Formation, morphology and internal structure of one-dimensional nanostructuresof the ferroelectric polymer P(VDF-TrFE)

Nitin Shingne a,b, Markus Geuss b, Brigitte Hartmann-Azanza c, Martin Steinhart c,*,Thomas Thurn-Albrecht a,*aMartin-Luther-Universität Halle Wittenberg, Institut für Physik, 06099 Halle, GermanybMax-Planck-Institut für Mikrostrukturphysik, Weinberg 2, 06120 Halle, GermanycUniversität Osnabrück, Institut für Chemie neuer Materialien, Barbarastrasse 7, 49069 Osnabrück, Germany

a r t i c l e i n f o

Article history:Received 16 January 2013Received in revised form11 March 2013Accepted 18 March 2013Available online 26 March 2013

Keywords:PVDF-TrFECapillary wettingTemplated nanostructures

a b s t r a c t

We present an investigation of the formation, morphology, thermodynamic properties and crystaltexture of P(VDF-TrFE) nanostructures obtained by wetting of porous anodic aluminum oxide templates,using electron microscopy, DSC and X-ray scattering. Wetting occurs initially via precursor film formationproducing thin walled nanotubes. Formation of nanorods happens by subsequent thickening of the tubewalls with thickness undulations occurring as a transient state. While crystallization within the one-dimensional nanostructures during cooling leads to the formation of highly oriented crystals in thehigh temperature paraelectric phase, the subsequent transition to the ferroelectric phase goes along withpartial loss of orientation. Nevertheless, at room temperature the samples show a dominant orientationof the chains perpendicular to the long axis of the nanostructures. For narrow pores the size of crystallinedomains is restricted by the strong confinement.

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1. Introduction

The preparation of functional 1D polymer nanostructures byinfiltration of nanoporous templates is a well-established method.However, understanding of important aspects of the wetting pro-cess is incomplete, in particular as to the nature of the transitionbetween tubular segments formed by precursor films and solidsegments. Similarly, our knowledge is incomplete about the effectof confinement in 1D polymer nanostructures on ordering pro-cesses and phase transitions. In this paper we report about in-vestigations of both aspects using ferroelectric polymers as anexemplary functional system. Nanoporous anodic aluminum oxide(AAO) is frequently used as rigid shape-defining template to pro-duce polymeric nanotubes and nanorods [1e3]. Due to the highsurface energy of the AAO nanopores, in equilibrium they arecompletely filled by the polymer forming thus nanorods. However,infiltration kinetics is complex and in general two infiltrationmechanisms, precursor wetting and capillary wetting, may occur[4]. Capillary wetting of AAO was observed for some very viscoushomopolymers and block copolymer melts [5e8]. Precursor

wetting on the other hand involves at first the formation of annularprecursor films [9e11] on the walls of cylindrical pores. The sub-sequent transition to complete filling has been much less studiedfor polymeric systems than for liquids consisting of small mole-cules. For the latter case it was reported that filling can occur via theformation of menisci which span across the entire channel crosssection [12,13] and move along the channel as more liquid istransported to the menisci. This mechanism seems to be oftenfaster than the alternative process, namely filling via the formationof Rayleigh instabilities [14,15]. In any case, for highly viscouspolymer melts filling is slow and thus solidified polymer nanotubesare accessible by thermal quenching [2,3]. Generally, in semi-crystalline polymers lamellar thickness and spherulitic growth arecharacterized by length scales affected by the two-dimensionalconfinement imposed by the rigid AAO nanopore walls [16]. As aconsequence the formation of spherulites is suppressed in AAOnanopores, and directed crystal growth resulting in preferentialcrystal orientation occurs [17]. Typically, the chains are orientednormal to the nanopore axes so that the lamellae can grow alongthe nanopores. This mechanism was initially described for poly(-vinylidene difluoride) (PVDF) [18,19] and later shown to apply alsofor other semicrystalline polymers [20,21]. As a common functionalpolymer we here use the ferroelectric random copolymer poly(-vinylidene difluoride-ran-trifluoroethylene), P(VDF-TrFE). Ferro-electric polymer nanostructures have many potential applications

* Corresponding authors.E-mail addresses: martin.steinhart@uos.de (M. Steinhart), thurn-albrecht@

physik.uni-halle.de (T. Thurn-Albrecht).

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Polymer 54 (2013) 2737e2744

in ferroelectric memory devices, actuators and sensors [22,23]. Theferroelectric properties of thin films, nanotubes and nanorods, aswell as dot-like nanostructures consisting of ferroelectric polymers[19,24e27] are determined by outer shape, i.e. morphology andcrystal orientation within the nanostructure. P(VDF-TrFE)commonly forms a stable orthorhombic and ferroelectric low-temperature phase, in which the chains adopt an all-trans confor-mation so that all dipole moments of a chain point along the b-axis[28,29]. At the Curie-temperature TC, the material shows a solidesolid transition between the ferroelectric low-temperature phaseand a hexagonal paraelectric high-temperature phase, character-ized by highly mobile chains with transetrans and trans-gaucheconformations [29,30]. As P(VDF-TrFE) develops relatively thicklamellar crystals, confinement in narrow AAO nanopores affects thecrystal size of P(VDF-TrFE) leading to changes in melting temper-ature [25]. Wang et al. demonstrated polarization switching inP(VDF-TrFE) nanotubes by external electric fields [26]. Similarphenomena could be observed under 3D confinement, as revealedby piezoresponse force microscopy [27]. Here we present a detailedstudy of formation, morphology, thermodynamic properties, andcrystal texture of one-dimensional P(VDF-TrFE) nanostructuresobtained by wetting of AAO templates, using electron microscopy,differential scanning calorimetry (DSC) and X-ray scattering. Weshow that during wetting precursor films, covering the walls of thetemplate pores, become gradually thicker and show thickness un-dulations. Eventually the precursor film becomes thick enough tofill the entire pore volume. In addition, filling via Rayleigh in-stabilities in narrower pores could be deduced from the observa-tion of periodic voids in the nanorods. Concerning the internalstructure of nanorods and nanotubes we present a qualitativethermodynamic scheme rationalizing the observed confinementeffects on transition temperatures and analyze how the transitionfrom the paraelectric high-temperature phase to the ferroelectriclow-temperature phase affects crystal orientation inside AAOnanopores.

2. Experimental section

2.1. Materials and sample preparation

P(VDF-TrFE) granules were obtained from Piezotech, France. Themolecular weight was determined by GPC on a Viscotek VE2001system with Viscotek VE3580 refractive index detector usingpolystyrene as standard for external calibration and dime-thylformamide as solvent. The following result was obtained:Mn ¼ 100 kg/mol, Mw ¼ 210 kg/mol. The content of VDF and TrFEunits in the P(VDF-TrFE) was determined by using 1H NMR. ForNMR measurements P(VDF-TrFE) was dissolved in deuteratedAcetone and the 1H NMR measurements were performed on aVarion Gemini 400 MHz FT-NMR spectrometer, MestRec (4.9.9.9)was used for the data evaluation. The NMR spectrum shows twoprominent peaks, one between 5 and 6 ppm related toTrFE unit andsecond between 2 and 2.3 ppm related to VDF unit. The intensityratio between TrFE:VDF is 1:6.7 and considering that TrFE has onlyone Hydrogen atom and VDF two hydrogen atoms, the actual in-tensity ratio for TrFE:VDF is 1:3.35 which corresponds to TrFE:VDFcontent of 23:77 mol% in P(VDF-TrFE).

P(VDF-TrFE) nanostructures were prepared by melt infiltrationinto self-ordered AAO templates prepared by Masuda’s two-stepanodization procedure [31e33]. Self-ordered AAO contains arraysof aligned cylindrical nanopores with narrow diameter distributionand uniform nanopore depth. Using precision furnaces operableunder vacuum or argon, the AAO was at first heated to the chosenwetting temperature Tw (either 200 "C or 250 "C) for 10 min. Then,P(VDF-TrFE) was spread on the AAO surface using a spatula. The

samples were then heated to Tw for a specific wetting time tw(either 10 min or 12 h) and then cooled to room temperature at acooling rate of 1 K/min. The spreading of the P(VDF-TrFE) on theAAO surface was conducted under argon. All other high-temperature steps were carried out under vacuum. Before furthercharacterization, the bulk P(VDF-TrFE) film on the AAO surface wasremoved using scalpels. Higher Tws led to the formation ofconsiderably longer P(VDF-TrFE) nanostructures (e.g., w16 mm at200 "C and w31 mm at 250 "C as determined by evaluation ofexemplary TEM images).

2.2. Electron microscopy

For transmission electron microscopy (TEM), we used a JEOL1010 electron microscope operated at an acceleration voltage of100 kV. The P(VDF-TrFE) nanostructures were released by etchingthe AAO with 40 wt-% KOH solution. The obtained suspension ofP(VDF-TrFE) nanostructures was then neutralized by severalwashing cycles including centrifugation, removal of supernatantliquid and redispersion in deionized water.

2.3. DSC

A Perkin Elmer DSC 7 was used for calorimetric experimentsperformed at a scan rate of 20 K/min. Before any measurement, theP(VDF-TrFE) surface film was removed from the AAO surface.Subsequently, the aluminum substrate to which the AAO layerswere attached were etched with a solution of 1.7 g copper chloride(CuCl2$H2O, 50ml fuming HCl(aq) and 50ml deionizedwater) undercooling with water/ice mixtures.

2.4. X-ray scattering

Crystal textures of P(VDF-TrFE) nanostructures in AAO werestudied by q/2q scans and Schulz scans [34] in reflection geometryusing a diffractometer PANalytical X’Pert and a PANalytical X’PertPro MRD diffractometer (high temperature measurements), bothoperated with Cu Ka radiation equipped with a Eulerian cradle.Orientation distributions of specific sets of lattice planes were ob-tained by Schulz scans. For the high temperature q#2q measure-ments, as well as for Schulz scans, the sample was heated to 180 "Cand cooled to measuring temperature with a cooling rate of 2 K/min, at each measuring temperature the sample was equilibratedfor 3 min before starting themeasurement. For the Schulz scans the2q position was fixed based on the temperature dependent q#2qmeasurements. Fig. 1 shows the X-ray diffraction set-up for q/2qand Schulz scans.

3. Results and discussion

3.1. Formation of P(VDF-TrFE) nanostructures by template wetting

Wetting of a flat surface by a liquid is generally classified intotwo thermodynamic regimes which can be distinguished by thespreading parameter S [35], given by

S ¼ gsv # ðgsl þ glvÞ (1)

Here gsv, gsl, and glv are the interfacial energies between substrateand vapor, substrate and liquid, and liquid and vapor, respectively.In case of positive spreading parameter, S > 0, the total surfaceenergy is lowered if the liquid spreads on the substrate. This case iscalled complete wetting. The liquid spreads on the substrate byadvancing a precursor film of microscopic thickness [35]. Thecontact angle in this case is 0". If on the other hand S < 0, the liquid

N. Shingne et al. / Polymer 54 (2013) 2737e27442738

droplet rests on the substrate with an equilibrium shape with acontact angle qe, given by Young’s equation.

gsv ¼ gsl þ glvcos qe (2)

If qe < 90" one speaks of partial wetting. Different from the caseof flat surfaces, for narrow pores in contact with a liquid equilib-rium corresponds in both of the above cases to completely filledpores. Here we assume that the effect of gravity can be neglected.The pathway though, inwhich filling can be achieved is different forthe two cases. In case of partial wetting the pores are filled bycapillary rise where a liquid meniscus moves into the capillary andfills the complete pore [7]. The driving force is given by the lowervalue of gsl in comparison to gsv. In the complete wetting regime aprecursor film can first cover the pore surface and if the radius ofthe pore is smaller than the thickness of the film, nanotubes areformed, which can be stabilized by solidification upon cooling[3,4,7,36]. However, such a nanotube is not an equilibrium state, asthe inner surface of the tube and therefore the surface energy de-creases with increasing thickness of the tubewall. The inner surfacevanishes completely when the pore is completely filled resultingagain in a nanorod as an equilibrium state. Once the pore wall iscovered by a precursor film the driving force for further filling, i.e.thickening of the film, is in general reduced as the surface energy ofthe polymer is typically considerably lower than that of the barepore wall. Therefore it is expected that the second step of porefilling will in general be slower.

Fig. 2 shows TEM images of selected parts of released nano-structures obtained by wetting of AAO with a pore diameterDp¼ 400 nmdisclosing interesting details of thewetting process. Inall cases the TEM investigations revealed the presence of threebasic morphology types, as schematically displayed in Fig. 3.Segment A is solid; the P(VDF-TrFE) fills the entire cross section ofthe AAO nanopores. It is reasonable to assume that these solidsegments were located close to the mouths of the AAO nanopores.Tubular segments are located ahead of the solid segments, i.e.further away form the AAO nanopore mouths. In the vicinity of thesolid segments, thick precursor films (diameter in the 100 nmrange) with thickness undulations encompass relatively narrowhollow spaces in the center (Segment B in Fig. 3; Fig. 2a and b). It iswell known that the surface area of liquid films wetting the wall ofcylindrical channels may be minimized by the formation of thick-ness undulations [37,38]. Away from the transition from solidsegments to tubular segments with thick walls, the tube walls

become thinner and smoother, as for thinner annular films thedevelopment of undulations is suppressed by intermolecular in-teractions [15]. Fig. 2c shows a segment of a (VDF-TrFE) nano-structure with thicker, undulating walls (bottom) and anothersegment with thinner smooth walls at the top.

The picture is somewhat different for P(VDF-TrFE) nano-structures prepared in AAO with Dp ¼ 60 nm and Dp ¼ 35 nm.While TEM investigations revealed that at least for long wetting

Fig. 2. TEM images of P(VDF-TrFE) nanostructures inside AAO with Dp ¼ 400 nm. a)Portion of a P(VDF-TrFE) nanostructure (Tw ¼ 250 "C, tw ¼ 10 min) corresponding tosegments A and B in Fig. 3. b) Portion of a P(VDF-TrFE) nanostructure (Tw ¼ 250 "C,tw ¼ 12 h) corresponding to segment B in Fig. 3. c) Portions of two P(VDF-TrFE)nanostructures (Tw ¼ 180 "C, tw ¼ 12 h); the upper P(VDF-TrFE) nanostructure cor-responds to segment C in Fig. 3, the lower P(VDF-TrFE) nanostructure to segment B.

Fig. 3. Schematic diagram of the morphology of P(VDF-TrFE) nanostructures insideAAO with Dp ¼ 400 nm (orange, P(VDF-TrFE); blue, AAO nanopore walls). In segment Aclose to the AAO nanopore mouths, the P(VDF-TrFE) nanostructures are solid. Fartheraway from the AAO nanopore mouths, tubular segments with thick P(VDF-TrFE) filmson the AAO nanopore walls exhibiting thickness undulations follows (segment B).Towards the AAO nanopore bottoms, segment B transforms into tubular segment Ccharacterized by annular P(VDF-TrFE) precursor films covering the AAO nanopore wallsstabilized by Van der Waals forces, which are smooth and thinner than the P(VDF-TrFE)layers in segment B. (For interpretation of the references to color in this figure legend,the reader is referred to the web version of this article.)

Fig. 1. Sketch of the X-ray scattering set-up. During Schulz scans, q and 2qare fixed. TheAAO membrane is tilted around the j axis. The j axis lies in the plane of the AAOsurface as well as in the scattering plane and is oriented perpendicularly to the AAOnanopore axes and to the rotation axis of q/2q scans. The measured I(j) profilesrepresent the orientation distribution of the corresponding lattice planes with respectto the AAO surface.

N. Shingne et al. / Polymer 54 (2013) 2737e2744 2739

times solid structures are prevailing, a significant portion of theP(VDF-TrFE) nanostructures show linear arrays of elongated voids(Fig. 4), which are often located between two solid segments. Werationalize this observation as follows. At first, a mesoscopic P(VDF-TrFE) precursor film forms on the walls of the AAO nanopores(Fig. 5a). Probably at defect sites, menisci spanning across the AAOnanopore cross section develop (Fig. 5b). As more P(VDF-TrFE)flows into the AAO nanopores, the interfaces of the menisci movein opposite directions so that the AAO nanopores are partially filled(Fig. 5c). Apparently, “random” meniscus formation accompaniedby interface movement is a faster process than development ofRayleigh instabilities. However, in some of the plug segmentslocated between solid segments, Rayleigh instabilities may haveenough time to develop and grow resulting in linear arrays of voidssurrounded by P(VDF-TrFE) (Fig. 5d).

Up to now, Rayleigh instabilities inside AAO nanopores havebeen reported to form in the course of post-infiltration annealingafter solution wetting [39,40]. However, the occurrence of Rayleighinstabilities as a transient state during melt infiltration of polymersinto AAO has, to the best of our knowledge, not been reported up tonow. It should be noted that structure formation of annular filmscovering the walls of cylindrical channels or nanopores is complex,and identification of criteria for the stability of such films is far frombeing trivial. It is, however, obvious that formation of menisci andoccurrence of Rayleigh instabilities are by far more important forsmaller nanopore diameters.

All together our results indicate that the filling of the nanoporestakes place in two steps, in the first step tubemorphology is formed

by precursor wetting and over a long wetting period the tube wallsthicken and might even fill completely if the nanopores are narrowenough. These results are different from the result obtained byZhang et al. [7], where the observation of nanorods resp. nanotubesat different temperatures was explained by a transition from partialto complete wetting. To check the plausibility of our experimentalresults on P(VDF-TrFE) we tried to estimate the spreading param-eter S. Generally, the surface energy of inorganic materials likealumina is one order of magnitude larger than the surface energy ofa polymeric melt [41]. As direct experimental data for surface en-ergy of P(VDF-TrFE) and for the interfacial energies between P(VDF-TrFE) and alumina were not available in the literature we usedsurface energy data of PVDF and Equation (3) to estimate the cor-responding quantities.

gsl ¼ glv þ gsv # 2fffiffiffiffiffiffiffiffiffiffiffiffiffiglvgsv

p(3)

where f is Good’s interaction parameter given byf ¼ 4ðVsVlÞ

1=3=ðV1=3s þ V1=3

l Þ2 and Vs and Vl are molar volume ofsolid and liquid, respectively [42]. For the calculation the followingvalues were used, PVDF: gsv z glv ¼ 32.7 ' 10#3 J/m2 [43],Vl¼ 32 cm3/mol; Alumina: gsv¼ 1.59 J/m2 [44], Vs¼ 12.84 cm3/mol.The resulting value for gsl z 1.17 J/m2. Using Equation (1), weobtain the spreading parameter for PVDF-AAO nanoporeS ¼ 0.387 J/m2. The positive value is consistent with the observedprecursor wetting mechanism. Due to the large differences in sur-face energy for the polymer and alumina the sign of S is ratherinsensitive to small changes or errors in surface energies of the twomaterials.

Fig. 4. TEM images of P(VDF-TrFE) Dp ¼ 60 nm nanostructures and Dp ¼ 35 nmnanostructures (bottom) released from AAO; Tw amounted to 250 "C and tw was 12 h.

Fig. 5. Schematic diagram of the morphology evolution of P(VDF-TrFE) nanostructuresinside AAO with Dp ¼ 60 nm and Dp ¼ 35 nm (orange, P(VDF-TrFE); blue, AAOnanopore walls). a) A mesoscopic P(VDF-TrFE) precursor film moves into the AAOnanopores and covers the AAO nanopore walls. b) At defects, menisci form. c) As moreliquid P(VDF-TrFE) moves into the AAO nanopore, the interfaces of the menisci eventmove in opposite directions. d) In the plug-like segment resulting from the meniscusformation Rayleigh instabilities lead to the formation of linear arrays of voids sur-rounded by P(VDF-TrFE). (For interpretation of the references to color in this figurelegend, the reader is referred to the web version of this article.)

N. Shingne et al. / Polymer 54 (2013) 2737e27442740

3.2. Phase transitions in P(VDF-TrFE) nanostructures

3.2.1. Effects of confinement on the phase transition temperaturesThe effect of confinement on the thermal properties and the

phase diagram of P(VDF-TrFE) was studied by differential scanningcalorimetry (DSC). All the nanostructures studied here were pre-pared at a Tw of 200 "C with a tw of 10 min Fig. 6a shows the heatingand cooling scans of bulk P(VDF-TrFE) and nanostructures of400 nm, 180 nm, 60 nm and 35 nm diameters. All measurementswere carried out with a scan rate of 20 K/min. The data obtainedfrom the bulk sample show the well known properties of P(VDF-TrFE): Melting around 150 "C and a solidesolid transition at theCurie temperature, TC z 125 "C, between a low temperatureferroelectric and a high temperature paraelectric phase. Bothtransitions show supercooling, as it is typical for first order tran-sitions. Table 1 shows the peak temperatures of the transitions fromall measurements.

The peak melting temperature, Tm, is similar for bulk and400 nm and 180 nm nanostructures, there might be small changesin onset temperature. For the smaller nanostructures (60 nm and35 nm) there is a clear decrease in Tm, indicative of a reduction ofcrystal size due to confinement. The diameter of the nanostructureis of similar size or smaller in this case than the typical lamellarthickness of the crystals. As a second effect we observe a decreasein the crystallization temperature, Tcry, for the smaller nano-structure. For 35 nm structures Tcry is not even separated anymorefrom the para- to ferroelectric transition. This general effect hasbeen observed before in other systems under 3D as well as 2Dconfinement [18,45,46]. Crystallization under strong confinementin many small separated compartments requires high nucleationdensities which can only be realized by homogeneous nucleation atlow temperatures. We can also conclude that obviously the surfaceof the alumina nanopores does not act as a heterogeneous nucle-ation site, otherwise an increase in Tcry should be observed. Ofcourse, the reduced crystallization temperature might alsocontribute to the reduction in melting temperature for the smallernanostructures.

It is interesting, how little the Curie transition, TC, is affected byconfinement. During cooling of the nanostructures, the para- toferro- transition, Tp#f, covers a broad temperature range, for thenanostructures even without a clear peak. (Tp#f for the nano-structures is therefore not indicated in Table 1.) However, the widthof the Tp#f transition is similar in bulk and for the nanostructures.The DSC scans also show that the TC observed during heating re-mains basically unaffected by confinement. At first this seemssurprising as the decrease in Tm shows that the crystal thicknessmust decrease with increasing confinement. More specifically, thisobservation means that the stability of the semicrystalline struc-ture in the ferroelectric and the paraelectric phase is similarlyaffected by the confinement as illustrated in Fig. 6b. If the freeenergy of both phases increases by the same amount the transitiontemperature remains unchanged. A possible explanation would bethat the size of the ferroelectric domains is equal to the crystal sizealong the chain direction. The free energy of the melt state on theother hand is not affected by confinementwhich leads to a decreaseof the melting temperature, Tm.

3.2.2. Effect of thermal annealing on the phase transitiontemperatures

In bulk P(VDF-TrFE) the effects of annealing on the phase tran-sitions were extensively studied [47e50]. We here analyze howconfinement modifies the annealing effects during cooling by aseries of DSC measurements. The samples were cooled withdifferent rates (40 K/min to 1 K/min) and afterwards DSC heatingscans were measured with a constant rate of 20 K/min to study theeffect of the thermal history on TC and Tm. By decreasing the coolingrate from 40 K/min to 1 K/min, the time which the samples effec-tively stay in the paraelectric and the ferroelectric phase duringcooling is effectively increased.

Fig. 6. a) DSC heating and cooling scans (scan rate of 20 K/min) for bulk P(VDF-TrFE)and nanostructures with different diameters; b) Schematic of the Gibbs free energy forferroelectric, paraelectric and melt phase. Due to confinement there is an increase inthe Gibbs free energy, shown by dashed lines (Tm#b: melting temperature of bulk,Tm#p: melting temperature in pores, TC: Curie temperature).

Table 1Peak transition temperatures for bulk P(VDF-TrFE) and the nanostructures from theDSC scans shown in Fig. 6 (TC: ferroelectric to paraelectric transition temperature,Tm: melting temperature, Tcry: crystallization temperature, Tp#f: paraelectric toferroelectric phase transition).

Pore diameters TC("C) Tm("C) Tcry("C) Tp#f("C)

Bulk 125 151 135 76400 nm 125 150 127 e

180 nm 125 151 122 e

60 nm 125 145 101 e

35 nm 125 135 e e

N. Shingne et al. / Polymer 54 (2013) 2737e2744 2741

The DSC heating scans after cooling with different rates areshown in the supporting information (Figs. S1 and S2). Fig. 7ashows a quantitative analysis of cooling rate effects, TC and Tm areplotted against cooling rate for bulk, 400 nm and 35 nm nano-structures. With increasing cooling rate TC increases and Tm de-creases for bulk and 400 nm nanostructures, but for 35 nmnanostructures TC and Tm are basically unaffected. During cooling,due to the high mobility of the chains in the paraelectric phase, theannealing effects in the paraelectric phase will dominate. The

reduced Gibbs free energy of the paraelectric phase intersects withthe free energy of the ferroelectric phase at lower temperature andthe free energy of the melt at higher temperature, causing adecrease in TC and an increase in Tm with decreasing cooling rate.For 35 nm nanostructures however, the crystallization temperature(Tcry) is strongly reduced due to confinement and takes place in thesame temperature range as the paraelectric to ferroelectric phasetransition. Therefore no extended annealing in the paraelectricphase takes place. In addition a possible increase in crystal size islimited by confinement.

3.3. Texture of P(VDF-TrFE) nanostructures

It was shown before that crystallization within the 1Dconfinement of nanotubes leads to oriented crystal growth result-ing in macroscopic alignment of the crystal growth direction par-allel to the axis of the tube [18]. For PVDF-TrFE the question ariseshow this scenario is modified by the additional phase transitionfrom the paraelectric to ferroelectric state. We addressed thisquestion by temperature dependent X-ray diffraction experiments.

Figs. 8 and 9 show the q#2q pattern of bulk PVDF-TrFE and of400 nm nanostructures at selected temperatures measured duringcooling. Upon cooling PVDF-TrFE crystallizes in the paraelectricstate giving rise to a Bragg reflection at 2q ¼ 17.1", corresponding tothe (110/200) lattice planes. The measurement on the bulk sampleshows that in the ferroelectric phase this peak shifts to higherangles and two additional reflections show up at 2q ¼ 35.1" and2q ¼ 40.8". All reflections could be indexed based on the knownorthorhombic unit cell of ferroelectric P(VDF-TrFE) having a molarVDF:TrFE ratio of 80:20 (Table 2) [51]. The broadening of the re-flections in the ferroelectric phase is at least partially caused by thefact that the crystal lattice looses its hexagonal symmetry, the (110)and the (200) split up into two different reflections at slightlydifferent angles. Higher order reflections are suppressed in theparaelectric phase due to thermal disorder. While crystallizationand the transition to the paraelectric phase take place in the sametemperature range as in the bulk sample the additional Bragg re-flections are absent in the nanostructured sample suggesting

Fig. 7. a) Melting temperature, Tm, and Curie temperature, TC, for bulk samples,400 nm nanostructures and 35 nm nanostructures measured during heating aftercooling with different scan rates; b) free energy diagram of the ferroelectric, para-electric and melt phases of P(VDF-TrFE). The dashed line shows the shift due toannealing in the paraelectric phase.

Fig. 8. q#2q scan of bulk PVDF-TrFE during cooling. The inset shows the unit cell ofP(VDF-TrFE) in the ferroelectric phase.

N. Shingne et al. / Polymer 54 (2013) 2737e27442742

orientation. This assumption could be proven by performing Schulzscans on both samples [34]. Fig. 10 shows Schulz scans (cf. MethodsSection) of the (110/200) reflections of 400 nm nanostructures,measured at 110 "C and 25 "C, respectively. In the paraelectricphase, the narrow peak in the Schulz scan with a FWHM of 9.25"

indicates strong alignment of the (110/200) planes parallel to thetemplate surface. We attribute this orientation to the same mech-anism as it was observed for PVDF [18]. Only crystals with a latticeplane (110/200) parallel to the substrate can grow along the poreaxis over longer distances, the growth of other crystals is soonstopped as they impinge on the pore wall. This also means that thechains are lying parallel to the plane of the substrate, i.e. perpen-dicular to the pore axis, which also explains why the crystalthickness is affected by the 1D confinement (cf. DSC results). In thescan measured in the ferroelectric phase, the intensity peaks againat j ¼ 0", but the peak is much broader with FWHM of 39.9". Weinterpret this change in crystal orientation upon cooling as anindication for stresses caused by the deformation of the crystallattice at the transition from the paraelectric to ferroelectric state.They obviously destroy the integrity of the crystals and changetheir original orientation.

4. Conclusions

In conclusion our experiments showed that infiltration ofPVDF-TrFE into AAO nanopores can be used to prepare ferro-electric nanostructures with a controlled crystal orientation. Togain broader understanding we studied the processes occurringduring wetting, crystallization and the transition to the ferro-electric state. The morphology results from precursor film wet-ting followed by much slower thickening processes of the wallsof the nanotubes. Details of the thickening process could beclarified. Similarly as it was shown for PVDF, crystallization in theone dimensional confinement of nanopores leads to orientationresulting from selective growth along the direction of the poreaxis [18]. As observed before [25] ferroelectricity is retained inthe nanopores and the transition temperature from the para-electric to the ferroelectric phase is not much affected byconfinement. The high degree of alignment which is obtainedafter crystallization in the paraelectric phase is reduced at roomtemperature. We attribute this loss of orientation to the internalstresses occurring at the transition to the ferroelectric state dueto the lattice deformation. Nevertheless the samples at roomtemperature show a dominant orientation of the chainsperpendicular to the axis of the pore. As a consequence also thethermodynamics of the nanostructures is affected by confine-ment. These effects are strong as the thickness of the crystallinelamellae is of comparable scale as the diameter of the nanoporeswhich affects the transitions occurring during heating, i.e. fromthe ferroelectric to the paraelectric state and from the latter tothe melt. As we could show, changes in the transition ferro-electric to paraelectric are related to annealing effects occurringdominantly in the highly mobile paraelectric phase. In PVDF-TrFEpolarization points along the b-axis, i.e. perpendicular to thechain axis which itself is perpendicular to the pore axis. Thismakes PVDF-TrFE nanotubes and nanorods potentially inter-esting objects for experiments in which polarization is manipu-lated on a microscopic level.

15 20 25 30 35 40

(110/200)

(110/200)

para(110/200)

Inte

ns

ity

18 20 22

para(110/200)

θ

Fig. 9. a) q#2q scan of 400 nm nanostructures during cooling. The inset shows the XRDpattern on a smaller scale around the (110/200) peaks.

Table 2Calculated positions of Bragg reflections for 80:20 mol% VDF:TrFE [51], measuredpositions for 77:23 mol% VDF:TrFE.

Lattice plane q (calculated) 2q (calculated) 2q (measured)

(110) 1.430 20.19 19.91(200) 1.412 19.93 19.91(310) 2.456 35.03 35.11(020) 2.488 35.50 35.11(001) 2.464 35.15 35.11(220) 2.860 41.03 40.8(400) 2.823 40.48 40.8(111) 2.849 40.869 40.8(201) 2.839 40.719 40.8

0 20 40 600.0

0.5

1.0

Re

la

tiv

e In

te

ns

ity

ΨFig. 10. Schulz scans of P(VDF-TrFE) nanostructures in AAO with Dp of 400 nm, pre-pared at Tw ¼ 200 "C and tw ¼ 10 min. The solid curve is the Schulz scan belonging tothe (110/200) reflection of the paraelectric high-temperature phase measured at110 "C; the dotted curve is the Schulz scan belonging to the (110/200) reflection of theferroelectric low-temperature phase measured at 25 "C. Both curves are normalized tounity at j ¼ 0" .

N. Shingne et al. / Polymer 54 (2013) 2737e2744 2743

Acknowledgments

The authors thank K. Sklarek and S. Grimm for the preparationof AAO membranes, B. Pulamgatta for GPC measurements, and Y.Luo for preliminary X-ray investigations. Financial support by theResearch Network Saxony-Anhalt “Nanostructured Materials” andthe German Research Foundation (STE 1127/13, INST 190/134-1, SFBTRR 102) is gratefully acknowledged.

Appendix A. Supplementary data

Supplementary data related to this article can be found at http://dx.doi.org/10.1016/j.polymer.2013.03.034.

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