Magnetic properties of BaTiO3/La0.7Sr0.3MnO3 thin films integrated on Si(100)Srinivasa Rao Singamaneni, Wu Fan, J. T. Prater, and J. Narayan Citation: Journal of Applied Physics 116, 224104 (2014); doi: 10.1063/1.4903322 View online: http://dx.doi.org/10.1063/1.4903322 View Table of Contents: http://scitation.aip.org/content/aip/journal/jap/116/22?ver=pdfcov Published by the AIP Publishing

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Magnetic properties of BaTiO3/La0.7Sr0.3MnO3 thin films integrated on Si(100)

Srinivasa Rao Singamaneni,1,2,a) Wu Fan,2 J. T. Prater,1,2 and J. Narayan2

1Materials Science Division, Army Research Office, Research Triangle Park, North Carolina 27709, USA2Department of Materials Science and Engineering, North Carolina State University, Raleigh, North Carolina27695, USA

(Received 24 September 2014; accepted 22 November 2014; published online 10 December 2014)

Two-phase multiferroic heterostructures composed of room-temperature ferroelectric BaTiO3 (BTO)

and ferromagnetic La0.7Sr0.3MnO3 (LSMO) epitaxial thin films were grown on technologically

important substrate Si (100). Bilayers of BTO/LSMO thin films display ferromagnetic Curie

transition temperatures of �350 K, close to the bulk value, which are independent of BTO films

thickness in the range of 25–100 nm. Discontinuous magnetization jumps associated with BTO

structural transitions were suppressed in M(T) curves, probably due to substrate clamping effect.

Interestingly, at cryogenic temperatures, the BTO/LSMO structure with BTO layer thickness

of 100 nm shows almost 2-fold higher magnetic coercive field, 3-fold reduction in saturation

magnetization, and improved squareness compared to the sample without BTO. We believe that the

strong in-plane spin pinning of the ferromagnetic layer induced by BTO layer at BTO/LSMO

interface could cause such changes in magnetic properties. This work forms a significant step forward

in the integration of two-phase multiferroic heterostructures for CMOS applications. VC 2014AIP Publishing LLC. [http://dx.doi.org/10.1063/1.4903322]

INTRODUCTION

Multiferroic materials offer the possibility of switching

the magnetization with an electric field and switching of

electric polarization with a magnetic field, which would rep-

resent a major advance for information storage and low

power non-volatile device applications. In that respect, there

have been numerous recent reports of ferromagnetic/ferro-

electric (FM/FE) heterostructures such as La0.7Sr0.3MnO3/

BaTiO3 (LSMO/BTO),1–3 La0.7Ca0.3MnO3/BTO (LCMO/

BTO),3,4 Fe/BTO,5 EuO/BTO,6 and Fe3O4/BTO7 hetero-

structures which show the coupling effect manifested as a

change in magnetization, magnetic anisotropy, and electrical

resistance due to strain- and interface-charge mediated cou-

pling. Very recently,8,9 Udalov et al. have discovered that

long range Coulomb interaction could explain the magneto-

electric coupling in composite granular multiferroics.

Among several materials, LSMO/BTO has emerged as the

prototypical system as both materials show strong ferroic

order parameters well above room temperature. Of particular

interest, LSMO deposited on BTO single crystals exhibited

large magneto-electric coupling due to strain associated

with BTO structural transitions as a function of tempera-

ture.1–3 For instance, Eerenstein et al.1 also demonstrated

electrically induced giant, sharp, and persistent magnetic

changes at the interface of LSMO/BTO. In another intrigu-

ing work,2 E-field induced magnetic modulation was

reported on BTO/LSMO heterostructure. BTO exhibits3

three distinct structural phase transitions: cubic-tetragonal

(C-T) at a temperature T¼ 393 K (also, ferroelectric

Curie point), tetragonal-orthorhombic (T-O) at T¼ 278 K,

and orthorhombic-rhombohedral (O-R) at T¼ 190 K,

respectively.

These materials possess a negligible lattice mismatch

which favours epitaxial growth of BTO/LSMO interfaces

when the thin films of those were grown on lattice matched

SrTiO3 (STO) substrates. In the present day microelectronics

industry, STO is unsuitable for fabricating the devices based

on the multiferroics. The integration of this important class

of multiferroic heterostructures on the CMOS compatible

substrate Si (100) is needed to advance the field and realize

its tremendous future prospects for data storage applications.

To address this, a few studies10,11 have reported on the depo-

sition of BTO thin films on buffered (with SrTiO3) Si (100)

by taking the advantage of molecular beam epitaxy (MBE)

technique to avoid the formation of an interfacial SiOx layer

that would disrupt epitaxial growth. For instance, recently, a

group12 from IBM have reported on the deposition of a BTO

thin film on Si (100), in which, they have used MBE to

deposit a STO buffer layer on Si (100). As early13 as 1995,

the epitaxial growth of BTO on Si (100) has been reported,

deposited by PLD using TiN as a buffer layer, in which, TiN

was used as both top and bottom electrodes. In another

report,14 vertically aligned thin film nanostructures of BTO-

CeO2 have been deposited on silicon substrates. However,

none of the above studies reported on the magnetic proper-

ties of two-phase epitaxial multiferroic BTO/LSMO on

Si(100). In that spirit, we demonstrate the epitaxial integra-

tion of BTO/LSMO thin film heterostructures on Si (100)

using TiN/MgO as buffer layers and report on their ferro-

magnetic properties as a function of BTO layer thickness.

Magnetization jumps associated with the BTO structural

transitions were suppressed in M(T) curves, probably due to

substrate clamping. The ferromagnetic Curie temperature

(TC) of the bilayers is close to that of the parent LSMO film,

and, found to be independent of the BTO layer. More inter-

estingly, we find that the BTO/LSMO bilayer structure hav-

ing a BTO layer thickness of 100 nm shows almost 2-timesa)ssingam@ncsu.edu

0021-8979/2014/116(22)/224104/6/$30.00 VC 2014 AIP Publishing LLC116, 224104-1

JOURNAL OF APPLIED PHYSICS 116, 224104 (2014)

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higher magnetic coercive field, 3-fold reduction in saturation

magnetization, and improved squareness compared to the

sample without BTO.

EXPERIMENTAL DETAILS

To describe the deposition procedure briefly, TiN, MgO,

LSMO, and BTO targets were ablated sequentially in the

same chamber using a rotating target assembly. The deposi-

tion of TiN layer was performed at 625 �C in vacuum

(1� 10�6 Torr). Following TiN deposition, the first few

mono layers (for about 500 pulses) of MgO were deposited

under vacuum (1� 10�6 Torr) at 575 �C. The remaining

MgO was then deposited at the same temperature under an

oxygen pressure of 6� 10�4 Torr. Finally, LSMO and BTO

depositions were carried out under a O2 partial pressure of

2� 10�1 Torr at a substrate temperature of 750 �C. The

energy density and pulse frequency were 1.5–2.5 J/cm2 and

10 Hz, respectively. Once completed, the samples were

cooled slowly to ambient temperature under the same O2

partial pressure. For ferroelectric measurements, capacitor

structures were made separately using SrRuO3 (SRO) and Pt

as bottom and top electrodes, respectively. Three samples of

BTO/LSMO were prepared by varying the thicknesses

(25, 50, and 100 nm) of BTO layer. We found that the mag-

netization data of 25- and 50-nm thick BTO samples are

similar. Hence, in this letter, we present the data pertaining

to two samples, in which, the thicknesses of BTO layer are

100 and 50 nm, labeled as samples A and B, respectively.

For comparison, the magnetization data of the sample with-

out BTO is also included, labeled as sample C. The thickness

of the LSMO layer is kept constant in all the studied samples

at 217 nm. No post growth oxygen annealing was performed

on any of the structures discussed here. All the magneti-

zation measurements are performed in the plane of the film.

The structures of BTO samples were characterized by

XRD h-2h scans using a Rigaku x-ray diffractometer with

Cu Ka radiation (k¼ 1.5418 A�). XPS was conducted on a

SPECS FlexMod system equipped with an Al Ka

monochromatic x-ray source (1486.7 eV) to identify (if

present) any surface metal contamination and to probe the

oxidation states of the respective elements from BTO. High

temperature XRD measurements were performed by heating

the sample in a high temperature displex up to a maximum

of 700 �C, and a temperature control of 60.5 �C. The micro-

structures of these films were characterized using a JEOL-

2000FX transmission electron microscope (TEM). A detailed

atomic-resolution study at BTO/LSMO/MgO interfaces was

performed, using a JEOL-2010F high resolution TEM

(HRTEM), equipped with a Gatan image filter tuning attach-

ment, which has a point-to-point resolution of 0.18 nm.

RESULTS AND DISCUSSION

In Fig. 1(a), we present the h-2h XRD pattern of the

BTO/LSMO/MgO/TiN/Si (100) heterostructure. For clarity,

we show the zoomed version of the Bragg peaks correspond-

ing to the LSMO and BTO (002) reflections in Figs. S1(a)

and S1(b).15 It is evident from this pattern that all the layers

show preferential (00l) orientation, suggesting either the tex-

tured or epitaxial growth of the multilayered structure sug-

gesting the presence of 90� ferroelectric domains in BTO.

From the 2h XRD data for the (002) peak, we determined the

out-of-plane (OOP) lattice parameter of BTO layer for sam-

ple A ands B as 3.993 and 3.985 A, respectively, less than

the bulk value16 of 4.036 A, indicating that these BTO layers

are under compression in OOP, suggesting that the film is

under tension in the plane of the film. The c-axis lattice con-

stants of LSMO layers in samples A, B, and C are estimated

as 3.869, 3.857, and 3.854 A, respectively, while the bulk

LSMO shows 3.870 A, with the residual strain of about

1.5%–0.5%. The epitaxial growth and the in-plane (IP) ori-

entation of all the three layers were studied in detail by

means of u-scan XRD. As depicted in Fig. 1, the u-scan

patterns of (111) reflection for BTO, LSMO, and Si were

collected. This pattern shows 4 peaks separated by �90�

indicating its cubic symmetry and establishing the cube-on-

cube relationship of the BTO with the underlying substrate

FIG. 1. (a) Typical h-2h (out of plane) XRD patterns of BTO/LSMO/MgO/TiN/Si (100) heterostructures (samples A, B, and C) showing high quality, single

phase and only (00l) reflections. (b) u-scan XRD patterns of BTO, LSMO, and Si of (111) reflection collected from sample A at 2h¼ 38.90�, x¼ 19.45�, and

v¼ 55.07� for BTO; 2h¼ 40.20�, x¼ 20.10�, and v¼ 55.07� for LSMO; and 2h¼ 28.46�, x¼ 14.23�, and v¼ 54.74� for Si (100). This pattern shows 4 peaks

separated by �90� indicating its pseudo cubic/rhombohedral symmetry, establishing the cube-on-cube relationship with the underlying substrate Si (100).

224104-2 Singamaneni et al. J. Appl. Phys. 116, 224104 (2014)

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Si (100), confirming that all 3 layers are grown epitaxially

cube-on-cube, i.e., (001)epilayer//(001)buffer and [110]epi-

layer//[110]buffer. Figure S1(c) shows15 the rocking curve

(to evaluate mosaicity, which is, angular dispersion along the

growth direction) of the BTO (002) peak; where the full

width at half maximum (FWHM) is about 1� to 1.3�.Fig. 2(a) is a typical bright-field cross-section TEM

image of the sample B, in which the BTO/LSMO/MgO/TiN/

Si (100) layers are labeled. The thicknesses of BTO, LSMO,

MgO, and TiN are determined as �50, 217, 175, and

133 nm, respectively. It should be emphasized that the

epitaxial growth of BTO/MgO on Si (100) is made possible

due to the epitaxial growth of large mismatched system

based on the DME paradigm, e.g., TiN on Si (100) where

four lattice constants of TiN match with three of Si

(100).10,11 An important feature of DME concept10,11,17 is

that most of the strain is relieved almost immediately upon

initiation of growth, i.e., within the first couple of mono-

layers of growth. In this way, lattice misfit strain accommo-

dation is confined to the interface making it possible for the

rest of the film to be grown free of defects and lattice strain.

TiN has good lattice match with MgO, which in turn, has

good lattice match with LSMO. In addition, TiN also pro-

vides a good diffusion barrier. More details on TiN/Si depo-

sition can be found in our earlier work.10,11 Figures 2(b) and

2(c) present typical high resolution electron microscopic

(HREM) images taken at the LSMO/MgO and BTO/LSMO

interfaces, showing clean and coherent growth. The tempera-

ture dependent x-ray diffraction, Raman, and ferroelectric

polarization measurements clearly show the room tempera-

ture ferroelectric nature of BTO films. Temperature depend-

ent I-V measurements show hysteric non-leaky ferroelectric

behavior. The in-detail experimental findings are communi-

cated in another paper.18 For clarity, we present the

polarization-voltage (P-V) hysteresis measurements col-

lected for several BTO devices are shown in Fig. S2.15

The magnetic response of the BTO/LSMO heterostruc-

ture with variation in temperature, measured by the use of

a superconducting quantum interference device (SQUID)

magnetometer (Quantum Design, MPMS-XL), is presented

in Fig. 3, for all the three samples. The temperature depend-

ence of the in-plane magnetization (M vs T), was measured

in small magnetic fields of 20 and 100 Oe after cooling the

samples from T¼ 400 K to T¼ 4 K in the plane of the sam-

ple, i.e., in the (100) plane of BTO under no magnetic field

(zero field cooled). As M(T) curves measured both at 20

and 100 Oe show similar behavior, we show the data corre-

sponding only to 100 Oe measuring field. The data were

collected15 (see Fig. S3) while cooling and heating the sam-

ple and found no thermal hysteresis. All the samples show

clear saturation at low temperature, though with varying

saturation magnetization (discussed later). The saturation

FIG. 2. (a) Bright field cross-section TEM image taken from sample B, where BTO (�50 nm) film was grown at 750 �C. All 4-layers are marked. The scale

bar is 100 nm (b) HRTEM image of LSMO (217 nm)/MgO (175 nm) interface. The scale bar is 5 nm (c) HRTEM image of BTO (50 nm)/LSMO (217 nm) inter-

face. The scale bar is 2 nm. The two interfaces are clean and sharp without inter diffusion and secondary phases.

224104-3 Singamaneni et al. J. Appl. Phys. 116, 224104 (2014)

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magnetization (�380 emu/cc) of LSMO without BTO is in

close agreement with the reported values.3 Two worthy

observations can be noted: (a) the ferromagnetic Curie tem-

perature (TC), the point at which the magnetization rises, is

found to be independent of BTO layer thickness, consistent

with the earlier reports and (b) surprisingly, the large, sharp,

and discontinuous jumps in the temperature dependent mag-

netization data associated with BTO structural transforma-

tions are absent in both samples A and B. We noticed that

such behavior is observed when the data were collected dur-

ing both the heating and cooling cycles. In contrast, such

jumps were observed when thin films of Fe,5 Fe3O4,7

LCMO3,4 and LSMO1–3 were deposited on BTO single crys-

tal substrates.

In Figure 4, we present the isothermal M-H data col-

lected on three samples at 4 K, when the samples were

cooled under 100 Oe. Three profound observations can be

noted: (i) the squareness (ratio of saturation to remnant mag-

netization) of M-H loop is improved with the increase of

BTO film thickness varying from 0–100 nm, (ii) the coercive

field (Hc) of sample A (�365 Oe) is almost twice that of

sample C (�163 Oe), and (iii) the saturation magnetization

of sample A is reduced to about 1/3 that of sample C. We

have noticed that the same trend (data not shown) is followed

when the measurements were performed at several higher

cooling fields such as 6300, 6500, 61000, and 62000 Oe

collected at 4 K, 170 K, 190 K, 270 K, and 290 K, below the

ferromagnetic Curie temperature of (�350 K) LSMO. The

observed trend is found to be independent of the polarity of

the field cooling. This indicates that the observed magnetic

features as a function of BTO film thickness are not associ-

ated with the BTO structural variations caused by the tem-

perature sweep and applied magnetic fields.

Now, we discuss the possible sources that might lead to

strong BTO-thickness dependence on LSMO magnetic prop-

erties. M vs T heating and cooling branches do not exhibit

distinct anomalies and concomitant thermal hysteresis at

BTO T-O and O-R structural transitions. The arrest of mag-

netization jumps could be due to the clamping of BTO/

LSMO heterostructure to the underlying Si (100) as it was

recently19 argued in the case of Fe/BTO fully epitaxial films.

The substrate may constrain the expansion accompanying

the structural transitions. This is a kind of clamping effect is

common for thin films. If the BTO goes into a multi-domain

state during the cubic to tetragonal transition where 50% of

the structural domains have a c-axis in the plane and 50% of

the structural domains have a c-axis normal to the plane, the

effect may be diminished or absent.

In addition, the cross-sectional HREM images (see

Figs. 2(b) and 2(c)) did not manifest the appearance of any

defects close to the interface in the investigated films. In

addition, the temperature-dependent magnetization curves

suggest the existence of just a single-phase magnetic state. If

there was any phase separation or magnetic inhomogeneity

in the bulk of the LSMO film, one would expect a clear ther-

mal hysteresis between the cooling and heating runs of M(T)

curves. Oxygen vacancies are known to severely affect the

magnetic and magneto transport properties of LSMO/BTO

bilayers and superlattices deposited on STO substrate.20

Oxygen vacancies are not believed to explain the observed

change in magnetic properties since the same O2 partial pres-

sure has been used in all the bilayer samples deposition. The

newly discovered8,9 long range Coulomb interaction mecha-

nism recently forwarded to explain magneto electric cou-

pling for granular multiferroic materials could not be applied

in the present case as the geometry of our thin film hetero-

structures is significantly different.

Similar to our case, Alberca and co-authors previously

have reported21 that magnetization at 10 K was reduced

almost 2-fold and coercive field was enhanced 8-fold when

12 nm-thick ferromagnetic LCMO layer is in contact with a

(001)-oriented ferroelectric BTO substrate measured, com-

pared to LCMO deposited on bare STO substrates. In addi-

tion, they found that the ferromagnetic Curie temperature of

FIG. 3. Magnetization vs. temperature (M-T) curves of sample A (in black),

sample B (in red), and sample C (in blue) for all the structures. The data

were collected during the cooling cycle with the measuring field of 100 Oe.

As it can be noticed, the Curie temperature (TC) of all the samples is found

to be the same at �350 K. The magnetic field is applied along h100i direc-

tion of the sample.

FIG. 4. Comparison of isothermal (4 K) M-H curves measured on all the

samples cooled from 400 K under 100 Oe. As one can notice, sample A

shows 2-times higher Hc, much improved squareness (Ms/Mr) and 3-times

lesser saturation magnetization (Ms) in comparison with sample C (with no

BTO). All the experiments were conducted under the same conditions as

mentioned above. The magnetic field is applied along h100i direction of the

sample.

224104-4 Singamaneni et al. J. Appl. Phys. 116, 224104 (2014)

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LCMO was reduced from 270 K to 200 K. They have

explained these magnetic features to be a result of non-

uniform strain field in LCMO/BTO caused by the corruga-

tion of the ferroelectric domains in the rhombohedric phase

of the BTO bulk crystal. Though at the microscopic level,

we cannot rule out the above possibility, this situation is

unlikely in the present case since we did not observe thermal

hysteresis between cooling and heating M(T) runs (see

Fig. S3), caused by inhomogeneous strain fields that would

generate inhomogeneous magnetic phases. The bright field

transmission electron microscopy and high resolution elec-

tron microscopy images (see Figs. 2(a)–2(c)) show continu-

ous epitaxial films without granularity. We observed the

Curie temperature (350 K) close to the value reported

(360 K) for bulk single crystalline LSMO.22 Furthermore, we

performed zero field cooled (ZFC) and field cooled (FC)

temperature dependent magnetization (M-T) measurements

on sample A (see Fig. S4). We see no indication of a block-

ing temperature, a characteristic feature of inhomogeneous

magnetic phases. Thus, our comprehensive experimental

findings provide no evidence for the presence of inhomoge-

neous phases/grain boundary effects. If a significant inhomo-

geneous phase is present, the Curie temperature is

expected23 to be widely different from the bulk. Hence, in

the present case, our experimental observations led us to

believe that grain boundaries do not affect the magnetic na-

ture of the sample. The nature of thin film epitaxy minimizes

the large angle grain boundaries unlike bulk polycrystalline

materials.

In addition, we did not see a clear strain effect of the

BTO layer on the ferromagnetic properties of the LSMO

layer either in the form of discontinuous magnetization

jumps or varying ferromagnetic Curie temperature as the

structural transitions of BTO layer seem to be arrested or

broadened. Our results are consistent with the findings of a

recent spin wave resonance study,24 reported on

LSMO(330 nm)/BTO(270 nm) epitaxial heterostructure de-

posited on LaAlO3. This particular study concludes that the

BTO layer itself induces in-plane surface pinning at the

BTO/LSMO interface that modifies the in-plane bulk mag-

netic anisotropy of the LSMO film. Due to strong in-plane

pinning between BTO and LSMO induced by the ferroelec-

tric BTO layer, a strong reduction in saturation magnetiza-

tion and enhanced coercive field are observed, and appear to

vary as a function of the stress state of the BTO. We believe

that the stress state of the BTO layer directly influences the

reordering of the atoms at the BTO-LSMO interface and

thus affects the extent of magnetic pinning in the LSMO

layer. To support our hypothesis and to prove that the ferroe-

lectricity in BTO causing the changes in LSMO magnetic

properties, we have inserted a thin (�10–15 nm) insulating-

non-magnetic-non-ferroelectric SrTiO3 layer between BTO

and LSMO and measured its magnetic properties (data not

shown). We observed that the corresponding M-H and M-T

curves behave very similar to the sample C without BTO

layer and retains all the pristine (magnetization, ferromag-

netic Curie temperature, and coercive field) characteristics of

LSMO layer, thus confirming that it is most likely that the

ferroelectric nature of BTO causing the changes in the

magnetic properties of LSMO. To explore this further, we

are planning to perform temperature dependent MOKE

measurements on our samples to probe the super exchange

coupling between Mn and Ti ions, which could lead to the

observed reduction in magnetization and enhancement in co-

ercive field.

SUMMARY

In summary, in the present work, we have demonstrated

the effect of ferroelectric BTO layer thickness on the ferro-

magnetic properties of LSMO layer, which have been epitax-

ially deposited on the technologically important substrate Si

(100). Bilayers of BTO/LSMO thin films preserve good fer-

romagnetic properties with a ferromagnetic Curie tempera-

ture of �350 K, close to the expected value, and found to be

independent of BTO films thickness of 25–100 nm. Our data

show that the discontinuous magnetization jumps associated

with BTO structural transitions were suppressed in M(T)

curves, probably due to substrate clamping effect. More

interestingly, the bilayer structure with BTO layer thickness

of 100 nm shows almost 2-fold higher magnetic coercive

field, 3-fold reduction in saturation magnetization, and

improved squareness compared to the sample without BTO.

We believe that the strong in-plane spin pinning by BTO

ferroelectric layer at the BTO/LSMO interface could lead to

the observed magnetic properties of LSMO in contact with

BTO. This work demonstrates that we could manipulate the

magnetic properties of ferromagnetic layer by conjoining

with a ferroelectric layer. Our work makes a promising step

in the realization of multiferroic materials for CMOS

applications.

ACKNOWLEDGMENTS

S.R.S. acknowledges financial support from the

National Academy of Science (NAS), USA in the form of

NRC postdoctoral research associate fellowship. We thank

Chrstian Binek, University of Nebraska and N. D. Mathur,

University of Cambridge for stimulating discussion. Also,

the authors are pleased to acknowledge the support of the

Army Research Office under Grant W911NF-04-D-0003.

Also, the authors acknowledge the use of the Analytical

Instrumentation Facility (AIF) at North Carolina State

University, which was supported by the State of North

Carolina and the National Science Foundation.

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