Epitaxial growth of superconducting oxides

Epitaxial growth of superconducting oxides

5

H. Yamamoto1, Y. Krockenberger1, M. Naito2 1 NTT Basic Research Laboratories, NTT Corporation, Atsugi, Kanagawa, Japan; 2 Tokyo University of Agriculture and Technology, Koganei, Tokyo, Japan

5.1

Introduction

Epitaxial growth of superconducting films is one of the most extensively studied subjects in the research of complex metal-oxide thin films, especially after the discovery of high-Tc cuprates (Bednorz & M€ uller, 1986). Generally, the setup for the synthesis of high-quality oxide thin films does not relate to physical properties of the oxide, but is rather common among complex metal oxides—whether they are superconducting or not. The synthesis of high-quality superconducting cuprates is an exceptional challenge, especially when it comes to epitaxial thin films. At least three criteria may be identified to support this view: (1) three or more cations are present, and their stoichiometry requires precise tuning, so that undesired phases are suppressed; (2) a fairly strong oxidizing environment is required to oxidize copper, and excessive oxidation may result in films of inferior quality; and (3) off-stoichiometries in oxygen are often detrimental to superconducting properties (Naito, Sato, & Yamamoto, 1997; Sato, Tsukada, Krockenberger, Yamamoto, & Naito, 2011). In spite of those obstacles, significant progress in thin-film growth techniques is driven by the fascination of the very nature of high-Tc superconductivity (Dijkkamp et al., 1987; Naito et al., 1987; Terashima, Iijima, Yamamoto, Bando, & Mazaki, 1988). A steady increase in comprehension of the growth of complex metal oxides and technology not only ameliorated superconducting oxide materials but also other complex-oxide thin films, for example, ruthenates (Siemons et al., 2007). The scope of this chapter is to describe the present status of advanced thin-film growth methods for superconducting complex metal oxides, mainly shedding a spotlight on molecular-beam epitaxy (MBE). With regard to materials, discussions here are particularly dedicated to cuprates, since so far our research efforts have been devoted in large part to cuprates. In fact, the challenges in synthesizing cuprate superconductors might be considered as blueprints for other complex-oxide materials. We prioritize core areas of high-quality thin-film synthesis and subsequently present examples and results of investigations performed on high-quality oxide thin films. Finally, examples are given of mastering materials by design, in particular novel superconductors, by harnessing the powers of thin-film growth methods.

Epitaxial Growth of Complex Metal Oxides. http://dx.doi.org/10.1016/B978-1-78242-245-7.00005-1 Copyright © 2015 Elsevier Ltd. All rights reserved.

96

5.2

Epitaxial Growth of Complex Metal Oxides

Overview of epitaxial growth of superconducting oxides

Epitaxial growth of superconducting complex metal oxides has been carried out by several methods: for example, MBE also referred to as reactive co-evaporation (Bozovic et al., 2009; Ingle, Hammond, & Beasley, 2002; Krockenberger et al., 2008; Krockenberger, Sakuma, & Yamamoto, 2012; Krockenberger, Yamamoto, Mitsuhashi, & Naito, 2012; Krockenberger, Yamamoto, Tsukada, Mitsuhashi, & Naito, 2012; Maritato et al., 2013; Naito et al., 1997; Terashima et al., 1988; Tsukada, Greibe, & Naito, 2002; Tsukada, Kobayashi, Koyama, Watanabe, & Uchinokura, 1992; Tsukada et al., 2005; Tsukada, Noda, Yamamoto, & Naito, 2005; Tsukada, Yamamoto, & Naito, 2006; Yamamoto, Krockenberger, & Naito, 2013, 2014); pulsed-laser deposition (PLD) (Adachi et al., 2002; Kawasaki et al., 1993; Krockenberger et al., 2010; Leca, Blank, Rijnders, Bals, & van Tendeloo, 2006; Maiser et al., 1998; Mukaida & Miyazawa, 1993; Tomaschko et al., 2012; Tsukada, 2004; Venkatesan et al., 1989); and sputtering (Ihara et al., 1999; Jovanovic, Li, & Raffy, 2011; Karimoto, Kubo, Tsuru, & Suzuki, 1996; Michikami, Asahi, & Asano, 1990). Table 5.1 compares pros and cons of the three methods. In addition to those gas-phase epitaxial growth methods, liquid-phase epitaxy (Klemenz & Scheel, 1993; Miura, Hashimoto, Wang, Enomoto, & Morishita, 1997), solid-phase epitaxy (post-deposition methods), for example, electron-beam codeposition (Naito et al., 1987) and metal-organic decomposition (Matsumoto et al., 2009a; Obradors et al., 2004), and vaporeliquidesolid triphase epitaxy (Yun et al., 2002) have also been adopted. By analogy to well-matured semiconductor technologies for the growth of Si and GaAs thin films, MBE might be suitable to prepare the highest-quality thin films. This is, however, not always the case with complex metal oxides because of the difficulties in composition control. In PLD, obstacles due to composition variation can be readily circumvented. The nearly perfect stoichiometric transfer of cations from target to substrate together with a low cost-of-ownership is, to some extent, responsible for its widespread usage. Similarly to PLD, the success of sputtering methods is also partly based on stoichiometric cation transfer mechanisms from target to substrate. However, investigations by Ohnishi, Lippmaa, Yamamoto, Meguro, and Koinuma (2005) led to conclusions showing that an exact stoichiometric transfer from target to substrate is deceptive. Finally, the presumed concept of stoichiometric transfer of cations from target to substrate was rendered ad absurdum and target materials with defined off-stoichiometric compositions have been deployed for the synthesis of high-quality oxide thin films (Hitosugi, Packwood, & Shiraki, 2014; Roberge et al., 2009). Instead of bluntly relying on methods where exact stoichiometric transfer appears to be inadmissible, methods with active composition control allow the growth of crystalline films of unparalleled quality, at least in terms of cation arrangement. In fact, high-precision deposition monitoring-and-control technologies perform unprecedented reproducibility. Among complex-oxide materials, transition-metal oxides, notably cuprates, nickelates, and cobaltates, are well-known for their large

Epitaxial growth of superconducting oxides

Table 5.1

Comparison table of the vapor-phase epitaxy methods Molecular-beam epitaxy

Pulsed-laser deposition

Sputtering

Superconducting properties (Tc, Hc2, Jc)

Excellent

Excellent

Fair

Heating setup

Resistive laser

Resistive, lamp laser

Resistive

Growth rate

Medium

Low (high per pulse)

Low

Composition control

AASa, EIESb, shuttering

Limited by target and growth parameters

Limited by target

Thickness control

Excellent

Excellent

Fair

Surface flatness

Good (if stoichiometric)

Good (droplets)

Good

Ultra-thin film

Possible

Difficult

Poor

Large-area growth

Possible

Bad

Limited

Method

a

97

AAS: Atomic absorption spectroscopy. EIES: Electron impact emission spectroscopy.

b

oxygen-diffusion coefficients. While the oxygen stoichiometry in nickelates and cobaltates definitely rules over macroscopic physical quantities, for example, Curieor Néel temperatures, the influence on macroscopic physical properties of cuprates is subject to subtle changes in oxygen sub-lattice imperfections and/or oxygen offstoichiometry. In cuprates, a well-known example is YBa2Cu3O7ed, where d describes the degree of oxygen deficiency. Particularly, these oxygen ions are loosely bound to the lattice and their mobilities are high even at low thermodynamic excitations. However, it is the occupation of those sites driving electronic responses. Accordingly, for the synthesis of complex transition metal oxides, particularly cuprates, “oxygen engineering” is a crucial factor. Moreover, oxygen engineering is not limited to single crystalline thin films but to a much higher extent on surfaces and interfaces of thin films. Properties of superconducting tunnel junctions composed of cuprates are vulnerable to the junction preparation conditions, partly due to redox reactions often taking place at interfaces (Naito, Yamamoto, & Sato, 2000). Again, subtle variations of oxygen occupations are known to dress the entire electronic response, and one might be advised to caution on experimental results and their interpretations for heterostructures and superlattices, including FET structures (Bollinger et al., 2011; Nojima et al., 2011). Even in the presence of atomically smooth and homogeneous interfaces, a subtle deviation in oxygen arrangement (reversible or irreversible) may alter the overall electronic states in these systems.

98

5.3

Epitaxial Growth of Complex Metal Oxides

Requirements for growth of high-quality complex metal-oxide films by molecular-beam epitaxy (MBE)

The ultimate challenges in the synthesis of complex metal-oxide thin films are reproducibility in quality and physical properties. Fulfilling these challenges requires (1) a precise and individual control of evaporation beam fluxes (elemental discrimination), (2) tunable oxidizing agents, and (3) real-time monitoring of phase formation (Naito et al., 1997; Yamamoto et al., 2014). In addition to these three requirements, the choice of substrate material appears to strongly influence physical properties of the film brought about by strain or lattice mismatching conditions. While the presence of suitable substrate materials is an extrinsic prerequisite to the epitaxial growth of complex metal oxides, it allows for fine tuning of the solubility of doped/substituted ions or promotes stabilization of competing crystallographic phases. The second requirement has been rather easily fulfilled, in terms of designated phase formation, by the development of activated oxygen (O*) and O3 sources. Nevertheless, a proper choice of either O* or O3, along with operation conditions of the sources, may play a dominant role in achieving perfection of the oxygen sub-lattice. The first requirement has been the largest challenge in the growth of complex metal-oxide films by MBE since extra metal atoms do not evaporate from the growing surface of complex-oxide films unlike As or Sb in the growth of IIIeV compounds. Small values of off-stoichiometry can be taken up by interstitial atoms and lattice vacancies that exist as point defects. If the composition mismatch is large enough, grains of one or more impurity phase nucleate are embedded in the single-crystal film of the target phase. Nowadays, real-time in situ composition control in oxidizing environments is achieved by, for example, atomic absorption spectroscopy (AAS) and electron impact emission spectroscopy (EIES) rate control systems, which enable deposition vapor monitoring and on-the-fly feedback to the evaporation source power supply (Matias & Hammond, 2011; Yamamoto et al., 2013). Here, the MBE setup is described taking the MBE system design of the authors as an example. The MBE growth of the multi-cation metal oxide films is carried out by electron-beam co-evaporation from metal sources in a custom-designed ultra-high vacuum (UHV) chamber (Figure 5.1). As we discussed above, one of the most crucial parts of the MBE system is a precise control of evaporation fluxes, and we installed an EIES feedback system (Lu, Blissett, & Diehl, 2008). In the MBE chamber, EIES sensors are measuring optical emission spectra specific to each evaporation chemical element. As the emission intensity is directly proportional to the fluxes, the optoelectronic converted signal is fed back to control the emission current of e-beam evaporators. These EIES sensors are located next to the sample holder, in the same horizontal plane. The sensor head is equipped with a filament, generating thermal electrons for the excitation of atoms. Since the emitted light spectra are characteristic for each element, optical band-pass filters are used for element-specific detection of the optical signals. For example, Figure 5.2(a) shows atomic emission spectra for La and Cu taken by using the EIES sensors. A wavelength of the optical band-pass filters suitable for detection of La or Cu flux is also indicated. Photomultipliers (PMTs), located outside

Epitaxial growth of superconducting oxides

EIES sensor head

99

EIES

Heater

EIES

Substrate

QCM 1

Electrons Light emission

ED RHE Shutter

Filament

O*

Vapor

O3

RF QCM2 Movable PMT detector

6 E-guns

Optical filter PC

Movable

Signal processor and (For real-time feedback control) microcontroller

10 sources

Evaporation source power supply

Figure 5.1 Schematic illustration of our multisource molecular-beam epitaxy setup for the growth of complex metal-oxide thin films. (EIES, electron impact emission spectroscopy sensor; QCM, quartz crystal microbalance; RHEED, reflection high-energy electron diffraction; PMT, photomultiplier tube.) From Yamamoto, H., Krockenberger, Y., & Naito, M. (2014). Proceedings of SPIE 8987, 89870V.

of the vacuum chamber (Figure 5.1), amplify optical signals and convert them into electronic signals. The signal generated by oxygen can be subtracted. Such setup allows simultaneous stabilization of the La and Cu fluxes (Figure 5.2(b)), and, for example, La2CuO4 thin films can be reproducibly grown. Initially, the EIES signals are calibrated against a quartz crystal microbalance (QCM 1), which can be inserted at the substrate position while the substrate holder is lifted. Typical accuracy values are w10
3 Å/s for Cu and 10
1 to 10
3 Å/s for rare-earth elements other than Ce, as Ce does not show any significant emission lines within the optical spectral range used by EIES. Therefore, the flux rate of Ce is controlled by an additional QCM 2, located near the Ce source. Independent verification of the atomic ratio of thin films is occasionally performed using inductively coupled plasma spectrometry (ICPS). Several methods are used to independently determine thickness values of films grown by EIES-assisted MBE, for example, (1) cation supply rates measured by EIES, (2) total number of cations measured by ICPS, (3) period of reflection high-energy electron diffraction oscillation, (4) interval of Laue fringes in the X-ray diffraction (XRD) spectra, and (5) cross-sectional transmission electron microscope images. Overall, the results of those methods coincide well with each other.

100

Epitaxial Growth of Complex Metal Oxides

(a)

(b) 3.0 Shutter open

Shutter close

2.0

La

Rate (Å/s)

Intensity (arb. units)

BP filter (La)

BP filter (Cu)

La 1.0

Cu Cu

200

300

400

500

600

0.0

0

200

400

Wavelength (nm)

600 800 1000 1200 1400 Time (s)

Figure 5.2 (a) Atomic emission spectra for La and Cu taken by using the electron impact emission spectroscopy sensors. Wavelengths of optical band-pass (BP) filters appropriate for element-specific detection of the optical signals from La and Cu fluxes are also shown. (b) Profile of the supplied flux rates before and during the growth of a La2CuO4 film.

5.4

Case studies

In this section, some case studies will be provided with respect to material families. While the vast majority of research efforts have been performed over more than two decades, we are limited to discussions on exemplary case studies and regret any important omissions.

5.4.1

Cuprates with a K2NiF4 (T) structure

Cuprates having the K2NiF4 (T) structure, as represented by La2
xSrxCuO4 and La2
xBaxCuO4, have been prototype materials since the dawn of high-Tc superconductivity (Bednorz & M€ uller, 1986). In those cuprates, the dopant concentration x can be widely varied, and hence, the hole-doped side of the electronic phase diagram of hightemperature superconductors was prepared and has been discussed on the basis of experimental results of those materials. These T-structure materials are also recognized as model materials from the viewpoint of thin-film growth methodology. In addition to their complex crystal structure, 3- or 4-constituent element stoichiometries need to be accurately controlled, which allows for investigations of electronic states over a dopant concentration for x, that is, 0  x  2 (Sato, Tsukada, Naito, & Matsuda, 2000a). In the research of thin-film growth of T-structure cuprates, (1) enhancement of Tc due to epitaxial strain; (2) extra-oxygen-induced/enhanced superconductivity (hole doping by extra oxygen) in La2
xSrxCuO4þd and La2
xBaxCuO4þd, especially, x > 0 in addition to x ¼ 0; and (3) superconductivity in ultrathin films and their thickness dependence are remarkable, and these would not have been accomplished without high-quality thin films.

Epitaxial growth of superconducting oxides

101

Figure 5.3 shows temperature dependences of resistivity for La1.85Sr0.15CuO4 (LSCO), La1.85Ba0.15CuO4 (LBCO), and La2CuO4þd (LCOþ) thin films grown on LaSrAlO4 (LSAO) substrates by MBE. Sharp superconducting transitions are observed at 45, 47, and 56 K for LSCO, LBCO, and LCOþ, respectively (Naito, Tsukada, Greibe, & Sato, 2002). Comparably high Tc is accomplished also by Bozovic, Logvenov, Belca, Narimbetov, & Sveklo, 2002. Those Tc values are significantly higher than the values for corresponding bulk samples, and such Tc enhancements are associated with epitaxial strain (Sato & Naito, 1997; Sato, Tsukada, Naito, & Matsuda, 2000a; Sato et al., 2011). In short, films grown on LaSrAlO4 substrates (a0 ¼ 3.755 Å) are subject to in-plane compressive and out-of-plane tensile strains due to the smaller in-plane lattice constant of (001) LaSrAlO4 than that of LSCO. Detailed analyses indicate that the increase of Tc originates mainly from the c-axis expansion. Conversely, films grown on (001) SrTiO3 (a0 ¼ 3.905 Å) substrates are under in-plane tensile and out-of-plane compressive strain and their Tc (<30 K) is substantially lower than that of bulk specimens. Figures 5.4 (a) and (b) show variations of Tc as a function of Sr/Ba concentrations (x) for the LSCO and LBCO films. Regarding the stoichiometric films grown on LSAO substrates, a well-known dome shape relation between Tc and x is observed while Tc is higher than that of bulk samples at each x due to epitaxial strain. Especially, there is no singularity at x ¼ 0.125, indicating absence of the so-called “1/8 anomaly” (Sato, Tsukada, Naito, & Matsuda, 2000b). Another remarkable point is that, in the under- and zero-doped region (0  x < 0.15), LSCO and LBCO can accept extra oxygen, and Tc values of such samples with extra oxygen (La2
xSrxCuO4þd and La2
xBaxCuO4þd) are comparable to or even higher than those of the optimally

600

100

Resistivity (μΩ cm)

500 50 400 300

0 40

50

60

70

200 La1.85Sr0.15CuO4 La1.85Ba0.15CuO4 La2CuO4+ δ

100 0 0

50

100 150 200 Temperature (K)

250

300

Figure 5.3 Plots of resistivity versus temperature of La1.85Sr0.15CuO4 (LSCO), La1.85Ba0.15CuO4 (LBCO), and La2CuO4þd (LCOþ) films prepared by molecular-beam epitaxy. The inset shows superconducting transitions. Data from Naito, M., Tsukada, A., Greibe, T., & Sato, H. (2002). Proceedings of SPIE 4811, 140.

102

Epitaxial Growth of Complex Metal Oxides

(a)

(b)

La2–xSrxCuO4+δ

80

Thin films on LaSrAlO4 (δ ~ 0) Thin films on LaSrAlO4 (δ > 0) Thin films on SrTiO3 (δ ~ 0) Bulk (δ ~ 0)

60

Tc (K)

Tc (K)

60

40

40

20

20

0

La2–xBaxCuO4+δ 80

Thin films on LaSrAlO4 (δ ~ 0) Thin films on LaSrAlO4 (δ > 0) Bulk (δ ~ 0)

0

0.3 0.1 0.2 Sr composition x

0.4

0

0

0.1 0.2 0.3 Ba composition x

0.4

Figure 5.4 (a) Plots of Tc as a function of Sr composition x for La2
xSrxCuO4þd thin films grown on (001) LaSrAlO4 substrates in comparison with that for bulk samples. Filled and open circles denote d w 0 and d > 0, respectively (Sato et al., 2000a). The dashed line represents the data for bulk polycrystalline samples (Takagi et al., 1989a). (b) Plots of Tc as a function of Ba composition x for La2
xBaxCuO4þd thin films grown on (001) LaSrAlO4 or (001) SrTiO3 substrates in comparison with that for bulk samples. and are for films on (001) LaSrAlO4 substrates with d w 0 and d > 0, respectively. are for films on (001) substrates with d w 0 (Sato et al., 2000b). The dashed lines represent data for bulk polycrystalline samples with d w 0 (Yamada & Ido, 1992).

doped ones (La1.85Sr0.15CuO4 and La1.85Ba0.15CuO4). It is well known that superconductivity can be induced to the end-member La2CuO4 by insertion of extra oxygen ions by, for example, electrochemical oxidation (Arrouy et al., 1996; Chou, Johnston, Cheong, & Canfield, 1993). A homogeneous oxidation of La2CuO4 thin films by ozone exposure was demonstrated by Sato, Naito, and Yamamoto (1997) by fully utilizing large surface-to-volume ratios. The insertion of the extra oxygen takes place also for Sr/Ba doped ones in the underdoped range, leading to a fairly high Tc of w45 K. The statement that Tc can be doubled by epitaxial strain alone (Loquet et al.,1998) had to be reexamined: while Loquet et al. (1998) claimed that Tc of 49.1 K in La1.9Sr0.1CuO4 thin films is almost twice as high as Tc in the corresponding bulk compound (x ¼ 0.1), Bozovic et al. (2002) concluded that Tc cannot be doubled by compressive epitaxial strain alone but by additive effect of extra oxygen. Since the interface between substrate and thin film plays a significant role in epitaxial strain, the Tc increase by epitaxial strain suggests that ultra-thin films are of very high quality. Figure 5.5 shows resistivity-temperature characteristics of La1.85Sr0.15CuO4 thin films with various thicknesses (Sato, Yamamoto, & Naito, 1997). Films of 1 or 1.5 unit cell thickness (UCT) do not show superconducting

Epitaxial growth of superconducting oxides

103

3

La1.85Sr0.15CuO4

Resistivity (mΩ cm)

(× 1/5)

1.5-UCT

2

1-UCT

(× 1/5) 1 2-UCT

4-UCT

3-UCT

140-UCT 0 0

100

200

300

Temperature (K)

Figure 5.5 Temperature dependence of the resistivity for La1.85Sr0.15CuO4 thin films grown on (001) LaSrAlO4 substrates with various unit cell thickness (UCT). The resistivity for the 1 and 1.5-UCT films is scaled down by a factor of 5. Data from Sato, Yamamoto, and Naito (1997).

transition while metallic behavior is observed in the temperature range between 150 and 300 K. Films of 2 UCT or thicker show a superconducting transition, and Tc increases while resistivity values decrease with increasing thickness. Clearly, from these results one can rule out that the enhancement of Tc is an interface phenomenon and this enhancement is subject to deformations of the CuO6 octahedrons brought about by strain. In summary, epitaxial strain induced by differences in lattice constants of substrates is prominent on epitaxial growth of cuprates having the T-structure. Moreover, oxidation control also plays an essential role to clarify inherent properties of the T-cuprates at each cation substitution level.

5.4.2

Cuprates with an Nd2CuO4 (T 0 ) structure

The symmetry of the electronic phase diagram of cuprate superconductors has been under intense debate since the beginning of high-Tc superconductivity research as it provides fundamental input parameters to the theory. In this context, cuprates having the Nd2CuO4 (T0 ) structure, also known as electron-doped cuprates, are of great importance as counter compounds to the T-cuprates (Maple, 1990). Unlike T-cuprates, the choice of substrate materials is considered not to be crucial in the growth of T0 -cuprate thin films. This is because T0 -cuprates readily grow epitaxially on lattice-mismatched substrates though not in a coherent way but relaxed, as shown in Figure 5.6 (Krockenberger, Yamamoto, Mitsuhashi, et al., 2012). Moreover, our

104

Epitaxial Growth of Complex Metal Oxides

0.80

(a)

>104 103

0.78 (103) SrTiO3 0.76

qz(001) (Å–1)

0.74

102 101

(109) Nd1.85Ce0.15CuO4

(b) 0.78 (103) SrTiO3 0.76 0.74 0.72 0.23

(109) Nd2CuO4 qx(001) (Å–1)

0.28

Figure 5.6 Reciprocal space maps of Nd1.85Ce0.15CuO4 (a) and Nd2CuO4 (b) thin films grown on (001) SrTiO3 substrates. From Krockenberger, Y., Yamamoto, H., Mitsuhashi, M., & Naito, M. (2012). Japanese Journal of Applied Physics, 51, 010106. Copyright ©2012 by The Japan Society of Applied Physics.

research efforts highlight the ubiquitous importance of oxygen engineering to seclude the influence of copper coordination from doping dependence on the induction of superconductivity. Further details on Ce-doped/substituted T0 -cuprates are published elsewhere (Sato et al., 2011). The following discussions here are confined to results of superconducting, and undoped T0 -cuprates and related topics (Ikeda, Matsumoto, Yamamoto, Manabe, & Naito, 2011; Krockenberger et al., 2013; Krockenberger, Eleazer, Irea, & Yamamoto, 2014; Krockenberger, Irie, et al., 2014; Krockenberger, Yamamoto, Mitsuhashi, et al., 2012; Krockenberger, Yamamoto, Tsukada, et al., 2012; Matsumoto et al., 2008; Matsumoto et al, 2009a, 2009b, 2009c; Matsumoto, Tsukada, Yamamoto, Manabe, & Naito, 2010; Naito et al., 2008; Noda, Tsukada, Yamamoto, & Naito, 2005; Tsukada et al., 2005; Tsukada, Noda, et al., 2005; Yamamoto, Matsumoto, Krockenberger, Yamagami, & Naito, 2011). In the electronic phase diagram for the cuprate superconductors, the appearance of superconductivity with respect to hole and electron doping is—although this is generally ignored—associated with the asymmetry of copper coordination. While hole doping into five- and/or six-fold coordinated CuO2 planes rapidly induces metallic conduction and hence superconductivity, electron doping into CuO2 planes with

Epitaxial growth of superconducting oxides

105

square-planar-coordinated Cu ions alone is insufficient for the induction of superconductivity. Instead, it is necessary to treat the as-grown specimens of T0 -RE2
xCexCuO4 (RE ¼ rare-earth elements) under reducing environments, irrespective of the dopant (Ce) concentration x. Such a reduction treatment is vital for elimination of defects, for example, removal of excess oxygen at apical sites while simultaneously regular oxygen sites remain occupied (Radaelli, Jorgensen, Schultz, Peng, & Greene, 1994; Takagi, Uchida, & Tokura, 1989; Tokura, Takagi, & Uchida, 1989). As the annealing process is a diffusion process, thin-film samples are advantageous for achieving an ideal and uniform oxygen configuration owing to their large surface-to-volume ratios. Figure 5.7 shows (a) temperature(T)-dependence of resistivity (r), and (b) in situ photoemission spectra, for Nd2
xCexCuO4 thin films with various Ce concentration x. The thin films were grown on (001) SrTiO3 substrates. The post-growth annealing (reduction) conditions were kept almost constant over the entire doping range (Naito et al., 1997). The r(T) responses show a systematic Ce doping dependence, which is roughly consistent with those reported for bulk materials. However, for thin films the superconducting transitions are sharper and resistivity values are lower. It is remarkable that undoped Nd2CuO4 thin films show metallic conduction (dr/dT > 0) for T  180 K. In Figure 5.7(b), the photoemission spectra taken at 300 K (measurements were carried out in 1996) do not vary with respect to x, and the Fermi-edge is observed even in Nd2CuO4. Those experimental results argue against the scenario that the undoped materials are charge-transfer insulators. Subsequently, relationships between Tc and x were investigated for MBE-grown T0 -RE2
xCexCuO4 (RE ¼ La, Pr, Nd, Sm, Eu) films as shown in Figure 5.8(a) (Tsukada, Krockenbeger, et al., 2005). Obviously, the phase diagram of

Nd2–xCexCuO4 Thin films

250 Resistivity (μΩ cm)

× 1/100 200

x=0

(b) x= 0.128 0.131 0.137 0.150 0.166 0.187

150 100

Intensity (arb.units)

(a) 300

Nd2–xCexCuO4 x= 0.20 0.15 0.11

1.0

0.0 EF

0 at~300 K

50 12 0

He I UPS

0

50 100 150 200 250 300 Temperature (K)

8 4 0 Binding energy (eV)

Figure 5.7 (a) Plots of resistivity versus temperature of Nd2exCexCuO4 thin films grown on (001) SrTiO3 substrates with various Ce concentrations. (Data from Naito et al (1997).) (b) In situ ultraviolet photoemission spectra of Nd2
xCexCuO4 thin films with various Ce concentrations. An enlarged view around the Fermi energy is also shown.

106

Epitaxial Growth of Complex Metal Oxides

(a) 40 35

Tc (K)

30

La Pr Nd Sm Eu

25 20 15 10 T

5

Tʹ

0 0

0.05

0.1

0.15

0.2

0.25

0.1 0.15 0.2 Ce content x

0.25

0.3

Tc (K)

20

10

Semiconducting

(b) 30

0 0

0.05

Figure 5.8 (a) Plots of Tc versus x for molecular-beam epitaxy grown T0 -RE2
xCexCuO4 thin films with various RE. The symbols represent Tczero , and dashed lines Tconset . (b) Tc versus x for bulk T0 -Pr2
xCexCuO4 single crystals. (Data from Brinkmann et al. (1995).) The solid and dashed lines represent Tc achieved by improved and standard annealing processes, respectively.

Nd2
xCexCuO4 is idiosyncratic and a strong RE dependence is observed. Figure 5.8(b) compares Tc(x) of Pr2
xCexCuO4 single crystals reduced under standard or improved annealing conditions (Brinkmann, Rex, Bach, & Westerholt, 1995). In the improved annealing procedure, the single crystals were sandwiched by two polycrystalline pellets of the same composition, in order to avoid chemical decomposition by annealing even at 1080  C; this temperature is 130  C higher than that for the standard reduction process. While a very narrow superconducting range close to x ¼ 0.15 is observed for the crystals subjected to the standard reduction procedure, the crystals reduced by the improved method show superconductivity in a much wider x range. Also, Tc monotonically increases with decreasing x down to 0.04, whereas Tc at x ¼ 0.15 stays almost constant, independently of the annealing procedure. Appropriate annealing conditions for the induction of superconductivity themselves are likely doping dependent. If one takes a look at Figure 5.8(a) and (b) with unbiased eyes, one may conclude that undoped T0 -RE2CuO4 compounds are also superconducting per se, but their inherent properties are altered by a still insufficiently optimized reduction recipe. For T0 -La2
xCexCuO4, Tc abruptly

Epitaxial growth of superconducting oxides

107

vanishes at x w 0.04 (Figure 5.8(a)). The disappearance of superconductivity for T0 -La2
xCexCuO4 at x w 0.04 clearly stems from a structural phase transition (from T0 to T-structure). We examined isovalent substitution of La3þ by other RE3þ in order to stabilize the T0 -structure without doping. This isovalent substitution actually led us to the discovery of superconductivity in the undoped T0 -cuprates (Tsukada, Krockenberger, et al., 2005). For example, Figure 5.9 shows temperature dependence of resistivity for La1.9Y0.1CuO4. Superconducting transition is observed at Tc w 21 K and resistivity value at 300 K is r(300 K) ¼ 470 mU cm. The bulk nature of superconductivity has been confirmed by muon spin rotation (Kojima et al., 2014) as well as magnetization and magnetoresistivity measurements (Krockenberger, Yamamoto, Tsukada, et al., 2012). Recently, we succeeded in inducing superconductivity to pristine T0 -La2CuO4, thus without any cation substitution (Krockenberger, Eleazer, Irea, & Yamamoto, 2014; Krockenberger, Irie, et al., 2014). Though the T-structure resembles thermodynamically the most stable allotrope in the La2CuO4 system, geometric rearrangements of O around Cu can be triggered by state-of-the-art thin-film growth methods (Naito, Tsukada, Greibe, & Sato, 2002; Tsukada et al., 2002; Yamamoto et al., 2013), resulting in either pyramidal (T*) or square-planar (T0 ) coordination of copper. Regarding electronic correlations associated with T-, T*-, and T0 -structures of La2CuO4, one can claim an insulator-to-metal phase transition within an identical chemical formula (isomer) without charge carrier doping. The three different phases have been selectively stabilized by combined optimization of growth temperature and substrate, namely, in-plane lattice constant (as): T-, T*-, and T0 -La2CuO4 thin films (100 nm) were coherently grown on (001) SrLaAlO4 (as ¼ 3.755 Å), (110) DyScO3 (as ¼ 3.944 Å), and (110) PrScO3 (as ¼ 4.022 Å) substrates, respectively. As shown in Figure 5.10, the T0 -La2CuO4 thin film is a metal with a superconducting transition at 28 K, quite in contrast to T- or T*-La2CuO4. It is noteworthy that extremely (<0.1 sccm O3) weak oxidization conditions during the growth, together 0.6

ρ (mΩ cm)

T'-La1.9Y0.1CuO4 / (110)YAlO3 0.4

0.2

0.0

0

50

100

150 T (K)

200

250

300

Figure 5.9 Plot of resistivity versus temperature of a 2700-Å thick T0 -La1.9Y0.1CuO4 thin film grown on a (110) YAlO3 substrate, with which low-energy muon spin rotation experiments have been recently carried out.

108

Epitaxial Growth of Complex Metal Oxides

La2CuO4 Thin films

ρ (Ω cm)

106

T

102

10–2

10–6

T*

T'

0

100

200 300 T (K)

400

Figure 5.10 Electric transport behavior of three La2CuO4 isomers (T-, T*-, and T0 -La2CuO4) as a function of temperature. While T- and T*-La2CuO4 show semiconducting behavior, the T0 phase is a superconductor.

with in situ post-growth annealing under UHV, are essential in eliminating defects (removal of irregular oxygen ions). Under such synthesis conditions superconductivity can be induced to T0 -La2CuO4. In contrast to T0 -La2CuO4, for the induction of superconductivity to MBE-grown thin films of other parent compounds, for example, Pr2CuO4 and Nd2CuO4, an elaborate two-step annealing process was found necessary to meet stringent thermodynamic constraints (Figure 5.11). In contrast, for optimally doped Pr1.85Ce0.15CuO4 and

(b) 10–2 Ta

Step I P ao 2 [flowing mixture of N2 and O2]

Tred

Resistivity (Ω cm)

Temperature

(a)

Step II red

P [vacuum]

ta

tred

Time

10–3

or

One-step annealing

10–4 Elaborate two-step annealing

10–5 10–6

Tc > 26 K

0

100 200 300 Temperature (K)

Figure 5.11 (a) Schematic plot of the post-annealing sequence performed in a vacuum tubular furnace to induce superconductivity to, for example, Pr2CuO4 and Nd2CuO4 thin films. The ex situ annealing process is composed of two steps: the films are annealed in a reducing atmosphere at relatively high-temperature ta (step I), followed by a reduction in vacuum, typically at a lower temperature tred (step II). (b) Temperature dependence of resistivity r(T) in Pr2CuO4 thin films. The film annealed by a conventional one-step process shows insulating behavior, while that annealed by an elaborate two-step process shows metallic conduction with superconducting transition at Tc > 26 K. From Yamamoto, H., Krockenberger, Y., & Naito, M. (2014). Proceedings of SPIE 8987, 89870V.

Epitaxial growth of superconducting oxides

109

Nd1.85Ce0.15CuO4 (Krockenbereger et al., 2013; Krockenberger, Yamamoto, Mitsuhashi, et al., 2012; Yamamoto et al., 2011, 2014), a one-step annealing process is sufficient. Needless to say, since fine tuning of parameters of the annealing process is considered to be tedious, only limited efforts on its improvement have been undertaken. When comparing the superconducting transition temperatures it is noteworthy that Tc is even higher at x ¼ 0.00 than for “optimally doped” Pr1.85Ce0.15CuO4. As our results ostensibly are in contrast to the widely reported phase diagrams of electron-doped cuprate superconductors, questions on the microscopic mechanisms during the annealing processes arise. Our investigations show that during the first step of the annealing process, regular oxygen sites of the CuO2 planes are evacuated (being consistent to Raman observations, neutron scattering results). Hence, those defective CuO2 planes are being re-occupied and “healed” during the second annealing step by an oxygen rearrangement from the apical sites to regular in-plane sites (Krockenberger et al., 2013). Unfortunately, this very complicated material issue has obscured the greater physical relevance of important aspects of orbital and electronic states in cuprates with square-planar-coordinated copper for decades while it still holds the key for a profound understanding of high-Tc superconductivity. In summary, it has been assumed that the high-Tc cuprates universally have an insulating ground state in their undoped state due to electronic correlations, and doping is necessary to change the correlations and induce superconductivity. However, our research efforts have revealed that the cuprates with the square-planar-coordinated Cu have a metallic ground state per se and superconductivity is induced after removal of defects, perturbing the playground of high-Tc superconductivity, that is, the copperoxygen planes. In addition, superconducting T0 -(La, Sm)2CuO4 (Asai, Ueda, & Naito, 2011) and T0 -(La, Eu)2CuO4 (Takamatsu, Kato, Noji, & Koike, 2012) bulk samples have already been synthesized. These observations are well aligned with our results on thin films and oppose general paradigms of theoretical models of high-Tc superconductivity. Nonetheless, predictions by first principle methods (Das & Saha-Dasgupta, 2009; Weber, Haule, & Kotliar, 2010a, 2010b) may imply that a significant overhaul in theoretical comprehension is advised. Although at the moment this conclusion is still under intense debate, the very result of our research will be a turning point in the quest for high-Tc superconductivity once further investigations support it.

5.4.3

Cuprates with an infinite-layer structure

Infinite-layer (IL) cuprates are the second-best-known compounds of cuprates with square-planar-coordinated copper. Also for IL-cuprates, the degree of perfection of the oxygen sub-lattice is essential to electronic responses of the samples. Since the crystal structure of IL-cuprates only consists of the essential building blocks required for superconductivity, they are of significant importance for the understanding of highTc cuprates. However, single crystals of infinite-layer cuprates remain elusive as this phase can be formed only by high-pressure synthesis except at a singularity composition of Ca0.86Sr0.14CuO2. In contrast, several thin-film growth techniques allow the formation of single crystalline IL-phase itself with proper combinations of substrate

110

Epitaxial Growth of Complex Metal Oxides

and growth temperature (Adachi, Sato, Ichikawa, Setsune, & Wasa, 1992; Di Castro et al., 2012; Jovanovic et al., 2011; Karimoto & Naito, 2004a, 2004b; Karimoto, Ueda, Naito, & Imai, 2001; Krockenberger, Sakuma, et al., 2012; Leca et al., 2006; Maritato et al., 2013; Naito, Karimoto, Karimoto & Tsukada, 2002; Nie et al., 2003; Terada et al., 1991; Tomaschko et al., 2012; Yamamoto et al., 2013). However, similar to the T0 -cuprates, defects within the oxygen sub-lattice play a predominant role in metallic- and superconducting properties, for example, residual resistivity ratio (RRR), Tc, transition width (DTc), etc. For example, PLD-grown films by Tomaschko et al., 2012 show sharp diffraction peaks with clear Laue fringes in the XRD patterns. With a small structure factor and a negligible scattering cross-section of oxygen ions, information about oxygen disorder in crystals remains murky. The as-grown films were insulating and superconductivity was induced by post-growth reduction procedure. Their rather small RRR of 2, low Tc of 20 K, and high resistivity values of r(300 K) ¼ 160 mU cm most likely arise from still a non-ideal oxygen sub-lattice. Several routes have been proposed and examined to achieve an (nearly) ideal oxygen sub-lattice. Maritato et al. (2013) carried out the reduction process at a higher temperature than the growth temperature while Leca et al. (2006) grew the films in a less oxidizing atmosphere and oxidized them after the growth. Nevertheless, obtaining IL-cuprate films having comparable Tc to bulk powder samples (>40 K) synthesized by high-pressure synthesis methods is still not a straightforward task. Nearly bulk-like Tc has been achieved for samples grown by MBE (Karimoto, Ueda, et al., 2001; Krockenberger, Sakuma, et al., 2012; Karimoto & Naito, 2004a, 2004b; Naito, Karimoto, et al., 2002; Yamamoto et al., 2013). In the case of MBE synthesis of IL-cuprates, several competing crystal structures, for example, Sr2CuO3, Sr14Cu24O41, and orthorhombic SrCuO2, are impeding the synthesis of single-phase samples. The growth of single-phase infinite-layer thin films is of primary importance and can only be achieved by extremely fine tuning of the cation stoichiometry and growth temperature (Ts). Actually, our research group has prepared coherently grown thin films of IL-Sr1
xLaxCuO2 by MBE on (110) DyScO3 substrates, as confirmed by high-resolution reciprocal space maps as well as high-resolution transmission electron microscopy. With respect to the oxygen sub-lattice issue, it turned out that the oxidizing conditions during the growth of infinite-layer cuprates have a tremendous influence. Figure 5.12(a) shows resistivity as a function of temperature r(T) of Sr0.9La0.1CuO2 thin films synthesized by using an atomic oxygen source with RF power varied. In terms of crystalline quality examined by XRD (i.e., degree of cation arrangement perfection), those three samples are comparable, though their electronic transport properties are entirely different. Sr0.9La0.1CuO2 thin films grown with an RF power of 400 W show insulating behavior, partly due to the formation of IL-Sr0.9La0.1CuO2þd with a substantial amount of excess oxygen (d z 0.1). This phase is often referred to as the “long-c” phase of IL-Sr0.9La0.1CuO2þd due to an expanded c-axis lattice constant (Leca et al., 2006; Tomaschko et al., 2012) when compared to the superconducting IL phase. We note that the usage of ozone as an oxidizing agent instead of atomic oxygen favors the formation of IL long-c phase, while Maritato et al. (2013) has succeeded in growing superconducting IL films by

Epitaxial growth of superconducting oxides

(a) 300

111

(b)

400 W

ρ (μΩ cm)

200 IL-Sr0.9La0.1CuO2 Thin films

100

7Å

c=

300 W

18 3.4

160 W 5Å

c

0

0

100

T (K)

411 = 3.

200

at 1.5 sccm O*

+

O2 104

O2+

RF power +

c = 3.6449 Å RF power (Long-c phase)

Optical intensity (arb. units)

105

O2+

600 W 550 W 500 W 450 W 400 W 350 W 300 W 250 W 200 W 180 W 160 W

O2

103

580

590

600 610 λ (nm)

620

630

300

Figure 5.12 (a) Influence of RF power of the radical oxygen (O*) source on the resistivitytemperature characteristics of the IL-Sr0.9La0.1CuO2 thin films grown on (110) DyScO3 substrates. (b) Emission spectra for the oxygen plasma in the radical oxygen source for various RF powers. For the RF power 300 W, no peak is discernible at l w 590 nm, while it becomes more prominent with increasing RF power (350 W). From Yamamoto, H., Krockenberger, Y., & Naito, M. (2014). Proceedings of SPIE 8987, 89870V.

using ozone. A reduction of RF power from 400 to 300 W gives rise to superconductivity, and Tc increases while the values of r decrease with decreasing the power to 160 W. This behavior is reflected in the c-axis lattice parameter as it decreases as well, implying that a defect-free oxygen configuration is stabilized at lower RF powers. In Figure 5.12(b), emission spectra of oxygen plasma generated by our RF oxygen source are plotted for a wavelength range of 580 < l < 630 nm. All of the spectra show a sharp peak at w616 nm, originating from O* (Striganov & Sventitskii, 1968) and its intensity marginally scales with the applied RF power. In contrast, the RF power dependence of the emission peak at w590 nm is more prominent and it is associated with the excitation of Oþ 2 (Pearse & Gayden, 1976). The presence of Oþ 2 species at high RF powers suggests that its existence stabilizes an oxygen sub-lattice with excess oxygen at apical sites of Cu causing an expansion of the c-axis length. Since the infinite-layer phase is, per se, not a thermodynamically stable phase, little can be done to rearrange the oxygen sub-lattice by post-growth annealing. In our experience, superconducting, infinite-layer cuprate films with Tc > 35 K are only obtained by adoption of appropriate RF powers, irrespective of the post-growth annealing conditions. Further refining of growth parameters led us to IL-cuprate films with RRRs exceeding 4 and Tc of 41 K (Figure 5.13). In addition to high crystalline quality, a defect-free oxygen configuration is necessary in order to track the inherent electronic properties of IL-cuprate superconductors. Keeping this tender oxygen configuration issue in mind, reexamination of the x

112

Epitaxial Growth of Complex Metal Oxides

120 IL-(Sr, La)CuO2 / (110)DyScO3 100

ρ (300 K) = 80 μΩ cm

ρ (μΩ cm)

80 60 40 20 0

Tc = 41 K 0

100

200

300

400

T (K)

Figure 5.13 Resistivity as a function of temperature r(T) of an IL-Sr0.9La0.1CuO2 thin film grown by molecular-beam epitaxy (MBE). Metallic conduction (RRR > 4) and superconductivity (Tc ¼ 41 K) are achieved.

dependence in IL-Sr1
xLaxCuO2 (Karimoto & Naito, 2004a, 2004b) for a comprehensive electronic phase diagram will be of great importance for future research topics.

5.4.4

Cuprates with an REBa2Cu3O7
d structure

Thin films of cuprates having a general formula of REBa2Cu3O7
d (RE ¼ rare-earth elements), often coined as RE-123, are the most extensively studied from the viewpoint of coated conductor (Goyal, 2005; Shiohara, Taneda, & Yoshizumi, 2012), microwave (Cole et al., 1992; Kurian & Naito, 2005; Lorenz et al., 1997; Michikami, Yososawa, Wakana, & Kashiwaba, 1997; Moeckley Peng, & Fischer, 2003; Naito, Karimoto, Yamamoto, Nakada, & Suzuki, 2001; Utz et al., 1997), and Josephson junction (Moeckley & Char, 1997; Usagawa, Utagawa, Koyama, Tanabe, & Shiohara, 2002) applications. Some critical experiments for basic research, for example, pairing symmetry investigation (Kirtley et al., 1995; Tsuei & Kirtley, 2000), have been also carried out by using Y-123 thin films. Due to space limitations, here, we focus on the MBE growth of RE-123 thin films having the highest Tc, as Tc is an intrinsic parameter that directly reflects the quality of samples, unlike such externally controllable parameters as Jc. Similarly to the MBE growth of T, T0 , and IL cuprates, a strict control of composition is vital for induction of the highest Tc values also for RE-123. This requirement is more difficult to fulfill in the growth of RE-123 thin films. In the T, T0 , and IL cuprates, what is of primary importance is to adjust the ratio of Cu to the summation of the remaining cations (Naito & Sato, 1995): for example, Cu/(La þ Sr) ratio in T-La2
xSrxCuO4, Cu/(Nd þ Ce) ratio in T0 -Nd2
xCexCuO4, and Cu/(Sr þ La) ratio in IL-Sr1
xLaxCuO2. On the other hand, in RE-123, the ratio of the three cations RE:Ba:Cu has to be adjusted to 1:2:3 (stoichiometric ratio). Generally, a precise control

Epitaxial growth of superconducting oxides

113

of the individual rates of the constituent cations by EIES, along with additional finetuning procedures, may result in stoichiometry. However, in practice, it is challenging (at least, very time-consuming) to adjust the RE:Ba:Cu ratio strictly to the stoichiometric ones. In a practical sense, information on allowance for deviation from the stoichiometry is very helpful for constantly yielding the high-quality RE-123 thin films. Kurian and Naito (2004, 2005) and Kurian, Naito, Sato, and Cho (2004), systematically investigated the influence of Cu/RE and Ba/RE ratios on Tc and resistivity values of RE-123 films. Figure 5.14 shows results for Nd-123. Regarding the Cu/ Nd ratio, even a slight deviation from 3 to the Cu-poor side causes substantial deterioration of the film quality: Tc decreases and the resistivity values increase rapidly. Also, XRD peak intensities are low for Cu-poor films. When the Cu/Nd ratios are between 3 and 3.6, Tc values of the Nd-123 films stay at the highest levels, around 94 K. In that Cu/Nd range, resistivity values at room temperature are also low and remain constant, insensitive to Cu/Nd ratios. When the Cu/Nd ratio exceeds 3.6, Tc starts to decrease with an increase of room temperature resistivity values. The existence of the plateau regime may somewhat ease the burden of strict composition adjustment. However, precipitates originating from the off-stoichiometric composition may play a critical role for the fabrication of heterostructures, including junctions. Subsequently, the influence of the Ba/Nd ratio was also investigated (Figure 5.14(b)). Resistivity values monotonically decrease for increasing Ba/Nd ratios, while Tc remains constant in the Ba-poor and stoichiometric regime and steeply decreases on the Ba-rich side. This is in contrast to the Cu/Nd ratio dependence, where Tc and resistivity show a simple correlation. Replacing Nd by other RE elements, the resulting dependencies remain comparable. Influences by the other two growth parameters, substrate temperature (Ts) and oxidizing conditions, are RE-dependent. For constant oxidizing conditions, the optimum growth temperatures and thermodynamic stability lines vary significantly with RE. Generally, as the ionic radius of RE3þ becomes larger, the higher Ts becomes more preferable. Comparison between Y-123 and Dy-123 revealed that the latter can be grown at a higher temperature while the ionic radii of Y3þ and Dy3þ are almost

(a)

(b)

85

400

80

200 2.8

3.2 3.6 Cu:Nd ratio

0 4.0

90

230

Tczero

190

80 1.7

1.9 2.1 Ba:Nd ratio

2.3

(μΩ cm)

600

RT

Tczero

270

Tconset

ρ

90

Tc (K)

800

RT

Tconset

(μΩ cm)

100

95

ρ

Tc (K)

100

150

Figure 5.14 Variation of Tconset , Tczero , and room temperature resistivity of Nd-123 films with (a) Cu:Nd ratio and (b) Ba:Nd ratio. The lines provided are a guide for the eye. From Kurian, J., & Naito, M. (2004). Japanese Journal of Applied Physics, 43, L1502. Copyright ©2012 by The Japan Society of Applied Physics.

114

Epitaxial Growth of Complex Metal Oxides

Resistivity (μΩ cm)

300 100

La Eu Gd Dy Nd

50 200

0

85 90

La

Eu Gd

95 100

Dy Nd

100

0 0

100

200

300

Temperature (K)

Figure 5.15 Plots of resistivity versus temperature of RE-123 thin films grown on MgO substrates for various REs. The inset shows the enlarged portion of the superconducting transition. From Kurian, J., & Naito, M. (2005). IEEE Transactions on Applied Superconductivity, 15, 2966.

same. This indicates that the growth of Y-123 requires stronger oxidization conditions than Dy-123 does. After combined optimization of those growth parameters, RE-123 films with high Tc and low resistivity values are reproducibly synthesized (Figure 5.15). Note that the MBE-grown films were subjected to oxygen annealing in an external furnace at a temperature between 300 and 350  C depending on RE for fully loading oxygen. Among the different RE-123 films (RE ¼ La, Nd, Gd, Eu, Dy) grown on MgO substrates, Nd-123 films yield the highest Tc of 94 K. For Nd-123 films, a room temperature resistivity values as low as 140 mU cm and r(300 K)/r(100 K) w 3 have been synthesized. These values are comparable to those of non-de-twinned Nd-123 single crystals. Finally, it is noteworthy that preparation of La-123 thin films with Tc values higher than 90 K is rather straightforward by MBE, although such samples are not subject to bulk synthesis methods. Also, it is inferred that Tc of La-123 thin films can be even higher than 92 K, the value shown in Figure 5.15, by increasing the growth temperature, as the Tc-versus-Ts curve does not appear to saturate at Ts of 710  C (Kurian & Naito, 2005).

5.4.5

Non-cuprate superconductors

The Ba1
xKxBiO3 (BKBO) system is well known for its high Tc (30 K) among the oxide superconductors. A prominent difference between BKBO and cuprates is that BKBO compounds are isotropic without half-filled d-orbitals, quite in contrast to cuprates, which are two-dimensional compounds and Cu is in a 3 d9 state. Highquality BKBO single crystals remain elusive due to high reactivity and volatility of K2O: bulk samples have to be prepared in a vacuum or dry environment as well as

Epitaxial growth of superconducting oxides

115

at a low temperature (Hinks et al., 1988; Nagata, Mishiro, Uchida, Ohtsuka, & Samata, 1999; Nishio & Uwe, 2003), or by high-pressure and high-temperature technique (Kim et al., 2002, 2003; Khasanova, Yoshida, Yamamoto, & Tajima, 2001). Such constraints can be overcome by thin-film synthesis methods, for example, sputtering (Hu, Lee, Chan, & Pettiette-Hall, 1993; Sato et al., 1993), PLD (Lacoe, Wendt, & Adams, 1993; Moon, Platt, Schweinfurth, & Van Harlingen, 1991), and MBE (Hellman, Hartford, & Gyorgy, 1991; Utz, Wiest, Prusseit, Berberich, & Kinder, 1995; Yamamoto et al., 2004). In MBE, we examined the synthesis of BKBO at growth temperatures of 300e400  C. Typical Tc and resistivity reported for films are Tc ¼ 25e29 K, r(300 K) w 220e250 mU cm, and r(150 K) w 150 mU cm. Though handling of K is cumbersome by any method, a lot can be done for the improvement of the crystalline quality of BKBO films. Another topic that has been attracting increasing attention is transparent superconducting thin films of complex oxides. Hitosugi and colleagues (Hitosugi et al., 2014; Kumatani et al., 2012) have grown LiTi2O4 films on (111) MgAl2O4 substrates by PLD. Although the target stoichiometry was Li4Ti5O12, it turned out that the Li/Ti stoichiometry arriving at the substrate during deposition was w0.5. Such stoichiometry allows formation of high-quality LiTi2O4 films under certain oxygen partial-pressure values. Superconducting polycrystalline LiTi2O4 thin films (Tc w 13 K) (Inukai, Murakami, & Inamura, 1982; Inukai & Murakami, 1985) as well as superconducting epitaxial LiTi2O4 thin films (Tc w 10 K) (Chopdekar, Wong, Takamura, Arenholz, & Suzuki, 2009) themselves had been reported. The synthesis method introduced by Hitosugi et al. allows for high-quality LiTi2O4 thin films with Tc as high as 13 K (highest ever reported for epitaxial thin films of LiTi2O4) and r(300 K) as low as 330 mU cm (one order of magnitude lower than previously reported). Together with a negligible absorption coefficient in the visible range, LiTi2O4 thin films seem to be promising as transparent electrode materials. In general, low carrier densities, for example, w1020e1021 cm
3, give rise to optical transparency, and conductivity arises from high-charge carrier mobilities. However, for LiTi2O4, a rather high carrier density of w1  1022 cm
3 (calculated from the Hall coefficient) was reported (Inukai & Murakami, 1983, 1985), which is consistent with its fairly high Tc of w13 K, as high density of states at the Fermi level, N(0), enhances Tc. Hitosugi et al. speculated that the plasma frequency of this material is shifted into an infrared regime by a large dielectric constant or a large effective mass, and hence it is transparent. As a final example of non-cuprate-oxide superconductors we discuss the growth of superconducting Sr2RuO4 thin films. While the Tc of Sr2RuO4 is low (w1 K), this superconductor has received attention mainly due to a unique order parameter (likely p-wave symmetry). For further investigation of the paring symmetry and utilization of the unique character of the order parameter, order parameter phase-sensitive experiments, that is, Josephson junctions, are required, and therefore epitaxial thin films. Since superconducting single crystals were reported by Maeno et al. (1994), considerable efforts have been made to create superconducting Sr2RuO4 thin films (Madhaven, Schlom, Dabkowski, Dabkowska, & Liu, 1996; Tian et al., 2007). Although epitaxial Sr2RuO4 thin films had been prepared, superconductivity had not been achieved before the success by Krockenberger et al. (2010) by means of the

116

Epitaxial Growth of Complex Metal Oxides

PLD technique. They reported the growth window for high-quality films is quite narrow (e.g., allowance of growth temperature is <10  C at around 920  C, and the use of a 19.6-mTorr-Ar and 0.4-mTorr-O2 mixture is mandatory). From single-crystal work, it is known that both impurities and structural disorders can easily quench superconductivity in this material, and a low defect density is key to inducing superconductivity. MBE might be a tool of choice for the synthesis of ultra-high-quality, superconducting Sr2RuO4 thin films, as such high-quality SrRuO3 has already been grown by MBE (Koster et al., 2012; Siemons et al., 2007).

5.5

Synthesis of new superconductors by thin-film growth methods

The search for new superconductors with higher Tc is the most challenging subject in materials sciences. Here, we demonstrate that multisource MBE with a high-precision rate control is a promising approach not only for the preparation of high-quality films of known materials but also for the synthesis of brand-new complex metal oxides beyond approaches of alternate stacking of existing lattices.

5.5.1

Advantages of MBE as a synthesis method of new materials

Besides the superior crystalline quality of resultant samples, the list of merits of the MBE synthesis can be extended to (1) low temperature reaction by ultimately small reactants (atoms, ions, etc.), (2) quasi-stable phase formation by epitaxy (pseudomorphic stabilization effect), (3) contamination-free environment under UHV utilizing pure metal sources, and (4) high throughput screening of synthesis conditions. In addition, a key factor (5) large surface-to-volume ratio, can be readily exploited in film materials, as oxygen engineering plays a fundamental role in certain oxides. It is noteworthy that the idea of artificially stacking the building blocks of cuprates has been unsuccessful since each building block is not charge-neutral. Redox reaction at the interface and/or change of crystal structure (altered phase formation) take place and render the approach of artificial stacking inconclusive.

5.5.2

Case studies

We have already indicated that the annealing process necessary for T0 and IL cuprates is for the stabilization of a defect-free oxygen sub-lattice. Thermodynamically, the annealing process can be reversed for the introduction of excess oxygen. For T-La2CuO4 it is well known that superconductivity is induced by introducing excessive oxygen into the lattice. This process, however, is very time-consuming and requires high-pressure or electrochemical methods in bulk samples. In contrast, in our MBE films, oxygenation is easily achieved by ozone exposure (Sato, Naito, & Yamamoto, 1997) and the resultant T-La2CuO4þd films show Tc as high as 56 K (Naito, Tsukada, Greibe, & Sato, 2002). We further utilized the strong oxidizing power

Epitaxial growth of superconducting oxides

117

for the synthesis of Ba2CuO4d and Sr2CuO4d (Karimoto, Yamamoto, Greibe, & Naito, 2001, Karimoto, Yamamoto, Sato, Tsukada, & Naito, 2003; Yamamoto, Naito, & Sato, 1997, 2000) and discovered that these materials are also superconducting (Figure 5.16). In Ba2CuO4d and Sr2CuO4d, structural phase transitions from orthorhombic Ba2CuO3þd (c ¼ 13.1 Å)/Sr2CuO3þd (c ¼ 12.71 Å) to tetragonal Ba2CuO4d (c ¼ 14.6 Å)/Sr2CuO4d (c ¼ 13.55 Å) take place during ozone exposure. Since Ba2CuO3þd and Sr2CuO3þd are very unstable in air—especially, Ba2CuO3þd is formed only under a CO2-free and moisture-free environment—keys for the successful synthesis of superconducting Ba2CuO4d and Sr2CuO4d are contamination-free environment achieved by MBE and strong oxidation by utilizing the large surface-to-volume ratio. It is worth noting that the c-axis lattice constant of superconducting, tetragonal Sr2CuO4d (n = 1) prepared by this method well coincides with the extrapolation of the lattice constants of Sr2Can
1CunOy for n ¼ 2e4. However, the unit cell of tetragonal Sr2CuO3þd (c ¼ 12.43 Å) cannot be extrapolated from the Sr2Can
1CunOy series, as pointed out in a previous report (Kawashima & Takayama-Muromachi, 1996). Next, the synthesis of a new superconductor PbSr2CuO5þd (Pb-1201) by MBE is described (Karimoto & Naito, 1999, 2000; Naito, Karimoto, & Yamamoto, 2000). When atomic beam fluxes of the constituent cations, that is, Pb, Sr, and Cu, are supplied in an MBE setup with the appropriate stoichiometric ratio while simultaneously supplying O3 or atomic oxygen, SrPbO3, having the perovskite structure, is formed, which is thermodynamically the most stable phase. We demonstrated, however, that the resultant phase can be altered depending on substrate materials due to pseudomorphic stabilization effects. Figure 5.17(a) shows XRD patterns for 750-Å thick PbeSreCueO films on (a) (001) SrTiO3 (a ¼ 3.905 Å) and (b) (001) LaAlO3 (a ¼ 3.821 Å) substrates. The SrPbO3 (a ¼ 4.15 Å) phase dominates in the film grown on SrTiO3, whereas the film grown on LaAlO3 is single-phase c-axis 2.0

Resistivity (mΩ cm)

Ba2CuO4±δ

1.0

Sr2CuO4±δ

La2CuO4+δ 0

0

50

100 150 200 Temperature (K)

250

300

Figure 5.16 Plots of resistivity versus temperature of Ba2CuO4d, Sr2CuO4d, and La2CuO4þd thin films. Data from Karimoto et al. (2003).

118

Epitaxial Growth of Complex Metal Oxides

Intensity (arb. units)

Substrate

(c)

on SrTiO3

(110) (220)

(020)

SrPbO3 Substrate (002)

Pb-1201

Substrate (004)

on LaAlO3

(003)

(001)

0

10

(006)

20

(b) 0.5

30 40 50 2θ (degree)

60

70

0.4 0.3 0.2 0.1 0

0

50

100 150 200 250 300 Temperature (K)

Pb-1201 (002)

(003)

Substrate YAlO3

(3.72 Å)

LaSrAlO4 (3.76 Å) LaAlO3

(3.82 Å)

LaSrGaO4 (3.84 Å) SrPbO3

*

20

15

NdGaO3 (3.84 Å)

*

SrTiO3

25 30 2θ (degree)

(d) Intensity (arb. units)

Resistivity (mΩ cm)

Substrate

Intensity (arb. units)

(a)

35

Pb-1201 Sr2CuO3 SrCuO2

10

20

(3.91 Å)

Ts = 650 °C 625 °C 600 °C 575 °C 550 °C 525 °C 500 °C

30

40

50

60

2θ (degree)

Figure 5.17 (a) XRD patterns of PbeSreCueO films grown on SrTiO3 and LaAlO3 substrates. (b) Plot of resistivity versus temperature of Pb-1201 thin films grown on a LaAlO3 substrate. (c) XRD patterns indicating influence of substrate on the resultant phase formation in the PbeSreCueO films. (d) X-ray diffraction patterns indicating influence of growth temperature (Ts) on the resultant phase formation in the PbeSreCueO films. From Karimoto and Naito (1999, 2000) and Naito, Karimoto, et al. (2000).

oriented Pb-1201 (a ¼ 3.81 Å). Pb-1201 films grown on LaAlO3 show superconductivity at Tc w 40 K (Figure 5.17(b)). Further investigations on how substrate materials influence the growth revealed that the Pb-1201 phase is stabilized on substrates whose in-plane lattice constants are smaller than 3.85 Å (Figure 5.17(c)). In other words, the substrates whose lattice constant is smaller than 3.85 Å prevent formation of the thermodynamically stable SrPbO3 phase. This behavior shows that chemical reactions or resultant compounds may be tuned/altered by utilizing epitaxy well beyond a simple epitaxial strain effect. Hence, another important factor in the synthesis of Pb-1201 is the growth temperature. Figure 5.17(d) shows the growth temperature dependence of the XRD patterns for the films grown on LaAlO3 substrates. The Pb-1201 phase is exclusively stabilized in the low temperature range below 550  C. Accordingly, an appropriate choice of substrate and low-temperature synthesis, both of which can be achieved by MBE, are the key for the successful formation of Pb-1201. Last but not least, recent discoveries of IL-CaCuO2/SrTiO3 superconducting artificial superlattices, based on two insulating oxides, are quite intriguing (Di Castro et al., 2012, 2014). Although electric-field-induced superconductivity in insulating oxides (Ueno et al., 2008, 2011) as well as superconductivity emerged at the interface

Epitaxial growth of superconducting oxides

119

between two insulators, that is, the LaAlO3eSrTiO3 interface (Richter et al., 2013) was reported, their Tc values are rather low (<1 K). In contrast, Di Castro et al. reported Tc of 50 K for the CaCuO2/SrTiO3 superlattices. While the origin of the 50 K superconducting phase has not yet been clarified, it might have some relevance to previously reported results by a different research group: IL-Sr1
xCaxCuO2
d thin films prepared by sputtering showed an onset Tc of w50 K (Nie et al., 2003). Since cuprates having the IL structure are composed of CuO2 planes with square-planar coordinated copper, those observations are also interesting in relation to superconductivity in parent compounds of T0 -cuprates.

5.6

Conclusions and future trends

The preparation of high-quality epitaxial films of superconducting oxides has been described, mainly focusing on MBE growth of cuprate thin films. We have demonstrated that such research-grade thin-film specimens can further promote fundamental physical investigations, including the quest for the superconducting pairing mechanism itself. In addition to growth of research-grade thin films, synthesis of new superconducting oxides by advanced thin-film growth methods has been reviewed. Since outstanding developments in measurement techniques and equipment are enabling cutting-edge characterization methods to be used for thin-film specimens, the trend of physical property investigation using thin-film specimens will be further accelerated. Research on novel properties that high-quality ultrathin films and/or interfaces exhibit will also be actively pursued, at least partially intended to search for new superconductors with higher Tc. Finally, thin-film research on complex metal oxides containing 4d and/or 5d transition metal elements have been neglected but will play a trailblazing role in the search for novel superconducting materials.

5.7

Sources of further information and advice

With respect to cuprates, still holding the record Tc, there has been an enormous amount of experimental results accumulated. Although we cannot hope to cover all the sources of further information, here we list a few articles for so-called electron-doped cuprates. Unlike hole-doped cuprates, for electron-doped cuprates, reported experimental results themselves have been sometimes controversial, especially depending on sample dimensions, that is, thin films or bulk crystals. Understanding the origin of such an enigma requires vast knowledge on materials-researchrelated issues beyond thin-film growth, and the following comprehensive review articles are helpful. Armitage, N. P., Fournier, P., & Greene, R. L. (2010). Progress and perspectives on electrondoped cuprates. Reviews on Modern Physics, 82, 2421. Naito, M. (2011). Electron-doped cuprates as high-temperature superconductors. In X. G. Qiu (Ed.), High-temperature superconductors (pp. 208e274). Cambridge: Woodhead Publishing Limited.

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Acknowledgments The authors are thankful for fruitful discussions with Dr A. Tsukada, Dr H. Sato, Dr J. Kurian, Dr S. Shibata, and Dr S. Karimoto. They also thank Dr C. Lu and Dr R. H. Hammond for great lessons on rate-controlling systems. HY expresses deep appreciation to Prof. N. Terada for helpful discussions on the assignment of oxygen species formed in the RF plasma source. The authors also acknowledge the successive directors of NTT Basic Research Laboratories and the division of Materials Science of the laboratories for their continuous support.

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