- Email: [email protected]
Electron microscopy studies of epitaxial MgB2 superconducting thin films grown by in situ reactive evaporation
ARTICLE IN PRESS
Journal of Crystal Growth 280 (2005) 602–611 www.elsevier.com/locate/jcrysgro
Electron microscopy studies of epitaxial MgB2 superconducting thin films grown by in situ reactive evaporation Lin Gua, Brian H. Moecklyb, David J. Smitha, a
Center for Solid State Science, Arizona State University, Tempe, AZ 85287-1704, USA b Superconductor Technologies, Inc., Santa Barbara, CA 93111, USA Received 12 February 2005; accepted 14 March 2005 Available online 23 May 2005 Communicated by T.F. Kuech
Abstract The morphology and chemistry of epitaxial MgB2 thin films grown using reactive Mg evaporation on different substrates have been characterized by transmission electron microscopy methods. For polycrystalline alumina and sapphire substrates with different surface planes, an MgO transition layer was found at the interface region. No such layer was present for films grown on MgO and 4-H SiC substrates, and none of the MgB2 films had any detectable oxygen incorporation nor MgO inclusions. High-resolution electron microscopy revealed that the growth orientation of the MgB2 thin films was closely related to the substrate orientation and the nature of the intermediary layer. Electrical measurements showed that very low resistivities (several mO cm at 300 K) and high superconducting transition temperatures (38 to 40 K) could be achieved. The correlation of electrical properties with film microstructure is briefly discussed. r 2005 Elsevier B.V. All rights reserved. PACS: 68.37.Lp; 68.55.Jk; 74.70.Ad Keywords: A1. Characterization; A1. Crystal morphology; B2. Superconducting materials
1. Introduction The recent discovery of superconducting MgB2 [1] holds much promise for the development of Corresponding author. Tel.: +1 4809654540;
fax: +1 4809659004. E-mail address: [email protected] (D.J. Smith).
superconducting applications and devices operating at medium temperatures. The simple crystal structure, relatively high superconducting transition temperature, Tc, of about 40 K [2], long coherence length of about 50 A˚ [3], metallic charge carrier density [4], and strongly linked current flow in polycrystalline forms [5,6] are all characteristics that facilitate the use of this material for active and
0022-0248/$ - see front matter r 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.jcrysgro.2005.03.066
ARTICLE IN PRESS L. Gu et al. / Journal of Crystal Growth 280 (2005) 602–611
passive devices, detectors, and current carrying applications. Despite intensive efforts and the conspicuous progress that has recently been made, the growth of single-crystal MgB2 remains a non-trivial challenge because of its tendency to oxidize under certain conditions and the volatility of the Mg component over a large temperature range. Several different shapes of MgB2 samples have been successfully developed, including bulk [1], wire [6], and thin films [4]. For most thin film applications, it is preferable to deposit MgB2 completely in situ without involving in situ or ex situ annealing processes under high Mg vapor pressure at elevated temperature [7]. Earlier MgB2 thin films grown by pulsed-laser deposition (PLD) with ex situ annealing had Tcs ranging from 29–34 K [8] to 39 K [4]. In situ post-annealed films deposited by PLD had Tcs of 22–24 K [9]. Electron beam evaporation (e-beam) followed by postgrowth annealing provided a path towards fabrication of large-area thin films [10]. In order to generate the high Mg vapor pressure required for the thermodynamically stable MgB2 phase, hybrid physical–chemical vapor deposition (HPCVD) techniques were developed by combining physical vapor deposition (PVD) and chemical vapor deposition (CVD) to produce high-quality epitaxial thin films on selected substrates [11,12]. In situ film growth using molecular beam epitaxy (MBE) has also been reported [13–16], but the growth temperature must be low to ensure that Mg does not desorb rapidly from the film surface. Contrary to growth of high-temperature superconductors, highly epitaxial growth of MgB2 films is not always necessary in order to achieve good electrical properties [17]. Some substrates are, however, amenable to good quality epitaxial growth. MgB2 has a hexagonal structure with space group P6/mmm and lattice constants of ( c ¼ 3:524 A ( [18]. Sapphire (Al O ) a ¼ 3:086 A; 2 3 ( c ¼ 12:991 A ( with lattice constants of a ¼ 4:758 A; [18] serves as one of the major substrates for MgB2 thin film deposition, despite the relatively large lattice mismatch [4,10,11]. MgO [9,10] and Mg [19] have been used for MgB2 thin film growth because of their compatibility with elemental magnesium. 6H-SiC [9] with an a lattice spacing of 3.08 A˚, and
603
4H-SiC [11,20] with an a lattice spacing of 3.073 A˚, have a lattice mismatch with MgB2 of less than 1%, and should be ideal for high-quality epitaxial growth. Si and SrTiO3 substrates have also been used in studies of the electrical properties of MgB2 thin films [9]. The growth of high-quality MgB2 films on a number of single crystalline and polycrystalline substrates has also been demonstrated [17]. It has been shown that the use of different substrates affects the interfacial chemistry and affects the crystal growth [16,20,21]. The transmission electron microscope (TEM) provides a range of techniques for studying the local morphology and chemistry of MgB2 thin films. Previous analysis of MgB2 thin films has mainly concentrated on whether or not interfacial regions had a thin MgO or MgAl2O4 layer [20,21]. Bulk MgB2 studies have revealed that MgO as the major secondary phase is responsible for the formation of dislocations and other defects [22]. Systematic studies of changes in MgB2 film morphology, interfacial structure and local chemistry for several different substrates are presented and discussed in this work.
2. Experimental details 2.1. Sample growth We have developed a reactive evaporation technique for MgB2 thin film deposition that directly addresses the major problems associated with complete in situ growth of this material [17]. The method centers on the development of a rotating blackbody heater with an internal pocket filled with Mg vapor. This design maintains a localized source of Mg vapor near the substrate(s) while maintaining high vacuum in the rest of the deposition chamber. Boron from a single source is deposited by electron beam evaporation. The MgB2 phase then forms naturally by adsorptionlimited growth. Large-area film growth and double-sided deposition are both possible using this method, and an intermediate growth temperature range of 400–600 1C is accessible. MgB2 films have been deposited on a multitude of singlecrystal, polycrystalline, and metallic substrates up
ARTICLE IN PRESS 604
L. Gu et al. / Journal of Crystal Growth 280 (2005) 602–611
to 400 in diameter. This technique enables growth of high-quality MgB2 films on any substrate for which there is not a chemical reaction, including a large number of single-crystal substrates as well as polycrystalline alumina, stainless steel, and buffered Si. These films regularly display Tc values over 39 K, and their resistivity values are among the lowest reported in the literature. The films are also very clean as exhibited by low microwave loss [23], low Hc2 values [24], and a lack of oxygen or carbon contamination. In addition, the films are relatively stable with respect to patterning and upon exposure to water. All the films in this study were deposited at a growth rate of about 1 A˚/s and at a substrate temperature of 550 1C, a temperature which is generally inaccessible by other MgB2 thin film growth techniques. Table 1 lists some key electrical properties of these films. 2.2. TEM sample preparation Since MgB2 thin films can be highly susceptible to hydro-reaction, the procedures and chemicals normally used for TEM cross-section sample preparation had to be modified. The isocut fluid used during wafer slicing was replaced by polishing oil for cutting lubrication, and chloroform was used instead of acetone for general cleaning before the samples were glued together for cross-sectioning. Samples were ground to about 100 mm in the presence of polishing oil, followed by double-sided dimple polishing. Further thinning was carried out using a precision ion polishing system at 4 keV ion energy followed by a short period of final cleaning at 2.5 keV. In this work, bright-field diffraction-contrast imaging, selected-area electron diffraction and
high-resolution electron microscopy for general morphology and defect characterization were performed using a JEOL 4000EX (operated at 400 keV) with an interpretable resolution of 1.7 A˚. Interface chemistry and film uniformity were analyzed by electron-energy-loss spectroscopy (EELS) and energy-dispersive X-ray spectroscopy (EDXS) using a Philips CM200 operated in the focused probe mode with a spot size of 1 nm.
3. Experimental results In this study, we have performed a systematic investigation of the morphology of epitaxial MgB2 thin films grown on five different substrates. Polycrystalline material and columnar growth was generally observed throughout the epilayer. The presence of an Mg-O transition layer at the interfaces with c-plane sapphire, r-plane sapphire and polycrystalline alumina substrates had a significant impact on the growth orientation of the subsequent MgB2 layer. 3.1. c-plane sapphire MgB2 deposited on c-plane sapphire showed a growth discontinuity with the presence of an oxide layer between the film and the substrate, consistent with other reports [20,21]. Fig. 1(a) is a bright-field diffraction-contrast image showing the morphology of a 300-nm-thick MgB2 film and the intermediate layer, and the inset is the corresponding selected-area electron diffraction (SAED) pattern. It is clear from the image that the epitaxial MgB2 thin film is columnar and polycrystalline.
Table 1 Properties of MgB2 films discussed in the text, including the zero-resistance Tc, DC resistivity values at 300 and at 40 K, and the film thickness Substrate
Tc (K)
r300 K (mO cm)
r40 K (mO cm)
Thickness (nm)
c-plane sapphire r-plane sapphire Polycrystalline alumina MgO 4H-SiC
38.2 39.1 39.0 38.5 39.5
14.5 10.2 9.5 12.3 25.5
6.5 2.6 2.2 4.4 12.3
300 500 500 500 500
All films were deposited by reactive evaporation at 550 1C.
ARTICLE IN PRESS L. Gu et al. / Journal of Crystal Growth 280 (2005) 602–611
605
or so. This morphology is presumably due to the graphite-like layered nature of MgB2 and the relatively high compatibility of elemental Mg with respect to the Mg distribution around the interface [9,10,19]. The high-resolution electron micrograph in Fig. 2 shows the interface microstructure in greater detail. A cubic phase aligned in a [1 1 0]-type orientation and a structurally coherent grain boundary with slight off-axis tilt are observed. The intermediate layer has an overall darker appearance suggesting either greater defect density than the sapphire substrate and/or higher packing density with heavier elements overall than those present in the MgB2 layer above. Fig. 3 shows an annular-dark-field (ADF) electron micrograph (a) of MgB2 grown on c-plane sapphire, together with the corresponding EDXS (b) and EELS (c) line profiles. Spectra were acquired at the location indicated in (a). In the ADF imaging mode, image contrast is proportional to atomic number, and heavier material would thus appear to be brighter. The Al signal
Fig. 1. (a) Electron micrograph showing morphology of MgB2 thin film grown on c-plane sapphire with corresponding selected-area electron diffraction pattern inset; and (b) enlargement showing faceting of MgO intermediary layer and initial MgB2 growth region.
Grain sizes range in width from 50 to 100 nm, and the top surface is relatively bumpy with height variations of as much as 20 nm and a rootmean-square roughness of a few nm. As shown by the enlargement in Fig. 1(b), the initial growth of MgB2 is highly faceted with respect to the underlying transition layer for the first 100 nm
Fig. 2. High-resolution lattice image showing interface between sapphire substrate and unintentional MgO intermediary layer.
ARTICLE IN PRESS L. Gu et al. / Journal of Crystal Growth 280 (2005) 602–611
606
(a)
EDX profile EELS profile
200 nm (b)
500
Mg Al O
Counts
400 300 200 100 0 0
50 Position (nm)
100
(c) B O
Counts (a.u.)
70000
was found to diminish rapidly across the intermediate layer as the probe was moved away from the sapphire substrate with step sizes of about 4 nm. At almost the same location where the Al signal dropped to zero, the O profile also exhibited a drop before finally decreasing to zero after another 50 nm. The rapid increase of the Mg signal combined with the O signal indicated that the major component of the intermediate layer was MgO. The EELS line profile confirmed this result by excluding the presence of boron at least within the detection limit of the EELS technique. The presence of the MgB2 phase became evident after termination of the unintentionally grown MgO transition layer. The relative lower concentration of Mg in the MgB2 film compared with that measured in the MgO layer is consistent with the ADF and bright-field image contrast. Contrary to the observation of MgO as a major secondary phase in bulk MgB2 [22], it appears that the epitaxial MgB2 material here was not impacted by any deterioration by oxidation. This is likely because little or no MgO exists in our Mg vapor pocket during growth, with any gettering of residual oxygen by Mg escaping the heater and depositing on cooler surfaces in the vacuum chamber. Electrical measurements of this film indicated a Tc value of 38.2 K, and low resistivity values of 14.5 mO cm at 300 K, and 6.5 mO cm at 40 K. However, higher Tc values and even lower resistivities were observed for films deposited on other substrates. 3.2. r-plane sapphire
35000
0 0
15
30
45
59
Position (nm) Fig. 3. (a) Annular-dark-field electron micrograph of MgB2 grown on c-plane sapphire, with corresponding EDX (b), and EELS (c), line profiles showing quantitative chemical distribution.
Fig. 4(a) shows the microstructure of an epitaxial 500-nm-thick MgB2 thin film grown on r-plane sapphire, and the inset shows the corresponding SAED pattern. An intermediate MgO layer is again present at the substrate interface. The thickness of the MgO layer is larger than that observed for growth on c-plane sapphire. The surface roughness and grain sizes ranging from 80 to 120 nm are also visibly larger. There is a significant difference of the columnar growth morphology relative to the films grown on c-plane sapphire in that the growth direction is tilted away from the substrate normal, leading to an inclined
ARTICLE IN PRESS L. Gu et al. / Journal of Crystal Growth 280 (2005) 602–611
607
texture of the subsequent MgB2 film. Comparison between the MgO lattice planes visible in Fig. 4(b) and in Fig. 2 reveals that the MgO intermediate layer here has been rotated by 301 relative to the case for the c-plane sapphire in Fig. 2. These different orientations of the intermediate layer have determined the subsequent MgB2 growth orientation. Thus, the difference in MgB2 growth morphology observed for the sapphire substrates is primarily because of the different cutting plane. For MgB2 thin films grown on non-latticematched substrates, growth on r-plane sapphire exhibits excellent properties. Electrical measurements indicated a Tc value of 39.1 K for this film, which was one of the highest among the five films studied here. The resistivity values of this film were rð300 KÞ ¼ 10:2 mO cm and rð40 KÞ ¼ 2:6 mO cm. This residual resistivity value is among the lowest reported in the literature for thin films, indicating that the films are characterized by excellent connectivity and low scattering from impurities or grain boundaries [25].
3.3. Polycrystalline alumina
Fig. 4. (a) Cross-sectional electron micrograph showing morphology of MgB2 thin film grown on r-plane sapphire with corresponding electron diffraction pattern inset; and (b) highresolution lattice image of interfacial oxygen-diffusion region.
film morphology. The origin of the non-vertical microstructure can be attributed to the MgO intermediate layer which establishes the growth
Fig. 5(a) was the microstructure of an MgB2 thin film grown on polycrystalline alumina. An MgO transition layer with a moderate thickness of 50 nm was observed at the substrate surface. The surface roughness of the upper film is excellent and would be favorable for graded-structure device applications. Moire´ fringes were observed in the MgB2 film indicating overlapping grains with different lattice spacings and/or orientations. The growth direction was again tilted away from the substrate normal and the grain size was larger than that observed for the different sapphire cutting planes, ranging from 200 to 400 nm. This larger grain size and consequent reduced intergrain scattering is likely responsible for our lowest observed residual resistivity of rð40 KÞ ¼ 2:2 mO cm. A Tc value of 39.0 K was measured for these epitaxial films grown on alumina, indicating high quality material. Fig. 5(b) shows a high-resolution lattice-image of the interface and the MgO transition region. It is clear that the growth direction of MgO on the alumina substrate is different from the two previous sapphire
ARTICLE IN PRESS 608
L. Gu et al. / Journal of Crystal Growth 280 (2005) 602–611
layer, where the oxygen appears to have been supplied by the O-containing substrate. For growth on an MgO substrate, an MgO intermediate layer would not be expected. Indeed, no apparent transition layer is visible in Fig. 6(a), though the region close to the substrate is highly defective. SAED patterns reveal excellent alignment of the diffraction spots for the MgB2 thin film and the MgO substrate. No extra ring-like polycrystalline patterns were visible except for slight elongation of the MgB2 spots. Threading defects were observed throughout the film, and no major Moire´ fringes were detected. These are promising characteristics for epitaxial growth of high-quality superconducting films. Fig. 6(b) is a high-resolution lattice image from the region of the interface which shows no sign of any secondary phase. The growth direction is aligned exactly parallel to the substrate normal shown in Fig. 6(c) and is aligned according to the MgO [1 0 0] orientation. Electrical characterization of this film grown on MgO showed a small decrease of Tc to 38.5 K and slight increase of rð40 KÞ to 4.4 mO cm [rð300 KÞ ¼ 12:3 mO cm]. 3.5. 4-H SiC
Fig. 5. (a) Electron micrograph showing morphology of MgB2 thin film grown on a polycrystalline alumina substrate with corresponding electron diffraction pattern inset; and (b) highresolution lattice image of the interfacial oxygen-diffusion region.
substrates, and establishes the texture for growth of the subsequent MgB2 layer. 3.4. MgO A common observation for the three previous substrates was the presence of an MgO transition
4-H SiC has a very close lattice match with MgB2. From the diffraction-contrast micrograph shown in Fig. 7(a), no obvious MgO intermediate layer was present. However, the MgB2 film exhibits a granular morphology with fine grains of MgB2 separating the substrate and the upper columnar grains. Almost one-third of the total film thickness is occupied by the granular transition region of about 100 nm. The width of the columnar grains above this region is the smallest of all five films studied. However, the film surface is comparatively flat and should be suitable for growth of graded structures. Fig. 7(b) shows a high-resolution lattice-image of the MgB2 faceted microstructure close to the interface region, which was not observed for the growth on the MgO [1 0 0] substrate. A possible explanation is the different substrate structure and orientation, similar to inclined growth of MgB2 on the intermediate layer, for which the rotated grains are mainly due to the different MgO transition layer orientation.
ARTICLE IN PRESS L. Gu et al. / Journal of Crystal Growth 280 (2005) 602–611
609
Electrical measurements revealed that the Tc of this film is 39.5 K, the highest value of the films studied here. However, the resistivity values are the worst, with rð300 KÞ ¼ 25:5 mO cm and rð40 KÞ ¼ 12:3 mO cm. The high rð40 KÞ value is likely to be due to increased intergrain scattering resulting from the fine grain morphology. The higher Dr value [rð300 KÞ rð40 KÞ ¼ 13:2 mO cm] of this film in comparison to the vast majority of our films [Dr ¼ 7 to 9 mO cm] may be due to decreased connectivity throughout the fine-grained region [25].
4. Discussion
Fig. 6. (a) Electron micrograph showing epitaxial MgB2 thin film grown on an MgO substrate with corresponding electron diffraction pattern inset. High-resolution lattice images of: (b) interface without distinctive extra layers; (c) top MgB2 layer with [0 0 0 1] growth direction.
High-quality MgB2 thin films have been successfully deposited on five different substrates using the reactive Mg evaporation technique. The morphology and chemistry of the highly epitaxial films were characterized using transmission electron microscopy. An MgO transition layer was detected immediately above the substrate for films grown on c-plane and r-plane sapphire and polycrystalline alumina, but no such layer was observed for MgO and 4-H SiC substrates. The layer does not appear to form as a result of reaction with residual oxygen in the deposition chamber, since this layer is not observed for growth on the non-oxide substrates. Moreover, MgO inclusions are not observed in the bulk of the MgB2 films. However, during our normal growth process, the heated substrates are exposed to Mg vapor prior to B evaporation. We would therefore surmise that the MgO layer growth during this stage of the process via out-diffusion of oxygen from the substrate. We have since tested this idea by first evaporating a small amount of B onto the substrate surface prior to exposure to Mg vapor. In this case, analysis by Rutherford backscattering detected no MgO layer at the interface of MgB2 films grown on sapphire [26]. Furthermore, we have not yet observed any negative consequences of this MgO layer, since the demanding microwave properties of our films are excellent [23]. The epitaxial films had columnar morphology in all cases. The growth direction of the films was strongly influenced by the MgO intermediate layer orientation, and the width of the grains varied
ARTICLE IN PRESS 610
L. Gu et al. / Journal of Crystal Growth 280 (2005) 602–611
conducting gaps [27]. Furthermore, electrical measurements on these films reveal very low resistivities and high superconducting transition temperatures, indicating a compact and stronglylinked film morphology. These results are consistent with the TEM observations.
Acknowledgments We acknowledge use of facilities in the John M. Cowley Center for High Resolution Electron Microscopy at Arizona State University.
References
Fig. 7. (a) Electron micrograph showing the morphology of MgB2 thin film grown on 4H-SiC substrate; and (b) highresolution lattice-image adjacent to the interface showing faceted growth morphology.
depending on the different substrate and the transition layer thickness and orientation. Pointcontact tunneling spectra taken across the surface of MgB2 films grown on r-plane sapphire have indicated the clear presence of both MgB2 super-
[1] J. Nagamatsu, N. Nakagawa, T. Muranaka, Y. Zenitani, J. Akimitsu, Nature (London) 410 (2001) 63. [2] S.L. Bud’ko, G. Lapertot, C. Petrovic, C.E. Cunningham, N. Anderson, P.C. Canfield, Phys. Rev. Lett. 86 (2001) 1877. [3] D.K. Finnemore, J.E. Osterson, S.L. Bud’ko, G. Lapertot, P.C. Canfield, Phys. Rev. Lett. 86 (2001) 2420. [4] W.N. Kang, H.J. Kim, E.M. Choi, C.U. Jung, S.L. Lee, Science 292 (2001) 1521. [5] D.C. Larbalestier, L.D. Cooley, M.D. Rikel, A.A. Polyanskii, J. Jiang, S. Patnaik, X.Y. Cai, D.M. Feldmann, A. Gurevich, A.A. Squitieri, M.T. Naus, C.B. Eom, E.E. Hellstrom, R.J. Cava, K.A. Regan, N. Rogado, M.A. Hayward, T. He, J.S. Slusky, R. Khalifah, K. Inumaru, M. Haas, Nature 410 (2001) 186. [6] P.C. Canfield, D.K. Finnemore, S.L. Bud’ko, J.E. Ostenson, G. Lapertot, C.E. Cunningham, C. Petrovic, Phys. Rev. Lett. 86 (2001) 2423. [7] Z.K. Liu, D.G. Schlom, Q. Li, X.X. Xi, Appl. Phys. Lett. 78 (2001) 3678. [8] A. Berenov, Z. Lockman, X. Qi, J.L. MacManus-Driscoll, Y. Bugoslavsky, L.F. Cohen, M.H. Jo, N.A. Stelmashenko, V.N. Tsaneva, M. Kambara, N. Hari Babu, D.A. Cardwell, M.G. Blamire, Appl. Phys. Lett. 79 (2001) 4001. [9] D.H.A. Blank, H. Hilgenkamp, A. Brinkman, D. Mijatovic, G. Rijnders, H. Rogalla, Appl. Phys. Lett. 79 (2001) 394. [10] S.H. Moon, J.H. Yun, H.N. Lee, J.I. Kye, H.G. Kim, W. Chung, B. Oh, Appl. Phys. Lett. 79 (2001) 2429. [11] X.H. Zeng, A.V. Pogrebnyakov, A. Kotcharov, J.E. Jones, X.X. Xi, E.M. Lysczek, J.M. Redwing, S.Y. Xu, Q. Li, J. Lettieri, D.G. Schlom, W. Tian, X.Q. Pan, Z.K. Liu, Nat. Mater. 1 (2002) 1. [12] X.H. Zeng, A.V. Pogrebnyakov, M.H. Zhu, J.E. Jones, X.X. Xi, S.Y. Xu, E. Wertz, Q. Li, J.M. Redwing, J. Lettieri, V. Vaithyanathan, D.G. Schlom, Z.K. Liu, O. Trithaveesak, J. Schubert, Appl. Phys. Lett. 82 (2003) 2097. [13] W. Jo, J.U. Huh, T. Ohnishi, A.F. Marshall, M.R. Beasley, R.H. Hammond, Appl. Phys. Lett. 80 (2002) 3563.
ARTICLE IN PRESS L. Gu et al. / Journal of Crystal Growth 280 (2005) 602–611 [14] A. Erven, T.H. Kim, M. Muenzenberg, J. Moodera, Appl. Phys. Lett. 81 (2002) 4982. [15] K. Ueda, M. Naito, Appl. Phys. Lett. 79 (2003) 2046. [16] J. Kim, R.K. Singh, J.M. Rowell, N. Newman, L. Gu, D.J. Smith, J. Crystal Growth 270 (2004) 107. [17] B.H. Moeckly, W.S. Ruby, unpublished. [18] Wyckoff, W.G. Ralph, Crystal Structures, vol. 3, Interscience, New York, 1963. [19] C. Cepek, R. Macovez, M. Sancrotti, L. Petaccia, R. Larciprete, S. Lizzit, A. Goldoni, Appl. Phys. Lett. 85 (2004) 976. [20] J.S. Wu, N. Jiang, B. Jiang, J.C.H. Spence, A.V. Pogrebnyakov, J.M. Redwing, X.X. Xi, Appl. Phys. Lett. 85 (2004) 1155. [21] W. Tian, X.Q. Pan, S.D. Bu, D.M. Kim, J.H. Choi, S. Patnaik, C.B. Eom, Appl. Phys. Lett. 81 (2002) 685.
611
[22] Y. Zhu, L. Wu, V. Volkov, Q. Li, G. Gu, A.R. Moodenbaugh, M. Malac, M. Suenaga, J. Tranquada, Physica C 356 (2001) 239. [23] B.H. Moeckly, K.E. Kihlstrom, A. Findikoglu, D.E. Oates, IEEE Trans. Appl. Supercond., in press (June 2005). [24] V. Braccini, A. Gurevich, J. Giencke, M. Jewell, C.B. Eom, D. Larbalestier, A.V. Pogrebnyakov, Y. Cui, B.T. Liu, Y.F. Hu, J.M. Redwing, Q. Li, X.X. Xi, R. Singh, R. Gandikota, J. Kim, B. Wilkens, N. Newman, J. Rowell, B. Moeckly, V. Ferrando, C. Tarantini, D. Marr, M. Putti, C. Ferdeghini, R. Vaglio, E. Haanappel, Phys. Rev. B 71 (2005) 12504. [25] J.M. Rowell, Nature Mater. 1 (2002) 5. [26] R.K. Singh, N. Newman, B.H. Moeckly, unpublished. [27] W.K. Park, L.H. Greene, B.H. Moeckly, J.M. Rowell, unpublished.
Journal of Crystal Growth 280 (2005) 602–611 www.elsevier.com/locate/jcrysgro
Electron microscopy studies of epitaxial MgB2 superconducting thin films grown by in situ reactive evaporation Lin Gua, Brian H. Moecklyb, David J. Smitha, a
Center for Solid State Science, Arizona State University, Tempe, AZ 85287-1704, USA b Superconductor Technologies, Inc., Santa Barbara, CA 93111, USA Received 12 February 2005; accepted 14 March 2005 Available online 23 May 2005 Communicated by T.F. Kuech
Abstract The morphology and chemistry of epitaxial MgB2 thin films grown using reactive Mg evaporation on different substrates have been characterized by transmission electron microscopy methods. For polycrystalline alumina and sapphire substrates with different surface planes, an MgO transition layer was found at the interface region. No such layer was present for films grown on MgO and 4-H SiC substrates, and none of the MgB2 films had any detectable oxygen incorporation nor MgO inclusions. High-resolution electron microscopy revealed that the growth orientation of the MgB2 thin films was closely related to the substrate orientation and the nature of the intermediary layer. Electrical measurements showed that very low resistivities (several mO cm at 300 K) and high superconducting transition temperatures (38 to 40 K) could be achieved. The correlation of electrical properties with film microstructure is briefly discussed. r 2005 Elsevier B.V. All rights reserved. PACS: 68.37.Lp; 68.55.Jk; 74.70.Ad Keywords: A1. Characterization; A1. Crystal morphology; B2. Superconducting materials
1. Introduction The recent discovery of superconducting MgB2 [1] holds much promise for the development of Corresponding author. Tel.: +1 4809654540;
fax: +1 4809659004. E-mail address: [email protected] (D.J. Smith).
superconducting applications and devices operating at medium temperatures. The simple crystal structure, relatively high superconducting transition temperature, Tc, of about 40 K [2], long coherence length of about 50 A˚ [3], metallic charge carrier density [4], and strongly linked current flow in polycrystalline forms [5,6] are all characteristics that facilitate the use of this material for active and
0022-0248/$ - see front matter r 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.jcrysgro.2005.03.066
ARTICLE IN PRESS L. Gu et al. / Journal of Crystal Growth 280 (2005) 602–611
passive devices, detectors, and current carrying applications. Despite intensive efforts and the conspicuous progress that has recently been made, the growth of single-crystal MgB2 remains a non-trivial challenge because of its tendency to oxidize under certain conditions and the volatility of the Mg component over a large temperature range. Several different shapes of MgB2 samples have been successfully developed, including bulk [1], wire [6], and thin films [4]. For most thin film applications, it is preferable to deposit MgB2 completely in situ without involving in situ or ex situ annealing processes under high Mg vapor pressure at elevated temperature [7]. Earlier MgB2 thin films grown by pulsed-laser deposition (PLD) with ex situ annealing had Tcs ranging from 29–34 K [8] to 39 K [4]. In situ post-annealed films deposited by PLD had Tcs of 22–24 K [9]. Electron beam evaporation (e-beam) followed by postgrowth annealing provided a path towards fabrication of large-area thin films [10]. In order to generate the high Mg vapor pressure required for the thermodynamically stable MgB2 phase, hybrid physical–chemical vapor deposition (HPCVD) techniques were developed by combining physical vapor deposition (PVD) and chemical vapor deposition (CVD) to produce high-quality epitaxial thin films on selected substrates [11,12]. In situ film growth using molecular beam epitaxy (MBE) has also been reported [13–16], but the growth temperature must be low to ensure that Mg does not desorb rapidly from the film surface. Contrary to growth of high-temperature superconductors, highly epitaxial growth of MgB2 films is not always necessary in order to achieve good electrical properties [17]. Some substrates are, however, amenable to good quality epitaxial growth. MgB2 has a hexagonal structure with space group P6/mmm and lattice constants of ( c ¼ 3:524 A ( [18]. Sapphire (Al O ) a ¼ 3:086 A; 2 3 ( c ¼ 12:991 A ( with lattice constants of a ¼ 4:758 A; [18] serves as one of the major substrates for MgB2 thin film deposition, despite the relatively large lattice mismatch [4,10,11]. MgO [9,10] and Mg [19] have been used for MgB2 thin film growth because of their compatibility with elemental magnesium. 6H-SiC [9] with an a lattice spacing of 3.08 A˚, and
603
4H-SiC [11,20] with an a lattice spacing of 3.073 A˚, have a lattice mismatch with MgB2 of less than 1%, and should be ideal for high-quality epitaxial growth. Si and SrTiO3 substrates have also been used in studies of the electrical properties of MgB2 thin films [9]. The growth of high-quality MgB2 films on a number of single crystalline and polycrystalline substrates has also been demonstrated [17]. It has been shown that the use of different substrates affects the interfacial chemistry and affects the crystal growth [16,20,21]. The transmission electron microscope (TEM) provides a range of techniques for studying the local morphology and chemistry of MgB2 thin films. Previous analysis of MgB2 thin films has mainly concentrated on whether or not interfacial regions had a thin MgO or MgAl2O4 layer [20,21]. Bulk MgB2 studies have revealed that MgO as the major secondary phase is responsible for the formation of dislocations and other defects [22]. Systematic studies of changes in MgB2 film morphology, interfacial structure and local chemistry for several different substrates are presented and discussed in this work.
2. Experimental details 2.1. Sample growth We have developed a reactive evaporation technique for MgB2 thin film deposition that directly addresses the major problems associated with complete in situ growth of this material [17]. The method centers on the development of a rotating blackbody heater with an internal pocket filled with Mg vapor. This design maintains a localized source of Mg vapor near the substrate(s) while maintaining high vacuum in the rest of the deposition chamber. Boron from a single source is deposited by electron beam evaporation. The MgB2 phase then forms naturally by adsorptionlimited growth. Large-area film growth and double-sided deposition are both possible using this method, and an intermediate growth temperature range of 400–600 1C is accessible. MgB2 films have been deposited on a multitude of singlecrystal, polycrystalline, and metallic substrates up
ARTICLE IN PRESS 604
L. Gu et al. / Journal of Crystal Growth 280 (2005) 602–611
to 400 in diameter. This technique enables growth of high-quality MgB2 films on any substrate for which there is not a chemical reaction, including a large number of single-crystal substrates as well as polycrystalline alumina, stainless steel, and buffered Si. These films regularly display Tc values over 39 K, and their resistivity values are among the lowest reported in the literature. The films are also very clean as exhibited by low microwave loss [23], low Hc2 values [24], and a lack of oxygen or carbon contamination. In addition, the films are relatively stable with respect to patterning and upon exposure to water. All the films in this study were deposited at a growth rate of about 1 A˚/s and at a substrate temperature of 550 1C, a temperature which is generally inaccessible by other MgB2 thin film growth techniques. Table 1 lists some key electrical properties of these films. 2.2. TEM sample preparation Since MgB2 thin films can be highly susceptible to hydro-reaction, the procedures and chemicals normally used for TEM cross-section sample preparation had to be modified. The isocut fluid used during wafer slicing was replaced by polishing oil for cutting lubrication, and chloroform was used instead of acetone for general cleaning before the samples were glued together for cross-sectioning. Samples were ground to about 100 mm in the presence of polishing oil, followed by double-sided dimple polishing. Further thinning was carried out using a precision ion polishing system at 4 keV ion energy followed by a short period of final cleaning at 2.5 keV. In this work, bright-field diffraction-contrast imaging, selected-area electron diffraction and
high-resolution electron microscopy for general morphology and defect characterization were performed using a JEOL 4000EX (operated at 400 keV) with an interpretable resolution of 1.7 A˚. Interface chemistry and film uniformity were analyzed by electron-energy-loss spectroscopy (EELS) and energy-dispersive X-ray spectroscopy (EDXS) using a Philips CM200 operated in the focused probe mode with a spot size of 1 nm.
3. Experimental results In this study, we have performed a systematic investigation of the morphology of epitaxial MgB2 thin films grown on five different substrates. Polycrystalline material and columnar growth was generally observed throughout the epilayer. The presence of an Mg-O transition layer at the interfaces with c-plane sapphire, r-plane sapphire and polycrystalline alumina substrates had a significant impact on the growth orientation of the subsequent MgB2 layer. 3.1. c-plane sapphire MgB2 deposited on c-plane sapphire showed a growth discontinuity with the presence of an oxide layer between the film and the substrate, consistent with other reports [20,21]. Fig. 1(a) is a bright-field diffraction-contrast image showing the morphology of a 300-nm-thick MgB2 film and the intermediate layer, and the inset is the corresponding selected-area electron diffraction (SAED) pattern. It is clear from the image that the epitaxial MgB2 thin film is columnar and polycrystalline.
Table 1 Properties of MgB2 films discussed in the text, including the zero-resistance Tc, DC resistivity values at 300 and at 40 K, and the film thickness Substrate
Tc (K)
r300 K (mO cm)
r40 K (mO cm)
Thickness (nm)
c-plane sapphire r-plane sapphire Polycrystalline alumina MgO 4H-SiC
38.2 39.1 39.0 38.5 39.5
14.5 10.2 9.5 12.3 25.5
6.5 2.6 2.2 4.4 12.3
300 500 500 500 500
All films were deposited by reactive evaporation at 550 1C.
ARTICLE IN PRESS L. Gu et al. / Journal of Crystal Growth 280 (2005) 602–611
605
or so. This morphology is presumably due to the graphite-like layered nature of MgB2 and the relatively high compatibility of elemental Mg with respect to the Mg distribution around the interface [9,10,19]. The high-resolution electron micrograph in Fig. 2 shows the interface microstructure in greater detail. A cubic phase aligned in a [1 1 0]-type orientation and a structurally coherent grain boundary with slight off-axis tilt are observed. The intermediate layer has an overall darker appearance suggesting either greater defect density than the sapphire substrate and/or higher packing density with heavier elements overall than those present in the MgB2 layer above. Fig. 3 shows an annular-dark-field (ADF) electron micrograph (a) of MgB2 grown on c-plane sapphire, together with the corresponding EDXS (b) and EELS (c) line profiles. Spectra were acquired at the location indicated in (a). In the ADF imaging mode, image contrast is proportional to atomic number, and heavier material would thus appear to be brighter. The Al signal
Fig. 1. (a) Electron micrograph showing morphology of MgB2 thin film grown on c-plane sapphire with corresponding selected-area electron diffraction pattern inset; and (b) enlargement showing faceting of MgO intermediary layer and initial MgB2 growth region.
Grain sizes range in width from 50 to 100 nm, and the top surface is relatively bumpy with height variations of as much as 20 nm and a rootmean-square roughness of a few nm. As shown by the enlargement in Fig. 1(b), the initial growth of MgB2 is highly faceted with respect to the underlying transition layer for the first 100 nm
Fig. 2. High-resolution lattice image showing interface between sapphire substrate and unintentional MgO intermediary layer.
ARTICLE IN PRESS L. Gu et al. / Journal of Crystal Growth 280 (2005) 602–611
606
(a)
EDX profile EELS profile
200 nm (b)
500
Mg Al O
Counts
400 300 200 100 0 0
50 Position (nm)
100
(c) B O
Counts (a.u.)
70000
was found to diminish rapidly across the intermediate layer as the probe was moved away from the sapphire substrate with step sizes of about 4 nm. At almost the same location where the Al signal dropped to zero, the O profile also exhibited a drop before finally decreasing to zero after another 50 nm. The rapid increase of the Mg signal combined with the O signal indicated that the major component of the intermediate layer was MgO. The EELS line profile confirmed this result by excluding the presence of boron at least within the detection limit of the EELS technique. The presence of the MgB2 phase became evident after termination of the unintentionally grown MgO transition layer. The relative lower concentration of Mg in the MgB2 film compared with that measured in the MgO layer is consistent with the ADF and bright-field image contrast. Contrary to the observation of MgO as a major secondary phase in bulk MgB2 [22], it appears that the epitaxial MgB2 material here was not impacted by any deterioration by oxidation. This is likely because little or no MgO exists in our Mg vapor pocket during growth, with any gettering of residual oxygen by Mg escaping the heater and depositing on cooler surfaces in the vacuum chamber. Electrical measurements of this film indicated a Tc value of 38.2 K, and low resistivity values of 14.5 mO cm at 300 K, and 6.5 mO cm at 40 K. However, higher Tc values and even lower resistivities were observed for films deposited on other substrates. 3.2. r-plane sapphire
35000
0 0
15
30
45
59
Position (nm) Fig. 3. (a) Annular-dark-field electron micrograph of MgB2 grown on c-plane sapphire, with corresponding EDX (b), and EELS (c), line profiles showing quantitative chemical distribution.
Fig. 4(a) shows the microstructure of an epitaxial 500-nm-thick MgB2 thin film grown on r-plane sapphire, and the inset shows the corresponding SAED pattern. An intermediate MgO layer is again present at the substrate interface. The thickness of the MgO layer is larger than that observed for growth on c-plane sapphire. The surface roughness and grain sizes ranging from 80 to 120 nm are also visibly larger. There is a significant difference of the columnar growth morphology relative to the films grown on c-plane sapphire in that the growth direction is tilted away from the substrate normal, leading to an inclined
ARTICLE IN PRESS L. Gu et al. / Journal of Crystal Growth 280 (2005) 602–611
607
texture of the subsequent MgB2 film. Comparison between the MgO lattice planes visible in Fig. 4(b) and in Fig. 2 reveals that the MgO intermediate layer here has been rotated by 301 relative to the case for the c-plane sapphire in Fig. 2. These different orientations of the intermediate layer have determined the subsequent MgB2 growth orientation. Thus, the difference in MgB2 growth morphology observed for the sapphire substrates is primarily because of the different cutting plane. For MgB2 thin films grown on non-latticematched substrates, growth on r-plane sapphire exhibits excellent properties. Electrical measurements indicated a Tc value of 39.1 K for this film, which was one of the highest among the five films studied here. The resistivity values of this film were rð300 KÞ ¼ 10:2 mO cm and rð40 KÞ ¼ 2:6 mO cm. This residual resistivity value is among the lowest reported in the literature for thin films, indicating that the films are characterized by excellent connectivity and low scattering from impurities or grain boundaries [25].
3.3. Polycrystalline alumina
Fig. 4. (a) Cross-sectional electron micrograph showing morphology of MgB2 thin film grown on r-plane sapphire with corresponding electron diffraction pattern inset; and (b) highresolution lattice image of interfacial oxygen-diffusion region.
film morphology. The origin of the non-vertical microstructure can be attributed to the MgO intermediate layer which establishes the growth
Fig. 5(a) was the microstructure of an MgB2 thin film grown on polycrystalline alumina. An MgO transition layer with a moderate thickness of 50 nm was observed at the substrate surface. The surface roughness of the upper film is excellent and would be favorable for graded-structure device applications. Moire´ fringes were observed in the MgB2 film indicating overlapping grains with different lattice spacings and/or orientations. The growth direction was again tilted away from the substrate normal and the grain size was larger than that observed for the different sapphire cutting planes, ranging from 200 to 400 nm. This larger grain size and consequent reduced intergrain scattering is likely responsible for our lowest observed residual resistivity of rð40 KÞ ¼ 2:2 mO cm. A Tc value of 39.0 K was measured for these epitaxial films grown on alumina, indicating high quality material. Fig. 5(b) shows a high-resolution lattice-image of the interface and the MgO transition region. It is clear that the growth direction of MgO on the alumina substrate is different from the two previous sapphire
ARTICLE IN PRESS 608
L. Gu et al. / Journal of Crystal Growth 280 (2005) 602–611
layer, where the oxygen appears to have been supplied by the O-containing substrate. For growth on an MgO substrate, an MgO intermediate layer would not be expected. Indeed, no apparent transition layer is visible in Fig. 6(a), though the region close to the substrate is highly defective. SAED patterns reveal excellent alignment of the diffraction spots for the MgB2 thin film and the MgO substrate. No extra ring-like polycrystalline patterns were visible except for slight elongation of the MgB2 spots. Threading defects were observed throughout the film, and no major Moire´ fringes were detected. These are promising characteristics for epitaxial growth of high-quality superconducting films. Fig. 6(b) is a high-resolution lattice image from the region of the interface which shows no sign of any secondary phase. The growth direction is aligned exactly parallel to the substrate normal shown in Fig. 6(c) and is aligned according to the MgO [1 0 0] orientation. Electrical characterization of this film grown on MgO showed a small decrease of Tc to 38.5 K and slight increase of rð40 KÞ to 4.4 mO cm [rð300 KÞ ¼ 12:3 mO cm]. 3.5. 4-H SiC
Fig. 5. (a) Electron micrograph showing morphology of MgB2 thin film grown on a polycrystalline alumina substrate with corresponding electron diffraction pattern inset; and (b) highresolution lattice image of the interfacial oxygen-diffusion region.
substrates, and establishes the texture for growth of the subsequent MgB2 layer. 3.4. MgO A common observation for the three previous substrates was the presence of an MgO transition
4-H SiC has a very close lattice match with MgB2. From the diffraction-contrast micrograph shown in Fig. 7(a), no obvious MgO intermediate layer was present. However, the MgB2 film exhibits a granular morphology with fine grains of MgB2 separating the substrate and the upper columnar grains. Almost one-third of the total film thickness is occupied by the granular transition region of about 100 nm. The width of the columnar grains above this region is the smallest of all five films studied. However, the film surface is comparatively flat and should be suitable for growth of graded structures. Fig. 7(b) shows a high-resolution lattice-image of the MgB2 faceted microstructure close to the interface region, which was not observed for the growth on the MgO [1 0 0] substrate. A possible explanation is the different substrate structure and orientation, similar to inclined growth of MgB2 on the intermediate layer, for which the rotated grains are mainly due to the different MgO transition layer orientation.
ARTICLE IN PRESS L. Gu et al. / Journal of Crystal Growth 280 (2005) 602–611
609
Electrical measurements revealed that the Tc of this film is 39.5 K, the highest value of the films studied here. However, the resistivity values are the worst, with rð300 KÞ ¼ 25:5 mO cm and rð40 KÞ ¼ 12:3 mO cm. The high rð40 KÞ value is likely to be due to increased intergrain scattering resulting from the fine grain morphology. The higher Dr value [rð300 KÞ rð40 KÞ ¼ 13:2 mO cm] of this film in comparison to the vast majority of our films [Dr ¼ 7 to 9 mO cm] may be due to decreased connectivity throughout the fine-grained region [25].
4. Discussion
Fig. 6. (a) Electron micrograph showing epitaxial MgB2 thin film grown on an MgO substrate with corresponding electron diffraction pattern inset. High-resolution lattice images of: (b) interface without distinctive extra layers; (c) top MgB2 layer with [0 0 0 1] growth direction.
High-quality MgB2 thin films have been successfully deposited on five different substrates using the reactive Mg evaporation technique. The morphology and chemistry of the highly epitaxial films were characterized using transmission electron microscopy. An MgO transition layer was detected immediately above the substrate for films grown on c-plane and r-plane sapphire and polycrystalline alumina, but no such layer was observed for MgO and 4-H SiC substrates. The layer does not appear to form as a result of reaction with residual oxygen in the deposition chamber, since this layer is not observed for growth on the non-oxide substrates. Moreover, MgO inclusions are not observed in the bulk of the MgB2 films. However, during our normal growth process, the heated substrates are exposed to Mg vapor prior to B evaporation. We would therefore surmise that the MgO layer growth during this stage of the process via out-diffusion of oxygen from the substrate. We have since tested this idea by first evaporating a small amount of B onto the substrate surface prior to exposure to Mg vapor. In this case, analysis by Rutherford backscattering detected no MgO layer at the interface of MgB2 films grown on sapphire [26]. Furthermore, we have not yet observed any negative consequences of this MgO layer, since the demanding microwave properties of our films are excellent [23]. The epitaxial films had columnar morphology in all cases. The growth direction of the films was strongly influenced by the MgO intermediate layer orientation, and the width of the grains varied
ARTICLE IN PRESS 610
L. Gu et al. / Journal of Crystal Growth 280 (2005) 602–611
conducting gaps [27]. Furthermore, electrical measurements on these films reveal very low resistivities and high superconducting transition temperatures, indicating a compact and stronglylinked film morphology. These results are consistent with the TEM observations.
Acknowledgments We acknowledge use of facilities in the John M. Cowley Center for High Resolution Electron Microscopy at Arizona State University.
References
Fig. 7. (a) Electron micrograph showing the morphology of MgB2 thin film grown on 4H-SiC substrate; and (b) highresolution lattice-image adjacent to the interface showing faceted growth morphology.
depending on the different substrate and the transition layer thickness and orientation. Pointcontact tunneling spectra taken across the surface of MgB2 films grown on r-plane sapphire have indicated the clear presence of both MgB2 super-
[1] J. Nagamatsu, N. Nakagawa, T. Muranaka, Y. Zenitani, J. Akimitsu, Nature (London) 410 (2001) 63. [2] S.L. Bud’ko, G. Lapertot, C. Petrovic, C.E. Cunningham, N. Anderson, P.C. Canfield, Phys. Rev. Lett. 86 (2001) 1877. [3] D.K. Finnemore, J.E. Osterson, S.L. Bud’ko, G. Lapertot, P.C. Canfield, Phys. Rev. Lett. 86 (2001) 2420. [4] W.N. Kang, H.J. Kim, E.M. Choi, C.U. Jung, S.L. Lee, Science 292 (2001) 1521. [5] D.C. Larbalestier, L.D. Cooley, M.D. Rikel, A.A. Polyanskii, J. Jiang, S. Patnaik, X.Y. Cai, D.M. Feldmann, A. Gurevich, A.A. Squitieri, M.T. Naus, C.B. Eom, E.E. Hellstrom, R.J. Cava, K.A. Regan, N. Rogado, M.A. Hayward, T. He, J.S. Slusky, R. Khalifah, K. Inumaru, M. Haas, Nature 410 (2001) 186. [6] P.C. Canfield, D.K. Finnemore, S.L. Bud’ko, J.E. Ostenson, G. Lapertot, C.E. Cunningham, C. Petrovic, Phys. Rev. Lett. 86 (2001) 2423. [7] Z.K. Liu, D.G. Schlom, Q. Li, X.X. Xi, Appl. Phys. Lett. 78 (2001) 3678. [8] A. Berenov, Z. Lockman, X. Qi, J.L. MacManus-Driscoll, Y. Bugoslavsky, L.F. Cohen, M.H. Jo, N.A. Stelmashenko, V.N. Tsaneva, M. Kambara, N. Hari Babu, D.A. Cardwell, M.G. Blamire, Appl. Phys. Lett. 79 (2001) 4001. [9] D.H.A. Blank, H. Hilgenkamp, A. Brinkman, D. Mijatovic, G. Rijnders, H. Rogalla, Appl. Phys. Lett. 79 (2001) 394. [10] S.H. Moon, J.H. Yun, H.N. Lee, J.I. Kye, H.G. Kim, W. Chung, B. Oh, Appl. Phys. Lett. 79 (2001) 2429. [11] X.H. Zeng, A.V. Pogrebnyakov, A. Kotcharov, J.E. Jones, X.X. Xi, E.M. Lysczek, J.M. Redwing, S.Y. Xu, Q. Li, J. Lettieri, D.G. Schlom, W. Tian, X.Q. Pan, Z.K. Liu, Nat. Mater. 1 (2002) 1. [12] X.H. Zeng, A.V. Pogrebnyakov, M.H. Zhu, J.E. Jones, X.X. Xi, S.Y. Xu, E. Wertz, Q. Li, J.M. Redwing, J. Lettieri, V. Vaithyanathan, D.G. Schlom, Z.K. Liu, O. Trithaveesak, J. Schubert, Appl. Phys. Lett. 82 (2003) 2097. [13] W. Jo, J.U. Huh, T. Ohnishi, A.F. Marshall, M.R. Beasley, R.H. Hammond, Appl. Phys. Lett. 80 (2002) 3563.
ARTICLE IN PRESS L. Gu et al. / Journal of Crystal Growth 280 (2005) 602–611 [14] A. Erven, T.H. Kim, M. Muenzenberg, J. Moodera, Appl. Phys. Lett. 81 (2002) 4982. [15] K. Ueda, M. Naito, Appl. Phys. Lett. 79 (2003) 2046. [16] J. Kim, R.K. Singh, J.M. Rowell, N. Newman, L. Gu, D.J. Smith, J. Crystal Growth 270 (2004) 107. [17] B.H. Moeckly, W.S. Ruby, unpublished. [18] Wyckoff, W.G. Ralph, Crystal Structures, vol. 3, Interscience, New York, 1963. [19] C. Cepek, R. Macovez, M. Sancrotti, L. Petaccia, R. Larciprete, S. Lizzit, A. Goldoni, Appl. Phys. Lett. 85 (2004) 976. [20] J.S. Wu, N. Jiang, B. Jiang, J.C.H. Spence, A.V. Pogrebnyakov, J.M. Redwing, X.X. Xi, Appl. Phys. Lett. 85 (2004) 1155. [21] W. Tian, X.Q. Pan, S.D. Bu, D.M. Kim, J.H. Choi, S. Patnaik, C.B. Eom, Appl. Phys. Lett. 81 (2002) 685.
611
[22] Y. Zhu, L. Wu, V. Volkov, Q. Li, G. Gu, A.R. Moodenbaugh, M. Malac, M. Suenaga, J. Tranquada, Physica C 356 (2001) 239. [23] B.H. Moeckly, K.E. Kihlstrom, A. Findikoglu, D.E. Oates, IEEE Trans. Appl. Supercond., in press (June 2005). [24] V. Braccini, A. Gurevich, J. Giencke, M. Jewell, C.B. Eom, D. Larbalestier, A.V. Pogrebnyakov, Y. Cui, B.T. Liu, Y.F. Hu, J.M. Redwing, Q. Li, X.X. Xi, R. Singh, R. Gandikota, J. Kim, B. Wilkens, N. Newman, J. Rowell, B. Moeckly, V. Ferrando, C. Tarantini, D. Marr, M. Putti, C. Ferdeghini, R. Vaglio, E. Haanappel, Phys. Rev. B 71 (2005) 12504. [25] J.M. Rowell, Nature Mater. 1 (2002) 5. [26] R.K. Singh, N. Newman, B.H. Moeckly, unpublished. [27] W.K. Park, L.H. Greene, B.H. Moeckly, J.M. Rowell, unpublished.