Tg ∼ 120−140 °C of PCDTBT with a Mn ∼ 22 kg mol−1 was

(Figure 8).42,54−57,80−82,93−95 Unsubstituted polythiophene (PT) is intractable in organic solvents and features a Tg ∼ 120 °C as measured by Chen and Ni with DMA at a frequency of 1 Hz.54 The Tg, i.e., the α-transition temperature, strongly decreases with increasing alkyl side-chain length, reaches values just below room temperature for P3HT and drops below 0 °C for P3OT. Any further decrease in Tg for longer side chains, i.e., decyl and dodecyl, is less pronounced. Moreover, as discussed in section 3.2, the Tg appears to only weakly depend on the degree of regioregularity for a given side-chain length. Instead, for the β-transition temperature, which is due to alkyl sidechain relaxation, a markedly different trend can be deduced. For P3HT, Pankaj et al. have measured a value around −100 °C using DMA at a frequency of ∼1.6 Hz, which increases to −40 °C for P3DDT.56,57 The strong influence of side-chain length on Tg has been observed for a number of other conjugated polymers. For instance, unsubstituted poly(p-phenylenevinylene) (PPV) has a Tg ∼ 220 °C.96 Instead, the two most common derivatives MEH-PPV and MDMO-PPV feature a significantly lower Tg because they carry an ethyl-hexyloxy and a longer dimethyloctyloxy side chain, respectively (Table 2). Due to its slightly shorter side chain MEH-PPV displays a somewhat higher Tg ∼ 65−66 °C10,82,86 as compared to MDMO-PPV with a Tg ∼ 18− 50 °C.82,97 Furthermore, Teetsov and Fox investigated the thermal behavior of the polyfluorenes F6, F8, and F12 with comparable molecular weight and observed a decrease from Tg ∼ 94 to 72 °C and 47 °C with increasing dialkyl side-chain length.89 In contrast, the use of branched side chains appears to have a minor effect on the glass transition temperature as indicated by the similar range of values reported for F8 and F2/ 6 with Tg ∼ 65−72 °C88,89 and Tg ∼ 56−80 °C,5,49,87 respectively (Table 1). Segmental relaxation is hindered by adjacent polymer chains. Thus, the glass transition temperature scales with the volume that a chain segment requires for rotation around the backbone axis. As a result, heavy pendant units such as phenyl groups strongly raise the Tg, which is the reason why atactic PS has a I


Review


glass.15 According to several DSC and DMA studies regioregular P3HT has a Tg ∼ 12−14 °C (cf. Figure 8). In contrast, disordered PCBM features a much higher Tg ∼ 110− 150 °C, which has been measured with variable-temperature ellipsometry,18 grazing-angle X-ray scattering,106 DSC,42,80,107 and DMA. 59 A number of studies have investigated P3HT:PCBM blends across the full composition range and found a single Tblend with values that lie in between those of the g neat components (Figure 9).59,80,82,97,107,108 Evidently, addition

found to be influenced by thermal annealing as well as the casting solvent, i.e., chloroform or o-dichlorobenzene, and the underlying substrate, i.e., PEDOT:PSS or Si/SiO2. Evidently, although classical polymer science rationales can be employed to explain the glass transition of polymer semiconductors, special care must be taken to account for effects such as intermolecular interactions in the form of π−π stacking, which more readily occurs in conjugated materials.

4. GLASS TRANSITION TEMPERATURE AND THERMAL STABILITY OF POLYMER:FULLERENE BLENDS The most promising active layer material for organic solar cells is composed of bulk-heterojunction blends of an electrondonating (hole-conducting) conjugated polymer and an electron-accepting (electron-conducting) fullerene derivative. Typically, optimized blends feature phase separation on the nanoscale, which provides (i) a large donor/acceptor contact area, needed for charge generation, as well as (ii) percolating and mostly phase-pure domains, required for charge extraction. Solution processing of polymer:fullerene blends tends to result in nonequilibrium nanostructures. The high glass transition temperature of donor polymers and fullerene acceptors leads to f reezing-in of mixed phases that are largely disordered. As a result, solution-cast polymer:fullerene blends typically only display a single glass transition temperature, which is associated with the mixed polymer:fullerene phase. Upon annealing above the blend glass transition temperature Tblend rapid coarsening g can occur, which is detrimental for the photovoltaic performance. In addition, ordered polymer as well as fullerene domains can develop, which however do not feature a glass transition since the latter is associated with an amorphous, disordered state. Because of the nonequilibrium character of polymer:fullerene bulk-heterojunction blends, Tblend is a critical parameter for the g thermal stability of polymer solar cells, which must be ensured during both device fabrication as well as operation. One critical aspect of large-area, continuous solar cell fabrication is the coating speed, which must be maximized in order to enable an economic process. Therefore, high-throughput roll-to-roll coating processes require several heating cycles to ensure rapid solvent removal during deposition of each layer. Thus, the bulk-heterojunction blend nanostructure must be able to withstand repeated heating without coarsening or degradation. Since the cost of the substrate constitutes a significant fraction of the total material cost, polyethylene terephthalate (PET) foil is typically chosen. PET foil starts to deform at temperatures above 140 °C, which hence is the highest possible processing temperature.32,103 A second critical temperature is the operating temperature of a solar cell, which is likely to heat up considerably when exposed to full sunlight but may also experience very low temperatures in, e.g., a desert at night or arctic countries such as Sweden. Here, industrial standards that have been developed for inorganic solar cells provide a useful reference temperature: in order to certify a solar cell as stable, the device must be able to withstand thermal cycling between −40 and 85 °C.104 In 2011, the organic photovoltaic research community has published a consensus on stability testing, which proposes to use the same temperature range for organic solar cells.105 4.1. Glass Transition of Semicrystalline P3HT:PCBM Blends. The most widely studied polymer:fullerene photovoltaic blend comprises regioregular P3HT and PCBM, which has the strong tendency to vitrify, i.e., it forms a molecular

Figure 9. (a) Glass transition temperatures of P3HT:PCBM blends as a function of stoichiometry measured with DSC by Bruner et al. (□),42 Ngo et al. (▼),107 Kim and Frisbie (△),82 and Zhao et al. (●),80,97 and with DMA by Hopkinson et al. (◇),59 as well as Tblend predicted g by the Fox equation (−) based on a Tg ∼ 12 °C for P3HT and Tg ∼ 131 °C for PCBM; (b) open-circuit voltage Voc, short-circuit current Jsc, and fill factor FF of a 1:1 P3HT:PCBM solar cell measured in situ as a function of annealing temperature. Figure 9b adapted from ref 12. Copyright 2011 The Royal Society of Chemistry.

of PCBM to the polymer raises the glass transition temperature and hence of the resulting blend. A further increase in Tblend g thermal stability of the blend nanostructure can be anticipated if the fullerene derivative phenyl-C71-butyric acid methyl ester (PC71BM) is used instead, which has an even higher Tg ∼ 163 °C.109 The observation of a single Tblend is consistent with studies by g Treat et al. and Leman et al., who have deduced the existence of a mixed, amorphous P3HT:PCBM phase.109,110 In contrast, a careful variable-temperature ellipsometry study by Pearson et al. revealed two glass transition temperatures at 40−60 °C and 70−80 °C for intermediate P3HT:PCBM stoichiometries, which indicates that two mixed phases with different compositions can exist.108 It is important to note that, besides the disordered, amorphous phase(s), additional ordered, crystalline phases exist, which however do not display a glass transition. J


Review


Figure 10. Transmission electron micrographs and electron diffraction patterns (insets) of 1:4 MDMO-PPV:PCBM (top) and 1:4 high-Tg PPV:PCBM active layers (bottom) annealed for 0−16 h at 110 °C. Note the formation of fullerene clusters in MDMO-PPV:PCBM, whereas high-Tg PPV:PCBM displays a more stable nanostructure. Adapted with permission from ref 17. Copyright 2008 Elsevier.

compared two polymer:PCBM blends based on (1) MDMOPPV with a Tg ∼ 50 °C and (2) a random copolymer based on three PPV monomers with a high Tg ∼ 138 °C.16,17 Annealing at 110 °C resulted in a rapid loss of performance in the case of MDMO-PPV:PCBM solar cells whereas devices based on highTg PPV:PCBM were considerably more stable, which coincided with a strongly reduced rate of fullerene crystal formation (Figure 10). Similar results have been obtained for polymer:fullerene blends based on MPE-PPV, which has a high Tg ∼ 111 °C,101 as well as APFO-3 and TQ1,18,20 which both display liquid-crystalline order at best77,114 and feature a high Tg ∼ 123 °C and Tg ∼ 100 °C, respectively.77,115 The glass transition temperature of APFO-3:PCBM blends was investigated with variable-temperature ellipsometry. Due to the was found to be similar Tg of the two blend components, Tblend g largely independent of stoichiometry.18 This composition invariance of Tblend was also reported for blends comprising g the structurally similar PCDTBT, which features a slightly higher Tg ∼ 130 °C (cf. Figure 6).70 Besides the stoichiometry, the glass transition of bulk-heterojunction blends depends on the molecular weight of the donor polymer. 1:4 APFO3:PCBM thin films that comprise a low molecular-weight ∼ 98 °C.18 APFO-3 batch with Mn ∼ 5 kg mol−1 display a Tblend g Annealing at temperatures above 70 °C resulted in a rapid loss of photovoltaic performance. Instead, the use of a higher molecular-weight APFO-3 with Mn ∼ 36 kg mol−1 yielded a ∼ 104 °C, which resulted in good thermal slightly higher Tblend g stability of the blend nanostructure up to 90 °C. For noncrystalline polymer:PCBM blends the decrease in photovoltaic performance upon annealing above T gblend correlates with changes in the blend nano- and microstructure, i.e., coarsening and the growth of micrometer-sized fullerene crystals. A detailed study concerning the formation of fullerene crystals was carried out for TQ1:PCBM thin films, which revealed that above Tblend ∼ 110 °C this blend suffers from a g very low nucleation rate of not more than a few hundred fullerene crystals mm−2 s−1 but a high growth rate of e.g., 20 nm s−1 at 170 °C.19 Therefore, besides detrimental coarsening of amorphous, PCBM-rich domains, annealing above Tgblend inevitably resulted in the growth of millimeter-sized fullerene crystals (Figure 2), which led to a rapid decrease in photovoltaic performance (Figure 11).20 It is important to

Several relationships have been proposed that predict the glass transition temperature of miscible polymer blends,111 which generally display a single Tblend . The empirical Fox g equation is particularly appealing due to its simplicity:112 1 Tgblend

=

w1 w + 2 Tg1 Tg2

(3)

where w1 and w2 are the weight fractions and Tg1 and Tg2 the glass transition temperatures of the two blend components. When applied to the P3HT:PCBM blend system the Fox equation accurately describes Tblend of P3HT:PCBM blends g that are rich in either component but fails for compositions of 40 to 70 wt % P3HT (Figure 9). Note that this simple analysis neglects the presence of crystalline domains, which is likely to affect the composition of the miscible blend fraction. Evidently, although helpful for anticipating general trends, the Fox equation should be used with caution since deviations are likely to occur. Annealing of as-cast P3HT:PCBM above the glass transition temperature, which encourages in particular crystallization of P3HT, is a necessary step for the optimization of the blend nanostructure. Treat et al. have monitored the photovoltaic performance of P3HT:PCBM solar cells in situ and found that the short-circuit current starts to strongly increase upon annealing above 50 °C, which lies just above the glass transition temperature of the blend.12 Gratifyingly, the formation of P3HT crystallites impedes fullerene diffusion and hence improves the thermal stability of the blend even at temperatures above Tblend .13 g 4.2. Thermal Stability of Noncrystalline Polymer:Fullerene Blends. Many high-performance donor polymers that are currently investigated lack the ability to crystallize. When these materials are used, control of the blend nanostructure is particularly challenging since the degree of phase separation rapidly and continuously increases during heating above Tblend . This is because above Tblend fullerene g g molecules diffuse in the polymer matrix, slowly aggregate, and eventually form micrometer-sized crystals, which are detrimental for the performance of polymer solar cells. Yang et al. demonstrated that the thermal stability of MDMO-PPV:PCBM bulk-heterojunction blends is governed by their glass transition temperature.113 This was confirmed by Bertho et al., who K


Review


Figure 11. Short-circuit current density Jsc of 1:1 TQ1:PCBM solar cells as a function of annealing temperature Tanneal. Figure adapted with permission from ref 20. Copyright 2014 The Royal Society of Chemistry.

Figure 12. Relative diffusion coefficient D(T)/D(Tblend ) of 1 vol % g deuterated dPCBM in P3HT (●)119 and of PCBM in 1:1 TQ1:PCBM thin films (◇)19 as a function of annealing temperature relative to ∼ 50 °C and Tblend ∼ 110 °C, respectively. The here used Tblend Tblend g g g values correspond to the onset of detectable structural changes (cf. refs 19 and 119; Figure 9b). The solid line represents a fit based on a diffusion-limited growth model of PCBM crystals in TQ1:PCBM according to the Stokes−Einstein relation Drel = kT·(6πaη)−1 where k is the Boltzmann constant, a is the particle size, and η is the viscosity, which is given by the VFT equation.19

consider that all here discussed polymer:fullerene blends feature a Tblend above room temperature. In spite of being g beneficial for the long-term stability of the blend nanostructure, a high Tblend is undesirable for flexible solar cells, which requires g a more ductile or even elastic material that can accommodate large strains (cf. Section 1.4).41 Although below Tgblend extended molecular motion is suppressed, physical aging can occur over an extended period of time. Solution-cast bulk-heterojunction blends typically adopt a nanostructure far away from thermodynamic equilibrium and rapid solvent removal leads to an excess of free volume, which promotes gradual aging of the blend nanostructure. For instance, sub-Tblend annealing of spin-coated g TQ1:PCBM blends at temperatures as low as 40 °C, i.e., as much as 70 °C below Tblend , resulted in red-shifted UV−vis g absorption and photoluminescence spectra, which was rationalized with local segmental relaxation of the donor polymer that is most likely accompanied by short-range diffusion of PCBM.20 4.3. Fullerene Diffusion in Bulk-Heterojunction Blends. Fullerene diffusion determines the rate of coarsening of a polymer:fullerene bulk-heterojunction blend19,116,117 because of the higher molecular mobility of the acceptor compared to (entangled) donor polymer chains. Watts et al. have studied the depletion zone that typically surrounds PCBM crystals in thermally annealed P3HT:PCBM thin films.118 Diffusion of PCBM molecules to the crystal growth front typically results in a lateral concentration profile, which could be quantified with scanning transmission X-ray microscopy (STXM). The measured concentration profile was fitted with Fick’s diffusion law, yielding an estimate of the diffusion coefficient D ∼ 2.5 × 10−14 m2 s−1 at 140 °C for PCBM in a saturated blend, i.e., a blend that contains an excess of PCBM, which thus tends to phase-separate. Above Tblend the PCBM g volume fraction of about 20% that remains close to the crystal edge was associated with the PCBM fraction that is soluble in (the amorphous fraction of) P3HT at 140 °C. In another study, Treat et al. probed the lateral diffusion of low volume fractions of deuterated fullerene (dPCBM) in a regioregular P3HT/ dPCBM bilayer−monolayer architecture and found that the diffusion coefficient varies by almost 2 orders of magnitude between 50 and 110 °C and increases with temperature from about 2.2 × 10−15 m2 s−1 to 1 × 10−13 m2 s−1 for a dPCBM volume fraction of 1% (Figure 12).119 A strong temperature dependence of the fullerene diffusion coefficient can also be

expected for other polymer:fullerene blends. For instance, Lindqvist et al. examined the growth rate of fullerene crystals in 1:1 TQ1:PCBM thin films across a wide range of temperatures from 110 to 230 °C, which could be described with a diffusionlimited growth model.19 According to the Stokes−Einstein relation the diffusion coefficient D is inversely proportional to the viscosity η. Since η ∝ τα, according to the VFT equation D exponentially increases above Tblend (Figure 12). As a result, the g growth rate of fullerene crystals in TQ1:PCBM thin films displays an exponential temperature dependence above the relatively high Tblend ∼ 110 °C.19 Moreover, it is interesting to g note that the increase in the relative diffusion coefficient D(T)/ D(Tblend ) is comparable for TQ1:PCBM and P3HT:dPCBM; g e.g., for both blends the diffusion coefficient increases by about one order of magnitude 40 °C above Tblend (cf. Figure 12). g Hence, other bulk-heterojunction blends can be anticipated to behave in a similar way. Several strategies have been proposed that permit a reduction of the rate of fullerene diffusion and thus enhance the thermal stability. Fischer et al. used an ingenious setup based on photoluminescence quenching that allowed probing the vertical diffusion of C60 through a F2/6 thin film.120 Cross-linking of F2/6 resulted in a decrease in diffusion coefficient by up to 3 orders of magnitude from D ∼ 10−15 m2 s−1 to D ∼ 10−18 m2 s−1 at 140 °C. Since the glass transition temperature increases with the degree of cross-linking, this approach is well suited to enhancing the thermal stability of polymer solar cell blends.121−123 Other strategies that are likely to raise Tblend g and hence impede fullerene diffusion include polymer sidechain removal by thermal cleavage,124,125 which again raises the polymer Tg (cf. Figure 8), cross-linking of fullerenes,126−128 and light-induced dimerization of fullerenes,129,130 but also the use of PC71BM as the acceptor material because of its high Tg ∼ 163 °C. Moreover, routes that enhance the thermal stability of bulk-heterojunction blends without sacrificing ductility are desirable, which rules out the use of high-Tg semiconductors. Instead, excessive coarsening of the blend nanostructure can be prevented at temperatures above Tblend through (1) controlled g fullerene crystallization with nucleating agents such as C60, L


Review


(9) Blouin, N.; Michaud, A.; Gendron, D.; Wakim, S.; Blair, E.; Neagu-Plesu, R.; Belletête, M.; Durocher, G.; Tao, Y.; Leclerc, M. J. Am. Chem. Soc. 2008, 130, 732. (10) Lee, T.-W.; Park, O. O. Adv. Mater. 2000, 12, 801. (11) Donley, C. L.; Zaumseil, J.; Andreasen, J. W.; Nielsen, M. M.; Sirringhaus, H.; Friend, R. H.; Kim, J. S. J. Am. Chem. Soc. 2005, 127, 12890. (12) Treat, N. D.; Shuttle, C. G.; Toney, M. F.; Hawker, C. J.; Chabinyc, M. L. J. Mater. Chem. 2011, 21, 15224. (13) Ma, W. L.; Yang, C. Y.; Gong, X.; Lee, K.; Heeger, A. J. Adv. Funct. Mater. 2005, 15, 1617. (14) Li, G.; Shrotriya, V.; Huang, J. S.; Yao, Y.; Moriarty, T.; Emery, K.; Yang, Y. Nat. Mater. 2005, 4, 864. (15) Müller, C.; Ferenczi, T. A. M.; Campoy-Quiles, M.; Frost, J. M.; Bradley, D. D. C.; Smith, P.; Stingelin-Stutzmann, N.; Nelson, J. Adv. Mater. 2008, 20, 3510. (16) Bertho, S.; Haeldermans, I.; Swinnen, A.; Moons, W.; Martens, T.; Lutsen, L.; Vanderzande, D.; Manca, J.; Senes, A.; Bonfiglio, A. Sol. Energy Mater. Sol. Cells 2007, 91, 385. (17) Bertho, S.; Janssen, G.; Cleij, T. J.; Conings, B.; Moons, W.; Gadisa, A.; D’Haen, J.; Goovaerts, E.; Lutsen, L.; Manca, J.; Vanderzande, D. Sol. Energy Mater. Sol. Cells 2008, 92, 753. (18) Müller, C.; Bergqvist, J.; Vandewal, K.; Tvingstedt, K.; Anselmo, A. S.; Magnusson, R.; Alsonso, M. I.; Moons, E.; Arwin, H.; CampoyQuiles, M.; Inganäs, O. J. Mater. Chem. 2011, 21, 10676. (19) Lindqvist, C.; Sanz-Velasco, A.; Wang, E. G.; Bäcke, O.; Gustafsson, S.; Olsson, E.; Andersson, M. R.; Müller, C. J. Mater. Chem. A 2013, 1, 7174. (20) Bergqvist, J.; Lindqvist, C.; Bäcke, O.; Ma, Z.; Tang, Z.; Tress, W.; Gustafsson, S.; Wang, E.; Olsson, E.; Andersson, M. R.; Inganäs, O.; Müller, C. J. Mater. Chem. A 2014, 2, 6146. (21) Sachs-Quintana, I. T.; Heumüller, T.; Mateker, W. R.; Orozco, D. E.; Cheacharoen, R.; Sweetnam, S.; Brabec, C. J.; McGehee, M. D. Adv. Funct. Mater. 2014, 24, 3978. (22) Botiz, I.; Freyberg, P.; Leordean, C.; Gabudean, A. M.; Astilean, S.; Yang, A. C. M.; Stingelin, N. ACS Appl. Mater. Interfaces 2014, 6, 4974. (23) Kumar, J.; Li, L.; Jiang, X. L.; Kim, D. Y.; Lee, T. S.; Tripathy, S. Appl. Phys. Lett. 1998, 72, 2096. (24) Karageorgiev, P.; Neher, D.; Schulz, B.; Stiller, B.; Pietsch, U.; Giersig, M.; Brehmer, L. Nat. Mater. 2005, 4, 699. (25) Reiter, G. Phys. Rev. Lett. 1992, 68, 75. (26) Reiter, G. Europhys. Lett. 1993, 23, 579. (27) Reiter, G. Langmuir 1993, 9, 1344. (28) Jacobs, K.; Herminghaus, S.; Mecke, K. R. Langmuir 1998, 14, 965. (29) Botiz, I.; Freyberg, P.; Stingelin, N.; Yang, A. C.-M.; Reiter, G. Macromolecules 2013, 46, 2352−2356. (30) Poly(3,4-ethylenedioxythiophene):poly(4-styrenesulfonate). (31) Dupont, S. R.; Oliver, M.; Krebs, F. C.; Dauskardt, R. H. Sol. Energy Mater. Sol. Cells 2012, 97, 171. (32) Jorgensen, M.; Norrman, K.; Gevorgyan, S. A.; Tromholt, T.; Andreasen, B.; Krebs, F. C. Adv. Mater. 2012, 24, 580. (33) Dupuis, A.; Tournebize, A.; Bussiere, P. O.; Rivaton, A.; Gardette, J. L. Eur. Phys. J. Appl. Phys. 2011, 56, 34104. (34) Moulton, J.; Smith, P. Polymer 1992, 33, 2340. (35) Koch, F. P. V.; Rivnay, J.; Foster, S.; Müller, C.; Downing, J. M.; Buchaca-Domingo, E.; Westacott, P.; Yu, L. Y.; Yuan, M. J.; Baklar, M.; Fei, Z. P.; Luscombe, C.; McLachlan, M. A.; Heeney, M.; Rumbles, G.; Silva, C.; Salleo, A.; Nelson, J.; Smith, P.; Stingelin, N. Prog. Polym. Sci. 2013, 38, 1978. (36) O’Connor, B.; Chan, E. P.; Chan, C.; Conrad, B. R.; Richter, L. J.; Kline, R. J.; Heeney, M.; McCulloch, I.; Soles, C. L.; DeLongchamp, D. M. ACS Nano 2010, 4, 7538. (37) Savagatrup, S.; Makaram, A. S.; Burke, D. J.; Lipomi, D. J. Adv. Funct. Mater. 2014, 24, 1169. (38) Awartani, O.; Lemanski, B. I.; Ro, H. W.; Richter, L. J.; DeLongchamp, D. M.; O’Connor, B. T. Adv. Energy Mater. 2013, 3, 399.

which results in the formation of a metastable nanostructure composed of fullerene nanocrystals embedded in a polymerrich phase,131,132 or (2) the use of suitable fullerene mixtures, which hinders crystallization of the acceptor material.133−137

5. CONCLUSIONS The glass transition temperature of polymer semiconductors critically impacts (1) the choice of thermal annealing protocols, (2) the thermal stability of the semiconductor nanostructure, (3) dewetting and delamination of thin films, and (4) the mechanical properties. The same structure−property relationships that apply to insulating polymers can be used to rationalize the thermal behavior of polymer semiconductors. It can be anticipated that the here discussed design principles will guide the synthesis of new conjugated polymers, which feature a glass transition temperature that is optimized for particular optoelectronic applications. Future work concerning the glass transition temperature of polymer semiconductors should elucidate the influence of parameters such as the entanglement density, crystallinity, degree of cross-linking, and intramolecular forces such as π−π stacking, which more readily occur in conjugated materials, as well as interaction with light and doping, which have received little attention. To realize truly flexible solar cells based on polymer:fullerene bulk-heterojunction blends a compromise must be struck between (1) adequate mechanical robustness and ductility, which requires that the glass transition temperature of the blend is situated below the processing as well as operating temperature, and (2) a high degree of thermal stability of the bulk-heterojunction nanostructure, which is facilitated by a high glass transition temperature of the donor polymer.

■

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Notes

The author declares no competing financial interest.

■

ACKNOWLEDGMENTS Financial support from the Swedish Research Council and the Chalmers Area of Advance Energy is gratefully acknowledged. Table of Contents and Abstract graphic adapted with permission from ref 19. Copyright 2013 The Royal Society of Chemistry.

■

REFERENCES

(1) Strobl, G. The Physics of Polymers, 3rd ed.; Springer Verlag: Berlin, Heidelberg, 2007. (2) Grell, M.; Bradley, D. D. C.; Inbasekaran, M.; Woo, E. P. Adv. Mater. 1997, 9, 798. (3) Whitehead, K. S.; Grell, M.; Bradley, D. D. C.; Jandke, M.; Strohriegl, P. Appl. Phys. Lett. 2000, 76, 2946. (4) Donald, A. M.; Windle, A. H.; Hanna, S. Liquid Crystalline Polymers, 2nd ed.; Cambridge University Press: Cambridge, 2005. (5) Knaapila, M.; Stepanyan, R.; Torkkeli, M.; Lyons, B. P.; Ikonen, T. P.; Almasy, L.; Foreman, J. P.; Serimaa, R.; Guntner, R.; Scherf, U.; Monkman, A. P. Phys. Rev. E 2005, 71, 041802. (6) Knaapila, M.; Stepanyan, R.; Lyons, B. P.; Torkkeli, M.; Hase, T. P. A.; Serimaa, R.; Guntner, R.; Seeck, O. H.; Scherf, U.; Monkman, A. P. Macromolecules 2005, 38, 2744. (7) Zheng, Z.; Yim, K. H.; Saifullah, M. S.; Welland, M. E.; Friend, R. H.; Kim, J. S.; Huck, W. T. Nano Lett. 2007, 7, 987. (8) Lu, X. H.; Hlaing, H.; Germack, D. S.; Peet, J.; Jo, W. H.; Andrienko, D.; Kremer, K.; Ocko, B. M. Nat. Commun. 2012, 3, 795. M


Review

Chemistry of Materials (39) Brand, V.; Bruner, C.; Dauskardt, R. H. Sol. Energy Mater. Sol. Cells 2012, 99, 182. (40) Brand, V.; Levi, K.; McGehee, M. D.; Dauskardt, R. H. Sol. Energy Mater. Sol. Cells 2012, 103, 80. (41) Savagatrup, S.; Printz, A. D.; O’Connor, T. F.; Zaretski, A. V.; Rodriquez, D.; Sawyer, E. J.; Rajan, K. M.; Acosta, R. I.; Root, S. E.; Lipomi, D. J. Energy Environ. Sci. 2015, 8, 55. (42) Bruner, C.; Novoa, F.; Dupont, S.; Dauskardt, R. ACS Appl. Mater. Interfaces 2014, 6, 21474−21483. (43) Savagatrup, S.; Printz, A. D.; Rodriquez, D.; Lipomi, D. J. Macromolecules 2014, 47, 1981. (44) O’Connor, T. F.; Zaretski, A. V.; Shiravi, B. A.; Savagatrup, S.; Printz, A. D.; Diaz, M. I.; Lipomi, D. J. Energy Environ. Sci. 2014, 7, 370. (45) Lipomi, D. J.; Tee, B. C. K.; Vosgueritchian, M.; Bao, Z. N. Adv. Mater. 2011, 23, 1771. (46) Ferenczi, T. A. M.; Müller, C.; Bradley, D. D. C.; Smith, P.; Nelson, J.; Stingelin, N. Adv. Mater. 2011, 23, 4093. (47) Yu, G.; Gao, J.; Hummelen, J. C.; Wudl, F.; Heeger, A. J. Science 1995, 270, 1789. (48) Liu, Y.; Zhao, J.; Li, Z.; Mu, C.; Ma, W.; Hu, H.; Jiang, K.; Lin, H.; Ade, H.; Yan, H. Nat. Commun. 2014, 5, 5293. (49) Papadopoulos, P.; Floudas, G.; Chi, C.; Wegner, G. J. Chem. Phys. 2004, 120, 2368. (50) Note that the VFT equation is mathematically equivalent to the Williams−Landel−Ferry equation, which is typically used in conjunction with the time−temperature superposition principle (cf. refs 1 and 53). (51) Gitsas, A.; Floudas, G.; Wegner, G. Phys. Rev. E 2004, 69, 041802. (52) Mierzwa, M.; Floudas, G.; Neidhöfer, M.; Graf, R.; Spiess, H. W.; Meyer, W. H.; Wegner, G. J. Chem. Phys. 2002, 117, 6289. (53) Gedde, U. W. Polymer Physics; Kluwer Academic Publishers: Dordrecht, 1995. (54) Chen, S.-A.; Ni, J.-M. Macromolecules 1992, 25, 6081. (55) van de Leur, R. H. M.; de Ruiter, B.; Breen, J. Synth. Met. 1993, 54, 203. (56) Pankaj, S.; Hempel, E.; Beiner, M. Macromolecules 2009, 42, 716. (57) Pankaj, S.; Beiner, M. J. Phys. Chem. B 2010, 114, 15459. (58) Barard, S.; Heeney, M.; Chen, L.; Cölle, M.; Shkunov, M.; McCulloch, I.; Stingelin, N.; Philips, M.; Kreouzis, T. J. Appl. Phys. 2009, 105, 013701. (59) Hopkinson, P. E.; Staniec, P. A.; Pearson, A. J.; Dunbar, A. D. F.; Wang, T.; Ryan, A. J.; Jones, R. A. L.; Lidzey, D. G.; Donald, A. M. Macromolecules 2011, 44, 2908. (60) Kesters, J.; Verstappen, P.; Raymakers, J.; Vanormelingen, W.; Drijkoningen, J.; D’Haen, J.; Manca, J.; Lutsen, L.; Vanderzande, D.; Maes, W. Chem. Mater. 2015, 27, 1332. (61) Beaucage, G.; Composto, R.; Stein, R. S. J. Polym. Sci., Part B: Polym. Phys. 1993, 31, 319. (62) Keddie, J. L.; Jones, R. A. L.; Cory, R. A. Europhys. Lett. 1994, 27, 59. (63) Kawana, S.; Jones, R. A. L. Phys. Rev. E 2001, 63, 021501. (64) Dalnoki-Veress, K.; Forrest, J. A.; Murray, C.; Gigault, C.; Dutcher, J. R. Phys. Rev. E 2001, 63, 031801. (65) Singh, L.; Ludovice, P. J.; Henderson, C. L. Thin Solid Films 2004, 449, 231. (66) Sims, M.; Zheng, K.; Campoy-Quiles, M.; Xia, R.; Stavrinou, P. N.; Bradley, D. D. C.; Etchegoin, P. J. Phys.: Condens. Matter 2005, 17, 1. (67) Campoy-Quiles, M.; Sims, M.; Etchegoin, P. G.; Bradley, D. D. C. Macromolecules 2006, 39, 7673. (68) Campoy-Quiles, M.; Alonso, M. I.; Bradley, D. D. C.; Richter, L. J. Adv. Funct. Mater. 2014, 24, 2116. (69) Wang, T.; Pearson, A. J.; Dunbar, A. D. F.; Staniec, P. A.; Watters, D. C.; Coles, D.; Yi, H. N.; Iraqi, A.; Lidzey, D. G.; Jones, R. A. L. Eur. Phys. J. E 2012, 35, 129.

(70) Wang, T.; Pearson, A. J.; Dunbar, A. D. F.; Staniec, P. A.; Watters, D. C.; Yi, H. N.; Ryan, A. J.; Jones, R. A. L.; Iraqi, A.; Lidzey, D. G. Adv. Funct. Mater. 2012, 22, 1399. (71) Liu, D.; Orozco, R. O.; Wang, T. Phys. Rev. E 2013, 88, 022601. (72) Müller, C.; Andersson, L. M.; Peña-Rodríguez, O.; Garriga, M.; Inganäs, O.; Campoy-Quiles, M. Macromolecules 2013, 46, 7325. (73) Labram, J. G.; Domingo, E. B.; Stingelin, N.; Bradley, D. D. C.; Anthopoulos, T. D. Adv. Funct. Mater. 2011, 21, 356. (74) Andersson, L. M.; Hedström, S.; Persson, P. Appl. Phys. Lett. 2013, 103, 213303. (75) Lindqvist, C.; Wang, E. G.; Andersson, M. R.; Müller, C. Macromol. Chem. Phys. 2014, 215, 530. (76) Fox, T. G.; Flory, P. J. J. Appl. Phys. 1950, 21, 581. (77) Müller, C.; Wang, E.; Andersson, L. M.; Tvingstedt, K.; Zhou, Y.; Andersson, M. R.; Inganäs, O. Adv. Funct. Mater. 2010, 20, 2124. (78) Banach, M. J.; Friend, R. H.; Sirringhaus, H. Macromolecules 2003, 36, 2838. (79) McCulloch, B.; Ho, V.; Hoarfrost, M.; Stanley, C.; Do, C.; Heller, W. T.; Segalman, R. A. Macromolecules 2013, 46, 1899. (80) Zhao, J.; Swinnen, A.; Van Assche, G.; Manca, J.; Vanderzande, D.; Mele, B. V. J. Phys. Chem. B 2009, 113, 1587. (81) Hugger, S.; Thomann, R.; Heinzel, T.; Thurn-Albrecht, T. Colloid Polym. Sci. 2004, 282, 932. (82) Kim, J. Y.; Frisbie, C. D. J. Phys. Chem. C 2008, 112, 17726. (83) Gettinger, C. L.; Heeger, A. J.; Drake, J. M.; Pine, D. J. J. Chem. Phys. 1994, 101, 1673. (84) Fytas, G.; Nothofer, H. G.; Scherf, U.; Vlassopoulos, D.; Meier, G. Macromolecules 2002, 35, 481. (85) Grell, M.; Bradley, D. D. C.; Long, X.; Chamberlain, T.; Inbasekaran, M.; Woo, E. P.; Soliman, M. Acta Polym. 1998, 49, 439. (86) Liu, Y.; Liu, M. S.; Li, X.-C.; Jen, A. K.-Y. Chem. Mater. 1998, 10, 3301. (87) Tanto, B.; Guha, S.; Martin, C. M.; Scherf, U.; Winokur, M. J. Macromolecules 2004, 37, 9438. (88) Blondin, P.; Bouchard, J.; Beaupre, S.; Belletete, M.; Durocher, G.; Leclerc, M. Macromolecules 2000, 33, 5874. (89) Teetsov, J.; Fox, M. A. J. Mater. Chem. 1999, 9, 2117. (90) Donat-Bouillud, A.; Lévesque, I.; Tao, Y.; D’Iorio, M.; Beaupré, S.; Blondin, P.; Ranger, M.; Bouchard, J.; Leclerc, M. Chem. Mater. 2000, 12, 1931. (91) Kinder, L.; Kanicki, J.; Petroff, P. Synth. Met. 2004, 146, 181. (92) Wei, S. K. H.; Zeng, L. C.; Marshall, K. L.; Chen, S. H. J. Polym. Sci., Part B: Polym. Phys. 2011, 49, 725. (93) Dang, X. D.; Intelmann, C. M.; Rammelt, U.; Plieth, W. J. Solid State Electrochem. 2005, 9, 706. (94) Wu, P.-T.; Xin, H.; Kim, F. S.; Ren, G.; Jenekhe, S. A. Macromolecules 2009, 42, 8817. (95) Kuila, B. K.; Nandi, A. K. J. Phys. Chem. B 2006, 110, 1621. (96) Schäfer, O.; Greiner, A.; Pommerehne, J.; Guss, W.; Vestweber, H.; Tak, H. Y.; Bässler, H.; Schmidt, C.; Lüssem, G.; Schartel, B.; Stümpflen, V.; Wendorff, J. H.; Spiegel, S.; Möller, C.; Spiess, H. W. Synth. Met. 1996, 82, 1. (97) Zhao, J.; Bertho, S.; Vandenbergh, J.; Van Assche, G.; Manca, J.; Vanderzande, D.; Yin, X. Q.; Shi, J. D.; Cleij, T.; Lutsen, L.; Van Mele, B. Phys. Chem. Chem. Phys. 2011, 13, 12285. (98) Müller, C.; Esmaeli, M.; Riekel, C.; Breiby, D. W.; Inganäs, O. Polymer 2013, 54, 805. (99) Kroon, R. Synthesis and properties of pi-conjugated polymers for organic photovoltaics. Chalmers University of Technology, 2013. (100) Shakutsui, M.; Matsuura, H.; Fujita, K. Org. Electronics 2009, 10, 834. (101) Vandenbergh, J.; Conings, B.; Bertho, S.; Kesters, J.; Spoltore, D.; Esiner, S.; Zhao, J.; Van Assche, G.; Wienk, M. M.; Maes, W.; Lutsen, L.; Van Mele, B.; Janssen, R. A. J.; Manca, J.; Vanderzande, D. J. M. Macromolecules 2011, 44, 8470. (102) Chen, S. H.; Su, A. C.; Su, C. H.; Chen, S. A. J. Phys. Chem. B 2006, 110, 4007. (103) Krebs, F. C. Org. Electronics 2009, 10, 761. N


Review

Chemistry of Materials (104) Standard Test Methods for Photovoltaic Modules in Cyclic Temperature and Humidity Environments, ASTM Standard E 1171. (105) Reese, M. O.; Gevorgyan, S. A.; Jørgensen, M.; Bundgaard, E.; Kurtz, S. R.; Ginley, D. S.; Olson, D. C.; Lloyd, M. T.; Moryillo, P.; Katz, E. A.; Elschner, A.; Haillant, O.; Currier, T. R.; Shrotriya, V.; Hermenau, M.; Riede, M.; Kirov, K. R.; Trimmel, G.; Rath, T.; Inganas, O.; Zhang, F. L.; Andersson, M.; Tvingstedt, K.; Lira-Cantu, M.; Laird, D.; McGuiness, C.; Gowrisanker, S.; Pannone, M.; Xiao, M.; Hauch, J.; Steim, R.; DeLongchamp, D. M.; Rösch, R.; Hoppe, H.; Espinosa, N.; Urbina, A.; Yaman-Uzunoglu, G.; Bonekamp, J. B.; van Breemen, A. J. J. M.; Girotto, C.; Voroshazi, E.; Krebs, F. C. Sol. Energy Mater. Sol. Cells 2011, 95, 1253. (106) Verploegen, E.; Mondal, R.; Bettinger, C. J.; Sok, S.; Toney, M. F.; Bao, Z. A. Adv. Funct. Mater. 2010, 20, 3519. (107) Ngo, T. T.; Nguyen, D. N.; Nguyen, V. T. Adv. Nat. Sci.: Nanosci. Nanotechnol. 2012, 3, 045001. (108) Pearson, A. J.; Wang, T.; Jones, R. A. L.; Lidzey, D. G.; Staniec, P. A.; Hopkinson, P. E.; Donald, A. M. Macromolecules 2012, 45, 1499. (109) Leman, D.; Kelly, M. A.; Ness, S.; Engmann, S.; Herzing, A.; Snyder, C.; Ro, H. W.; Kline, R. J.; DeLongchamp, D. M.; Richter, L. J. Macromolecules 2015, 48, 383. (110) Treat, N. D.; Brady, M. A.; Smith, G.; Toney, M. F.; Kramer, E. J.; Hawker, C. J.; Chabinyc, M. L. Adv. Energy Mater. 2011, 1, 82. (111) Lu, X.; Weiss, R. A. Macromolecules 1992, 25, 3242. (112) Fox, T. G. Bull. Am. Phys. Soc. 1956, 1, 123. (113) Yang, X.; Duren, J. K. J. v.; Janssen, R. A. J.; Michels, M. A. J.; Loos, J. Macromolecules 2004, 37, 2151. (114) Wang, E. G.; Bergqvist, J.; Vandewal, K.; Ma, Z. F.; Hou, L. T.; Lundin, A.; Himmelberger, S.; Salleo, A.; Müller, C.; Inganäs, O.; Zhang, F. L.; Andersson, M. R. Adv. Energy Mater. 2013, 3, 806. (115) Kroon, R.; Gehlhaar, R.; Steckler, T. T.; Henriksson, P.; Müller, C.; Bergqvist, J.; Hadipour, A.; Heremans, P.; Andersson, M. R. Sol. Energy Mater. Sol. Cells 2012, 105, 280. (116) Swinnen, A.; Haeldermans, I.; Vanlaeke, P.; D’Haen, J.; Poortmans, J.; D’Olieslaeger, M.; Manca, J. V. Eur. Phys. J. Appl. Phys. 2006, 36, 251. (117) Conings, B.; Bertho, S.; Vandewal, K.; Senes, A.; D’Haen, J.; Manca, J.; Janssen, R. A. J. Appl. Phys. Lett. 2010, 96, 163301. (118) Watts, B.; Belcher, W. J.; Thomsen, L.; Ade, H.; Dastoor, P. C. Macromolecules 2009, 42, 8392. (119) Treat, N. D.; Mates, T. E.; Hawker, C. J.; Kramer, E. J.; Chabinyc, M. L. Macromolecules 2013, 46, 1002. (120) Fischer, F.; Hahn, T.; Bässler, H.; Bauer, I.; Strohriegl, P.; Köhler, A. Adv. Funct. Mater. 2014, 24, 6172. (121) Kim, B. J.; Miyamoto, Y.; Ma, B. W.; Fréchet, J. M. J. Adv. Funct. Mater. 2009, 19, 2273. (122) Miyanishi, S.; Tajima, K.; Hashimoto, K. Macromolecules 2009, 42, 1610. (123) Carlé, J. E.; Andreasen, B.; Tromholt, T.; Madsen, M. V.; Norrman, K.; Jørgensen, M.; Krebs, F. C. J. Mater. Chem. 2012, 22, 24417. (124) Krebs, F. C.; Spanggaard, H. Chem. Mater. 2005, 17, 5235. (125) Liu, J. S.; Kadnikova, E. N.; Liu, Y. X.; McGehee, M. D.; Frechet, J. M. J. J. Am. Chem. Soc. 2004, 126, 9486. (126) Drees, M.; Hoppe, H.; Winder, C.; Neugebauer, H.; Sariciftci, N. S.; Schwinger, W.; Schaffler, F.; Topf, C.; Scharber, M. C.; Zhu, Z. G.; Gaudiana, R. J. Mater. Chem. 2005, 15, 5158. (127) Cheng, Y. J.; Hsieh, C. H.; Li, P. J.; Hsu, C. S. Adv. Funct. Mater. 2011, 21, 1723. (128) Derue, L.; Dautel, O.; Tournebize, A.; Drees, M.; Pan, H. L.; Berthumeyrie, S.; Pavageau, B.; Cloutet, E.; Chambon, S.; Hirsch, L.; Rivaton, A.; Hudhomme, P.; Facchetti, A.; Wantz, G. Adv. Mater. 2014, 26, 5831. (129) Li, Z.; Wong, H. C.; Huang, Z. G.; Zhong, H. L.; Tan, C. H.; Tsoi, W. C.; Kim, J. S.; Durrant, J. R.; Cabral, J. T. Nat. Commun. 2013, 4, 2227. (130) Piersimoni, F.; Degutis, G.; Bertho, S.; Vandewal, K.; Spoltore, D.; Vangerven, T.; Drijkoningen, J.; Van Bael, M. K.; Hardy, A.;

D’Haen, J.; Maes, W.; Vanderzande, D.; Nesladek, M.; Manca, J. J. Polym. Sci., Part B: Polym. Phys. 2013, 51, 1209. (131) Richards, J. J.; Rice, A. H.; Nelson, R. D.; Kim, F. S.; Jenekhe, S. A.; Luscombe, C. K.; Pozzo, D. C. Adv. Funct. Mater. 2013, 23, 514. (132) Lindqvist, C.; Bergqvist, J.; Feng, C. C.; Gustafsson, S.; Bäcke, O.; Treat, N. D.; Bounioux, C.; Henriksson, P.; Kroon, R.; Wang, E. G.; Sanz-Velasco, A.; Kristiansen, P. M.; Stingelin, N.; Olsson, E.; Inganäs, O.; Andersson, M. R.; Müller, C. Adv. Energy Mater. 2014, 4, 1301437. (133) Liu, H. W.; Chang, D. Y.; Chiu, W. Y.; Rwei, S. P.; Wang, L. J. Mater. Chem. 2012, 22, 15586. (134) Santo, Y.; Jeon, I.; Yeo, K. S.; Nakagawa, T.; Matsuo, Y. Appl. Phys. Lett. 2013, 103, 073306. (135) Chen, C. Y.; Tsao, C. S.; Huang, Y. C.; Liu, H. W.; Chiu, W. Y.; Chuang, C. M.; Jeng, U. S.; Su, C. J.; Wu, W. R.; Su, W. F.; Wang, L. Nanoscale 2013, 5, 7629. (136) Lindqvist, C.; Bergqvist, J.; Backe, O.; Gustafsson, S.; Wang, E.; Olsson, E.; Inganäs, O.; Andersson, M. R.; Müller, C. Appl. Phys. Lett. 2014, 104, 153301. (137) Schroeder, B. C.; Li, Z.; Brady, M. A.; Faria, G. C.; Ashraf, R. S.; Takacs, C. J.; Cowart, J. S.; Duong, D. T.; Chiu, K. H.; Tan, C. H.; Cabral, J. T.; Salleo, A.; Chabinyc, M. L.; Durrant, J. R.; McCulloch, I. Angew. Chem., Int. Ed. 2014, 53, 12870.

O


On the Glass Transition of Polymer Semiconductors and Its Impact on

Recommend Documents