Ramanathan Sh. (Ed.) Thin Film Metal-Oxides: Fundamentals and Applications in Electronics and Energy

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1 In Situ Synchrotron Characterization of Complex Oxide Heterostructures 27

0 250 500 750 1000

3.85

3.9

3.95

4.05

Film Thickness (Å)

H (recip. latt. units)

0 500 1000 1500 2000

Time after Growth Start (s)

Coherent Position

Relaxed PZT

0 250 500 750 1000

3.93

3.94

3.95

3.96

3.97

3.98

Film Thickness (Å)

Surface In plane Lattice Par. (Å)

0 1000 2000

0.02

0.04

0.06

0.08

0.1

0.12

Time after Growth Start (s)

Surface Composition x

Fig. 1.15 (a) Scattered intensity during repeated radial .H / scans as a function of time near the

PZT 402 during growth at vapor Zr mole fraction x

vap

D 8:6% at 0.04 nm/s. The reciprocal lattice

units are based on room temperature cubic SrTiO

.(b) Evolution of the in-plane lattice parameter

and surface composition during the growth shown in (a) (Reprinted with permission from [135].

signiﬁcantly alter material properties and produce complex thickness-dependent

behavior. Another implication of this study is that the strain state affects incorpora-

tion behavior even for uniformly strained ﬁlms (which may be true for oxygen as

well as cation incorporation [141]). This example illustrates the importance of car-

rying out spectroscopic measurements in parallel with structural studies for complex

oxide systems.

1.4 Complex Oxide Film Structure

1.4.1 Interfaces

In Fig. 1.1, we illustrated some of the structural instabilities that can occur in per-

ovskites as the size of the cation is varied. Regions of symmetry-breaking can also

give rise to instabilities such as the tilted octahedra shown in Fig. 1.1borc.Such

octahedra have been observed at the surface of PbTiO

[139] and are expected at

the coherently strained PbTiO

=SrTiO

interface [142]. Cation displacements and

octahedral distortions (Fig. 1.1d) are also possible at perovskite interfaces and have

been found at the LaAlO

=SrTiO

(001) buried interface [143].

Polar unit cells like that shown in Fig. 1.1d are not typically stable on their own.

The depolarizing ﬁeld induced by ferroelectric polarization must be compensated

at the interfaces by charged adsorbates [144], charged point defects or free carriers,

or by the formation of ferroelectric domains [145]. A ferroelectric material also has

28 T.T. Fister and D.D. Fong

the option of returning to a centrosymmetric unit cell. Non-ferroelectric materials

like LaAlO

(001), however, also require interfacial compensation as, in the ionic

limit, they are comprised of interleaved .LaO/

–.AlO

e

planes, and therefore

have a non-zero dipole moment. These “type 3” polar interfaces [146] can also be

compensated by charged adsorbates, charged point defects, or by a modiﬁcation in

surface charges (Fermi level pinning) [147]. If the ﬁlm is thin enough, however, it

may be possible to sustain an uncompensated surface [26].

Therefore, oxide heterointerfaces, even when free of impurities and dislocations,

can have structures and properties deviating substantially from the adjacent bulk

materials. In this section, we discuss SrTiO

; PbTiO

, and LSMO interfaces.

1.4.1.1 The SrTiO

(001) Surface

A great variety of surface symmetries have been observed on the SrTiO

(001) sur-

face including .2  1/; .2  2/; c.4  2/,andc.6 2/ [148]. Using SXRD, Herger

et al. [48] found that after performing the standard treatments for TiO

surface

termination, the room temperature SrTiO

(001) surface consists of a mixture of

.1  1/; .2  1/,and.2  2/ surface domains. The measurements were carried out

in ultra-high vacuum. However, once heated to 750

CinP

D 7:5  10

6

Torr,

only the .1  1/ structure remains, perhaps due to disordering of the .2  2/ and

.2  1/ domains. As previously found for the .2  1/ structure [149], the average

.1  1/ structure consists of a TiO

double-layer, as shown in Fig. 1.16a. Herger

et al. suggest that the double-layer structure may help to compensate the weakly po-

lar SrTiO

(001) surface [112] and could be the stable surface structure during PLD

growth under similar T and P

conditions.

Fig. 1.16 The structure of TiO

-terminated SrTiO

.001/.(a)TheTiO

double-layer model for

asampleat750

CandP

D 7:5  10

6

Torr. (b) The oxygen overlayer model for room

temperature and in air

1 In Situ Synchrotron Characterization of Complex Oxide Heterostructures 29

In contrast, Vonk et al. [150] examined TiO

-terminated SrTiO

substrates at

room temperature and in air and found no evidence for a TiO

double-layer (nor

any evidence for a surface reconstruction). Their best ﬁt model is one with an

oxygen overlayer possibly formed by dissociative water adsorption, as shown in

Fig. 1.16b. More recent work suggests that this model remains valid at high tem-

peratures [143]. It appears that despite the development of standard procedures for

SrTiO

(001) surface preparation, the surface structure can still exhibit considerable

variation. A recent study has shown that the standard buffered HF etching procedure

itself can produce a high level of defects at the surface [105]. The authors suggest

use of an alternative etchant to mitigate possible surface damage.

1.4.1.2 The LaAlO

=SrTiO

(001) Buried Interface

After growing LaAlO

on SrTiO

(001) by PLD (T D 800

CandP

D 10

4

to 10

6

Torr), Ohtomo and Hwang found that the n-type AlO

–LaO–TiO

–SrO

interface exhibited high electrical conductivity; the p-type LaO–AlO

–SrO–TiO

interface, however, remained insulating [28]. The discovery of quasi-2D electron

gas (q2DEG) behavior at the interface between two band insulators has attracted

a tremendous amount of interest, and several groups around the world are investi-

gating its underlying physics [1, 151]. One possible cause is the polar discontinuity

at the interface. Unlike SrTiO

; LaAlO

(001) has a strongly polar surface. When

LaAlO

is grown on SrTiO

(001) past a critical thickness [152], compensation is

necessary at the interfaces. For TiO

-terminated substrates, this may be accom-

plished by surface charge modiﬁcation such that the terminating TiO

plane has

a net charge of e=2; since Ti is multivalent, this can occur by proportioning the

normally Ti

into half Ti

and half Ti

[153]. To determine the atomic-scale

structure of this buried interface, Willmott et al. [50] performed SXRD measure-

ments on 14 inequivalent CTRs and 9 additional symmetry-equivalent CTRs on a

5-unit-cell thick LaAlO

ﬁlm grown on TiO

-terminated SrTiO

(001). The ﬁlm was

prepared at 770

CandinP

D 3:75  10

6

Torr and cooled to room temperature

for measurement. Both COBRA [154]andFIT[155] were used to determine the

atomic positions throughout the ﬁlm thickness.

A summary of their results is shown in Fig. 1.17a. The nominal interface is in-

dicated by the dashed vertical line. As seen, the transition between Sr and La lies

at 7 unit cells. Thus, the expected transition between Ti and Al is at 7.5 unit cells

rather than at the observed 8.5 unit cells, suggesting that the interfacial layers are

alloys with an approximate composition of .La

0:7

0:3

/.Ti

0:6

0:4

. Using EELS,

Nakagawa et al. [153] also observed that the n-type interface was relatively rough

(about twice as rough as the p-type interface). They suggested that Sr and La in-

termix at the interface in response to the large interface dipole energy induced by

the delocalized electrons. Compensation at the p-type interface, however, may be

accomplished by an atomic rather than electronic rearrangement (e.g., an accumu-

lation of oxygen vacancies at the interface) and no intermixing is necessary.

30 T.T. Fister and D.D. Fong

Fig. 1.17 Summary of

COBRA and FIT results

across the ﬁlm-substrate

interface for a ﬁve-unit-cell

thick LaAlO

ﬁlm grown on

TiO

-terminated

SrTiO

.001/.(a)The

occupancies of each atom as

functions of position. Error

bars are small compared with

the data-point circles. (b)The

occupancies of Ti

and

as functions of position.

The nominal interface is

indicated by the dashed line

(Reprinted with permission

the American Physical

Society)

0.0

0.2

0.4

0.6

0.8

1.0

Occupancy of layer

bulk

surface

0.2

0.4

0.6

0.8

Occupation

10 12

Position [unit cell]

By requiring that the ratio of Ti

to Ti

minimizes the electric ﬁeld across each

unit cell and that the electric potential be zero at the ﬁlm boundaries, Willmott et al.

[50] partitioned the Ti proﬁle into Ti

and Ti

proﬁles, as shown in Fig. 1.17b.

The abnormally large out-of-plane expansion found at the interface could be ex-

plained by the presence of Ti

, which has a larger radius than Ti

Vonk et al. [143] also used SXRD to examine submonolayer thick LaAlO

TiO

-terminated SrTiO

(001) at both 850 and 200

C .P

D 7:510

5

Torr/.After

growth at 850

C, the atomic positions in the surface layers .AlO

–LaO–TiO

–SrO/

correspond well to bulk SrTiO

lattice positions. However, once the sample is

cooled to 200

C, the Sr, La, and Al cations shift away from the interface while

the O anions shift toward it, in good agreement with HRTEM results [156]. Vonk

et al. found that the TiO

octahedra are compressed along the out-of-plane direc-

tion and the AlO

,LaO,andTiO

planes display large cation displacements. The

octahedral (Jahn–Teller) distortion may be related to the presence of Ti

at the

interface, while the large dependence of surface structure on temperature may stem

from the cubic to rhombohedral transition (at 550

CforbulkLaAlO

[157]).

Similar q2DEG behavior has been observed for the LaVO

=SrTiO

(001) n-type

interface [158] but not for the p-type interface, nor for heterostructures with (110)-

interfaces that have no polar discontinuity. A good review on this topic has recently

been published [151].

1 In Situ Synchrotron Characterization of Complex Oxide Heterostructures 31

Fig. 1.18 (a) Depiction of the reconstructed c.22/ PbTiO

surface. The subsurface TiO

square

is rotated by 10 ˙ 1

. Adjacent unit cells are counter-rotated. (b) Depiction of the reconstructed

PbTiO

=SrTiO

interface with an AFD

structure. Arrows indicate the rotation of oxygen octahe-

dra around Ti

1.4.1.3 The PbTiO

(001) Surface

Using in situ SXRD after MOCVD growth of PbTiO

on SrTiO

(001), Munkholm

et al. [139] found that the PbO-terminated surface undergoes a c.2  2/ surface re-

construction, irrespective of the ﬁlms being in the paraelectric or ferroelectric phase.

The c.2  2/-afﬁliated SRs had narrow rocking curve widths (e.g., 0:08

for the

0), corresponding to a surface domain size larger than 180 nm. After measuring

the integrated intensities of 14 independent SRs at L  0, they found that the best

ﬁt model consists of alternating clockwise/counterclockwise rotations .10 ˙ 1

/ of

the Ti-centered oxygen squares in the topmost PbTiO

unit cell (Fig. 1.18a).

This antiferrodistortive (AFD) reconstruction occurs only for PbO-terminated

surfaces and is motivated by the reduction in bond length between the Pb ions in

the surface plane and O ions in the subsurface plane. First principles calculations

suggest that the TiO

squares in underlying unit cells rotate out-of-phase with each

other, as for the AFD

instability. The rotation magnitude, however, decays quickly

away from the surface [159].

If insufﬁcient lead oxide vapor pressure is supplied at high temperature, the

PbTiO

begins to decompose into TiO

(Fig. 1.14). Consequently, the c.22/ struc-

ture lifts and is replaced by a slowly forming .1  6/ surface that is most probably

TiO

-terminated [160]. The .1  6/ surface reconstruction has not been solved.

1.4.1.4 The PbTiO

=SrTiO

(001) Buried Interface

Studies on PbTiO

=SrTiO

superlattices show that the buried interface may also ex-

hibit counter-rotated octahedra, as shown in Fig. 1.18b[142]. Here, the ferroelectric

polarization couples with the AFD instabilities AFD

and AFD

. Successive TiO

squares along the [001] rotate in- or out-of-phase with each other in the AFD

and AFD

modes, respectively. First principles calculations again indicate that

32 T.T. Fister and D.D. Fong

the rotation magnitudes decay rapidly away from the interface, demonstrating that

this behavior is purely an interfacial effect. Since the ferroelectric ground state

is coupled with the AFD

and AFD

distortions at the interface, short-period

PbTiO

=SrTiO

superlattices behave as improper ferroelectrics – materials in which

polarization is no longer the primary order parameter. Synchrotron X-ray studies

may also demonstrate evidence of c.2  2/-related SRs for these superlattices.

1.4.1.5 The La

0:7

0:3

MnO

.001/

Surface

Unlike the surfaces described above, no atomic reconstructions have been ob-

served on La

0:7

0:3

MnO

.001/

, i.e., the pseudocubic (001) surface [51]. However,

changes in the average composition [51,161–163] and the manganese valence state

[164] at the surface and interface of LSMO thin ﬁlms have been observed. Since

the electronic and magnetic properties of LSMO depend on the interplay between

strontium concentration, the manganese charge state, and oxygen nonstoichiometry,

these surface effects are potentially important in controlling the transport properties

as well as catalytic behavior [165,166].

Until recently, measurements of strontium surface segregation have required the

use of ultrahigh vacuum and/or room temperature conditions that are far from

equilibrium. Using high energy X-rays, strontium surface segregation can also be

measured at high T and P

using total reﬂection X-ray ﬂuorescence (TXRF).

While TXRF has primarily been used for trace-element detection in the semi-

conductor industry [167], the technique can be easily adapted for compositional

depth-proﬁling by ﬁtting the dependence of the ﬂuorescence signal on the incident

angle ˛ [168]. TXRF probes the top 2–4 nm for ˛<˛

and is primarily sensi-

tive to the bulk above the critical angle, which can be seen in the ﬂuorescence

spectra shown in Fig. 1.19a. Below ˛

, there is a clear increase in the strontium-

to-lanthanum ﬂuorescence signal. Shown in Fig. 1.19b, the ˛ -dependence of this

ﬂuorescence proﬁle can be used to extract an average strontium concentration in the

top 2 nm of the ﬁlm.

Using this approach, Fister et al. [169] found strontium surface segregation in

PLD-grown LSMO ﬁlms on NdGaO

and DyScO

(110)-substrates, which pro-

duce in-plane compression and tension in the ﬁlms, respectively. As discussed in

Sect. 2, Herger et al. [51] also observed Sr segregation in LSMO ﬁlms grown

on SrTiO

, and Sr segregation was found to produce asymmetric interface proﬁles

in LaVO

=SrTiO

heterostructures [134]. Clearly, the potential impact of Sr segre-

gation on heterostructure properties should be a consideration whenever it is present

in the ﬁlm or substrate.

Fister et al. did not observe any thickness or strain dependence on segregation

behavior. The strontium surface enhancement was, however, strongly affected by

and T . Focusing on a 16 nm LSMO=DyScO

ﬁlm, the strontium segregation

was found to persist over a wide range of temperatures .300–900

C/ and oxygen

partial pressures (0.15–150 Torr). As seen in Fig. 1.20a, the Sr surface concentration

generally increases with decreasing T and P

1 In Situ Synchrotron Characterization of Complex Oxide Heterostructures 33

Fig. 1.19 (a) X-ray ﬂuorescence spectrum above and below critical angle for total reﬂection for

16 nm LSMO on DyScO3 (110) at 700

C and 150 Torr P

. The spectra have been normalized

to the La and Mn intensity for comparison. Note the increase in the Sr intensity below ˛

and the

increase in the substrate signal above ˛

.(b) Fit to Sr/(Sr C La) TXRF ratio (left axis) as a function

of incidence angle ˛. For reference, the penetration length is plotted on the top horizontal axis. To

emphasize the position of ˛

, the normalized Sr and La TXRF intensities are shown on the right

Fig. 1.20 (a) Sr surface region concentration vs. oxygen partial pressure at T D 300–900

C. (b)

Sr surface enrichment as a function of inverse temperature for a variety of oxygen partial pressures.

The quantity 

denotes the Sr/La ratio integrated over the ﬁrst 2 nm, while 

is the bulk ratio, i.e.,

3/7. The lines are linear ﬁts to equilibrium data taken between 500 and 900

C. The deviation from

linearity of the outlying 300

C data indicates that achieving thermodynamic equilibrium requires

longer equilibration time than the typical 1 h used in the experiment (Reprinted with permission

The P

-dependent enthalpy of segregation, H

seg

, can be extracted from the

Arrhenius plot shown in Fig. 1.20b. In the standard models for equilibrium seg-

regation [27, 170], the enthalpy of segregation is proportional to the slope of the

isobars plotted in Fig. 1.20b and was found to increase from 9:5 to 2:0 kJ=mol

for P

D 0:15–150 Torr.

The enhanced strontium surface segregation at low oxygen partial pressures

could result from electrostatic interactions between charged point defects, the con-

centrations of which depend strongly on P

in bulk LSMO. For example, a large

34 T.T. Fister and D.D. Fong

concentration of positively charged oxygen vacancies near the surface may favor a

high concentration of Sr near the surface for local electroneutrality [171]. While the

TXRF measurements were not sensitive to changes in the manganese charge state,

additional in situ grazing incidence X-ray studies near the Mn K-edge may provide

a more complete description of the surface defect chemistry.

1.4.2 Internal Structure: Non-Cubic Films on SrTiO

(001)

The deposition of epitaxially oriented complex oxide thin ﬁlms permits non-cubic

perovskites to be grown epitaxially onto cubic substrates like SrTiO

(001). For

(001) epitaxy, the orthorhombic, ferroelectric tetragonal, and rhombohedral struc-

tures may have the orientations displayed in Fig. 1.21. The orthorhombic and

rhombohedral unit cells are shown in their respective 2  2  2 pseudocubes

(producing half-order peaks in the scattering [16]), and the directions of cation

displacements are indicated in the ferroelectric tetragonal structures. As depicted,

orthorhombic and ferroelectric tetragonal perovskites have six potential variants,

while rhombohedral perovskites have four. The number of possible variants for

different symmetries and interfaces can be determined rigorously using bicrystallog-

raphy theory [172], which relies on the principles of symmetry compensation [173].

As demonstrated by Eom et al. for both orthorhombic [174,175] and rhombohedral

Fig. 1.21 Depiction of the possible variants that may arise when epitaxially growing a pseudocu-

bic cell on a cubic substrate for orthorhombic, tetragonal ferroelectric, and rhombohedral systems

1 In Situ Synchrotron Characterization of Complex Oxide Heterostructures 35

[176] perovskites, it is possible to isolate particular variants by growing in the

step-ﬂow growth mode, where preferred variants attach to the step edges. Mon-

odomain structures may also be stabilized in ultrathin ﬁlms [177,178].

In the thick-ﬁlm limit, ﬁlms may relax their coherency strain by the formation

of dislocation arrays and/or domains, each having an associated critical thickness

that depends on the balance between the reduction of elastic energy and the cost of

dislocation arrays or domain walls. For ferroelectric ﬁlms, domains may also form

to reduce the electrostatic energy. Interfacial coherency between the substrate and

a ﬁlm comprised of tilted domains can be described by a distribution of continuous

dislocations. For example, Farag et al. [179] have investigated the case of rhom-

bohedral LSMO variants grown on (001) cubic substrates. The stress ﬁelds of the

continuous dislocations can be used to determine the non-uniform strains caused

by domain formation. In this section, we survey recent studies of polydomain and

monodomain structures using synchrotron X-ray scattering.

1.4.2.1 Polydomain Structures

PbTiO

=MgO(001)

Several studies have been carried out on ferroelectric/ferroelastic domain forma-

tion and evolution in PbTiO

thin ﬁlms [38, 180]. Using in situ synchrotron X-ray

scattering, Lee and Baik [181] examined the evolution of the c=a .90

/ domain

structure in a PbTiO

ﬁlm grown on MgO(001). At room temperature, the lattice

parameter of the substrate, 0.421 nm, is larger than both the a-andc-lattice param-

eters of bulk ferroelectric PbTiO

(0.390 and 0.415nm, respectively). Growth of

the 300 nm PbTiO

ﬁlm by PLD was carried out at 650

CandP

D 0:1 Torr,

conditions appropriate for deposition in the paraelectric phase. As shown by the HL

map in Fig. 1.22, a ferroelectric phase transition occurs in the PbTiO

when cooled

from 650

C to room temperature. At 600

C, only the Bragg peak from the (par-

tially relaxed) cubic PbTiO

lattice parameter is evident. Just below the transition,

at 460

C, the ﬁlm transforms primarily into a-domains having in-plane polarization;

the intensity at higher q

originates from the associated (100) planes. The small frac-

tion of c-domains has a larger out-of-plane lattice constant, as seen by the contours

on the low q

side of the Bragg peak. Further cooling leads to larger tetragonal-

ity and increased separation between the (100) and (001) peaks at 430 and 360

As expected from domain modeling [182–184], the fraction of c-domains also in-

creases during cooling to reduce the total elastic energy of the ﬁlm. The splitting

of the (100) Bragg peak at 200 and 25

C is caused by tetragonality-induced tilting

of the domains. In the HK-plane, the splitting exhibits fourfold symmetry because

the tilting is equally likely along each h100i direction. The temperature and strain

dependence of the c-domain fraction can be compared directly with theory to better

understand the factors determining 90

domain stability [185].

36 T.T. Fister and D.D. Fong

Fig. 1.22 In situ high-temperature X-ray scattering contour maps of reciprocal space along the HL

plane showing domain evolution in PbTiO

/MgO(001) as a function of temperature during cooling

PbTiO

=SrTiO

(001)

While thick PbTiO

ﬁlms grown on SrTiO

(001) also show evidence for 90

c=a

domains [185], only 180



domains are found in ﬁlms less than 40 nm thick

[137]. These domains do not form to reduce the elastic energy of the ﬁlm but rather

to reduce the depolarizing electric ﬁeld. The domains were ﬁrst observed by the ap-

pearance of diffuse intensity around each PbTiO

peak (except those at L  0), as

shown in Fig. 1.23a. The appearance of satellites below the ferroelectric transition

indicates the presence of an in-plane structural modulation with a well-deﬁned pe-

riod, while the lack of satellites at L  0 implies that the atomic displacements lie

only along the out-of-plane direction, as expected for 180

c-domains. Furthermore,

L-scans through the satellites show thickness fringes with a spacing identical to that

along the CTR, demonstrating that the stripe domains extend through the thickness