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Jun 27, 2017 - rheometer offer a solution to the challenges mentioned above for the measurement of the glass transition temperature of conjugated poly...
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Glass Transition Temperature of Conjugated Polymers by Oscillatory Shear Rheometry Renxuan Xie,† Youngmin Lee,† Melissa P. Aplan,† Nicholas J. Caggiano,† Christian Müller,∥ Ralph H. Colby,*,†,‡,§ and Enrique D. Gomez*,†,§ †

Department of Chemical Engineering, ‡Department of Materials Science and Engineering, and §The Materials Research Institute, The Pennsylvania State University, University Park, Pennsylvania 16802, United States ∥ Department of Chemistry and Chemical Engineering, Chalmers University of Technology, 41296 Göteborg, Sweden S Supporting Information *

ABSTRACT: The stiff backbones of conjugated polymers can lead to a rich phase behavior that includes both crystalline and liquid crystalline phases, making measurements of the glass transition challenging. In this work, the glass transitions of regioregular poly(3-hexylthiophene-2,5-diyl) (RR P3HT), regiorandom (RRa) P3HT, and poly((9,9-bis(2-octyl)-fluorene-2,7-diyl)-alt-(4,7-di(thiophene-2-yl)-2,1,3benzothiadiazole)-5′,5″-diyl) (PFTBT) are probed by linear viscoelastic measurements as a function of molecular weight. We find two glass transition temperatures (Tg’s) for both RR and RRa P3HT and one for PFTBT. The higher Tg, Tα, is associated with the backbone segmental motion and depends on the molecular weight, such that the Flory−Fox model yields Tα = 22 and 6 °C in the long chain limit for RR and RRa P3HT, respectively. For RR P3HT, a different molecular weight dependence of Tα is seen below Mn = 14 kg/mol, suggesting this is the typical molecular weight of intercrystal tie chains. The lower Tg (TαPE ≈ −100 °C) is associated with the side chains and is independent of molecular weight. RRa P3HT exhibits a lower Tα and higher TαPE than RR P3HT, possibly due to a different degree of nanophase separation between the side chains and the backbones. In contrast, PFTBT only exhibits one Tg above −120 °C, at 144 °C in the long chain limit.

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scanning calorimetry (DSC)6,7,11,18 and dynamic mechanical analysis (DMA).10,19−21 The application of either DSC or DMA to measure the Tg of semicrystalline conjugated polymers has significant limitations. The glass transition is challenging to measure when a crystalline phase is present,10,11 such that Tg relies on the response from the amorphous phase in semicrystalline polymers. Rapid heat/ cool DSC (ΔT/Δt ∼ 500 K/s) allows amplification of the signal but requires that the frequency dependence is taken into account.22−25 The step change in heat flow by DSC could be either too small to accurately locate Tg or be mistaken for other transitions (e.g., cold crystallization, liquid-crystal-to-isotropic transition, and melting). Typical DMA experiments require 100 mg of sample or more (except for thin film tensile DMA, which requires about 20 mg). Nevertheless, when trying to load conjugated polymer thin films under tension at room temperature, the sample could break easily if Tg is either at or above room temperature and if the molecular weight is too low, as such samples are brittle. This limitation has been addressed with a recently developed unconventional sample loading technique for DMA, which reinforces the conjugated polymer with either a steel pocket10 or woven glass fibers.26

INTRODUCTION Conjugated polymers have great potential in future flexible electronics applications because of their delocalized π orbitals along a covalently linked backbone that allows either photoinduced or electric-field-induced charges to conduct within a stretchable network.1−3 The importance of intermolecular electronic coupling that allows for macroscopic charge transport leads to a strong dependence on the degree of order, such that tuning processing can be important to maximize performance when incorporating conjugated polymers in electronic devices. Thus, a key challenge in controlling the microstructure of the active layer and the degree of phase separation between donor and acceptor molecules4 in polymeric electronic devices is identifying the glass transition temperature (Tg), because chain motion is governed by the amount of superheating above Tg. Nevertheless, reported values for the glass transition temperature vary significantly, for example, from −14 to 140 °C for poly(3-hexylthiophene-2,5-diyl) (P3HT).5−15 Part of the challenge is that the dependence of Tg on the molecular weight, which is predicted by the Flory−Fox equation,16 has not been previously considered. Such a large range of reported Tg values cannot be solely attributed to the effect of molecular weight, but is perhaps a result of inconsistent thermal history,5,8,9,11 the effects of regioregularity,17 and the inherent limitations of the common Tg characterization techniques, namely differential © XXXX American Chemical Society

Received: April 4, 2017 Revised: June 6, 2017

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DOI: 10.1021/acs.macromol.7b00712 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules Table 1. Characteristics of the Eight Regioregular and Four Regiorandom P3HTs Used in This Study polymer

1

2

3

4

5

6

7

8

RRa 1

RRa 2

RRa 3

RRa 4

Mwa [kg/mol] Đ from GPC Mnb [kg/mol] regioregularity xcc

3.1 1.14 2.7 88% 0.25

6.7 1.26 5.3 93% 0.59

16.7 1.15 14.5 93% 0.48

20.7 1.35 15.3 94% 0.47

29.5 1.22 24.2 95% 0.55

38.3 1.78 21.5 94% 0.45

50.9 1.84 27.7 98% 0.45

67.0 1.83 36.6 96% 0.47

11.4 2.06 5.5 61% 0

25.2 2.04 12.3 61% 0

101 2.42 41.7 58% 0

110 2.45 44.9 59% 0

Mw is determined by a combination of Mn from 1H NMR and Đ ≡ Mw/Mn from GPC for batches 1 and 2 and by the static light scattering (SLS) for all other P3HTs. bMn is determined by 1H NMR directly for batches 1 and 2 and by a combination of Mw from SLS and Đ from GPC for all other P3HTs. cCrystallinity

a

precipitations in methanol followed by Soxhlet extractions in methanol, acetone, and hexane. Using a previously reported method,29 regioregularities for both RR and RRa P3HTs were calculated based on integration of α-methylene protons on the alkyl side chain for both head-to-tail and head-to-head configuration via solution 500 MHz 1H NMR as shown in Table 1. PFTBT was synthesized by a Suzuki polycondensation reaction.30 Molecular weight was modulated by varying the molar ratio of monomers, 9,9-dioctylfluorene-2,7-diboronic acid bis(1,3-propanediol) ester (F) and 4,7-bis(2-bromo-5-thienly)-2,1,3-benzothiadiazole (TBT), according to Carothers’ equation.31 All batches were synthesized by the same procedure, adjusting only the amount of TBT added in order to modulate the molecular weight. A 100 mL Schlenk flask equipped with a stir bar was charged with 12 mL of anhydrous toluene, 168 mg (0.3 mmol) of F, and TBT. For low molecular weight PFTBT, 103 mg (0.23 mmol) of TBT was added; for medium molecular weight PFTBT, 121 mg (0.26 mmol) of TBT was added; for high molecular weight PFTBT, 138 mg (0.3 mmol) of TBT was added. The solution was stirred and purged with argon for approximately 20 min, at which time 3.6 mL of tetraethylammonium hydroxide (20% in water), 2 drops of Aliquat 336, and 26 mg (0.02 mmol) of tetrakis(triphenylphosphine)palladium(0) (Pd(PPh3)4) were added. The mixture was purged for an additional 20 min and then degassed by three freeze−pump−thaw cycles. The solution was heated to 80 °C and allowed to react for 24 h. The polymerization was quenched with 4 mL of bromobenzene and stirred at 80 °C for an additional hour. The reaction mixture was then cooled to room temperature, and the aqueous phase was separated and discarded. The raw product was precipitated into 500 mL of methanol, filtered, and purified by sequential Soxhlet washes with methanol, acetone, and hexane. The purified product was dried under vacuum for 24 h. Yields were measured as 148 mg (80%), 170 mg (87%), and 143 mg (69%) for low, medium, and high molecular weight PFTBT, respectively. The alkyl side chains are either regularly (RR P3HT) or randomly (RRa P3HT) attached to the thiophene backbone. As a result, RR P3HT can crystallize via π−π stacking of the thiophene rings while RRa P3HT remains completely amorphous. Nevertheless, in both cases, some amorphous chains are still present, although it is harder to detect the glass transition for RR P3HT than for RRa P3HT. All RR P3HT samples in this work melt in the range of 210−240 °C except for batch 1 (Tm = 160 °C) but show no clear indication of glass transitions in DSC traces (Figure S1 of the Supporting Information). Weight-average molecular weights for all P3HTs except for 1 and 2 were determined via static light scattering using a Brookhaven BI200SM SLS/DLS instrument over an angular range of 40°−140° using a 637 nm diode laser (Figure 1). All solutions of P3HT in chlorobenzene (CB) were prepared at concentrations from 3.0 to 0.8 mg/mL, well below their overlap concentrations, and then filtered through 0.2 μm PVDF filters multiple times. The solutions were maintained at 30.0 ± 0.2 °C during the scattering experiment. The differential refractive index (dn/dc) was determined to be 0.179 and 0.155 mL/g respectively for RR and RRa P3HTs in CB solutions at 30.0 °C using a Brookhaven differential refractometer in static mode with a laser source of 620 nm (Figure S2a). The expected molecular weights for 1 and 2 are well below 10 kg/mol, which usually is the lower limit for SLS experiments. For these two polymers, high resolution 1H NMR was used instead to determine the number-

This approach, however, leads to challenges in accurately extracting values for the modulus due to the undefined sample geometry and possible void space within the sample. Based only on the phase angle, it would be difficult to distinguish among various possible thermal transitions for conjugated polymers. UV−vis spectra have also been used to determine the Tg of conjugated polymers by carefully tracking changes with temperature of the absorption spectrum due to the molecular-scale rearrangement,27 but the origin of such changes may not be limited to the glass transition process. Linear viscoelastic measurements made in a rotational rheometer offer a solution to the challenges mentioned above for the measurement of the glass transition temperature of conjugated polymers. Only 10 mg of sample is needed using a 3 mm diameter parallel plate geometry. Also, loading the sample only requires placing a molded sample puck between the parallel plates followed by annealing above Tm to ensure good adhesion between the plates and the sample. Finally, rheometers can not only easily distinguish between glass transitions, melting, cold crystallization, and nematic to isotropic transitions based on the absolute moduli values and their temperature dependences, but also can sometimes probe the extent of nanophase separation within conjugated polymers. In this work, the number-average molecular weights (Mn) of various P3HTs were determined by the combination of dilute solution static light scattering, 1H NMR, and gel permeation chromatography. Through subambient linear shear rheology, the dependence of Tg on Mn is found to follow the Flory−Fox equation, from which the infinite molecular weight limit of Tg (a material-specific constant) is extrapolated for both regioregular (RR) and regiorandom (RRa) P3HT. Also explored by linear viscoelastic measurements is the impact of regioregularity on the extent of formation of separate domains by the side chains and backbones (i.e., nanoscale phase separation) in P3HTs. Finally, we illustrate that this technique is also applicable to other conjugated polymers, such as poly((9,9-dioctylfluorene)-2,7-diyl-alt-[4,7-bis(thiophen-5-yl)2,1,3-benzothiadiazole]-2′,2″-diyl) (PFTBT), where rheometry clearly identifies the glass transition despite the complex phase behavior. As a result, we are able to clarify the ambiguities in the literature regarding the glass transition temperatures for P3HTs through a systematic study of molecular weight and regioregularity effects.

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MATERIALS AND METHODS

Higher molecular weight RR P3HT batches (4, 5, 6, 7, 8) were purchased from Merck and used without further purification. Lower molecular weight RR P3HTs (1, 2, 3) were synthesized using established Grignard metathesis procedures.28 Regiorandom (RRa) P3HTs (RRa 3 and RRa 4) were purchased from Sigma-Aldrich while the lower molecular weight RRa 1 and RRa 2 were synthesized using the Rieke method.29 All RRa P3HTs were purified by multiple B

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of fusion per crystalline repeat unit of 49 J/g and an equilibrium melting temperature of 272 °C as previously described.32 Samples for rheometry require molding into bubble-free 1 mm thick and 3 mm or 8 mm diameter disks. To limit the potential for either photo or thermal degradation,33 samples were molded inside a nitrogen-filled glovebox (H2O < 1 ppm and O2 < 10 ppm). Heating samples above the melting temperature of our polymers (260 °C) under vacuum eliminated air bubbles. Then, samples were compressed into disks with about 1 MPa of pressure. Molded polymer pucks were loaded into an ARES-LS operating under a nitrogen environment by first heating to 250 °C to ensure good contacts. Samples were then stretched to an hourglass shape with a center diameter of roughly twothirds of the plates in order to avoid compliance effects with the transducer when cooling down to −120 °C. The oscillatory frequency sweep test was rerun at 40 °C to determine the effective rheometry geometry constant (stress/torque) of this stretched sample by matching with the response from the 3 mm diameter unstretched sample.34

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RESULTS AND DISCUSSION The mechanical response of polymers can provide clear signatures of various processes, such as the melting from a crystalline or glassy state. A crucial component is to combine measurements at various frequencies and temperatures. We then generate master curves at Tref = 0 °C by frequency-scale shifting the isothermal frequency sweep results at strain amplitude of 0.001 from 40 to −120 °C (−10 °C increments), as shown in Figure 2 for RR P3HT 6 and RRa 4. Two relaxation processes are clear. The lower frequency process (around 10−2 rad/s) has been attributed to the chain segmental relaxation (α process), while the higher frequency relaxation (around 109 rad/s) is associated with the hexyl side chain relaxation (αPE process); this classification was verified by studying the dependence of interlayer spacing on the side chain length from 6 to 12 carbons of P3ATs.20 Furthermore, the temperature range for the αPE process in P3HT is similar to that from alkyl side-chain relaxations of other polymers.35 Time−temperature superposition (tTs) works well for the α process (between 40 and −30 °C) and then fails for the αPE process (between −30 and −120 °C) as shown in Figures 2a and 2b. The mismatch of tan(δ) between the successive isothermal frequency sweep results is a clear signature for tTs failure. This is caused by the drastic difference in temperature dependence of the thiophene backbone relaxation versus that of the alkyl side chain relaxation. Because tTs is based on the assumption that the relaxation time associated with each mode has the same temperature dependence, tTs fails when both the backbone and the side chain relaxations contribute significantly to the measured torque. Therefore, for the α process between 40 and −30 °C, the side chain nanophase is still very “soft” (roughly 90 °C above the αPE process), such that the measured torque due to the side chains is small in comparison. When temperature is decreased below −40 °C, the side chain motion slows down, so both relaxations contribute significantly to the torque, and tTs fails as expected. The backbone segmental relaxation is confirmed to be associated with the observed α process by fitting the temperature dependence of the frequency-scale shift factors (aT) with the Williams−Landel−Ferry (WLF) equation:36

Figure 1. Zimm plots generated by static light scattering of dilute chlorobenzene solutions of (a) RRa 4 (110 ± 2 kg/mol, Rg = 15.4 ± 1.1 nm) and (b) polymer 8 (67 ± 2 kg/mol, Rg = 14.2 ± 2.1 nm). The weight-average molecular weight (Mw) was calculated from the intercept of zero angle and zero concentration extrapolations in the plot, while Rg was obtained from the slope of the zero concentration extrapolation (K: optical constants, 2.89 × 10−7 for RR and 2.17 × 10−7 mol mL/g2 for RRa; c: polymer concentration (g/mL); Rθ: Rayleigh ratio; θ: scattering angle). average molecular weight by calculating the area ratio of the αmethylene protons on the alkyl side chain between the backbone and the end groups (both H-end and Br-end). NMR-derived molecular weights are consistent with SLS measurements of RR P3HTs 3−8. The larger dn/dc value for RR P3HT suggests that the polymer is more polarizable than RRa P3HT, confirmed by the larger dielectric constant at −40 °C in Figure S2b (εRR/εRRa = 3.25/2.82). Dielectric relaxation spectroscopy (DRS) measurements for P3HTs were measured with thicknesses of 50 μm in between two polished brass electrodes. A voltage amplitude of 1.5 V for RRa or 0.5 V for RR was used at various temperatures under nitrogen. Frequency range spans from 107 to 0.1 Hz at each temperature. Dispersities of P3HTs were measured using gel permeation chromatography (GPC) calibrated relative to narrow polystyrene standards with chlorobenzene as the eluent as shown in Figure S3. A high temperature GPC coupled with a viscometer detector was used to characterize the molecular weight distribution of PFTBT at 150 °C in 1,2,4-tricholorobenzene using a universal calibration after complete dissolution overnight and filtering, as shown in Figure S4 (SLS experiments for PFTBTs were not possible due to the absorption of the polymer). Mn values for P3HTs were calculated based on the Mw from SLS and the dispersity (Đ ≡ M w /M n ) from GPC (chlorobenzene) as summarized in Table 1. Furthermore, the crystalline mass fraction for RR P3HT was estimated from the second heating DSC trace at 20 °C/min shown in Figure S1 using an enthalpy

log(aT ) =

C(Tr − T ) T − T∞

(1)

where Tr is the reference temperature at 0 °C, C is an empirical fitting parameter, and T∞ is the Vogel temperature where the C

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Figure 2. Two examples of master curves for (a) storage and loss moduli and (b) tan(δ) for RR P3HT 6 and RRa 4. Tref of 0 °C. Strain amplitude of 0.001. (c) Frequency shift factors (aT) used to generate master curves shown in (a), (b), and (d). Equation 1 is fitted to aT for RRa 3 at temperatures just above Tα, which should be between 0 and 10 °C where the data deviate from the WLF fit. Below Tα the activation energy is 65 ± 3 kJ/mol. (d) Master curve for storage and loss moduli for RRa 3 P3HT from the terminal region to the glassy region. Tref of 0 °C.

glass transition temperatures for all P3HTs in this work are shown in Table S2. From Figure 3b, it is clear that the lower Tg does not depend on molecular weight, supporting the conclusion that the side chain relaxation is solely responsible for the αPE process. The temperature range of the αPE process is similar between RR and RRa P3HTs, suggesting a similar origin of the dynamics within the hexyl nanodomains. Temperature-dependent charge mobility studies on RR P3HTs have shown a significant deviation from an Arrhenius temperature dependence near 180 K; often charge mobilities are higher than expected at low temperatures.40−43 We speculate that arresting the side chain motion below 180 K can affect charge transport by affecting electronic coupling and charge transfer integrals between chains. For example, freezing the side chains could lead to a decrease in the thermal expansion coefficient or a drop in thermal fluctuations. The dynamics of side chains and backbones are not completely independent from each other because side chains can plasticize the backbone. 13C solid state NMR data suggest that hexyl side chains may induce a twist like transition of the thiophene rings

free volume is zero. The master curve that spans from the terminal region to the side chain relaxation is constructed for regiorandom P3HT in Figure 2d. The frequency shift factors between 260 and 10 °C are well fitted by eq 1 with C = 15.9 ± 0.1 and T∞ = 209 ± 1 K in Figure 2c. The deviation between the fit and the data starts between 0 and 10 °C at the glass transition (or Tα) and when the shift factors transition from WLF to an Arrhenius dependence associated with monomer relaxation. Nevertheless, we cannot construct a similar master curve from the terminal region to the glassy region for RR P3HTs due to the semicrystalline response that masks Rouse relaxation and chain entanglement. After equilibration at −120 °C, Figure 3 shows the temperature ramp result from −120 to 40 °C at a fixed frequency of 10 rad/s, heating rate of 5 °C/min, and commanded strain amplitude of 0.001. The glass transition temperature can be quantified from rheometry data by the local maximum in G′′37 because the loss modulus defines the relaxation process.38 Sometimes the peak in tan δ is used, leading to higher Tg values.10,39 A tabulated summary of the D

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relative intensity of the α process for RRa P3HTs because the α process occurs only in the amorphous phase. The nanoscale separation of the side chains can also explain the lower value for G′ observed near −40 °C for RR P3HT in Figure 3a. At temperature between Tα and TαPE, the thiophene backbone segmental motion is “frozen” while the hexyl side chains are mobile, such that stress can only be transmitted through stacked backbones within crystals and intercrystalline tie chains and entanglements. Thus, G′ for RR depends on contributions from the crystalline structure and the glassy modulus. Nevertheless, the mobile side chain domains could “soften” the storage modulus. The stronger nanophase separation of the side chains, denoted by the higher αPE peak in G′′ of RR P3HT, leads to a 70% lower G′ after the αPE side chain relaxation compared to RRa P3HT, as shown in Figure 4. Coarsening of immiscible polymer blends45 and macroscopic alignment of block copolymer by LAOS46 have also been shown to lower G′.

Figure 4. ΔG′ for the αPE side chain glass transition as a function of molecular weight. ΔG′ was calculated between the two local minima of G″ of the master curves generated for RR and RRa P3HTs shown in Figure 2a.

Figure 3. (a) Storage modulus and (b) loss modulus as a function of temperature during heating from −120 to 40 °C at fixed frequency of 10 rad/s, strain amplitude of 0.001, and heating rate of 5 °C/min. Only from −35 to 40 °C for RR P3HT 1 and 2.

that affects π−π stacking.44 Furthermore, as shown in Figure 3b, the relative intensity of G′′ for the αPE process differs between RR and RRa P3HTs. The larger magnitude of G″ in the αPE process for semicrystalline RR P3HTs apparent from Figure 3b implies stronger nanophase separation of the hexyl nanodomain from the thiophene ring backbones than that of RRa P3HTs. The larger G″ could be attributed to the strong π−π interchain stacking within crystals in RR P3HTs, which itself induces formation of larger and more regular packing of hexyl domains between the stacked layers of thiophene rings. Moreover, the broader αPE transition region for RRa P3HT (fwhm = 76 ± 5 °C versus 42 ± 4 °C for RR P3HT) is consistent with more diffuse interfacial regions of the hexyl side chain domains. RRa 1 with the lowest molecular weight displays the broadest αPE transition region, indicating even weaker nanophase separation between the side chain and the backbone. Nevertheless, the larger amount of amorphous material also leads to a stronger

When cooling further down to −120 °C, nearly all side chain motion is frozen into a glassy state, leading to a rise in G′. Consistent with our hypothesis that the stronger side chain nanophase separation leads to a stronger drop in G′ (larger ΔG′ in Figure 4) in RR P3HT, once side chain motion is arrested, G′ increases more strongly in RR P3HT. The mechanical response of RR P3HT is composed of contributions from crystalline domains in addition to the glassy modulus, leading to consistently higher G′ for RR P3HT at temperatures below −120 °C regardless of molecular weight (Figure 3a). In contrast, Figure 3b shows that the α glass transition temperature depends on molecular weight. Figure 5 plots Tα and TαPE against the inverse of the number-average molecular weight (Mn) for both RR and RRa P3HTs. Tα clearly depends on molecular weight, and this dependence is well described by the Flory−Fox equation16 based on chain ends having extra free volume E

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mobilities saturate after about Mn = 10−20 kg/mol.47−53 Furthermore, RR P3HT 1 and 2 (Mn < 14 kg/mol) fracture under very small compression force when cooling below −40 °C in the rheometer, possibly due to the lack of tie chains. Therefore, the elevated Tα below 14 kg/mol for RR P3HT is not likely due to tie chains but may instead be due to incorporation of chain ends in crystals that lowers the number density of mobile chain ends and slows dynamics significantly. In addition, the presence of a rigid amorphous phase at a crystal−amorphous interface region could further raise Tα, and its response can be separated from the mobile amorphous phase from analysis of TMDSC data.54 Tα’s for RRa P3HTs from rheology agree with those determined by dielectric relaxation spectroscopy (DRS) as shown in Figure 5 and Table S2. The dielectric strength for the α transition of RRa P3HT is 0.16. For RR P3HT, a peak in the dielectric loss is not apparent, likely due to the higher conductivity response that is about 3 orders of magnitude higher than for RRa P3HT (see details in Figure S5). Figure 5 also summarizes TαPE for RR P3HTs and RRa P3HTs at various molecular weights. As expected for a relaxation attributed to the side chains, TαPE is independent of molecular weight. Low molecular weight samples were too brittle for rheology measurements of TαPE because of fracture below about −40 °C. TαPE is about 5−6 °C lower for RR P3HT when compared to RRa P3HT, which we attribute to stronger side chain nanophase separation in regioregular polymers. To illustrate the effectiveness of probing the glass transition of conjugated polymers through linear viscoelastic measurements, we also examined a push−pull alternating copolymer, poly((9,9-dioctylfluorene)-2,7-diyl-alt-[4,7-bis(thiophen-5-yl)2,1,3-benzothiadiazole]-2′,2″-diyl) (PFTBT). Figure 6 shows

Figure 5. Molecular weight dependence of glass transition temperatures for both RR and RRa P3HTs. Star symbols represent Tα measured by DRS, while square symbols represent rheology results measured at 10 rad/s and heating rate of 5 °C/min. Green symbols correspond to the backbone segmental relaxation temperature, Tα, and orange symbols represent the side chain relaxation temperature, TαPE. Open symbols refer to RR P3HTs, and filled symbols refer to RRa P3HTs. Green lines (both dash and solid) correspond to the linear fits of the Flory−Fox eq 2 for RR and RRa P3HTs. Orange dashed and solid lines represent the average TαPE’s that are independent of Mn for RR P3HT (−101 ± 2 °C) and RRa P3HT (−94 ± 3 °C), respectively.

Tα = Tα , ∞ −

K Mn

(2)

where Tα,∞ is Tα extrapolated to the infinite molecular weight limit and K is a material specific constant. For RR P3HTs, the molecular weight dependence of Tα can be clearly divided into two regions, which are modeled separately with two Flory−Fox equations with fitting parameters shown in Table 2. One Table 2. Flory−Fox Equation Parameters for RR P3HT and RRa P3HT RR P3HTs (Mn > 14 kg/mol) RR P3HTs (Mn < 14 kg/mol) RRa P3HTs

Tα,∞ [°C]

K [K kg/mol]

22.1 ± 1.2 2.9 ± 0.2 5.9 ± 1.0

291 ± 23 22 ± 1 100 ± 10

Flory−Fox equation is sufficient to describe the Mn dependence of Tα for RRa P3HT. Although we expect glass transition temperatures to depend on frequency and thermal history, such studies are beyond the scope of this work and are not included in Figure 5. Figure 5 shows that Tα for RR P3HT is consistently higher than Tα for RRa P3HT. We hypothesize that the crystallization of RR P3HT leads to some chain ends incorporated within crystals and stretched tie chains with reduced conformational entropy, leading to higher glass transition temperatures. Because the crystallinity is similar at about 50% for RR P3HTs 3, 4, 5, 6, 7, and 8 (Mn > 14 kg/mol), we attribute the abrupt change in the temperature dependence of Tα at Mn of 14 kg/mol for RR P3HT to the formation of tie chains at this molecular weight. The formation of chains interconnecting crystallites at 14 kg/mol is consistent with the dependence of the charge mobility on molecular weight, where charge

Figure 6. DSC traces for the three PFTBTs after isothermal crystallization at 195 °C for 30 min, at a heating rate of 20 °C/min.

the DSC traces of PFTBT with three different molecular weight distributions (see Table 3 and Figure S4 for details), which were isothermally crystallized at 195 °C for 30 min in the DSC prior to the heating scans. Both in situ heating wide-angle X-ray Table 3. Glass Transition Temperatures of PFTBTs

F

polymers

PFTBT 1

PFTBT 2

PFTBT 3

Mn by GPC [kg/mol] Đ by GPC Tg by DSC, 20 °C/min [°C] Tg by rheology, 10 rad/s [°C] ± 1 Tg by rheology, 1.0 rad/s [°C] ± 2

16.0 4.23 144 140 132

11.7 4.49 133 132 126

9.3 3.93 129 128 122

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Macromolecules scattering (WAXS) and rheology confirm the cold crystallization process of PFTBT at 195 °C during heating for the sample that was previously quenched from 270 °C to room temperature (Figure S6 and Figure 7). Three features in the DSC scans in Figure 6 are clearly identified, near 130, 210, and 270 °C; the highest is Tm, but the other two are more challenging to assign.

Figure 7. Storage and loss moduli versus temperature for PFTBT 2 at 1 rad/s for samples with two different crystallization procedures: quenched from 270 °C to room temperature (Quenched) and isothermally crystallized at 195 °C for 30 min before cooling to room temperature (Crystallized).

Figure 8a shows G′ and G″ for three PFTBT samples as a function of temperature. An α transition between 120 and 150 °C can be assigned as the glass transition due to the relatively sharp change in G′ between 600 and 30 MPa. This Tg is consistent with previously reported values obtained from amorphous PFTBT.55 Nevertheless, previous reports did not consider the effect of crystal confinement on the glass transition, and this effect is shown to be significant for PFTBT in Figure 7, where Tg differs by 15 °C depending on the thermal history of the sample. Specifically, quenching the sample to room temperature to suppress crystallization leads to a lower Tg when compared to a sample that was isothermally crystallized above Tg. This again suggests that stretching of tie chains reduces the entropy of the intercrystalline amorphous chains and thereby increases Tg. Thus, we assign the transition between 120 and 144 °C in DSC to a Tg and compare with values obtained from rheometry in Table 3. Values from DSC and rheometry are in reasonable agreement. With the Flory− Fox equation, we find Tg = 144 ± 3 °C at infinite molecular weight for semicrystalline PFTBT as shown in Figure 8b. At the transition temperature of around 210 °C identified by DSC (Figure 6), G″ shows a local maximum while G′ keeps decreasing from 80 MPa to 20 MPa. We speculate that these values for G′ are too low for PFTBT to be a glass at these temperatures. Therefore, this transition might be due to melting of the smaller crystals, perhaps due to large dispersities with bimodal distributions of the polymers, as shown in Figure S4. A key difference between P3HT and PFTBT is the absence of TαPE in PFTBT despite the longer alkyl side chains in

Figure 8. (a) Temperature ramp results at frequency of 1.0 rad/s and heating rate of 5 °C/min, after isothermally crystallizing at 195 °C for 30 min. For temperatures below 220 °C, 3 mm diameter parallel plates were used, while 8 mm plates were used above 220 °C. Only data above 80 °C are shown for PFTBT 1 due to the loss of adhesion on the parallel plates at lower temperatures. (b) Molecular weight dependence of Tg determined from the G″ peak at frequency of 1.0 rad/s for isothermally crystallized PFTBTs. The dashed line is a fit to the Flory−Fox equation. Tg determined from rheomery at 10 rad/s and by DSC are also shown for comparison.

PFTBT. The mass fraction of side chains for P3HT and PFTBT are 0.52 and 0.33, respectively; the absence of TαPE in PFTBT might be due to smaller nanophase separated side chain domains than P3HT due to the differences in side chain fraction. Nevertheless, PFTBT appears to exhibit nanophase separation, as an amorphous halo at 1.4 Å−1 from WAXS data with a backbone spacing peak at 0.4 Å−1 (Figure S6). We therefore propose three possible explanations. The side chain nanodomains for PFTBT may not be large enough to exhibit its own TαPE, TαPE in PFTBT may be below −120 °C, or the side chains may have more crystalline order than in P3HT, such that αPE is not observed. G

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CONCLUSIONS In summary, we have demonstrated that linear viscoelastic rotational rheometry offers unique advantages over traditional DSC and DMA experiments to examine the glass transition of conjugated polymers. The small sample quantity requirement and straightforward sample loading coupled with a sensitive response to mechanical properties as a function of temperature are well suited for clearly identifying subtle changes in chain dynamics and thus revealing glass transitions due to vitrification of the backbone (α process) and side chain motion (αPE process). Tα for RR P3HT depends on the molecular weight and follows the Flory−Fox equation with two distinct regions, above and below Mn = 14 kg/mol. Above this Mn, Tα varies between 1 and 15 °C, yielding Tα,∞ = 22 °C in the long chain limit. Below Mn = 14 kg/mol, Tα exhibits a weaker dependence on Mn. We speculate that this critical Mn is the threshold value for the formation of tie chains between P3HT crystals. Tα for RRa P3HT is lower than RR P3HT across all measured molecular weights and also follows the Flory−Fox equation, yielding Tα,∞ of 5.9 °C in the long chain limit. The glass transition temperature of the side chains, TαPE, is independent of molecular weight and is −101 °C for RR P3HT and −94 °C for RRa P3HT, where we attribute the difference to weaker nanophase separation in RRa P3HT. We also examined a semicrystalline push−pull alternating copolymer, PFTBT, where we found Tg between 122 and 144 °C and Tg,∞ = Tα,∞ = 144 °C in the long chain limit.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.7b00712.

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Figures S1−S6 and Tables S1, S2 (PDF)

AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected] (R.H.C.). *E-mail: [email protected] (E.D.G.). ORCID

Renxuan Xie: 0000-0001-8015-8550 Christian Müller: 0000-0001-7859-7909 Ralph H. Colby: 0000-0002-5492-6189 Enrique D. Gomez: 0000-0001-8942-4480 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors gratefully acknowledge the financial support of the National Science Foundation under grant DMR-1629006. This research used resources of the Advanced Light Source, which is a DOE Office of Science User Facility under contract DEAC02-05CH11231. We thank Anders Mårtensson for help with GPC measurements.

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