Enhanced dielectric properties of Pb0.92La0.08 Zr0.52Ti0.48O3 films withcompressive stressBeihai Ma, Shanshan Liu, Sheng Tong, Manoj Narayanan, and U. (Balu) Balachandran Citation: J. Appl. Phys. 112, 114117 (2012); doi: 10.1063/1.4768926 View online: http://dx.doi.org/10.1063/1.4768926 View Table of Contents: http://jap.aip.org/resource/1/JAPIAU/v112/i11 Published by the American Institute of Physics. Related ArticlesHomogeneous pinhole free 1nm Al2O3 tunnel barriers on graphene Appl. Phys. Lett. 101, 203104 (2012) Single-crystalline BiFeO3 nanowires and their ferroelectric behavior Appl. Phys. Lett. 101, 192903 (2012) Influence of target composition and deposition temperature on the domain structure of BiFeO3 thin films AIP Advances 2, 042104 (2012) The structural and piezoresponse properties of c-axis-oriented Aurivillius phase Bi5Ti3FeO15 thin films depositedby atomic vapor deposition Appl. Phys. Lett. 101, 112903 (2012) Accessing intermediate ferroelectric switching regimes with time-resolved transmission electron microscopy J. Appl. Phys. 112, 052013 (2012) Additional information on J. Appl. Phys.Journal Homepage: http://jap.aip.org/ Journal Information: http://jap.aip.org/about/about_the_journal Top downloads: http://jap.aip.org/features/most_downloaded Information for Authors: http://jap.aip.org/authors

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Enhanced dielectric properties of Pb0.92La0.08 Zr0.52Ti0.48O3 filmswith compressive stress

Beihai Ma,1,a) Shanshan Liu,1 Sheng Tong,1,2 Manoj Narayanan,1

and U. (Balu) Balachandran1

1Energy Systems Division, Argonne National Laboratory, Argonne, Illinois 60439, USA2School of Electronic and Computing Systems, College of Engineering and Applied Science, University ofCincinnati, Cincinnati, Ohio 45221, USA

(Received 21 September 2012; accepted 6 November 2012; published online 10 December 2012)

We deposited ferroelectric (Pb0.92La0.08)(Zr0.52Ti0.48)O3 (PLZT 8/52/48) films on nickel foils and

platinized silicon (PtSi) substrates by chemical solution deposition. Prior to the deposition of PLZT, a

conductive oxide buffer layer of LaNiO3 (LNO) was deposited on the nickel foil. Residual stresses of

the films were determined by x-ray diffraction. Compressive stress of ��370 MPa and tensile stress

of �250 MPa were measured in �2-lm-thick PLZT grown on LNO-buffered Ni foil and PtSi

substrate, respectively. We also measured the following electrical properties for the PLZT films

grown on LNO-buffered Ni and PtSi substrates, respectively: remanent polarization, �23.5 lC/cm2

and �10.1 lC/cm2; coercive electric field, �23.8 kV/cm and �27.9 kV/cm; dielectric constant at

room temperature, �1300 and �1350; and dielectric loss at room temperature, �0.06 and �0.05.

Weibull analysis determined the mean breakdown strength to be 2.6 MV/cm and 1.5 MV/cm for

PLZT films grown on LNO-buffered Ni and PtSi substrates, respectively. The difference in dielectric

properties and breakdown strength can be attributed to the residual stress in the PLZT films. Our

results suggest that compressive stress enhances the dielectric breakdown strength of the PLZT films.VC 2012 American Institute of Physics. [http://dx.doi.org/10.1063/1.4768926]

INTRODUCTION

In recent years, significant interest has been focused on

the development of ferroelectric (FE) thin film materials

because of their potential applications for decoupling capaci-

tors, micro-actuators, electro-optical components, and digital

memories.1–4 Applications are also possible in advanced

power electronics, which require capacitors that can operate

under high voltage and yet have minimal footprint. This

requirement can be fulfilled by stacking or embedding high-

permittivity ceramic film capacitors into a printed circuit

board (PCB).5,6 Such a technology would free up surface

space and reduce the number of solder joints on the PCB, and

this improvement would lead to increased device reliability

and minimization of electromagnetic interference and induct-

ance loss. Although this technology has primarily received

attention for decoupling capacitors in microelectronic appli-

cations, it can potentially be extended to high-power energy

storage, such as plug-in hybrid electric vehicles.

The maximum energy storage capability of a capacitor is

proportional to the dielectric constant times the square of the

breakdown strength. Therefore, ferroelectric materials with

high dielectric constant and breakdown strength provide sig-

nificant advantages in reducing the volume and weight of an

energy storage device. However, at present, integration of

high-permittivity ceramic dielectric films into PCBs is chal-

lenging because of the incompatibility in the processing con-

ditions for the different materials involved. Polymer layers in

a PCB cannot withstand the high temperatures (600–700 �C)

required for processing the ceramic film to the desired crys-

talline structures. Development of these crystalline structures

is extremely difficult at reduced processing temperatures.7

Success has been demonstrated through a film-on-foil

approach, where a ceramic is first coated on a thin base metal

foil by chemical solution deposition and then crystallized at

high temperature.8,9 These coated foils can subsequently be

embedded into a PCB. By far, the vast majority of FE films

are fabricated on single crystals or silicon wafers.10–13 Use of

metallic substrates, however, is more desirable for the fabri-

cation of large-area devices with low cost. Metal foils coated

with dielectric material, the so-called “film-on-foil” dielectric

sheets, could be laminated onto PCBs to create embedded

capacitors that have broad applications in power electronics.

Alternatively, these film-on-foils could be stacked to form a

capacitor cartridge with minimal footprint on a circuit board.

However, substantial technical challenges still exist in the

fabrication of FE film-on-foil capacitors that exhibit high

dielectric strength and are crack free. Problems include ther-

mal expansion mismatch between films and metal substrates,

formation of a low-dielectric-constant parasitic layer at the

film/substrate interface, and diffusion of cations from the sub-

strate into the dielectric film.

Recently, several research groups have reported that

stress or strain can significantly affect the physical and elec-

tric properties of ferroelectric materials.14–16 In epitaxial

BaTiO3 thin films, for example, biaxial compressive strain

resulted in a higher ferroelectric transition temperature

(approximately 500 �C) and enhancement of remanent polar-

ization (at least 250%) compared to bulk BaTiO3 ceramics.

The stress magnitude and sign can be controlled by modify-

ing the interface between the film and substrate which have

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

bma@anl.gov. Telephone: 630-252-9961. Fax: 630-252-3604.

0021-8979/2012/112(11)/114117/9/$30.00 VC 2012 American Institute of Physics112, 114117-1

JOURNAL OF APPLIED PHYSICS 112, 114117 (2012)

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small lattice misfit strain. Regarding the electrical properties

with residual stress level, Lee et al.15 reported that the ferro-

electric properties of (111)-oriented Pb(Zr,Ti)O3 (PZT) thin

films were enhanced by inducing compressive stress, which

was introduced during annealing by bending the substrate.

Xu et al.17 have also shown that highly oriented PZT thin

films under compressive residual stress exhibited superior

ferroelectric properties compared to bulk polycrystalline fer-

roelectrics. However, prior studies are mostly based on the

epitaxial thin films or require a special fixture to control the

stress. Epitaxial films with thickness over 1 lm are difficult

to fabricate because of the critical stress-relaxing thickness.

We had earlier reported18–20 fabrication of high-quality

polycrystalline ferroelectric films grown on base metal

(nickel and copper) foils. These films exhibit superior ferro-

electric properties and dielectric breakdown strength under

high (>2� 106 V/cm) applied field. To produce film-on-foil

sheets with high quality, we found it necessary to apply a

conductive oxide LaNiO3 (LNO) buffer to the nickel foil

prior to the deposition of Pb0.92La0.08Zr0.52Ti0.48O3 (PLZT).

This architecture allows crystallization of dielectric ceramic

films grown on nickel foils in air without the formation of

deleterious interfacial secondary phases. An alternative

approach is to process the ceramic dielectric films in a con-

trolled atmosphere with reduced oxygen partial pressure.20

In this paper, we report our recent results of dielectric char-

acterization and residual stress analysis of PLZT films grown

on LNO-buffered nickel and platinized silicon (PtSi)

substrates.

EXPERIMENT

High-purity (99.98% pure) nickel substrates with dimen-

sions of 25 mm� 25 mm� 0.4 mm were obtained from MTI

Corp. (Richmond, CA). They were polished by a chemical-

mechanical planarization process. A root-mean-square sur-

face roughness of �2 nm was measured by atomic force

microscopy (AFM) in the tapping mode with 5 lm� 5 lm

scan size. Platinized silicon wafers with 500-nm-thick ther-

mally oxidized silicon dioxide coated with 20-nm-thick tita-

nium and 200-nm-thick platinum were acquired from Nova

Electronic Materials (Flower Mound, TX). Prior to being

coated, the substrates were ultrasonically cleaned in distilled

water, then wipe-cleaned with acetone and methanol in

sequence.

Nickel substrates were first coated with a conductive ox-

ide film of LNO by chemical solution deposition to serve as a

buffer layer. Precursor solutions of 0.2 M LNO were prepared

by dissolving an appropriate amount of lanthanum nitrate hex-

ahydrate and nickel acetate tetrahydrate in 2-methoxyethanol

(all from Sigma-Aldrich) and refluxing for 2 h inside a chemi-

cal glove box. In addition, precursor solutions of 0.5 M PLZT

were prepared by a modified 2-methoxyethanol synthesis

route5,21 using an appropriate amount of titanium isopropox-

ide, zirconium n-propoxide, lead acetate trihydrate, and lan-

thanum nitrate hexahydrate (all from Sigma-Aldrich). The

resulting stock solution contained 20% excess lead to com-

pensate for lead loss during the heat treatments described

below. Before being used for coating, the LNO and PLZT pre-

cursor solutions were filtered through Whatman (Whatman

Inc., Sanford, ME) syringe filters of 0.02-lm pore size. The

filtered LNO precursor solution was spin coated with a Laurell

WS400 spin processor (Laurell Technologies, North Wales,

PA) at 3000 rpm for 30 s on the Ni substrates, pyrolyzed at

450 �C for 5 min, and annealed at 625 �C for 2–5 min in air.

This process was repeated three to five times to build the

desired �0.4-lm-thick LNO buffer film. Subsequently, to pro-

duce a single coating layer, filtered PLZT precursor solution

was spin coated on the LNO-buffered Ni or PtSi substrates at

3000 rpm for 30 s, followed by pyrolysis at 450 �C for 5 min

and annealing at 650 �C for 5–10 min. After every three

layers, additional annealing was performed at 650 �C for

15 min. Solution coating and firing were repeated to produce

films of desired thickness. All pyrolysis and annealing were

performed in air in Lindburg tube furnaces. The final PLZT

coating thickness after pyrolysis and crystallization was

�0.115 lm per coating. Film thickness was determined from

cross-sectional images with scanning electron microscopy

(SEM). Residual stress in the film was studied by a Bruker

(Madison, WI) AXS D8 diffraction system equipped with

a General Area Detector Diffraction System (GADDS) detec-

tor. Film microstructures were characterized by a Hitachi

S-4700-II field-emission electron microscope. The surface

roughness was determined by a Veeco Instruments D3100

scanning probe microscope operated in the tapping mode.

Platinum (Pt) top electrodes with thickness of 100 nm

were deposited by electron-beam evaporation through a

shadow mask to define �250-lm-diameter capacitors. Sam-

ples with Pt top electrodes were annealed at 450 �C in air for

2 min for electrode conditioning. A Signatone QuieTempVR

probe system with heatable vacuum chuck (Lucas Signatone

Corp., Gilroy, CA) was used for electrical characterization.

An Agilent E4980A precision LCR Meter was used to deter-

mine the capacitance and dissipation factor under an applied

bias field. A Radiant Technologies Precision Premier II tester

measured the hysteresis loops. The capacitor samples were

immersed in Fluka silicone oil (Sigma-Aldrich) during meas-

urements of high-field hysteresis loops and dielectric break-

down. A Keithley 237 high-voltage source meter measured

the current-voltage characteristics for determination of the

steady-state leakage current density.

RESULTS AND DISCUSSION

Figures 1(a) and 1(b) show the AFM height images

measured over a 1 lm� 1 lm scan area on �2-lm-thick

PLZT films deposited on LNO-buffered nickel and on PtSi

substrate, respectively. We measured the average grain size

to be �45 nm and �25 nm on PLZT films grown on LNO/Ni

and PtSi substrates, respectively. The root mean square

(RMS) surface roughness determined from AFM was

roughly the same value (�2 to 3 nm) for both samples grown

on LNO/Ni and PtSi substrates.

Figure 2 shows the 2h x-ray diffraction (XRD) patterns

measured on �2-lm-thick PLZT films deposited on LNO-

buffered Ni and PtSi substrates. The patterns indicate that

both samples are well crystallized without preferred orienta-

tion. All peaks can be indexed accordingly (JCPDS 56-0900).

114117-2 Ma et al. J. Appl. Phys. 112, 114117 (2012)

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The use of LNO prevented the formation of a parasitic inter-

facial nickel oxide layer. The LNO also compensated for the

roughness of the Ni foil and provided a smooth interface for

the PLZT film to grow on. Diffraction peaks from the nickel

substrate and the LNO buffer film can be identified from the

diffraction pattern shown in Fig. 2(a). The tiny peak at

2h� 36�, as indicated by a small open triangle on the XRD

pattern for PLZT grown on PtSi (Fig. 2(b)), is from Cu-Kb

diffraction of Pt(111). No crack or delamination was

observed from microscopic analysis.5

A close look at the two diffraction patterns in Fig. 2

revealed that the PLZT peaks in Fig. 2(a) shifted to lower

angle when compared to those peaks for the same index

shown in Fig. 2(b). With the regular h-2h scan configuration,

diffraction patterns are measured on these crystallites with

the diffraction plane parallel to the substrate surface.

According to Bragg’s equation, a peak shift to lower 2h

angle indicates an expansion of d-spacing in the out-of-plane

direction, which is a result of compressive strain in the in-

plane directions that are parallel to the substrate surface.

Residual stress analysis by XRD indicated compressive

stress in the PLZT grown on LNO-buffered nickel substrate,

and tensile strain in the PLZT grown on PtSi substrate, as a

result of the difference in the coefficient of thermal expan-

sion (CTE).

Figure 3 shows XRD peaks measured under various tilt

angles on PLZT deposited on LNO/Ni and PtSi substrates

for three high angle diffractions: PLZT (211), (220), and

(310). It is clear that peaks measured on PLZT film grown

on LNO/Ni substrate shift to higher 2h angles with increas-

ing tilt angle, while peaks for PLZT film grown on PtSi sub-

strate shift to lower 2h angles with increasing tilt angle.

Lattice d-spacing for a specific group of planes (hkl) can be

calculated from Bragg’s diffraction equation

2d sin h ¼ nk; (1)

where n is an integer (the diffraction order, n¼ 1 for this

study), k is the wavelength of the incident x-ray, and d is the

interplanar spacing. A change in Bragg angle, Dh, is related

to the change in d-spacing, Dd. Taking the derivative of

Eq. (1), the lattice strain in the film under investigation can

be calculated by the following equation:

e ¼ Dd

d0

¼ � cos hsin h

� Dh ¼ �cot h � Dð2hÞ2

; (2)

where d0 is the unstrained lattice d-spacing and e is the me-

chanical strain related to the change in diffraction angle Dh.

Considering that the Bragg angle satisfies 0� h� 90� (thus

cot h> 0), shifting to higher 2h angle (Dh> 0) with increas-

ing tilt w indicates a compressive strain (e< 0) in the film in-

plane direction. Such residual stress in the PLZT films can

be determined from XRD data measured under various tilt

angles22 and the equationFIG. 2. X-ray diffraction patterns of PLZT grown on (a) LNO-buffered Ni

and (b) PtSi substrates.

FIG. 1. AFM images measured with 1 lm� 1 lm scan area on (a) LNO-buffered Ni and (b) PtSi substrates.

114117-3 Ma et al. J. Appl. Phys. 112, 114117 (2012)

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ru ¼E

1þ �

� �ðhklÞ� 1

du0

� @duw

@sin2w; (3)

where E and � are Young’s modulus and Poisson ratio of the

thin film, d/0 and d/w are the lattice d-spacing in a stress-

free sample and d-spacing measured in a stressed sample

that was tilted at angle w, respectively. We measured XRD

patterns at various rotation angles (/¼ 0�, 45�, and 90�) and

observed no measurable difference. This finding indicates

that crystallites in the polycrystalline films are randomly ori-

ented without preferred orientation.

Figure 4 shows lattice d-spacing as a function of sin2wmeasured by XRD on PLZT films grown on LNO/Ni and

PtSi substrates. Data measured from PLZT (211), (220), and

(310) diffraction peaks are plotted in Figs. 4(a)–4(c), respec-

tively. Slopes determined from linear fitting to the data were

plugged into Eq. (3) to determine the residual stresses in the

PLZT films. Results are summarized in Table I. The elastic

modulus and Poisson ratio values used for calculation of the

residual stresses were taken from those listed in Table II.

Our measurements of residual stresses are in good agreement

with those reported for PZT films.27,28 We noticed that resid-

ual stresses measured using (211) and (220) diffraction peaks

are virtually the same, while the values determined from

(310) diffraction peaks are �40% higher than those deter-

mined from (211) and (220) peaks. This difference is likely a

result of rhombohedral distortion that causes a split for the

PLZT (310) peak in strained PLZT thin films.

Figure 5 shows the polarization-field (P-E) hysteresis

loops measured at room temperature on �2-lm-thick PLZT

films grown on LNO/Ni and PtSi substrates. Both samples

were coated with Pt top electrodes and measured using a field

sweeping period of 10 ms (frequency of 100 Hz). From the

P-E loop with a maximum applied voltage of 30 V (corre-

sponding to an electric field of 150 kV/cm), we measured a re-

manent polarization (Pr) of �23.5 lC/cm2 and �10.1 lC/cm2,

and coercive electric field (Ec) of �23.8 kV/cm and �27.9 kV/

cm, for PLZT films grown on LNO/Ni and PtSi substrates,

respectively. This difference in P-E loops can be attributed to

the total strain in PLZT films. Lee et al.29 reported that the re-

manent and saturation polarizations decrease with an increase

in tensile residual stress and increase with an increase in

compressive stress (decrease in tensile stress), as in-plane

compressive stress is beneficial to dipole alignment along the

out-of-plane direction that is parallel to the applied field. Our

experimental results revealed the same trend.

Remanent polarization and coercive field determined

from P-E hysteresis loops measured with various maximum

applied fields are plotted in Figs. 6(a) and 6(b). Both Pr and

Ec increases with increasing maximum applied field because

larger applied field leads to increased domain wall move-

ment and better dipole alignment in the PLZT films. With a

maximum applied field of 150 kV/cm, Pr measured on PLZT

grown on LNO/Ni substrate is more than two times of the

value measured on PLZT grown on PtSi substrate. The dif-

ference in Pr decreases with increasing maximum applied

field. With a maximum applied field of 1000 kV/cm, the dif-

ference in Pr drops to less than 25% as shown in Fig. 6(a).

On the other hand, Ec measured on PLZT grown on LNO/Ni

and that measured on PLZT grown on PtSi substrates are

nearly the same when measured with a maximum applied

field of 150 kV/cm. The difference in Ec increases with

increasing maximum applied field. With a maximum applied

field of 1000 kV/cm, Ec measured on PLZT grown on PtSi is

almost twice as much as that measured on PLZT grown on

LNO/Ni as shown in Fig. 6(b). This behavior can be under-

stood as follows: dipoles in the PLZT film grown on PtSi are

not so well aligned because of the tensile stress in PLZT

film. When the PLZT film is subjected to external field, those

dipoles can be aligned at a faster rate than those better

aligned ones in the PLZT grown on LNO/Ni substrate which

is under compressive stress. This leads to a greater change in

remanent polarization with increased applied field. Due to

the exact same reason, dipoles in PLZT film under tensile

stress are more difficult to get aligned. When the applied

field changed its direction, it is more difficult to get the

dipoles flipped and higher field is required to get the job

done. This resulted in an increased Ec when higher maxi-

mum field is applied for the hysteresis loop measurement.

FIG. 3. X-ray diffraction peaks: (a) PLZT (211), (b) PLZT (220), (c) PLZT

(310) measured under various tilt angles on films grown on LNO/Ni; (d)

PLZT (211), (e) PLZT (220), (f) PLZT (310) measured under various tilt

angles on films grown on PtSi substrates.

114117-4 Ma et al. J. Appl. Phys. 112, 114117 (2012)

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Figure 7 shows the dielectric constant and loss as a func-

tion of bias field measured at room temperature for samples

of PLZT grown on LNO/Ni and PtSi substrates. The data

conformed to butterfly shaped curves, which are typical for

FE materials measured under the Curie temperature. At zero

bias field, a dielectric constant of �1300 and dielectric loss

of �0.06 were measured for PLZT grown on LNO/Ni sub-

strate, and a dielectric constant of �1350 and dielectric loss

of �0.05 were found for samples deposited on PtSi sub-

strates. Since both samples have PLZT films of the same

thickness (�2 lm) but different materials beneath, they ex-

hibit different field-dependent change, i.e., “tunability,”

which can be expressed as

g ¼ 1� erðEÞerð0Þ

� �� 100%; (4)

where er(E) is the dielectric response in the presence of an

external field E. We observed a tunability of �55% and

FIG. 4. Lattice d-spacing as a function of sin2w measured from (a) PLZT (211), (b) PLZT (220), and (c) PLZT (310) diffraction peaks.

TABLE I. Summary of strain and stress determined from XRD.

Substrate

type

Peak

index

(hkl)

2h(�)

@d/w/@

(sin2w) (A)

d/0

(A)

r(MPa)

LNO/Ni (211) 55.2 �0.0111 1.665 �373

LNO/Ni (220) 64.6 �0.0095 1.442 �372

LNO/Ni (310) 73.2 �0.0122 1.293 �530

PtSi (211) 55.5 0.0075 1.655 255

PtSi (220) 65.0 0.0064 1.434 251

PtSi (310) 73.9 0.0078 1.282 343

TABLE II. Selected physical properties of relevant materials.

Material E (GPa) �

CTE, a(�10�6/k) References

Ni 200 0.31 13.4 23

Pt 168 0.39 8.8 24

Si 130 0.28 3.6 25

LNOa 70 0.30 8.2 26

PLZT 72 0.30 5.4 27 and 28

aYoung’s modulus (E) and Poisson ratio (�) values for LNO are not avail-

able. We estimated these values based upon those for other perovskites.

CTE¼ coefficient of thermal expansion.)

114117-5 Ma et al. J. Appl. Phys. 112, 114117 (2012)

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�40% at room temperature under 100 kV/cm applied field

for PLZT deposited on LNO/Ni and PtSi substrates, respec-

tively. This difference in tunability can be attributed to dif-

ference in residual stress for PLZT films grown on different

substrates. Because the dipoles in PLZT film under compres-

sive stress have better tendency of getting aligned in the out-

of-plane direction, the dipole momentum is easier to get

changed under small applied field (or near zero-field); but

under high applied field, dipoles are more clamped and

dipole momentum is more difficult to get changed. This is

why we measured sharper peaks in the butterfly curve of

dielectric constant as a function of applied field for PLZT

grown LNO/Ni substrate and broader peaks for PLZT grown

on PtSi substrate.

Figure 8 shows the dielectric constant and loss measured

as a function of frequency for PLZT grown on LNO/Ni and

PtSi substrates. The dielectric constant decreases while

dielectric loss increases with increasing frequency for both

samples. No obvious effect observed due to different stress

levels in the PLZT films grown on LNO/Ni and PtSi

substrates.

Figure 9 shows the time relaxation for the current den-

sity measured at 25 �C on PLZT films grown on LNO/Ni and

PtSi substrates with a constant bias field of �100 kV/cm

applied across the top and bottom electrodes. The measure-

ments were conducted by keeping the top Pt electrode posi-

tive and the bottom Ni electrode grounded. Both curves

show strong initial time dependence with comparable magni-

tude in current density. The decay in dielectric relaxation

current obeys the Curie-von Schweidler law30

J ¼ Js þ J0 � t�n; (5)

where Js is the steady-state current density, J0 is a fitting

constant, t is the relaxation time in seconds, and n is the

slope of the log-log plot. Fitting the data to Eq. (5), we

FIG. 5. Polarization-field hysteresis loops measured at room temperature on

PLZT films grown on LNO/Ni and PtSi substrates.

FIG. 6. (a) Remanent polarization and (b) coercive electric field as a function of maximum applied field measured at room temperature on PLZT films grown

on LNO/Ni and PtSi substrates.

FIG. 7. Field-dependent dielectric properties measured at room temperature

on PLZT films grown on LNO/Ni and PtSi substrates.

114117-6 Ma et al. J. Appl. Phys. 112, 114117 (2012)

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determined steady-state leakage current densities of �9.0

� 10�9 A/cm2 and �5.1� 10�8 A/cm2, and n values of 0.7

and 0.8 for �2-lm-thick PLZT grown on PtSi and LNO/Ni

substrates, respectively. Higher steady state leakage current

density measured for PLZT grown on LNO/Ni substrates is

likely due to larger grain size and therefore increased grain

boundary conduction when compared to PLZT film grown

on PtSi substrate.

The dielectric breakdown strength was measured on

250-lm-diameter capacitors with Pt top electrodes using the

“top-to-bottom” electrode configuration. The applied voltage

was increased by 5 V per second, and the breakdown

voltage was determined with a 1-lA criterion. Weibull statis-

tics31,32 was employed for failure behavior analysis. The em-

pirical two-parameter Weibull function for failure

probability distribution, also known as the “cumulative dis-

tribution function,” can be written as

FðtÞ ¼ 1� exp � t

g

� �b" #

; (6)

where b and g are fitting parameters, the Weibull modulus

and mean value of experimental observable t, respectively.

The Weibull modulus is a measure of sample distribution,

high values of b corresponding to smaller variation. For

breakdown field measurements, the parameter of interest is

applied field. Mean breakdown strength can be extracted

from points where the fitting lines intersect with the horizon-

tal line through lnðlnð1=ð1� pÞÞÞ¼ 0. We calculated the

sampling cumulative distribution parameter p, using Ber-

nard’s approximation, from the following formula33 for each

set of experimental data

p ¼ r � 0:3

N þ 0:4; (7)

where N is the total sample number and r is the sample se-

quential number.

Figure 10 shows the Weibull plot of breakdown field

strength obtained at room temperature from 30 capacitor

samples of PLZT grown on LNO/Ni and another 30 capac-

itor samples of PLZT grown on PtSi substrates. The solid

straight line is a least-square fitting to the two-parameter

distribution function given in Eq. (6). We measured mean

breakdown field strengths of 2.6� 106 V/cm and 1.5� 106

V/cm, and Weibull moduli b of 7.8 and 4.0, for the

�2-lm-thick PLZT films grown on LNO/Ni and PtSi sub-

strates, respectively. Note that PLZT samples grown on

LNO/Ni and PtSi are bearing compressive stress of

��370 MPa and tensile stress of �250 MPa, respectively,

as determined from XRD analysis. When an electrical

potential is applied between the top and the bottom elec-

trodes, the electrostrictive strain causes expansion in the

longitudinal direction (same direction as the applied elec-

tric field) and contraction in the radial direction (film

FIG. 8. Frequency-dependent dielectric properties measured at room tem-

perature on PLZT films grown on LNO/Ni and PtSi substrates.

FIG. 9. Time-relaxation current density measured on PLZT films grown on

PtSi and LNO/Ni substrates.

FIG. 10. Weibull plot for breakdown field strength of PLZT deposited on

LNO/Ni and PtSi substrates.

114117-7 Ma et al. J. Appl. Phys. 112, 114117 (2012)

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in-plane direction). This condition leads to additional ten-

sile stress in the PLZT film that is subjected to the applied

field. A threshold value of external applied field (Eb) can

be defined to characterize the breakdown field strength in

a PLZT film that is free of residual stress. When there is

tensile residual stress in a PLZT film, such as the one

grown on a PtSi substrate, it needs less additional tensile

strain to reach the threshold strain value that will cause

dielectric breakdown. Therefore, this sample bearing resid-

ual tensile stress needs only an electric field of Eb � DE1

(with DE1> 0) to reach the breakdown threshold value.

For PLZT grown on LNO/Ni substrate, due to the existing

compressive stress, additional field strength DE2 is

required to compensate the residual compressive strain. In

other words, for the film to reach the threshold tensile

strain level value that will cause dielectric breakdown, a

total electric field of EbþDE2 (with DE2> 0) must be

applied. For this reason, we observed higher mean break-

down field strength in PLZT films grown on LNO/Ni when

compared to those grown on PtSi substrates. This differ-

ence in breakdown strength is a direct result of the differ-

ence in residual stresses inside these two types of samples,

PLZT films grown on LNO/Ni and PtSi substrates. Resid-

ual compressive strain in the PLZT film was found to be

beneficial to enhanced dielectric strength.

CONCLUSIONS

We have grown high-breakdown-strength ferroelectric

PLZT films on LNO-buffered Ni foils and PtSi substrates by

chemical solution deposition. Residual stress analysis by

XRD revealed compressive stress of ��370 MPa in the

PLZT grown on nickel foil but a tensile stress of �250 MPa

in the PLZT grown on PtSi substrates. We measured Pr of

�23.5 lC/cm2 and �10.1 lC/cm2, and Ec of �23.8 kV/cm

and �27.9 kV/cm, for PLZT films grown on LNO/Ni and

PtSi substrates, respectively. At room temperature, we meas-

ured a dielectric constant of �1300 and �1350 and dielectric

loss of �0.06 and �0.05 for PLZT films grown on LNO-

buffered Ni foils and PtSi substrates, respectively. Weibull

analysis determined the mean breakdown strength to be

2.6� 106 V/cm and 1.5� 106 V/cm for samples grown on

LNO-buffered Ni and PtSi, respectively. The difference in

dielectric breakdown strength can be attributed to the differ-

ence in residual stress for PLZT grown on LNO-buffered Ni

foil and PtSi substrates.

ACKNOWLEDGMENTS

This work was funded by the U.S. Department of

Energy, Vehicle Technologies Program, under Contract No.

DE-AC02-06CH11357.

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