Effect of PbZr0.52Ti0.48O3 thin layer on structure, electronic and magneticproperties of La0.65Sr0.35 MnO3 and La0.65Ca0.30MnO3 thin-films

T. Hezareh,1 F. S. Razavi,1,a) R. K. Kremer,2 H.-U. Habermeier,2 O. I. Lebedev,3

D. Kirilenko,4 and G. Van Tendeloo41Department of Physics, Brock University, St. Catharines, Ontario, Canada, L2S 3A12Max-Planck-Institut für Festkörperforschung, Heisenbergstrasse 1, D70506 Stuttgart, Germany3Laboratoire CRISMAT, ENSICAEN, Université de Caen, CNRS, 6 Bd Maréchal Juin,F-14050 Caen 4, France4EMAT, Universiteit Antwerpen (RUCA), B 2020 Antwerpen, Belgium

(Received 16 November 2010; accepted 25 April 2011; published online 6 June 2011)

Epitaxial thin film heterostructures of high dielectric PbZr1�xTixO3 (PZT) and La1�xAxMnO3(A-divalent alkaline earth metals such as Sr (LSMO) and Ca (LCMO)) were grown on SrTiO3substrates and their structure, temperature dependence of electrical resistivity, and magnetization

were investigated as a function of the thickness of the LSMO(LCMO) layer. The microstructures of

the samples were analyzed by TEM. By applying an electric field across the PZT layer, we applied a

ferrodistortive pressure on the manganite layer and studied the correlations between lattice

distortion and electric transport and magnetic properties of the CMR materials. VC 2011 AmericanInstitute of Physics. [doi:10.1063/1.3592660]

I. INTRODUCTION

Manganese based oxide perovskites, T1�xDxMnO3(where T is a trivalent rare-earth ion and D a divalent ion

such as Ca, Sr, Ba, or Pb) have long been known for their

unique electrical and magnetic properties. The Curie and

charge ordering transition temperatures of these materials

depend on the kind of dopant and the value of x, the pressureand magnetic field. An external magnetic field and pressure,

as well as internal pressure (e.g., induced by ion size mis-

match) alter the Mn-Mn electronic transfer integral which

can lead to significant changes of the magnetic and electric

transport properties of doped manganese perovskites.1 Thin

films, ceramics, and single crystals of these compounds with

carefully controlled compositions have revealed an enormous

magnetoresistance (MR).2 A large negative MR is usually

observed near the Curie temperature, Tc, (the so-called co-

lossal magnetoresistance (CMR) effect), associated with the

ferromagnetic to paramagnetic phase transition.3 These mate-

rials are of high interest for magnetic recording and memory

devices and transistors due to the CMR effect.

Among a number of perovskite type manganese oxides

with various combinations of (T, D), La1�xSrxMnO3 andLa1�xCaxMnO3 are considered to be prototypical and refer-ence materials. The parent compound LaMnO3 is an antifer-

romagnetic insulator. Substitution of the trivalent La3þ

cations by a divalent cation (such as Sr2þ or Ca2þ in thisstudy) constitutes the coexistence of Mn3þ and Mn4þ and thelanthanum manganite becomes a conducting ferromagnet.

For La1�xSrxMnO3 (LSMO) crystals, the FM state below Tcappears above a critical doping level xc � 0.17 (composi-tional metal insulator (MI) phase boundary) and increases up

to Tc � 380 K at the doping level x¼ 0.3–0.5.2

The transport properties of LCMO are optimized for a

doping value of x � 0.3. For this value, the compound’sstructure has an optimal value of ferromagnetic spin mix-

ture of Mn3þ (3d4) and Mn4þ (3d3) ions,4 and as a result,the electrical resistivity is minimized and the sensitivity to

magnetic field and pressure is increased. This composition

yielded the highest reported Curie temperature and the low-

est resistivities in the metallic state. The transport properties

of these materials are understood on the basis of the double

exchange mechanism. According to the double exchange

interaction, the transfer of an itinerant eg electron between

neighboring Mn ions (local t2g spins) through the O2� ion

results in ferromagnetic interaction due to on site Hund cou-

pling. An external magnetic field forces the local t2g spins

to align. It reduces spin scattering, therefore decreases the

resistivity, supports the ferromagnetic phase, and results in

an increase of Tc.1 External pressure has a similar effect on

the resistivity and Curie temperature. External pressure

moves the Mn-O-Mn bond angle toward 180� with anincrease of the Mn-O bond length. This results in a more

effective transfer of the eg electrons and causes the resistiv-

ity to drop and Tc to grow.4

Recently, heterostructure thin films of ferromagnetic

and ferroelectric material, so called multiferroic materials,

have attracted particular interest for their properties and their

possible technological application.5–7 The focus of our paper

is to study heterostructures of La1�xSrxMnO3(LSMO),La1�xCaxMnO3 (LCMO), and piezoelectric perovskite oxidePbZr1�xTixO3 (PZT) deposited on SrTiO3 (STO) substrates.Piezoelectric materials become strained when placed in an

electric field, and conversely, piezoelectric materials are

electrically polarized when exposed to mechanical stress.8

Consequently, application of a bias voltage across the PZT

layer causes the PZT structure either to stretch or to com-

press.9 This effect acts as an external pressure on the epitax-

ially connected manganite layer, thus altering its lattice

a)Author to whom correspondence should be addressed. Electronic mail:

frazavi@brocku.ca.

0021-8979/2011/109(11)/113707/8/$30.00 VC 2011 American Institute of Physics109, 113707-1

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parameters and in due course affecting its electrical transport

and magnetic properties. We investigated these effects as a

function of the applied bias voltage across the PZT layer and

the thickness of LSMO/LCMO layer.

II. EXPERIMENTAL METHODS

Ceramic target pellets of La0:65Sr0:35MnO3, 13 mm in

diameter and about 3 mm thick, were prepared following the

citrate synthesis route or the solgel method followed by a

final pyrolysis process. The polycrystalline powders obtained

by this method are ultra-fine and already have correct com-

positions (for details of preparation see, e. g., Ref. 10).

Thin film samples were grown by the pulsed laser deposi-

tion (PLD) technique, using an excimer laser at wavelengths

of 308 nm or 248 nm. The laser was usually recharged after

growing every two samples. Several thin-film samples were

fabricated at different temperatures with different LSMO stoi-

chiometries and, finally, the growth conditions were opti-

mized for doping level x¼ 0.35. For PZT and LSMOdeposition the substrate STO was kept at 580 �C and 680 �C,respectively.

The LSMO and PZT ceramic targets surfaces were pol-

ished and pasted on target holders with silver paste. The sub-

strates were cleaned ultrasonically in ethanol and acetone

and dried with argon gas. In order to keep the growth cham-

ber free from any dust or soot it was flushed with argon gas

before starting each growth. Initially, the growth chamber

was evacuated to a pressure of the order of 10�5 Torr and thesubstrate was heated up to 580 �C. The temperature was sta-bilized by a temperature controller. Once the chamber is

evacuated and the temperature is stabilized, oxygen gas was

allowed inside the chamber to a pressure of 0.4 mbar. The

PZT deposition was started at 580 �C. Once the depositionwas completed a stepping motor rotated and substituted the

LSMO target for PZT at the ablation position. The tempera-

ture was then increased to 680 �C for deposition of theLSMO layer. After growing the LSMO layer on PZT, the

sample was annealed in situ at 730 �C in an oxygen atmos-phere (1 bar) for an hour. The temperature was lowered fast

enough (about 50 K/min) to maintain the structure adapted

by the film yet also to avoid poor crystallization if the cool-

ing rate is too low.4

All PZT-LSMO heterostructure samples were grown

with the growth parameters as close as possible. The thick-

ness of the PZT layer was kept the same (2000 Å) for all

samples, while the LSMO layer thicknesses were varied

from 1000 Å to 500 Å in order to investigate the effect on

the LSMO transport properties. The notation PZT-LSMO

specifies the order of the layers: first the PZT layer is ablated

on the STO substrate followed by the LSMO layer deposited

on top of the PZT layer.

We also fabricated thin films of LSMO (LCMO) on

10� 10 mm niobium-doped STO (which is a conductingsubstrate). The PZT layer was grown on LSMO (LCMO)

using a mask, such that it would only cover the central area

of about 5� 5 mm on LSMO (LCMO). Also, for the growthof the manganite layer on Nb-STO another mask was applied

to leave the near-edge parts of the substrate surface uncov-

ered in order to apply contacts between the PZT layer and

the substrate.

A. Characterization of the samples

Four different heterostructures of PZT and LSMO

(LCMO) samples were deposited on (001) SrTiO3 substrates.

The diffraction patterns of PZT-LCMO, LCMO-PZT, PZT-

LSMO and LSMO-PZT thin films grown on SrTiO3 substrates

were determined by x-ray diffraction (XRD) (see Figs. 1–4).

The x-ray wavelengths used were Ka1 ¼ 1.540 Å, Ka2¼ 1.544 Å and Kb¼ 1.392 Å. As seen in the XRD patterns,the dominant wavelength is Ka. The intensity ratio of Ka2 to

Ka1 is 0.5. The double (00l) peaks observed for STO are dueto diffraction of x-rays with these two wavelengths. The (00l)peaks in the diffraction patterns reveal c-axis-oriented growthof the PZT and LSMO (LCMO) layers on the STO substrates.

Kb peaks of the substrate and the PZT layer are visible in all

of the patterns. No impurity phases have been detected. In all

samples, the PZT peaks appear at smaller angles than STO,

and the LSMO or LCMO peaks appear at larger angles due to

the lattice mismatch of the three films. Assuming the hetero-

structures to crystallize with a pseudocubic structure, the lat-

tice parameters of PZT, LSMO, LCMO, and STO were

determined to be 5.805 Å, 5.450 Å, 5.436 Å, and 3.904 Å,

respectively, in good agreement with the corresponding bulk

values.

1. TEM investigations of the thin-film samples

In multilayer perovskite oxide thin films grown on sin-

gle crystal substrates, the stress induced on the films due to

the lattice misfit with the substrate and the adjacent layers

affects their structure. Their lattice spacing and interatomic

distances can be significantly different from the bulk equiva-

lent. Therefore, the electronic structure and the physical and

chemical properties can be very different.12

Although the XRD patterns confirmed that the PZT and

LSMO(LCMO) layers had grown epitaxially on the STO

substrates, the microstructure of the films were carefully

inspected by transition electron microscopy (TEM) in order

to detect any imperfections that could affect the conductivity

and transport properties of the thin-film samples.

FIG. 1. h� 2h x-ray diffraction scan taken on LCMO-PZT thin film.

113707-2 Hezareh et al. J. Appl. Phys. 109, 113707 (2011)

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Figure 5 displays cross-section bright field images of

four different double layer films and corresponding electron

diffraction (ED) patterns. All films demonstrated well

defined layers and sharp and flat interfaces. The diffraction

contrast reveals a uniform homogeneous structure of

LC(S)MO layers while PZT layers exhibit columnar struc-

ture in most samples. Diffraction patterns confirm epitaxial

growth of all layers and evidence of stress relaxed films, if

the crystal structure of each layer is close to the crystal struc-

ture of the corresponding bulk material. Such ED patterns

are superpositions of the patterns of the STO substrate (indi-

cated by white squares in Fig. 5) and PZT and LC(S)MO

layers. The structure of PZT for all samples is found to be

close to the structure of the bulk (P4 mm), and LCMO andLSMO crystallized in space groups Pnma and R-3 C, respec-tively. Extra streaking spots on the ED pattern of the LSMO-

PZT sample (Fig. 5(c)) at position of 2ap along the growthdirection evidences the presence of stacking faults (layer

interruption in the stacking sequence) at the STO/LSMO

interface, which are clearly revealed as a dark line in the

image (marked by white arrows in (Fig. 5(c)). These stacking

faults are also confirmed by HRTEM images of the STO/

LSMO interface, which are not presented here.

An HRTEM image of the STO/PZT interface is shown

in Fig. 6. This interface looks very similar for both the PZT-

LCMO and the PZT-LSMO samples and it exhibits some

inclusions and misfit dislocation distributed along the inter-

face. It has been suggested that the dislocations and inclu-

sions coexist, and one entails the other.13 The bright contrast

inclusions seen within the PZT layer (see arrows in Fig. 6)

indicate Ti-rich regions. Such Ti-rich inclusions are respon-

sible for columnar diffraction contrasts, which are clearly

FIG. 2. h� 2h x-ray diffraction scan taken on PZT-LCMO thin film.

FIG. 3. h� 2h x-ray diffraction scan taken on LSMO-PZT thin film.

FIG. 4. h� 2h x-ray diffraction scan taken on PZT-LSMO thin film.

FIG. 5. Bright field low magnification cross-section TEM images and corre-

sponding selected area ED patterns of four different double layer films

grown on STO(001) substrate: (a) PZT-LCMO; (b) PZT-LSMO; (c) LSMO-

PZT; (d) LCMO-PZT.

113707-3 Hezareh et al. J. Appl. Phys. 109, 113707 (2011)

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visible in Fig. 5 (most prominent in Fig. 5(c)). The conse-

quences of these Ti-rich columns on the piezoelectricity of

the PZT layer will be discussed in the next section. The

STO/LSMO (cf. Figure 7) and STO/LCMO interfaces appear

to be perfectly coherent, free of dislocations, inclusions, and

amorphous or secondary phases.

The PZT/LSMO and LCMO/PZT interfaces, as shown in

Fig. 8(a) and Fig. 9(a), respectively, mainly contain periodi-

cal misfit dislocations along the interface and have a Burgers

vector of ap (ap � 0.39 nm of the order of the STO latticeparameter). The dislocations along the interface are periodi-

cal, with a periodicity of �11.5 nm for PZT/LSMO interfaceand �10.5 nm for the LCMO/PZT interface. These valuesindicate complete relaxation of the misfit between LC(S)MO

and PZT. The results of geometrical phase analysis (GPA)

shown in Fig. 8(b) and Fig. 9(b) confirm a completely relaxed

structure with well defined LC(S)MO and PZT layers. The

strain field is localized around a core of a dislocation (�2 nmarea) and does not extend into the layers. As a consequence

of the roughness at the interface between PZT/LSMO and

LCMO/PZT the electronic and magnetic properties of films

with thicknesses less than 500 Å are affected and their prop-

erties will be discussed in the following sections.

B. Measurements

1. Electrical Resistivity

The electrical resistivity of the thin-film samples was

measured by the four point contact technique using the Van

der Pauw method.14

Out of several samples of LSMO on PZT grown at dif-

ferent ablation temperatures and different stoichiometry for

LSMO(LCMO), three samples with doping level x¼ 0.35and thicknesses 1000(2000) Å, 800(500) Å and 500(100) Å

revealed satisfactory conducting behavior.

The plots of electrical resistivity versus temperature for

three samples with different thicknesses of PZT-LSMO

and PZT-LCMO are illustrated in Fig. 10 and Fig. 11,

respectively.

In both cases, the resistivity increases and the metal-in-

sulator transition temperature TMI decreases as the thickness

is reduced. In Fig. 10, for the thinnest sample, TMI decreases

to 297 K. The increase in resistivity is ascribed to the effect

of the PZT layer underneath that causes strain and distorts

the thin layer of LSMO deposited on top of it. The effect

becomes stronger as the thickness of the LSMO film is

reduced. Samples with LSMO layers thinner than 500Å are

insulators.

Overall, thicker films exhibit a higher TMI due to reduced

strain. In the layers more distant to the PZT interface, the

FIG. 6. Cross-section HRTEM image of STO/PZT interface. The white

arrows point to Ti rich inclusions.

FIG. 7. Cross-section HRTEM image of STO/LSMO interface.

FIG. 8. (Color online) Cross-section HRTEM image of PZT/LSMO inter-

face and corresponding GPA image shows strain field distribution perpen-

dicular to the interface direction.

113707-4 Hezareh et al. J. Appl. Phys. 109, 113707 (2011)

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MnO6 octahedra are structurally more relaxed and electron

transfer is more effective, thus leading to a higher TMI.

In the case of PZT-LCMO thin films the transitions

occur at lower temperatures compared with PZT-LSMO

samples and a much sharper drop in resistivity is observed

just below TMI. The 100 Å sample becomes insulating, again

due to the strain effects of the PZT layer on LCMO’s

structure.12

Another sample studied was LCMO-PZT-3, where the

PZT layer was grown on top of the LCMO layer deposited

on the Nb-doped SrTiO3 substrate. To prepare this film we

first grew LCMO-PZT layers and then using the lithography

technique and ion milling we arranged the contact configura-

tion as shown in Fig. 12; the lead configuration and applied

bias voltage is also shown in the figure. Although the TEM

results for this sample showed that the structure of the film is

without any defects and dislocations the resistivity results

indicate a semiconducting behavior with appearance of two

peaks at about 80 K and 165 K, as can be seen in Fig. 13.

This might be caused by oxygen deficiency in the LCMO

layers (see growth conditions) as well as by additional strain

induced by the top PZT layers. Compare this to the PZT-

LCMO sample where the LCMO layer relaxed for film thick-

nesses greater than 500 Å, and where it was directly con-

tacted to oxygen during 30 min of oxygen annealing.

However, film with a thickness of 100 Å shows semiconduc-

tor and insulating behavior due to the strain effect12 of the

PZT layer, as shown in Fig. 11.

Figure 13 shows the behavior of the resistivity versus

temperature for three different applied bias voltages com-

pared with zero applied voltage. The resistivity is revealed

with peaks at �100 K, �165 K, and �75 K. There is anoticeable initial drop in resistivity by about 15% once a bias

voltage of 0.185 V is applied, however, no further decrease in

FIG. 9. (Color online) Cross-section HRTEM image of PZT/LCMO inter-

face and corresponding GPA image shows strain field distribution perpen-

dicular to the interface direction.

FIG. 10. (Color online) Resistivity vs temperature for three samples with

different thicknesses of the PZT-LSMO thin films. (Thickness of PZT layer

is 2000 Å).

FIG. 11. (Color online) Resistivity vs temperature for PZT-LCMO thin-film

samples with different film thicknesses as indicated in the inset. (Thickness

of PZT layer is 2000 Å).

113707-5 Hezareh et al. J. Appl. Phys. 109, 113707 (2011)

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resistivity was observed when the bias voltage was increased.

Application of a bias voltage across the PZT layer causes the

PZT structure to either stretch or compress, depending on the

polarity of the applied voltage, thus affecting its lattice pa-

rameters.9 This exerts an external strain on the manganite

layer, affecting its transport properties by altering its lattice

parameters and, thus, the transfer integral between Mn-O-Mn

bond.

In order to understand the reason for the multiple transi-

tions, the injected current was measured at an applied bias

voltage of VB¼ 1.035 V, as a function of temperature.As seen in Fig. 14, there is a large leakage current of the

order of 10�2 A flowing through the PZT layer at low tem-peratures. This leakage is larger by four orders of magnitude

than the very low current applied to the sample to measure

its resistivity (10�6 A). The large leakage current impedesfollowing up the resistivity measurements to higher bias vol-

tages. The origin of this leakage current is ascribed to the Ti

columns in the PZT layer identified by TEM and described

before in the TEM section. When the external voltage is

applied to the sample, the Ti chains conduct current through

the PZT layer; therefore, the piezoelectric effect of PZT will

not be significantly diminished. The LSMO-PZT and PZT-

LSMO samples were also examined, but the leakage current

was even larger in those samples, and no drop in resistivity

was observed.

By measuring the resistance in two different directions

across the sample we found a noticeable difference. In Fig.

15 and Fig. 16 we have plotted R1 and R2 versus tempera-

ture, separately. The difference in the behavior of the resist-

ance in the two graphs is as a result of anisotropy in the thin-

film sample.

2. Magnetization

The magnetic response of the PZT-LSMO and PZT-

LCMO thin films was studied as a function of temperature to

confirm that the manganite layer retained its magnetic proper-

ties in the presence of the ferroelectric layer. The magnetic

response of the PZT layer and the SrTiO3 substrate have been

taken into account in order to identify the correct response of

the lanthanum manganite layer to the magnetic field. Insulat-

ing STO shows a magnetic moment that is partly diamagnetic

and partly paramagnetic.15 PZT reveals a weak temperature-

dependent paramagnetic response,16 which has also been

deduced by the comparison of the temperature dependence of

the magnetic moment of STO and STO-PZT. The magnetic

moment of STO-PZT was subtracted from the moment of

FIG. 12. (Color online) (a) Schematic view from above of LCMO-PZT-3

heterostructure grown on Nb-doped SrTiO3 substrate with the gold wire con-

tacts to the bias voltage and four contacts for resistivity configuration; (b)

cross-section view of sample.

FIG. 13. (Color online) Resistivity vs temperature for LCMO-PZT-3. The

inset indicates the bias voltages applied across the PZT film. (Thickness of

LCMO is 1000 Å and that of PZT layer is 2000 Å).

FIG. 14. Injecting current vs temperature for LCMO-PZT-3. LCMO layer

1000 Å.

113707-6 Hezareh et al. J. Appl. Phys. 109, 113707 (2011)

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PZT-LSMO and PZT-LCMO samples to yield the magnetic

response of the LSMO(LCMO) layer only. It was noticed that

the STO-PZT film had a very weak response in the case of

H¼ 100 G, but its response became significantly strong in thecase of H¼ 50 000 G. Figures 17 and 18 show the tempera-ture dependence of the magnetization ratio for PZT-LSMO

samples in 100 G and 50 000 G fields respectively, and Fig.

19 illustrates the behavior of the magnetization ratio versus

temperature for the PZT-LCMO samples in H¼ 100 G.Although the 100Å thin film is an insulator it retains its ferro-

magnetism; this effect was observed previously and has been

attributed to the epitaxial strain-induced metal insulator tran-

sition in thin films by the substrate.12

As illustrated in Fig. 18, in a much stronger field of

50 000 G, the splitting of the field-cooled (FC) and zero

field-cooled (ZFC) measurements almost vanishes, i.e., rather

strong fields are required to align the moments in the samples

in order to overcome the domain structure. Although the

magnetic response of the STO substrate and the PZT layer

were estimated and subtracted from the total signal, a dia-

magnetic response was still observed above 300 K after cor-

rection, which might be due to the strong influence of the

STO substrate versus the very thin layers of LSMO. This

effect is more pronounced in thinner films and in strong mag-

netic fields. Also, Tc has moved to a temperature higher than

375 K. In Fig. 19, the leap in the magnetization curve for the

100 Å thick sample is due to very low signals and error in

measurement.

In thick films (> 900 Å) or bulk matter, many domainsexist within which the spins are all aligned by very strong

exchange forces with Bloch-type walls separating the

domains. The spin orientation changes from one domain to

that in a neighboring domain by rotating around an axis per-

pendicular to the plane of the wall. In thin film (< 450 Å),the domain walls are of the Néel type; hence, the spin

FIG. 15. (Color online) R1 vs temperature for LCMO-PZT-3. LCMO layer

1000 Å.

FIG. 16. (Color online) R2 vs temperature for LCMO-PZT-3.

FIG. 17. (Color online) Magnetization ratio vs temperature for three sam-

ples with different film thicknesses of PZT-LSMO measured in an external

field of 100 G applied parallel to the films.

FIG. 18. (Color online) Magnetization ratio vs temperature for three sam-

ples with different film thicknesses of PZT-LSMO measured in an external

field of 50 000 G applied in the film plane.

113707-7 Hezareh et al. J. Appl. Phys. 109, 113707 (2011)

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direction changes from a given domain to that in a neighbor-

ing domain by rotating around an axis in the plane of the

wall. Bloch walls behave as three-dimensional structures,

while Néel walls behave as two-dimensional ones.18 In a

Bloch wall, when a magnetic field is applied, all the domains

force one another to align either by each domain aligning

itself with the field or by moving the domain wall. Moving a

domain wall means that a domain grows as another adjacent

domain shrinks. In thin films, the same thing happens, but

the reinforcement from other domains is limited to immedi-

ate in-plane domains.17

III. CONCLUSIONS

Epitaxial thin films of PZT-LSMO with different doping

levels for La1�xSrxMnO3 were fabricated using the pulsedlaser deposition technique. We could not grow LSMO thin-

films at the ideal temperature (750 �C) because the PZT layerwould decompose at temperatures higher than 750 �C. There-fore, with low doping levels of LSMO targets, the PZT-

LSMO thin films became insulators. It was observed that for

x¼ 0.35, the thin film samples revealed the desired metallicbehavior. Therefore, PbZr0:52Ti0:48O3-La0:65Sr0:35MnO3 thin

film samples were grown on SrTiO3 substrates (001) with

different thicknesses of LSMO. The temperature dependence

of the resistivity and magnetization as a function of thickness

carried out on PZT-LSMO as well as PZT-LCMO samples

deduced similar results as in Ref. 4.

Film thickness plays a decisive role in the transport

properties of the manganite layers. Thinner samples showed

higher resistivity and lower transition temperatures. This is

ascribed to strain affecting the structure of the manganite

layer due to lattice mismatch with the PZT layer. Addition-

ally, a higher number of oxygen defects present in thinner

films reduces the number of carriers within the material.

Thicker films exhibit higher magnetization. This is

ascribed to the difference in the nature of their domain walls.

PZT had a very weak response to the magnetic field, so the

magnetization data were very similar to those of STO-

LSMO(LCMO) samples.

As seen in the TEM images, defect formation in the

growth of the PZT layer gave rise to a large leakage current

that increased drastically with voltage. The same measure-

ments were performed on LSMO-PZT and PZT-LSMO het-

erostructures but due to growth defects in the PZT layer, this

effect was not observable.

In summary, our investigations show that in the studied

heterostructures, the manganites largely maintained their

electrical transport properties in the presence of PZT. LSMO

seems to be a better candidate than LCMO in engineering

piezoelectric and manganite heterostructures for device pur-

poses, such as miniaturized transistors and pressure sensors,

due to their higher transition temperature and broader metal-

lic phase. We have demonstrated that the thickness of the

manganite layer must be taken into account, and growth con-

ditions for such heterostructures must be optimized in order

to minimize defects in the structures.

ACKNOWLEDGMENTS

Financial support for this work was partially provided

by the Natural Sciences and Engineering Research Council

of Canada and Canada Foundation for Innovation.

1A. N. Ulyanov, I. S. Maksimov, E. B. Nyeanchi, Y.V. Medvedev, S. C.

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FIG. 19. (Color online) Magnetization ratio vs temperature for three differ-

ent thicknesses of PZT-LCMO samples with H¼ 100 G.

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http://dx.doi.org/10.1063/1.1447293http://dx.doi.org/10.1103/PhysRevB.54.1716http://dx.doi.org/10.1103/PhysRevB.54.1716http://dx.doi.org/10.1088/0953-8984/9/39/005http://dx.doi.org/10.1038/nmat2189http://dx.doi.org/10.1038/nmat1860http://dx.doi.org/10.1063/1.3327889http://dx.doi.org/10.1111/j.1551-2916.2008.02421.xhttp://dx.doi.org/10.1088/0960-1317/10/2/3070http://dx.doi.org/10.1016/j.jmmm.2004.06.005http://dx.doi.org/10.1016/j.jmmm.2004.06.005http://dx.doi.org/10.1103/PhysRevB.68.134415http://dx.doi.org/10.1063/1.125687http://dx.doi.org/10.1080/01418610151133230http://dx.doi.org/10.1103/PhysRev.147.583http://dx.doi.org/10.1063/1.363436http://dx.doi.org/10.1103/PhysRevB.44.12395

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