Growth studies of heteroepitaxial oxide thin films using reflection high-energy electron diffraction (RHEED)

Growth studies of heteroepitaxial oxide thin films using reflection high-energy electron diffraction (RHEED)

1

G. Koster, M. Huijben, A. Janssen, G. Rijnders MESAþ Institute for Nanotechnology, University of Twente, Enschede, The Netherlands

1.1

Introduction: reflection high-energy electron diffraction and pulsed laser deposition

Reflection high-energy electron diffraction (RHEED) was limited to low background pressures until the development of high-pressure RHEED, which makes it possible to monitor in situ surface structure during deposition of oxides at higher pressures, opened new possibilities (Rijnders, Koster, Blank, & Rogalla, 1997). In addition to intensity oscillations caused by layer-by-layer growth, enabling accurate control of the growth rate, it became clear that intensity relaxation caused by typical pulsed deposition leads to a wealth of information about growth parameters (Blank, Rijnders, Koster, & Rogalla, 1998). Pulsed laser deposition (PLD) has become an important technique in the fabrication of novel materials. Although use of PLD started in the mid-1960s (Ready, 1963), when initial attempts to produce high-quality thin films showed the promise of this technique, it was not until the discovery of high-Tc superconductors that PLD became widespread. The main benefits of PLD that are often quoted in literature are the relative easy stoichiometric transfer of material from target to the substrate, the flexibility of choice of materials and an almost free choice of (relatively high) background pressure. For instance, during the deposition of oxides, an oxygen background pressure up to 1 mbar is usually used. Here, we demonstrate the use of RHEED during PLD by means of four examples showing how the effect of initial growth on the recorded RHEED signal leads to information on factors such as crystal termination on growth kinetics. For an extensive overview of the various growth models that apply to highly kinetic deposition, refer to the article by Christen and Eres (2008) or Rijnders and Blank (2007).

1.2

Basic principles of RHEED1

In a typical RHEED system, a high-energy electron beam (10e50 keV) arrives at a surface under a grazing incident angle (0.1 e5 ) (Figure 1.1(a)). At these energies 1

Taken from Koster et al. (2011).

Epitaxial Growth of Complex Metal Oxides. http://dx.doi.org/10.1016/B978-1-78242-245-7.00001-4 Copyright © 2015 Elsevier Ltd. All rights reserved.

4

Epitaxial Growth of Complex Metal Oxides

(a)

Screen

ϕ i,f

θi

θf

Sample

(b) (0,0)

(2h,2k)

2 π/d

Figure 1.1 Typical setup for a reflection high-energy electron diffraction experiment in real space (a) and reciprocal space (b).

the electrons penetrate any material for several hundreds of nanometers. Because of a grazing angle of incidence, however, the electrons interact with only the topmost layer of atoms (1e2 nm) at the surface, which makes the technique very surface sensitive. (By contrast, low-energy electron diffraction is surface sensitive because of the shallow penetration depth of low-energy electrons [100e500 eV].) The scattered electrons collected on a phosphorus screen form a diffraction pattern characteristic of the crystal structure of the surface and also contain information concerning the morphology of the surface. As we will see, RHEED is remarkably sensitive to variations in morphology and roughness during thin film growth and therefore is often used as a method to monitor thickness in real time.

1.3

Variations of the specular intensity during deposition

Here we briefly discuss the variations of specular intensity as a function of the variation of the surface roughness during deposition and growth of thin films, considering basic kinetics in the case of homoepitaxial growth (e.g., strontium titanate [SrTiO3] on SrTiO3 [001]). For an in-depth review of the use of RHEED during PLD, see Koster et al. (2011).

Growth studies of heteroepitaxial oxide thin films using RHEED

5

θ=1/2

θ∼0.4

(d)

l / lo

(a)

(b) (e)

(c)

(f)

0

1

2

3

θ

Figure 1.2 (a) Intensity oscillations during homoepitaxial growth of strontium titanate (SrTiO3) at 850  C and 3 Pa, indicating true layer-by-layer growth. (d) Calculated intensity oscillations using the diffraction model (b) (a schematic representation of which is given in (e)), and a stepdensity model (c) (a schematic representation of which is given in (f)). The number of pulses needed to complete one unit cell layer is estimated to be 27.

1.3.1

Specular intensity oscillations

The intensity oscillations recorded for homoepitaxy of SrTiO3 are depicted in Figure 1.2(a), together with the intensities calculated using two different models. Similar oscillations have been observed during deposition of GaAs (Van Hove, Lent, Pukite, & Cohen, 1983) and silicon (Si) (Sakamoto et al., 1986). From the shape and amplitude we conclude that under the conditions used here, SrTiO3 deposition proceeds in a true layer-by-layer growth mode (Figure 1.2(d)).

1.3.1.1

Step density model

A widely used model is the step density model, where the specular intensity depends negatively on the number of up and down steps on the surface. Every step can act as a diffuse scatterer, and the entailed intensity is “lost” for specular reflection (Figure 1.2(f)). Although this model is highly empirical and there exists no diffraction model as a physical explanation, it turns out to qualitatively describe some of our

6

Epitaxial Growth of Complex Metal Oxides

results well. Following Stoyanov and Michailov (1988), and assuming that nucleation of islands takes place only when t ¼ T (the period for coverage of one unit cell layer), the step density evolution as a function of the coverage is given by: pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi pffiffiffiffiffiffiffiffiffi sðtÞ ¼ 2 pN0 ð1  qÞ lnð1  qÞ

(1.1)

where N0 is the initial number of nuclei when t ¼ T. When growth takes place at only one level and with a constant supply of particles (q f t), the step density oscillations again have a cusp-like shape (Figure 1.2(c)). A maximum step density s is expected at q w 0.4 (minimum intensity). For a multilevel system, qm are coupled through a series of m rate equations, where m represents the mth unit cell level. The total step density is given by summing all participating levels. The basic outcome is again a decrease of the amplitude of step density oscillations as the number of involved levels m is increased. The shape and the occurrence of a minimum for the measured intensity oscillation suggest that here the step density model is applicable, as indicated in Figure 1.2(a)e(c). The fact that we used an incident angle corresponding to the inphase condition and still observe strong intensity oscillations also favors the step density model.

1.4

RHEED intensity variations during heteroepitaxy: examples

1.4.1 1.4.1.1

Fabrication of high-quality LaAlO3eSrTiO3 interfaces2 Introduction: atomic interface ordering

The materials investigated here, LaAlO3 and SrTiO3, are band insulators with bandgaps of w5.6 and w3.2 eV, respectively, and both belong to the perovskite structural family. The SrTiO3 compound consists at room temperature of a simple cubic structure. The lattice parameters are 3.905 Å with the titanium (Ti) atoms located at the corners and the strontium (Sr) atoms at the centers of the cubes (Abramov, Tsirelson, Zavodnik, Ivanov, & Brown, 1995; Yamanaka, Hirai, & Komatsu, 2002) The oxygen atoms are placed at the centers of the 12 cube edges, giving strings of TiO6 octahedra that share corners and extend in three dimensions. The TiO6 octahedra are perfect, with 90 angles and six equal TieO bonds at 1.952 Å. Each Sr atom is surrounded by 12 equidistant oxygen atoms at 2.760 Å. The SrTiO3 compound undergoes a secondorder phase transition from cubic (spacegroup Pm3 m) to tetragonal (spacegroup I4/ mcm) at a temperature of w110 K because of the rotation of neighboring TiO6 octahedra in opposite directions (Cowley, 1964; M€uller, Berlinger, & Waldner, 1968; Shirane & Yamada, 1969; Unoki & Sakudo, 1967). In the ionic limit SrTiO3 can be described as Sr2þTi4þO2 3 . 2

Most of this section was taken with permission from Huijben et al. (2009).

Growth studies of heteroepitaxial oxide thin films using RHEED

7

On the other hand, the LaAlO3 compound consists at room temperature of a rhombohedrally distorted perovskite structure (spacegroup R3 c) that undergoes a transition to the ideal cubic perovskite structure (spacegroup Pm3 m) at w813 K (Lehnert, Boysen, Dreier, & Yu, 2000; Lehnert, Boysen, Schneider, et al., 2000). The lowtemperature rhombohedral structure can be described as a perovskite structure with an antiphase rotation of the AlO6 octahedra. Using diffraction analysis, this can be observed from a subtle splitting of the main peaks, attributed to a distortion from the cubic structure. This splitting is very small, however, and can be observed only in high-resolution experiments. There have been many investigations of this phase transition (Bueble, Knorr, Brecht, & Schamhl, 1998; Geller & Bala, 1956; Geller & Raccah, 1970; M€ uller, Berlinger, & Waldner, 1968; O’Bryan, Gallagher, Berkstresser, & Brandle, 1990; Scott, 1969), including recent ones using neutron powder methods to examine the thermal evolution of the structure (Chakoumakos, Schlom, Urbanik, & Luine, 1998; Hayward et al., 2005; Howard, Kennedy, & Chakoumakos, 2000). The rhombohedral structure at room temperature can be described as pseudocubic with lattice parameters of 3.791 Å with aluminum (Al) atoms located at the corners and the lanthanum (La) atoms at the center of the cube (Hayward et al., 2005). This compound can be described in the ionic limit as La3þAl3þO2 3 . The charge states in the LaAlO3 are positive for La3þO2 and negative for 2þ 2 Al3þO2 and Ti4þO2 2 . On the contrary, for SrTiO3, the Sr O 2 layers have a neutral charge. Because perovskite heterostructures’ AO-BO2 stacking sequence is maintained along the [001] direction, a polarity discontinuity arises at the LaAlO3eSrTiO3 interface. The Ti ion compensates for mixed valence charge, and this results in the net transfer of electrons (nominally 0.5 electron per twodimensional [2D] unit cell) from LaAlO3 to SrTiO3 across the interface (Figure 1.3(a)). The extra electrons at the LaOetitanium dioxide (TiO2) interface were confirmed by metallic conductivity and Hall measurements by Ohtomo & Hwang (2004, 2006). The interface charges at this “n-type” interface are induced by electronic reconstruction, conceivably through mixed-valence Ti states (Ti4þ to Ti3þ) that place extra electrons in the SrTiO3 conduction band. The analogous construction of the aluminum oxide (AlO2)estrontium oxide (SrO) interface, as shown in Figure 1.3(b), must now acquire extra holes per 2D unit cell to maintain charge neutrality. This interface is formally called “p-type.” Electrically, however, this interface is insulating (Ohtomo & Hwang, 2004). Because this p-charging is still conceivable and there are no available mixed-valence states to accommodate the holes, an atomic reconstruction is required and most likely is formed by the introduction of oxygen vacancies.

1.4.1.2

Control of substrate surface termination

A prerequisite to obtaining high-quality interfaces between SrTiO3 and LaAlO3—or any other hetero layer—with control on the atomic scale is that the starting surface of the substrate has to be atomically smooth. For perovskites, however, the substrate surface obtained by cleaving or cutting typically consists of an equal amount of AO- and BO2-terminated domains separated by half unit cell steps (Figure 1.4(a)).

8

Epitaxial Growth of Complex Metal Oxides

(a)

(b)

(LaO)+

)-

(AIO2)-

(LaO)+

(LaO)+

)-

(AIO2)-

(LaO)+

(SrO)0

(TiO2)0

(TiO2)0

(SrO)0

(SrO)0

(TiO2)0

(TiO2)0

(AIO2 (AIO2

(SrO)0

Figure 1.3 Schematic models of the two possible interfaces between strontium titanate (SrTiO3) and LaAlO3 in the (001) direction. The resulting (LaO)þe(TiO2)0 (a) and (AlO2)e(SrO)0 interfaces (b) show the composition and the ionic charge state of each layer. The schematic models are taken from Ohtomo and Hwang (2004). Reprinted by permission from Macmillan Publishers Ltd: Nature 427: 423e426, © 2004.

Figure 1.4 Surface analysis of strontium titanate substrates by atomic force microscopy (AFM). AFM micrographs (top) and surface roughness analysis results (bottom) of an as-received (ethanol cleaned) double-terminated surface (a), a chemically and thermally treated single titanium dioxideeterminated surface (b), and a pulsed laser deposited single SrO-terminated surface (c). From M. Huijben et al. (2009). Copyright 2009, Wiley-VCH Verlag GmbH & Co. KGaA, Weinheim.

Thin film growth on these as-received substrates results in an interface with a mixture of LaOeTiO2 and AlO2eSrO interfaces. To fabricate a single type of interface, the initial substrate has to be single terminated by either AO or BO2. A polar material with charged layers becomes problematic. For example, surface studies of singlecrystalline LaAlO3 substrates have given contradicting results about surface termination

Growth studies of heteroepitaxial oxide thin films using RHEED

9

by a thermal treatment (Kim et al., 1999; Wang & Shapiro, 1995; Yao, Merrill, Perry, Marton, & Rabalais, 1998). Furthermore, surface reconstruction (Lanier et al., 2007) is present for La3þAl3þO3 crystals because of the presence of a polar surface of (LaO)þ or (AlO2). Similar polar surfaces are always present for various other substrates, such as Ndþ3Gaþ3O3, Kþ1Taþ5O3, and Dyþ3Scþ3O3. As a result, only nonpolar SrTiO3 singlecrystalline substrates, expected to render single-terminated surfaces with charge-neutral single surface terminations of either TiO2 or SrO, are therefore used to investigate atomically controlled LaAlO3eSrTiO3 interfaces. A chemical route was suggested to achieve this single termination for SrTiO3 substrates, whereby a chemical treatment and a thermal treatment were combined (Kawasaki et al., 1994). The etching mechanism was later analyzed in more detail, and a two-step chemical treatment was developed to form perfectly crystalline TiO2-terminated SrTiO3 surfaces (Koster, Kropman, Rijnders, Blank, & Rogalla, 1998) (Figure 1.4(b)). A few refinements with an additional etch procedure were recently reported (Ohnishi et al., 2004); this was investigated by high-resolution synchrotron radiation photoemission spectroscopy and shown to result in very stable TiO2-terminated surfaces (Kobayashi et al., 2004). To date, no chemical treatments have been reported to produce opposite single-terminated SrO surfaces, whereas heat treatment of as-received SrTiO3 substrates usually results in mixed termination. Single-terminated SrO surfaces can, however, be obtained by depositing an SrO monolayer on a single-terminated TiO2 surface. Epitaxial growth of SrO has occurred in a layer-by-layer mode for molecular beam epitaxy (Migita, Kasai, & Sakai, 1996), as well as for PLD (Takahashi et al., 2002) at relatively low temperatures (400e500  C). For SrO monolayer growth, interval PLD has to be applied at normal SrTiO3 deposition temperatures (850  C) (Koster, 1999; Koster, Rijnders, Blank, & Rogalla, 2000). In this deposition technique the total number of laser pulses for one monolayer has to be provided rapidly (50 Hz) to stabilize the correct SrO layer without multilevel islands.3 This results in crystalline SrO-terminated SrTiO3 surfaces with perfectly straight step ledges (Figure 1.4(c)).

1.4.1.3

Growth of atomically controlled interfaces

PLD has been used for the homoepitaxial growth of SrTiO3 (Eres et al., 2002; Kawasaki et al., 1994; Koster, Kropman, Rijnders, Blank, & Rogalla, 1998; Lee et al., 2000; Lippmaa, Nakagawa, Kawasaki, Ohashi, & Koinuma, 2000; Li et al., 2004; Song & Jeong, 2003) as well as LaAlO3 (Jia, Findikoglu, Reagor, & Lu, 1998; Lu, Jia, & Findikoglu, 1999). High-quality heteroepitaxial growth of SrTiO3 and LaAlO3 also has been obtained by PLD, but before 2004 this rarely occurred by combining both materials. An example of where it was applied is the growth of SrTiO3 thin films on LaAlO3 substrates to produce electronically tunable microwave devices, such as resonators, filters, and phase shifters (Bouzehouane et al., 2002; Dalberth, Stauber, Price, Rogers, & Galt, 1998; Petrov, Carlsson, Larsson, 3

Concerning the deposition conditions, a single-crystal SrO target is ablated with an energy density of 1.3 J/cm2. During growth, the substrate is held at 850  C in an oxygen environment at 0.13 mbar.

10

Epitaxial Growth of Complex Metal Oxides

Friesel, & Ivanov, 1998; Treece, Thompson, Mueller, Rivkin, & Cromar, 1997; Van Keuls et al., 1997; Yamada et al., 2005). Following the initial publications of Ohtomo and Hwang (2004, 2006), various groups have grown thin films of LaAlO3 by PLD on single-terminated SrTiO3 substrates to investigate the properties of the two possible heteroepitaxial interface configurations. To obtain well-controllable layer-by-layer growth, use of a single-crystal LaAlO3 target is suitable. Most groups have used a krypton fluoride excimer laser at a repetition rate of 1 Hz and a laser fluency of 1e2 J/cm2. A typical deposition temperature range is 750e850  C, whereas the oxygen pressure can be varied between 106 and 103 mbar to control the oxidation level; this is discussed in more detail in Section 1.4. The oxygen pressure has to be limited to this range to ensure the quality of the interface structure; a transition from 2D layer-by-layer growth to island growth was observed for oxygen pressures of 102 mbar and higher. The annealing procedure after the thin film growth plays an important role in the fabrication process. To carefully study the oxidation level of the LaAlO3eSrTiO3 heterostructures, the oxygen pressure has to be kept at the deposition pressure during cooldown. On the other hand, annealing under high oxygen pressure has been used by various groups to fully oxidize the fabricated heterostructure and, presumably, to remove all oxygen vacancies. It must be noted that the oxygen pressure during growth determines the growth mode as well as the oxidation level. Subsequent exposure to high-pressure molecular oxygen diminishes the number of oxygen vacancies, but full stoichiometry is hard to achieve. Surface quality was monitored by RHEED during the growth of LaAlO3 thin films on single TiO2-terminated and single SrO-terminated SrTiO3 substrates. The fluctuations in RHEED intensity during the initial growth of the first unit cells for both types of surface terminations are shown in Figure 1.5. Oscillations in the RHEED intensity can be observed in both cases, which indicate 2D layer-by-layer growth of LaAlO3 for both types of SrTiO3 surface terminations. The clear 2D spots in the RHEED pattern, shown in the insets, confirmed this growth behavior. The sharp decrease in RHEED intensity for the first LaAlO3 unit cells in both cases can be explained by the difference in the optimal diffraction conditions for both materials because the RHEED monitoring was initially aligned with the SrTiO3 unit cell of the substrate. The difference in c-axis length between the initial SrTiO3 unit cell (w3.905 Å) and the deposited LaAlO3 unit cell (w3.791 Å) requires a new alignment of the RHEED monitoring for optimal surface analysis. For well-aligned RHEED analysis, the 2D layer-by-layer growth of individual LaAlO3 unit cells can be observed up to thicknesses of w20 nm. The oscillations in RHEED intensity were investigated to indicate the growth of individual unit cells. The constant number of laser pulses, which is required to form one unit cell, and the constant RHEED intensity at the maximum oscillations suggest the growth of individual unit cells of LaAlO3 with a constant surface roughness (Figure 1.6(a)). This was confirmed by fluctuations in the full width at half maximum of the specular RHEED spot, which exhibit identical oscillations but was inverted compared with the specular spot amplitude. The constant full width at half maximum value after

Growth studies of heteroepitaxial oxide thin films using RHEED

(b)

100

RHEED intensity (a.u.)

RHEED intensity (a.u.)

(a)

80 60 40 20 0 0

50

100 150 Time (s)

200

11

100 80 60 40 20 0 0

50

100 150 Time (s)

200

Figure 1.5 Monitoring of the reflection high-energy electron diffraction (RHEED) intensity during the initial growth of LaAlO3 unit cells on single-terminated strontium titanate substrates with a titanium dioxide (TiO2)eterminated surface (a) and a strontium oxide (SrO)-terminated surface (b). The insets show the RHEED patterns; clear two-dimensional RHEED spots after growth of 26 unit cells of LaAlO3 are seen. a.u., arbitrary units. From M. Huijben et al. (2009). Copyright 2009, Wiley-VCH Verlag GmbH & Co. KGaA, Weinheim.

the growth of each LaAlO3 unit cell indicates a constant surface roughness without the formation of islands. The low level of surface roughness was confirmed by AFM of LaAlO3 film, 26 unit cells (w10 nm) thick, on a TiO2-terminated SrTiO3 substrate (Figure 1.6(b)). The micrograph and the roughness analysis show smooth terraces with clear unit cell steps.

1.4.2 1.4.2.1

SrRuO3: termination switch Introduction: initial growth of SrRuO3 on SrTiO3

First, depositions at different substrate temperatures were performed. The specular RHEED intensity, recorded during growth at 500, 600, and 700  C, is shown in Figure 1.7. The number of pulses at the first RHEED intensity maximum greatly depends on the deposition temperature. At 700  C, the number of pulses is almost equal to the number of pulses needed to deposit two unit cell layers (as can be seen from the dashed lines in Figure 1.7). The shape of the first RHEED oscillation is asymmetric. A minimum intensity is reached after approximately 27 pulses, equivalent to approximately 1.5 unit cell layers. Figure 1.8(a) shows an AFM micrograph of an SrRuO3 film deposited at 700  C, interrupting the growth at this minimum. Many small islands can be observed. The height of these islands is determined to be approximately 4 Å (wc-lattice parameter of the pseudo-cubic cell). During subsequent growth, the islands coalesce until an almost closed layer is obtained (Figure 1.8(b)). Here, the growth is interrupted at the first RHEED intensity maximum after 37 pulses, equivalent to the deposition of approximately two unit cell layers.

12

Epitaxial Growth of Complex Metal Oxides

Intensity (a.u.)

(a) 70 60 50 40 30

FWHM (a.u.)

8.0 7.5 7.0 6.5 6.0

0

50

(b)

100

150

Time (s) 3.0

nm

2.0 1.0 0

0

0.5

1.0

1.5

2.0

μm

Figure 1.6 Reflection high-energy electron diffraction intensity and full width at half maximum (FWHM) monitoring during growth of LaAlO3 unit cells on a strontium titanate (SrTiO3) substrate (a) and surface analysis by atomic force microscopy of a 26-unit cellethick LaAlO3 thin film on a titanium dioxideeterminated SrTiO3 substrate (b, left). The roughness analysis (b, right) shows smooth terraces with unit cell steps. a.u., arbitrary units. From M. Huijben et al. (2009). Copyright 2009, Wiley-VCH Verlag GmbH & Co. KGaA, Weinheim.

1.4.2.2

Discussion

Based on the RHEED and AFM data, the following conclusions can be made: 1. At 700  C, the first RHEED minimum is observed after deposition of material needed for approximately 1.5 unit cell layers, whereas the first RHEED maximum is observed after deposition of material needed for approximately two unit cell layers. 2. The position of the minimum and the maximum RHEED intensities depends on the growth temperature. 3. AFM micrographs taken after interrupting growth at 700  C at the first RHEED minimum show islands only about 4 Å high. Step heights comparable to 1.5 unit cells have never been observed.

Growth studies of heteroepitaxial oxide thin films using RHEED

(a)

13

120

700 ºC

100 80 60 40 20

(b)

0

RHEED intensity (a.u.)

120

(c)

600 ºC

100 80 60 40 20 0 120

500 ºC

100 80 60 40 20 0

0

10

20

30

40

50

60

70

Time (s)

Figure 1.7 Reflection high-energy electron diffraction intensity of specular spots recorded during the growth of SrRuO3 on titanium dioxideeterminated strontium titanate using deposition temperatures of 700  C (a), 600  C (b), and 500  C (c). The vicinal angle b was 0.11 . The distance between the dashed lines corresponds to the time necessary for deposition of one pseudo-cubic unit cell layer at 700  C. a.u., arbitrary units.

Combining the experimental results, the following initial growth behavior is proposed, visualized in Figure 1.8: 1. During the growth of SrRuO3, no RuO2 termination occurs, that is, SrO is the termination atomic layer.

14

Epitaxial Growth of Complex Metal Oxides

Figure 1.8 Atomic force microscopy micrographs of SrRuO3 deposited on titanium dioxide (TiO2)eterminated strontium titanate (SrTiO3) at a deposition temperature of 700  C (top) and schematic cross-sectional views (bottom). Corresponding to the first reflection high-energy electron diffraction (RHEED) minimum, 27 pulses were applied (a); corresponding to the first RHEED maximum, 40 pulses were applied (b). The arrows indicate the SrTiO3eSrRuO3 interface. RuO2, ruthenium(II) oxide; SrO, strontium oxide. 2. At 700  C, after deposition of the equivalent of two unit cell layers, a closed layer of SrRuO3 with SrO termination is observed (Figure 1.8(b)). 3. After the completion of this SrRuO3 unit cell layer, subsequent stoichiometric deposition leads to layer-by-layer growth of the unit cell, indicated by the equidistant RHEED intensity oscillations after the first maximum (Figure 1.7(a)). 4. Depending on the average terrace length, the growth mode converts to a steady-state growth mode.

The desorption of RuO2 is most likely due to the relatively low oxidation power at the PLD conditions used. Under these conditions, the RuO2 terminating layer cannot be stabilized, leading to the formation of RuxOy. Among the ruthenium (Ru) oxides, only RuO2 is stable (Lee, Kim, Min, & Choh, 1995), whereas other forms, such as Ru2O, are volatile (Vijay, Desu, & Pan, 1993).4 The temperature has a large influence on the evaporation of the RuxOy (Jia et al., 1996), causing the distributed transition in the oscillation period at lower temperatures (Figure 1.7).

4

Deposition of pure RuO2 from a stoichiometric target did not result in a measurable deposition rate at the mentioned deposition conditions of SrRuO3. Many groups have proposed RuO2 as a metallic-like conducting oxide for use as electrode material. Using deposition techniques such as sputter deposition and PLD, the deposition temperature is limited to 600e700  C. Higher temperatures lead to re-evaporation of the deposited material.

Growth studies of heteroepitaxial oxide thin films using RHEED

Figure 1.9 Reflection high-energy electron diffraction intensity of a specular spot recorded during the growth of SrRuO3 on strontium oxideeterminated strontium titanate using a deposition temperature of 600  C. The distance between the dashed lines is the same as in Figure 1.7. a.u., arbitrary units.

RHEED intensity (a.u.)

120 100 80 60 40 20 0 0

15

10

20

30 40 Time (s)

50

60

70

To verify the above-proposed termination conversion, deposition was performed on SrO-terminated5 SrTiO3 substrates. From Figure 1.9 it can be seen that the first maximum of the RHEED oscillation is indeed completed after deposition of one unit cell layer. Here, the termination conversion and the accompanying evaporation of RuxOy at the initial growth stage of SrRuO3 do not occur.

1.4.3 1.4.3.1

Lead titanate Introduction

Lead titanate (PbTiO3) is known as a ferroelectric material. It is frequently used in the thin film growth of ferroelectrics because of its well-controlled layer-by-layer growth. In this chapter we describe PbTiO3 grown on SrRuO3 to understand the initial growth on this oxide electrode material. SrRuO3 (SRO) on SrTiO3 (STO) implies a termination switch from the B site to the A site because of the volatility of Ru. The growth of PbTiO3 on SrRuO3 shows a different initial growth than on SrTiO3. Dependent on the thickness of SRO, it starts with an island-like growth, which converts to layer-by-layer growth after w4 monolayers of material. The island-like growth has a typical fingerprint structure, often seen in ferroelectric domains. The islands have constant heights (w1 nm) and grow only in a lateral direction until they coalesce. Experiments using SRO with different thicknesses indicate that it is not just a termination effect. The literature suggests that the elasticity of SrRuO3 reduces the depolarizing field by adapting a ferroelectric structure. These results indicate that PbTiO3 grows ferroelectric with a predefined polarization direction. The minimum ferroelectric (critical) thickness on SrRuO3 is w1 nm. The impact of this finding is double-sided: What are properties of the polarized PbTiO3 islands? Even more interesting, what are the electrical properties of polarized monolayers of SrRuO3 below the polarized PbTiO3? 5

To obtain this termination, a monolayer of SrO was deposited on TiO2-terminated SrTiO3 before the deposition of SrRuO3.

16

Epitaxial Growth of Complex Metal Oxides

It was already shown that an oxide electrode does have significant advantages over a metallic electrode for oxide ferroelectrics. The conductive oxide SrRuO3 was used as an electrode material. Because a part of this research focuses on multilayer growth, controlling the growth on a monolayer level is important. PbTiO3 is a good research candidate because of its good controllable layer-by-layer growth. Therefore we chose PbTiO3 growth on SrRuO3 to research the nucleation behavior. PbTiO3 grown on single-terminated SrTiO3 already showed a controllable layer-by-layer growth (see above). With the same settings, PbTiO3 is grown on SrRuO3. Differences are expected because SrRuO3 is a conductive oxide (Koster et al., 2012) and has a different type of termination (see previous section). Buffered hydrofluoric acid treated SrTiO3 has a B-site termination (TiO2), whereas SrRuO3 grown under these conditions shows an A-site termination: SrO. As was discussed earlier, this occurs because of the high volatility of RuO at the temperatures used.

1.4.3.2

Stable growth of PbTiO3 on TiO2-terminated SrTiO3 substrates

As mentioned before, the growth of PbTiO3 on TiO2-terminated SrTiO3 shows a controllable layer-by-layer growth, as depicted in Figure 1.10. PbTiO3 is grown on high-frequency-treated SrTiO3 in a 0.13-mbar oxygen environment at 570  C, with a laser frequency of 0.25 Hz. This low frequency is used because of the long relaxation times of PbTiO3 at these temperatures. The deposition settings are listed in Table 1.1. Figure 1.10 indicates layer-by-layer growth, although the first monolayer does show an irregularity seen in more measurements of PbTiO3 on SrTiO3 in this research. Layer-by-layer growth of PbTiO3 on TiO2 terminated SrTiO3 100

2.0 nm

Intensity (a.u.)

90 80 70 60 50

AFM

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500.0 nm

Time (s)

Figure 1.10 Left: Reflection high-energy electron diffraction (RHEED) signal at 0.25 Hz during growth of lead titanate on strontium titanate. Right: Atomic force microscopy image (AFM) of a similar film stopped at the point indicated by the arrow in the RHEED on the left. The AFM image shows typical layer-by-layer growth. Top right corner: RHEED reflection corresponding to the same point showing diffraction spots indicating a two-dimensional surface. a.u., arbitrary units; TiO2, titanium dioxide.

Growth studies of heteroepitaxial oxide thin films using RHEED

Table 1.1

17

Deposition settings for PbTiO3 and SrRuO3

Material

Pressure (mbar O2)

Temperature ( C)

Frequency (Hz)

PbTiO3

0.13

570

0.25

SrRuO3

0.13

700

1

AFM of a similar film deposition stopped at this point is shown (arrow in Figure 1.10), confirming the layer-by-layer growth of the first monolayers. The RHEED reflection of this point is also depicted in Figure 1.10, supporting the 2D growth.

1.4.3.3

Growth of PbTiO3 on an SrO-terminated SrRuO3 buffer layer

For the growth of PbTiO3 on SrRuO3, a thin layer of SrRuO3 (about four monolayers) was deposited on the TiO2-terminated SrTiO3. The deposition settings are depicted in Table 1.1. Because of the volatility of the Ru under these conditions, growth of SrRuO3 at these temperatures implies a termination switch (Vijay, Desu, & Pan, 1993). Figure 1.11 schematically illustrates the different terminations of treated SrTiO3 and of treated SrTiO3 with a couple of SrRuO3 monolayers. PbTiO3 is grown on the surface shown on the right in Figure 1.12 under the same conditions as in Figure 1.10. Figure 1.12 shows the corresponding intensity of the RHEED signal during the growth of PbTiO3 on four monolayers of SrRuO3. The growth of PbTiO3 on the SrO-terminated SrRuO3 is clearly different than on the TiO2-terminated SrTiO3. It starts with an island-like structure, which, after approximately four monolayers of material are deposited, transforms into layer-by-layer growth. This difference is confirmed by the corresponding AFM image in Figure 1.13. The AFM data in Figure 1.13 show the typical island-like structure of the PbTiO3 in Termination differences between bare-treated SrTiO3 and covered with three monolayers of SrRuO3

Ti Sr Ti Sr

Sr Ru Sr Ru Sr Ti Sr Ti Sr

Figure 1.11 Left: Titanium dioxide termination of treated strontium titanate (SrTiO3). Right: Termination switch from the B site to the A site of three monolayers of SrRuO3. This occurs because of the volatility of ruthenium under the deposition conditions.

18

Epitaxial Growth of Complex Metal Oxides Island-like growth of PbTiO3 on SrO terminated SrRuO3 on SrTiO3

100

2.0 nm

AFM 1

90

Intensity (a.u.)

80 70 60 50

AFM 1

40

AFM 2

AFM 3

30 20 10 0

0

100

200

300

400

500

600

700

800

1: Height

900 1000 0.0

500.0 nm

Time (s)

Figure 1.12 Left: Reflection high-energy electron diffraction (RHEED) signal at 0.25 Hz during growth of lead titanate (PbTiO3) on strontium oxideeterminated SrRuO3 on strontium titanate (SrTiO3). Right: Atomic force microscopy (AFM) of a similar film stopped at the point indicated by the arrow in the RHEED plot on the left, showing island-like growth. Top right corner: RHEED reflection corresponding to the same point and showing diffraction spots, indicating a three-dimensional surface. a.u., arbitrary units.

2.0 nm

2.0 nm

AFM 1

0.0

AFM 2

1: Height

500.0 nm 0.0

2.0 nm

AFM 3

1: Height

500.0 nm

0.0

1: Height

500.0 nm

Figure 1.13 Atomic force microscopy (AFM) image of 0.25-Hz growth of lead titanate (PbTiO3) on strontium oxide (SrO)eterminated SrRuO3 on strontium titanate (SrTiO3) taken at different moments in time (indicated in Figure 1.3). These AFM data show the corresponding topography of the island-like growth toward the layer-by-layer growth of PbTiO3 on SrOeterminated SrRuO3 on SrTiO3.

the initial growth on an SrO-terminated surface. This island-like growth shows specific characteristics: 1. The height of the islands in AFM 1 and AFM 2 seems constant and about 1 nm in height; they grow only in lateral size until they coalesce. 2. PbTiO3 shows a typical structure in AFM 1 and 2, where the “islands” seem to repel each other. A balance occurs between island length and spacing between the islands. In the literature this structure is often referred to as a “fingerprint” or “labyrinth” structure.

After approximately four monolayers of material the PbTiO3 growth seems to transform into a layer-by-layer mode, whereby the initial fingerprint structure disappears.

Growth studies of heteroepitaxial oxide thin films using RHEED

19

Comparison of interfaces between PbTiO3 on SrTiO3 and on SrRuO3

PbTiO3

PbTiO3

SrTiO3

SrRuO3

2 nm

2 nm

Figure 1.14 Transmission electron microscopy of 0.25-Hz growth of lead titanate (PbTiO3) on titanium dioxide (TiO2)eterminated strontium titanate (SrTiO3) (left) and PbTiO3 on strontium oxideeterminated SrRuO3 on SrTiO3 (right). No structural differences in the PbTiO3 appear.

To determine whether this typical initial growth of PbTiO3 on SrRuO3 shows a different interface than on SrTiO3, transmission electron microscopy images were taken of the interfaces in both cases. Figure 1.14 shows high magnification of the two interfaces. With the available resolution and contrast, no significant differences between the two interfaces can be detected. To understand the growth behavior, the dependability of the surface termination has to be investigated in more detail. The next two paragraphs describe two experiments where termination-dependent growth is researched in more detail.

1.4.3.4

Growth of PbTiO3 on direct termination switch by a TiO2 monolayer on SrRuO3

As described above, PbTiO3 on SrRuO3 has a typical initial growth. To determine whether this growth is only termination dependent, another experiment was conducted. The termination of SrRuO3 on SrTiO3 was switched from SrO to TiO2 with the deposition of one monolayer of TiO2. TiO2 itself has a tendency toward island-like growth, forming three-dimensional (3D) islands (Koster, 1999; Sakama, 2006). Furthermore, the normally used deposition temperature of 850  C is undesired for the stability of the SrRuO3 layer. A 2D monolayer of TiO2 is therefore deposited at 700  C with interval PLD (Koster, 1999; Rijnders, 2001). Interval PLD is used to increase the nucleation states at the start of the growth and force a 2D growth at this temperature. Furthermore, TiO2 was grown under the same conditions as SrRuO3 with 68 pulses with a laser frequency of 60 Hz. Figure 1.15 shows the RHEED intensity signal of TiO2 depositing using interval PLD with the corresponding 2D RHEED pattern. Figure 1.15 indicates that when the termination of SrRuO3 is switched toward TiO2, the growth of PbTiO3 is in the layer-by-layer mode again. The results indicate that the growth behavior of PbTiO3 is indeed termination dependent. Although the growth mode is similar to that of PbTiO3 on SrTiO3, one can see a difference. Compared with Figure 1.10, the relaxation peaks at every laser pulse seen in Figure 1.15 are more pronounced. In fact, they resemble the relaxation peaks of the growth of PbTiO3 on SrRuO3, as shown in

20

Epitaxial Growth of Complex Metal Oxides

100

100

90

90

80

80

Intensity (a.u.)

Intensity (a.u.)

PbTiO3 shows layer-by-layer growth on SrRuO3 after termination switch by TiO2

70 60

TiO2

50 40

70 60

40

30

30

20

20

10

10

0

0

50

Time (s)

100

150

PbTiO3

50

0

0

50

100

150

200

250

300

Time (s)

350

400

450

500

Figure 1.15 Left: Reflection high-energy electron diffraction (RHEED) intensity of interval pulsed laser deposition of titanium dioxide (TiO2) (68 pulses at 60 Hz) on SrRuO3. The recovery of the TiO2 monolayer is clearly seen. The inset shows the RHEED reflection of the surface of the TiO2 monolayer. Right: Lead titanate (PbTiO3) growth at 0.25 Hz on the TiO2 monolayer, showing layer-by-layer growth as if it was grown directly on TiO2eterminated strontium titanate. a.u., arbitrary units.

Figure 1.12. No direct clear explanation for this phenomenon can be found. In addition to these results of a direct termination switch between SrO and the TiO2 surface, the opposite result should be found if the termination is the other way around, from TiO2 directly to SrO.

1.4.3.5

Growth of PbTiO3 on direct termination switch by an SrRuO3 monolayer on TiO2-terminated SrTiO3

To strengthen the evidence for surface termination-dependent growth of PbTiO3, an experiment opposite to the one described above was conducted. On a TiO2-terminated SrTiO3 surface, SrRuO3 was grown at 700  C. As mentioned in the previous paragraph, at this temperature the SrRuO3 makes a direct termination switch because of the volatility of RuO2 at these temperatures. Logically, if the growth of PbTiO3 is only termination dependent, it should show the same island-like growth as that which occurs on thicker SrRuO3 layers, as shown in Figure 1.12. Figure 1.16 shows the growth of one monolayer of SrRuO3 followed by the growth of PbTiO3. Contrary to the expected result, RHEED indicates layer-by-layer growth comparable to growth on a TiO2 terminated surface. The AFM image taken after growth shows layer-bylayer growth, although the morphology does show a fingerprint structure, as seen before. To look into this result further, the same experiment was repeated with different thicknesses of SrRuO3. Figure 1.17 shows the growth of PbTiO3 on one, two, and three monolayers of SrRuO3 under the same conditions. This figure shows that the

Growth studies of heteroepitaxial oxide thin films using RHEED

21

100

90

90

80

80

70

70

Intensity (a.u.)

Intensity (a.u.)

PbTiO3 shows layer-by-layer growth on 1-ml SrRuO3 100

60 SrRuO3

50 40 30

60

40 30 20

10

10 0

50

100

150

PbTiO3

50

20 0

3.0 nm

0

1.0 μm

0

100

200

300

400

500

600

Time (s)

Time (s)

Figure 1.16 Left: Reflection high-energy electron diffraction intensity of the growth of one monolayer of SrRuO3 switching the titanium dioxide (TiO2) strontium titanate (SrTiO3) surface into a strontium oxide termination. Right: Lead titanate (PbTiO3) growth at 0.25 Hz on an SrRuO3 monolayer, showing layer-by-layer growth comparable to growth on a TiO2 terminated surface. In the inset AFM shows the surface topography and RHEED the corresponding diffraction pattern both recorded after deposition. Comparison of PbTiO3 grown on 1, 2 and 3 monolayers of SrRuO3

Intensity (a.u.)

On 3-ml SrRuO3

On 2-ml SrRuO3

On 1-ml SrRuO3 0

100

200

300 Time (s)

400

500

600

Figure 1.17 Lead titanate (PbTiO3) grown under the same conditions on one, two, and three monolayers of SrRuO3 with different thicknesses. A transformation of growth mode can be observed around three monolayers of SrRuO3. a.u., arbitrary units.

transformation of layer-by-layer growth to island-like growth starts after approximately three monolayers of SrRuO3. The results in Figures 1.15 and 1.17 therefore indicate that growth is not purely dependent on termination; apparently there is another reason for these two different growth modes of PbTiO3 seen in Figures 1.10 and 1.12. For this discussion, verifying whether ferroelectric properties could play a role in the growth dynamics observed is crucial.

22

Epitaxial Growth of Complex Metal Oxides

1.4.3.6

Ferroelectric Curie temperature at the growth temperature of PbTiO3 on SrTiO3

Ferroelectric properties might influence the growth mode. The bulk ferroelectric transition temperature of PbTiO3 is 490  C. This implies that PbTiO3 is cubic at a growth temperature of 570  C, not ferroelectric. Thin film properties can, however, change due to strain and stress. The misfit strain is determined (Pertsev, 1998) by um ¼

ðb  a0 Þ b

where b is the substrate lattice parameter and a0 the cubic cell constant of the freestanding film. In this case PbTiO3 (cell parameters (Shirane, Hoshino, & Suzuki, 1950): 3.89, 3.89, 4.14 Å) is grown on an SrTiO3 (3.90, 3.90, 3.90 Å) substrate, resulting in a compressive strain of PbTiO3 of 0.3% at room temperature and 0.9% at growth temperature (SrTiO3: 3.93, 3.93, 3.93 Å and PbTiO3: 3.97, 3.97, 3.97 Å at 570  C) (Biegalski et al., 2005). A shift of the Curie transition temperature can be expected with these compressive strains. To determine the transition temperature of these specific PbTiO3 films grown at the specified conditions on SrTiO3, temperaturedependent X-ray diffraction (XRD) was performed at the (001) peak of the SrTiO3 and the PbTiO3. The results are shown in Figure 1.18. The figure also shows the RHEED data for PbTiO3 grown on SrRuO3-buffered SrTiO3 at different temperatures. The XRD data indicate that the transition temperature shifted from 490  C toward w600  C. The PbTiO3 grown at 570  C could therefore still be ferroelectric. 4.04

Transition temperature of PbTiO3 on SrTiO3

4.03

610°C

c-axis (Å)

4.02

620°C 630°C 630°C

4.01

640°C

4 3.99

650°C

3.98 3.97 300

Growth temperature 350

400

450

500 550 600 Temperature (°C)

650

700

750

Figure 1.18 C-axis length of lead titanate (PbTiO3) grown on strontium titanate (SrTiO3) with an SrRuO3 buffer layer versus temperature. Insets show reflection high-energy electron diffraction reflection spots of PbTiO3 grown on SrTiO3 with an SrRuO3 buffer layer after 7 min of deposition at these temperatures.

Growth studies of heteroepitaxial oxide thin films using RHEED

23

To determine whether these XRD data are comparable with the growth observed, another PbTiO3 growth series was conducted. PbTiO3 was grown on SrTiO3 with a SrRuO3 buffer layer under the same conditions as those described in Table 1.1 but at different temperatures. The RHEED reflection spots after 7 min of growth are depicted as insets in Figure 1.18. At 7 min the RHEED reflection spots at 590  C should show island-like spots as in AFM 2 in Figure 1.17. The pictures at 610 and 620  C in Figure 1.18 still represent the same morphology, but at 630  C a transition is observed where the 3D peaks decrease until a 2D surface appears for growth at 650  C. Although a transition in the RHEED data is observed, one has to be careful when drawing conclusions. PbTiO3 stability around these temperatures in these conditions is questionable.

1.4.3.7

Discussion: a summary of results

The results of the experiments described in this chapter indicate that: 1. PbTiO3 grows in a layer-by-layer mode on TiO2-terminated SrTiO3, and PbTiO3 grows in an island-like mode; after approximately four monolayers of material, this is followed by a layer-by-layer mode if grown on SrRuO3 on SrTiO3, which is SrO terminated. 2. The island-like growth observed during PbTiO3 growth on SrRuO3 shows a typical “fingerprint” structure, where the islands have a certain height (w1 nm) and grow only in lateral size until the islands coalesce. Transmission electron microscopy analysis does not show differences in the interface between PbTiO3 grown on SrTiO3 or grown on SrRuO3 on SrTiO3. With a direct termination switch of SrO-terminated SrRuO3 with a monolayer of TiO2, PbTiO3 grows layer by layer, as if it was grown on TiO2-terminated SrTiO3. Only a difference in relaxation behavior is observed. 3. By contrast, a termination switch of TiO2-terminated SrTiO3 with one monolayer of SrRuO3 does not change the growth behavior as seen on a “thick” layer of SrRuO3. PbTiO3 starts to grow in a island-like mode after only three monolayers of SrRuO3. This switches while the termination switch after one monolayer of SrRuO3 is complete.

1.4.4 1.4.4.1

Yttrium barium copper oxide Initial growth of RE123 (RE¼rare earth) on TiO2-terminated SrTiO3

RHEED patterns recorded after the deposition of one and four unit cell layers of yttrium barium copper oxide (YBa2Cu3O7; abbreviated here as Y123) on TiO2-terminated SrTiO3 are depicted in Figure 1.19(a) and (b), respectively. During the initial growth, that is, the deposition of the first few unit cell layers, the sharp 2D spots originating from the SrTiO3 substrate are blurred into streaks. These streaks originate from the unit cell of Y123. The corresponding specular RHEED intensity shows a sharp decrease during deposition of the first unit cell layer, without a large recovery (Figure 1.19(c)). Intensity oscillations are damped after deposition of approximately four unit cell layers. The observed streaks in the RHEED pattern and loss of RHEED

24

Epitaxial Growth of Complex Metal Oxides

(a)

RHEED intensity (a.u.)

(c)

(b)

250 200 150 100 50 0

0

20 Time (s)

40

60

80

Figure 1.19 Reflection high-energy electron diffraction (RHEED) patterns recorded after deposition of one (a) and four (b) unit cell layers of yttrium barium copper oxide on titanium dioxideeterminated strontium titanate and the specular RHEED intensity recorded during initial growth (c). a.u., arbitrary units. The arrows in both images indicate diffraction spots corresponding to a so-called 3D transmission pattern.

intensity are indications of a roughened surface, that is, the step density at the surface is increased.6 This increase can be seen in the AFM micrographs depicted in Figure 1.20(a) and (b), showing the surface morphology of 4- and 20-unit cellethick Y123 layers, respectively. Clearly, the growth proceeds by 2D nucleation and subsequent spreading into larger 2D islands. In addition to steps with heights corresponding to the c-axis lattice parameter, smaller steps, corresponding to subeunit cell heights, are visible. Note that the average island size is small compared with the terrace width of the vicinal substrate. The vicinal substrate steps are, therefore, not responsible for most 6

Bulk lattice parameters of substrate and film influence the specular RHEED intensity because of the shallow penetration depth of the electron beam and multiple scattering effects. The angle of incidence used in the experiments is set to w1 . At this angle, the specular RHEED intensity is maximum on SrTiO3. A small increase (w10%) of the specular RHEED is always observed by a small change in the angle of incidence of the electron beam after deposition.

Growth studies of heteroepitaxial oxide thin films using RHEED

(a)

25

(b)

10 0 n m

100 n m

Figure 1.20 Atomic force microscopy micrographs of a 4- (a) and 20-unit cellethick (b) yttrium barium copper oxide film on titanium dioxideeterminated strontium titanate.

of the subeunit cell steps at the surface of the Y123 film. The morphology of the 20-unit cellethick film is a result of 2D nucleation and growth. Here, nucleation and incorporation of ad atoms at step edges proceed on an increasing number of unit cell levels. For example, in Figure 1.20(b), up to four unit cell levels can be seen. In Figure 1.19(a) and (b), in addition to the observed streaks, clear 3D spots are visible (arrows). From the intensity and sharpness of the spots, one can conclude that many small crystallites are formed on top of the Y123 film. From the horizontal and vertical distance between the spots,7 a lattice parameter of w4.3 Å can be determined. From the applied constituents, that is, dysprosium (Dy), barium, copper, and oxygen, only cubic cuprite, Cu2O, can be identified to be responsible for these 3D spots. The lattice parameter of Cu2O is 4.27 Å, which fits with the observed lattice parameter. The position of the spots revealed that the Cu2O precipitates are aligned with the [010] Cu2O//[010] Y123, and [001] Cu2O//[001] Y123. Assuming a perovskite stacking sequence, stoichiometric deposition of Y123 on TiO2 (B-site)eterminated SrTiO3 should lead to B-site-terminated Y123, for example, either a CuO2 plane or the CuO chain layer. Furthermore, a symmetric unit cell with stacking sequence bulkebarium oxide (BaO)eCuO2eYeCuO2eBaOeCuOxebulk leads to the CuO chain layer as the terminating atomic layer. As mentioned earlier, it is expected that the small oxidation power at standard PLD conditions gives rise to an unstable CuOx chain layer. As a consequence, BaO becomes the terminating atomic layer of Y123 and Cu2O precipitates are formed. The RHEED patterns recorded after the growth of SrO and, subsequently, four unit cell layers of Y123 are depicted in Figure 1.21(a) and (b), respectively. Now, only streaks originating from the Y123 layer can be seen. No indication of Cu2O formation is found, that is, no 3D transmission spots are observed. During the deposition of Y123, the supplied CuOx is incorporated in the unit cell layer and A-site termination is preserved. Still, a large decrease in the RHEED intensity is observed, without large recovery, indicating an increased step density (Figure 1.21(c)). In fact, the 7

Not all spots are observed in the RHEED pattern. Some spots are extinct and therefore not visible.

26

Epitaxial Growth of Complex Metal Oxides

(b)

(a)

RHEED intensity (a.u.)

(c) 150

100

50

0

0

20

40

60

80

Time (s)

Figure 1.21 Reflection high-energy electron diffraction (RHEED) patterns recorded after deposition of 1 ml strontium oxide (a) and, subsequently, four unit cell layers of yttrium barium copper oxide(Y123) (b) on titanium dioxideeterminated strontium titanate and the specular RHEED intensity recorded during Y123 growth (c). a.u., arbitrary units.

RHEED intensity resembles the RHEED intensity recorded during growth on TiO2-terminated SrTiO3. So, no difference in the growth mode is expected. Here, the occurrence of two stacking sequences is also likely, that is, bulk substratee SrOeTiO2eSrOeCuOeBaOeCuO2eDyeCuO2eBaOebulk film and bulk substratee SrOeTiO2eSrOeCuO2eDyeCuO2eBaOeCuOeBaOebulk film. The probability of equal nucleation during the initial growth leads to an increase in the step density and the formation of antiphase boundaries. From both RHEED and AFM measurements we conclude that the growth mode on SrO-terminated SrTiO3 resembles the growth mode on TiO-terminated SrTiO3.

1.5

Conclusions

With the examples discussed in this chapter we give a broad overview of the application of in situ RHEED during the epitaxial growth of heterostructures of complex oxides.

Growth studies of heteroepitaxial oxide thin films using RHEED

27

Acknowledgments The author acknowledges Dave H. A. Blank for discussions and help with graphics.

References Abramov, Y. A., Tsirelson, V. G., Zavodnik, V. E., Ivanov, S. A., & Brown, I. D. (1995). Acta Crystallographica, Section B, 51, 942. Biegalski, M. D., Haeni, J. H., Trolier-McKinstry, S., Schlom, D. G., Brandle, C. D., Brandle, C. D., & VenGraitis, A. J. (2005). Thermal expansion of the new perovskite substrates DySCO3 and GdSCO3, 20, 952e958. Blank, D. H. A., Rijnders, A. J. H. M., Koster, G., & Rogalla, H. (1998). Applied Surface Science, 129, 633. Bouzehouane, K., Woodall, P., Marcilhac, B., Khodan, A. N., Crété, D., Jacquet, E., et al. (2002). Applied Physics Letters, 80, 109. Bueble, S., Knorr, K., Brecht, E., & Schamhl, W. W. (1998). Surface Science, 400, 345. Chakoumakos, B. C., Schlom, D. G., Urbanik, M., & Luine, J. (1998). Journal of Applied Physics, 83, 1979. Christen, H. M., & Eres, G. (2008). Recent advances in pulsed-laser deposition of complex oxides, Journal of Physics: Condensed Matter, 20, 264005. Cowley, R. A. (1964). Physical Review, 134, A981. Dalberth, M. J., Stauber, R. E., Price, J. C., Rogers, C. T., & Galt, D. (1998). Applied Physics Letters, 72, 507. Dam, B. (1998). Journal of Materials Research, 9, 217. Eres, G., tischler, J. Z., Yoon, M., Larson, B. C., Rouleau, C. M., Lowndes, D. H., et al. (2002). Applied Physics Letters, 80, 3379. Geller, S., & Bala, V. B. (1956). Acta Crystallographica, 9, 1019. Geller, S., & Raccah, P. M. (1970). Physical Review B, 2, 1167. Hayward, S. A., Morrison, F. D., Redfern, S. A. T., Salje, E. K. H., Scott, J. F., Knight, K. S., et al. (2005). Physical Review B, 72, 054110. Howard, C. J., Kennedy, B. J., & Chakoumakos, B. C. (2000). Journal of Physics, 12, 349. Huijben, M., Brinkman, A., Koster, G., Rijnders, G., Hilgenkamp, H., & Blank, D. H. A. (2009). Structure-property relation of SrTiO3/LaAlO3 interfaces. Advanced Materials, 21, 1665e1677. Jia, Q. X., Findikoglu, A. T., Reagor, D., & Lu, P. (1998). Applied Physics Letters, 73, 897. Jia, Q. X., Song, S. G., Wu, X. D., Cho, J. H., Foltyn, S. R., Findikoglu, A. T., et al. (1996). Applied Physics Letters, 68, 1069. Kawasaki, M., Takahashi, K., Maeda, T., Tsuchiya, R., Shinohara, M., Ishiyama, O., et al. (1994). Science, 266, 1540. Kim, D.-W., Kim, D.-H., Kang, B.-S., Woh, T. W., Lee, D. R., & Lee, K.-B. (1999). Applied Physics Letters, 74, 2176. Kobayashi, D., Kumigashira, H., Oshima, M., Ohnishi, T., Lippmaa, M., Ono, K., et al. (2004). Journal of Applied Physics, 96, 7183. Korte, U. (1997). Physical Review Letters, 78, 2381. Koster, G. (1999). Artificially layered oxides by pulsed laser deposition, ISBN 9036513367. Koster, G. (2011). Reflection high energy electron diffraction. In G. Koster, & G. Rijnders (Eds.), In situ characterization of thin film growth. Woodhead Publishing Ltd., pp. 3e28.

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Epitaxial Growth of Complex Metal Oxides

Koster, G., Klein, L., Siemons, W., Rijnders, G., Dodge, J. S., Eom, C.-B., et al. (2012). Structure, physical properties, and applications of SrRuO3 thin films. Reviews of Modern Physics, 84, 253e298. Koster, G., Kropman, B. L., Rijnders, A. J. H. M., Blank, D. H. A., & Rogalla, H. (1998). Applied Physics Letters, 73, 2920. Koster, G., Rijnders, G., Blank, D. H. A., & Rogalla, H. (2000). Physica C, 339, 215. Lanier, C. H., Rondinelli, J. M., Deng, B., Kilaas, R., Poeppelmeier, K. R., & Marks, L. D. (2007). Physical Review Letters, 98, 086102. Lee, J. Y., Juang, J. Y., Ou, J. H., Chen, Y. F., Wu, K. H., Uen, T. M., et al. (2000). Physica B, 284e288, 2099. Lee, J.-G., Kim, Y. T., Min, S.-K., & Choh, S. H. (1995). Journal of Applied Physics, 77, 5473. Lehnert, H., Boysen, H., Dreier, P., & Yu, Y. (2000). Zeitschrift f€ ur Kristallographie, 215, 145. Lehnert, H., Boysen, H., Schneider, J., Frey, F., Hohlwein, D., Radaelli, P., et al. (2000). Zeitschrift f€ur Kristallographie, 215, 536. Li, Y. R., Li, J. L., Zhang, Y., Wei, X. H., Deng, X. W., & Liu, X. Z. (2004). Journal of Applied Physics, 96, 1640. Lippmaa, M., Nakagawa, N., Kawasaki, M., Ohashi, S., & Koinuma, H. (2000). Applied Physics Letters, 76, 2439. Lu, P., Jia, Q. X., & Findikoglu, A. T. (1999). Thin Solid Films, 348, 38. Migita, S., Kasai, Y., & Sakai, S. (1996). Journal of Low Temperature Physics, 105, 1337. M€ uller, K. A., Berlinger, W., & Waldner, F. (1968). Physical Review Letters, 21, 814. O’Bryan, H. M., Gallagher, P. K., Berkstresser, G. W., & Brandle, C. D. (1990). Journal of Materials Research, 5, 183. Ohnishi, T., Shibuya, K., Lippmaa, M., Kobayashi, D., Kumigashira, H., Oshima, M., et al. (2004). Applied Physics Letters, 85, 272. Ohtomo, A., & Hwang, H. Y. (2004). A high-mobility electron gas at the LaAlO3/SrTiO3 heterointerface. Nature, 427, 423e426. Ohtomo, A., & Hwang, H. Y. (2006). Nature, 441, 120. Pertsev, N. A. (1998). Physical Review Letters, 80(4), 1988. Petrov, P. K., Carlsson, E. F., Larsson, P., Friesel, M., & Ivanov, Z. G. (1998). Journal of Applied Physics, 84, 3134. Ready, J. F. (1963). Applied Physics Letters, 3, 11. Rijnders, G. (2001). The initial growth of complex oxides, study and manipulation, ISBN 90365-1657-9. Rijnders, G., & Blank, D. H. A. (2007). Growth kinetics during PLD, In R. Eason (Ed.), Pulsed laser deposition of thin films: Applications-led growth of functional materials, Wiley, pp. 177e190. Rijnders, A. J. H. M., Koster, G., Blank, D. H. A., & Rogalla, H. (1997). Applied Physics Letters, 70, 1888. Sakama, H. (2006). Thin Solid Films, 515, 535e539. Sakamoto, T., Kawai, N. J., Nakagawa, T., Ohta, K., Kojima, T., & Hashiguchi, G. (1986). Surface Science, 174, 651. Scott, J. F. (1969). Physical Review, 183, 823. Shirane, G., Hoshino, S., & Suzuki, K. (1950). Physical Review, 80, 1105. Shirane, G., & Yamada, Y. (1969). Physical Review, 177, 858. Song, J. H., & Jeong, Y. H. (2003). Solid State Communications, 125, 563. Stoyanov, S., & Michailov, M. (1988). Surface Science, 202, 109. Takahashi, R., Matsumoto, Y., Ohsawa, T., Lippmaa, M., Kawasaki, M., & Koinuma, H. (2002). Journal of Crystal Growth, 234, 505.

Growth studies of heteroepitaxial oxide thin films using RHEED

29

Treece, R. E., Thompson, J. B., Mueller, C. H., Rivkin, T., & Cromar, M. W. (1997). IEEE Transactions on Applied Superconductivity, 7, 2363. Unoki, H., & Sakudo, T. (1967). Journal of the Physical Society of Japan, 23, 546. Van Hove, J. M., Lent, C. S., Pukite, P. R., & Cohen, P. I. (1983). Journal of Vacuum Science & Technology B, 1(3), 741. Van Keuls, F. W., Romanofsky, R. R., Bohman, D. Y., Winters, M. D., Miranda, F. A., Mueller, C. H., et al. (1997). Applied Physics Letters, 71, 3075. Vijay, D. P., Desu, S. B., & Pan, W. (1993). In E. R. Myers, B. A. Tuttle, S. B. Desu, & P. K. Larsen (Eds.), Ferroelectric thin films III (p. 133). Pittsburg, PA: Materials Research Society. Wang, Z. L., & Shapiro, A. J. (1995). Surface Science, 328, 141. Yamanaka, T., Hirai, N., & Komatsu, Y. (2002). American Mineralogist, 87, 1183. Yao, J., Merrill, P. B., Perry, S. S., Marton, D., & Rabalais, J. W. (1998). Journal of Chemical Physics, 108, 1645. Zandvliet, H. J. W., Elswijk, H. B., Dijkkamp, D., van Loenen, E. J., & Dieleman, J. (1991). Journal of Applied Physics, 70, 2614.