J. Phys. Chem. B 2009, 113, 11403–11413
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A Multitechnique Study of Structure and Dynamics of Polyfluorene Cast Films and the Influence on Their Photoluminescence G. C. Faria,† T. S. Plivelic,‡,⊥ R. F. Cossiello,§ A. A. Souza,† T. D. Z. Atvars,§ I. L. Torriani,‡,| and E. R. deAzevedo*,† Instituto de Fı´sica de Sa˜o Carlos, UniVersidade de Sa˜o Paulo, Caixa Postal 369, 13560-970 Sa˜o Carlos, SP, Brazil, Laborato´rio Nacional de Luz Sı´ncrotron, Caixa Postal 6192, 13083-970 Campinas, SP, Brazil, Instituto de Quı´mica, UniVersidade Estadual de Campinas, Caixa Postal 6154, 13084-971 Campinas, SP, Brazil, Instituto de Fı´sica, UniVersidade Estadual de Campinas, Caixa Postal 6165, 13084-971 Campinas, SP, Brazil, and MAXLab, Lund UniVersity, P.O Box 118, SE-22100 Lund, Sweden ReceiVed: May 9, 2009; ReVised Manuscript ReceiVed: June 15, 2009
This article describes the microstructure and dynamics in the solid state of polyfluorene-based polymers, poly(9,9-dioctylfluorenyl-2,7-diyl) (PFO), a semicrystalline polymer, and poly[(9,9-dioctyl-2,7-divinylenefluorenylene)-alt-co-{2-methoxy-5-(2-ethyl-hexyloxy)-1,4-phenylene vinylene}, a copolymer with mesomorphic phase properties. These structures were determined by wide-angle X-ray scattering (WAXS) measurements. Assuming a packing model for the copolymer structure, where the planes of the phenyl rings are stacked and separated by an average distance of ∼4.5 Å and laterally spaced by about ∼16 Å, we followed the evolution of these distances as a function of temperature using WAXS and associated the changes observed to the polymer relaxation processes identified by dynamical mechanical thermal analysis. Specific molecular motions were studied by solid-state nuclear magnetic resonance. The onset of the side-chain motion at about 213 K (β-relaxation) produced a small increase in the lateral spacing and in the stacking distance of the phenyl rings in the aggregated structures. Besides, at about 383 K (R-relaxation) there occurs a significant increase in the amplitude of the torsion motion in the backbone, producing a greater increase in the stacking distance of the phenyl rings. Similar results were observed in the semicrystalline phase of PFO, but in this case the presence of the crystalline structure affects considerably the overall dynamics, which tends to be more hindered. Put together, our data explain many features of the temperature dependence of the photoluminescence of these two polymers. 1. Introduction Electroluminescent (EL) polymers have been the subject of many studies as a result of their potential applications in passive/ active optical devices. Molecular dynamics and chain packing play an important role on the polymers’ optoelectronic properties.1-3 For example, the temperature dependence of the transport properties of electroluminescent devices with MEHPPV showed changes associated with specific movements of the polymer chain.4-6 Besides that, molecular packing is another feature of great importance for the electroluminescence performance, and in many cases, the formation and the type of crystalline phases or molecular aggregates dictates essential functional properties. It has been proven that different packing usually induces or quenches specific luminescent processes,7 a feature that makes thermal treatments usually mandatory when constructing reproducible polymer-based devices. Among the conjugated polymers, polyfluorene-based polymers form a large category of thermal and chemical stable polymers emitting from the blue to the red,8-12 depending on the chemical structure, substituent, and side chain.13 These chemical modifications control the supramolecular structure and * To whom correspondence should be addressed. E-mail: azevedo@ ifsc.usp.br. Fax: 55-1633739876. † Universidade de Sa˜o Paulo. ‡ Laborato´rio Nacional de Luz Sı´ncrotron. § Instituto de Quı´mica, Universidade Estadual de Campinas. | Instituto de Fı´sica, Universidade Estadual de Campinas. ⊥ Lund University.
their optical properties.14 Thus, in addition to the relaxation processes involving segments of the polymer chains, the emission of polyfluorenes is also related to the solid-state phase transitions because some of them crystallize in more than one phase and others form molecular aggregates with mesomorphic phases.15 These structural characteristics probably have significant effects on the polymer dynamics. There are many studies about the structural features of polyfluorene derivatives. In particular, a variety of crystalline and nematic phases have been reported for poly(9,9-dioctylfluorenyl-2,7-diyl) (PFO)16 as well as other derivatives.15-19 For toluene as cast PFO films, a lamellar mesormorphic metastable β phase was observed, sometimes together with a solventinduced clathrate crystalline phase.19 These phases could be converted to two polymorphic crystalline phases (referred to as R and R′ phases) upon proper thermal treatments. The structures of the R and R′ phases were found to be similar,16 but the R′ has lower melting and crystallization temperatures. Above their respective melting temperatures, these crystalline phases transform into a nematic liquid crystalline phase, which keeps this type of order up to 573 K.16-19 The most stable R phase was described by an orthorhombic unit cell with parameters a ) 2.56 nm, b ) 2.34 nm, and c ) 3.32 nm, belonging to the space group P212121, with eight chains per unit cell (four fluorine repeat units per cell). In this structure, the fluorine groups were found to be twisted by 15-24° and the close side group packing is such that the backbone groups of different chains are kept apart, supporting the observation that
10.1021/jp9043368 CCC: $40.75 2009 American Chemical Society Published on Web 07/23/2009
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Faria et al. are used mainly as a reference for comparison with the structural behavior of the PFO-altco-MEHPV copolymer. 2. Experimental Methods and Materials
Figure 1. (a) Chemical structures of PFO and PFO-altco-MEHPV; (b) room temperature 2D WAXS images; (c) WAXS 1D pattern obtained from the corresponding 2D images. For all, (left) PFO and (right) PFO-altco-MEHPV.
the emission of crystalline PFO is dominated by intrachain excitons.15 On the other hand, the presence of interchain species was observed in the β phase due to the close contact between the conjugated rings.20 In the present study we compare the packing structure, the phase transitions, and the dynamics of the phase transitions of PFO and poly{(9,9-dioctyl-2,7-divinylene-fluorenylene)-alt-co[2-methoxy-5-(2-ethyl-hexyloxy)-1,4-phenylene vinylene]} (PFOaltco-MEHPV) (chemical structure sketched in Figure 1a). As mentioned, PFOs exhibit different crystalline structures, mesomorphic or nematic phases.9 The presence of the MEHPV unit in the PFO-altco-MEHPV may prevent crystallization but may also favor the formation of a mesomorphic phase due to the rigidity of the backbone. For this study we initially analyzed the structural features of these two polymers using wide-angle X-ray scattering (WAXS) measurements. The temperature dependence of the structural parameters is established, and the main variations are discussed in detail. Then, the temperatures of the thermal relaxation processes were properly determined by dynamical mechanical thermal analysis (DMTA). Solid-sate nuclear magnetic resonance (NMR) measurements were used to describe the specific movements of the polymer chains. The main differences observed between both polymers are emphasized using results from NMR techniques. Finally, put together, these results are used to explain the temperature dependence of the polymer photoluminescence. It should be explicitly mentioned that the microstructure and phase behavior of PFO is already well-known,16,17,19 and our WAXS data show only small differences in comparison with the previously reported results. Thus, in this article, our PFO WAXS results
Both samples were purchased from American Dye Source (ADS) Company under the trademarks ADS129BE (PFO) and ADS108GE (PFO-altco-MEHPV). The molecular weights and the polydispersity were measured by gel permeation chromatography using polystyrene standards. The measured values were j w/M j n ) 2.1 and M j n) 71 kg/mol, M j w/M jn ) j n) 58 kg/mol, M M 2.4 for PFO and PFO-altco-MEHPV, respectively. Self-standing 15 µm thick polymer films were obtained by drop-casting from a 20 mg/mL chloroform solution and dried in a saturated solvent environment. After preparation, the samples were stored under dark and vacuum for several days before WAXS and NMR measurements were performed. For the X-ray experiments, multilayer samples of up to 500 µm thickness were prepared by stacking pieces cut from a single film. Steady-state fluorescence spectra were recorded using a PC1 Photon Counting Spectrofluorimeter from ISS Inc. The spectral range was from 400 to 700 nm for the emission spectra. Slits were selected for a spectral resolution of ∼0.5 nm. Excitation was performed using the 337 nm emission of a He:Cd laser. DMTA measurements were performed in a Netzsch 242 C instrument using the tension mode at frequencies ranging from 0.01 to 100 Hz and heating rate of 1 K/min from 150 to 400 K. WAXS experiments were performed at the D02A-SAXS beamline of the LNLS (Brazilian Synchrotron Light Laboratory). The wavelength used was 1.488 Å, and the sample detector distance was approximately 189.62 mm in all cases. The films were set up in two configurations: with the incident X-ray beam perpendicular (⊥) and near-parallel (//) to the film plane. The sample scattering was registered using a two-dimensional (2D) MAR-CCD detector with 5 min of data acquisition in all cases. Average radial intensity profiles were obtained by integrating an arbitrary 5° angular sector in the case of the isotropic scattering pattern (⊥ incidence). For // incidence, 5° angular sectors were taken around the maximum in the oriented scattering ring (meridional sector) and also perpendicular to this direction (equatorial sector). Intensities were normalized by the integrated intensity incident on the sample during the exposure and by sample absorption. Parasitic scattering was subtracted from each pattern. Films were placed in a specific sample support adapted to a hot stage cell designed for X-ray scattering measurements described in details in ref 21. The measurements were performed allowing 10 min of stabilization and 5 min of data acquisition for each desired temperature. After a first measurement at room temperature, the samples were rapidly quenched (60 Κ/min) down to 173 K. Next, after each step of a heating ramp (10 K/min), WAXS patterns were obtained for several temperatures. NMR experiments were performed using a VARIAN INOVA spectrometer at 13C and 1H frequencies of 100.5 and 400.0 MHz, respectively. A VARIAN 7-mm MAS double-resonance probe head with variable temperature (VT) was used. The spinning speeds, varying between 4 and 6 kHz, were controlled by a VARIAN pneumatic system that ensures a rotation stability of (2 Hz. Typical π/2 pulse 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 70 kHz, crosspolarization time of 1 ms, and recycle delays varying between 3 and 5 s were used. Amplitude of slow molecular motions were investigated using the centerband-only detection of exchange (CODEX) technique1,22 using a constant time scheme (CON-
Multitechnique Study of Polyfluorene Cast Films TRA)23 with mixing time tm of 300 ms and maximum evolution times of 800 µs. Here the evolution time is the total duration of periods where the 13C spins evolves under the chemical shift anisotropy (CSA). During this period the CSA is reintroduced by a set of synchronized π pulses applied in a REDOR fashion during four MAS rotation periods (tr) (two before and two after the mixing time making a total evolution time of 4tr ) 800 µs, see ref 23 for details). The temperature dependence of C-H dipolar coupling was measured using the dipolar chemical shift correlation (DIPSHIFT) technique,24 where H-H homonuclear decoupling was achieved by the phase modulated Lee Goldburg (PMLG) homonuclear decoupling sequence, using field strengths of approximately 60 kHz. 3. Results 3.1. Structural Characterization. The wide-angle X-ray scattering results reveal the presence of an anisotropic morphology for both studied materials (Figure 1b) as usually observed for films of conjugated polymers.25,26 The 2D-WAXS images are composed of scattering rings with roughly uniform intensity when the X-ray beam is perpendicular to the film surface, but we observe short arcs and rings of nonuniform intensity when the X-ray incidence is nearly parallel. This shows that the scattering from certain periodic structures in the sample have a preferential orientation, with a tendency to be arranged parallel to the film surface and randomly distributed in the plane. Despite anisotropy being a common feature in both samples, the corresponding WAXS patterns are distinct. The 2D-WAXS image, or the corresponding 1D pattern of PFO, Figure 1c (left), shows sharp peaks, revealing the presence of organized structures in this polymer. In contrast, the 2D-WAXS images of PFOaltco-MEHPV, Figure 1c (right), are composed of broad scattering peaks, indicating the presence of only local ordering in the PFO-altco-MEHPV. The three defined broad maxima in both diffractograms, taken from both meridian and equatorial sectors as described in the Methods section, correspond to characteristic lengths (di ) 2π/qi) of d1 ) 15.9 Å, d2 ) 7.9 Å, d3 ) 4.5 Å suggesting an aggregation structure similar to that observed for alkyl side-chain substituted PPV,26 where the backbone rings are stacked parallel to each other with interplanar distances of 4.5 Å and laterally spaced by the side chains with average distances of 15.9 Å. In this supramolecular structure, the other important characteristic distance d2 ) 7.9 Å corresponds to the separation between monomeric units (which is about the size of a fluorene unit). A sketch of the packing structure is shown in Figure 2b (right). Note that the tacticity of the polymer is probably random, which is more consistent with the observation of a single d1 distance value. Figure 1c shows two PFO diffractograms corresponding to meridional and equatorial cuts taken from the 2D images in the near parallel incidence. The equatorial and meridional intensity profiles have some sharp reflections, corresponding to characteristic distances of 12.6, 9.1, 6.6, 6.1, 5.2, 4.8, 4.4, and 4.2 Å. Interestingly, the peaks associated to the distances 12.6, 6.6, and 4.2 Å are fairly equally spaced in q-space, suggesting that they arise from different order reflections of a lamellar phase. The existence of a lamellar β phase in PFO as cast films has been already suggested in reference,19 but higher order reflections were not clearly observed. Thus, it seems that in our case the β phase was better organized, which can be a consequence of the different solvent used or an aging effect (our measurements were performed a few days after the film preparation). Several other peaks were also well-defined in the profile corresponding to the near parallel indidence meridional profile, confirming that in our samples the β phase had a higher
J. Phys. Chem. B, Vol. 113, No. 33, 2009 11405 degree of organization than in the samples studied in the reference cited above. WAXS measurements were also performed as a function of temperature, giving some insight about the changes in the microstructure of both polymers. For PFO, the temperature dependence of the equatorial WAXS profile is shown in Figure 2a (left). The profiles are identical bellow 373 K, showing that the β phase is stable in this temperature range. Above 373 K the characteristic diffraction peaks corresponding to the β phase have their intensity reduced, indicating the dissipation of this phase. At 413 K a series of sharp peaks start showing up, accounting for the onset of the sample crystallization, which results in a diffractogram with well-defined peaks at 433 K, where the crystalline phase seems to be fully formed. Note that the peaks at q ) 3.5 nm-1 (d ) 17.9 Å), q ) 4.5 nm-1 (d ) 13.9 Å), q ) 5.2 nm-1 (d ) 12.0 Å), and q ) 7.3 nm-1 (d ) 8.6 Å) correspond to the (110), (200), (210), and (310) reflections of the PFO R phase.16 At 453 K the crystalline R phase is melted, giving rise to a pattern typical of a nematic liquid crystalline phase. This pattern is similar to that observed at room temperature for PFO-altco-MEHPV, showing that the mesomorphic phase in PFO-altco-MEHPV has a packing structure similar to that of the nematic phase of PFO. Note that even before the melting of the R phase, there is a background contribution (“amorphous halo”) in the PFO diffractogram (see also the corresponding image in Figure 1b) that resembles the pattern of the liquid crystalline phase also suggesting a similar structural packing. Figure 2a (right) show the PFO-altco-MEHPV 1D diffraction profiles corresponding to equatorial and meridional sectors of the 2D Images obtained for near parallel incidence as a function of temperature. These X-ray profiles basically show the changes in the three characteristic distances d1, d2, and d3, plotted as a function of temperature in Figure 2b (left). No temperature dependence is observed for the distance d2, but clear trends are observed for d1 and d3. Note that whereas d1 generally increases as the temperature is raised, there is a jump in the slope of the d1 versus T curve at about 383 K. The slope of the d3 versus T curve also changes at about 383 K, bending downward in contrast with the behavior of d1. These results show that, although the aggregated structures are present at all temperatures explored in these measurements, the intermolecular distances inside these structures depend on temperature, suggesting the existence of thermal relaxations. 3.2. Polymer Relaxation and Molecular Dynamics Characterization. Here we give a more detailed and consistent picture of the molecular dynamics in PFO and PFO-altcoMEHPV responsible by the solid-state polymer relaxation processes, including the determination of specific features such as correlation times, activation energies, and amplitude of molecular motions. The identification of the molecular relaxation processes in the samples was initially achieved using DMTA experiments. Figure 3a shows the temperature dependence of the tan δ ) E′′/E′ for excitation frequencies ranging from 1 to 50 Hz. For PFO-altco-MEHPV, the lower temperature peak (β-relaxation), composed by a single broad peak for all excitation frequencies, was observed from 140 to 240 K. For PFO two overlapped peaks are observed between 160 and 280 K, being distinguished only for frequencies lower than 5 Hz. This suggests the existence of at least two molecular processes that contribute to the β-relaxation in PFO, while for PFO-altco-MEHPV only a single contribution can be distinguished. In addition to this peak, we also observed for PFO a peak at about 130 K, which we will
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Figure 2. (a) WAXS 1D diffractograms of PFO (right) and PFO-altco-MEHPV (left) as a function of temperature. For the sake of simplicity, only those diffractograms corresponding to cuts and incidences that bring distinct information are shown. (b) Temperature dependence of the d1, d2, and d3 parameters extracted from the WAXS 1D diffractograms of PFO-altco-MEHPV for near parallel incidence (left). Molecular model for the supramolecular aggregation used in the data interpretation (right).
denominate γ-relaxation process. Two other overlapped peaks (at 330 and 380 K) are observed at higher temperatures. For PFO their relative intensity and position depend on the frequency; for PFO-altco-MEHPV only the peak position depends on the frequency. These two peaks are nominated RA and RB in order of increasing temperature. Figure 3b shows the Arrhenius plot of the frequency at the maximum DMTA intensity (fmax) for the peaks corresponding to the β-relaxation. The activation energy was calculated by fitting the frequency f versus 1/T dependence of the tan δ according to the Arrhenius law f ) f0 exp(-E/RT) where f is the frequency (Hz), f0 is the pre-exponential factor (Hz), E is the activation energy (J mol-1), T is the absolute temperature where the maximum of tan δ occurs for each relaxation process, and R is 8.32 J K-1 mol-1. The calculated activation energies were Ea ) (16 ( 4) kJ/mol for PFO and Ea ) (23 ( 4) kJ/mol
for PFO-altco-MEHPV, respectively. These values are typical of local motions of specific segments.27,28 The higher activation energy determined for PFO-altco-MEHPV may indicates that the molecular motions are more hindered or they involve larger macromolecular segments. Note, however, that because of the superposition of the two peaks in the PFO DMTA curves, the activation energy reflects an average value, which is probably attributed to two different processes or to processes involving two different phases. As a result of the spectral superposition of the RA peak, we were unable to obtain the Arrhenius plot and then not much can be concluded about the origin of this relaxation process from the analysis of the DMTA results. Fortunately, this is not the case of the RB processes, whose the Arrhenius plots (Figure 3c) are clearly nonlinear, exhibiting a typical Vogel-Fulcher behavior.29,30 This profile is usually observed in the relaxation phenomena of glass-forming and
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TABLE 1: WLF Parameters Extracted from the DMTA Arrhenius Plot WLF parameters
PFO
PFO-altco-MEHPV
C1 C2 Tg f(Tg)
23 68.3 K 368 K 0.01 Hz
17 42.3 K 362 K 0.01 Hz
viscous systems which can be fitted by way of the known Williams-Landel-Ferry (WLF)31 equation, i.e,
( )
log
C1(T - Tg) f(T) ) f(Tg) C2 + T - Tg
(1)
Where the C1 and C2 parameters are characteristic of the material; Tg is a reference temperature (glass transition temperature in the case of glass forming system). The obtained parameters are shown in Table 1 for both samples. A common derivation of the WLF equation considers that the free volume decreases linearly with decreasing temperature.31 These are in agreement with the WAXS results for both polymers since we noted that only the interplanar ring distance d3 changes significantly at temperatures above 383 K, pointing to a correlation between the increase in the free volume and the temperature. However, it should also be pointed out that there might be also nonordered (amorphous) regions in the sample, which also contributes to the glass transition process. Therefore, the behavior observed in Figure 3c strongly supports the
Figure 3. (a) DMTA curves for selected excitation frequencies. (b) Arrhenius plot of the β-relaxation peak maximum. (c) Arrhenius plot of the RA-relaxation peak maximum. PFO (left) and PFO-altco-MEHPV (right).
attribution of the RB peaks as a result of a glass transition in both polymers. Despite DMTA indicating the presence of macromolecular relaxations in the polymer, this experiment does not provide information about the molecular nature of the motions. For that, a technique capable of providing site specificity and sensibility to molecular rotations at different frequency scales is need. Solid-state NMR techniques are extensively used to probe molecular reorientations occurring in a wide frequency range, from 0.1 Hz to 100 MHz. Thus, a required frequency may be selected for the study of every macromolecular motion. This frequency range can be selected using the DMTA information. The DMTA data showed that while the maximum of the β relaxation peak for excitation at 10 Hz occurs at 223 K for PFO and 203 K for PFO-altco-MEHPV, for the RA and RB processes it occurs at approximately 383 K for both polymers. Thus, by extrapolating the Arrhenius plot of Figures 3b and c, one can predict that in the temperature range of 233 K to about 373 K (temperature range accessible to our MAS NMR probehead), the frequency scale of the molecular motion associated to the β-relaxation is in the kHz to MHz range (intermediate motional regime), whereas for the RA and RB processes it is in the Hz to kHz range (slow motion regime). Hence, three NMR methods were applied to probe local molecular motions in these frequency scales: a simple analysis of the cross-polarization magic angle spinning (CPMAS) NMR37 spectra as a function of temperature (to identify the onset of molecular motions with rates higher than kHz); the DIPSHIFT method32 (to obtain specific information about these motions, for example, molecular order parameter,38 correlation times, and activation energies34); and the CODEX method (to study motions in the Hz to kHz frequency scale).23 Figure 4 shows the CPMAS spectra for PFO (left) and PFOaltco-MEHPV (right) at several temperatures together with the corresponding line assignments according to the carbon numbering of Figure 1. At temperatures higher than 293 K, the narrowing of the lines associated to side-chain carbons is evident in both samples, indicating that these segments execute molecular rotations with rates higher than a few kHz. The absence of line narrowing for lines attributed to the backbone carbons shows that, even at 373 K, there is no intermediate regime dynamics in the polymer backbone. More specific information about the molecular motions observed in the CPMAS versus temperature experiments were obtained using the DIPSHIFT method,24 which provides a measurement of the CH magnetic dipolar coupling of each chemical group. This is done by measuring the dependence of the signal amplitude (each line in the CPMAS spectrum) with an evolution period (t1),32 where the nuclear spins evolve under the action of the CH magnetic dipolar coupling, producing a curve S(t1)/S(0) that depends on the dipolar coupling strength. Motions with rates between 103-107 Hz average the CH coupling, changing the shape of the DIPSHIFT curves and making possible to distinguish rigid from mobile groups.32 Figure 5a shows the DIPSHIFT curves obtained as a function of temperature for several lines corresponding to some selected side-chain carbons in PFO (left) and PFO-altco-MEHPV (right). The curves corresponding to backbone carbons or those attached to it (not shown) do not present significant changes as a function of temperature, confirming the absence of dynamics in the kHz frequency scale from 213 to 353 K as already observed by the CPMAS results. In contrast, a progressive increase in the minimum of the DIPSHIFT curves associated with side-chain carbons can be observed, revealing the presence of molecular reorientations in the side chain. The modification on the
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Figure 5. (a) DIPSHIFT curves of Cγ carbons of PFO (carbons 20 and 28) and PFO-altco-MEHPV (carbons 20, 28, 45, and 47). (b) DIPSHIFT curves of Cβ carbons of PFO (carbons 15-19 and 23-27) and PFO-altco-MEHPV (carbons 15-19, 23-27, 43, and 44). The lines correspond to AW fits of the experimental data. (c) Arrhenius plot of the correlation times extracted from fitting the DIPSHIFT curves of carbons Cβ and Cγ using a formula based on the Anderson and Weiss approximation. For all, (left) PFO and (right) PFO-altco-MEHPV. The measured temperatures are those corresponding to the points in the Arrhenius plot.
Figure 4. CPMAS NMR spectra as a function of temperature: (a) PFO, (b) PFO-altco-MEHPV.
DIPSHIFT curves starts at 233 K, meaning that at lower temperatures the motional rates are in the Hz frequency scale, so that they can be ascribed as responsible for the β peak in the DMTA measurements. The same behavior is observed for both samples. Another important feature of Figure 5a is that the curves are identical above 343 K, meaning that the fast limit motion regime (motional rates . dipolar coupling frequencies) has been achieved.39 In the fast limit regime, the DIPSHIFT curves have no dependence on the motional rates, being completely defined by the residual CH dipolar coupling.32 Thus, for isotropic rotations, which average the CH
couplings to zero, the DIPSHIFT curves have no decay. In contrast, for anisotropic type of motions, such as jumps between discrete sites (except for tetrahedral sites) or rotations around a specific axis, there is a residual CH coupling that depends on the geometry of the molecular motions, leading to some decay in the DIPSHIFT curves. Assuming a rotation around a specific axis (rotation inside a cone), the reorientation angle of the molecular motions are related to the order parameter S as S ) cos θ (1 + cos θ)/2,33 which is obtained as the ratio between the motional averaged dipolar coupling (in the fast limit) and the full coupling (in the absence of motion). Thus, DIPSHIFT allows the determination of the order parameter, which can be related to the reorientation angles. To facilitate the quantification of the DIPSHIFT results, we will name the group of carbons in the end of the side chains as Cγ and those in the middle as Cβ. Thus, for PFO the group Cγ corresponds to carbons 20 and 28 (line at ∼20 ppm in the CPMAS spectrum) and Cβ to carbons 15-19 and 23-27 (line at ∼30 ppm in the CPMAS spectrum). In the same manner, for PFO-altco-MEHPV Cγ corresponds to carbons 20, 28, 45, and 47 (line at ∼20 ppm in the CPMAS spectrum) and Cβ to carbons 15-19, 23-27, 43, and 44 (line at ∼30 ppm in the CPMAS spectrum). The order parameter and the rotation angles extracted
Multitechnique Study of Polyfluorene Cast Films from the fast limit DIPSHIFT curves for PFO-altco-MEHPV were S ) 0.39 and θ ) 59° for both Cγ and Cβ carbons. This shows that despite the fact that the side chain executes relatively large angle motions, there is some restriction on the side-chain mobility as a consequence of the packing in the mesomorphic phase. Furthermore, the similarities between the DIPSHIFT curves of carbons Cγ and Cβ show that the mobility is similar along the full side chain. For the PFO the interpretation is more complicated. The DIPSHIFT curves are deeper for carbons Cβ than for Cγ, showing that overall the carbons in the end of the side chain experience a less restrict dynamics than those near the backbone. This may be because Cγ carbons are located in the end of the side chains, where differences in packing structure are probably not very effective in restricting their motions (predominance of local conformational rotations). Thus they experience a dynamics similar to that in PFO-altco-MEHPV. The fast limit DIPSHIFT curves for the Cβ carbons show that these groups are more dynamically restricted than the corresponding ones of PFO-altco-MEHPV. This is indeed an effect related to the presence of regions in PFO where the side chains are closely packed, which is consistent with the existence of an organized β phase, as suggested by the WAXS results. To support this interpretation, the DIPSHIFT curves for Cβ carbons in PFO were simulated considering two components. The first is associated to the side chains in the organized β phase, which were assumed to be fully hindered so that they contribute as rigid segments (this is of course only an approximation, since local molecular rotations in this region cannot be ruled out). The second component was assumed to execute the same motion as in the sample PFO-altco-MEHPV, which is justified by the behavior of the Cγ carbons. Combining these two components with different percentage contributions, the DIPSHIFT curve for Cβ carbons in PFO was reproduced considering that 40% of the side chains are rigid in the DIPSHIFT time scale. This compares favorably with the ratio between the peaks integral of the organized β phase and the “amorphous halo” (mesomorphic phase), which was found to be about 50% for the as cast films. The observed difference can be due to the presence of some restricted mobility in the PFO side chains in the organized β phase. This suggests that side chains in the organized β phase of PFO are more hindered as a result of the close packing, whereas those in the mesomorphic phase portion move similarly as in PFO-altco-MEHPV. This agrees with the observation of two β-relaxation peaks in DMTA, the first being associated to side-chain motions in the mesosmorphic phase and the second to side-chain motions in the β phase regions of the polymer. Other important characteristics of the molecular motion that can be determined via DIPSHIFT experiments are the correlation times. As discussed in ref 34, the more precise method to do so is based on the simulation of the 13C spin dynamics under the simultaneous presence of CH coupling and molecular motions. This is a computational costly method and requires a good knowledge of the motional geometry. Thus, because of the line superposition and the lack of information about the details of the motional geometry, we decided to analyze the data using a formula based on the so-called Anderson-Weiss (AW) approximation,34 which considers a Gaussian distribution of local dipolar fields and diffusive anisotropic motions (diffusion on a cone), so that the correlation times are the unique fitting parameter. In ref 34 it was also shown that, for DIPSHIFT evolution times shorter than half of the rotation period, AW can be a good approximation even for nondiffusive motions.34 Thus, the fittings of the DIPSHIFT curves using the AW formula (dotted lines in Figure 5a) were performed only for the first six
J. Phys. Chem. B, Vol. 113, No. 33, 2009 11409 points. The Arrhenius plots of the correlation times extracted from the curves for Cγ and Cβ carbons versus temperature of both samples are shown in Figure 5b. The corresponding activation energies are also shown. As can be observed, although the activation energies are similar in both samples for Cγ carbons, they are different for Cβ carbons. Note that the obtained values are in good agreement with those estimated from the DMTA measurements, confirming that these motions are those responsible by the β-relaxation. Interestingly, a lower Ea value was obtained for Cβ carbons in PFO. In this sense, it turns out that local rotations are predominant in the Cβ motion. Once the details of the molecular dynamics in the side chain of PFO and PFO-altco-MEHPV were characterized and associated to the β-relaxation observed by DMTA, the immediate question about the molecular nature of the processes associated with the RA- and RB-relaxation arises. From the results shown in Figure 3, it seems unambiguous that the RB-relaxation is associated to a glass transition process, but not much can be said about the RA process. To investigate the molecular nature of the RA-relaxation, the CODEX NMR method in its constant time version (CONTRA, from constant time recoupling of the anisotropies) was used. This experiment is capable of probing molecular motions with rates between Hz and kHz,23 providing reliable information about the amplitude of the molecular motions. The experiment detects the signal reduction due to changes in the orientation-dependent chemical-shift frequencies (encoded by an evolution period with duration of t1 ) s2tr, s ) 0,.., 1) that takes place during a mixing time (tm),23 usually in the order of hundreds of miliseconds. Quantification is performed by measuring the line intensities of the NMR spectrum obtained after applying the pulse sequence, S, and subtracting from the corresponding intensities obtained in a control spectrum, S0. The control spectrum is acquired in such a way that no molecular motion effects are encoded,22 but it has the same intensity reduction due to relaxation effects as the S spectra. The intensity difference obtained for each individual line in the spectra is then normalized ∆S ) (S0 - S)/S0 and plotted as a function of t1, giving rise to a curve that depends on the reorientation angle of the slow molecular motion. A first identification of the mobile groups can be obtained by simply observing the difference CONTRA spectra (S0 - S) acquired with mixing and evolution time long enough to encode the molecular motion23 (tm ) 300 ms and sNtr ) 800 µs in our case). This spectrum only has nonvanishing signals corresponding to groups presenting slow mobility. The CONTRA spectra of PFO and PFO-altco-MEHPV at 295 K revealed that all backbone carbon groups undergo slow reorientations (not shown). Thus, the RA peaks in the DMTA curves can be ascribed as being the result of rotations in the backbone of the polymers. To obtain the reorientation angles using CONTRA experiments, the chemical shift (CS) tensor (principal values and orientation of principal axes, PAS) must be known. The CS principal values can be estimated from a Hersfeld and Berger analysis40 of the spinning sidebands, but the direct measurement of the principal axis orientation is not a trivial task for chemically complex systems as the one reported here. Fortunately, for aromatic para-carbons the principal values orientation is similar for different compounds,41 being σzz nearly perpendicular to the ring plane and σxx along the symmetry ring axis. Thus, to estimate the amplitude of the backbone rotations in PFO and PFO-altco-MEHPV, it is advantageous to specifically probe the intensity of the line corresponding to carbons
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Figure 6. (a) Plot of ∆S/S0 CONTRA intensities as a function of t1 ) sNtr at 298 and 348 K for carbons 8 and 12 in PFO and 8, 12, and 30-32 in PFO-altco-MEHPV. (b) Plot of ∆S/S0 CONTRA intensities as a function of t1 ) sNtr at 298 and 348 K for carbons 2 and 5 in PFO and 2, 5, and 33-36 in PFO-altco-MEHPV.
8 and 12, where the CS PAS can be approximated to σzz nearly perpendicular to the ring plane and σxx along the symmetry ring axis. Figure 6 shows the CONTRA curves for the line corresponding to carbons 8 and 12 at 295 and 348 K for PFO and PFOaltco-MEHPV. The curves simulations were performed using chemical shift principal values of σxx ) 15 ppm, σyy ) 150 ppm and σzz ) 220 ppm, which were obtained by a Hersfeld and Berger analysis of two first order spinning side bands (νr ) 4 and 5 kHz) of the corresponding carbons. Using these values and the above-mentioned principal axis tensor orientation, the curves were calculated considering different reorientation angles for rotation about the fluorenyl group axis. Average reorientation angles for PFO of (6 ( 3)° and (15 ( 5)° were obtained at 295 and 348 K, respectively. For PFO-altco-MEHPV, the average reorientation angles increased from (7 ( 3)° to (25 ( 5)°. Similar results were obtained for all backbone carbons, showing that this is indeed a motion that involves rotation of the complete fluorenyl group. Note that the actual values for the reorientation angles may not be accurate because of the approximation in the CSA orientation, but the increase observed in the reorientation angle is quite reliable. The CONTRA NMR results of both samples revealed a similar general behavior, but the asymptotic values and the hightemperature reorientation angles of the curves are different. This difference can be explained considering that the fraction of segments that reorients in the Hz to kHz frequency scale is smaller in PFO than in PFO-altco-MEHPV, which again indicates the presence of motionally hindered segments in PFO. 3.3. Temperature Dependence of the Photoluminescence. Figure 7a shows the steady-state photoluminescence spectra of PFO and PFO-altco-MEHPV films as a function of temperature, covering the range from 20 to 380 K. The temperature dependence of the emission of PFO is very complex, because this is a morphologically complex semicrystalline material. The spectra are shaper at lower temperatures, with a higher intensity peak at 425 nm, two overlapped peaks at around 470-480 nm, and a green emission around 510 nm. Polyfluorene emission
Figure 7. (a) Fluorescence spectra as a function of temperature. (b) Integrated spectral intensity as a function of temperature. (c) Maximum position of the 0-0 band as a function of temperature. (left) PFO and (right) PFO-altco-MEHPV.
around 420 and 500 nm can be assigned to the isolated fluorenyl moieties and aggregated species, respectively.35 This behavior is indicating that more than one species are emitting. Moreover, these spectra are blue-shifted under the sample heating, and the vibronic structure is lost as reported for other polyfluorenes.14,15 The temperature dependence of the integrated and normalized spectra IF(T)/IF(T0) can be used to determine the temperatures of the polymer relaxation processes. For PFO, there is a pronounced change of the spectral profile with the increase of temperature, corroborating the assumption that multiple species are emitting and the relative amount of every one is changing during the heating. Furthermore, both the integrated intensity versus temperature curve (Figure 7b) and the relative intensity of the 0-0 band change upon heating, with changes in the slope of both curves. From both curves there is an initial decrease of the intensity from 20 to 100 K, and then the intensity becomes approximately constant from 110 to 220 K, and for T > 220 K there is an increase of the intensity with the temperature for the two peaks at 330 and 380 K. In addition, the 0-0 band peak position presents clear blue shifts where the curve profiles changes, i.e, at the temperature similar to the changes of the emission intensities: it decreases almost monotonically until 210 K, remains almost constant from 220 to 350 K, and then decreases again. For PFO-altco-MEHPV, the spectral behavior is different (Figure 7a). The spectral profile at lower temperatures is redshifted compared with the PFO. This emission centered at 510 nm originates probably from the short segments of the MEHPV groups. In addition, the broad spectra at lower temperatures indicate the presence of conformational disorder of the main chains that confers to every group in the electronic excited state
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TABLE 2: Temperatures and Assignments of Relaxation Processes Probed by WAXS, DMTA, PL, and NMR PFO NMR Tγ: motions of short Segments (γ-relaxation) Tβ: local motions of specific segments (β-relaxation) TRA: slow backbone rotations TRB: glass transition, slow motions, increase of the interchain distances melting of the R phase
experiences different environments, which then are unable to relax within the excited state lifetime.36 The temperature dependence of the fluorescence intensity decays monotonically but not monoexponentially from 20 to 290 K, consistently with the decay of systems with more than one value of activation energy (Figure 7b).37 Around 320 K, the intensity becomes approximately constant. Differently, the position of the 0-0 band (Figure 7c) undergoes a continuous blue shift with the increase of the temperature, but the curve shows slope changes around 140, 220, 300, and 370 K. 4. Discussion In general, the emission of conjugated polymers in the solid state is strongly affected by modifications in their local structure and packing. Because such changes are usually dynamic related processes, it is important to understand at the molecular level how molecular motions affect the structural features. This can be achieved by correlating the temperature dependence of the photoluminescence processes with the microstructure and dynamic processes as revealed by X-ray, NMR, and DMTA, which is the main aim of this section. The DMTA data show that both polymers have a lowtemperature relaxation at ∼220 K and a high-temperature relaxation at ∼370 K (Table 2). The microscopic nature of these relaxations was elucidated by solid-state NMR methods capable of detecting molecular dynamics of individual chemical groups.38 It was shown that the low-temperature relaxation is associated with motions in the side chains, occurring with activation energies of ∼20 kJ/mol in both polymers. Despite the similar activation energies, it was also observed that the fraction of mobile side chains is higher in the PFO-altco-MEHPV sample, pointing to the presence of rigid side chains in PFO, probably in the organized phase. The NMR measurements also revealed the presence of slow motions (miliseconds correlation time range) in the polymers backbones, the amplitude of which increases as a function of temperature. These results show that the onset of the side-chain motion produces an increase of the lateral spacing of the phenyl rings in the aggregated structures, facilitating the onset of torsional motion in the backbone, which were found to have ∼7° at 298 K, increasing to ∼25° at 348 K. This torsional motion can be considered responsible for the slightly blue shift observed in the fluorescence spectra of the PFO-altco-MEHPV copolymer. Besides, at ∼373 K the increase in the amplitude of the backbone motions induces an increase in the conformation disorder of the phenyl rings, producing stronger blue shift above this temperature. Despite the motional behavior in PFO and PFO-altco-MEHPV seems to be similar, the fluorescence spectra of PFO cannot be fully explained only considering the dynamic aspects, confirming that the phase transitions observed for PFO play a key role in its luminescence. Despite that, the Vogel-Fulcher-like temperature dependence observed for the DMTA peak at ∼363 K indicated that this process is likely to be attributed to a glass transition process. Of course, these molecular motions favor the expansion
WAXS
370 K 430 K
PFO-coalt-MEHPV
DMTA
PL
130 K 220 K 330 K 380 K
130 K 220 K 330 K 380 K
WAXS
DMTA
PL
380 K
200 K 330 K 370 K
140 K 220 K 300 K 370 K
of the aggregated structures, justifying the increase observed in the d1 and d3 WAXS characteristic lengths. As soon as the aggregates are destroyed (in an optical point of view), the photoluminescence efficiency increases because the emission quantum yield of the isolated groups is higher than the interchain or the intrachain aggregates.36 The increase of the intensity is a result of the greater quantum yield of the emission from the isolated groups, which compensates the increase of the nonradiative rate constants at higher temperatures. The DMTA results of the PFO sample show similar behavior as compared to PFO-altco-MEHPV. The NMR results show that the β-relaxation at low temperatures is also due to movements in the side chains, which are indeed very similar in both samples. Thus, as for PFO-altco-MEHPV, this explains the initial decrease in the fluorescence intensity at lower temperatures and the initial increase at 133 K (onset of conformational motions in the backbone after the β-relaxation). However, this explanation is not satisfactory for the strong intensity increase observed already at 213 K, because the amplitude of the backbone motions would be too small (they are only 10° at 298 K) for producing such a strong dissociation of the inter chain species. A possible explanation for this intensity increase after the β-relaxation process might be the change of the polymer local conformations induced by the motion of the side-chain groups, leading to the formation of species with higher emission quantum yield.7 However, further measurements will be necessary to confirm such statements. The DIPSHIFT NMR results for PFO-altco-MEHPV indicated that, at lower temperatures, the main dynamic processes are the side-chain motions produced by the β-relaxation process. Despite being local rotations, these are high amplitude motions involving almost the entire side chains, which require higher free volumes as the motional rates increase as a function of temperature. Thus, taking into account the presence of π-stacked aggregated structures, as shown by the WAXS results, it is possible to consider these motions as being responsible for the monotonically temperature increase of the WAXS characteristic length d1, which is associated with the lateral spacing of the rings. Note also that the increase in the side-chain motion may produce enough free volume to allow local torsions motions in the polymer backbone, producing a small increase in the interplanar ring distances as observed by WAXS. The presence of these torsion motions were confirmed by the CONTRA NMR results that detected slow rotations of the backbone rings with average amplitudes of ∼7° at 298 K. The high-temperature behavior of the PFO-altco-MEHPV fluorescence spectra also correlates well with the observed dynamic process. First, the increase in the backbone torsion motions as a function of temperature directly explains the increasing blue shift observed in the PL spectra as a function of temperature. Second, the stabilization of the PL intensity in the 300-340 K range can be understood as a competing
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effect of the increase in the system overall dynamics, which induces a decrease in the fluorescent emission, and the dissociation of the interchain species as a result of the increasing of torsional disorder, which favors the radiative emission. This is actually justified by theoretical ab initio calculations, which showed that deviation for the planar packing of the rings can lead to the dissociation of the interchain species.7,39 This effect is enhanced at ∼350 K, so it produces an increase in the PL intensity and a higher blue shift above this temperature. It is worth mention that the increase observed in the d3 parameter is much smaller than the Fo¨rster length, so it is too small for producing the dissociation of the interchain species, showing that this dissociation is indeed due to the modifications in the chain conformation. Finally we would like to explicitly correlate the dynamic results with temperature behavior of the characteristic WAXS length parameters d1, d2, and d3 observed in the PFO-altcoMEHPV. Below ∼380 K a clear variation of d1 is observed, while d3 remains almost constant. This behavior can be directly associated with the dynamic results, which showed that in this temperature range the side-chain motion is predominant. Of course, this side-chain motion may produce a substantial increase in the lateral spacing of the rings (d1 parameter) but should not produce a significant change in the interplanar ring distance (d3 parameter). This is in agreement with the results observed. For temperatures higher than 380 K, the side-chain motion is already fully established, but the backbone motion start to be more effective, producing a more pronounced increase in the interplanar ring distance as function of temperature. This is again in fully agreement with the results observed in Figure 2b. However, an “anomalous” behavior is the fact that at 380 K the variation of d1 with temperature decreases. This can be explained, considering that as a result of the gain of mobility of the backbone, the free volume of the side-chain region is already sufficient for the side chains to move almost freely, so no significant change is induced in the lateral spacing of the rings, as observed in Figure 2b (left). It is also worth mentioning that the temperature at which the changes in d1 and d3 occur (∼380 K) is slightly higher than the glass transition observed by DMTA (370 K). This is actually an expected result since for the Tg measured by DMTA, the backbone molecular motions is slow (∼10 Hz), and it is not surprising that they become sufficient to produce the changes in the spacings only at temperatures that are just a bit higher than the Tg as measured by DMTA. Taken together (Table 2), the results explains the two main aspects of the low-temperature side of the PFO-altco-MEHPV emission: (a) the decrease in the total intensity, which can be associated with the increase in the efficiency of nonradiative pathways induced by the molecular dynamics, and (b) the blue shifts observed as a function of temperature, which can be attributed to the appearing of the backbone torsional motion, whose amplitude increases from 7° at 300 K to 25° at 340 K as observed in the CONTRA experimentss. This relaxation process is likely to be the one responsible for the shoulder observed in the high-temperature peak observed in the DMTA and assigned as the RA-relaxation. 5. Conclusions In this article we presented an investigation of the microstructure and dynamics of two representative polyfluorene-based polymers, poly(9,9-dioctylfluorenyl-2,7-diyl) (PFO)
Faria et al. and poly[(9,9-dioctyl-2,7-divinylene-fluorenylene)-alt-co-{2methoxy-5-(2-ethyl-hexyloxy)-1,4-phenylene vinylene} (PFOaltco-MEHPV), using WAXS, solid-state NMR, DMTA, and PL spectroscopy. The combined results of this multitechnique study allowed the identification of four polymer relaxation processes, and we were able to describe the type of macromolecular relaxation, molecular motion and supramolecular structure ascribed to every one of them. Shorter segmental motions from the lateral groups of the polymer chains were observed around 130 K (γ-relaxation), local motions were observed at 200-220 K (β-relaxation), glass transitions were observed at 330 and 380 K, and the melting temperature of the ordered phase was observed for PFO at 430 K. Excluding the last, all processes change the photoluminescence processes in terms of the emission profile, peak positions, and relative intensity of the emission bands. Acknowledgment. Financial support of the Brazilian funding agencies FAPESP, CNPq, MCT/INEO, and CAPES are gratefully acknowledged. We also thank the Brazilian Synchrotron Light Laboratory for the use of the D02ASAXS2 beamline in the framework of project D11A-SASX5379. The authors thank Prof. R.M. Faria for providing the DMTA facility. E.R.dA thanks G. L. Mantovani for useful suggestions. References and Notes (1) deAzevedo, E. R.; Franco, R. W. A.; Marletta, A.; Faria, R. M.; Bonagamba, T. J. J. Chem. Phys. 2003, 119, 2923. (2) Cossiello, R. F.; Akcelrud, L.; Atvars, D. Z. J. Braz. Chem. Soc. 2005, 16, 74. (3) Bloise, A. C.; deAzevedo, E. R.; Cossiello, R. F.; Bianchi, R. F.; Balogh, D.; Faria, R. M.; Atvars, T.; Bonagamba, T. J. Phys. ReV. B 2004, 71, 1712001. (4) Bozano, L.; Tutle, S. E.; Carter, S. A.; Brock, P. J. Appl. Phys. Lett. 1998, 73, 3911. (5) Bozano, L.; Carter, S. A.; Scott, J. C.; Malliaras, G. C.; Brock, P. J. Appl. Phys. Lett. 1999, 74, 1132. (6) Lupton, J. M.; Samuel, I. D. W. Synth. Met. 2000, 111, 381. (7) Ferretti, A.; Ruini, A.; Molinari, E.; Caldas, M. J. Phys. ReV. Lett. 2003, 90. (8) Machado, A. M.; Neto, J. D. D.; Cossiello, R. F.; Atvars, T. D. Z.; Ding, L.; Karasz, F. E.; Akcelrud, L. Polymer 2005, 46, 2452. (9) Machado, A. M.; Munaro, M.; Martins, T. D.; Davila, L. Y. A.; Giro, R.; Caldas, M. J.; Atvars, T. D. Z.; Akcelrud, L. Macromolecules 2006, 39, 3398. (10) Yang, S. H.; Lee, C. F. J. Optoelectron. AdV. Mater. 2007, 9, 2078. (11) Chang, L. H.; Lee, Y. D.; Chen, C. T. Macromolecules 2006, 39, 3262. (12) Zhang, H. A.; Li, Y.; Jinag, Q.; Xie, M. G.; Peng, J. B.; Cao, Y. J. Mater. Sci. 2007, 42, 4476. (13) Akcelrud, L. Prog. Polym. Sci. 2003, 28, 875. (14) Grey, J. K.; Kim, D. Y.; Donley, C. L.; Miller, W. L.; Kim, J. S.; Silva, C.; Friend, R. H.; Barbara, P. F. J. Phys. Chem. B 2006, 110, 18898. (15) Grell, M.; Bradley, D. D. C.; Ungar, G.; Hill, J.; Whitehead, K. S. M. Macromolecules 1999, 32, 5810. (16) Chen, S.; Su, A.; Su, C.; Chen, S. A. Macromolecules 2005, 38, 379. (17) Chen, S. H.; Chou, H. L.; Su, A. C.; Chen, S. A. Macromolecules 2004, 37, 6833. (18) Chen, S.; Su, A.; Su, C.; Chen, S. J. Phys. Chem. B 2006, 110, 4007. (19) Chen, S. H.; Su, A. C.; Chen, S. A. J. Phys. Chem. B 2005, 109, 10067. (20) Winokur, M.; Slinker, J.; DL, H. Phys. ReV. B 2003, 67, 184106. (21) Souza, A. A.; Cossiello, R. R.; 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, 4063. (22) deAzevedo, E. R.; Hu, W. G.; Bonagamba, T. J.; Schmidt-Rohr, K. J. Am. Chem. Soc. 1999, 121, 8411. (23) Reichert, D.; Pascui, O.; Bonagamba, T.; Belton, P.; Schmidt, A.; deAzevedo, E. J. Magn. Reson. 2008, 191, 141.
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J. Phys. Chem. B, Vol. 113, No. 33, 2009 11413 (35) Oliveira, H. P. M.; Martins, T. D.; Hono´rio, K. M.; Rodrigues, P. C.; Akcelrud, L.; Silva, A. B. F.; Atvars, T. D. Z. J. Braz. Chem. Soc. 2009, 20, 160–166. (36) Lakowicz, J. R. Principles of Fluorescence Spectroscopy; Academic Press: New York, 1999. (37) Guillet, J. Polymer Photophysics and Photochemistry; Cambridge University Press: Cambridge, 1985. (38) deAzevedo, E. R.; Bonagamba, T. J.; Reichert, D. Prog. Nucl. Magn. Reson. Spectrosc. 2005, 47, 137. (39) Ruini, A.; Ferretti, A.; Bussi, G.; Molinari, E.; Caldas, M. J. Semicond. Sci. Technol. 2004, 19, 362. (40) Herzfeld, J.; Berger, A. E. J. Chem. Phys. 1980, 73, 6021. (41) Veeman, W. S. Prog. Nucl. Magn. Reson. Spectr. 1984, 16, 193.
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