Molecular Origin of Charge Traps in Polyfluorene-Based

Sep 16, 2013 - The comprehensive control of morphology and structure is of extreme importance in semiconducting polymers when used as active layers in...
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Molecular Origin of Charge Traps in Polyfluorene-Based Semiconductors Gregório C. Faria,†,‡,* Eduardo R. deAzevedo,† and Heinz von Seggern‡ †

Instituto de Física de São Carlos, Universidade de São Paulo, P.O. Box: 369, 13560-970, São Carlos, SP, Brazil Institute of Materials Science, Technische Universitaet Darmstadt, Petersenstrasse 23, 64287 Darmstadt, Germany

‡

ABSTRACT: The comprehensive control of morphology and structure is of extreme importance in semiconducting polymers when used as active layers in optoelectronic devices. In the work reported here, a systematic investigation of the structural and dynamical properties of poly(9,9-di-n-octyl-fluorene-alt-benzothiadiazole), known as F8BT, and their correlation with electrical properties is presented when the material is used as an active layer in optoelectronic devices. By means of X-ray diffraction, one observes that in thick layer films (thickness of about 4 μm) grown by drop-cast deposition, a solvent induced crystalline phase exists which evolves to a stable phase as the temperature is raised. This was not observed in thin films (thickness of about 250 nm) prepared by spin-coating within the investigated temperature range. By modeling the current−voltages characteristics of both thick and thin film devices, important information on the influence of crystallization on the trapping states could be drawn. Furthermore, the temperature dependence of the charge carrier mobility was found to be closely related to that of the molecular relaxation processes. The understanding of the nature of such molecular relaxations, measured by solid-state nuclear magnetic resonance methods, allows one to understand the importance of molecular relaxations and microstructure changes on the trap states of the system.

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INTRODUCTION Among polyfluorene derivatives, poly(9,9-di-n-octylfluorenealt-benzothiadiazole) (F8BT) has attracted special attention in recent years due to its high performance and good stability as active layer in polymer based devices. F8BT is a fluorene copolymer characterized by an alternation of a biphenyl unit with a benzothiadiazole molecule with chemical structure shown in Figure 1a. It has been extensively used as active layer in highly efficient light-emitting diodes,1,2 as acceptor materials for photovoltaic cells,3−5 and more recently in light-emitting transistors.6,7 In all cases, F8BT has demonstrated relative high mobility for holes and electrons8 and an overall highly efficient performance.1,2 However, mechanisms which control the optical properties and conduction of charge carriers as well as its relation to the materials supramolecular properties are not yet fully understood.9−14 As a semicrystalline polymer, the bulk structure of F8BT exhibits a rather complex packing structure, which impacts on its electronic properties. It seems reasonable that the optoelectronic properties of conjugated polymers depend on the physical conformation of the polymer chains. Generally, it is expected that the polymer chain packs in such a way that π orbitals between adjacent molecules overlap efficiently. In this situation the π-orbitals are indeed delocalized over neighboring polymer chains and the charge carriers can be readily driven through the device by an external electric field.15 As a result, the charge carrier mobility increases and an improved output performance can be expected. © 2013 American Chemical Society

However, this is not the situation one faces experimentally. Kline and co-worker,16 have published a paper comparing the charge transport of poly(3-hexylthiophene) (P3HT) with different molecular weights (MW). The morphology analysis of those films showed that the low MW films were substantially more crystalline than those of the high MW material. Nonetheless, despite being more crystalline and ordered, the low-MW films presented a lower mobility than the high-MW films. The answer to this apparent contradiction was based on the possible missing orientation of neighboring crystals in lowMW films due to the flexibility of the amorphous spacer, showing that morphology is key to charge transport. On top of that, molecular relaxations, such as the glass transition, also play a distinguished role on the charge carrier transport behavior as already suggested before.9,17,18 The influence of polymer molecular packing and dynamics on the transport properties can also be viewed in terms of the appearance of structural defects and trapping states, affecting the charge transport properties of conjugated polymers as a whole. For instance, those phenomena can modulate the interchain π-spacing, which controls the hopping of charges between delocalized sites and also the exciton transfer between molecules; whereas the torsion angle between neighboring Received: March 28, 2013 Revised: July 23, 2013 Published: September 16, 2013 7865

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Figure 1. (a) Molecular structure of poly(9,9-dioctylfluoren-2,7-diyl-co-benzothiadiazole). (b) DSC results revealing a glass transition at T = (385 ± 10) K, an endothermic peak at T = (470 ± 10) K due to the melting of a pre-existent crystalline phase, an exothermic process at T = (450 ± 10) K, indicating a recrystallization, and, finally, a second endothermic transition T = (490 ± 15) K is due to the melting point (Tm) of the recrystallized phase. (c and d) F8BT thick film structural evolution probed by WAXD diffractograms as a function of temperature for near-parallel and normal incidence.

atures at which molecular relaxations and microstructure changes occur, when the F8BT is embedded within a diodelike device, were determined by performing dielectric relaxation measurements. Finally, the temperature dependence of the charge carrier mobility as well as the current−voltage characteristics of the diode-like devices were analyzed based on the presence of the assigned molecular relaxations and microstructure transitions and interpreted by means of changing of trap states properties.

units affects planarity and consequently the effective conjugation length of the backbone. Indeed, both the polymer structure and main-chain molecular dynamics are coupled with relaxation processes in polymers.19,20 Moreover, changes in the local structure due crystal-to-crystal phase transitions, commonly observed in semiconducting polymers10,21−25 and the onset of collective molecular motion in the amorphous phase due to Tg may have a major effect on the transport properties. In this article, a multitechnique study has been conducted to understand the molecular structure and chain dynamics of F8BT films and their influence on electronic properties, mainly on the trap states physics, when used as active layers in diodelike devices. To achieve this understanding X-ray diffraction (XRD) and differential scanning calorimetry (DSC) has been used to characterize the temperature evolution of the film’s microstructure. Dynamic mechanical thermal analysis (DMTA) and solid-state nuclear magnetic resonance (NMR) were used to probe and characterize the glass transition and secondary molecular relaxations of the material.26−28 Then, the temper-

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MATERIALS AND METHODS

Purified poly(9,9-dioctylfluoren-2,7-diyl-co-benzothiadiazole) (F8BT) was purchased from the Fraunhofer Institute in Golm. The chemical structure is displayed in Figure 1a. The molecular weight and the polydispersivity were measured by Gel Permeation Chromatography (GPC) using polystyrene as reference. The measured values were Mn = 206 kg/mol and Mw/Mn = 3. HOMO and LUMO levels were already measured elsewhere,29 and exhibit values of 5.8 and 3.3 eV, respectively. A 4 μm thick self-standing polymer film was obtained by drop-cast deposition of a 20 mg/mL toluene solution with subsequent drying 7866

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of the measurements were 70 to 490 K and 10−1 to 106 Hz, respectively. The TOF technique was performed using a 9-ns pulse duration and wavelength-tunable Nd:YAG LPS-1500 Solar Laser allowing excitations from 335 to 900 nm. A Keithley 6517A was used as a DC-voltage source. The photocurrents were amplified by the high-frequency amplifier FEMTO DHPCA-100 and recorded using a Tektronix TDS5052 oscilloscope. In order to increase the signal-to-noise ratio, at least 64 transients were averaged. The laser was tuned to λ = 470 nm corresponding to the F8BT fundamental absorption. Current−voltage measurements were carried out using a Keithley 6517A electrometer. The temperature range was the same as used by the other techniques (70−490 K).

inside a glovebox with a saturated solvent environment. After solvent evaporation, samples were stored in the dark and vacuum for several days before measurements were performed. For electrical measurements, drop-casted light-emitting diodes with thickness of 4 μm and spin coated devices with a thickness of 250 nm were assembled according to the following steps: Initially patterned ITO (20 Ω/cm2) on glass was cleaned by ultrasonication at 60 °C for 15 min in an alkaline soap solution (Denonex); then the substrates were rinsed in ultrapure water with subsequent drying followed by a second ultrasonication step for 15 min in acetone and subsequent drying in a N2 stream. Thereafter, cleaned substrates were treated in an ozone atmosphere for 15 min. The subsequent film deposition was then achieved for thick film device production by drop-casting the 20 mg/mL toluene polymer solution on cleaned substrates and for thin film production by spin-coating for 60 s at 2000 rpm from the same solution. The devices were completed by a 120 nm thermal evaporation of aluminum as cathode with an evaporation rate of 2 Å/s at 10−7 Torr. Differential scanning calorimetry (DSC) was conducted using the differential scanning calorimeter DSC-910 from TA Instrument in two subsequent steps: a heating step from room temperature to 570 K followed by a cooling step down to 230 K. All heating and cooling rates were fixed to 10 K/min. During the measurements the samples were kept in argon atmosphere with a flow rate of 100 mL/min. WAXD experiments were performed at the D11A-SAXS beamline of the LNLS (Brazilian Synchrotron Light Laboratory). The wavelength used was 1.488 Å and the sample detector distance was approximately 139 mm in all cases. The films were set up in two configurations: (1) with the incident X-ray beam perpendicular (⊥) and (2) near-parallel (∥) to the film plane. The samples were examined using a two-dimensional (2D) CCD detector. Average radial intensity profiles were obtained integrating a 15° angular sector in the case of the isotropic scattering pattern (⊥ incidence) and a similar sector centered around the maximum in the anisotropic scattering ring (∥ incidence). Intensities were normalized to the integrated intensity of the incident beam and by the sample absorption. For temperaturedependent measurements the samples were placed in a hot stage cell specially designed for X-ray scattering measurements (THM 600, Linkam Ltda). For both experimental geometries used (⊥ and ∥ incidences), polymer films were placed in a sample support adapted to the hot stage. The difference between the temperature values of the sample and the values set on the controller was less than 5 K for all scans in the range of 123 to 523 K, allowing a sufficiently precise determination of the thermal state of the sample. In situ measurements were performed allowing about 15 min for thermal stabilization followed by 15 s of data acquisition at each desired temperature. The heating rate was 10 K/min between measuring temperatures. Dynamical mechanical thermal analysis (DMTA) was performed to determine the viscoelastic properties of F8BT by using a tension mode Dynamic Mechanical Analysis setup (Netzsch DMA 242C) with frequencies ranging from 0.01 to 100 Hz and a heating rate of 2 K/min from 100 to 400 K. Solid-state NMR experiments were accomplished using the VARIAN INOVA spectrometer at 13C and 1H frequencies of 100.5 and 400.0 MHz, respectively. A VARIAN 7-mm MAS doubleresonance probe head with variable temperature (VT) was used. The spinning frequencies, varying between 4 and 10 kHz, were controlled by a VARIAN pneumatic system which ensures a rotation stability of ±2 Hz. Typical π/2 pulses lengths of 3.5 and 4.5 μs were applied for 13C and 1H, respectively. Time proportional phase modulated (TPPM) proton decoupling with field strength of 60 kHz, cross-polarization time of 1 ms and recycle delays varying between 3 and 5 s were used. The temperature dependence of C−H dipolar coupling was measured by the dipolar chemical-shift correlation technique, DIPSHIFT,30 in which H−H homonuclear decoupling was achieved by the phase-modulated Lee−Goldburg (PMLG) sequence,20 using field strengths of approximately 60 kHz. An impedance analyzer Solartron SI 1260 connected to a dielectric interface Solartron SI 1296 was used to conduct the dielectric relaxation measurements. The temperature and the frequency ranges

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RESULTS Structure and Dynamics of F8BT Films. Differential Scanning Calorimetry (DSC). DSC results are illustrated in Figure 1b. The baseline shift at T = (385 ± 10) K indicates the glass transition of the material.9,10 The peak at T = (460 ± 10) K has been registered only during the heating stage, being an endothermic process due to the melting of a pre-existent crystalline phase. A negative peak (exothermic) is observed at T = (475 ± 10) K indicating a recrystallization process. The second endothermic transition T = (490 ± 15) K is due to melting of the recrystallized phase (Tm). The fusion enthalpies for the melting peaks at 460 and 490 K were 9.1 and 8.6 J/g, respectively. Wide Angle X-ray Diffraction (WAXD). Parts c and d of Figure 1 show the results of the WAXD measurements performed as a function of temperature. At room temperature (298 K), three equally spaced diffraction peaks can be seen at q1 = 3.5 nm−1, q2 = 7 nm−1, and q3 =10.5 nm−1, for both incidence angles corresponding to characteristic lengths of (di= 2π/qi), d1 = 16, d2 =9.8, and d3 =5.6 Ǻ (see markers at Figure 1d at 398 K). Those peaks were earlier attributed to different reflection orders of a lamellar phase21−24 with interplanar distances of 5.6 Ǻ . Such lamellar structure will be referred here as β-phase. The β-phase remains unchanged up to nearly 398 K, suggesting stability up to this temperature. However, at T > 398 K the lines start to narrow and a collection of new sharp peaks start to show up still on top of the amorphous background, especially between 10 nm−1 and 20 nm−1. Note that 398 K is the region of the appearance of the glass transition as determined by the DSC method. This suggests that the onset of collective movement of the backbone in the amorphous region at Tg (T ≈ 385 K) may act as a trigger for the dissipation of the pre-existent β-phase and the appearance of a new semicrystalline phase. It is worth mentioning that the overall intensity of the background contribution (amorphous halo) is not substantially modified during this transition meaning that no substantial change in the sample’s crystalline portion seems to occur. This fact suggests that such process is a phase transformation from the solvent induced β-phase to a more stable phase (below called α-phase), i.e., a crystal-to-crystal structural transition. This is in good agreement with the DSC results, where the area under the first melting peak (460 K) and the recrystallization peak (475 K) are similar and even seem to overlap in the sense that the crystal-to-crystal transition is smeared out over a wider temperature range. Also, as presented, the results for the fusion enthalpies for both phases are quite similar, (9.1 J/g and 8.6 J/g). From parts c and d of Figure 1, one notices that the crystalto-crystal transition starts at around 400 K and continues up to 498 K with the appearance of a series of well-defined sharp 7867

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the Boltzmann constant. The calculated activation energy is EA = (0.27 ± 0.04) eV. This value is typical for local motions of specific chain segments, as reported previously.24 13 C Nuclear Magnetic Resonance. Figure 3a shows the 13C NMR CPMAS spectra for F8BT at room temperature with the peak assignments given according to the carbon numbering shown in the chemical structure of Figure 1a. The assignments were achieved by comparing the 13C NMR spectra with the Dipolar Dephasing13C NMR spectra, Figure 3b, which suppresses the signals from the carbons not bound to hydrogen, and also by comparison with the spectra of similar compounds reported in the literature.24,31,32 Information about segmental motions in the kHz to MHz frequency range was obtained by the NMR technique known as dipolar chemical shift correlation (DIPSHIFT). The DIPSHIFT technique provides a measurement of the 13C−1H magnetic dipolar coupling (CH coupling) by reducing the intensity of each line in the 13C CPMAS spectra by a factor that depends on the a typical evolution time t1, the transverse 13C relaxation time, the MAS frequency, and the dipolar coupling strength. Since the CH coupling is orientation dependent, the presence of molecular motions with rates between 103 and 107 Hz produces the averaging of the coupling, which is detected as specific change in the curve of the line intensity versus the t1 evolution time. Parts c and d of Figure 3 show the DIPSHIFT curves obtained at several temperatures for side-chain carbons. The steady increase in the minimum of the DIPSHIFT curve with rising temperature suggests the progressive averaging of the CH coupling due to molecular reorientations in the segments where the carbon nuclei corresponding to the observed signal is sited on. From 193 to 213 K, the minimum of the DIPSHIFT curve remains almost constant, due to the lack of molecular dynamics on the side-chain. However, as the β-relaxation temperature is reached (≈ 210 K), the minimum of the DIPSHIFT curve starts to increase, due to the onset of molecular movement responsible for the β-relaxation. At increasing temperatures, a reciprocal increase in the minimum is observed, suggesting that the motion rate is also constantly increasing. Nevertheless, for carbons 9−13, the last two temperatures (353 and 373 K) produces almost identical DIPSHIFT curves. This feature implies that the fast limit motion regime (motional rates ≫ dipolar coupling frequencies) has been reached. Such a regime occurs when the DIPSHIFT curves have no dependence on the motional rates, being completely defined by the residual CH dipolar coupling, which is defined only by the motion geometry. Another important characteristic of the molecular dynamics which can be extracted from the DIPSHIFT curves is the correlation time. This is achieved by using a similar treatment described in detail in reference30. In this approximation the local CH dipolar fields are considered to be Gaussian and the anisotropic motion geometry is modeled by diffusion on a cone. Under this assumption the motion correlation time becomes the single fitting parameter, so its temperature dependence can be extracted from the DIPSHIFT curves. Arrhenius plots of the obtained correlation times for the sidechain CH 2 motion are illustrated, together with the corresponding activation energies in Figure 3e. The activation energy for the motions experienced by carbon 14 at the extremity of the side chains is clearly smaller, which corroborates the presence of a molecular mobility gradient along the side chain. Note also, that the activation energies,

peaks indicating the end of the transition. The sharp peak positions are consistent with the formation of a helical like structure, usually referred as the α-phase. Such a phase was already reported elsewhere for polyfluorene homopolymer.21−23 One recognizes that the local order within the crystallites is much higher in the α-phase than in the β-phase as can be seen from the much sharper peaks of higher temperature WAXD diffractograms. The crystalline α-phase finally melts between 498 and 523 K, which can be seen by the loss of the diffraction peaks and seems to be in reasonable agreement with the temperature registered by DSC for the melt (≈500 K). Dynamic-Mechanical Thermal Analysis (DMTA). Figure 2a shows the temperature dependence of tan δ obtained from

Figure 2. (a) Temperature dependence of the tan δ obtained from DMTA measurements with excitation frequencies ranging from 1 to 50 Hz. (b) Arrhenius plot of the frequency at the maximum DMTA intensity for the peaks corresponding to the β-relaxation.

DMTA measurements with excitation frequencies ranging from 1 to 50 Hz. The lower temperature peak was observed in the temperature range from 160 to 250 K and is assigned as a βrelaxation. The higher temperature peak at T = (390 ± 15) K is assigned to the glass transition being consistent with DSC measurements. Figure 2b displays an Arrhenius plot of the frequency versus inverse temperature, latter taken as inverse of the temperature at the maximum DMTA tan δ intensity for the peaks corresponding to the β-relaxation. The activation energy was calculated by fitting the frequency versus 1/T dependence of the tan δ according to:

f = fo e−EA / kT

(1)

where f is the frequency (Hz), f 0 the pre-exponential factor (Hz), EA the activation energy (eV), T the absolute temperature at which the maximum of tan δ occurs, and k 7868

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Figure 3. (a) CPMAS NMR spectra at room temperature of F8BT. (b) 13C NMR spectra with the dipolar dephasing13C NMR spectra. (c) and (d) DIPSHIFT curves for several temperatures of carbons 9−13 and 14, respectively. (e) Arrhenius plots of the obtained correlation times for the sidechain CH2 motion.

(0.23 and 0.30 eV), are in the same order of those estimated from the DMTA measurements, confirming that these motions are indeed responsible for the β-relaxation process. Summarizing the results from this section, DSC revealed a glass transition at 380−390 K, which was confirmed by the DMTA measurement. DMTA also pointed to a secondary relaxation at approximately 200 K, whose molecular nature, revealed by DIPSHIFT NMR, was associated with the onset of side-chain motions. From WAXD, a crystal-to-crystal transition was observed above the glass transition, suggesting that the collective movement of the backbone in the amorphous region above Tg acts as a trigger for the dissipation of a solventinduced phase and the appearance of a new stable semicrystalline phase with distinct molecular structure. Electrical Properties of Thick-Film Devices. Dielectric Relaxation (DR). The main molecular relaxations of F8BT thick films were identified by thermal and structural analysis above. In this section, it will be shown that those transitions can also be identified by electrical techniques and, more important, they can influence the dielectric behavior.33 We start with dielectric relaxation, since molecular reorientation due to the above mentioned relaxations may change the dielectric properties of our material. Figure 4a shows the temperature dependence of the tan δ measured using a thick film ITO/ F8BT/Al diode for several frequencies (from 0.15 to 1 Hz). The tan δ curves, representing the dielectric losses due to dielectric relaxations, exhibit three peaks, with maxima observed between 130 and 200 K, 280 and 360 K, and 400 and 470 K, which are in a similar temperature interval assigned above to the β-relaxation, glass-transition (Tg) and the crystal-to-crystal structural transition (β-phase to α-phase). It is expected that a changing of the dielectric properties may also influence the materials conductivity. Accordingly, we proceed to TOF measurements to evaluate changes in the

Figure 4. (a) Temperature dependence of tan δ measured by dielectric spectroscopy and (b) charge carrier mobility as a function temperature. 7869

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charge carrier mobility due to those assigned molecular relaxations and transitions. Time of Flight Technique (TOF). The time-of-flight technique (TOF) is performed by generating a thin sheet of charge carriers close to one of the electrodes of the samples. The excitation is produced by a short laser pulse with wavelengths at the fundamental absorption band of the F8BT, leading to a charge plasma close to the illuminated electrode. According to the applied DC electric field charges are separated, where one carrier type is collected immediately by the illuminated electrode and the other is traversing as a charge packet through the sample leading to a measurable current. When the leading edge of the packet reaches the opposite electrode, a transit time transient (tT) is measured, from which it is possible to extract the average transit time and the effective charge-carrier mobility (μ). The relation between μ and tT is simply given by the following equation17

μ=

L tT E

(2)

where L is the sample thickness and E is the applied electric field. The mobility has been determined for several temperatures utilizing a positive voltage at the ITO electrode (E = 0.5 MV/ cm). Since the light exposure occurs through the ITO interface the resulting photocurrent is attributed to a hole transient (see traces as inset in Figure 4b). The hole mobilities calculated were in the range of 10−5 to 10−4 cm2/(V s), which is roughly 1 order of magnitude lower than previously reported.34 The mobility curve reveals two characteristic features: a clear change in the slope at 190 K and a deviation from a straight line at around 350 K. Those temperatures are the same as those measured for the β-relaxation and glass-transition found by DR measurements. Despite the small deviation at 350 K, a straight line can be fitted in order to estimate the activation energy of the mobility. The values obtained below and above 190 K are 10 and 100 meV, respectively, which are in the range of energies expected for this class of materials.9,17 Similar results were published using thin-film ITO/F8BT/Al diodes, measured by the Photo-CELIV technique.9 Current−Voltage Characteristics (I−V). In order to get more inside in the traps and their energetic distribution involved in transport, the current-voltage characteristics were measured at several temperatures and are presented in Figure 5(a) as log−log plot. At low temperatures and low voltages the current is almost constant becoming Ohmic at higher temperatures. The Ohmic regime extends up to a certain voltage, from which on a strong increase in the current is observed. The constant current, which dominates the low voltage region until about 230 K, can be understood in the following way: For thick devices, the conduction current at low fields and low temperatures is expected to be ohmic but in the present case too small to be detected by the measurement setup. Therefore, at lower temperatures only the capacitor C will be charged and the observed current is a pure displacement current described by I = C(dV/dt). For the utilized voltage ramp of 10 V per second, the currents can be estimated to be approximately 1 pA, which is in agreement with experimental data for T = 50 K. For higher temperatures, the conduction current becomes noticeable and evolves into an Ohmic current at T > 340 K, assuming a slope of one in a log−log plot. In this case, the displacement current can be neglected. A prerequisite of an Ohmic current is a neutral bulk of the sample, which in

Figure 5. Temperature dependent results obtained from the modeled current−voltage characteristics: (a) original current−voltage data in log−log scale; (b) the widths of the Gaussian distribution of trap states (σ) for thick films devices (full square) and thin films devices (open circles); (c) the density of traps (Hd) for thick film devices.

the present case is most likely established by impurity doping where the dopant delivers a charge carrier of one sign and itself assumed the opposite charge sign. These ohmic currents thereby dominate until a transition point (VT), which is temperature-dependent and moves toward high voltages as the temperature increases. Above VT, the curves follow a power law of J ∝ Vn with 7.8 ≤ n ≤ 9, which is commonly attributed to space charge limited conduction or currents (SCLC). Space charge effects thereby occur when excess charge is injected from the electrodes and accumulates in the bulk of the sample leading to strong deformation of the average applied electric field Eappl = V/L, where V is the applied voltage and L is the sample thickness. If injection barriers become very low compared to kT (Ohmic contact) the injection currents become very large and excess charge is driven into the sample. Thereby the electric field resulting from these charges starts to suppress the average applied electric field Eappl at the injecting electrode down to E(0) = 0, an often used boundary condition for SCLC transport. On the other hand, the electric field at the opposite electrode starts to exceed the average applied electric field in order to maintain the constant applied voltage. In case of the absence of electronic traps, Child’s law 7870

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9 V2 εε0μ 3 8 L

magnitude. The density of traps calculated ranges from 1017 to 1018 cm−3, which is also in the range of published results.38,39 Summarizing the findings from this section, a connection between the temperature dependence of the dielectric relaxation measurement, the onset of molecular relaxation processes and changes on molecular structure were established. In addition, a strong relation between the temperature dependence of the charge carrier mobility measured by TOF and the molecular relaxation and structural changes was observed which indicates a physical relation between the materials morphological and structural properties and the nature and origin of traps. This is supported by electrical J/V measurements at several temperatures whose modeling by trapmodulated SCLC indicates remarkable changes in depths and number of traps close to the glass and crystal-to-crystal transition. Discussion: Correlation between Molecular Dynamics, Structure and Electrical Properties. One very interesting aspect of the results above is that changes observed on both, the widths of the Gaussian distribution and the densities of traps occur in the same temperature ranges as the molecular transitions probed by thermal and structural methods, i.e., β-relaxation, glass transition, and crystal-to-crystal transition. In addition, similar molecular relaxation temperatures for polyfluorene materials have been published elsewhere.10,21−24,40 In those publications, however, no systematic correlations between the temperature dependence of the molecular relaxations and those of the electrical performance have been reported. Such dependence will now be highlighted in the following. The β relaxation is associated with the side-chain movements and the glass transition to segmental motions on the polymer backbone. WAXD measurements reveal that the glass transition precedes the dissipation of the solvent induced lamellar βphase, known to be metastable. In this sense, the onset of the molecular motions in the amorphous regions allows for the reorganization of the chains, so that a crystal-to-crystal transition (β-phase to α-phase) can occur. A stable α-phase is considered to be more ordered, as the polymer backbone organizes in a helical arrangement. At higher temperatures, the polymer reaches the melting point of the α-phase, as revealed by DSC and WAXD. The narrowing of the density of states σ in the vicinity of the glass transition (270−360 K) seen in Figure 5b, was already reported by Bässler et al.17 and Gyro et al.41 The X-ray data indicate the existence of the ordered β-phase, and an amorphous contribution. Since the glass transition occurs in the amorphous phase, the corresponding effect observed for charge carriers is most likely associated with the nonordered phase of the polymer. The motional effect due to the glass transition promotes a narrowing of the density of traps width,17 as can be seen from the thin film results in Figure 5a (open circles), where charges experience a smoother energy landscape due to the shacking backbones of the polymer. For thick film devices, (dark square), the narrowing is experienced by a plateauing of the width of the DOS. The reason therefore is most probably due to the superposition of the narrowing occurring in the amorphous phase (glass transition) and a widening due to the crystal-to-crystal transition occurring at the crystalline phase as can be seen for higher temperatures. In the same temperature range, the number of traps duplicates for thick films as can be seen in Figure 5c, probably due to an increase in torsion of the backbone induced by the glass

(3)

is obeyed with the usual meaning of the parameters. If, however, electronic traps are present in the bulk of the material the response changes from a simple quadratic dependence to a more complicated one which depends on the actual density of states (DOS) of the trap states.35,36 Limitations of this modeling procedure are related to the chosen distribution of the DOS. Exponential and Gaussian distribution are frequently used with success, nonetheless, the Gaussian distribution of traps is a more general model, including the exponential distribution in the case of charges populating the tail of the Gaussian DOS. For a Gaussian distribution of traps within the bandgap the SCLC current can be modeled by the following expression:35−37 m ⎛ 2m + 1 ⎞m + 1⎛ m εε0 ⎞ V m + 1 ⎟ J = q1 − mμp Nv⎜ ⎜ ⎟ 2m + 1 ⎝ m + 1 ⎠ ⎝ m + 1 Hd ⎠ L

(4)

where q is the electronic charge, μp the hole mobility, Nv the effective densities of states in the valence band, ε the dielectric constant of the material, ε0 the vacuum permittivity, and Hd the trap density for holes traps. The exponent m is related to the widths of the Gaussian distribution of the DOS (σ) by the following formula:35,36 1/2 ⎛ 2πσ 2 ⎞ m = ⎜1 + ⎟ ⎝ 16k 2T 2 ⎠

(5)

Here, k is the Boltzmann constant and T the temperature. Hence, σ can be directly calculated from the exponent of the SCLC region of the experimental data and is displayed as a function of inverse temperature in Figure 5a for thick and thin films. Nν has been estimated from the ohmic region (from 230 until 430 K) assuming that the current density J is given by35,36 J(T ) = qμp p(T )

V L

(6)

where p(T) = Nνe((EV−EF)/(kT)), EV is the energy of the valence band edge,36 and EF is the Fermi energy. Fitting of the temperature dependent J(T) resulted in Nv = 6.5 × 1019 cm−3 and EV − EF = 0.9 ± 0.3 eV, which agrees with the expectation that the Fermi level is located somewhere in the middle of the band gap (EF ≈ 4.6 eV) and the HOMO is located at approximately 5.8 eV.29 The value for Nv is consistent with values reported in literature for similar materials.38,39 By using the so determined Nν and the experimentally determined m one can uniquely extract Hd as a function of temperature from experimental data. The result is displayed in Figure 5b. The widths of the Gaussian, (Figure 5b), for thick filmsfull squareincreases almost linearly from 0.15 eV (120 K) to 0.25 eV (260 K) on the 1/T abscissa . In the vicinity of the glass transition (above 270 K), the width levels off until about 370 K and then increases again linearly up to 430 K on the 1/T abscissa. This second increase is absent for thin films. Concerning the density of traps (Hd) displayed in Figure 5c, it increases almost linearly from 120 to 270 K followed by a steeper slope until 370 K again on the 1/T scale. From this point upward, the density of traps Hd decreases roughly by three times. Within the analyzed range of temperatures the density of traps changed by approximately 1 order of 7871

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Interestingly, the widths of the DOS for thin film devices has also not increased as it did for thick film devices (Figure 5b). This indicates the absence of a well-organized crystalline phase in thin film devices and allows the conclusion that the increase in the width of the Gaussian distribution is related to the crystal-to-crystal transition.

transition and the increasing α-phase evoking shorter delocalized sections of the polymer and therewith a larger number of traps. Both effects, the narrowing of the width of the DOS and an increasing of the trap number are competitive processes resulting in the changes of mobility as can be seen from Figure 4b by the slight alterations of the slope. Another fact which deserves attention is that the interfaces between amorphous and crystalline regions in the sample are critical for the formation of charge traps. In this sense, it is expected that this interface becomes more prevalent due to the appearance of the well-organized α-phase above Tg. This leads to an increase of the widths of the DOS as suggested from Figure 5b. Despite the increasing widths of the DOS, the number of occupied traps decreases (Figure 5c). A likely explanation for this fact seems to be the increased thermal release from shallower traps at these already relatively high temperatures. To support the above statements, the same analysis of the current−voltage characteristics was performed for the spincasted device (thickness of 250 nm) prepared under the same conditions and solutions as the thick film device. Comparing the results of Figure 5b to the thick film device, a similar behavior is observed until 370 K. However, no sudden increase of σ is observed above this temperature for the thin film devices. This suggests that the broadening of the Gaussian distribution observed for thick film seems to be restricted to those films. Parts a and b for Figure 6 show the WAXD diffractograms for thin and thick films at room temperature and at 473 K,

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CONCLUSION In this article, an investigation of the temperature evolution of the microstructure and molecular relaxation of thick films and its correlation to the electrical properties of a representative polyfluorene-based polymer, known as poly(9,9-di-n-octylfluorene-alt-benzothiadiazole) (F8BT) has been presented. Several techniques have been employed, such as DSC, WAXD, DMTA and NMR to unravel the temperature dependent evolution of the microstructure and its correlation with molecular relaxations and dynamics as well as DS, TOF, IV to correlate them to the electrical transitions. Essentially, WAXD measurements revealed that the glass transition seems to precede the dissipation of the solvent induced lamellar βphase. In this sense, the onset of the molecular motion in the amorphous regions allows for the reorganization of the chains, so that a crystal-to-crystal transition can occur. A stable α-phase is considered ordered, when the polymer backbones are organized in a helical arrangement. At higher temperatures, the polymer reaches the melting point of the α-phase, as revealed by DSC and WAXD. The knowledge of the microstructure and molecular dynamics evolution enables one to correlate those results with the electrical properties of thick film devices. The comparison revealed that the increase of the widths of the Gaussian distribution of trap states and the decrease of the effective number of traps are connected with the crystallization of the α-phase. Furthermore, thin film devices were investigated in order to compare with thick film devices. The modeling of the current−voltage characteristics of thin film devices did not show any increase of the widths of the Gaussian distribution of trap states. A further analysis of thin film WAXD diffractograms for representative temperatures revealed no existence of the crystal-to-crystal transition as was observed in thick films. Such result strengthens the hypothesis that the increase of the width of the Gaussian distribution of trap states and the decrease of the effective number of traps are in fact connected with the crystallization of the α-phase.

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AUTHOR INFORMATION

Corresponding Author

*(G.C.F.) E-mail: [email protected]. Telephone/Fax: 55-163373-9825.

Figure 6. WAXD diffractograms for thin and thick films (a) at room temperature and (b) at 473 K. WAXD measurements for thin film were done using the same experimental procedures as that of thick films.

Author Contributions

The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript.

respectively, where the α phase is already formed in case of the thick films. The WAXD measurements for thin film were done using the same experimental procedures as those for thick films. Figure 6a shows that a β phase is also induced for spin-coated thin films of F8BT. Nonetheless the structure for the thick film seems to indicate a slightly higher order. At 473 K the crystalto-crystal transition has occurred for the thick film, however, no any substantial changes are observed for the thin film. The consequence is that the fast solvent evaporation during the spin coating accounts for a different structural kinetics, avoiding or retarding the formation of the well-organized α-phase.

Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS G.C.F. acknowledges a fellowship awarded by the “Fundaçaõ de Amparo á Pesquisa do Estado de São Paulo” (FAPESP) (Proc. number 2008/01935-5 and 2012/01303-4) and by the “Deutscher Akademischer Austauschdienst” (DAAD) (A/09/ 72945; ref 415). The authors are also grateful to the Brazilian funding agencies FAPESP (Proc. number 2009/18354-8 and 7872

dx.doi.org/10.1021/ma400648g | Macromolecules 2013, 46, 7865−7873

Macromolecules

Article

(26) Tsao, H. N.; Cho, D. M.; Park, I.; Hansen, M. R.; Mavrinskiy, A.; Yoon, D. Y.; Graf, R.; Pisula, W.; Spiess, H. W.; Mullen, K. J. Am. Chem. Soc. 2011, 133 (8), 2605−2612. (27) Dudenko, D.; Kiersnowski, A.; Shu, J.; Pisula, W.; Sebastiani, D.; Spiess, H. W.; Hansen, M. R. Angew. Chem., Int. Ed. 2012, 51 (44), 11068−11072. (28) Hansen, M. R.; Graf, R.; Spiess, H. W. Acc. Chem. Res. 2013, DOI: 10.1021/ar300338b. (29) Zaumseil, J.; Donley, C. L.; Kim, J. S.; Friend, R. H.; Sirringhaus, H. Adv. Mater. 2006, 18 (20), 2708−2712. (30) Deazevedo, E. R.; Saalwachter, K.; Pascui, O.; De Souza, A. A.; Bonagamba, T. J.; Reichert, D. J. Chem. Phys. 2008, 128−140, 10. (31) Reichert, D.; Pascui, O.; Bonagamba, T. J.; Belton, P.; Schmidt, A.; Deazevedo, E. R. J. Magn. Reson. 2008, 191 (1), 141−147. (32) Souza, A. A.; Cossiello, R. F.; Plivelic, T. S.; Mantovani, G. L.; Faria, G. C.; Atvars, T. D. Z.; Torriani, I. L.; Bonagamba, T. J.; Deazevedo, E. R. Eur. Polym. J. 2008, 44 (12), 4063−4073. (33) Graff, M. S.; Boyd, R. H. Polymer 1994, 35 (9), 1797−1805. (34) Kim, Y.; Cook, S.; Choulis, S. A.; Nelson, J.; Durrant, J. R.; Bradley, D. D. C. Chem. Mater. 2004, 16 (23), 4812−4818. (35) Kao, K. C.; Hwang, W. Electrical transport in solids: with particular reference to organic semiconductors; Pergamon: Oxford, U.K., 1981. (36) Lampert, M. A.; Mark, P. Current injection in solids; Academic Press: New York and London, 1970. (37) Bechtold, I. H.; Eccher, J.; Faria, G. C.; Gallardo, H.; Molin, F.; Gobo, N. R. S.; de Oliveira, K. T.; von Seggern, H. J. Phys. Chem. B. 2012, 116 (45), 13554−13560. (38) Blom, P. W. M.; de Jong, M. J. M. IEEE J. Sel. Top. Quantum Electron. 1998, 4 (1), 105−112. (39) Blom, P. W. M.; Vissenberg, M. C. J. M. Mater. Sci. Eng. R-Rep. 2000, 27 (3−4), 53−94. (40) Arif, M.; Volz, C.; Guha, S. Phys. Rev. Lett. 2006, 96 (2), 025503. (41) Giro, G.; Dimarco, P. G. Chem. Phys. Lett. 1989, 162 (3), 221− 226.

2007/08688-0), CNPq and MCT/INEO for funding, and to the Brazilian Synchrotron Light Source, especially to Débora Magalhães and Mateus Cardoso, for helpful assistance during WAXD measurements (Projects: SAXS-1/8138 and SAXS-1/ 13602). H.v.S. acknowledges the scholarship “Special Visiting Researcher” by the Brazilian Science without Borders Program (CNPq and Capes, Proc. number 400133/2012-1). We are also grateful to Débora Balogh for helping with DMTA measurements.

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REFERENCES

(1) Kabra, D.; Lu, L. P.; Song, M. H.; Snaith, H. J.; Friend, R. H. Adv. Mater. 2010, 22 (29), 3194−3198. (2) Nakayama, Y.; Morii, K.; Suzuki, Y.; Machida, H.; Kera, S.; Ueno, N.; Kitagawa, H.; Noguchi, Y.; Ishii, H. Adv. Funct. Mater. 2009, 19 (23), 3746−3752. (3) Peters, C. H.; Sachs-Quintana, I. T.; Kastrop, J. P.; Beaupre, S.; Leclerc, M.; McGehee, M. D. Adv. Energy Mater. 2011, 1 (4), 491− 494. (4) Snaith, H. J.; Arias, A. C.; Morteani, A. C.; Silva, C.; Friend, R. H. Nano Lett. 2002, 2 (12), 1353−1357. (5) Yan, H. P.; Swaraj, S.; Wang, C.; Hwang, I.; Greenham, N. C.; Groves, C.; Ade, H.; McNeill, C. R. Adv. Funct. Mater. 2010, 20 (24), 4329−4337. (6) Schidleja, M.; Melzer, C.; von Seggern, H. Frequenz 2008, 62 (3− 4), 100−103. (7) Zaumseil, J.; Friend, R. H.; Sirringhaus, H. Nat. Mater. 2006, 5 (1), 69−74. (8) Chua, L. L.; Zaumseil, J.; Chang, J. F.; Ou, E. C. W.; Ho, P. K. H.; Sirringhaus, H.; Friend, R. H. Nature 2005, 434 (7030), 194−199. (9) Faria, G. C.; Faria, R. M.; deAzevedo, E. R.; von Seggern, H. J. Phys. Chem. C 2011, 115 (51), 25479−25483. (10) 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 (37), 12890−12899. (11) Kohler, A.; Hoffmann, S. T.; Bassler, H. J. Am. Chem. Soc. 2012, 134 (28), 11594−11601. (12) Qin, R. P.; Li, W. W.; Li, C. H.; Du, C.; Veit, C.; Schleiermacher, H. F.; Andersson, M.; Bo, Z. S.; Liu, Z. P.; Inganas, O.; Wuerfel, U.; Zhang, F. L. J. Am. Chem. Soc. 2009, 131 (41), 14612−14613. (13) Wang, S.; Kiersnowski, A.; Pisula, W.; Mullen, K. J. Am. Chem. Soc. 2012, 134 (9), 4015−4018. (14) Woo, C. H.; Thompson, B. C.; Kim, B. J.; Toney, M. F.; Frechet, J. M. J. J. Am. Chem. Soc. 2008, 130 (48), 16324−16329. (15) Seanor, D. A. Electrical properties of polymers; Academic Press: New York ; London, 1982. (16) Kline, R. J.; McGehee, M. D. Polym. Rev. 2006, 46 (1), 27−45. (17) Bassler, H.; Borsenberger, P. M. Chem. Phys. 1993, 177 (3), 763−771. (18) Laquai, F.; Wegner, G.; Bassler, H. Philos. Trans. R. Soc. London, A: Math. Phys. Eng. Sci. 2007, 365 (1855), 1473−1487. (19) Boyd, R. H. Polymer 1985, 26 (8), 1123−1133. (20) Schmidt-Rohr, K.; Spiess, H. W. Multidimensional solid-state NMR and polymers; Academic Press: London, 1994. (21) Chen, S. H.; Su, A. C.; Chen, S. A. J. Phys. Chem. B. 2005, 109 (20), 10067−10072. (22) Chen, S. H.; Su, A. C.; Su, C. H.; Chen, S. A. J. Phys. Chem. B. 2006, 110 (9), 4007−4013. (23) Chen, S. H.; Su, A. C.; Su, C. H.; Chen, S. A. Macromolecules 2005, 38 (2), 379−385. (24) Faria, G. C.; Plivelic, T. S.; Cossiello, R. F.; Souza, A. A.; Atvars, T. D. Z.; Torriani, I. L.; deAzevedo, E. R. J. Phys. Chem. B. 2009, 113 (33), 11403−11413. (25) Pascui, O. F.; Lohwasser, R.; Sommer, M.; Thelakkat, M.; Thurn-Albrecht, T.; Saalwachter, K. Macromolecules 2010, 43 (22), 9401−9410. 7873

dx.doi.org/10.1021/ma400648g | Macromolecules 2013, 46, 7865−7873