Unique Crystal Orientation of Poly(ethylene oxide) Thin Films by

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Unique Crystal Orientation of Poly(ethylene oxide) Thin Films by Crystallization Using a Thermal Gradient Gabin Gbabode,*,† Maxime Delvaux,‡ Guillaume Schweicher,‡ Jens W. Andreasen,§ Martin M. Nielsen,∥ and Yves H. Geerts‡ †

Normandie univ, Université de Rouen Normandie, Laboratoire Sciences et Méthodes Séparatives, Place Emile Blondel, Mont-Saint-Aignan 76821 Cedex, France ‡ Laboratoire de Chimie des Polymères, Faculté des Sciences, Université libre de Bruxelles (ULB), Boulevard du Triomphe, 1050 Brussels, Belgium § Department of Energy Conversion and Storage, Technical University of Denmark P.O. Box 49, 4000 Roskilde, Denmark ∥ Department of Physics, Technical University of Denmark, DK-2800 Kgs. Lyngby, Denmark S Supporting Information *

ABSTRACT: Poly(ethylene oxide) (PEO) thin films of different thicknesses (220, 450, and 1500 nm) and molecular masses (4000, 8000, and 20 000 g/mol) have been fabricated by spin-coating of methanol solutions onto glass substrates. All these samples have been recrystallized from the melt using a directional thermal gradient technique. Millimeter-size domains with crystallites uniformly oriented in the direction of the thermal gradient are observed. Futhermore, the crystallites size and orientation distribution are enhanced (e.g., increases and decreases, respectively) when film thickness is decreased, ultimately leading to a single-crystal-like behavior for 220 nm thick PEO films of mass 8000 g/mol. Interestingly, this fine microstructure is partially retained after melting and subsequent cooling back to ambient temperature for the highest molecular weight polymer allowing, in this particular case, to significantly decrease the distribution of crystal orientation obtained after crystallization using the thermal gradient technique. Conventionally, polymer thin films are produced by solventbased deposition techniques (less expensive and more environment-friendly) that typically lead to the formation of small polycrystalline domains with random in-plane orientation. Thus, methods enabling a precise control of polymer crystallization in thin films to obtain uniformly oriented crystalline domains on large scales (hundreds of micrometers typically) have to be implemented. Most employed strategies are based on the influence of the underlying solid substrate on the crystallization of the deposited polymer thin film19 through epitaxial relationship20−24 or by using prealigned polymeric substrates25−29 or nanoimprint lithography.30−32 Those methods certainly proved to be efficient in tailoring the in-plane orientation of polymer crystallites on large scales, but they unavoidably necessitate a substrate preparation step that can introduce additional defects such as interfacial roughness that can be detrimental to device quality in particular for applications in electronics or optoelectronics.33,34 Mechanical rubbing of the polymer film35,36 is an alternative method which does not require the presence of an alignment substrate but which can suffer from similar roughness defects due to the

1. INTRODUCTION A comprehensive knowledge of the crystallization behavior of semicrystalline polymers is a crucial step for a rational design of products or devices based on these materials. Indeed, the physical properties of semicrystalline polymers greatly depend on the way they arrange in the solid state and on crystallization kinetics.1−3 In particular, owing to their hierarchical structural arrangement that ranges from the atomic scale (monomers arranged in a crystallographic lattice) to the macroscopic scale (large networks such as spherulites consisting of crystalline lamellae, i.e., crystallites, and amorphous parts), highly oriented polymeric materials can be produced with enhanced physical properties. These are the basic units of several manufactured plastic products of everyday life4−8 and natural fibers.1 Since the past decades, a wide range of novel advanced technologies based on polymer thin films such as coatings,9 electronics,10−13 and optoelectronics14−16 have attracted much interest from researchers and industrials. Generally in those applications, device performance is critically dependent on the structure17 (packing arrangement of crystalline moieties) and/ or microstructure18 (size and orientation of crystalline and amorphous parts) of the thin film polymeric materials. A profound knowledge of the crystallization of polymers in thin films deposited onto a solid substrate is then needed to rationalize and further optimize materials properties.19 © XXXX American Chemical Society

Received: February 28, 2017 Revised: May 26, 2017

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DOI: 10.1021/acs.macromol.7b00441 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules mechanical interaction between the rubbing tool and the polymer material.35 A different approach which does not bring any additional roughness issues consists in inducing high degrees of orientation of the polymer crystallites by recrystallization of the polymer film using a careful control of solvent vapor stream33,37,38 or temperature.39 We propose here a crystallization technique using a thermal gradient, derived from the well-known Bridgman method,40,41 that we previously implemented on a model π-conjugated small molecule with quite good success as single-crystal-like films were produced as proven by X-ray pole figure experiments.42 Crystallization of polymer ingots using a thermal gradient was already studied in the mid-1970s by Lovinger et al.43−48 These authors could prove that this method allowed the growth of polymer chains from a unique nucleus and then achieving highly oriented crystallites along the direction of the ingot displacement. More recent studies on this crystallization method performed on polymers include works on isotactic poly(butane-1), poly(vinylidene fluoride), and polyethylene by Toda et al. with the aim of a better fundamental understanding of spherulitic growth from melt crystallization.49 In a similar way, use of a directional thermal gradient in the so-called “zone annealing technique” has been proved to enhance both orientation and size of ordered mesoscopic networks formed by block copolymers.50−52 In the latter works, however, amorphous block copolymers were investigated; therefore, the ordering effect is based upon polymer relaxation occurring above the glass transition as in a conventional annealing procedure. In the present study, we are particularly interested in tailoring the microstructure of semicrystalline polymer thin films to obtain single-crystal-like orientation of crystallites, which might lead to improved anisotropic physical properties.19,33 The archetype crystalline polymer, poly(ethylene oxide) (of formula (C2H4O)n, hereafter abbreviated PEO), was then chosen as a model polymer since its crystallization behavior from the melt has been carefully characterized in powder samples53−58 as well as in single crystals.59−63 Typically, spin-coated PEO thin films are recrystallized using a thermal gradient method whereby the sample is displaced at a constant speed from a “hot” to a “cold” zone (see the Experimental Section for a detailed description of our setup). PEO is a well-suited candidate for this study because its radial crystal growth rates are comparable to the sample displacement velocities enabled by our thermal gradient setup (at least at moderate degrees of undercooling such as those imposed in our experiments; see Table 2). According to our previous studies, this criterion appeared to be important for a good morphological control of the sample.42,64 A thorough investigation of the influence of thermal gradient parameters (gradient magnitude, sample velocity) and sample parameters (polymer mass, film thickness) on the size and orientation of crystalline domains is carried out. This ultimately leads to the production of films of single-crystal character with precise orientation of the crystallographic unit cell with respect to the thermal gradient direction.

Table 1. Characteristics of the Three PEO Mass Fractions PEO 4K number-average molecular weight, Mna (g/mol) polydispersitya mean degree of polymerization, p average length of fully extended chains in the crystal lattice, lcb (nm) melting temperature, Tmc (°C) enthalpy of melting, ΔHmc (kJ/mol) degree of crystallinityd (%)

PEO 8K

PEO 20K

3640

7670

19990

1.07 82

1.10 174

1.12 454

22.8

48.4

126.3

58.9 ± 0.1 5.3 ± 0.1

60.1 ± 0.1 6.8 ± 0.1

61.2 ± 0.1 7.7 ± 0.1

86

87

88

a

Mn and polydispersity have been determined by gel permeation chromatography. blc = p × lu with lu the length of a PEO monomer inferred from the crystal structure of PEO.65 cPhase transition temperatures and enthalpies were determined from differential scanning calorimetry measurements (PerkinElmer Diamond DSC) performed at a heating/cooling rate of 5 K/min. dThe degree of crystallinity has been estimated from the ratio between the measured enthalpy of melting and the reference enthalpy of melting of a hypothetical 100% crystallized polymer.66 2.2. Thin Film Fabrication. Methanol solutions of PEO with different concentrations (10, 20, and 50 mg/mL) were deposited onto 20 × 20 × 0.16 mm3 thin glass substrates (Marienfeld cover glasses Cat. No. 0101040) by spin-coating with the conditions: 300 rpm for 5 s, then 1500 rpm for 5 s, and a last spinning step of 3000 rpm for 30 s. Glass substrates were cleaned prior to film deposition by plunging them into a “piranha” solution (one-third hydrogen peroxide/twothirds sulfuric acid 18 M volume ratio) for 1 min, then rinsed with distilled water, and dried by spin-coating (2000 rpm for 10 s, then 4000 rpm for 15 s, and then 6000 rpm for 15 s). Film thicknesses of 220 ± 30, 450 ± 50, and 1500 ± 200 nm were obtained following this procedure as measured by AFM (Digital Instruments Nanoscope III microscope in tapping mode) on films previously scratched with a razor blade. No significant film thickness variation was measured for the different mass fractions at each given solution concentration. 2.3. Thermal Gradient Setup. A detailed description of the apparatus has previously been presented elsewhere.42 To summarize, it consists of a Linkam GS350 temperature gradient system heating stage made of two independent heating devices separated by a distance of 2.5 mm where the thermal gradient is installed (Figure 1). One is set at a temperature Th above the melting temperature (hot side) and the other at a temperature Tc below the crystallization temperature (cold side) of PEO. The whole system is enclosed in a hermetic metallic container so that the sample is thermally insulated from the laboratory environment. Thin films of PEO are initially placed entirely at the hot side (that is, in the molten state) and are slowly translated to the cold side at a constant speed v (down to 1 μm/s) until the whole sample is at the cold side, then allowing directional crystallization of PEO crystallites. Several combinations of Th, Tc, and v have been tested in order to probe the influence of the gradient magnitude as well as the sample velocity on the microstructure of PEO thin films recrystallized using the thermal gradient technique. They are displayed in Table 2. 2.4. Polarized Optical Microscopy (POM). The thermal gradient apparatus is mounted on a polarized optical microscope (Nikon Eclipse 80i) so that images can be taken before, during, and after thermal gradient to obtain qualitative information on the crystallization behavior and also to evaluate the size of uniformly oriented crystalline domains. 2.5. X-ray Diffraction Measurements. Different experimental geometries were used not only to obtain structural information on PEO thin films crystallized using the thermal gradient technique but also to quantitatively determine the orientation of PEO crystallites with respect to the thermal gradient direction (Figure 2). The latter measurements bring invaluable insight into the crystallization behavior

2. EXPERIMENTAL SECTION 2.1. Materials. PEO fractions with three distinct number-average molecular weights Mn, namely 4000, 8000, and 20 000 g/mol (hereafter PEO 4K, PEO 8K, and PEO 20K, respectively), have been purchased from Sigma-Aldrich and used without further purification. The characteristics of the three samples are summarized in Table 1. B


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Figure 1. Schematic representation of the thermal gradient setup emphasizing (a) the sample and (b) the two heating stages above which the sample is translated. In (c) are sketched the anticipated shape of the crystalline domains obtained by this directional crystallization method starting each from a unique nucleation spot (in red) (see Figure 2 and Figure S2 for direct visualization of these domains by microscopy). A zoom is made on the crystalline domains which shows the face-on arrangement of PEO lamellae as will be demonstrated further in the text.

Table 2. Summary of the Different Thermal Gradient Parameters Applied in the Present Studya entry

Th (°C)

Tc (°C)

sample velocity v (μm/s)

gradient magnitude G (°C/mm)

cooling rate C (°C/min)

1 2 3 4 5 6 7

70 70 70 60 70 70 70

20 30 40 40 40 40 40

5 5 5 5 2 10 20

20 16 12 8 12 12 12

6.00 4.80 3.60 2.40 1.44 7.20 14.4

Gradient magnitude G and cooling rate C are calculated as G = (Th − Tc)/x and C = [(Th − Tc) × V]/x, with x = 2.5 mm, the gap between the hot and cold sides. a

of PEO thin films submitted to the thermal gradient treatment. To do so, the direction of sample translation during the crystallization process (hereafter called the “thermal gradient direction” which is opposite to the crystal growth direction as shown in Figure 1) was indicated by an arrow on each thin film sample as a reference, prior to X-ray diffraction experiments. Importantly, all X-ray diffraction measurements which are referred to in the text and Supporting Information have been performed at room temperature (i.e., 25 °C) unless otherwise stated. 2.5.1. Specular X-ray Diffraction (sXRD). The sXRD experiments were performed on a Bruker D8 Advance diffractometer in θ/θ reflection geometry using Cu Kα radiation (wavelength λ = 0.154 18 nm) and equipped with a Material Research Instruments heating stage. In this configuration, the source and the detector continuously move from the plane of the sample holder, which is fixed, at the same angle θ. Typical specular scans were measured in the 2θ range [1.6°−40°] (where 2θ is the angular deviation between the incident and diffracted rays) with an angular resolution of 0.02° and an integration time of 10 s per step. These measurements provide information on families of reticular planes oriented parallel to the thin film surface (out-of-plane arrangement).

Figure 2. Schematic representation of the (a) GIWAXS and (b, c) inplane X-ray measurements setup. In (a), αi is the incident angle (kept at 0.18°) and αf is the diffracted angle which can be separated into its in-plane (αfi) and out-of-plane (αfo) components. As mentioned in the text, the position of a reflection is rather expressed as the in-plane (qxy) and out-of-plane (qz) components of the scattering vector q. The two possible orientations of the sample during GIWAXS measurements are also highlighted by blue dashed arrows (thermal gradient direction either parallel (||) or perpendicular (⊥) to the specular reflection plane). In (b) and (c), S represents the X-ray source, D the point detector, αi the incident angle (kept at 0.2°), αf the out-of-plane angle of the detector (kept at 0.2°), and ϕ, ϕ′, and 2θχ are described in the text. In (c) is depicted the mathematical change of origin (ϕ ⇒ ϕ′) which allows the direct quantification of angle η between [hkl]* and the thermal gradient direction. 2.5.2. Grazing Incidence Wide-Angle X-ray Scattering (GIWAXS). GIWAXS experiments were performed at the X-ray facility at DTU C


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Macromolecules Energy, Denmark.67 Here, Cu Kα,β radiation is generated by a copper rotating anode (Rigaku) operating at 10 kW and is focused, collimated, and filtered (only Cu Kα is selected, λ = 0.154 18 nm) by a multilayer X-ray mirror. The hence produced X-ray beam irradiates the thin film sample at an incident angle of 0.18°, which is less than the critical angle for total reflection of the underlying glass substrate, thereby suppressing the scattering from the latter and thus effectively enhancing the signal from PEO thin films. Diffracted X-rays were collected on X-ray photostimulable image plates upon exposure time of 8 h (one-day measurement) or 16 h (overnight measurement). Two-dimensional GIWAXS images were then obtained, representing the diffracted intensity (logarithmic scale) in false color scale as a function of the scattering vector components qz (vertical axis) and qxy (horizontal axis) respectively perpendicular and parallel to the sample surface (q-range up to about 30 nm−1). Additionally, measurements with the thin film sample oriented so that the thermal gradient direction is either parallel (||) or perpendicular (⊥) to the specular reflection plane (Figure 2a) were both performed. 2.5.3. In-Plane Grazing Incidence X-ray Measurements. A Rigaku Ultima IV diffractometer equipped with a DHS 1100 heating stage from Anton Paar was used. As for the GIWAXS setup, the incoming parallel X-ray beam (Cu Kα radiation) was set at a grazing incidence angle (0.2°). The main difference is that the detector (a conventional point detector) is allowed to move along a circle parallel to the sample surface (see Figure 2b), thus enabling the diffraction of crystallographic planes whose normal is parallel to the sample surface (in-plane reflections) to be detected. Two Soller slits sets were intercalated in the path of the X-ray beam before and after the sample to ensure its efficient collimation (the lateral divergence of the incident and diffracted beam is by these means reduced to 0.5°). Two types of measurements were possible (Figure 2b): (a) in-plane grazing incidence X-ray diffraction (inGIXD) measurements or 2θχ/ϕ scans (the sample continuously rotates around its normal at an angle ϕ while the detector moves around the same axis at an angle 2θχ from the direct beam, so that ϕ = 2θχ/2 at all times) which reveals the diffraction of reticular planes which are oriented so that their normal is perpendicular to the initial position of the sample holder (ϕ = 0°); (b) in-plane grazing incidence (partial) pole figures (inGIPFs) or ϕ scans (the sample continuously rotates around its normal at an angle ϕ while the detector is fixed at an angle 2θχ corresponding to the conditions of observation of a specific hkl in-plane reflection) which probes the orientation of the normal to a particular family of (hkl) reticular planes (that is, [hkl]* direction) in the plane of the film. The in-plane pole figure is designated as “partial” as the rotation of the sample holder do not span an entire circle (−175° < ϕ < 90°, see Figure 2b). For both types of experiments, PEO thin films were positioned at the beginning of each measurement with the thermal gradient direction approximately coinciding with the ϕ = 0° position of the sample holder (with an uncertainty of a few degrees as this operation was made manually and controlled by eye) in order to relate directly the [hkl]* direction with that of the thermal gradient. In the case of inGIPF measurements, when diffraction of a (hkl) plane is measured at a specific ϕ0 azimuthal direction, the angle η between [hkl]* and the thermal gradient direction is directly given by the relation η = ϕ0 − (π − 2θχ)/2, where θχ is the angle of diffraction for this particular family of reticular planes (hkl) as shown in Figure 2c. Then, inGIPFs plotted as a function of ϕ′ = ϕ − (π − 2θχ)/2, as will be the case in the present paper, directly give the distribution of orientation of [hkl]* direction with respect to the thermal gradient direction and therefore allow a straightforward quantitative determination of the preferred in-plane orientation of PEO crystallites.

Figure 3. Polarized optical microscopy images obtained at room temperature for (a) an as-cast PEO 8K 1500 nm thick film and (b) the same sample during the thermal gradient treatment (conditions: Th = 60 °C, Tc = 40 °C, and v = 5 μm/s). In (a) a nucleation site is visible at the center of a spherulite. In (b) the thermal gradient direction is indicated as well as the position of the vertical growth front (PEO crystallizes starting from this plane). Scale bars represent 200 μm.

that they could even be spotted by eye on the surface of the film. This large size of spherulites indicates relatively low nucleation density, which is in contrast to the very fast deposition process of spin-coating. This suggests that the latter deposition method has reduced influence on the crystallization behavior of PEO with the deposition conditions used. Indeed, for thinner films (450 and 220 nm thick), the size of spherulites increases further (see Figure S1 of the Supporting Information), reflecting the lower amount of PEO which consequently leads to a decrease of the nucleation density. Importantly, the presence of spherulites emphasizes a random orientation of crystallites for as-cast PEO thin films. In contrast to the previous observations, after thermal gradient treatment (conditions: Th = 60 °C, Tc = 40 °C, and v = 5 μm/s), large uniform domains oriented parallel to the thermal gradient direction are clearly visible in POM images (Figure 3b, part at the right of the vertical dashed line). These domains, a few hundreds of micrometers wide and some millimeters long, seem to originate from a unique nucleation site (see Figure S2), suggesting a unique crystal orientation. In Figure 3b also, the growth front, at which crystallization of PEO starts, is well visible and is perfectly vertical (indicated by dotted lines). 3.1.2. Investigation of the Out-of-Plane Crystal Arrangement of PEO Thin Films. To confirm this assumption, sXRD measurements were first performed on these oriented thin films. First, a representative sXRD pattern of a bulk sample of PEO 35K68 is shown in Figure 4c for comparison. Sharp intense peaks are clearly distinguished in the whole 2θ range, confirming the high crystallinity of PEO fractions used in this study. The hkl indexations indicated for each reflection are in agreement with the monoclinic structure first solved by Takahashi and Tadokoro.65 As far as PEO thin films after thermal gradient treatment are concerned, no such peaks are observed at wide angle and are replaced by a broad diffuse halo originating from the glass substrate as shown for PEO 8K 1500 nm thick films in Figure 4a. This is in fact inherent to the specular reflection geometry since for incidence angle significantly larger than the critical angle for total reflection, the penetration depth of the incident X-ray beam in the thin film is too high (much higher than the film thickness), resulting in the masking of the signal from the thin film by that from the substrate. However, precious information is collected at low diffraction angle (Figure 4b). Indeed, several reflections are observed which are not present for as-cast PEO thin films (see Figure S3). These all are correlated reflections (from the third

3. RESULTS 3.1. Investigation of the Crystal Structure and Microstructure of PEO Thin Films Crystallized Using the Thermal Gradient Technique. 3.1.1. Polarized Optical Microscopy Observations. As-cast PEO thin films present a typical spherulitic microstructure as observed by POM (Figure 3a). The spherulites are large with diameter exceeding 1 mm so D


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Figure 4. (a, b) sXRD patterns measured at room temperature for PEO 8K 1500 nm thick films crystallized using the thermal gradient technique and represented in the (a) [1.2°−41°] 2θ range (gradient conditions: Th = 60 °C, Tc = 40 °C, and v = 2 μm/s) and (b) [1.2°−5°] 2θ range (gradient conditions: Th = 70 °C, Tc = 40 °C, and v = 5 μm/s). In (b) reticular distances corresponding to the 2θ positions of the observed reflections are indicated and are fairly equal to the first-order reticular distance (e.g., the lamellar thickness) divided by the corresponding order of diffraction (third to eight order from left to right). Inset schematically represents two stacked PEO face-on lamellae with the lamellar thickness extracted from the positions of the Bragg peaks. (c) sXRD pattern measured at room temperature for a bulk sample of PEO 35K. Reflections are indexed in agreement with the monoclinic structure determined by Takahashi and Tadokoro.65 Indexations highlighted in red with their 2θ positions also indicated point to hk0 reflections which are useful for the determination of the in-plane preferred orientation of PEO crystallites (see Results section 3).

to the eighth order) corresponding to a repeat distance of 19.5 nm for PEO 8K thin films crystallized using the thermal gradient technique (Figure 4b). This result indicates that PEO chains stand on top of the glass substrate and pack into thick lamellae (so around 20 nm thick for PEO 8K) which superimpose periodically in a direction perpendicular to the substrate surface (face-on arrangement, see inset of Figure 4b). Considering the degree of crystallinity of PEO 8K fractions, it follows that the lamellar thickness achieved corresponds to twice-folded chains.58 3.1.3. Investigation of the In-Plane Crystal Arrangement of PEO Thin Films. sXRD measurements give only information on the crystal packing in a direction perpendicular to the substrate surface. Hence, GIWAXS measurements were performed on PEO thin films before and after thermal gradient treatment (gradient conditions: Th = 70 °C, Tc = 40 °C, and v = 5 μm/s) as they provide diffraction data in (qxy, qz) reciprocal space. GIWAXS images collected for a PEO 8K 1500 nm thick film are shown in Figure 5. Before thermal gradient treatment (Figure 5a), a symmetrical 2D pattern (toward the central meridian line as shown in Figure 5a) composed of several intense ring-like reflections is obtained. This indicates a significant angular distribution of the corresponding diffraction planes, i.e., a powder-like arrangement of crystallites, in good agreement with the spherulitic morphology observed by POM for as-cast thin films (see Figure 3a). The observed reflections could be readily indexed using the monoclinic structure of PEO, demonstrating that the same crystal structure is present in thin films and bulk samples. Furthermore, exploitation of diffracted intensities using Sim Diffraction software69 (Figure 5c) showed that PEO chains preferentially lie down parallel to

Figure 5. Two-dimensional GIWAXS images, showing reflections in the (qxy, qz) reciprocal plane, measured at room temperature for a PEO 8K 1500 nm thick film (a,c) before and (b, d) after thermal gradient treatment (gradient conditions: Th = 70 °C, Tc = 40 °C, and v = 5 μm/ s). In (c) are represented calculated intensities using Sim Diffraction software (see text) to be compared to the experimental ones in (a). In particular, indexations of the most intense reflections (calculated from the lattice constants of the monoclinic structure of Takahashi and Tadokoro65) are indicated for (a) and (c). (b, d) GIWAXS images were obtained with the sample placed with the thermal gradient direction oriented either (b) parallel (“||”) or (d) perpendicular (“⊥”) to the X-ray reflection plane (see Experimental Section) as indicated in the upper right corner of each image.

E


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Figure 6. (a, c) sXRD patterns measured at room temperature for PEO 8K 1500 nm thick films crystallized using the thermal gradient technique with different values of (a) the gradient magnitude (v constant = 5 μm/s) and (c) the sample displacement speed (G constant = 12 °C/mm). Thermal gradient magnitude as well as sample displacement speed increases from the top curve to the bottom one, and integers 4, 5, and 6 indicate the order of diffraction of the corresponding reflections. In particular, reflections whose position is not correlated with that of the others are pointed by black arrows. (b, d) Correlation length determined from the first diffraction peak observed at around 1.8° in 2θ in (a) and (c) plotted as a function of (b) the gradient magnitude and (d) the sample displacement speed.

Table 2). Each sample has been analyzed by sXRD, and the results are given in Figure 6a. First of all, whatever the thermal gradient magnitude used, correlated peaks at the same positions are observed at low 2θ angle corresponding to a lamellar thickness of around 20 nm as shown in first section. This indicates that the lamellar thickness does not change with the thermal gradient at least in the range of values accessible in this study. The differences between the sXRD patterns rather lie in the number of peaks observed and their width, i.e., in the outof-plane order of the lamellar packing. Figure 6b shows the evolution of the correlation length (determined from the width of the first diffraction peak observed at 1.8° in 2θ using the Scherrer formula71) with the gradient magnitude G. This quantity gives an evaluation of crystallite size in the direction perpendicular to the lamellae and was taken as a means to compare the effectiveness of directional crystallization for samples submitted to different thermal gradient conditions. It appears that it slightly increases up to a value of around 100 nm (equivalent to 5 lamellae) for G = 12 °C/mm and then dramatically decreases for higher values of G. Then G = 12 °C/ mm (Th = 70 °C and Tc = 40 °C) was subsequently chosen as the optimal gradient magnitude for most efficient directional crystallization. It has to be mentioned that for G = 8 °C/mm (Th = 60 °C and Tc = 40 °C) a large crystallite size is achieved close to 90 nm (equivalent to 4−5 lamellae). However, another peak appeared in the sXRD pattern whose position is not correlated with that of the others (pointed by a black arrow in Figure 6a) and which likely corresponds to another lamellar thickness and therefore to the appearance of other nuclei during the crystallization process. This situation is obviously not suitable for our purpose to obtain films with single-crystallike orientation, and these thermal gradient conditions were then discarded.

the substrate surface (edge-on orientation) which explains the presence of localized intensity maxima along the ring-like reflections observed in Figure 5a. The presence of this preferential lamellar orientation is in agreement with previous studies on PEO which shows that it is favored for thin films with thickness larger than 300 nm.70 After thermal gradient treatment (Figures 5b,d) the 2D images are now completely different and display several localized thin reflections with reduced angular distribution. Moreover, the symmetry toward the central meridian line put forward for as-cast films no longer exists for thin films after thermal gradient treatment. These results reveal that within the thick ordered lamellae put forward by sXRD measurements preferred orientation of crystalline domains is induced by the thermal gradient treatment and maintained further after it is completed. Noticeably, reflections observed with the sample in the “⊥” direction are much broader (in the radial direction) than those observed with the sample in the “||” direction (see also Figure S4). This suggests crystallites with high aspect ratio, i.e., extended along the thermal gradient direction but thinner perpendicular to it, in good agreement with the shape of uniform domains revealed for a 1500 nm thick PEO 8K sample by POM (Figure 3b). As for as-cast films, all reflections could be indexed using the known monoclinic crystal structure (Figure S4), thus showing that only the crystal orientation of PEO is modified by the thermal gradient treatment, not its crystal structure. 3.2. Influence of Thermal Gradient and Sample Parameters on the Observed Crystal Structure and Microstructure. 3.2.1. Influence of Thermal Gradient Parameters. Four different gradient magnitudes ranging from 8 to 20 °C/mm were tested on PEO 8K 1500 nm thick films with a constant sample velocity of 5 μm/s (conditions 1−4 in F


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Figure 7. (a) sXRD patterns measured at room temperature for 450 nm thick films of PEO 4K, 8K, and 20 K (from bottom to top) crystallized using the thermal gradient technique. For PEO 4K and 8K, integers indicate the order of each correlated reflection and the corresponding lamellar thickness is also given (same color as the sXRD pattern). (b, c) GIWAXS images measured at room temperature for a PEO 20K 450 nm thick film crystallized using the thermal gradient technique with the sample placed in the (b) parallel (“||”) or (c) perpendicular (“⊥”) orientation as indicated in the upper right corner of each image.

Figure 8. (a, b) GIWAXS images measured at room temperature for a PEO 8K 220 nm thick film crystallized using the thermal gradient technique (conditions: Th = 70 °C, Tc = 40 °C, and v = 5 μm/s) and (c, d) corresponding integrated 1D XRD patterns. For (a, c) the sample was positioned in the parallel (“||”) orientation while for (b, d) it was in the perpendicular (“⊥”) one.

A similar trend was pointed out for the evolution of crystallite size as a function of sample displacement velocities (ranging from 2 to 20 μm/s with a constant thermal gradient of 12 °C/mm = conditions 3, 5, 6, and 7 in Table 2) with a maximum reached for v = 5 μm/s followed by a marked decrease for higher velocities (Figure 6d). Similarly, for too low values of v, nucleation rate seems to be higher although crystallite size is high (Figure 6c). These findings can be explained by the fact that at low cooling rate, as it occurs for

low values of either thermal gradient magnitude or sample displacement speed (see Table 2), decoupling of nucleation and growth might no longer be efficient,43 then resulting in crystal growth from different nuclei which might ultimately lead to crystallites exhibiting different lamellar thickness. Finally, the following thermal gradient conditions Th = 70 °C, Tc = 40 °C, and v = 5 μm/s were chosen as optimal parameters for directional crystallization of PEO thin films and will be kept in G


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Figure 9. (a) inGIXD patterns and (b) inGIPFs (2θχ fixed at 19.2° corresponding to the diffraction of (120) planes) measured at room temperature for PEO 4K, 8K, and 20K [from bottom to top in (a) and from top to bottom in (b)] 220 nm thick films crystallized using the thermal gradient technique (conditions: Th = 70 °C, Tc = 40 °C, and v = 5 μm/s). Indexations of observed reflections are given in (a). In (b) ϕ′ positions of observed major maxima are indicated for each curve. (c) Schematic representation of the two possible preferred orientations of the unit cell toward the thermal gradient direction (viewed in the plane of the film) inferred from the results given in (a) and (b).

The effect of film thickness is hardly visible by sXRD as peak intensities monotonically decrease with decreasing film thickness until vanishing almost completely for 220 nm thick films (Figure S5). This only reflects the fact that the number of coherent X-ray scatterers in the direction perpendicular to the glass substrate diminishes with lower film thickness until being not sufficient to produce a detectable signal for 220 nm thick films, at least with a conventional laboratory X-ray tube as is used in this study. However, peak positions remain the same at all film thickness showing that the same lamellar packing is preserved. By contrast, a marked change is put forward in the in-plane microstructure. Indeed, GIWAXS images of 220 nm thick PEO 8K films after thermal gradient treatment (Figures 8a,b) display spot-like reflections with reduced angular distribution, suggesting enhanced in-plane crystal orientation compared to thicker films (compare with Figures 5b,d). Moreover, reflections are quite thin in the radial direction independently of the sample orientation (“||” or “⊥”), indicating the presence of large uniformly oriented domains extending in directions both parallel and perpendicular to the thermal gradient direction. This result is all the more significant since the entire length of the sample is analyzed during GIWAXS measurements (beam footprint on the sample is much larger than the sample at +100 mm, with the setup used). The overall small size of reflections in the GIWAXS patterns could allow the integration of intensities as a function of q to yield conventional 1D X-ray diffraction patterns representing diffracted intensities as a function of 2θ angle (using the relation q = (4π/λ) sin θ with q = (qxy2 + qz2)1/2) as exposed in Figures 8c,d (indexations of the reflections of the GIXD patterns shown in Figures 8a,b can alternatively be found in Figure S6). Indexations of the observed reflections with the known monoclinic structure not only demonstrates that no change in the crystal structure occurs with the reduced film

the following of the paper even if not explicitly mentioned in the text. 3.2.2. Influence of Sample Parameters. Polymer mass and film thickness were the two adjustable sample parameters that were taken into account. In Figure 7a are superimposed the sXRD patterns measured for 450 nm thick PEO 4K, 8K, and 20K films crystallized using the thermal gradient technique. In the case of PEO 4K, a similar lamellar structure as for PEO 8K is revealed as shown by the presence of several correlated peaks at low angle. However, the lamellar thickness obtained is 13.7 nm, that is, much shorter than for PEO 8K thin films. Actually, it corresponds to lamellae composed of once-folded PEO chains. For PEO 20K thin films, no correlated peaks are observed at low angle in the sXRD pattern even though directional crystallization did take place as demonstrated by the GIWAXS image obtained for the same sample and showing discrete reflections spanning the whole (qxy, qz) domain at similar positions than for PEO 8K (Figures 7b,c to be compared to Figures 5b,d). The absence of diffraction peaks in the sXRD pattern can be explained by the presence of a too large lamellar thickness to be detected. Indeed, assuming that even 4 times folded lamellae are crystallized after thermal gradient treatment for PEO 20K thin films, a lamellar thickness of around 25 nm is expected. In this case, only reflections of the fifth order and more could be detected using our experimental sXRD setup, which would then be of very low intensity, hardly distinguishable from the background. To summarize, polymer mass seems not to significantly alter the in-plane microstructure of the oriented films (for more quantitative confirmation see section 3.3) but rather impacts on the out-of-plane lamellar packing by modifying lamellar thickness, the latter likely increasing with polymer mass. H


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Figure 10. inGIPFs (2θχ fixed at 19.2° corresponding to the diffraction of (120) planes) measured at room temperature for PEO (a, d) 4K, (b, e) 8K, and (c, f) 20K 220 nm thick films crystallized using the thermal gradient technique (conditions: Th = 70 °C, Tc = 40 °C, and v = 5 μm/s) (a−c) in a time lag of one month and (d−f) before and after melting. ϕ′ positions of observed major maxima are indicated. In (a) and (f) positions corresponding to the two coexisting different crystal orientations put forward for PEO 4K after 1 month and PEO 20K after melting (see text) are highlighted in different colors (red and blue).

parallel to the film surface.72 This then restricts the detection to only hk0 reflections in the in-plane geometry, that is, according to the crystal structure of PEO, to 020, 110, 120, and 040 reflections in the 2θ range scanned, occurring at 13.6°, 15.1°, 19.2°, and 27.2° in 2θ, respectively (see Figure 4c). Figure 9a superimposes the inGIXD patterns measured for PEO 4K, 8K, and 20K thin films previously submitted to the thermal gradient technique. These in-plane X-ray diffraction measurements show the diffraction of crystallographic planes oriented with their normal perpendicular to the thermal gradient direction (see Experimental Section). In the case of PEO 20K, 020, 040, and 120 reflections are observed while for PEO 8K and 4K only one family of reflections is observed, e.g., 020 (accompanied by 040, its second-order reflection) and 120, respectively. Hence, for PEO 20K two in-plane orientations of the unit cell are revealed, namely [020]* and [120]* perpendicular to the thermal gradient direction while a unique unit cell orientation is put forward for PEO 8K ([020]* perpendicular to the thermal gradient direction) and 4K ([120]* perpendicular to the thermal gradient direction). These results certainly illustrate the high in-plane orientation conferred to PEO thin films after thermal gradient treatment but are not sufficient to determine

thickness but also suggests the presence of unique crystal orientation as almost no overlap of the observed reflections is pointed out between the X-ray diffraction patterns obtained for the two orthogonal sample orientations (compare Figures 8c and 8d). The influence of reducing film thickness on the microstructure of thin films crystallized using the thermal gradient technique can be rationalized as a concomitant decrease of nucleation rate due to the lower amount of material available, resulting in significantly enhanced in-plane order and crystal orientation. 3.3. Investigation of the In-Plane Orientation of Crystallites. 3.3.1. Determination of the Preferred Orientation of the Unit Cell toward the Thermal Gradient Direction. Based on the results given in sections 3.1 and 3.2, quantification of crystal orientation has been undertaken for 220 nm thick films submitted to the thermal gradient conditions: Th = 70 °C, Tc = 40 °C, and v = 5 μm/s by means of in-plane XRD measurements which allow the detection of in-plane reflections at different azimuthal angles (see Experimental Section). Since PEO chains pack into face-on oriented lamellae (see previous sections), that is, c-axis is perpendicular to the film surface, then a* and b lattice vectors of the monoclinic structure are oriented I


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after annealing, indicating a fine distribution of orientation in clear contrast with the broad continuums of peaks observed before the annealing procedure. Two major preferred orientations of [120]* toward the thermal gradient direction, e.g., making angles of 129° and 164° with it (plus or minus 180°) are inferred from the peak positions. These indicate two possible orientations of the unit cell; that is, a* vector is either parallel to the thermal gradient direction (same as the “before melting” case) or making an angle of around 30° with it which is actually consistent with [110]* crystallographic direction being parallel to the thermal gradient direction.

the preferred in-plane orientation of the unit cell as only one azimuthal position toward the thermal gradient direction is investigated. Then, inGIPF measurements centered on 120 in-plane reflection (as it is the most intense, see Figures 4c and 9a), e.g., with 2θχ fixed at 19.2°, have been carried out for the three PEO samples and are presented in Figure 9b. Three maxima are clearly visible for PEO 4K, 8K, and 20K at approximately the same positions, namely, an intense one at ϕ′ = −135° and two less intense secondary ones around 90° apart from the major peak, e.g., at ϕ′ = −225° and ϕ′ = −45°. We attribute the principal maxima to the diffraction of (120) (or (−1−20)) planes while the two others correspond to the diffraction of (−120) and (1−20) planes which all occur at the same and only 2θχ angle (2θχ = 19.2°). These results unambiguously prove the existence of the same preferred in-plane orientation of crystallites which is induced by the thermal gradient directional crystallization technique independently of the polymer mass (at least, in the range of molecular weight investigated, e.g., from 4000 to 20 000 g/mol). It is characterized by [120]* crystallographic direction, making an angle of either 135° or 45° with the thermal gradient direction. According to PEO crystal structure, this leads to a uniaxial orientation of the unit cell with a* vector either in the same or opposite direction to the thermal gradient direction as sketched in Figure 9c. Particularly noteworthy is that for PEO 8K the same unit cell orientation is inferred from the inGIXD pattern shown in Figure 9a which indicates that b lattice vector is perpendicular to the thermal gradient direction (as [020]* = b) which is the same as a* being parallel to the thermal gradient direction as the unit cell is monoclinic. 3.3.2. Investigation of the Stability of the Oriented Thin Films. The stability with time of the determined crystal orientation was investigated by checking the evolution of the inGIPFs after one month (Figures 10a−c). For PEO 8K and 20K, no significant change is noticed after 1 month storage at ambient conditions indicating that the same crystal orientation is preserved during this time. The constant peak shift observed (up to 3° for PEO 8K, for example) is only due to the different position of the sample at the beginning of the experiment as this procedure is operated manually and controlled by eye. However, for PEO 4K, drastically different peak positions are observed, showing the presence of two coexisting preferred inplane unit cell orientations. The first one, characterized by maxima at ϕ′ = −84°, −175°, and −265°, corresponds to a unit cell oriented so that [120]* is perpendicular to the thermal gradient direction consistently with results obtained by inGIXD shown in Figure 8a while the second one, pointed out by the presence of a maximum at ϕ′ = −196°, would agree with [140]* of the monoclinic unit cell being parallel to the thermal gradient direction. We will come back on this point in the Discussion section. A similar strategy was adopted to assess the stability of the inplane crystal orientation after melting the samples and cooling them back to room temperature (see Figure S7). Figures 10d−f show the inGIPFs of PEO 4K, 8K, and 20K measured at room temperature before and after melting. For PEO 4K and 8K, the in-plane orientation is totally lost and replaced by a random orientation of crystallites as demonstrated by the quasi-constant intensity of 120 reflection independently of the azimuthal angle (Figures 10d and 10e). On the contrary, for PEO 20K, the preferred crystal orientation is preserved after annealing and even enhanced. Indeed, a few sharp localized peaks are revealed

4. DISCUSSION One important goal of this work is to use the profound knowledge gained throughout nearly 50 years by several researchers on the crystallization of PEO fractions to understand more deeply how directional crystallization proceeds with the thermal gradient technique herein considered. 4.1. Crystallization of PEO in Large Superimposed Lamellae. The first striking feature common to PEO samples crystallized using the thermal gradient technique is the presence of large lamellae composed of n times-folded PEO chains which superimpose on top of each other in the direction perpendicular to the film surface in a face-on arrangement. It is known from the pioneering works of Spegt et al.54,56 that PEO lamellar thickness evolves stepwise with increasing temperature of crystallization Tcryst or, which is the same, decreasing undercooling ΔT (= Tm − Tcryst), each step corresponding to a decreasing integral number of folds, down to n = 0 for fully extended chains at Tcryst very close to the melting temperature. Practically, however, the latter state is observed only for low molecular weight polymers (Mn < 4000 g/mol typically) as for higher polymer weights the time needed to obtain fully extended chains would be incredibly long (months or even years).53,54,57 The fact that for the same thermal gradient conditions n seemingly increases with polymer mass (n = 1 for PEO 4K, 2 for PEO 8K, and presumably ≈3−4 for PEO 20K) is consistent with a fixed position of the growth front for a given set of thermal gradient parameters (indeed, then ΔT increases with polymer mass, and so is n as Tm increases with polymer mass) as stated by other authors.49 Considering the works of Kovacs et al.59 on PEO single crystals, it is even possible to restrict the actual position of the growth front to the temperature interval (50.6− 54.6 °C) which corresponds to the domain of existence of once-folded PEO single crystals obtained for a polymer with similar mass and polydispersity as PEO 4K studied here. Furthermore, a crystallization temperature in this range leads, according to Spegt et al.,54,56 to the stabilization of twice-folded and 4−5 times-folded lamellae for PEO samples of mass Mn = 10 000 g/ mol and Mn = 20 000 g/mol, respectively, so in good agreement with our results for PEO 8K (n = 2) and PEO 20K (n ≈ 3−4). Importantly, a quite remarkable and unique peculiarity of PEO lamellae crystallized using the thermal gradient technique (more particularly for the thinnest films) is that they superimpose upon each other in registry; i.e., the orientation of the unit cell is essentially the same for all lamellae as shown in particular by in-plane X-ray diffraction measurements (Results section 3). Indeed for these measurements, the incident X-ray beam spans the whole thickness of the film so that the obtained diffraction pattern represents an average of the in-plane arrangement of all lamellae, yet with relatively J


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Macromolecules reduced orientation distribution, supporting our statement. Such long-range ordered lamellae have been already reported in the literature for polymer single crystals73−76 but, as far as we know, not for PEO. Crystallization of orientationally registered lamellae stacked on top of each other in polymer thin films necessarily originates from a unique primary lamella which transfers its structural information from lamella to lamella maybe through screw dislocation defects76,77 or branching of the primary lamella and diffusion of polymer chains at the surface of the amorphous layer.78 In our particular case, we hypothesize that with the chosen thermal gradient conditions production of a unique primary single crystalline lamella was achieved which was subsequently allowed to grow not only inplane but also out-of-plane (generating other lamellae on top which keep the same unique crystal orientation), thus leading to the growth of films with 3D single-crystal-like orientation. 4.2. Orientation of the Unit Cell for PEO Thin Films Crystallized Using the Thermal Gradient Technique. 4.2.1. Mechanism of Directional Crystallization. It has been shown that the PEO unit cell preferentially orients so that the a* direction is parallel to the thermal gradient direction; that is, {100} crystallographic faces grow in the direction of the thermal gradient direction (i.e., the latter is normal to {100} faces). Kovacs et al. have demonstrated for once-folded PEO single crystals (Mn = 6000 g/mol) that these faces possess the slowest growth rate,60,63 slightly slower than that of {140} faces, both families of faces constituting the sides of the classical hexagonal shape observed for melt-grown PEO single crystals independently of the polymer mass investigated (up to Mn = 10 000 g/mol). In these studies, growth along the {001} faces (parallel to the lamellar planes) was not considered as only monolayer single crystals were produced. In the present case, PEO lamellar planes are parallel to the film surface, and consequently only {hk0} faces can grow in the plane of the film. This is actually due to the fact that terminal OH groups of PEO chains tend to protrude from the outer surfaces of PEO lamellae,59 thus reducing the interface tension with the underlying substrate (glass substrate with SiOH groups at the surface) and also between subsequent lamellae due to favorable interactions. Then, considering the crystal orientation in the plane of the lamellae, it appears that crystallization occurs so that the slowest growing face is “selected” to grow along the sample displacement direction, allowing the formation of large size highly oriented crystals. This hypothesis was proposed in a previous study of the crystallization of a model small molecule, 3-terthiophene, using the thermal gradient technique42 and is fairly well confirmed here for PEO. Moreover, our results are also in line with previous works which showed theoretically and experimentally that the selection of a particular crystal orientation during directional solidification originate from surface tension anisotropy.79,80 4.2.2. Influence of Polymer Mass on the Unit Cell Orientation. Nevertheless, although the same preferred orientation of the unit cell is put forward for all PEO samples, some nuances are clearly visible as a function of polymer mass. In the case of PEO 4K, two distinct unit cell orientations seem to compete, namely, the first one O1 with (100) planes perpendicular to the thermal gradient direction (already mentioned above) and the second one O2 with one of the (120) planes which is parallel to the thermal gradient direction (see Figure 11e). Results obtained by inGIXD and inGIPF lead to the conclusion that PEO 4K samples are composed of several domains of one or the other unit cell orientation with

Figure 11. (a−d) Schematic representation of the in-plane orientation of the unit cell for (a) PEO 4K, (b) PEO 8K, (c) PEO 20K, and (d) PEO 20K (after melting) thin films after thermal gradient treatment. The thin films are supposed to be composed of several domains delimited by blue solid lines. Arrows represent the [100]* direction in each of those domains as referred to the thermal gradient direction schematized by the arrow in red to blue color gradient represented in (e) (the direction of crystal growth is also indicated by a same arrow but oriented in the opposite direction). (e) Hypothetical PEO crystal morphologies (composed of the relevant faces described in the text), designed to emphasize the four unit cell orientations (O1, O2, O3, and O4; see text) revealed for PEO thin films crystallized using the thermal gradient technique.

sharp boundaries between them. Indeed, results given in Figure 9a (inGIXD measurement), Figure 9b (inGIPF measurement), and Figure 10a (inGIPF measurement after 1 month) have been taken on the same sample but at slightly different locations as the sample was removed and replaced between the three measurements (they were not performed the same day). Hence, after one month, it seems that the unit cell orientation has been completely modified while it only appears that the signal corresponds to a domain exhibiting orientation O2. Besides, a closer look at Figure 9b shows that low intense maxima at ϕ′ = −180° and ϕ′ = −90°, so corresponding to orientation O2, are present together with much more intense maxima corresponding to orientation O1. Orientation O2 apparently contradicts the above-mentioned criteria imposing the growth of the slowest growing crystal face along the thermal gradient direction. However, it is known for PEO that (120) planes correspond to the planes of folding of the polymeric chains and are densely packed.44,61 Hence, it is expected that the growth of the corresponding {120} faces should be relatively fast as interactions between molecules is high in this direction.81 Then, as a corollary to our statement, the growth of K


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oriented crystalline domains.11,15 Combination of crystallization by the thermal gradient technique and self-seeding procedure could lead to large uniformly oriented crystallites allowing enhanced charge transport along and perpendicular to polymer chains. Lastly, PEO 8K appears to be the optimal choice for crystallization using the thermal gradient technique. Indeed, a unique crystal orientation O1 with fine orientation distribution is revealed by both inGIXD and inGIPF measurements. However, it is not possible from the latter experiments to distinguish between the two resulting possible directions of a* vector toward the thermal gradient direction, that is, either in the same or opposite direction to it (see Figure 9c). This uniaxial crystal orientation is schematized in Figure 11b by arrows pointing in both directions. The same is also made for PEO 4K (Figure 11a) and PEO 20K (Figures 11c,d). PEO 8K is finally a good compromise between too long polymer chains whose orientation can fluctuate more from the imposed direction of crystallization due to the higher overall viscosity of the polymer and too short chains which, to the contrary, are frozen too rapidly in a given orientation and are not able to further rearrange to follow the imposed direction of crystallization.

{120} faces should be oriented in a direction perpendicular to the thermal gradient direction. This can be achieved by orienting PEO single crystals so that either the apex joining two {120} faces is pointing in a direction perpendicular to the thermal gradient direction, so typically O1 (see Figure 11e), or one of the {120} faces is oriented parallel to the thermal gradient direction, so typically O2 (see Figure 11e). The latter configuration should be less stable as it implies that one of the other {120} faces should be oriented almost perpendicularly to the thermal gradient direction (see Figure 11e). However, this orientation can be trapped upon cooling for PEO 4K and further maintained, explaining why it is observed experimentally. Lastly, a third unit cell orientation O3 characterized by (140) planes perpendicular to the thermal gradient direction was also put forward in Figure 10a. As said above, {140} faces are slow growing faces of PEO single crystals with growth rate a little higher than {100} faces. However, it has also been demonstrated that the relative magnitude of those growth rates can be inverted in very reduced temperature range due to chain unfolding processes.60 Hence, maybe at some local points of PEO 4K sample crystallization proceeded within the temperature range where the growth rate of {140} faces was smaller than that of {100} faces and was then favored along the thermal gradient direction. The corresponding unit cell orientation O3 was then subsequently frozen out upon cooling to room temperature, thus explaining its observation by inGIPF. In the case of PEO 20K, the presence of both orientations O1 and O2 is also revealed by inGIXD measurements. However, here, PEO thin films are more likely composed of domains exhibiting a broad distribution of orientations around the major one O1 (Figure 11c). In clear contrast with the latter behavior, after melting, two unit cell orientations are put forward, namely O1 and a new one O4 characterized by {110} faces being perpendicular to the thermal gradient direction (see Figure 11d), with both very sharp distribution of orientations. This significant enhancement of crystal orientation distribution can be rationalized by a self-seeding mechanism which is known to be more efficient with increasing polymer mass.59,82 At the melting temperature, microcrystals exhibiting orientation O1 (as produced after thermal gradient treatment) still remain (maybe because distorted O1 crystals have a slightly lower melting point than O1 crystals due to lower lamellar thickness probably originating from tilted lamellae) and subsequently act as seeds during crystallization occurring upon cooling. The number of “cloned” crystals grown decreases exponentially with the self-seeding temperature. In the present case, samples were heated up to 72 °C before cooling back to room temperature (see Supporting Information), but it seems that an appreciable number of cloned crystals have been produced taking into account the sharp distribution of orientation of O1. However, the memory effect of the self-seeding process is not optimal as an additional unit cell orientation, O4, is generated in coexistence with O1. This might be due to the long time spent above the melting temperature (about 30 min) as it is known that the longer the time spent at the self-seeding temperature, the higher is the misorientation between the seed crystals and their clones.82 Hopefully then, crystal orientation of PEO 20K thin films could be even more improved by a careful adjustment of self-seeding temperature and time. This result is of high interest for applications in organic electronics for example as high molecular weight polymers with relatively high crystallinity are suitable choices for good device performance but usually yield thin films composed of small size randomly

5. CONCLUSION PEO thin films fabricated by spin-coating have been submitted to a thermal gradient treatment, and the resulting microstructure has been analyzed by microscopy and several X-ray diffraction techniques. Millimeter-size uniformly oriented domains (parallel to the thermal gradient direction) are put forward by polarized microscopy. Structural analysis of these films showed that PEO chains pack into large lamellae, their thickness varying with polymer mass, in a face-on arrangement (PEO chains stand upright onto the substrate). Suitable thermal gradient and sample parameters could be chosen to optimize the preferred intralamellar crystal orientation which happened to be the same independently of polymer mass, that is, characterized by a* reciprocal vector of the monoclinic unit cell which is parallel to the thermal gradient direction. Particularly noteworthy is that unique in-plane crystal orientation with very fine orientation distribution is revealed for PEO 8K, suggesting the formation of films with singlecrystal-like orientation. Significant enhancement of crystal orientation distribution could be obtained for PEO 20K after melting the films and subsequent cooling to room temperature. This “memory effect”, not observed for PEO 4K and PEO 8K, is a promising result that opens the way to efficient processing methods for applications demanding highly directional properties of high molecular weight polymers, providing that they can be crystallized from the melt. Lastly, PEO 4K is a good model to point out the complexity of directional crystallization of low molecular weight compounds using the thermal gradient technique due to relatively higher crystallization kinetics.

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ASSOCIATED CONTENT

* Supporting Information S

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.7b00441. Figures S1 and S2: additional polarized optical microscopy images of PEO thin films before and during thermal gradient treatment; Figure S3: the sXRD pattern L


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of an as-cast PEO 8K thin film; Figures S4 and S6: GIWAXS images of PEO 8K thin films obtained after thermal gradient treatment with indexed reflections; Figure S5: sXRD patterns of PEO 8K thin films crystallized using the thermal gradient technique with varying film thickness; Figure S7: inGIXD patterns measured at different temperatures for PEO 4K, PEO 8K, and PEO 20K thin films crystallized using the thermal gradient technique (PDF)

(13) Kim, G.; Kang, S.-J.; Dutta, G. K.; Han, Y.-K.; Shin, T. J.; Noh, Y.-Y.; Yang, C. A Thienoisoindigo-Naphthalene Polymer with Ultrahigh Mobility of 14.4 cm2/V·s That Substantially Exceeds Benchmark Values for Amorphous Silicon Semiconductors. J. Am. Chem. Soc. 2014, 136, 9477−9483. (14) Arias, A. C.; MacKenzie, J. D.; McCulloch, I.; Rivnay, J.; Salleo, A. Materials and Applications for Large Area Electronics: SolutionBased Approaches. Chem. Rev. 2010, 110, 3−24. (15) Biniek, L.; Schroeder, B. C.; Donaghey, J. E.; Yaacobi-Gross, N.; Ashraf, R. S.; Soon, Y. W.; Nielsen, C. B.; Durrant, J. R.; Anthopoulos, T. D.; McCulloch, I. New Fused Bis-Thienobenzothienothiophene Copolymers and Their Use in Organic Solar Cells and Transistors. Macromolecules 2013, 46, 727−735. (16) Guo, X.; Facchetti, A.; Marks, T. J. Imide- and AmideFunctionalized Polymer Semiconductors. Chem. Rev. 2014, 114, 8943−9021. (17) Lee, J.; Chung, J. W.; Kim, D. H.; Lee, B.-L.; Park, J.-I.; Lee, S.; Häusermann, R.; Batlogg, B.; Lee, S.-S.; Choi, I.; Kim, I. W.; Kang, M. S. Thin Films of Highly Planar Semiconductor Polymers Exhibiting Band-like Transport at Room Temperature. J. Am. Chem. Soc. 2015, 137, 7990−7993. (18) Zhang, R.; Li, B.; Iovu, M. C.; Jeffries-El, M.; Sauvé, G.; Cooper, J.; Jia, S.; Tristram-Nagle, S.; Smilgies, D. M.; Lambeth, D. N.; McCullough, R. D.; Kowalewski, T. Nanostructure Dependence of Field-Effect Mobility in Regioregular Poly(3-hexylthiophene) Thin Film Field Effect Transistors. J. Am. Chem. Soc. 2006, 128, 3480−3481. (19) Li, H.; Yan, S. Surface-Induced Polymer Crystallization and the Resultant Structures and Morphologies. Macromolecules 2011, 44, 417−428. (20) Brinkmann, M.; Wittmann, J. C. Orientation of Regioregular Poly(3-hexylthiophene) by Directional Solidification: A Simple Method to Reveal the Semicrystalline Structure of a Conjugated Polymer. Adv. Mater. 2006, 18, 860−863. (21) Brinkmann, M. Directional Epitaxial Crystallization and Tentative Crystal Structure of Poly(9,9′-di-n-octyl-2,7-fluorene). Macromolecules 2007, 40, 7532−7541. (22) Brinkmann, M.; Contal, C.; Kayunkid, N.; Djuric, T.; Resel, R. Highly Oriented and Nanotextured Films of Regioregular Poly(3hexylthiophene) Grown by Epitaxy on the Nanostructured Surface of an Aromatic Substrate. Macromolecules 2010, 43, 7604−7610. (23) Chang, H.; Guo, Q.; Shen, D.; Li, L.; Qiu, Z.; Wang, F.; Yan, S. Study on the Oriented Recrystallization of Carbon-Coated Polyethylene Oriented Ultrathin Films. J. Phys. Chem. B 2010, 114, 13104− 13109. (24) Chang, H.; Zhang, J.; Li, L.; Wang, Z.; Yang, C.; Takahashi, I.; Ozaki, Y.; Yan, S. A Study on the Epitaxial Ordering Process of the Polycaprolactone on the Highly Oriented Polyethylene Substrate. Macromolecules 2010, 43, 362−366. (25) Brinkmann, M.; Charoenthai, N.; Traiphol, R.; Piyakulawat, P.; Wlosnewski, J.; Asawapirom, U. Structure and Morphology in Highly Oriented Films of Poly(9,9-bis(n-octyl)fluorene-2,7-diyl) and Poly(9,9-bis(2-ethylhexyl)fluorene-2,7-diyl) Grown on Friction Transferred Poly(tetrafluoroethylene). Macromolecules 2009, 42, 8298− 8306. (26) Yan, C.; Li, H.; Zhang, J.; Ozaki, Y.; Shen, D.; Yan, D.; Shi, A.C.; Yan, S. Surface-Induced Anisotropic Chain Ordering of Polycarprolactone on Oriented Polyethylene Substrate: Epitaxy and Soft Epitaxy. Macromolecules 2006, 39, 8041−8048. (27) Werzer, O.; Matoy, K.; Smilgies, D.-M.; Rothmann, M. M.; Strohriegl, P.; Resel, R. Uniaxially aligned poly[(9,9-dioctylfluorenyl2,7-diyl)-co-bithiophene] thin films characterized by the X-ray diffraction pole figure technique. J. Appl. Polym. Sci. 2008, 107, 1817−1821. (28) van de Craats, A. M.; Stutzmann, N.; Bunk, O.; Nielsen, M. M.; Watson, M.; Müllen, K.; Chanzy, H. D.; Sirringhaus, H.; Friend, R. H. Meso-Epitaxial Solution-Growth of Self-Organizing Discotic LiquidCrystalline Semiconductors. Adv. Mater. 2003, 15, 495−499. (29) Bunk, O.; Nielsen, M. M.; Sølling, T. I.; van de Craats, A. M.; Stutzmann, N. Induced Alignment of a Solution-Cast Discotic

AUTHOR INFORMATION

Corresponding Author

*(G.G.) E-mail: [email protected]. ORCID

Gabin Gbabode: 0000-0001-6329-3769 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The research leading to these results has received funding from the European Community’s Seventh Framework Program (FP7/2007-2013) under grant agreement no. 212311 of the ONE-P project and the Belgian Federation Science Policy Office (PAI-SF2). G.S. kindly acknowledges support from the Belgian National Fund for Scientific Research (FNRS, Research Fellow) and postdoctoral fellowship support from the WienerAnspach Foundation. The authors gratefully thank Prof Roland Resel for helpful discussions on pole figures.

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REFERENCES

(1) Sperling, L. H. Introduction to Physical Polymer Science; John Wiley & Sons: 2005. (2) Strobl, G. R. The Physics of Polymers: Concepts for Understanding Their Structures and Behavior; Springer: 2007. (3) Mandelkern, L. Crystallization of Polymers; Cambridge University Press: 2002; Vol. 1. (4) Pennings, A. J.; van der Mark, J. M. A. A.; Booij, H. C. Hydrodynamically induced crystallization of polymers from solution. Colloid Polym. Sci. 1970, 236, 99−111. (5) Southern, J. H.; Porter, R. S. The properties of polyethylene crystallized under the orientation and pressure effects of a pressure capillary viscometer. J. Appl. Polym. Sci. 1970, 14, 2305−2317. (6) McHugh, A. J.; Schultz, J. M. Extensional flow induced crystallization from solution. Kolloid Z. Z. Polym. 1973, 251, 193−214. (7) Griswold, P. D.; Porter, R. S.; Desper, C. R.; Farris, R. J. Preparation of ultradrawn polyethylene by a novel radial-compression method. Polym. Eng. Sci. 1978, 18, 537−545. (8) Smith, P.; Lemstra, P. J. Ultra-high-strength polyethylene filaments by solution spinning/drawing. J. Mater. Sci. 1980, 15, 505−514. (9) Li, X.; Yu, X.; Han, Y. Polymer thin films for antireflection coatings. J. Mater. Chem. C 2013, 1, 2266−2285. (10) Ong, B. S.; Wu, Y.; Li, Y.; Liu, P.; Pan, H. Thiophene Polymer Semiconductors for Organic Thin-Film Transistors. Chem. - Eur. J. 2008, 14, 4766−4778. (11) McCulloch, I.; Heeney, M.; Chabinyc, M. L.; DeLongchamp, D. M.; Kline, R. J.; Cölle, M.; Duffy, W.; Fischer, D.; Gundlach, D.; Hamadani, B.; Hamilton, R.; Richter, L.; Salleo, A.; Shkunov, M.; Sparrowe, D.; Tierney, S.; Zhang, W. Semiconducting Thienothiophene Copolymers: Design, Synthesis, Morphology, and Performance in Thin-Film Organic Transistors. Adv. Mater. 2009, 21, 1091−1109. (12) Kang, I.; Yun, H.-J.; Chung, D. S.; Kwon, S.-K.; Kim, Y.-H. Record High Hole Mobility in Polymer Semiconductors via SideChain Engineering. J. Am. Chem. Soc. 2013, 135, 14896−14899. M


Article

Macromolecules Hexabenzocoronene Derivative for Electronic Devices Investigated by Surface X-ray Diffraction. J. Am. Chem. Soc. 2003, 125, 2252−2258. (30) Hu, Z.; Jonas, A. M. Control of crystal orientation in soft nanostructures by nanoimprint lithography. Soft Matter 2010, 6, 21− 28. (31) Aryal, M.; Trivedi, K.; Hu, W. Nano-Confinement Induced Chain Alignment in Ordered P3HT Nanostructures Defined by Nanoimprint Lithography. ACS Nano 2009, 3, 3085−3090. (32) Cavallini, M.; Stoliar, P.; Moulin, J.-F.; Surin, M.; Leclère, P.; Lazzaroni, R.; Breiby, D. W.; Andreasen, J. W.; Nielsen, M. M.; Sonar, P.; Grimsdale, A. C.; Müllen, K.; Biscarini, F. Field-Effect Transistors Based on Self-Organized Molecular Nanostripes. Nano Lett. 2005, 5, 2422−2425. (33) Crossland, E. J. W.; Tremel, K.; Fischer, F.; Rahimi, K.; Reiter, G.; Steiner, U.; Ludwigs, S. Anisotropic Charge Transport in Spherulitic Poly(3-hexylthiophene) Films. Adv. Mater. 2012, 24, 839−844. (34) Smith, J.; Hamilton, R.; McCulloch, I.; Stingelin-Stutzmann, N.; Heeney, M.; Bradley, D. D. C.; Anthopoulos, T. D. Solution-processed organic transistors based on semiconducting blends. J. Mater. Chem. 2010, 20, 2562−2574. (35) Derue, G.; Serban, D. A.; Leclère, P.; Melinte, S.; Damman, P.; Lazzaroni, R. Controlled nanorubbing of polythiophene thin films for field-effect transistors. Org. Electron. 2008, 9, 821−828. (36) Biniek, L.; Pouget, S.; Djurado, D.; Gonthier, E.; Tremel, K.; Kayunkid, N.; Zaborova, E.; Crespo-Monteiro, N.; Boyron, O.; Leclerc, N.; Ludwigs, S.; Brinkmann, M. High-Temperature Rubbing: A Versatile Method to Align π-Conjugated Polymers without Alignment Substrate. Macromolecules 2014, 47, 3871−3879. (37) DeLongchamp, D. M.; Kline, R. J.; Jung, Y.; Germack, D. S.; Lin, E. K.; Moad, A. J.; Richter, L. J.; Toney, M. F.; Heeney, M.; McCulloch, I. Controlling the Orientation of Terraced Nanoscale “Ribbons” of a Poly(thiophene) Semiconductor. ACS Nano 2009, 3, 780−787. (38) Lu, G. H.; Li, L. G.; Yang, X. N. Achieving Perpendicular Alignment of Rigid Polythiophene Backbones to the Substrate by Using Solvent Vapor Treatment. Adv. Mater. 2007, 19, 3594−3598. (39) Rahimi, K.; Botiz, I.; Stingelin, N.; Kayunkid, N.; Sommer, M.; Koch, F. P. V.; Nguyen, H.; Coulembier, O.; Dubois, P.; Brinkmann, M.; Reiter, G. Controllable Processes for Generating Large Single Crystals of Poly(3-hexylthiophene). Angew. Chem., Int. Ed. 2012, 51, 11131−11135. (40) Bridgman, P. W. Certain physical properties of single crystals of tungsten, antimony, bismuth, tellurium, cadmium, zinc and tin. Proc. Am. Acad. Arts Sci. 1925, 60, 305−383. (41) Stockbarger, D. C. The Production of Large Single Crystals of Lithium Fluoride. Rev. Sci. Instrum. 1936, 7, 133−136. (42) Schweicher, G.; Paquay, N.; Amato, C.; Resel, R.; Koini, M.; Talvy, S.; Lemaur, V.; Cornil, J.; Geerts, Y.; Gbabode, G. Toward Single Crystal Thin Films of Terthiophene by Directional Crystallization Using a Thermal Gradient. Cryst. Growth Des. 2011, 11, 3663− 3672. (43) Lovinger, A. J.; Gryte, C. C. The morphology of directionally solidified poly(ethylene oxide) spherulites. Macromolecules 1976, 9, 247−253. (44) Lovinger, A. J.; Lau, C.-M.; Gryte, C. C. Microstructure of directionally solidified poly(ethylene oxide). Polymer 1976, 17, 581− 586. (45) Lovinger, A. J.; Gryte, C. C. Model for the shape of polymer spherulites formed in a temperature gradient. J. Appl. Phys. 1976, 47, 1999−2004. (46) Lovinger, A. J. Crystallographic factors affecting the structure of polymeric spherulites. I. Morphology of directionally solidified polyamides. J. Appl. Phys. 1978, 49, 5003−5013. (47) Lovinger, A. J. Crystallographic factors affecting the structure of polymeric spherulites. II. X ray diffraction analysis of directionally solidified polyamides and general conclusions. J. Appl. Phys. 1978, 49, 5014−5028.

(48) Lovinger, A. J.; Davis, D. D. Electron-microscopic investigation of the morphology of a melt crystallized polyaryletherketone. J. Appl. Phys. 1985, 58, 2843−2853. (49) Toda, A.; Taguchi, K.; Kajioka, H. Structure Evolution in Directional Crystallization of Polymers under Temperature Gradient. Macromolecules 2012, 45, 852−861. (50) Mita, K.; Tanaka, H.; Saijo, K.; Takenaka, M.; Hashimoto, T. Cylindrical Domains of Block Copolymers Developed via Ordering under Moving Temperature Gradient. Macromolecules 2007, 40, 5923−5933. (51) Yager, K. G.; Fredin, N. J.; Zhang, X.; Berry, B. C.; Karim, A.; Jones, R. L. Evolution of block-copolymer order through a moving thermal zone. Soft Matter 2010, 6, 92−99. (52) Singh, G.; Yager, K. G.; Smilgies, D.-M.; Kulkarni, M. M.; Bucknall, D. G.; Karim, A. Tuning Molecular Relaxation for Vertical Orientation in Cylindrical Block Copolymer Films via Sharp Dynamic Zone Annealing. Macromolecules 2012, 45, 7107−7117. (53) Arlif, J. P.; Spegt, P. A.; Skoulios, A. E. Etude de la cristallisation des polymères I. Structure lamellaire de polyoxyéthylènes de faible masse moléculaire. Makromol. Chem. 1966, 99, 160−174. (54) Arlie, P. J. P.; Spegt, P.; Skoulios, A. Etude de la cristallisation ̂ des polymères. II. Structure lamellaire et repliement des chaines du polyoxyéthylène. Makromol. Chem. 1967, 104, 212−229. (55) Gilg, B.; Spegt, P.; Terrisse, J.; Skoulios, A. Étude de la cristallisation des polymères. IV. Influence de la distribution de la masse moléculaire sur l’épaisseur des lamelles d’un polyoxyéthylène. Makromol. Chem. 1967, 107, 39−45. (56) Spegt, P. Etude de l’évolution structurale de polyoxyéthylènes pendant leur recuit. Makromol. Chem. 1970, 139, 139−152. (57) Buckley, C. P.; Kovacs, A. J. Melting behaviour of low molecular weight poly (ethylene-oxide) fractions. 1. Extended chain crystals. Prog. Coll. Polym. Sci. 1975, 58, 44−52. (58) Buckley, C. P.; Kovacs, A. J. Melting behaviour of low molecular weight poly (ethylene-oxide) fractions. 2. Folded chain crystals. Colloid Polym. Sci. 1976, 254, 695−715. (59) Kovacs, A. J.; Gonthier, A. Crystallization and fusion of selfseeded polymers. II Growth rate, morphology and isothermal thickening of single crystals of low molecular weight poly(ethyleneoxide) fractions. Kolloid Z. Z. Polym. 1972, 250, 530−552. (60) Kovacs, A. J.; Gonthier, A.; Straupe, C. Isothermal growth, thickening, and melting of poly(ethylene oxide) single crystals in the bulk. J. Polym. Sci., Polym. Symp. 1975, 50, 283−325. (61) Kovacs, A. J.; Straupe, C.; Gonthier, A. Isothermal growth, thickening, and melting of poly(ethylene oxide) single crystals in the bulk. II. J. Polym. Sci., Polym. Symp. 1977, 59, 31−54. (62) Kovacs, A. J.; Straupe, C. Isothermal growth, thickening and melting of poly(ethylene-oxide) single crystals in the bulk: III. Bilayer crystals and the effect of chain ends. J. Cryst. Growth 1980, 48, 210− 226. (63) Kovacs, A. J.; Straupe, C. Isothermal growth, thickening and melting of poly(ethylene oxide) single crystals in the bulk. Part 4.Dependence of pathological crystal habits on temperature and thermal history. Faraday Discuss. Chem. Soc. 1979, 68, 225−238. (64) Schweicher, G. Alignment of organic semiconductors in a thermal gradient PhD thesis, Université Libre de Bruxelles, Brussels, Belgium, 2012. (65) Takahashi, Y.; Tadokoro, H. Structural studies of polyethers, (-(CH2)mO-)n. X. Crystal structure of poly(ethylene oxide). Macromolecules 1973, 6, 672−675. (66) Brandrup, J.; Immergut, E. H.; Grulke, E. A.; Abe, A.; Bloch, D. R. Polymer Handbook; Wiley: New York, 1999. (67) Apitz, D.; Bertram, R. P.; Benter, N.; Hieringer, W.; Andreasen, J. W.; Nielsen, M. M.; Johansen, P. M.; Buse, K. Investigation of chromophore-chromophore interaction by electro-optic measurements, linear dichroism, x-ray scattering, and density-functional calculations. Phys. Rev. E 2005, 72, 036610. (68) Asokan, K.; Anil Kumar, C. V.; Dasan, J.; Radhakrishnan, K.; Satheesh Kumar, K.; Ramamohan, T. R. Review of chaos in the dynamics and rheology of suspensions of orientable particles in simple N


Article

Macromolecules shear flow subject to an external periodic force. J. Non-Newtonian Fluid Mech. 2005, 129, 128−142. (69) Breiby, D. W.; Bunk, O.; Andreasen, J. W.; Lemke, H. T.; Nielsen, M. M. Simulating X-ray diffraction of textured films. J. Appl. Crystallogr. 2008, 41, 262−271. (70) Schönherr, H.; Frank, C. W. Ultrathin Films of Poly(ethylene oxides) on Oxidized Silicon. 1. Spectroscopic Characterization of Film Structure and Crystallization Kinetics. Macromolecules 2003, 36, 1188−1198. (71) Patterson, A. L. The Scherrer formula for X-ray particle size determination. Phys. Rev. 1939, 56, 978−982. (72) Bagheri, M.; Rad, R. Z. Synthesis and characterization of thermotropic liquid crystalline polyesters with biphenyl unit in the main chain. React. Funct. Polym. 2008, 68, 613−622. (73) Núñez, E.; Gedde, U. W. Single crystal morphology of starbranched polyesters with crystallisable poly(ε-caprolactone) arms. Polymer 2005, 46, 5992−6000. (74) Fernández, C. E.; Bermúdez, M.; Alla, A.; Muñoz-Guerra, S.; Tocha, E.; Vancso, G. J. Compared structure and morphology of nylon-12 and 10-polyurethane lamellar crystals. Polymer 2011, 52, 1515−1522. (75) Hong, S.; MacKnight, W. J.; Russell, T. P.; Gido, S. P. Orientationally Registered Crystals in Thin Film Crystalline/ Amorphous Block Copolymers. Macromolecules 2001, 34, 2398−2399. (76) Hong, S.; MacKnight, W. J.; Russell, T. P.; Gido, S. P. Structural Evolution of Multilayered, Crystalline−Amorphous Diblock Copolymer Thin Films. Macromolecules 2001, 34, 2876−2883. (77) Keith, H. D.; Chen, W. Y. On the origins of giant screw dislocations in polymer lamellae. Polymer 2002, 43, 6263−6272. (78) Zhang, H.; Yu, M.; Zhang, B.; Reiter, R.; Vielhauer, M.; Mülhaupt, R.; Xu, J.; Reiter, G. Correlating Polymer Crystals via SelfInduced Nucleation. Phys. Rev. Lett. 2014, 112, 237801. (79) Utter, B.; Bodenschatz, E. Dynamics of low anisotropy morphologies in directional solidification. Phys. Rev. E: Stat. Phys., Plasmas, Fluids, Relat. Interdiscip. Top. 2002, 66, 051604. (80) Amoorezaei, M.; Gurevich, S.; Provatas, N. Orientation selection in solidification patterning. Acta Mater. 2012, 60, 657−663. (81) Hartman, P.; Bennema, P. The attachment energy as a habit controlling factor: I. Theoretical considerations. J. Cryst. Growth 1980, 49, 145−156. (82) Reiter, G. Some unique features of polymer crystallisation. Chem. Soc. Rev. 2014, 43, 2055−2065.

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Unique Crystal Orientation of Poly(ethylene oxide) Thin Films by

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