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 37

Counts at 100 mA, offset by 20x

-0.2 -0.1 0.0 0.1 0.2

696 °C

464 °C

371 °C

276 °C

190 °C

109 °C

42 °C

H (reciprocal lattice units)

K (reciprocal lattice units)

12.1 nm

553 °C

3.9 nm

371 °C

10.4 nm

41 °C

21.1 nm

553 °C

-0.1 0 0.1 -0.1 0 0.1

-0.1

0.1

-0.1

0.1

DK [r.l.u.]

Fig. 1.23 (a) Sequence of scans through the PbTiO

304 peak at various temperatures, showing

the development of satellites during cooling for a 2.0 nm-thick ﬁlm. (b) Typical in-plane distri-

butions of diffuse X-ray intensity around the PbTiO

304 peak for various ﬁlm thicknesses and

temperatures. Reproduced with permission of the International Union of Crystallography

of the ﬁlm. The domain period is inversely proportional to

.H

sat

C .K

sat

the distance between the satellite and the central peak, and scales with the square-

root of ﬁlm thickness, as expected from Landau theory [186, 187]. The average

in-plane arrangement of the stripe domains may be deduced from HK maps such

as those shown in Fig. 1.23b. As seen, the stripe domain pattern depends on ﬁlm

thickness and temperature (a–c). It was also found that ﬁlms grown on substrates

with slightly larger miscut (0:25

from the (001)) exhibited stripe domains aligned

with the step edges (d).

PbTiO

=DyScO

(110)

The competition between domain formation for elastic vs. electrostatic reasons is

more intense for PbTiO

ﬁlms grown on orthorhombic DyScO

(110) substrates,

which have an in-plane lattice constant of 0:395 nm, in between the a-and

c-lattice constants of PbTiO

Noheda et al. [188, 189] used PLD to deposit such

ﬁlms, sometimes with a SrRuO

buffer layer. For PbTiO

ﬁlms only 5 nm thick, they

observed the formation of 180

domains minimizing the depolarizing ﬁeld [188].

For 30-nm-thick ﬁlms, however, they discovered the formation of 90

domains min-

imizing the elastic energy [189]. An HK map around the DyScO

pseudocubic 100

peakisshowninFig.1.24. Unlike the case of PbTiO

grown on MgO(001), the

PbTiO

cools into the ferroelectric phase without relaxation by misﬁt dislocations,

thus permitting the formation of highly periodic c/a domains. The satellites at mul-

38 T.T. Fister and D.D. Fong

Fig. 1.24 In-plane reciprocal space map at L  0, around the DyScO

100 (in pseudocubic

units). The four peaks surrounding the substrate peak stem from the SrRuO

layer (Reprinted with

tiples of H D 0:014 are evidence of regular c/a domains with a period of 28 nm,

while the broad peak at H D 0:952 stems from the in-plane lattice parameter of

the a-domains. The four peaks around the DyScO

100 are from the 30-nm-thick

SrRuO

, suggesting that domains also form in the buffer layer.

1.4.2.2 Monodomain Structures

The critical thickness for domain formation depends on the domain wall energy;

in some instances (particularly for rhombohedral domains [179, 190]), the critical

thickness can be zero. For instance, La

0:9

0:1

MnO

=SrTiO

(001) ﬁlms only 12 nm

thick show evidence for satellites and therefore twin formation [191]. However, for

ﬁlms with non-zero critical thickness, a monodomain structure can be established

in ﬁlms sufﬁciently thin.

SrRuO

=SrTiO

(001)

In bulk form, SrRuO

is orthorhombic Pbnm at room temperature, orthorhombic

Imma above 412

C, tetragonal I4=mcm above 552

C, and cubic above 677

[192]. Vailionis et al. [177] determined the effect of epitaxial strain on these

transitions, studying 20–30nm thick ﬁlms of SrRuO

grown on TiO

-terminated

SrTiO

(001) substrates by PLD or MBE at 700

1 In Situ Synchrotron Characterization of Complex Oxide Heterostructures 39

SrTiO

SrRuO

6.55

6.50

6.45

6.40

6.35

6.30

6.25

6.20

6.55

(260)

(204)

(260)

(444)

(44-4)(444)

(620)

(260)

(620)

(44-4)

(024) (-024)

(444)

(44-4)

(0-24)

(620)

6.50

6.45

6.40

6.35

6.30

6.25

6.20

SrRuO

on SrTiO

SrRuO

on DyScO

(rlu)

Fig. 1.25 (a) HL or KL maps of SrRuO

ﬁlmsgrownonSrTiO

.001/ or DyScO

.110/

substrates. SrRuO

on SrTiO

shows an orthorhombic unit cell, while SrRuO

on DyScO

tetragonal. Indices of the ﬁlm’s reﬂections are shown in bold. (Reprinted with permission from

for SrRuO

on SrTiO

along the HL plane. The degree of tilt is exaggerated for clarity

They found that at room temperature, the ﬁlm stabilized in a monodomain

conﬁguration with the following orientation: .110/

jj.001/

;Œ

110

jjŒ100

,and

Œ001

jjŒ010

. The subscripts o and c refer to the orthorhombic and cubic coordinate

systems. For reference, one may refer to the central two orthorhombic unit cells in

Fig. 1.21. When coherently strained to SrTiO

, the strain tensor for the SrRuO

contains not only normal strains but also shear strains, and the angle between the

orthorhombic a

and b

axes is no longer  D 90

but 89:41

, reﬂecting the

compressive strain state of the ﬁlm. This alters the Ru–O–Ru bond angles and

consequently the electronic and magnetic properties. As shown on the left side of

Fig. 1.25a, the inequality between a

and b

causes the SrRuO

peaks to tilt with

respect to the H axis in reciprocal space such that the 204-type positions appear at

different q

s. This is most easily visualized by the schematic drawing of reciprocal

space in Fig. 1.25b. It is seen that the distorted orthorhombic structure is actually

monoclinic.

By monitoring the 444 and 620 SrRuO

Bragg positions in situ while heating,

Vailionis et al. [177] observed a broad orthorhombic-to-tetragonal transition. How-

ever, the 221 intensity, which is also sensitive to orthorhombicity, disappears rapidly

above 310

C. This transition temperature is much less than the 552

C orthorhom-

bic to tetragonal transition in bulk SrRuO

. They also discovered that the 211 peak,

which is absent for cubic SrRuO

, does not vanish up to the highest temperatures

they could reach .730

C/, indicating that the stable crystal structure at the 700

growth temperature is tetragonal. The enlarged phase ﬁeld for tetragonal SrRuO

stems from the compressive biaxial strain state of the ﬁlm.

As shown in Fig. 1.21, there are two possible orthorhombic variants that have

the aforementioned epitaxial orientation with the substrate. It is unclear whether the

orthorhombic structure determined by Vailionis et al. consists of a single variant,

although it may be possible to tell by the intensities of the half-order peaks. The

isolation of these two variants as opposed to the other four may have been aided by

step ﬂow growth, with the c

axis favored to align along the step edge.

40 T.T. Fister and D.D. Fong

SrRuO

=DyScO

(110)

Interestingly, Vailionis et al. [178] found that 20–30nm thick SrRuO

grown on

orthorhombic DyScO

(110) does not form an orthorhombic structure but a slightly

distorted tetragonal one, with a

D b

and  D 90:49

. As shown on the right side

of Fig. 1.25a, the SrRuO

Bragg peaks appear at constant q

while the orthorhom-

bic DyScO

peaks vary considerably with q

. This change in behavior results from

the tensile strain state of the ﬁlm. Similar results were found for CaRuO

ﬁlms

grown in tension on SrTiO

(001) and NdGaO

(110) substrates. The dependence of

the Ru–O–Ru bond angles on epitaxial strain may be responsible for the observed

variations in electrical properties where increasing compressive strain leads to ﬁlms

with higher conductivity.

1.4.3 Spectroscopic Studies of Complex Oxides

While X-ray spectroscopy can provide information regarding changes in valence

and local bonding, few in situ studies have been carried out on complex oxide

thin ﬁlms. In the synchrotron community, hard X-ray surface spectroscopy has been

primarily focused on systems lacking long-range order, such as catalysts or quan-

tum dots, and techniques like XAFS are necessary to understand local structure.

The study of oxide thin ﬁlms can also be daunting theoretically, where near-edge

sensitivity to coordination and valence is a mixed blessing: charged defects, local

distortion of oxygen bonds, and chemical changes at the interfaces are often signif-

icant and competing effects in models of near-edge structure.

Nonetheless, X-ray spectroscopy can be a powerful complement to scattering

and resonant anomalous techniques for studying complex oxides under high T and

conditions. In this section, we discuss a few examples from the recent literature

that vividly illustrate the potential of X-ray spectroscopy on the study of complex

oxides.

1.4.3.1 A-Site Clustering in La

1x

MnO

In addition to the Sr surface segregation previously discussed for LSMO thin ﬁlms,

there is evidence that A-site inhomogeneity may be intrinsic to LSMO. Using

EXAFS, Shibata et al. [193] observed strontium and lanthanum clustering in bulk

samples with x

 0:3. This intrinsic clustering could act as nucleation sites for

magnetoelectronic phase separation and may alter the Curie temperature [194].

The ability of EXAFS to measure local, short-range order is especially critical

in this study, since X-ray diffraction would average out any nanoscale inhomo-

geneities. Changes in local atomic structure were correlated between data from the

Sr and La edges that were ﬁt simultaneously at compositions ranging from x

0:025–0:425. The magnitude of the Fourier-transformed ﬁne structure  from the

1 In Situ Synchrotron Characterization of Complex Oxide Heterostructures 41

123456

Distance from Sr ion (Å)

Fourier Transform Amplitude of k

Sr-Mn

Sr-La/Sr

Sr-O

x=0.075

x=0.100

x=0.175

x=0.300

x=0.375

x=0.050

x=0.025

x=0.425

x=0.225

3.1 3.2 3.3 3.4 3.5 3.6 3.7

Radial Distance (Å)

0.1

-1

Data x=0.100

Im ( (r))

Re ( r))

(r)

x =0.24

x =0.10

Fig. 1.26 (a) Fourier transform of the strontium K-edge EXAFS with varying strontium concen-

tration. Spectra are offset to aid comparison and the ﬁrst three coordination shells are labeled (b)

Detail of ﬁt at the third coordination shell for La

0:9

0:1

MnO

for a random distribution



D 0:1



and for the best ﬁt, x

.(b) Roughness as a function of time at 800

C for two different oxygen par-

Society)

0 0.1 0.2 0.3 0.4 0.5

0.1

0.2

0.3

0.4

0.5

Powder (First method)

Powder (Second method)

Crushed single crystal

Random distribution

Fig. 1.27 Strontium–strontium coordination x

as a function of strontium composition x

LSMO. The dashed line is the limit for a random distribution at the A-site. Several methods for

sample preparation were used, but the A-site clustering for x

 0:3 was found to persist in each

strontium K-edgeisshowninFig.1.26a for various strontium concentrations, with

the Sr–O; Sr–Mn, and Sr–La=Sr coordination shells labeled. In each case, the co-

ordination number can be extracted by ﬁtting the intensity of each peak. This ﬁt

includes parameters such as Debye–Waller factors and atomic scattering poten-

tials. Due to the large difference in scattering potentials between strontium and

42 T.T. Fister and D.D. Fong

lanthanum, the authors were able to extract the ratio of Sr–La and Sr–Sr nearest

neighbors. This difference can be seen in Fig. 1.27, in which a ﬁt of strontium’s

third coordination shell for La

0:9

0:1

MnO

yields a best ﬁt for x

D 0:24,rather

than the expected value of 0.1 for a purely homogeneous material. Repeating this

type of analysis, the authors found that x

tended to be statistically higher than that

expected for a random distribution at low x

, but eventually approached the random

distribution value near x

D 0:4. This overall trend is shown in Fig. 1.27.

While these ﬂuorescence-mode EXAFS measurements were performed on pow-

der and crushed-crystal pellets, the analysis could be extended to thin ﬁlms by

working at grazing incidence. Additionally, the data were measured at 10 K, but

could be extended to in situ conditions since the strontium and lanthanum K-edges

are well into the hard X-ray regime. While the authors carefully analyzed the oxygen

stoichiometry during sample fabrication, it would be interesting to analyze the effect

of ionic defects on strontium clustering in the same manner as in the analysis above.

1.4.3.2 B-Site Electronic Structure in La

1x

MnO

As discussed earlier, XANES is far more sensitive to changes to the local chemistry

than EXAFS, which is a more mature technique for measuring atomic structure. For

oxide perovskites, XANES is most often used to measure the oxidation state at the

B-site and to ﬁngerprint changes in the hybridized d-states that result from octahe-

dral distortions. One study that systematically tackled both of these themes was the

analysis of the Mn K-edge by Croft et al. [80]foraseriesofbulkLa

1x

MnO

(LCMO) samples. The samples were measured in ﬂuorescence and transmission

mode and were found to agree with the surface-sensitive total electron yield mode.

The systematic changes in XANES with varying calcium content are shown in

Fig. 1.28. The two most prominent features, the ﬁrst pre-edge peak and the ﬁrst

feature after the edge, are marked A and B, respectively. The position of feature B,

which correlates with the edge position, was compared with known standards and

stems from the evolution of Mn

to Mn

with increasing calcium concentration.

The area of the weakly allowed d-state, marked A, also correlates with the shift in

edge position. The abrupt discontinuity at x

D 0:2 is associated with the onset

of conductivity and ferromagnetism in LCMO and originates from the Mn

ions

necessary for double exchange. Below x

D 0:2, these effects also occur with a

transition from the cubic to rhombohedral crystal structure.

While this work was performed ex situ, the authors did compare samples that

had been annealed in air and in pure oxygen, which would presumably create an

excess of oxygen at cation vacancy sites. As shown in Fig. 1.29, they measured the

shift in the Mn K-edge as a function of x

, with the expectation that the oxygen

uptake might be most pronounced below x

D 0:2 and explain anomalous change

in Mn oxidation state. Instead, the Mn K-edge was most responsive at x

D 0:2

and 0.4. While this effect did not immediately address the trend seen in Fig. 1.29,it

likely reﬂects the defect chemistry of LCMO and should be readdressed with in situ

experimentation over a wider range of P

, where changes in the stoichiometry and

Mn valence state would be more pronounced.

1 In Situ Synchrotron Characterization of Complex Oxide Heterostructures 43

Fig. 1.28 (a) Change in La

1x

MnO

XANES with varying x

. Both the pre-edge feature (A)

and the position of the edge (B) were used to analyze the Mn oxidation state. The inset shows the

pre-edge feature at the two composition extremes. (b) Top: Correlation of edge position and pre-

edge area with Ca concentration x

. To verify the anomalous behavior for x

<0:2the authors

measured the edge position at several intermediate compositions. Bottom: Evolution of the pre-

edge feature with increasing Ca concentration (Reprinted with permission from [80]. Copyright

1997 by the American Physical Society)

Fig. 1.29 (a) Change in LCMO edge position for samples prepared in air and pure O

for x

D 0,

0.4, and 1.0. Despite the small change in P

, there is a noticeable shift in the 0.4 sample. (b)

Overall results indicate that only the x

D 0:2–0:4 samples had a statistically signiﬁcant response

Physical Society)

44 T.T. Fister and D.D. Fong

1.4.3.3 Polarization Sensitive Near-Edge Structure in Bi

4x

and La

0:7

0:3

MnO

While most surface sensitive XANES measurements have been on catalysts or other

disordered systems where a short-range order probe is necessary, it can also be

useful to pinpoint element-speciﬁc structural distortions in thin ﬁlm oxides. For

instance, the layered perovskite Bi

4x

(BNdT) is a candidate for fatigue-

free ferroelectric memory and is thought to have two unique titanium sites, which

are labeled Ti1 and Ti2 in Fig. 1.30. Chon et al. [84] used polarization-dependent

XANES to identify the direction and site of the “giant” ferroelectric distortion in

PLD and sol-gel grown ﬁlms on Pt/TiO

=SiO

=Si(100) substrates. The Ti K-edge,

showninFig.1.30, has several notable pre-edge features that result from weakly

allowed d-states. The difference in out-of-plane and in-plane XANES is most pro-

nounced at the features marked A and B, which are known to be weakly allowed

quadrupole transitions 1s ! 3dt

and 1s ! 3dt

, respectively. Distortions

in TiO

octahedra increase the intensity of these peaks due to orbital mixing with

dipole-allowed 4p states. The substantial increase in peak B in the out-of-plane

XANES conﬁrms that the ferroelectric distortion is along the c-axis. Further model-

ing using an ab initio multiple scattering code found that the distortion only occurred

at outer Ti2 sites.

Fig. 1.30 Polarization-dependent XANES data taken from the titanium K-edge of a 280 nm

3:15

0:15

ﬁlm (BNdT) grown on (100)-silicon. Ab initio calculations of the two indi-

vidual Ti sites shows that the observed giant polarization originates from the Ti2 site, which is

American Physical Society)

1 In Situ Synchrotron Characterization of Complex Oxide Heterostructures 45

1.5 Outlook

The recent developmentsin X-ray optics, detectors, and analytical tools [11–13,195]

have already left an indelible imprint on the ﬁeld of complex oxide heterostruc-

tures, and some existing techniques like resonant anomalous SXRD [53, 54]and

resonant phase-retrieval methods [196] will undoubtedly become more prevalent as

researchers continue to explore the intimate connections between structure and in-

terface chemistry. With the advent of fourth generation sources, probes will become

fully coherent and have shorter pulse widths. This will permit the full determina-

tion of structures with short correlation lengths (e.g., polydomain ﬁlms, relaxors,

or nanoparticle perovskites), and, with advances in detector technology, allow re-

searchers to monitor processes like phase transitions and domain nucleation with

vastly improved spatial and temporal resolution. It will be important to continue

the development of in situ growth and processing facilities so that we may take full

advantage of these revolutionary developments.

Acknowledgments We thank S. Nakhmanson, S. Hong, B. Ravel, A. Sorini, and J. Kas for

their comments and suggestions on the manuscript. The submitted work has been created by

UChicago Argonne, LLC, Operator of Argonne National Laboratory (“Argonne”). Argonne, a U.S.

Department of Energy Ofﬁce of Science laboratory, is operated under Contract No. DE-AC02–

06CH11357.

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