Probing the Relationship between Molecular Structures, Thermal

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Probing the Relationship Between Molecular Structures, Thermal Transitions and Morphology in Polymer Semiconductors Using a Woven Glass-Mesh Based DMTA Technique Anirudh Sharma, Xun Pan, Jonas M. Bjuggren, Desta Gedefaw, Xiaofeng Xu, Renee Kroon, Ergang Wang, Jonathan A Campbell, David A. Lewis, and Mats R Andersson Chem. Mater., Just Accepted Manuscript • Publication Date (Web): 23 May 2019 Downloaded from http://pubs.acs.org on May 23, 2019

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Chemistry of Materials

Probing

the

Relationship

Between

Molecular

Structures, Thermal Transitions and Morphology in Polymer Semiconductors Using a Woven Glass-Mesh Based DMTA Technique Anirudh Sharma,α,¥,* Xun Pan,α Jonas M. Bjuggren,α Desta Gedefaw,α,≠ Xiaofeng Xu,║ Renee Kroon,║ Ergang Wang,║ Jonathan A Campbell,α David A. Lewis,α Mats R. Anderssonα,* αFlinders

Institute for Nanoscale Science and Technology, Flinders University, Sturt Road,

Bedford Park, Adelaide, SA, 5042, Australia. ¥University

of Bordeaux, Laboratoire de Chimie des Polymères Organiques (LCPO), UMR 5629,

B8 Allée Geoffroy Saint Hilaire, 33615 Pessac Cedex, France. ≠School

of Biological and Chemical Sciences, The University of South Pacific, Laucala Campus,

Suva, Fiji ║Department

of Chemistry and Chemical Engineering, Chalmers University of Technology, SE-

41296 Göteborg, Sweden

ABSTRACT

1 Environment ACS Paragon Plus

Chemistry of Materials 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

The glass transition temperature (Tg) of polymers is an important parameter that determines the kinetics of molecular organization of polymeric chains. Understanding the Tg of conjugated polymers is critical in achieving a thermally stable and optimum morphology in polymer:polymer or polymer:small molecule blends in organic electronics. In this study, we have used the woven glass-mesh based method of dynamic mechanical thermal analysis (DMTA) to evaluate the Tg of polymer semiconductors, which is generally not easy to detect using conventional techniques such as differential scanning calorimetry (DSC). More importantly, we establish the relationship between the thermal transitions and the molecular structure of polymer semiconductors. For conjugated polymers with rigid conjugated backbone and large alkyl side chains, we report the presence of separate thermal transitions corresponding to the polymer backbone as well as transitions related to side chains, with latter being the most prominent. By systematically comparing polymer side chains, molecular weight and backbone structure, the origin of the Tg and a sub-Tg transitions have been successfully correlated to the polymer structures. The antiplastization effect of additives has also been used to further prove the origin of the different transitions. Thermal transitions of a range of high performing polymers applied in organic photovoltaics, including TQ1, PTNT, PTB7, PTB7-Th and N2200 has been systematically studied in this work. According to the measurements some of these polymers have a very small amorphous part, changing the way how the morphology should be described for these materials. We infer that the main phase in these polymers consists of hairy aggregates, with a few π-stacked rigid polymer chains forming the aggregates.

INTRODUCTION


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Recent innovation in the field of flexible electronics1 has been driven by considerable research on polymer semiconductors2 and their applications in electronic devices such as organic field effect transistors (OFETs),3-5 organic photodetectors (OPDs)6-7 and polymer solar cells (PSCs).8-9 The performance of these devices rely upon the basic structure of the polymer but additionally on the way these materials order, both with similar molecules and other materials, such as in bulk heterojunctions. Importantly, the as-deposited semiconducting material is in a ‘frozen in morphology’10 and is not in thermodynamic equilibrium and as a result, the morphology will continue to evolve over time as the polymer chains relax to their lowest energy conformation, altering the optoelectronic properties. To achieve efficient devices with a long lifetime, control over the morphology of such polymer films is essential. Noriega et al.11 have discussed the relationship between the microstructure and the charge transport in conjugated polymers and pointed out that intermolecular aggregation plays a key role for charge transport in many polymers. The thermal transition temperatures are critical for the materials used in different kinds of organic electronics. For example, during annealing processes the temperature must exceed the Tg to change the morphology and the final performance of the devices. In polymer solar cells a common way to optimize device properties is thermal treatment. Elevated temperatures can also be detrimental for the morphology and the device performance.12 Different materials behave differently and to be able to understand the material behavior the key issue is to know the thermal transitions of the used material and how the morphology of the material is affected. A good understanding of the Tg of conjugated polymers used in bulk hetero-junction PSCs is critical, in order to achieve an optimal and stable morphology.13-15 Therefore, the physical characteristics of conjugated polymers such as the Tg and other relaxations (eg sidechain motion) must be well understood in order to determine the optimal process conditions, including time and



temperature that each material or blend of materials can be exposed to. Better understanding of thermomechanical properties of semiconducting polymers is becoming even more important with the emergence of stretchable and wearable electronics16-17 involving a higher degree of mechanical strain and physical deformation.18 Differential scanning calorimetry (DSC) is one of the most commonly used techniques to determine the Tg of polymers.19 For conjugated polymers, which are commonly assumed to have a large fraction of amorphous phase, measuring the Tg using DSC can still be non-trivial and quite often no transition can be detected. For this reason, various methods such as variable temperature ellipsometry,20 plasmonic nanospectroscopy,21 oscillatory shear rheometry22 and more recently fast scanning calorimetry23 have been explored to measure the Tg of conjugated polymers. In addition, dynamical mechanical thermal analysis (DMTA) can also be used to measure the structural relaxation of polymers,24 however its application for conjugated polymers in traditional mode of operation has been quite limited as it normally requires free standing film of reasonable thickness for measurements. Due to high cost and complicated synthetic procedure25 of most conjugated polymers, making a freestanding thick film is not viable. According to Hopkins et al, this problem can be overcome using a DMTA material pocket method,26-27 but the results from this method had high noise to signal ratio. More recently, a polyimide template-based method has also been reported for measuring thermomechanical properties of conjugated polymers using DMTA.28 We have recently demonstrated a woven glassmesh based DMTA method which can be successfully utilized to measure the glass transition of conjugated polymers and their blends with PCBM, with high sensitivity.29 In this method, the polymer or polymer-small molecule blend films are reinforced using a woven glass-mesh and tested in tensile mode, making it a versatile method of precisely and sensitively measuring not only the Tg but also the melt/flow transition of conjugated polymers. Most importantly, this method


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allows DMTA measurements utilizing under 5 mg of polymer, instead of much thicker free standing solid samples which are often required for conventional DMTA measurements. P3HT is a thoroughly studied conjugated polymer using many different techniques. We and others have used DMTA to study its thermal transitions. The woven glass-mesh based DMTA technique agrees well with the previously generated DMTA measurements on P3HT. From the DMTA measurements it can be concluded that P3HT have a sub-Tg at -80 ºC from the side chains and a Tg at 34 ºC. When using the glass-mesh support a melting point at 235 ºC could also be determined.28 In this paper, we further utilize the woven glass-mesh based DMTA technique to develop an understanding of the structure-property relationship of conjugated polymers. We first measure the thermomechanical properties of polystyrene (PS), which is a well-studied polymer, as a free standing solid film with conventional DMTA as well as by using the woven glass-mesh method, to extend our understanding of the latter technique and to probe the limitations. We then use the woven glass-mesh method to systematically establish relationship between the molecular structure of a range of high-performing p-type polymers including TQ1, PTNT, PTB7 and PTB7-Th (Figure 1) to their resultant thermomechanical properties. For additional polymers such as P3TI, PBDTTTCT, PTB1, PBDTTT-E and PDPP3T, see Figure S1. In order to further unravel the complexity and interdependence between the main chain rigidity and the size of alkyl side chains on the thermal transitions of conjugated polymers, we also study the thermomechanical properties of a wellknown naphthalene diimide (NDI)-based n-type polymer N220030 and its copolymerized derivatives PNDI-T10. We establish correlation between the structure of the polymer backbone and their thermomechanical properties by probing a range of NDI based random polymers which were synthesized by systematic replacement of bithiophene units in N2200 with thiophene



moieties31 in order to give varied amount of flexibility to the polymer backbone and decreased crystallinity. From our work we can conclude that several of the studied conjugated polymers have an amorphous phase and therefore show a clear glass transition, whereas in a second group of polymers the glass transition is nearly absent which results from the presence of an aggregated phase instead of an amorphous phase.

Figure 1: Structures of the studied conjugated polymers.

EXPERIMENTAL Materials Polymers TQ1 (Mn= 53 kg mol-1, Mw = 132 kg mol-1), TQ1 with low molecular weight (LMWTQ1) (Mn= 10.8 kg mol-1, Mw= 21.8 kg mol-1) and PTNT (Mn= 55.7 kg mol-1, Mw= 163.2 kg mol-


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1)

were synthesized in our lab as per previously reported procedures.25, 32 PTB7 (Mn = 33.1 kg mol-

1,

Mw = 81.8 kg mol-1) and PTB7-Th (Mn = 46.9 kg mol-1, Mw = 102.5 kg mol-1) were generously

donated by Solarmer Materials Inc. N2200 (Mn = 42.3 kg mol-1, Mw = 140.2 kg mol-1) was synthesised as reported by Facchetti et al.30, 33 Molecular weights of all conjugated polymers were measured using a high temperature SEC at 150 ºC in 1,2,4-trichlorobenzene and all values are relative to polystyrene standards. PC61BM was purchased from Solenne BV. The random polymers PNDI-T10 (Mn = 41.9 kg mol-1, PDI = 2.8), PNDI-T20 (Mn = 67.7 kg mol1,

PDI = 5.0), PNDI-T50 (Mn = 41.9 kg mol-1, PDI = 3.1) were synthesized according to the

literature.31 Synthesis of PNDI-T10 with hexyldecyl side chain (PNDI-T10 HD) is described in the supplementary information.

Sample Preparation DMTA samples were prepared using the procedure described elsewhere.29 Polystyrene (Sigma Aldrich, atactic Mn =140 kg mol-1 and Mw = 230 kg mol-1) samples were prepared both via solution processing as well as by melt processing. Low molecular weight polystyrene (Agilent Technologies, GPC standarad, Mn = 34.8 kg/mol, Mw/Mn = 1.02) sample was prepared by drop casting polystyrene solution on glass-mesh. For sample prepared on woven glass-mesh, 20 mg mL1

of PS was dissolved in CHCl3 and deposited on the glass-mesh via drop casting, followed by a

drying in a vacuum oven at 140 ºC for 10 minutes, before cooling down to room temperature. The sample was then immediately measured on DMTA. For preparing a melt processed PS sample, PS pellets were compressed between two polished stainless steel plates in a hydraulic press at 180 ºC, and then cooled to room temperature resulting in a thin film of ~ 0.4 mm.



For all other polymers, samples were prepared by drop-casting the polymer solutions: TQ1 (20 mg mL-1), LMW-TQ1 (20 mg mL-1), PTNT (28 mg mL-1), PTB7 and PTB7-Th (25 mg mL-1), N2200 (25 mg mL-1), PNDI-Tx (~8 mg mL-1) and PNDI-T10 HD (10 mg mL-1) from CHCl3, and TQ1:PC61BM (1:2.5 weight ratio) (25 mg mL-1), PTNT:PC61BM (1:2 weight ratio) (25 mg mL-1) from ortho-dichlorobenzene (o-DCB) on the glass-mesh mesh until a continuous and uniform coverage is achieved as described earlier.29 The polymer film thickness on the glass-mesh was assumed/estimated to be 0.1 mm for all samples.

Dynamical Mechanical Thermal Analysis (DMTA) All measurements were performed on a Q800 DMTA machine supplied by TA Instruments. Samples were measured at a frequency of 1 Hz in strain-controlled mode, under a continuous flow of N2 at 60 mL min-1 and the rate of heating was set to 3 ˚C per minute. An initial drying run on all samples was performed from room temperature up to 80 ˚C in order to remove any remaining solvent and physiabsorbed water, followed by a second run from -110 ˚C up to 320 ˚C, unless specified otherwise. In this study, the peak temperature of the tan δ peak is used to define the Tg for all samples. In addition, the temperature corresponding to the tan δ peak at high temperatures accompanied by a drop or minimum in the storage modulus (E’), is used to define the melting temperature (Tm) of samples.

Differential Scanning Calorimetry (DSC) DSC measurements were performed on a TA Instruments Discovery DSC series DSC1-0039. In order for DMTA and DSC measurements to be comparable, all DSC samples were first dried by ramping the temperature from room temperature to 80 ˚C, at a rate of 10 ˚C per minute, followed by the second run from -80 ˚C to 300 ˚C. Data from the second run was used for all samples.


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RESULTS AND DISCUSSION In order to probe the thermal transitions observed from DMTA on films using conventional as well as the woven glass-mesh based technique, we first studied atactic polystyrene. Figure 2 shows the DMTA scan of polystyrene (PS) both as melt processed freestanding bar sample and a woven glass mesh supported sample. In the case of the melt processed sample, the Tg as determined by the tan δ peak was found to be at 112 ºC, and the large deformation of the sample due to softening precluded continuation of the measurement beyond 165 ºC. For the polystyrene sample prepared on woven glass-mesh, the Tg was also found to be at 112 ºC. It must be noted that due to the significantly smaller amount of polymer on the glass-mesh (~3-5 mg) as compared to that of the melt-processed freestanding PS film, the signal to noise ratio was relatively lower in the case of glass-mesh based sample compared to the latter. However, the glass-mesh supported sample was successfully measured until 250 ºC, as compared to the melt processed PS film which failed around ~ 165 ºC.

Figure 2: DMTA scan at 1 Hz showing the storage modulus (dashed lines) and normalized tan δ response (solid lines) from (a) high molecular weight polystyrene sample prepared by melt processing (b) low and (c) high molecular weight polystyrene, drop-cast on woven glass-mesh. Inset shows the chemical structure of PS.



In the case of high molecular weight (HMW) PS sample on glass-mesh, the tan δ peak at 220 ºC was also accompanied by a sharp drop in storage modulus, and is attributed to the viscous flow temperature of polystyrene. This was confirmed by measuring a relatively low molecular weight (LMW) PS sample (Mw - 34.8 kg mol-1), which shows the same Tg (112 °C)34-35 but the weak tan δ peak corresponding to the flow transition was found to be shifted to a much lower temperature (154 °C). The onset of the rubbery phase of both polystyrene samples, as seen from the significant drop in E’, was found at 95 °C. It must however be noted that the magnitude of the loss in E’ was over three orders of magnitude in the case of melt processed PS and only over two orders of magnitude for the glass-mesh based sample. This demonstrates that the glass-mesh is effective in reinforcing the polymer film to observe a range of relaxations that otherwise can not be observed with DMTA, but the magnitude of some relaxations is limited. It is thus important to emphasize that in the case of polymer samples measured using the woven glass-mesh, the identification and interpretation of the glass transition of polymers from the change in storage modulus should be done carefully, as E’ may not necessarily reduce by many orders of magnitude at the glass transition temperature, as conventionally known for amorphous polymers when using the DMTA technique.36 TQ1 is a simple, well performing conjugated polymer that has been used in polymer solar cells with efficiencies of up to 7% with conventional fullerene derivatives.8, 37 The DMTA for TQ1 (Figure 3), shows a small tan δ peak at 10 ºC, accompanied by a prominent peak at 100 ºC, which have previously been attributed to the sub-Tg and Tg of TQ1, respectively.29, 38 The prominent tan δ peak corresponding to the Tg is in agreement with the known amorphous nature of TQ1, which was also confirmed by XRD (Figure S2a). To experimentally correlate the origin of the sub-Tg to the side chain segments of the polymer, a TQ1 batch with low molecular weight (LMW-TQ1) was


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synthesized. The DMTA scan of LMW-TQ1 (Figure 3a) shows a significant reduction in the Tg from 100 °C to 80 °C, as seen from the shift in the tan δ peak. Meanwhile the sub-Tg is located at a simmilar temperature as that of high molecular weight TQ1, confirming that the observed low temperature features in the tan δ curve in both the cases are sub-Tg transitions due to the side chains in TQ1. For the active material in solar cells, TQ1:PC61BM (1:2.5 weight ratio) blend sample (Figure 3b), shows two tan δ peaks at 95 ºC and 147 ºC, attributed to the thermal transition of TQ1-rich domains and PC61BM-rich domains, respectively.29 However, the clear decrease of the sub-Tg transition in the blend is an indication that PC61BM reduced the mobility of TQ1 side chains due to antiplasticization. This is in agreement with previously reported fundamental studies on polymer properties39 further proving that this low temperature transition originates from the side chains on TQ1. The increase in E’ above 147 ºC for the blend originates from crystallization of PC61BM, resulting in an increased sample stiffness. Formation of large PCBM crystals in the active layer has been shown to significantly reduce the efficiency of solar cells.12

Figure 3: DMTA scan showing the changes in the storage modulus (dotted lines) and tan δ (solid lines) of TQ1 compared with that of (a) LMW -TQ1 and (b) TQ1:PC61BM blend.

In order to probe the influence of the long alkyl side chains on the thermal transitions of conjugated polymers, a donor polymer PTNT,25 which has a rigid backbone and long branched side chains 11 Environment ACS Paragon Plus


were studied. PTNT was found to soften at temperatures ranging between -40 ºC to 20 ºC, as seen from the drop in E’ with an onset at -40 ºC (Figure 4a). Unlike TQ1, the tan δ peak around -20 ºC showing the low temperature transition of PTNT was the most intense thermal transitions observed up to 270 ºC. This behavior is observed due to the large size of the polymer side chains, which impact sub-Tg relaxations and is in good agreement with recent findings on thermomechanical properties of rigid polymers reported by Z. Bao and collegues.28 These results are also consistent with the findings of Pankaj et al.40 who demonstrated that large alkyl side chains can contribute an additional relaxation process in addition to the main chain relaxation, in conjugated polymers such as poly(3-alkylthiophenes). A weak transition around 90 ºC was also observed which could be due to the relaxation of some main chain/conjugation segments. Importantly, DSC of PTNT only revealed a weak transition at ~ - 40 ºC while no other transitions were observed at higher temperatures (Figure S3). In order to verify that the tan δ peak around -20 ºC is not attributed to the glass transition of PTNT- which relates to the thermal relaxation of the polymer backbone, a blend film consisting of PTNT:PC61BM (1:2 weight ratio) was examined. Interestingly, a tan δ peak is clearly observed at 159 ºC, which is attributed to the Tg of the PC61BM rich phase. Since the blend-Tg in the case of PTNT:PC61BM is found to be higher than the thermal transition temperature (147 ºC) of PC61BM-rich phase observed in the case of TQ1:PC61BM blend, it indicates a higher Tg of PTNT or less PTNT in the PC61BM rich phase. Importantly, the low temperature tan δ peak of PTNT was found to be significantly suppressed when blended with PC61BM, similar to the TQ1:PC61BM blend. This proves that the thermal relaxation of PTNT observed at -20 ºC is a sub-Tg transition in nature and originates due to the motion of the side chains. The increase in E’ for the blend at 166 ºC is due to the crystallization of PC61BM, as


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described previously for TQ1:PC61BM blends. The tan δ peak observed at 260 ºC in neat PTNT sample accompanied by a step drop in E’ is attributed to the melting of PTNT crystallites.

Figure 4: DMTA scan showing the changes in the storage modulus (dotted lines) and tan δ (solid lines) of (a) PTNT and PTNT:PC61BM (1:2 weight ratio) blend (b) PTB7 and PTB7-Th. Since the melting temperature of PTNT was found to be 260 ºC, the Tg of PTNT, is expected to be below 260 ºC. The absence of any prominent DMTA signal in the temperature range between 50 ºC to 260 ºC is attributed to the rigidity of the backbone and its tendency to pack in a solid state film, as found by XRD measurements (Figure S2b). This is also consistent with the findings of Kroon et al.25 who reported the presence of local order in the polymer film due to π-stacking of the PTNT backbone. As reported by Noriega et al.11 π-stacking could also lead to a short-range order resulting in the formation of local intermolecular aggregates, which would explain the absence of a measureable backbone relaxations, that correspond to the lack of disordered amorphous phase in PTNT. The presence of aggregates already seen in PTNT solution were also confirmed using temperature dependent UV-vis absorption measurements41 (Figure S4 a). To explain the lack of measurable Tg by DMTA our hypothesis is that the aggregates dominates the morphology and there are no fully movable amorphous phase present in PTNT.



To further probe and understand the absence of a Tg in PTNT, we studied the high performing donor polymers PTB7 and PTB7-Th using DMTA, since PTB7 has a rigid backbone and is most likely paracrystalline42 with short-range order. While PTB7 and PTB7-Th possess a similar backbone conjugated structure, the difference in their chemical structure arises from the functional groups attached to the benzodithiophene groups: an ether in PTB7 and a thiophene in the case of PTB7-Th (Figure 1), making PTB7-Th thermally more stable than PTB7.43 For PTB7, a prominent tan δ peak was observed at -40 ºC along with a drop in the storage modulus. In addition, a small tan δ peak at 30 ºC and a large peak at 260 ºC were also observed. In the case of PTB7-Th, a drop in the storage modulus (Figure 4b) was observed between -50 ºC to 50 ºC accompanied with a prominent tan δ peak at -10 ºC, a small feature at 70 ºC and a peak at 260 ºC. Since both PTB7 and PTB7-Th have a rigid backbone similar to PTNT, the tan δ peak at -40 ºC and -10 ºC for the two polymers likely does not arise from the relaxation of the main chain segments. Thus, these features are attributed to a secondary relaxation related to the side chains of PTB7 and PTB7-Th. The tan δ peak at -40 ºC was found to shift to a relatively higher temperature of -10 ºC on replacing the ‘ether’ linkage in PTB7 with a ‘thiophene’ unit in PTB7-Th. This was expected as the alkyl side chain with a thiophene moiety is effectively stiffer compared to its ethercontaining counterpart. For both PTB7 and PTB7-Th, the DMTA tan δ scans between 120 to 180 ºC were found to be reproducibly noisy over multiple samples, but no clear transition was observed around 127 ºC, which was estimated to be the Tg of PTB7 by Root et al.44 For both PTB7 and PTB7-Th, XRD measurements revealed a broad feature along with a small diffraction peak around 15° (Figure S2c) corresponding to a packing distance of ~6.9 Å, indicating some ordered packing in both polymers. This is consistent with previous study by Bencheikh et al., where a strong aggregation


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behavior has been found for PTB7 and PTB7-Th (Figure S4b) in both solution as well as in solid films.45 Thus the negligible relaxation of the backbone seen in the DMTA measurements are due to the lack of disordered amorphous phase caused by presence of mainly aggregates. The tan δ peaks observed around 260 ºC for PTB7 and PTB7-Th are attributed to sample deformation, as a result of the onset of the melt transition of the polymer aggregates, which is also accompanied by a drop in the storage modulus at 260 ºC. It must be pointed out that the DSC measurements on PTB7 and PTB7-Th did not show any transitions between -50 ºC up to 350 ºC. (Figure S5) While it has been successfully established in the above experiments that large alkyl side chains contribute to a prominent sub-Tg transition in the low temperature range between 0 ºC to -50 ºC, the influence of the backbone rigidity on the non-existing Tg needs to be elucidated further. Our findings with regard to the Tg of PTB7-Th are in disagreement with a recently published report by Chen et al.46 who used DMTA and measured the Tg of PTB7-Th to be 138 ºC. In the same report, the sub-Tg of PTB7-Th was also claimed to have been reduced from 61.2 ºC to 36.4 ºC upon blending with PC71BM. To probe the correlation between the rigidity of polymer backbone with their thermal properties, an acceptor polymer N2200, which is substituted with octyldodecyl side groups, was investigated. DMTA scan of neat N2200 was found to have a broad transition between -60 to 40 ºC with a prominent tan δ peak at -15 ºC and a sharp drop in storage modulus in the same temperature range (Figure 5). This result is largely similar to the previously reported DMTA scan of N2200 by McNeill and colleagues.47 However, McNeill and colleagues have claimed -70 ºC to be the Tg of N2200,47 which is counter intuitive for a conjugated structure with a rigid backbone that can be expected to be stiffer than polymers such as PS48, PMMA (Tg ~ 107 ºC)49 and TQ1 (Tg ~95 ºC). On careful analysis, minor features at elevated temperature between 100 ºC and 180 ºC were also



observed in the tan δ of N2200 (Figure 5). However, the low signal intensity precludes definitive attribution to a glass transition. Even though XRD measurements revealed the semicrystalline nature of the polymer with a broad feature due to the amorphous phase and diffraction peaks at low angle (large d-spacing) attributed to the crystalline phase (Figure S2d), the limited signal in DMTA corresponding to the backbone relaxation can again be explained by a predominantly aggregated morphology and a fraction of larger crystals in agreement with previous reports stating that N2200 is crystalline.50-51 Furthermore, the tan δ peak at 315 ºC accompanied with a step drop in the E’ is attributed to the melt transition of N220047 and is also in close agreement with the previously reported melting temperature of N2200 (305 ºC).31

Figure 5: DMTA scan of N2200. In order to verify the origin of the thermal transition at -15 ºC to the polymer side chain relaxation of N2200, a set of random polymers with NDI-thiophene backbone similarly to N2200 (Figure 1) were studied. In the first set of random polymers, 10% (PNDI-T10), 20% (PNDI-T20) and 50% (PNDI-T50) of the bithiophene units were replaced with thiophene units. The introduction of the thiophene unit has previously been shown to increase the backbone flexibility and reduce the polymer crystallinity.31 Since only the amorphous phase of the polymer has a Tg, the reduction in the crystallinity is therefore expected to increase the intensity of the Tg transition.


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On replacing the bithiophene moieties with the thiophene units, a slightly more evident tan δ feature at 120 ºC was found, for the random polymers with 10% to 50% of bithiophene being replaced with thiophene (Figure 6a), suggesting the backbone relaxation to be the origin of this transition. It can be concluded that rigid backbone polymers such as N2200, due to the local order51 in the polymer chains have low extent of disordered amorphous domains making it difficult to probe the glass transition of such polymers. Also, from the tan δ peak at higher temperature, the melt transition for PNDI-T10 was found to be at 290 ºC, which shifted towards lower temperatures of 280 ºC and 240 ºC for PNDI-T20 and PNDI-T50, respectively. The nature of these melt transitions was also confirmed by the observed drop in E’ as a function of temperature, for all random polymers (Figure 6b) with different amounts of thiophene units.

Figure 6: (a) Tan δ transition and (b) E’ showing the melt transition of random polymers with NDIthiophene backbone. These results are in excellent agreement with the melting transitions that were reported earlier from the DSC measurements.31 A new polymer PNDI-T10 HD with the same backbone as PNDI-T10 but hexyldecyl side chain (Figure S6 a) was synthesized, in order to elucidate the origin of the low temperature tan δ feature observed for all random polymers. In the case of the random polymers, the most prominent tan δ



peak was observed at low temperature of 0 ºC. On altering the polymer side chain in PNDI-T10 from octyldodecyl to hexyldecyl as in the case of PNDI-T10 HD, the tan δ feature was found to shift from 0 ºC to -10 ºC (Figure S6). This confirms that the low temperature feature is a sub-Tg transition, and is associated with the modified side chains. When comparing all the polymers analyzed with DMTA in this and our previous work29 it is possible to divide the polymers into two groups (Figure 7). One group of polymers has an amorphous phase with a sub-Tg transition and a clear Tg. These polymers consist of P3HT, TQ1, PCDTBT29 and can be regarded to behave as classical amorphous or semi crystalline polymers. The other group, consisting of N2200, PNDI-Tx, PTB7, PTB7-Th, PTNT, P3TI, PBDTTT-CT, PTB1, PBDTTT-E, PDPP3T has a sub-Tg transition, no clearly detectable Tg and sometimes a crystalline phase. The conclusion is that the latter group of polymers have no clearly detectable disordered amorphous domains (Figure 7). Since they are not completely crystalline, for example the crystallinity of PTB7 is too low to be detectable using scattering techniques,42 the main phase must be different compared to classical amorphous or semicrystalline polymers. We propose that the main phase in the solid state is made up of “hairy aggregates” as shown in Figure 7. All of these polymers have a flat rigid backbone with long branched side chains, and it is well known that this type of polymers can π-stack and form aggregates (small ordered regions) even in solution. The series of small, relatively closely spaced aggregates explains the low apparent crystallinity by XRD, the inability to detect a Tg and the high temperature behaviour in the DMTA. The long side chains on the polymer backbones create a “hairy” layer of matter in between the aggregates (Figure 7, hence the name hairy aggregates) and has a thermal transition at sub-ambient temperatures. When fullerenes are added, such as in solar cell materials, they preferentially partition between the aggregates and reduce the mobility of the side chains,


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decreasing the intensity of the low temperature transition, as seen in Figure 4a for PTNT:PC61BM blends. This type of morphology is a variation of the so called paraffinic microstructures that has previously been observed for both classical as well as conjugated52-53 polymers, however, the proposed model overcomes the requirement that the molecular weights are relatively low, which is not the case for most of the studied polymers.

Amorphous

Crystalline

Aggregate

Hairy aggregates

Figure 7: A schematic illustration of the morphology corresponding to the two groups in which conjugated polymers can be divided into according to this study: polymers with amorphous and crystalline phase (left) and polymers having mainly hairy aggregate and crystalline phase (right). The crystalline fraction varies between different types of polymers.

Table 1: Summary of the thermal transitions measured using the glass-mesh based DMTA method and molecular weight of various conjugated polymers. Polymer TQ1 PTNT

Sub-Tg (°C)

Tg (°C)

Tm* (°C)

Mn (kg/mole)

PDI

10

100

-

53

2.5

10 -20

80 -

~260

10.8 55.7

2.0 2.9



PTB7

~260

33.1

2.5

~260

46.9

2.2

~315

42.3

3.3

-

290

41.9

2.8

0

-

280

67.7

5.0

0

-

240

41.9

3.1

-40

-

-10

-

-10

-

PNDI-T10

0

PNDI-T20 PNDI-T50

PTB7-Th N2200

Page 20 of 28

Polymers included in the supplementary information PNDI-T10 HD

-10

-

~330

-

-

P3TI

-30

-

265

54.7

3.8

PBDTTT-CT

-10

-

280

32.9

1.8

PTB1

-24

-

280

N/A

N/A

PBDTTT-E

-38

-

230

21.3

2.3

PDPP3T

-2

-

305

31.1

1.8

Polymers reported in previous publication29 Polymer

Sub-Tg (°C)

Tg (°C)

Tm (°C)

Mn (kg/mole)

Mw (kg/mole)

P3HT

-80

38

235

N/A

50-70

PCDTBT

-

126

-

N/A

20-100

*The high temperature maximum of tan δ accompanied with a drop or minimum in storage modulus (E’) is taken as the melting point of the crystallites or the aggregates.

Importantly, it must be emphasized that the glass-mesh based DMTA technique also makes it possible to measure the “melting point” of these aggregates, which is seen as a small decrease in


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the storage modulus and a maximum in the tan δ peak (Figure 4). In PTNT and N2200 that have higher crystallinity than PTB7 the modulus drop is larger due to the melting of the crystallites and not only that of the small aggregates (Figure 4a and 5). On the other hand, for polymers which are predominantly composed of aggregates of short range order such as PTB7 and PTB7-Th, the drop in the modulus is small, due to the melting of only aggregates. In Table 1 a summary of all the data we have acquired using the glass-mesh based DMTA is presented. In order to significantly change the morphology of the materials having a morphology mainly consisting of hairy aggregates annealing above the melting point of the aggregates or introducing local chain mobility using solvent annealing31 is necessary. For the polymers exhibiting a Tg thermal annealing above the Tg is sufficient. It is important to understand how the morphology can be modified and improved to achieve efficient and high performing organic electronics of different types of conjugated polymers. The thermal stability of the morphology will also be much higher for the polymers having an aggregated morphology.

CONCLUSIONS Understanding the thermomechanical properties of polymer semiconductors is crucial for their application in electronic devices, in order to achieve optimal and stable morphologies. This study has successfully demonstrated the woven glass-mesh based DMTA method to be a sensitive technique for measuring the relaxations of the polymer sidechains as well as that of the backbone of rigid polymer semiconductors, which are otherwise non-trivial to measure using conventional techniques such as DSC. The thermal transitions of the semiconducting polymers have been found to vary significantly with their structure and stiffness of the backbone. In the case of polymers with large alkyl side chains



and flexible backbone, two thermal transitions originating from the side chain as well as from the backbone are observed from the amorphous phase. Large alkyl side chains of semiconducting polymers contribute to a sub-Tg transition between -80 ºC to 0 ºC, and a Tg around 100 ºC from the backbone, which plays a significant role in defining the thermal properties of these materials. These polymers have morphologies as classical polymers including an amorphous and a crystalline phase in varying amounts. For rigid backbone polymers with long branched side chains, the sub-Tg transition can be the most dominant transition observed whereas the Tg transition is often too weak to be detected. This implies that the amorphous phase is very small in these polymers. Due to the varying crystallinity in these polymers the conclusion is that the main phase in these polymers consists of a hairy aggregated morphology. The aggregates consist of a few π-stacked chains forming small hairy aggregates. This has far-reaching consequences for the understanding on how these materials behave and how the morphology can be altered in the devices. The findings of this study provide new insights into the thermomechanical behavior of rigid backbone conjugated polymers and will be very relevant for a broader community working on flexible organic electronics.

ASSOCIATED CONTENT Supporting Information. XRD of polymers, experimental details of P3TI and PNDI-T10 HD synthesis, DMTA scan of polymers, DSC of PTNT, PTB7, PTBT-Th and PNDI-T10 HD, UV-vis absorption spectra of PTNT and N2200. This material is available free of charge via the Internet at http://pubs.acs.org.

AUTHOR INFORMATION


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Corresponding Author *Mats R. Andersson, Email: [email protected] *Anirudh Sharma, Email: [email protected] Author Contributions The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript.

ACKNOWLEDGMENT The authors would like to thank Flinders University and the Australian Research Council for financial support (DP170102467). Solarmer Materials Inc. is gratefully acknowledged for donating several polymers including PTB7, PTB7-Th for this work. Authors thank Helena Andersson for performing SEC measurements for determining the molecular weight of the polymers. EW further thanks the Swedish Research Council, the Swedish Research Council Formas and the Wallenberg Foundation for financial support.

ABBREVIATIONS PS: Polystyrene TQ1: poly[2,3-bis-(3-octyloxyphenyl)quinoxaline-5,8-diyl- alt -thiophene-2,5-diyl] TQ-HD: poly[2,3-bis-3-(2-hexyldecyl)-oxyphenyl)quinoxaline-5,8-diyl-alt-thiophene-2,5-diyl] PTNT: poly(2,5-thiophene-alt-4,9-bis(2-hexyldecyl)-4,9-dihydrodithieno[3,2c:3′,2′h][1,5] naphthyridie-5,10-dione) PTB7: poly[[4,8-bis[(2-ethylhexyl)oxy]benzo[1,2-b:4,5-b']dithiophene-2,6-diyl][3-fluoro-2-[(2ethylhexyl)carbonyl]thieno[3,4-b]thiophenediyl]] PTB7-Th: poly[4,8-bis(5-(2-ethylhexyl)thiophen-2-yl)benzo[1,2-b;4,5-b’]dithiophene-2,6diyl-alt-(4-(2-ethylhexyl)-3-fluorothieno[3,4-b]thiophene-)-2-carboxylate-2-6-diyl)]



P(NDI2OD-T2) or N2200: poly([N,N′-bis(2-octyldodecyl)-11 naphthalene-1,4,5,8bis(dicarboximide)-2,6-diyl]-alt-5,5′-(2,2′-12 biothiopene)) REFERENCES 1. Bao, Z.; Chen, X., Flexible and Stretchable Devices. Adv. Mater. 2016, 28 (22), 41774179. 2. Heeger, A. J., Semiconducting polymers: the Third Generation. Chem. Soc. Rev. 2010, 39 (7), 2354-2371. 3. Dimitrakopoulos, C. D.; Mascaro, D. J., Organic thin-film transistors: A review of recent advances. IBM J. Res. Dev. 2001, 45 (1), 11-27. 4. Choi, H. Y.; Kim, S. H.; Jang, J., Self-Organized Organic Thin-Film Transistors on Plastic. Adv. Mater. 2004, 16 (8), 732-736. 5. Dimitrakopoulos, C. D.; Malenfant, P. R. L., Organic Thin Film Transistors for Large Area Electronics. Adv. Mater. 2002, 14 (2), 99-117. 6. Benavides, C. M.; Murto, P.; Chochos, C. L.; Gregoriou, V. G.; Avgeropoulos, A.; Xu, X.; Bini, K.; Sharma, A.; Andersson, M. R.; Schmidt, O.; Brabec, C. J.; Wang, E.; Tedde, S. F., High-Performance Organic Photodetectors from a High-Bandgap Indacenodithiophene-Based πConjugated Donor–Acceptor Polymer. ACS Appl. Mater. Interfaces 2018, 10 (15), 12937-12946. 7. Murto, P.; Genene, Z.; Benavides, C. M.; Xu, X.; Sharma, A.; Pan, X.; Schmidt, O.; Brabec, C. J.; Andersson, M. R.; Tedde, S. F.; Mammo, W.; Wang, E., High Performance AllPolymer Photodetector Comprising a Donor–Acceptor–Acceptor Structured Indacenodithiophene–Bithieno[3,4-c]Pyrroletetrone Copolymer. ACS Macro Letters 2018, 7 (4), 395-400. 8. Sharma, A.; Kroon, R.; Lewis, D. A.; Andersson, G. G.; Andersson, M. R., Poly(4vinylpyridine): A New Interface Layer for Organic Solar Cells. ACS Appl. Mater. Interfaces 2017, 10929-10936. 9. Krebs, F. C., Fabrication and processing of polymer solar cells: A review of printing and coating techniques. Sol. Energy Mater. Sol. Cells 2009, 93 (4), 394-412. 10. Snyder, C. R.; DeLongchamp, D. M., Glassy phases in organic semiconductors. Curr. Opin. Solid State Mater. Sci. 2018, 22 (2), 41-48. 11. Noriega, R.; Rivnay, J.; Vandewal, K.; Koch, F. P. V.; Stingelin, N.; Smith, P.; Toney, M. F.; Salleo, A., A general relationship between disorder, aggregation and charge transport in conjugated polymers. Nature Materials 2013, 12, 1038. 12. Lindqvist, C.; Bergqvist, J.; Bäcke, O.; Gustafsson, S.; Wang, E.; Olsson, E.; Inganäs, O.; Andersson, M. R.; Müller, C., Fullerene mixtures enhance the thermal stability of a noncrystalline polymer solar cell blend. Appl. Phys. Lett. 2014, 104 (15), 153301. 13. Ma, W.; Yang, C.; Gong, X.; Lee, K.; Heeger, A. J., Thermally Stable, Efficient Polymer Solar Cells with Nanoscale Control of the Interpenetrating Network Morphology. Adv. Funct. Mater. 2005, 15 (10), 1617-1622. 14. Bergqvist, J.; Lindqvist, C.; Backe, O.; Ma, Z.; Tang, Z.; Tress, W.; Gustafsson, S.; Wang, E.; Olsson, E.; Andersson, M. R.; Inganas, O.; Muller, C., Sub-glass transition annealing enhances polymer solar cell performance. J. Mater. Chem. A 2014, 2 (17), 6146-6152. 15. Müller, C., On the Glass Transition of Polymer Semiconductors and Its Impact on Polymer Solar Cell Stability. Chem. Mater. 2015, 27 (8), 2740-2754.


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TOC Image

Polymer Morphology

-100

Sub-Tg

0

Tm

No-Tg 100

200

tan δ

Hairy Aggregates

E'

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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300

Temperature (°C)

DMTA Thermogram


Probing the Relationship between Molecular Structures, Thermal

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