Microstructure and ferroic properties of epitaxial†�-Fe2O3–BiFeO3‡−Bi3.25La0.75Ti3O12 composite bilayers

O. Gautreau,1 C. Harnagea,1,a� L. Gunawan,2 G. A. Botton,2 L. Pintilie,3 M. P. Singh,4 andA. Pignolet11INRS-Énergie, Matériaux et Télécommunications, 1650 Boulevard Lionel-Boulet, Varennes,Quebec J3X 1S2, Canada2Department of Materials Science and Engineering and Brockhouse Institute for Materials Research,McMaster University, 1280 Main Street West, Hamilton, Ontario L8S 4M1, Canada3National Institute of Materials Physics, P.O. Box MG.7, Magurelle 077125, Romania4Département de Physique, Regroupement Québécois sur les Matériaux de Pointe, Université deSherbrooke, Sherbrooke, Quebec J1K 2R1, Canada

�Received 29 June 2010; accepted 15 October 2010; published online 10 December 2010�

Epitaxial ��-Fe2O3–BiFeO3� /Bi3.25La0.75Ti3O12 and Bi3.25La0.75Ti3O12 / ��-Fe2O3–BiFeO3�composite bilayers were grown on SrRuO3 coated �111� SrTiO3 substrates in order to investigate theinfluence of the morphology of the �-Fe2O3–BiFeO3 self assembled nanocomposite layer on themultiferroic properties of the bilayer. Both types of bilayers exhibit high resistivity andsimultaneously ferroelectricity and ferrimagnetism at room temperature. When the�-Fe2O3–BiFeO3 composite layer is sandwiched between the Bi3.25La0.75Ti3O12 film and thesubstrate, the BiFeO3 component is not only subjected to epitaxial strain induced by the surface ontop of which it grows but also to elastic interactions with the Bi3.25La0.75Ti3O12 capping layer. Thelatter indeed reduce the amount of �-Fe2O3 inclusions, affects the morphology of the grains in the�-Fe2O3–BiFeO3 layer, and increases the shape anisotropy of the �-Fe2O3 inclusions. Additionally,this modification in the microstructure of the �-Fe2O3–BiFeO3 layer induces an imprint in theferroelectric hysteresis loop as well as a decrease in the saturation magnetization, and its magneticeasy axis direction changes from in-plane to out-of plane. © 2010 American Institute of Physics.�doi:10.1063/1.3514591�

I. INTRODUCTION

Multiferroic materials which exhibit both magnetic andferroelectric properties have increasingly become a subjectof investigation both because of their fundamental interestand for their potential in technological applications. Thepresence of well controlled ferroic properties in the materialsopen the possibility of their use in novel multifunctionaldevices,1–4 as well as understanding the genesis of multifer-roic coupling in oxides.

In this context, BiFeO3 �BFO� is one of the most inves-tigated multiferroics since it possesses high magnetic andferroelectric transition temperatures. Unlike in the bulkphase, where the helicoidal arrangement of the spins resultsin a zero average macroscopic magnetization, BFO in thinfilm form also exhibits both magnetic and ferroelectric orderparameters at room temperature. It exhibits a large value ofthe remanent ferroelectric polarization and a small value ofsaturation magnetization that arises from the spin cantingeffect. The presence of these unique properties in BFO atroom temperature prompts it to be a potential candidate indata storage media and multiple state memories.5–7 However,its use in real devices is hindered due to various issues notyet fully resolved. The most important are the leakage cur-rent arising from the presence of Fe2+ ions and oxygen va-cancies VO

2+ contributing to conduction, the high electric

coercive field, the low resistance to ferroelectric fatigue, aswell as the relatively low saturation magnetization.8–11

Despite extensive studies, a good control of thesetechnologically-important properties, in particular how notonly to control but to increase its magnetization, has beenelusive in BFO. In this context, we developed a unique wayto address some of these issues. Using a multilayer/composite approach, we recently demonstrated that the ferro-electric fatigue is consderably improved, the magnetizationis increased, and the leakage current is drastically reduced byusing an appropriate epitaxial buffer layer between the BFOlayer and the substrate.12 This was achieved by �i� depositingan epitaxial ferroelectric Bi3.25La0.75Ti3O12 �BLT� layer toprevent space charge migration from the film toward the bot-tom electrode, enhancing both the insulating and fatigueproperties while keeping an acceptable ferroelectric polariza-tion of the bilayer system, and �ii� choosing the appropriatedeposition conditions of BFO to form a self assembled�-Fe2O3–BiFeO3 epitaxial nanocomposite with �-Fe2O3

�FO� providing a high saturation magnetization. The self-assembled, all-epitaxial nanocomposite layer is the result ofa phase transformation of BFO under the epitaxial constrainimposed by both the underlying epitaxial BLT layer and thesubstrate.13 Due to elastic interactions between the BFO andBLT layers of the heterostructure, BFO transforms into alter-nating columnar grains of �-Fe2O3 and BiFeO3 phases.12,14

The epitaxial grains of �-Fe2O3 develop to reduce the effectof the elastic strain field in BFO and minimize the elastica�Electronic mail: harnagea@emt.inrs.ca.

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energy as described by the thermodynamic theory of elasticdomains.13 Such a composite is more stable than a singlephase BFO layer deposited epitaxially on BLT and its forma-tion is possible due to its growth at high temperature andunder high external stress. We further refer to this nanocom-posite heterostructure as BFO–FO/BLT. Recently, we dem-onstrated that the ferroelectric domain structure of theBFO-FO composite layer is ascribed to the BFO phase whilethe magnetic domain pattern can be attributed to the �-Fe2O3

phase, confirming the coexistence of ferroelectricity and fer-rimagnetism in the heterostructured composite at thenanoscale.14 Moreover, our studies underline the direct cor-relation between the nanostructure and the multiferroic prop-erties of the self-assembled epitaxial BFO–FO nanocompos-ite in our BFO–FO/BLT bilayers.

In such epitaxial heterostructured bilayers, the improve-ment in insulating properties while keeping a good ferroelec-tric behavior is governed by the presence of the BLT layer,whereas the magnetic properties are enhanced and mostlygoverned by the FO nanoscale inclusions.

To gain a deeper insight in the mechanisms that promotethe desired ferroic properties in these heterostructures, westudied more in detail the influence of epitaxial strain on themultiferroic properties of these unique bilayer heterostruc-tures. In particular, we compared the structural and physicalproperties of bilayers with different growth sequences, viz.,BLT/BFO–FO and BFO–FO/BLT. For reference, we alsogrew pure BFO and BLT films on identical substrates. In thispaper, we describe and discuss the impact of crystallinity,nanostructure, and morphology on the physical properties ofboth BFO–FO/BLT and BLT/BFO–FO bilayers.

II. EXPERIMENTAL

The composite bilayered thin film heterostructures:BFO–FO/BLT and BLT/BFO–FO were grown by pulsed la-ser deposition �PLD� on �111� oriented �40 nm� SrRuO3

�SRO� coated �111� SrTiO3 �STO� substrates under opti-mized growth conditions. In the following we use the term“substrate” for the fully strained epitaxial �40 nm� SRO elec-trically conductive film on the �111� oriented STO singlecrystal. The detailed deposition conditions have been re-ported elsewhere.12,15 The thicknesses of both BFO–FO andBLT layers were about 200 nm and were estimated using aDektak 3030 profilometer and confirmed by cross-sectiontransmission electron microscopy �XTEM�. The crystallo-graphic structure and orientation of the various phases werecharacterized by x-ray diffraction �XRD� �-2� and pole fig-ure measurements �PANalytical X’Pert MRD four-circle dif-fractometer using Cu K� radiation�. The sawtoothlike andcolumnar microstructure of grains within both BLT andBFO–FO films are visualized in XTEM bright field �BF�images. The BF images and corresponding selected areaelectron diffraction �SAED� patterns of both the bilayers andthe substrate were acquired by using a Philips CM12 TEM120 kV. The information on elemental distribution within thebilayers were recorded by energy dispersive x-ray spectros-copy �EDX� elemental mapping �INCA Microanalysissystem-Oxford Instrument� by using a JEOL 2010F TEM/

STEM 200kV. The high resolution low-angle annular darkfield scanning �ADF� STEM images of bilayers were re-corded by using FEI-Titan 80-300 “cubed” 300 kV withoutprobe corrector. For TEM sample preparation, the materialwas mechanically polished almost to electron transparencyby the low-angle wedge polishing method and the ion mill-ing was only used at low accelerating voltage for the finalstages �200 eV� for a few minutes to clean up the samplefrom the contamination and remove all possible damage inthe prior steps of preparation. To perform electrical charac-terization, Pt top electrodes were deposited by PLD througha metallic shadow mask. Macroscopic ferroelectric proper-ties were measured at 2 kHz and at room temperature using aferroelectric tester �TFA 2000, AixACCT Systems GmbH�.Before measuring the leakage current characteristics, theferroelectric polarization was saturated by applying a dcelectric field higher than the coercive field �except for theBLT/BFO–FO bilayers since the maximum voltage that canbe applied is smaller than the coercive field of the BLT/BFO–FO bilayers� and with the same polarity as that usedfor current measurements. In this way the polarization value,as well as its orientation, are known and the contribution ofthe current due to polarization reversal is minimized. Localout-of-plane converse piezoelectric coefficient dzz versus ap-plied electric field E hysteresis loops were performed usingpiezoresponse force microscopy �PFM�.16,17 In this study weused a DI-Enviroscope AFM �Veeco� equipped with aNSC36a �Mikromasch� cantilever and tips coated with Co/Cr. We applied an ac voltage of 0.5 V at 26 kHz between theconductive tip and the SRO layer and we detected the filmsurface induced piezoelectric vibrations using a lock-in am-plifier from Signal Recovery �model 7265�. Local hysteresismeasurements were achieved using a dc-source �Keithley2400� to apply bias voltages to the bottom electrode of thesample. Temperature dependence magnetization and mag-netic field dependence loops �i.e., M-H loops� were mea-sured using a superconducting quantum interference device�SQUID� magnetometer from Quantum Design. M-H loopsare measured in both in-plane and out-of plane directions.

III. RESULTS AND DISCUSSION

A. Crystallographic orientations and epitaxialrelationships

The configurations of the epitaxial heterostructures ana-lyzed in this paper are schematically presented in Fig. 1.XRD �-2� scans of BFO and BLT single layers, as well asBFO–FO/BLT and BLT/BFO–FO bilayers are shown in Fig.2. One should note that SRO is epitaxial and thin enough tobe fully strained. SRO and STO therefore have the samelattice parameters and cannot be distinguished on the XRD�-2� scans. For all the phases of our heterostructures, onlythe expected reflections corresponding to planes parallel tothe �111�-oriented substrate, namely BFO 111, FO 111, andBLT 104,12 are observed, indicating that no significantamount of extra phase, if any, is present and reveals the highdegree of crystallographic orientation of the heterostructures.

The interplanar distance between BFO �111� planes,dBFO�111�, the crystallite size of the BFO and FO phases,

114111-2 Gautreau et al. J. Appl. Phys. 108, 114111 �2010�

DBFO and DFO, as well as the amount of FO phase,AFO111 /ASTO111, have been estimated from XRD spectra �Fig.2� and are presented in Table I. The dBFO�111� values mea-sured indicate that the BFO phase is subjected to a signifi-cant in-plane compressive strain in BFO single layers onSTO�111�, but submitted to an in-plane tensile strain in bothBFO–FO/BLT and BLT/BFO–FO bilayers. Moreover, thecrystallite size values together with the estimated amount ofthe FO phase suggest that the amount of FO increases whenthe crystal quality of the epitaxial BFO phase degrades.These observations are consistent with our previousstudies12,18 and suggest that epitaxial FO inclusions developat the BFO grain boundaries in our bilayers while BFO re-laxes.

Pole figure analyses were performed to establish the in-plane crystallographic orientations of the different epitaxialphases of the bilayers. Figure 3 shows the pole figure of aBFO–FO/BLT bilayer on SRO-coated STO �111� substraterecorded at different 2� angles of 31.8°, 30.1°, and 43.3°,corresponding, respectively, to the BFO 110, BLT 117, andFO 400 reflections. For the BFO phase, the three diffractionpeaks every 120° at ��36° �Fig. 3�a�� corresponding to theangle 35.7° between the BFO �111� and BFO �110� planes,

indicate the expected threefold symmetry of epitaxial BFOfilms on SRO/STO�111� and verify the following cube oncube epitaxial relationship:

BiFeO3�111� SrTiO3�111�;BiFeO3�110� SrTiO3�110� .

�1�

For the BLT layer, the peaks appearing at ��36° and 84°

�Fig. 3�b�� correspond, respectively, to the 117/1̄17 and

1̄1̄7/11̄7 reflections of BLT, and indicate that the BLT layerhas the triple-twin-situation expected for epitaxial BLT filmson SRO/STO�111�,19 with the corresponding epitaxial rela-tionship as follows:

BLT�104� SrTiO3�111�;BLT�010� SrTiO3�11̄0� . �2�

For the FO phase, the three diffraction peaks every 120° at��54° �Fig. 3�c��, corresponding to the angle 54.7° betweenthe FO �111� and FO �400� planes, and observed at 180°from the corresponding BFO peaks in Fig. 3�a�, indicate thatthe �-Fe2O3 phase orientation with respect to the SrTiO3

substrate verify the following epitaxial relationship:

�-Fe2O3�111� SrTiO3�111�;

�-Fe2O3�1̄10� SrTiO3�1̄01� . �3�

B. Microstructure

SAED Pattern investigations confirmed the epitaxialrelationships estimated from XRD measurements

TABLE I. Interplanar distance between BFO �111� planes dBFO�111�, BFO and FO crystallites size �resp. DBFO,DFO�, area of the FO 111 reflection AFO111 normalized with respect to the area of the STO 111 reflection ASTO111

�AFO111 /ASTO111 gives an estimation of the relative amount of FO�, and interplanar distance between BLT �104�planes dBLT�104�, for BFO single layers, BFO–FO/BLT and BLT/BFO–FO bilayers, estimated from Fig. 2. Alllayers and multilayers are deposited on SRO-coated STO�111�.

HeterostructuredBFO�111� �bulk: 2.310�

��DBFO

�nm�DFO

�nm�AFO111 /ASTO111

�%�dBLT�104� �bulk: 4.519�

��

BFO 2.32�0.01 38�4 ¯ ¯ ¯

BFO–FO/BLT 2.31�0.01 33�4 38�8 �7.6�1.5��10−2 4.49�0.01BLT/BFO–FO 2.31�0.01 37�4 22�5 �1.5�0.3��10−2 4.50�0.01

FIG. 2. XRD patterns of BFO and BLT single layers and BFO–FO/BLT andBLT/BFO–FO bilayers. �, �, �, and � stand for SrTiO3 W L�, BiFeO3

W L�, SrTiO3 Cu K�, and BiFeO3 Cu K� reflections, respectively.

FIG. 1. �Color online� Schematic representation of the epitaxial heterostruc-tures: �a� BFO and �b� BLT single layers epitaxially grown on �111�-orientedSrTiO3 substrates coated with �40 nm� SrRuO3 bottom electrode; �c� BFO–FO/BLT and �d� BLT/BFO-FO bilayers in which the BFO-FO layer is theself assembled nanocomposite composed of BiFeO3 and �-Fe2O3 phases.

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�Figs. 2 and 3� and determined the local crystal structure ofthe various phases in the BFO–FO layer in more details. Thecross-section image-taken in BF image mode of the bilayersinvestigated, is presented in Figs. 4�a�–4�c� show SAEDpatterns of the substrate and of the BFO–FO layer, respec-tively, recorded along the STO �121� direction. The sharpdiffraction spots indicate the good crystalline quality of allphases of the bilayer. The indexes of the SAED pattern ofthe BFO–FO layer �Fig. 4�d�� show a good agreement with

the epitaxial relationships �1�–�3� established from XRDdata.

The crystal structure parameters of the BFO phase �i.e.,lattice parameter aBFO and angle between lattice vectors�BFO� have been estimated from the interatomic spacing val-ues dhkl extracted from the SAED pattern �Fig. 4�c��. Assum-ing that BFO has a rhombohedral crystal structure and usingthe expression of interatomic spacing in a rhombohedralsystem

1

dhkl2 =

�1 + cos ���h2 + k2 + l2� − �1 − tan2�

2��hk + kl + hl�

a2�1 + cos � − 2 cos2 ��, �4�

we obtain aBFO�3.96�0.01 Å and �BFO�89.2�0.1°.These values are very close to those of rhombohedral bulkBFO �i.e., aBFO=3.96 Å and �BFO=89.4°� and thus indicatethat the BFO phase in our bilayers is almost, if not totally,relaxed. For the maghemite �FO� phase, using the dhkl valuesextracted from the SAED pattern �Fig. 4�c��, we obtain thebest fit for the crystal structure parameters assuming a tetrag-onal cell with P4/mnm or �P4 /nnm space group, and latticeparameters aFO=7.87 Å, and cFO=8.07 Å, using the expres-sion of interatomic spacing in a tetragonal system

1

dhkl2 =

h2 + k2

a2 +l2

c2 . �5�

Thus, the crystal structure of the maghemite phase in ourBFO–FO layers is different from that of the bulk material

�cubic unit cell: a=8.3515 Å, space group Fd3̄m�,20 never-theless, a similar structural difference has already been ob-served in epitaxial maghemite thin films.21 The interplanarspacing values obtained suggest that the maghemite phase issubjected to both in-plane and out-of plane compressivestrain. Such out-of plane compressive strain of themaghemite phase in the bilayer is too large to be explainedby an in-plane epitaxial strain state caused by the underlyingBLT layer or STO substrate. However, the FO strain state isin well-agreement with a lateral epitaxial relationship be-tween the epitaxial maghemite grains and the BFO matrix.

These results suggest that strong elastic interactions exist atthe lateral interfaces between the BFO matrix and the epitax-ial maghemite inclusions in the BFO-FO layer of our bilay-ers. This is consistent with the conclusions discussed by Tan,Slutsker and Roytburd about epitaxial self-assembly ofperovskite-spinel nanostructures:22 elastic interactions be-tween phases along the vertical interfaces determine thestress state in the film, since the substrate induced stress isrelaxed for film thicker than 100 nm. In addition, the pos-sible presence of a significant number of vacancies inmaghemite, that would reduce the lattice parameters, cannotbe ruled out and could contribute to a further decrease involume of the maghemite unit cell. As can be seen in Table I,the interplanar distance between BLT �104� planes dBLT�104�,

indicates that - contrary to the almost relaxed BFOphase - the BLT phase is subjected to a significant in-planetensile strain in both BFO–FO/BLT and BLT/BFO–FO bi-layer configurations �a /abulk−1.3%�. The same mismatchis found between bulk BLT and fully strained SRO films onbulk STO �in the case of BFO–FO/BLT�, as well as betweenrelaxed �or bulk� BFO and strained FO �in case of BLT/BFO–FO�. Additionally, a careful analysis of the SAED pat-tern of the BFO–FO layer �Fig. 4�d�� reveals the presence of�-Fe2O3 together with �-Fe2O3. The �-Fe2O3 phase is alsoconfirmed to be oriented with respect to the STO substratewith the following relationship:

FIG. 3. �Color online� XRD pole figures �center: �=0°; rim: �=90°� of a BFO-FO/BLT bilayer recordedat different 2� angles of �a� 31.8°, �b� 30.1°, and �c�43.3° corresponding, respectively, to the BFO �110�,BLT �117�, and FO �400� reflections.

114111-4 Gautreau et al. J. Appl. Phys. 108, 114111 �2010�

�-Fe2O3�001� SrTiO3�111�;

�-Fe2O3�1̄20� SrTiO3�1̄01� . �6�

Since �-Fe2O3 is well crystallized yet undetected duringXRD measurements, we assume that the amount �-Fe2O3

within the BFO-FO layer is minute.Moreover, the difference between the values of hematite

crystal structure parameters in our bilayers �estimated fromFig. 4�c�� and bulk hematite crystal structure parameters, isless than 1%. Hence, we are assuming that the minuteamount of hematite phase present in our bilayers is relaxedand has the corresponding bulk crystal structure i.e., rhom-

bohedral �space group R3̄c� with a�-FO=5.03 Å, c�-FO

=13.75 Å, ��-FO=90°, and ��-FO=120°.The morphologies of the bilayers: grains shape and size,

and mutual arrangement of the different phases, have beenfurther probed by low-angle ADF-STEM and STEM-EDXelement mapping for Fe and Bi, and are presented in Fig. 5for the BLT/BFO–FO bilayers. Similar to the BFO–FO/BLTbilayers, already reported in Ref. 12, combining the resultsfrom the Fe and Bi chemical mappings �respectively Figs.5�c� and 5�d��, two distinct regions can be noticed: �i� Fe-richand Bi-depleted regions �ascribed to be the FO phase� and�ii� regions containing both Fe and Bi �ascribed to the BFOphase�. It can be clearly seen �as for BFO–FO/BLTbilayers�,12 that the FO phase develops within the BFO–FOlayer at BFO grain boundaries, confirming our previous re-sults discussed in Section III A. For both configurations ofbilayers, the nanocomposite layer exhibits a columnar grainstructure in cross sectional images. Similar to perovskite-spinel self assembled nanostructures on �111� SrTiO3,22

BFO–FO layer have the spinel �-Fe2O3 inclusions wettingthe underlayer and the perovskite BiFeO3 matrix grains as

inverted cone-shaped inclusions. In the case of BFO–FO/BLT bilayers, the pyramidal shape observed for BFO grainsis consistent with the morphology of perovskite phase inself-assembled perovskite-spinel nanostructures formed on�111� SrTiO3.22 Such morphology is determined by the low-est surface energy for the considered materials in this spe-cific orientation.

An important difference between the BFO–FO/BLT andBLT/BFO–FO bilayers is the surface roughness. The surfaceof the BFO–FO layer in BFO-FO/BLT bilayers is relativelyrough �root mean square roughness=18 nm�,12 and we at-tributed this to be due to elastic stress induced by the highlattice mismatch between BFO and BLT and between BFOand FO, as well as the triple-twin-situation of BLT.

In BLT/BFO–FO bilayers, the top of the BFO grains atthe interface between the two layers is flat. We attribute thesmoothing of the top of the BFO grains to �i� reduced latticemismatch between BFO and the material underneath �SRO/STO� and �ii� to the effect of BLT above it, which here playsthe role of a thick capping layer.23 The presence of a BLT caplayer on top of a BFO–FO slightly strained epitaxial film hastwo direct effects on the morphological stability of theBFO–FO layer underneath. First, it suppresses the masstransport on the otherwise free surface of the film. Second,the mechanical stiffness of the cap layer which increaseswith its increasing thickness, tends to stabilize the epitaxialfilm underneath �here BFO�. Furthermore, the cap layer itself�BLT� may effectively be stiffened against roughening whensubjected to a tensile residual stress,23 which is exactly hap-pening in our case due to the large amplitude of the latticemismatch between BLT and BFO �and FO�. The epitaxialBFO–FO layer is, therefore, effectively stabilized. In fact thecap layer suppresses the kinetic process of roughening, aprocess that takes place while still at high temperature, rightafter the deposition. Additionally, most of the FO might notdevelop fully during the growth of the BFO layer but onlyafter the deposition of the BLT layer, and even during the 1 h

FIG. 4. �Color online� BF-TEM image and selected area diffraction patterns�SADP� of a BFO–FO/BLT bilayer. �a� Cross-section image showing thearea selected �dashed circle� for the analysis of the phases of the BFO–FOlayer-orientations shown are STO substrate ones. SAD patterns taken alongthe STO �121� direction of �b� the STO substrate and �c� of the BFO–FOlayer. �d� Schematic of the BFO–FO phase diffraction pattern correspondingto �c�.

FIG. 5. �Color online� Low-angle ADF-TEM image �a�, TEM-EDX elementmappings for Fe �c�, and Bi �d�, of a BLT/BFO–FO bilayer cross-section. �b�Schematic representation of the phase’s localization in the BLT/BFO–FObilayer cross-section.

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in situ post deposition annealing at 400 °C.12 After the depo-sition of the BLT layer, elastic interactions may occur at theBLT/BFO interface, as indicated by the smoothing process ofthe top surface of the BFO grains. Due to the high latticemismatch between BFO and BLT, the FO phase may developwithin the BFO layer in a similar way than in BFO–FO/BLTbilayers, as a result of phase transformation in a constrainedepitaxial layer.13 However, in contrast to BFO–FO/BLTbilayers,12 there is a lesser amount of FO in the BFO–FOlayer of the BLT/BFO-FO bilayers. We attribute the reducedamount of FO phase to fewer BFO grain boundaries in theBFO–FO layers of BLT/BFO–FO compared to BFO–FO/BLT bilayers. The extent of the grain boundaries created dur-ing BFO film growth is governed by the surface roughnessand lattice mismatch with the underlying substrate. There-fore, since the SRO surface is much smoother than that of aBLT film, a lower amount of FO phase would form in theBFO film in the BLT/BFO–FO configuration than in theBFO–FO/BLT one�.12,24 Additionally, the lattice mismatchbetween BFO and SRO/STO is much smaller than that be-tween BFO and BLT, further reducing the number of grainboundaries.

Finally, by depositing BLT on top of BFO, evaporationof the volatile Bi from BFO,25 is considerably reduced, lim-iting if not completely stopping the transformation into theFO phase. Consequently, the lateral dimension of the FOgrains in the BFO–FO layer is smaller when the BLT isgrown on top of it, resulting in an aspect ratio �height/width)of approximately 5 �with major consequences on magnetiza-tion, as will be discussed in Sec. III D�, significantly higherthan that of the FO inclusions in the BFO–FO/BLT bilayers�approximately 1.6�.12

C. Ferroelectric properties

The ferroelectric hysteresis loops �polarization versuselectric field� for all types of heterostructures recorded fordifferent amplitudes of the maximum electric field are shownin Figs. 6�a�, 6�c�, 6�e�, and 6�g�. The polarization switchingcharacteristics, remanent polarization Pr and coercive fieldEc, extracted from the hysteresis loops measured are pre-sented in Figs. 6�b�, 6�d�, 6�f�, and 6�h�. For the BFO singlelayers we found distorted hysteresis loops with high coerciv-ity �2Ec=Ec+−Ec−�700 kV cm−1 for a maximum appliedelectric field Emax=610 kV cm−1� �Figs. 6�a� and 6�b��caused by the high leakage current contribution, as indicatedby the current versus electric field loops presented in Fig.7�a�. Nevertheless, the good ferroelectric properties of BFOare clearly demonstrated by its high remanent polarizationvalue, Pr+− Pr−�210 �C cm−2. In contrast, the ferroelectrichysteresis of BLT single layers �Figs. 6�c� and 6�d��, is welldefined with very low coercivity �2Ec−430 kV cm−1 forEmax�700 kV cm−1�, and a remanent polarization Pr+− Pr−

=20 �C cm−2 for Emax�700 kV cm−1. The leakage currentcontribution is not distinguishable even in the current versuselectric field loops �Fig. 7�b��, confirming the high insulatingpower of BLT films.12,18

The combination of BFO and BLT films in heterostruc-tures, however, results in different ferroelectric characteris-

tics: The BFO–FO/BLT bilayers exhibit �Figs. 6�e� and 6�f��a relatively high remanent polarization �Pr+− Pr−

�40 �C cm−2 at Emax=560 kV cm−1� and a moderate coer-civity �Ec+−Ec−395 kV cm−1 at the same Emax�. The wellsaturated character of the hysteresis loops together with aweak leakage current contribution �Fig. 7�c�� indubitablyconfirm the drastic improvement in the properties of BFO–FO/BLT bilayers compared to BFO single layers. The slightasymmetry of the hysteresis loops, identified in the coercivefield versus maximum applied field plot �Fig. 6�f�� can beexplained by the different nature of the top and bottom elec-trodes.

The ferroelectric behavior of BLT/BFO-FO bilayers issomewhat different: the amplitude of the remanent polariza-tion �Pr+− Pr−12 �C cm−2 for Emax�700 kV cm−1� isdrastically reduced and a pronounced imprint effect occurs inthe measured hysteresis loops, as indicated by the shift to-ward positive electric field values. Although the leakage cur-rent contribution is barely detected from current versus elec-tric field loops �Fig. 7�d�� and the coercivity is moderate�Ec+−Ec−440 kV cm−1 for Emax�700 kV cm−1�, the po-larization switching �Figs. 6�g� and 6�h�� is significantly dif-ferent. This degradation in polarization switching character-istics may be explained by a combination of the following:26

�i� stress present at both top and bottom interfaces of theBFO–FO layer, �ii� complex grain shape of the BFO grainsthat may cause domain pinning via intricate field distribu-tions at sharp edges between facets, and/or �iii� existence ofgraded surface layers �gradual change in composition, struc-ture and/or functional properties at the interfaces, differentfrom that of the bulk� likely to occur in the strained BFO–FOlayer of the BLT/BFO–FO heterostructure enduring asym-metric strains at its top and bottom interfaces. In addition,the electrode interfaces may play a role in the asymmetryobserved as well as in the difference in the ferroelectric prop-erties of the different bilayer configurations �in one case thetop interface is Pt/BLT and the bottom one is BFO–FO/SRO,and in the other case the top interface is Pt/BFO–FO and thebottom interface is BLT/SRO�.

To complete the electric characterization, we present inFigs. 8�a� and 8�b� the electric field dependence of the di-electric permittivity calculated from the capacitance mea-sured at 2 kHz and at room temperature, as well as the resultsof the leakage current tests �Fig. 8�c��. The permittivities inzero bias for single layers are comparable to the values foundin the literature, viz., 160 for BLT and 35 for BFO. Thedielectric permittivity values of the bilayers are betweenthose of BFO and BLT, consistent with a multilayered ca-pacitor model. For BLT single layers and BFO-FO/BLT bi-layers, no significant asymmetry is observed in contrast toBFO single layers and BLT/BFO–FO bilayers. The asymme-try suggests an important influence of the electrode inter-faces on the measured capacitances, most probably domi-nated by the BFO/SRO interface formed at high temperature.

The same tendency is observed for the leakage currentdensity versus electric field characteristics in Fig. 8�c�. Thedependencies suggest that the leakage current is bulk limitedin BLT single layers and BFO-FO/BLT bilayers while it isinterface limited in BFO single layers and BLT/BFO–FO

114111-6 Gautreau et al. J. Appl. Phys. 108, 114111 �2010�

bilayers.27 Again the results suggest that the limiting inter-face in the latter case is the BFO/SRO interface. The behav-ior of BFO single layers is consistent with the recent studyon leakage current in epitaxial Pt/BFO/SRO capacitors re-ported by Pintilie et al.,28 where leakage current has beenshown to be interface limited.

It is clear from the magnitude of the leakage current overthe whole range of electric field presented in Fig. 8�c�, thatBFO films have the highest leakage current density, whileBLT films as well as the BLT-based bilayers have the lowest.This confirms the assumption that the leakage current contri-bution to the switching current was minimal for BLT andboth bilayers as discussed previously. Analyzing the effect ofthe order of bilayers, we observe a higher leakage currentdensity of BFO–FO/BLT, which we attribute to an increaseddensity of BFO grain boundaries �according to DBFO valuesTable I�, reducing the overall resistivity of the film.

The local ferroelectric properties were characterized atthe nanoscale as well, using PFM, both imaging the ferro-electric domain structure and recording local piezoelectrichysteresis loops. The latter are shown in Fig. 9 for all typesof heterostructures. A strong correlation between macro-scopic characterization �Fig. 6� and characterization at thenanoscale can easily be observed: The BFO single layers doexhibit both high remanence and high coercivity �dzzr+

−dzzr−=20 pm V−1 and Ec+−Ec−�1050 kV cm−1, for Emax

�1600 kV cm−1, respectively� �Fig. 9�a��, the BLT singlelayers exhibit low remanence and low coercivity �dzzr+

−dzzr−=3.5 pm V−1 and Ec+−Ec−=490 kV cm−1, for Emax

�2000 kV cm−1, respectively� �Fig. 9�b��, while the BFO–FO/BLT bilayers present significant remanence and reducedcoercivity �respectively, dzzr+−dzzr−=20 pm V−1 and Ec+

−Ec−=600 kV cm−1, for Emax=600 kV cm−1� �Fig. 9�c��and BLT/BFO–FO bilayers exhibit an imprint effect also vis-

FIG. 6. �Color online� Room temperature ferroelectrichysteresis loops recorded at 2 kHz ��a�, �c�, �e�, and �g��together with electrical-field dependences of remanentpolarization Pr and coercive field Ec ��b�, �d�, �f�, and�h�� for BFO ��a�,�b�� and BLT ��c� and �d�� single lay-ers and for BFO–FO/BLT ��e� and �f��, and BLT/BFO-FO ��g� and �h�� bilayers.

114111-7 Gautreau et al. J. Appl. Phys. 108, 114111 �2010�

ible on the piezoelectric hysteresis loops �Fig. 9�d��. Thisstriking similitude in the macroscopic and nanoscale ferro-electric behaviors strongly suggests that ferroelectricity isspatially uniform over several orders of magnitude and origi-nate in the behavior of nanoscale regions, not only for ourBFO–FO/BLT bilayers,14 but for all the heterostructures dis-cussed as well.

D. Magnetic properties

The magnetic properties of both types of bilayers �Fig.10� were characterized by well defined in-plane �solidsquares� and out-of plane �solid circle� M-H hysteresis loopsindicating that both type of films are ferromagnetic in nature.However, the value of saturation magnetization and the re-sponse of the magnetic domains under an applied magnetic

field differ significantly from each other. The BFO–FO layersdeposited on BLT exhibit a relatively high value of in-planemagnetization, larger than the magnetization measured per-pendicular to the film �maximum in-plane magnetizationMIP max=45 emu cm−3 and maximum out-of-plane magneti-zation MOP max=20 emu cm−3�. It is important to note thatthe maghemite - primary source of M-H loops in our films -is embedded in the BFO matrix in form of columnar grains.14

Shape anisotropy will therefore play a crucial role in deter-mining the magnetic behavior of our films. In case ofBFO–FO layers grown on BLT, the maghemite grains havean aspect ratio �height/width� of about 1.6, a relatively lowvalue implying a magnetic anisotropy identical to that of thinfilms, with the direction of the magnetic easy axis lying inthe plane of the film, well-consistent with the structural prop-erties of these films.

In contrast, for BLT/BFO–FO bilayers, the maghemitegrain’s aspect ratio is significantly higher �viz., height/width

FIG. 7. �Color online� Current vs electric field loops for BFO �a� and BLT�b� single layers and for BFO–FO/BLT �c� and BLT/BFO–FO �d� bilayers.

FIG. 8. �Color online� Electric field dependences of ��a�and �b�� the dielectric constant �recorded at 2 kHz� and�c� the leakage dc current density of BFO and BLTsingle layers ��a� and �c�� and BFO-FO/BLT and BLT/BFO-FO bilayers ��b� and �c�.

FIG. 9. �Color online� Out-of-plane converse piezoelectric coefficient dzz vsapplied electric field E hysteresis loops of BFO �a� and BLT �b� single layersand BFO–FO/BLT �c� and BLT/BFO–FO �d� bilayers.

114111-8 Gautreau et al. J. Appl. Phys. 108, 114111 �2010�

�5�, indicating that the shape anisotropy of the magneticcolumnar inclusions is most likely the dominating factorgoverning the magnetic properties of these films. In such aconfiguration, the alignment of the magnetic domains paral-lel to an out-of-plane magnetic field is easier, leading to aneasy axis of magnetization direction along the normal to thesubstrate. Indeed, the M-H loop measurements attest �Fig.10�b�� that the value of saturation magnetization along theout-of plane direction is larger than the in-plane magnetiza-tion �MOP max=24 emu cm−3 and MIP max=17 emu cm−3�.As expected, the saturation magnetization along the respec-tive easy axes of the bilayers is lower in BLT/BFO–FO thanin BFO–FO/BLT, because the amount of FO in the first caseis lower than in the second one, as previously mentioned. Aprecise estimation of the amount �volume ratio� of the mag-netic phases �i.e., BFO and FO� from the magnetic responses�M-H loops� was not possible. Indeed we found that contri-butions from magnetocrystalline anisotropy and magneticcoupling between BFO and FO phases exist in addition to acombination of ferrimagnetism �from maghemite, Ms

=420 emu cm−3� and weak-ferromagnetism �arising fromcanted antiferromagnetism BFO, Ms=15 emu cm−3�.12,14

Hence, simply using the saturation magnetization values foreach phase �BFO and maghemite� to deduce the fraction ofeach phase from the measured saturation magnetization willnot provide an accurate estimation of the quantities of thedifferent phases.

IV. CONCLUSIONS

In conclusion, �self assembled BFO-FO nanocomposite�-BLT bilayers have been grown on SRO-coated STO �111�substrates by PLD for two different layer growth sequences,i.e., BFO–FO/BLT and BLT/BFO-FO. All phases and layers

are epitaxially grown. The maghemite phase is developed atBFO grain boundaries while BFO phase relaxes, and is sta-bilized by out-of plane compressive strain imposed by lateralepitaxial relationship at the transverse BFO/FO interfaces.For BLT/BFO–FO bilayers, the development of the FO phaseis reduced compared to the case of the BFO–FO/BLT bilay-ers, since the amount of BFO phase is subjected to a com-plex epitaxial strain field at both top and bottom interfacesand throughout the thickness, and cannot fully relax throughthe whole thickness of the BFO–FO layer. This induces agraded microstructure of the BFO–FO layer. The shape an-isotropy of the FO grains is also affected since the FO gainsare more elongated vertically and is higher for BLT/BFO–FObilayers compared to BFO–FO/BLT bilayers. For both typeof bilayers the presence of the BLT layer results in a signifi-cant reduction in the leakage current compared to BFOsingle thin films. For BFO–FO/BLT bilayers, well saturatedsymmetric hysteresis loops with high remanent polarizationand reduced coercivity are obtained at room temperaturewhile for BLT/BFO–FO bilayers the distorted shape of theBFO-FO layer grains causes an imprint effect on the ferro-electric hysteresis. The ferroelectric behavior of each bilayeris conserved down to the nanoscale. The high shape aniso-tropy of the FO inclusions in the BFO–FO layer of BLT/BFO–FO bilayers results in an out-of-plane magnetic easyaxis direction compared to BFO–FO/BLT bilayers where thelow shape anisotropy of the FO inclusions gives rise to anin-plane easy axis direction. Furthermore, a consequence ofthe higher amount of FO inclusions within the BFO–FOlayer of the BFO–FO/BLT bilayers is a higher saturationmagnetization along its easy axis compared to that observedfor BLT/BFO–FO bilayers. The �self assembled BFO–FOnanocomposite�-BLT bilayers are therefore a heterostruc-tured system with good multiferroic properties at room tem-perature, very promising for applications in integrated de-vices.

ACKNOWLEDGMENTS

The authors would like to acknowledge F. Normandinand T. Veres �NRC-IMI� for magnetic measurements and dis-cussions, M. Alexe �MPI-Halle� for discussions about ferro-electric properties, as well as the Natural Science and Engi-neering Research Council �NSERC� of Canada,NanoQuébec, the National Research Council of Canada andINRS startup funds for financial support. Access to the Ca-nadian Centre for Electron Microscopy, a national facilityfunded by NSERC and McMaster University is gratefullyacknowledged.

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