Impact of Confinement-Induced Cooperative Molecular Orientation

Feb 18, 2013 - The dipole orientation change showed a good correlation with the thickness dependence of remnant polarization, which were measured from...
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Impact of Confinement-Induced Cooperative Molecular Orientation Change on the Ferroelectric Size Effect in Ultrathin P(VDF-TrFE) Films Dong Guo*,†,‡ and Nava Setter† †

Ceramics Laboratory, Swiss Federal Institute of Technology (EPFL), CH-1015 Lausanne, Switzerland Institute of Acoustics, Chinese Academy of Sciences, Beijing 100190, China

‡

S Supporting Information *

ABSTRACT: We demonstrate a distinct confinement induced cooperative lattice orientation change in ultrathin P(VDF-TrFE) films by using various characterization techniques, including X-ray diffraction (XRD), two-dimensional X-ray diffraction (2D-XRD), grazing incidence X-ray diffraction (GIXD), and infrared reflection absorption spectra (IRRAS). Both the polymer chains and the molecular dipoles that are perpendicular to the polymer backbone showed a lying-down orientation change with decreasing thickness. The dipole orientation change showed a good correlation with the thickness dependence of remnant polarization, which were measured from highly reproducible ferroelectric loops in a wide range of thickness (30−250 nm) on inert electrode with suppressed “dead layer” effect. A simple microscopic molecular dipole interaction model revealed that the free energy of different orientation states was related to the thickness and lateral dimension of the polymer crystallites. The findings reveal a unique molecular orientation driven size effect in ferroelectric polymer films, providing new insights into the nature of ferroelectricity and orientation mechanisms in polymers relevant to the design of emerging flexible electronic devices.

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INTRODUCTION The copolymer of vinylidene fluoride and trifluoroethylene [P(VDF-TrFE)] is a typical organic ferroelectric material that has been widely used in sensors, actuators and other microelectromechanical systems. The rapidly growing flexible electronics industry has recently stimulated intense research on new devices based on thin P(VDF-TrFE) films. In addition to flexible nonvolatile memories,1,2 other novel device functionalities, such as photostrictive effect for actuators3 and enhanced charge separation in organic photovoltaic devices were demonstrated by using the ferroelectric layer.4 Because of its very large coercive field of about 50 MV/m, “ultrathin” (thickness t < 100 nm) P(VDF-TrFE) films are required to achieve low operation voltage. The continuous demand for device miniaturization also needs “ultrathin” films to maintain a suitable capacitance, which is inversely proportional to the film thickness. However, the reduction in structural dimension causes deteriorated ferroelectricity.5 Despite the significantly improved electrical performance achieved recently,6−8 the chemical and physical mechanisms underlying such a size effect in P(VDF-TrFE) films are still far from being clear. Elucidating the details is of great importance from both fundamental science and application points of view. The much more complicated structural characteristics of ferroelectric polymers relative to those of inorganic ferroelectrics, such as the partial crystallinity and the many possible conformations and configurations, etc., cause a high sensitivity of their microstructure to external factors. Consequently, © 2013 American Chemical Society

ferroelectric polymers show rather different size dependent properties from those of inorganic ferroelectrics. Polarization switching was observed in Langmuir−Blodgett (LB) P(VDFTrFE) films with a thickness of only one molecular monolayer.9 The driving forces for the size effect in inorganic ferroelectrics are well understood, which include long-range depolarizing field due to incomplete polarization screening in the electrodes,10 short-range coupling of the order parameter with the interface,11 and surface tension.12 In contrast, a comprehensive understanding of the driving force of size effect in ferroelectric polymer films is still lacking, although decreased crystallinity and electrode interface effects were proposed.5,13 In “ultrathin” polymer film, the mobility, packing order, and orientation of the polymer chains usually show a complicated or even conflicting thickness dependence due to the special enthalpic and entropic effects at the polymer/substrate interface under the confined geomtry.14,15 The ferroelectricity in P(VDF-TrFE) originates from the alignment of F−C−H covalent molecular dipoles that are perpendicular to the polymer backbone; therefore, the dipole orientation should play a crucial role in the size effect of the film if it is thickness dependent. However, this issue seems to have never received enough attention, although the lattice orientation of the copolymer films under certain conditions has been studied.16 Received: November 18, 2012 Revised: January 27, 2013 Published: February 18, 2013 1883

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Figure 1. (a and b) P−E hysteresis loops of P(VDF-TrFE) films on different substrates. In part b, a distorted loop of a 30 nm thick film on Si was also shown to illustrate the effect of the SiO2 “dead layer”. (c and d) Thickness dependence of Pr and Ec, respectively. spin coated on the substrates by using methyl ethyl ketone as the solvent. The film thickness was controlled by the solvent concentration. Nonleaky films were obtained after annealing at 135 °C in air for 70 min by using a Prazitherm PZ28−2 (Harry Gestigkeit Co., Germany) precision hot plate, which has a temperature accuracy of ±0.5 °C. The hot plate was then switched off, and the samples were naturally cooled to room temperature. The annealing temperature is slightly smaller than the melting point (∼145 °C). At this temperature the films show the highest Pr while still keep a dense and homogeneous morphology with small granular grains.19 Capacitor units were made by vacuum evaporation of 50 nm Au top electrodes through a shadow mask. Film thickness was repeatedly measured by atomic force microscopy (AFM) height profiles of scratches made on different places of the same sample. P−E loops were measured with a 100 Hz triangular voltage waveform. All the loops are highly reproducible and were obtained on P(VDF-TrFE) ultrathin films without any buffer layers. Out-of-plane X-ray diffraction (XRD) patterns were measured by a Bruker Advanced D8 diffractometer. Inplain grazing incidence X-ray diffraction (GIXD) patterns were measured by a Rigaku Smartlab diffractometer (Cu Kα radiation). 2D-XRD images were recorded by a Bruker D8 Discover diffractometer with a GADDS detector. Infrared reflection absorption spectra (IRRAS) were measured with a Bruker EQUINOX 55 FTIR spectrometer equipped with a liquid nitrogen-cooled mercury cadmium telluride detector. The spectra were recorded under a grazing incident angle of 85° with a Spectra-tech FT 85 reflection adsorption accessory. Au-coated Si wafers were used both as the background sample and as the substrates for the film samples, and 1200 scans were collected and averaged at a resolution of 2 cm−1 to ensure high signal-to-noise ratios. Large film samples (2 cm × 7 cm) were used to ensure full coverage of the rectangular sampling window under grazing incidence, and all the spectra were obtained in a same run to rule out the influence of environment or instrument variation for different runs. These procedures are crucial for getting high-quality IRRAS spectra. All electrical and structural characterizations were conducted at room temperature.

To clarify the details, a precise correlation between the thickness dependence of molecular orientation and ferroelectric properties of ultrathin P(VDF-TrFE) films is essential. The experimental difficulties come from preparation of homogeneous nonleaky ultrathin films with a wide range of thickness for direct measurement of remnant polarization (Pr) and coercive field (Ec) through polarization−electric field (P−E) hysteresis loops, as well as precise characterization of their macromolecular orientation change. Because of the difficulty in nonleaky film preparation, so far the relevant work on the LB9,13,17 and spin coated films6,7 have been focused on films thinner and thicker than 60 nm, respectively. However, an investigation on the same films with a wider thickness range is important because of the different microstructure of the two types of films.6,9,18,19 Spin coating seems more challenging than the LB technique in preparing ultrathin films with directly measurable loops. Although the loop of a 15 nm thick spin coated P(VDF-TrFE) film has been measured, the measurement was performed on “electrically” treated film capacitor that may have a special structure.20 Paying special attention to these problems, here we investigate the role of microstructural characteristics in controlling the size effect in spin coated (normal thermally treated) P(VDF-TrFE) films with a wide thickness range of 30−250 nm. The results demonstrate an unique molecular orientation origin of size effect in the ferroelectric polymer films, whose driving force is quantitatively explained by a simple microscopic dipole interaction model that considers the free energy of different orientation states.

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EXPERIMENTAL SECTION

Highly doped bare Si (100) wafers (resistivity of 0.001−0.005 Ω·cm) and 100 nm Al and Pt coated low doped Si wafers (hereafter referred to as Al and Pt for simplicity) were used as the substrates, which were ultrasonically cleaned in acetone, isopropanol, and deionized water, and finally blown dry with nitrogen before use. Ultrathin copolymer films with a VDF/TrFE mole ratio of 70/30 (Solvay Solexis S.A.) were 1884

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RESULTS AND DISCUSSION

P−E loops of various films are shown in Figure 1, parts a and b. To ease comparison, the loops of films with different thickness on the same substrate are measured under the same electric field (300 V/μm equals to a 9 V voltage for a 30 nm thick film). The different positive and negative coercive electric field of the loops is due to the built-in field caused by the different work functions of the top and the bottom electrodes. Films thicker than 100 nm on different substrates show similar regularly shaped loops. With decreasing t, the effect of the bottom electrodes becomes prominent. The gradual while substential tilt of the hysteresis loops for films on Al is a clear signature of a “dead layer” at the film/electrode interface. 21 Similar phenomena were observed for films on Si (see the loop of 30 nm thick film on Si in Figure 1b). The tilt of loops on the inert Pt starts relatively more abruptly for films with a thickness of 40 nm and below. Nevertheless, the loop of 30-nm thick film still keeps a regular shape with a high Pr of 5.43 μC/cm2. Also reflected from the thickness dependence of Pr (Figure 1c) and Ec (Figure 1d), thinner films (t < 60 nm) on Pt have higher Pr than those on Al and Si. These results indicate that the “dead layer” effect is largely suppressed for films on Pt. The “fish-like” loops of the 30-nm thick films on Al and Si are so distorted that the derived Ec data deviates from their Ec curves and can not reflect the exact field to reverse the polarization. In Figure 1c, the Pr decreases more rapidly with decreasing t, but the decrease is relatively slower than those reported in previous studies (e.g., Figure 4 of ref 5).5,22 This might be due to the different processing conditions. Evidently, the deteriorated ferroelectricity for films on Si and Al are largely due to depolarization field caused by native surface oxide “dead layers”, whose thickness is about 1−2 nm on deionized water cleaned Si (100)23 and about 4 nm on Al surface.24 Although it is unclear whether there is a surface oxide layer or not, Pt is more similar in nature to a “realistic” electrode compared to other electrodes as observed in previous studies.10,25,26 These results imply the existence of other types of driving forces of size effect for the films on Pt in addition to the depolarization effect caused by the “dead layer”. For electron diffraction characterization of polymer film structure, the difficulties lie in the preparation of electrontransparent specimens that cannot be destroyed during measurement. The specific problem in analyzing the quasihexagonal lattice orientation of P(VDF-TrFE) films by diffraction techniques is the serious peak overlap. Here several techniques were used for film structure analysis. As shown in Figure 2, the out-of-plane XRD patterns of the films have a broad peak at 2θ of ∼19.7° due to overlapped (200/110) reflections, revealing a β phase structure without any reconstruction in the tested thickness range. For textured films the crystallinity is difficult to be derived by the XRD intensity, while the strong diffraction signal of the 30 nm film should reflect its considerable crystallinity. The 2D-XRD images of 40-nm thick films and the GIXD patterns of 40 and 250-nm thick films are shown in Figure 3, parts a and b, respectively (The diffraction geometry is illustrated in Figure S1 in Supporting Information). There is a strong center part located at χ (the rotation angle in the plane perpendicular to the diffraction plane) of 90° in the (200/110) diffraction arc in the left panel of Figure 3a, indicating that the corresponding planes, namely, the 6 sides of the P(VDF-TrFE) hexagonal cell, are on average preferentially oriented parallel (or ±60° tilted)

Figure 2. Out-of-plane XRD patterns of the films on Si substrate.

Figure 3. (a) Typical 2D-XRD patterns of 40 nm thick films on Si obtained under a χ angle of 90° (left) and 60° (right). (b) GIXD patterns of two typical films with a thickness of 250 nm (upper) and 40 nm (lower).

in relation to the substrate (see Figure 6). As seen in the right panel of Figure 3a, the position of the strongest part changed correspondingly when χ is changed to 60°. This confirms the orientation revealed by the left image. 250-nm thick film shows almost the same 2D-XRD images except for a higher intensity (not shown), indicating that the (200/110) plane orientation state is independent of film thickness. In Figure 3b, two peaks that can not be detected by the out-of-plane XRD appear at 2θ of ∼35.2° and ∼40.8°, which are attributed to the overlapped (020/310/001) and (201/111) diffractions,18,27 respectively. The diffraction intensity of the 40-nm thick film is much weaker than that of the 250 nm one. What deserves noting is 1885

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that the strongest (200/110) diffraction peak in the 250-nm thick film becomes the weakest one in the 40-nm thick film. Since in in-plane GIXD only the lattice planes perpendicular to the substrate can be detected, a reduced relative intensity of the (200/110) diffraction implies a reduced ratio of the sides of the P(VDF-TrFE) hexagonal cell that are perpendicular to the substrate in the thinner film. While because the other two peaks are due to planes intersecting with each other and with the lattice axes, their overall intensity ratio is relatively insensitive to this lattice orientation change. Note that almost the same diffraction results have been observed for films on different electrodes, indicating that the film microstructure is insensitive to the bottom electrodes used here (see Figure 1 and Figure 7). This may be due to the weak van der Waals interaction between the copolymer and the inorganic substrate. The IRRAS technique that works through the “surface selection rule” and detects only the p-polarized IR component (with transition dipole moments perpendicular to the substrate plane) is particularly useful for analyzing the molecular orientation of ultrathin films. The IRRAS spectra of the films are shown in Figure 4a. The peak at ∼1205 cm−1 is attributed to the vibrational modes (va(CF2) coupled with r(CF2)) parallel to the a-axis of the lattice, and the ∼1297 cm−1 (vs(CF2) coupled with vs(CC) and δ(CCC)) and ∼1403 cm−1 (ω(CH2) coupled with va(CC)) bands are due to dipoles parallel to b-axis and c-axis, respectively.28,29 Since the dipoles corresponding to the three strongest bands at ∼1205, ∼1297, and ∼1403 cm−1 are in orthogonal directions, the “relative intensity” method was used to determine the lattice orientation. The corresponding dichroic ratios (intensity ratios) Dc/b, Db/a, and Dc/a are plotted in Figure 4b, which provides a quantitative estimation of the orientation degree. With decreasing t, Dc/b decreases only slightly while Db/a and Dc/a decrease sharply from 1.28 to 0.48 and from 0.94 to 0.28, respectively, and both curves show a turning point at around 40 nm. Because only the p-polarized IR component is active, a higher intensity of a band means a relatively more perpendicular orientation of the corresponding axis in relation to the film surface. Therefore, the Db/a and Dc/a curves in Figure 4b imply an interesting truth: the molecular dipole (parallel to b axis) and the polymer backbone (parallel to c axis) show a cooperative lying down movement with decreasing t, and the movement is relatively stronger below 40 nm. The overall macromolecular orientation change is schematically illustrated in Figure 5 through the movement of the polymer lattice cell in a three-dimensional coordinate system, where the included angle between the b axis and the surface plane (α) is visualized by the projection of dipole on the film plane, and α2 is smaller than α1. The backbones of the asspun polymer films are well-known to tend to orient parallel to the film plane.15 Previous study has shown that submelt annealing enhanced such a orientation tendency in P(VDFTrFE) films.19 Therefore, The drop in Dc/a suggests that the P(VDF-TrFE) backbone lies further down with decreasing t, tending to adopt an “edge-on” lamellar grain orientation. In addition, if the molecular dipoles in thick P(VDF-TrFE) films are randomly oriented as in a bulk sample, the substantial drop in Db/a suggest that they are on average largely (though not completely) oriented parallel to the film plane in ultrathin P(VDF-TrFE) films. Actually, the lying down movement of the polymer chains with decreasing t together with the 0° or ±60° tilted (200/110) plane (parallel to the polymer chains) reflected by 2D-XRD imply a reduced ratio of vertically oriented (200/110) planes

Figure 4. (a) IRRAS spectra of P(VDF-TrFE) films (30−250 nm). (b) Thickness dependence of dichroic ratios based on the three strongest bands located at ∼1205, ∼1297, and ∼1403 cm−1.

that are active in GIXD of thinner films. Hence, Figure 3 well agrees with the IRRAS spectra and the orientation change shown in Figure 5. Furthermore, because no significant microstructural and morphological change was found for films on different electrodes in the tested thickness range by using different characterization techniques (see Figure S1 of the Supporting Information), Figure 5 reflects the structural change of films on different substrates in used in this study. Comparing Figure 4b with Figure 1c, one can find that Db/a and Pr show a very similar thickness dependence and both data decrease more rapidly with decreasing thickness. In comparison, the large drop in Pr as shown in Figure 1c is hard to be well explained by a decreased crystallinity (see Figure 2), because the thinnest 30 nm thick film still show a considerable crystallinity. Also, the strain effect can be neglected since “lattice matching” is not applicable to the polymer/inorganic heterostructure bonded with weak van der Waals interaction. Therefore, it is reasonable to conclude that the confinement induced lying down molecular dipole orientation change with decreasing t constitutes a driving force that has never been recognized previously for the size effect in the thin P(VDF1886

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Figure 5. Schematic illustration of the confinement-induced overall P(VDF-TrFE) molecular orientation change with polymer lattice (a) and with a unit cell (b), where the arrows represent the dipoles. Dipoles with opposite directions have an identical orientation change as illustrated.

Figure 6. Schematic illustration of the dipole - dipole interactions with parallel dipoles (case 1, a) and 60° tilted dipoles in relation to the substrate (case 2, b). The red and green slabs represent the head-to-tail and side-by-side dipole interfaces, respectively, and d(200) is not equal to d(110).

TrFE) films. No doubt that during the ferroelectric switching the dipoles change their orientation. A question might therefore arise is why the initial dipole orientation prior to the application of the electric field is correlated with Pr. There may exist two main reasons: first, once the electric field is removed, the dipoles tend to back-switch to their energically more favored original orientation (the driving force will be discussed in the following section); second, because of the slightly different (110) and (200) interplanar distances (see Figure 6), tilting of dipoles away from their initial position after switching requires lattice dimension change that is associated with large steric barrier, while 180° changing of dipoles (e.g., no orientation change) has no such a barrier. Consequently, a good correlation between Db/a and Pr is observed. These two reasons also qualitatively explain the increased Ec with decreasing t, as the non-180° rotation of more flat lying dipoles in thinner films during polarization switching may induce a larger steric barrier. Because in thin films the available conformations are restricted under constrained geometries and the dimension of chains in the direction normal to the surfaces is reduced, the lying-down of the P(VDF-TrFE) backbone with decreasing t should be due to the geometry confinement effect, similar to that observed in other polymer ultrathin films. What deserves special attention is the driving force of the confinementinduced lying down of dipoles. Ferroelectricity in the copolymer arises from the alignment of molecular dipoles via the balance of short-range repulsive van der Waals interactions and long-range attractive dipole−dipole interactions between the macromolecular chains. The intermolecular van der Waals interaction is a constant for a specific lattice structure. The interactions between the permanent dipoles in the copolymer are one of the fundamental interactions that hold the macromolecules together as a condensed phase. The local dipole−dipole interactions have been used to describe the typical dielectric and ferroelectric properties in PVDF.30 Ab initio calculations also revealed that dipole−dipole interaction

was responsible for their complicated orientation states.31 Therefore, it seems reasonable to attribute the driving force of the dipole orientation change to a specific variation of the dipole−dipole interaction energy, and a simple model can be proposed in the following. For simplicity, we assume a polymer chain as a single dipole and the plane perpendicular to the polymer chain can be schematically illustrated in Figure 6. Because the molecular dipoles are on average preferentially oriented parallel or tilted at an angle of ±60° in relation to the substrate, two types of dipole interactions can be defined: the head-to-tail interactions (eht) and the side-by-side interactions (ess). Because of the Columbic nature of the dipole interactions, eht should be large than ess. This is analogous to the anisotropic hydrogen bonding. Actually, the aligned dipoles in a polymer chain may not switch simultaneously, while the relative strength of the interchain dipole interaction is not affected. For a film with a thickness of N and an average lateral crystallite size of W (N and W are the number of polymer chain monolayers), the sum of the dipole attraction energy for the two cases with the dipoles oriented parallel (Ea1) and tilted at an angle of ±60° in relation to the substrate (Ea2) can be derived: Ea1 = (W − 1)Neht + (2N − 2)(W − 1)ess + (N − 1)ess (1)

Ea2 =

2N − 1 N−1 × NWeht + × (W − 1)N ess N N

= (N − 1)Weht + (2N − 1)(W − 1)ess

(2)

The Gibbs free energy difference (inverse of the attraction energy difference) is: 1887

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corresponding to a lateral crystallite size in the range of 30− 50 nm. This value is close to the thickness where Pr and dichroic ratios (Figure 1c and 4b) start to change more rapidly. The values are reasonable compared to the grain size of about 70−80 nm (see Figure 7a, 7b and the inset of Figure 8) because the coherent single-crystalline region is smaller than a grain. The eht − ess of about 20−40 kJ/mol is in a normal range for dipole interactions.

ΔE = −(Ea1 − Ea2) = (N − W )eht + (W − 1)ess − (N − 1)ess = (N − W )eht + (W − N )ess = (N − W )(eht − ess)

(3)

The relatively weak interaction of the crystallites with the surrounding disordered regions or inorganic substrate should be similar for two cases and hence not included in eq 3. The equation implies that when N < W, a lying down dipole orientation (case 1) is energetically favorable, which has minimized net electric field at the surface. Then, following the Maxwell−Boltzmann statistics, Pr can be expressed via the ratio of crystallites with ±60° tilted dipoles: ∑ Pi ∑ Pit 2e−Ea2 / kT 2 = = P0 −Ea2 / kT V V 2e 2 + 2e−Ea1/ kbT 1 2 1 = P0 −ΔE / kbT (W − N )(eht − ess)/ kbT 2 1 + 0.5e 1 + 0.5e

Pr =

(4)

where Pit, kb, T and P0 are the number of total dipoles, the Boltzmann constant, the temperature and the polarization when all dipoles are vertically oriented (close to saturation polarization), respectively. A degeneracy of 2 is used considering the two identical states in case 2. The complicated structural characteristics of polymer make it difficult to give a clear correlation between N and W (W is different from the vertical crystallite size derived from out-ofplane XRD). The factors that affect W can be seen from the morphology of films prepared under different conditions. Figure 7 shows the morphology of films that have been

Figure 8. Fitting of the thickness dependence of Pr for the films on Pt as shown in Figure 1c according to eq 4. The inset is the AFM image of a 85-nm thick film on Pt.

In the forgoing calculation, the orientation is enforced only by the relative dipole interaction strength in the polymer crystallite itself and is independent of the depolarization field. Therefore, the results reveal a new driving force that is different from the depolarization field for the ferroelectric size effect. The model is based on the specific structural characteristics of the copolymer, such as the van der Waals interchain interaction, the preferential parallel hexagonal side orientation in relation to substrate, and the partial crystallinity (crystallites in amorphous matrix with randomly oriented chains). Therefore, the molecular orientation driven size effect is unique for P(VDFTrFE) films. The crystallites of LB films are formed basically in a layer by layer manner, which leads to different microstructural characteristics such as different crystallite size and crystallinity from those of spin coated films. Thus, the results may not be applicable to the LB films. It should be noted that the T in eq 4 implies a statistical physics (the well ordered equilibrium state is achieved by annealing) of the dipole orientation effect, although it is hard to evaluate the effect of T due to its possible influence on dipole interaction energy and lattice structure. Equation 4 reveals the critical role of lateral crystallite size W: for spin coated films with a thickness thinner than W the ferroelectric state is unstable. An implication is that in order to improve the ferroelectric performance of the films, it is necessary to modify the dipole orientation or the crystalline structure. This may be achieved by using a different deposition technique (e.g., the LB technique)8 or a source material with a different VDF/TrFE ratio,6 etc.

Figure 7. AFM images (5 μm × 5 μm) of different films. (a−c) 40 nm thick films on SiO2 that have been annealed at a temperature of 135, 140, and 145 °C, respectively. (d−f) 60 nm thick films on Al that have been annealed at a temperature of 135, 140, and 145 °C, respectively.

annealed at various temperatures on different substrates. Both films show an abrupt granular to needle-like large grain morphology change around 140 °C, demonstrating that the lateral grain size was dependent on the annealing conditions but independent of thickness and substrate. Therefore, it is reasonable to assume a constant W for the 135 °C annealed. Then eq 4 can give good fits of the Pr dependence of thickness for films on Pt (see Figure 7). Using a monolayer thickness of 0.43 nm, we obtained the best-fit W of 115 and 68,

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CONCLUSIONS In summary, by using various structural analysis techniques we found a remarkable confinement-induced cooperative molecular dipole lying down orientation change in spin coated 1888

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P(VDF-TrFE) films over a wide thickness range (30 to 250 nm). The thickness dependence of molecular dipole orientation degree shows a good correlation with that of the ferroelectric properties measured from highly reproducible P−E loops of the films on inert electrode with suppressed “depolarization effect”. The results reveal a unique molecular orientation origin of size effect in semicrystalline ferroelectric copolymer films, which is distinctly different from the well recognized physical type driving forces of size effect. The orientation mechanism was interpreted by a simplified microscopic model based on the molecular dipole interaction energy without considering the switching dynamics. The results also imply that the ferroelectric state is unstable when the film thickness is thinner than the lateral size of the crystallites. The findings shed some new light in understanding of the nature of ferroelectricity in polymers relevant to the design of emerging polymer based electronic devices.

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(14) Buck, E.; Fuhrmann, J. Macromolecules 2001, 34, 2172. (15) Jones, R. L.; Kumar, S. K.; Ho, D. L.; Briber, R. M.; Russell, T. P. Nature (London) 1999, 400, 146. (16) Park, Y. J.; Kang, S. J.; Park, C.; Lotz, B.; Thierry, A.; Kim, K. J.; Huh, J. Macromolecules 2008, 41, 109. (17) Naber, R. C. G.; Blom, P. W. M. J. Phys. D: Appl. Phys. 2006, 39, 1984. (18) Lee, K.; Zaitsu, S.; Ishibe, K.; Sekitani, T.; Someya, T. J. Appl. Phys. 2010, 107, 114506. (19) Guo, D.; Stolichnov, I.; Setter, N. J. Phys. Chem. B 2011, 115, 13455. (20) Mabuchi, Y.; Nakajima, T.; Furukawa, T.; Okamura, S. Appl. Phys. Expr. 2011, 4, 071501. (21) Tagantsev, A. K.; Landivar, M.; Colla, E.; Setter, N. J. Appl. Phys. 1995, 78, 2623. (22) Tsutsumi, N.; Ueyasu, A.; Sakai, W.; Chiang, C. K. Thin Solid Films. 2005, 483, 340. (23) Amir, H.; Al-Bayati, K. G.; Orrman-Rossiter, J. A.; Berg, V. D.; Armour, D. G. Surf. Sci. 1991, 241, 91. (24) Saif, M. T. A.; Zhang, S.; Haque, A.; Hsia, K. J. Acta Mater. 2002, 50, 2779. (25) Rabe, K. M. Nature Nanotechnol. 2006, 1, 171. (26) Sai, N.; Kolpak, A. M.; Rappe, A. M. Phys. Rev. B 2005, 72, 020101. (27) Bellet-Amalric, E.; Legrand, J. F. Eur. Phys. J. B 1998, 3, 225. (28) Blaudez, D.; Buffeteau, T.; Cornut, J. C.; Desbat, B.; Escafre, N.; Pezolet, M.; Turlet, J. M. Appl. Spectrosc. 1993, 47, 869. (29) Petzelt, J.; Legrand, J. F.; Pacesova, S.; Kamba, S.; Kozlov, G. V.; Volkov, A. A. Phase Transitions 1988, 12, 305. (30) Kuhn, M.; Kliem, H. Ferroelectrics 2008, 370, 207. (31) Duan, C. G.; Mei, W. N.; Yin, W. G.; Liu, J. J.; Hardy, J. R.; Ducharme, S.; Dowben, P. A. Phys. Rev. B 2004, 69, 235106.

ASSOCIATED CONTENT

S Supporting Information *

Schematic illustration of the 2D-XRD and GIXD diffraction geometry. This material is available free of charge via the Internet at http://pubs.acs.org.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected], [email protected]. cn. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by the Competence Centre for Materials Science and Technology (CCMX) of Switzerland, the “Bairen Programme” of Chinese Academy of Sciences and the National Science Foundation of China (No. 11074277). The authors also thank Dr. Terrettaz Samuel and Mr. Condemi Enrico for their help in IRRAS measurements.

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REFERENCES

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dx.doi.org/10.1021/ma302377q | Macromolecules 2013, 46, 1883−1889