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Molecular beam epitaxy growth and microstructure of thin superconducting Bi2Sr2CaCu2Ox films
PHYSICA ELSEVIER
Physica C 253 (1995) 383-390
Molecular beam epitaxy growth and microstructure of thin superconducting Bi2Sr2CaCu20 x films A. Brazdeikis "*, A. Vailionis a, A.S. Flodstr~Sm a, C. Tr~eholt b a Department of Physics, Materials Physics, Royal Institute of Technology, S-10044 Stockholm, Sweden b National Centre for HREM, Delft University, Rotterdamseweg 137, 2628 AL Delft, The Netherlands
Received 6 July 1995;revised manuscript received 2 August 1995
Abstract
The microstructure of molecular beam epitaxy (MBE) grown Bi2Sr2CaCu20 x films has been studied by reflection high-energy electron diffraction (RHEED), high-resolution electron microscopy (HREM) and X-ray diffraction (XRD). In situ recorded RHEED images of as-grown films show two-dimensional growth, 90° oriented domains (twist domains) and an incommensurate superstructure. For the first time, the presence of both 2201 and 2223 stacking faults in Bi2Sr2CaCu20 x thin films is reported. These local structural defects are not easily observed by standard XRD techniques. To examine the structure of the films quantitatively, a general one-dimensional XRD model was fitted to the experimental XRD data. The model considered changes in peak intensities, positions and line-widths, and thus allowed a quantitative determination of the structural properties of the high-Tc superconducting thin films.
1. I n t r o d u c t i o n
In recent years, the motivation to study epitaxially grown oxide films has increased due to the discovery of high-temperature superconductivity. Intrinsic properties of these superconductors and also of many other oxide compounds make them difficult to grow as epitaxial thin films. The compounds contain a relatively large number o f elements and hence they exhibit complex crystal structures and phase diagrams. The bismuth alkaline-earth cuprate system, Bi2Sr2Can_lCUnO4+2n+~ c a n be considered as a model system because it permits all aspects of growth to be studied. Different phases for this system are
* Corresponding author.
characterized by the number, n, of CHO 2 and Ca planes in the unit cell. The Bi2SrzCuO 6 (n = 1), Bi2Sr2CaCu20 8 (n = 2) and Bi2SrzCazCu3010 (n = 3) compounds are thermodynamically stable and show superconducting transition temperatures, T~, at 20, 85 and 110 K, respectively, as determined both from electrical resistivity and magnetization measurements. Different methods of synthesis have been used to prepare superconducting BiSrCaCuO films. Applications of superconductors have been limited to very specific areas, mainly because of poor reproducibility and lack of structurally " g o o d " films. To grow compact, homogeneous superconducting films and multilayers with smooth surfaces and planar interfaces, a detailed knowledge of the nucleation mechanisms, growth modes and resulting film microstructure is needed [1]. Molecular beam epitaxy (MBE) is an attractive synthesis technique because
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of its extreme dimensional control and the possibility of using in situ diagnostic techniques, such as reflection high-energy electron diffraction (RHEED). The development of the MBE technique to grow oxide films has resulted in epitaxial growth of superconducting thin films, novel oxide materials and heterostructures [2-7]. The advanced technique of high-resolution transmission electron microscopy (HREM) has enabled a better microstructural characterization of superconducting BiSrCaCuO thin films [8-11]. HREM has recently been used to successfully indentify non-unit-cell nucleation in the initial stage of film growth and the absence of incommensurate modulations at film/substrate interface [12,13]. Efforts to characterize local crystal defects and further reduce their density in superconducting films are important. This involves extensive research and the introduction of other complementary techniques. In the present work, MBE growth and the resulting microstructure of Bi2Sr2CaCu20 x (2212) films on MgO substrates have been investigated. Three issues of interest were considered: (1) in situ RHEED characterization of film growth; (2) thin film and film/substrate interface characterization by HREM; (3) quantitative characterization of stacking faults by a one-dimensional kinematic X-ray diffraction (XRD) model.
2. Experimental Experiments were conducted in an UHV MBE deposition chamber by coevaporation from Knudsen effusion cells of the metallic elements Bi, Sr, Ca and Cu. The MBE system was designed for the in situ synthesis of bismuth alkaline-earth superconducting cuprates using nitrogen dioxide, NO 2, as oxidant. The molecular beam evaporation rate was calibrated with a quartz-crystal thickness monitor in the substrate position. The MgO (100) substrates were glued to the centre of an Inconel alloy substrate holder with silver paste and heated by a halogen lamp with a copper reflector. To prevent changes in the substrate temperature during growth due to heat absorption in the film, the temperature was controlled by a thermocouple calibrated with an optical pyrometer. The heating arrangement allowed a reliable control
of the growth temperature and of the heating and cooling rates. Films were grown at typically 6 9 0 700°C and 1 × 10 -4 Torro N O 2. Film thicknesses ranged between 30-1000 A. Details of the experimental setup and of the growth conditions are reported elsewhere [14,15]. A 10 keV RHEED system was used to study growth orientation, crystal structure and the surface morphology of the films. For transmission electron microscopy investigations a cross section was made from the film/substrate sample by cutting it with a diamond blade saw, mechanically grinding and polishing the cut-off piece with a diamond paste down to a thickness of about 10 tzm. Finally, the cross section was ion-milled without rotation until electron transparency (less than 1000 A) was reached. The parameters used for ionmilling were first voltage 5 kV, gun current 0.5 mA and an angle of 15°, and finally 3 kV, 0.3 mA and 10°. Details of the preparation method can be found in Ref. [16]. HREM was performed with a Philips CM30ST electron microscope equipped with a (Schottky) field emission gun and a side entry goniometer with a tilt range x / y of 250/25 °. The microscope was operated at 300 kV giving a resolution of 2.0 A. The images were normally recorded close to Sherzer focus. XRD data were recorded utilizing an X-ray diffractometer in the Bragg-Brentano geometry using Cu Kct radiation (A = 1.5406 A) and equipped with a graphite monochromator. The XRD data were used to determine the c-axis lattice parameters, the epitaxial film quality, and to study the phase purity. For the quantitative analysis of the 2212 films, a one-dimensional kinematic X-ray diffraction model has been used. The interplanar distances in the unit cell, site occupancies and the number of stacking faults have been used as fitting parameters. The Debye-Waller coefficient, the Lorentz-polarization and absorptions factors are included in the fitting procedure. A detailed mathematical formalism has been published elsewhere [17].
3. Results and discussion In situ heated MgO substrates show RHEED surface patterns with sharp diffraction spots lying on the 0th Laue circle. Such a pattern is obtained because of
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modulations in the films. The domains have a common c-axis, but their a- and b-axes are rotated by 90 ° with respect to each other. The observed superstructure originates from an incommensurate modulation of the rock-salt type (BiO) 2 layers with respect to the S r O - C u O z - C a - C u O z - S r O perovskite slab. XRD analysis of the films shows a c-axis oriented film and the presence of a small amount of CaCuO impurities ( d = 2 . 3 9 3 A). The full-width-at-halfmaximum (FWHM) of the (0010) rocking curves is close to 0.3 °. Zero resistivity temperatures, Tc(R= 0), between 70 and 75 K were recorded for the as-grown 2212films. Critical current density values, Jc, between 1 × l0 s and 6 × 105 A / c m 2 were determined from
Fig. 1. RHEED surface patterns of as-grown 2212-film on MgO (100). Streak splinings (A-shape) (a) due to 90 ° oriented twist domains and the superstructure, and (b) due to incommensurate modulations, are visible. The incident beam is in the [1 I0] (a) and [100] (b) directions.
an atomically flat and periodic MgO (100) surface. In grazing-exit geometry, the diffraction spots become streaks due to surface disorder caused, for instance, by two-dimensional (2D) island film growth. It has already been shown that an increased oxygen activity at higher NO 2 pressures, for a given growth temperature, leads to a 2212-film surface almost free from second-phase precipitates [14]. Clear 2D RHEED surface patterns are usually obtained for as-grown 2212 films. Fig. 1 shows the in situ recorded RHEED pattern of 2212-film (a) along [110] and (b) [100] MgO substrate azimuths. Pronounced splitting of the reciprocal lattice rods and a satellite superstructure surrounding them are clearly visible. This suggest a superposition of 90 ° oriented (twist) domains and the presence of incommensurate
Fig. 2. Experimental HREM image of a 300 A thick MBE-grown 2212-f'flm on a (100)-oriented MgO substrate. Incommensurate modulations with a wave vector at 55 ° angle are clearly visible. The 2201 and 2223 intergrowth defects (stacking faults), double Bi-layers and the incommensurate superlattice are indicated.
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the transport measurements at 10 K and were consistent with arc values obtained from SQUID magnetic measurements [14]. HREM and selected area electron diffracfign images show that 2212-films are grown preferentially with their c-axes perpendicular to the MgO {100} surface, and that both the a- and b-axes are observed when the cross section is viewed in the direction parallel to [010] MgO. Fig. 2 shows an experimental HREM image of a 300 A thick 2212-film on a (100)-oriented MgO substrate. The modulations along the b-axis leading to a wavy pattern are due to displacements of the Bi atoms in adjacent BiO layers in the film [ 18,19]. The wave vectors of these modulations make an angle of 55 ° with the {001} layer planes and are preserved up to the film surface. The structure modulations are very strong, and the local structure deviates significantly from the average structure. It seems likely that the lattice misfit between the perovskite slabs and the (BiO) 2 layers is a driving force for the modulations [20]. The results on the twinning and the modulation are consistent with HREM data of 2212 bulk samples observed by several researchers [21-24]. HREM investigation of the interface shows large areas where the 2212-film grows epitaxially on MgO
and other regions where Ca-Cu-oxide grains precipitated at the film/substrate interface. A HREM image of the f i l m / M g O interface region is shown in Fig. 3. No impurity phases were observed in the bulk of the film as have often been observed in co-evaporated YBaCuO superconductors [25]. It is reasonable to expect that nucleation and growth of precipitates as the 2212 will be affected by growth temperature, oxidation, local surface composition and defects on the substrate surface. The details of the atomic structure at the 2212/MGO interface, however, depend on a number of factors such as bonding and chemistry. Several possible mechanisms can be proposed to explain the interface precipitates that appear in MBE co-evaporated growth. There are a large number of different phases possible in the BiSrCaCuO phase diagram [26]. If thermodynamics determine the reactions on the substrate surface, uncontrolled nucleation of stable impurity precipitates may occur. Such effects can be enhanced by different sticking coefficients of deposited species on a bare MgO surface [27] and on Ca impurities segregated to the MgO surfaces [28], and also by local variations in temperature across the substrate and instabilities in the molecular beam fluxes due to the suppressed sublimation of alkaline-earth from effu-
Fig. 3. ExperimentalHREM image of the 2212-film/MgO interface. The Ca-Cu-oxide grain (P) precipitated at the film/substrate and the different n intergrowth defects are indicated.
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sion cells [29]. Another important factor which can cause precipitation of other (unwanted) phases and is related to the growth kinetics, is the substrate surface defects. There will always be terraces, mono a n d / o r multiple atomic steps on the MgO surface acting as preferred binding sites. The steps are line defects and thus the kinetics of the nucleation will be inferred from the distribution of nuclei. From several lowmagnification HREM images, it was confirmed that the C a - C u - O precipitates appear only occasionally but as rather large grains. This suggests that the impurities possess relatively high growth speeds and, probably, act as a " s i n k " for Cu and Ca during 2212-film MBE co-evaporated growth. HREM investigations show that the 2212-film growth starts with a single B i t layer on the MgO substrate. Similar growth results have recently been reported [8,30]. The incommensurate modulations are found to be present near the f i l m / M g O interface. In contrast, HREM observations of the 2212-film grown on SrTiO 3 [13] show that no modulations are present at the interface. This may be related to the fact that the 2212 nucleates on MgO as a stable and charge-neutral; B i O - S r O - C u O 2 - C a - C u O z - S r O - B i O block, whereas the 2212 at the initial growth stages on SrTiO 3 exhibits non-unit-cell growth [12].
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A(20), deg. Fig. 5. Calculated shift (in 20 deg) of (008) and (0012) Bragg peaks as a function of the 2201/2223 proportion in a 2212-film. The total number of stacking faults is assumed to be 20% by volume.
The 2212-film is found to contain several stacking faults (SF). These are intergrowth defects in which some layers of Cu2CaO 4 in the perovskite slab are replaced by layers of CuaCa206 a n d / o r C u t 2. As a result, 2223 and 2201 stacking faults occur. The 2201 and 2223 intergrowth defects in the 2212-film matrix occur because of the small difference in free energy between the 2212 phase and these phases [31]. Thus, any local variations in the 2212 stoichiometry will result in a stacking fault, which in fact is found to be a common type of defect in these layered compounds. The presence, however, of both 2223 and 2201 intergrowth defects in a thin 2212 film is not trivial. The "defect" 2 2 0 1 / 2 2 1 2 / 2 2 2 3 / 2212 structure (schematically shown in Fi~. 4) has an average c-axis lattice parameter of 30.8 A and the same stoichiometry as pure 2212. In this case, the local disorder cannot easily be distinguished by normal XRD. If only 2201 or 2223 SF's are present as intergrowth defects in 2212-film, a shift of some (00l) XRD peaks in a 0 - 2 0 scan will occur [32]. Equal amounts of 2223 and 2201 stacking faults do not result in a shift of XRD peaks, but alter their intensities and line-widths. In order to quantitatively examine the structure of such films, more advanced XRD methods have to be applied. Here a general
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one-dimensional kinematic X-ray diffraction model is applied to fit the XRD spectra from the MBEgrown 2212-films [17]. The type and amount of SF as well as the unit cell disorder are determined by an iterative fitting of the calculated profiles to the experimental XRD spectra. The measured XRD profiles show Ganssian-type peak shapes. A Gaussian distribution function was used for the continuous interplanar distance fluctuations in the unit cell. The 2201 a n d / o r 2223 phase stacking faults were assumed to be randomly distributed within the films. Fig. 5 displays a calculated shift (in 2 0 deg) of the (008) and (0012) Bragg peaks as a function of the 2201/2223 proportion. The total number of stacking faults is assumed to be 20% of the 2212-film volume. The (008) and (0012) peaks are shifted only if the number of 2201 defects is not equal to the number of 2223 defects. The shifts are towards the (0010) peak when the amount of 2223 SF is higher than that of 2201. The (008) and (0012) peaks are moved apart when the amount of 2223 SF in the film is lower than that of 2201. As can be seen in the calculations in Fig. 5, there is no peak shift when the number of 2201 SF is equal to the number of 2223 SF. Such an interpretation by observing the shift of XRD peaks in 0 - 2 0 scans can only be used to
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Number of SF, % Fig. 6. Calculated F W H M (in 2 0 deg) of (008), (0010.__) and (001..~2) Bragg peaks of a 370 P, thick 2212-film as a function of the total number of stacking faults. The 2201/2223 ratio is kept equal to unity.
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determine which type of stacking fault, 2201 or 2223, dominates in the film. A complete characterization of the SF defects in the 2212 layer requires a determination of the line-widths of the (008), (0010) and (0012) Bragg peaks and their intensities. Fig. 6 gives the calculated (in 2 0 deg) FWHM of (008), (0010) and (0012) peaks of a 370 ,~ thick 2212-film when the SF 2201/2223 ratio is equal to unity. The FWHM of the (008) and (0012) peaks increases with the amount of SF, but does not change significantly for the (0010) peak. The sensitivity of the FWHM of the (008) and (0012) peaks to the presence of small amounts of SF's in the 2212 compound can easily be understood by comparing 0-2 0 spectra of the 2201, 2212 and 2223 phases. The positions of the Bragg peak ((008), (0010) and (0012) for the 2201, 2212 and 2223, respectively) at 20 = 29 ° remain mostly unchanged from one phase to the other. The superimposition of the surrounding peaks certainly broadens, e.g., the (008) and (0012) peaks in 2212 if small amounts of 2201 a n d / o r 2223 are present. The intensity distribution analysis in the 0 - 2 0 XRD profile is not straightforward. The intensities are not only affected by the presence of stacking
A. Brazdeikis et al. / Physica C 253 (1995) 383-390
Table 1 Calculated average c-axis length and the number of 2201 and 2223 stacking faults for three different MBE-grown2212-films Sample Average c-axis Number of stacking faults length (,~) (%) 2201 2223 #154 30.84 44.5 1.2 #159 30.79 2.1 34.7 #166 30.84 3.5 4.6 faults in the film. Structural disorder in the unit cell due to large substitutional effects may lead to intensity changes of some XRD peaks [17]. The fitting revealed that the continuous interploanar distance fluctuations in the 2212-films are 0.04 A. The substitution at Sr 2+ sites by Ca 2+ and by Bi 3+ is approximately 20% and 15%, respectively. The Ca 2+ site is populated by Sr 2+ to 40%. These substitutions increase the interplanar CuO2-Ca distance from 1.65 to 1.70 A, and thus the S r O - C u O 2 distance in the unit cell is contracted from 1.72 to 1.58 A. This contributes to the lower 0012 XRD peak intensity than expected for a perfect 2212-film. The changes in the interplanar distances may also explain why the line-width of this peak in the measured 2212 XRD spectrum is usually observed to be larger than the value calculated from Fig. 6. Experimental XRD spectra of three different MBE-grown 2212 films and the fitted XRD profiles are presented in Fig. 7. The average c-axis length and the number of stacking faults obtained by the iterative fitting are presented in Table l. A large number of the stacking faults are obtained during co-evaporated growth when large variations from the 2212 stoichiometry occur, mainly due to inaccurate calibration of molecular beams, while the presence of only a small number of 2201 and 2223 stacking faults indicates that the stoichiometric 2212 MBE growth is achieved. Any variation in temperature across the substrate and any changes in the molecular beam fluxes, as already discussed, will certainly promote nucleation of local intergrowth defects. Nevertheless, the fact that the 2201 and 2223 stacking faults are present in almost equal amounts indicates that the 2212 growth here is probably accompanied by an intrinsic generation of stacking faults, which is supported by the fact that the 2212 growth conditions are far from the thermodynamic equilibrium and by the energetic proximity o
389
of similar polytype phases. Once the stacking fault, e.g. 2201, is locally generated, the excess Ca and Cu changes the chemical potential in that region. Such a local "off-stoichiometry" condition will, with rather high probability, result in the formation of a 2223 stacking fault. A similar sequence of events can be observed when the 2223 SF is formed first, resulting in the local Ca and Cu deficiency and, thus, followed by the 2201 SF formation.
4. Conclusions MBE growth and the resulting microstructure of films have been studied by RHEED, HREM and XRD. In situ RHEED analysis showed that as-grown films are formed by two-dimensional island growth with 90 ° oriented twist domains and exhibit incommensurate modulations. The observed features were confirmed by HREM. The film-substrate interface studies showed large areas where the film grew epitaxially on the MgO (100) surface, and several regions where C a - C u - O impurity precipitated at the substrate-film interface. The nucleation of precipitates is explained as being governed by the different sticking coefficients of deposited species on MgO surfaces and also by Ca impurities segregated on the MgO surfaces. Such effects could be enhanced by a local variation of the temperature across the substrate and time-dependent changes of molecular beam fluxes. For the first time, both 2201 and 2223 stacking-fault defects were observed to be present in Bi2Sr2CaCu20 x thin films. These local structural defects were examined by a general one-dimensional XRD model. A characteristic feature of the defects was their occurrence in almost eqtial amounts. The appearance of the defects in the 2212-films was explained on the basis of the 2212-film growth at thermodynamic non-equilibrium and the energetic proximity of similar polytype phases. Bi 2S r 2 C a C u 2 0 x
Acknowledgements This work was supported by the Swedish Research Council for Engineering Sciences (TFR) and Swedish National Board for Industrial and Technical Development (NUTEK).
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References [1] For a review, see MRS Bulletin, XIX (9), issue on "Crystal Engineering of High-Tc-Related Oxide Films". [2] For a review, see I. Bozovic, J.N. Eckstein, G.F. Virshup, A. Chaiken, M. Wall, R. Howell and M. Fluss, J. Supercond. 7 (1994) 187, and references therein. [3] Y. Nakayama, I. Tsukada and K. Uchinokura, J. Appl. Phys. 70 (1991) 4371. [4] Y. Kasai and S. Sakai, J. Crystal Growth 115 (1991) 758. [5] M. Kawai, S. Watanabe and T. Hanada, J. Crystal Growth 112 (1991) 745. [6] M. Kanai, T. Kawai and S. Kawai, Jpn. J. Appl. Phys. 31 (1992) L331. [7] A. Brazdeikis, A. Vailionis and A.S. Flodstr~m, Physica C 235-240 (1994) 711. [8] K. Kojima, D.G. Schlom, K. Kuroda, M. Tanioku, K. Hamanaka, J.N. Eckstein and S.J. Harris Jr., Jpn. J. Appl. Phys. 29 (1990) L1638. [9] D.G. Schlom, A.F. Marshall, J.S. Harris Jr., I. Bozovic and J.N. Eckstein, Advances in Superconductivity III, eds. K. Kajimura and Hayakawa, (Springer, Tokyo, 1991) p. 1011. [10] T. Satoh, J. Fujita, T. Yoshitake and H. Igarashi, Physica C 183 (1991) 286. [11] T. Satoh, J. Fujita, T. Yoshitake and H. Igarashi, Physica C 191 (1992) 359. [12] J.G. Wen, T. Mofishita, N. Koshizuka, C. Traeholt and H.W. Zandbergen, Appl. Phys. Lett. 66 (1995) 1830. [13] N.D. Zakharov, D. Hesse, F. Nouvertnd, J. Auge, H. Hoffschulz, J. Drel3en, H.G. Roskos, H. Kurtz and G. Giintherodt, Physica C 245 (1995) 84. [14] A. Brazdeikis, A. Vailionis and A. Flodstrfim, Applied Superconductivity, ed. H.C. Freyhardt, (DGM Informationsgesellschaft, 1993) p. 563. [15] A. Brazdeikis, Licentiate thesis, Dept. of Physics, Royal Institute of Technology, Stockholm, 1994.
[16] C. Traeholt, J.G. Wen, V. Svetchnikov, A. Delsing and H.W. Zandbergen, Physica C 206 (1993) 318. [17] A. Vailionis, A. Brazdeikis and A.S. Flodstr~m, Phys. Rev. B 51 (1995) 3097. [18] H.W. Zandbergen, W.A. Groen, F.C. Mijhoff, G. van Tende1oo and S.A. Amelinckx, Physica C 156 (1988) 325. [19] Y. LePage, W.R. McKinnon, J.-M. Tarascon and P. Barboux, Phys. Rev. B 40 (1989) 6820. [20] P.L. Gai and P. Day, Physica C 152 (1988) 334. [21] Y. Matsui, H. Maeda, Y. Tanaka and S. Hofiuchi, Jpn. J. Appl. Phys. 27 (1988) L372. [22] E.A. Hewat, P. Bordet, J.J. Capponi, C. Chaillout, J.L. Hodeau and M. Marezio, Physica C 153-155 (1988) 619. [23] M. Hervieu, C. Michel, B. Domenges, Y. Laligant, A. Lebail, G. Ferey and B. Raveau, Mod. Phys. Lett. B 2 (1988) 491. [24] For a review, see C.H. Chen, Physical Properties of High Temperature Superconductors II, ed. D.M. Ginsberg, (World Scientific, Singapore, 1990) p. 242. [25] See e . g . V . Matijasevic, P. Rosenthal, K. Shinohara, A.F. Marshall, R.H. Hammond and M.R. Beasley, J. Mater. Res. 6 (1991) 682. [26] For a review, see I. Bozovic and J.N. Eckstein, MRS Bulletin, XX (5) (1995) 32. [27] E.S. Hellman, E.H. Hartford and E.A. Fitzgerald, J. Mater. Res. 7 (1992) 795. [28] R.V. Smilgys, S.W. Robey and C.K. Chiang, J. Vac. Sci. Technol. A 11 (1993) 1361. [29] E.S. HeUman and E.H. Hartford Jr., J. Vac. Sci. Technol. B 12 (1994) 1178. [30] C. Tr~eholt, Ph.D. Thesis, Technical University of Denmark, 1994. [31] J.M. Tarascon, W.R. McKinnon, P. Barboux, D.M. Hwang, B.G. Bagley, L.H. Greene, G.W. Hull, Y. LePage, N. Stoffel and M. Giroud, Phys. Rev. B 38 (1988) 8885. [32] L. Ranno, D. Martinez-Garcla, J. Perribre and P. Barboux, Phys. Rev. B 48 (1993) 13945.
Physica C 253 (1995) 383-390
Molecular beam epitaxy growth and microstructure of thin superconducting Bi2Sr2CaCu20 x films A. Brazdeikis "*, A. Vailionis a, A.S. Flodstr~Sm a, C. Tr~eholt b a Department of Physics, Materials Physics, Royal Institute of Technology, S-10044 Stockholm, Sweden b National Centre for HREM, Delft University, Rotterdamseweg 137, 2628 AL Delft, The Netherlands
Received 6 July 1995;revised manuscript received 2 August 1995
Abstract
The microstructure of molecular beam epitaxy (MBE) grown Bi2Sr2CaCu20 x films has been studied by reflection high-energy electron diffraction (RHEED), high-resolution electron microscopy (HREM) and X-ray diffraction (XRD). In situ recorded RHEED images of as-grown films show two-dimensional growth, 90° oriented domains (twist domains) and an incommensurate superstructure. For the first time, the presence of both 2201 and 2223 stacking faults in Bi2Sr2CaCu20 x thin films is reported. These local structural defects are not easily observed by standard XRD techniques. To examine the structure of the films quantitatively, a general one-dimensional XRD model was fitted to the experimental XRD data. The model considered changes in peak intensities, positions and line-widths, and thus allowed a quantitative determination of the structural properties of the high-Tc superconducting thin films.
1. I n t r o d u c t i o n
In recent years, the motivation to study epitaxially grown oxide films has increased due to the discovery of high-temperature superconductivity. Intrinsic properties of these superconductors and also of many other oxide compounds make them difficult to grow as epitaxial thin films. The compounds contain a relatively large number o f elements and hence they exhibit complex crystal structures and phase diagrams. The bismuth alkaline-earth cuprate system, Bi2Sr2Can_lCUnO4+2n+~ c a n be considered as a model system because it permits all aspects of growth to be studied. Different phases for this system are
* Corresponding author.
characterized by the number, n, of CHO 2 and Ca planes in the unit cell. The Bi2SrzCuO 6 (n = 1), Bi2Sr2CaCu20 8 (n = 2) and Bi2SrzCazCu3010 (n = 3) compounds are thermodynamically stable and show superconducting transition temperatures, T~, at 20, 85 and 110 K, respectively, as determined both from electrical resistivity and magnetization measurements. Different methods of synthesis have been used to prepare superconducting BiSrCaCuO films. Applications of superconductors have been limited to very specific areas, mainly because of poor reproducibility and lack of structurally " g o o d " films. To grow compact, homogeneous superconducting films and multilayers with smooth surfaces and planar interfaces, a detailed knowledge of the nucleation mechanisms, growth modes and resulting film microstructure is needed [1]. Molecular beam epitaxy (MBE) is an attractive synthesis technique because
0921-4534/95/$09.50 © 1995 Elsevier Science B.V. All rights reserved SSDI 0921-4534(95)00511-0
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of its extreme dimensional control and the possibility of using in situ diagnostic techniques, such as reflection high-energy electron diffraction (RHEED). The development of the MBE technique to grow oxide films has resulted in epitaxial growth of superconducting thin films, novel oxide materials and heterostructures [2-7]. The advanced technique of high-resolution transmission electron microscopy (HREM) has enabled a better microstructural characterization of superconducting BiSrCaCuO thin films [8-11]. HREM has recently been used to successfully indentify non-unit-cell nucleation in the initial stage of film growth and the absence of incommensurate modulations at film/substrate interface [12,13]. Efforts to characterize local crystal defects and further reduce their density in superconducting films are important. This involves extensive research and the introduction of other complementary techniques. In the present work, MBE growth and the resulting microstructure of Bi2Sr2CaCu20 x (2212) films on MgO substrates have been investigated. Three issues of interest were considered: (1) in situ RHEED characterization of film growth; (2) thin film and film/substrate interface characterization by HREM; (3) quantitative characterization of stacking faults by a one-dimensional kinematic X-ray diffraction (XRD) model.
2. Experimental Experiments were conducted in an UHV MBE deposition chamber by coevaporation from Knudsen effusion cells of the metallic elements Bi, Sr, Ca and Cu. The MBE system was designed for the in situ synthesis of bismuth alkaline-earth superconducting cuprates using nitrogen dioxide, NO 2, as oxidant. The molecular beam evaporation rate was calibrated with a quartz-crystal thickness monitor in the substrate position. The MgO (100) substrates were glued to the centre of an Inconel alloy substrate holder with silver paste and heated by a halogen lamp with a copper reflector. To prevent changes in the substrate temperature during growth due to heat absorption in the film, the temperature was controlled by a thermocouple calibrated with an optical pyrometer. The heating arrangement allowed a reliable control
of the growth temperature and of the heating and cooling rates. Films were grown at typically 6 9 0 700°C and 1 × 10 -4 Torro N O 2. Film thicknesses ranged between 30-1000 A. Details of the experimental setup and of the growth conditions are reported elsewhere [14,15]. A 10 keV RHEED system was used to study growth orientation, crystal structure and the surface morphology of the films. For transmission electron microscopy investigations a cross section was made from the film/substrate sample by cutting it with a diamond blade saw, mechanically grinding and polishing the cut-off piece with a diamond paste down to a thickness of about 10 tzm. Finally, the cross section was ion-milled without rotation until electron transparency (less than 1000 A) was reached. The parameters used for ionmilling were first voltage 5 kV, gun current 0.5 mA and an angle of 15°, and finally 3 kV, 0.3 mA and 10°. Details of the preparation method can be found in Ref. [16]. HREM was performed with a Philips CM30ST electron microscope equipped with a (Schottky) field emission gun and a side entry goniometer with a tilt range x / y of 250/25 °. The microscope was operated at 300 kV giving a resolution of 2.0 A. The images were normally recorded close to Sherzer focus. XRD data were recorded utilizing an X-ray diffractometer in the Bragg-Brentano geometry using Cu Kct radiation (A = 1.5406 A) and equipped with a graphite monochromator. The XRD data were used to determine the c-axis lattice parameters, the epitaxial film quality, and to study the phase purity. For the quantitative analysis of the 2212 films, a one-dimensional kinematic X-ray diffraction model has been used. The interplanar distances in the unit cell, site occupancies and the number of stacking faults have been used as fitting parameters. The Debye-Waller coefficient, the Lorentz-polarization and absorptions factors are included in the fitting procedure. A detailed mathematical formalism has been published elsewhere [17].
3. Results and discussion In situ heated MgO substrates show RHEED surface patterns with sharp diffraction spots lying on the 0th Laue circle. Such a pattern is obtained because of
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modulations in the films. The domains have a common c-axis, but their a- and b-axes are rotated by 90 ° with respect to each other. The observed superstructure originates from an incommensurate modulation of the rock-salt type (BiO) 2 layers with respect to the S r O - C u O z - C a - C u O z - S r O perovskite slab. XRD analysis of the films shows a c-axis oriented film and the presence of a small amount of CaCuO impurities ( d = 2 . 3 9 3 A). The full-width-at-halfmaximum (FWHM) of the (0010) rocking curves is close to 0.3 °. Zero resistivity temperatures, Tc(R= 0), between 70 and 75 K were recorded for the as-grown 2212films. Critical current density values, Jc, between 1 × l0 s and 6 × 105 A / c m 2 were determined from
Fig. 1. RHEED surface patterns of as-grown 2212-film on MgO (100). Streak splinings (A-shape) (a) due to 90 ° oriented twist domains and the superstructure, and (b) due to incommensurate modulations, are visible. The incident beam is in the [1 I0] (a) and [100] (b) directions.
an atomically flat and periodic MgO (100) surface. In grazing-exit geometry, the diffraction spots become streaks due to surface disorder caused, for instance, by two-dimensional (2D) island film growth. It has already been shown that an increased oxygen activity at higher NO 2 pressures, for a given growth temperature, leads to a 2212-film surface almost free from second-phase precipitates [14]. Clear 2D RHEED surface patterns are usually obtained for as-grown 2212 films. Fig. 1 shows the in situ recorded RHEED pattern of 2212-film (a) along [110] and (b) [100] MgO substrate azimuths. Pronounced splitting of the reciprocal lattice rods and a satellite superstructure surrounding them are clearly visible. This suggest a superposition of 90 ° oriented (twist) domains and the presence of incommensurate
Fig. 2. Experimental HREM image of a 300 A thick MBE-grown 2212-f'flm on a (100)-oriented MgO substrate. Incommensurate modulations with a wave vector at 55 ° angle are clearly visible. The 2201 and 2223 intergrowth defects (stacking faults), double Bi-layers and the incommensurate superlattice are indicated.
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the transport measurements at 10 K and were consistent with arc values obtained from SQUID magnetic measurements [14]. HREM and selected area electron diffracfign images show that 2212-films are grown preferentially with their c-axes perpendicular to the MgO {100} surface, and that both the a- and b-axes are observed when the cross section is viewed in the direction parallel to [010] MgO. Fig. 2 shows an experimental HREM image of a 300 A thick 2212-film on a (100)-oriented MgO substrate. The modulations along the b-axis leading to a wavy pattern are due to displacements of the Bi atoms in adjacent BiO layers in the film [ 18,19]. The wave vectors of these modulations make an angle of 55 ° with the {001} layer planes and are preserved up to the film surface. The structure modulations are very strong, and the local structure deviates significantly from the average structure. It seems likely that the lattice misfit between the perovskite slabs and the (BiO) 2 layers is a driving force for the modulations [20]. The results on the twinning and the modulation are consistent with HREM data of 2212 bulk samples observed by several researchers [21-24]. HREM investigation of the interface shows large areas where the 2212-film grows epitaxially on MgO
and other regions where Ca-Cu-oxide grains precipitated at the film/substrate interface. A HREM image of the f i l m / M g O interface region is shown in Fig. 3. No impurity phases were observed in the bulk of the film as have often been observed in co-evaporated YBaCuO superconductors [25]. It is reasonable to expect that nucleation and growth of precipitates as the 2212 will be affected by growth temperature, oxidation, local surface composition and defects on the substrate surface. The details of the atomic structure at the 2212/MGO interface, however, depend on a number of factors such as bonding and chemistry. Several possible mechanisms can be proposed to explain the interface precipitates that appear in MBE co-evaporated growth. There are a large number of different phases possible in the BiSrCaCuO phase diagram [26]. If thermodynamics determine the reactions on the substrate surface, uncontrolled nucleation of stable impurity precipitates may occur. Such effects can be enhanced by different sticking coefficients of deposited species on a bare MgO surface [27] and on Ca impurities segregated to the MgO surfaces [28], and also by local variations in temperature across the substrate and instabilities in the molecular beam fluxes due to the suppressed sublimation of alkaline-earth from effu-
Fig. 3. ExperimentalHREM image of the 2212-film/MgO interface. The Ca-Cu-oxide grain (P) precipitated at the film/substrate and the different n intergrowth defects are indicated.
A. Brazdeikis et al. /Physica C 253 (1995) 383-390
sion cells [29]. Another important factor which can cause precipitation of other (unwanted) phases and is related to the growth kinetics, is the substrate surface defects. There will always be terraces, mono a n d / o r multiple atomic steps on the MgO surface acting as preferred binding sites. The steps are line defects and thus the kinetics of the nucleation will be inferred from the distribution of nuclei. From several lowmagnification HREM images, it was confirmed that the C a - C u - O precipitates appear only occasionally but as rather large grains. This suggests that the impurities possess relatively high growth speeds and, probably, act as a " s i n k " for Cu and Ca during 2212-film MBE co-evaporated growth. HREM investigations show that the 2212-film growth starts with a single B i t layer on the MgO substrate. Similar growth results have recently been reported [8,30]. The incommensurate modulations are found to be present near the f i l m / M g O interface. In contrast, HREM observations of the 2212-film grown on SrTiO 3 [13] show that no modulations are present at the interface. This may be related to the fact that the 2212 nucleates on MgO as a stable and charge-neutral; B i O - S r O - C u O 2 - C a - C u O z - S r O - B i O block, whereas the 2212 at the initial growth stages on SrTiO 3 exhibits non-unit-cell growth [12].
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The 2212-film is found to contain several stacking faults (SF). These are intergrowth defects in which some layers of Cu2CaO 4 in the perovskite slab are replaced by layers of CuaCa206 a n d / o r C u t 2. As a result, 2223 and 2201 stacking faults occur. The 2201 and 2223 intergrowth defects in the 2212-film matrix occur because of the small difference in free energy between the 2212 phase and these phases [31]. Thus, any local variations in the 2212 stoichiometry will result in a stacking fault, which in fact is found to be a common type of defect in these layered compounds. The presence, however, of both 2223 and 2201 intergrowth defects in a thin 2212 film is not trivial. The "defect" 2 2 0 1 / 2 2 1 2 / 2 2 2 3 / 2212 structure (schematically shown in Fi~. 4) has an average c-axis lattice parameter of 30.8 A and the same stoichiometry as pure 2212. In this case, the local disorder cannot easily be distinguished by normal XRD. If only 2201 or 2223 SF's are present as intergrowth defects in 2212-film, a shift of some (00l) XRD peaks in a 0 - 2 0 scan will occur [32]. Equal amounts of 2223 and 2201 stacking faults do not result in a shift of XRD peaks, but alter their intensities and line-widths. In order to quantitatively examine the structure of such films, more advanced XRD methods have to be applied. Here a general
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388
one-dimensional kinematic X-ray diffraction model is applied to fit the XRD spectra from the MBEgrown 2212-films [17]. The type and amount of SF as well as the unit cell disorder are determined by an iterative fitting of the calculated profiles to the experimental XRD spectra. The measured XRD profiles show Ganssian-type peak shapes. A Gaussian distribution function was used for the continuous interplanar distance fluctuations in the unit cell. The 2201 a n d / o r 2223 phase stacking faults were assumed to be randomly distributed within the films. Fig. 5 displays a calculated shift (in 2 0 deg) of the (008) and (0012) Bragg peaks as a function of the 2201/2223 proportion. The total number of stacking faults is assumed to be 20% of the 2212-film volume. The (008) and (0012) peaks are shifted only if the number of 2201 defects is not equal to the number of 2223 defects. The shifts are towards the (0010) peak when the amount of 2223 SF is higher than that of 2201. The (008) and (0012) peaks are moved apart when the amount of 2223 SF in the film is lower than that of 2201. As can be seen in the calculations in Fig. 5, there is no peak shift when the number of 2201 SF is equal to the number of 2223 SF. Such an interpretation by observing the shift of XRD peaks in 0 - 2 0 scans can only be used to
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determine which type of stacking fault, 2201 or 2223, dominates in the film. A complete characterization of the SF defects in the 2212 layer requires a determination of the line-widths of the (008), (0010) and (0012) Bragg peaks and their intensities. Fig. 6 gives the calculated (in 2 0 deg) FWHM of (008), (0010) and (0012) peaks of a 370 ,~ thick 2212-film when the SF 2201/2223 ratio is equal to unity. The FWHM of the (008) and (0012) peaks increases with the amount of SF, but does not change significantly for the (0010) peak. The sensitivity of the FWHM of the (008) and (0012) peaks to the presence of small amounts of SF's in the 2212 compound can easily be understood by comparing 0-2 0 spectra of the 2201, 2212 and 2223 phases. The positions of the Bragg peak ((008), (0010) and (0012) for the 2201, 2212 and 2223, respectively) at 20 = 29 ° remain mostly unchanged from one phase to the other. The superimposition of the surrounding peaks certainly broadens, e.g., the (008) and (0012) peaks in 2212 if small amounts of 2201 a n d / o r 2223 are present. The intensity distribution analysis in the 0 - 2 0 XRD profile is not straightforward. The intensities are not only affected by the presence of stacking
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Table 1 Calculated average c-axis length and the number of 2201 and 2223 stacking faults for three different MBE-grown2212-films Sample Average c-axis Number of stacking faults length (,~) (%) 2201 2223 #154 30.84 44.5 1.2 #159 30.79 2.1 34.7 #166 30.84 3.5 4.6 faults in the film. Structural disorder in the unit cell due to large substitutional effects may lead to intensity changes of some XRD peaks [17]. The fitting revealed that the continuous interploanar distance fluctuations in the 2212-films are 0.04 A. The substitution at Sr 2+ sites by Ca 2+ and by Bi 3+ is approximately 20% and 15%, respectively. The Ca 2+ site is populated by Sr 2+ to 40%. These substitutions increase the interplanar CuO2-Ca distance from 1.65 to 1.70 A, and thus the S r O - C u O 2 distance in the unit cell is contracted from 1.72 to 1.58 A. This contributes to the lower 0012 XRD peak intensity than expected for a perfect 2212-film. The changes in the interplanar distances may also explain why the line-width of this peak in the measured 2212 XRD spectrum is usually observed to be larger than the value calculated from Fig. 6. Experimental XRD spectra of three different MBE-grown 2212 films and the fitted XRD profiles are presented in Fig. 7. The average c-axis length and the number of stacking faults obtained by the iterative fitting are presented in Table l. A large number of the stacking faults are obtained during co-evaporated growth when large variations from the 2212 stoichiometry occur, mainly due to inaccurate calibration of molecular beams, while the presence of only a small number of 2201 and 2223 stacking faults indicates that the stoichiometric 2212 MBE growth is achieved. Any variation in temperature across the substrate and any changes in the molecular beam fluxes, as already discussed, will certainly promote nucleation of local intergrowth defects. Nevertheless, the fact that the 2201 and 2223 stacking faults are present in almost equal amounts indicates that the 2212 growth here is probably accompanied by an intrinsic generation of stacking faults, which is supported by the fact that the 2212 growth conditions are far from the thermodynamic equilibrium and by the energetic proximity o
389
of similar polytype phases. Once the stacking fault, e.g. 2201, is locally generated, the excess Ca and Cu changes the chemical potential in that region. Such a local "off-stoichiometry" condition will, with rather high probability, result in the formation of a 2223 stacking fault. A similar sequence of events can be observed when the 2223 SF is formed first, resulting in the local Ca and Cu deficiency and, thus, followed by the 2201 SF formation.
4. Conclusions MBE growth and the resulting microstructure of films have been studied by RHEED, HREM and XRD. In situ RHEED analysis showed that as-grown films are formed by two-dimensional island growth with 90 ° oriented twist domains and exhibit incommensurate modulations. The observed features were confirmed by HREM. The film-substrate interface studies showed large areas where the film grew epitaxially on the MgO (100) surface, and several regions where C a - C u - O impurity precipitated at the substrate-film interface. The nucleation of precipitates is explained as being governed by the different sticking coefficients of deposited species on MgO surfaces and also by Ca impurities segregated on the MgO surfaces. Such effects could be enhanced by a local variation of the temperature across the substrate and time-dependent changes of molecular beam fluxes. For the first time, both 2201 and 2223 stacking-fault defects were observed to be present in Bi2Sr2CaCu20 x thin films. These local structural defects were examined by a general one-dimensional XRD model. A characteristic feature of the defects was their occurrence in almost eqtial amounts. The appearance of the defects in the 2212-films was explained on the basis of the 2212-film growth at thermodynamic non-equilibrium and the energetic proximity of similar polytype phases. Bi 2S r 2 C a C u 2 0 x
Acknowledgements This work was supported by the Swedish Research Council for Engineering Sciences (TFR) and Swedish National Board for Industrial and Technical Development (NUTEK).
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