DOI: 10.1021/cg100221w
Induced Epitaxy for Growth of Aligned Indium Nitride Nano- and Microrods
2010, Vol. 10 3887–3891
B. S. Simpkins,*,† Aman D. Kansal,‡ and P. E. Pehrsson† †
Chemistry Division, Naval Research Laboratory, Washington, DC, and ‡Thomas Jefferson High School for Science and Technology, Alexandria, Virginia Received February 12, 2010; Revised Manuscript Received June 1, 2010
ABSTRACT: Wurtzite indium nitride nano- and microscale crystals were grown on c-plane sapphire substrates in an oxide-, template-, and catalyst-free chemical vapor deposition reactor using the halogen-based precursor InCl3. The impact of growth temperature, In flux, and ammonia flow on the InN microstructure and epitaxial alignment was evaluated. Increased growth temperature resulted in nanorods with high-aspect ratio and enhanced epitaxial alignment of InN(001)//Al2O3(001). Faceting exhibited by our growth product confirmed that simple bond-counting does not determine the exhibited planes and phenomena such as dipole-dipole attractions and surface diffusion determined crystal morphology under certain conditions. An increase in growth product with increased In flux indicated our growths are carried out under heavily N-rich conditions with V/III ∼ 104. These results provide a stronger understanding of the growth mechanisms of this unique semiconductor.
Introduction Group III-nitride nanostructures are of great scientific interest due to their potential impact in fields ranging from biomedicine to optical communications. Indium nitride (InN) exhibits unique and particularly useful optical and electronic properties. Although the reported optical gap of InN has been as high as 1.89 eV,1,2 more recent studies indicate that this transition likely corresponds to an indium oxide and the true InN gap is 0.65-0.90 eV.3-5 In fact, spontaneous infrared emission in this energy range has been generated through impact ionization in InN.6 Accordingly, InN may be alloyed with gallium nitride (GaN) to tune optical emission or absorption over the entire visible spectrum and into the infrared, including the communications band at ∼1.5 μm. The electronic properties of InN include a drift velocity predicted to be larger than those of both GaN and GaAs7 and the formation of a surface electron accumulation region due to downward bandbending.8 These properties ease the fabrication of low-resistivity ohmic electrical contacts and may make the conductivity of InN nanostructures particularly sensitive to surface charge modulation by biological or chemical surface molecules. In order to produce InN nanostructures with technological impact, control of crystal size, morphology, and crystallography must be improved. InN nanostructures have been produced using metal catalyst particles9,10 or oxide precursors6,11 despite the potential to form the undesirable wide band-gap oxide and deep traps through incorporation of the metal catalyst material. Naturally, it would be preferable to produce InN nanostructures without oxide precursors or metal catalysts. Guidance toward this goal is gained by reviewing results of InN nanorod growth utilizing Cl-based MOCVD reactions in which an optimal In/Cl ratio of ∼1/3 was computationally identified12 and experimentally verified.13 These findings suggest that solid InCl3 is an effective precursor with several advantages over conventional metal-organic precursors. Although InCl3 is moisture-sensitive, it (i) requires less specialized equipment, (ii)
provides an effective In/Cl ratio for nanorod formation, (iii) bypasses the multiple decomposition reactions necessary with metal-organic precursors, and (iv) eliminates the need for high temperatures to drive the kinetically limited formation of InCl3.12 The current study describes the catalyst- and oxidefree growth of InN nano- and microstructures utilizing InCl3 as the group III precursor. Solid InCl3 has been previously used to grow InGaN alloy nanostructures;14 however, this study focused on the characterization of a single sample with spatially inhomogeneous stoichiometry and microstructure. In contrast, we identify the influences of several important growth parameters on material microstructure and epitaxial alignment and use the knowledge gained from these results to generate a desired InN rod morphology. Materials and Methods
*To whom correspondence should be addressed. E-mail: blake.simpkins@ nrl.navy.mil.
The InN structures were grown on c-plane Al2O3 substrates, which were chosen because they have a hexagonal crystal structure that resembles InN,15 although the lattice mismatch between a-Al2O3 and a-InN is large at ∼25%.2 Substrates were sonicated in acetone and isopropyl alcohol, each for 10 min. They were then exposed to a piranha solution (3:1 H2SO4/ H2O2) for 20 min to remove carbonaceous contamination, rinsed in distilled water, and dried under N2. Our chemical vapor deposition (CVD) process is a variation of the hydride metal-organic vapor phase epitaxy (H-MOVPE) technique used for growths of GaN and InN13,16,17 and was carried out in a 3-zone furnace capable of maintaining three independently temperature-controlled regions (Figure 1). Approximately 100 g of InCl3 powder in a precursor boat was loaded into zone 1 through a dry-bag. Cleaned deposition substrates were loaded into zone 3. During the temperature ramps and the 20 min growth, N2 gas was flowed through the outer tube and the InCl3-containing tube to deliver the In precursor to the growth substrate. NH3 was flowed through the outer tube once the growth temperature was reached. The NH3 was dissociated at the higher temperature of zone 2 and encountered the evaporated In precursor at the growth substrate in zone 3. In this
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Published on Web 07/28/2010
pubs.acs.org/crystal
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Figure 1. Schematic of a 3-zone growth furnace. Zones 1-3 serve to (i) evaporate In precursor, (ii) dissociate NH3, and (iii) combine reactants to grow InN, respectively.
study, The In flux, In/N flux ratio, and substrate temperature were varied by modifying the temperature of zone 1 (TIn), the NH3 flow rate (fNH3), and the growth temperature in zone 3 (Tgr), respectively. Our primary study centered on nominal growth conditions of TIn = 280 C, fNH3 = 100 sccm, and Tgr = 440 C. Each of these parameters was varied individually while keeping the others fixed at these nominal values. Each parameter was varied over 5 values as follows: TIn from 240 to 320 C, fNH3 from 60 to 140 sccm, and growth temperature Tgr from 400 to 500 C. Finally, the conclusions drawn from this growth series were used to optimize the growth conditions used in actual growths. Scanning electron microscopy (SEM), carried out on a LEO field emission instrument, was used to evaluate the morphology of the growth product. X-ray diffraction (XRD) was carried out in a Bruker AXS D8 Advance unit utilizing Cu KR X-rays (λ = 1.54 A˚) to identify the phase of the material and determine the preferred orientation of InN crystals relative to the c-Al2O3 substrate. Results Typical growth morphologies are shown in Figure 2. All of the crystallites exhibited hexagonal faceting and often a faceted pointed tip and were generally anisotropic in dimension, with aspect ratios from ∼1 to 10. A subset of the rods grew normal to the single-crystal substrate. The rod radius ranged from ∼200 to 700 nm. Crystals can be found as single rods (Figure 2a and b) or in 2-, 3-, or 4-fold multibranched structures (Figure 2c). In general, the observed morphologies can be categorized as (i) single rods oriented perpendicular to or (ii) at an inclined angle relative to the substrate and (iii) branched structures. Pointed InN rods with hexagonal cross sections have previously been grown using a hydride-based technique13,16 and identified as growing along the c-direction, with six m-type (100) facets comprising the rod walls. Justification for attributing this assignment to our material will be given in the Discussion section. The morphology of the growth product was controlled through variation of the growth conditions as summarized by the matrix of SEM images in Figure 3. Each row in Figure 3 highlights the variation of a single growth parameter. Images a-c show the morphological evolution associated with varying the growth temperature (Tgr) from 400 to 500 C. The material produced at the lowest growth temperature (Figure 3a) was characterized by increased nucleation and growth on the rod sidewalls, leading to polydispersion in size and orientation with respect to the substrate. Enhancement of the aspect ratio was seen at higher temperature, resulting in individual high-aspect ratio structures with a reduced amount of branched growth. Individual rods aligned normal to the growth substrate appear as simple hexagons in the plan view image of Figure 3c. Material
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Figure 2. Typical InN morphologies. Material (a) is hexagonally faceted, (b) grows normal to or at an angle to the substrates, and (c) may form a branched structure consisting of several crystallites. Scale bars are all 2 μm.
grown at all temperatures retained clear 6-fold faceting. An immediate and significant conclusion from these data is that InN nanorods could be grown at temperatures below that identified as the lower limit imposed by kinetically driven formation of InCl3 in hydride vapor-phase deposition processes (∼525 C).12 Reference 12 has identified the formation of InCl3 as influential in the formation of highly anisotropic InN structures. The formation of InCl3 within an HVPE environment is kinetically limited and therefore requires elevated temperatures. Our system bypasses these steps by starting with the InCl3 species rather than relying on its formation within the reactive environment. The presence of InCl3 molecules may not be a strict requirement for the growth of anisotropic crystallites; however, it has been identified as participating in gas-phase InN formation, which leads to many small nucleating crystals and anisotropic growth rather than uniform surface nucleation followed by thin film growth. The effect of In flux is shown in Figure 3d-f. Indium flux at the growth substrate was controlled through the temperature of the InCl3 precursor (TIn). The most significant trend observed was an increase in surface coverage with increasing In flux. Material morphology appeared essentially unchanged until a continuous film coalesced at high In flux (Figure 3f). Crystallites formed at the highest indium flux may be somewhat smaller due to constraints imposed through impingement with neighboring crystals, but this cannot be quantified from the data. This increase in surface coverage as a function of In flux confirms that In was the limiting resource in our growth process. This is expected since complete dissociation of all introduced NH3 would produce ∼10-1 mol of N while the mass loss of the InCl3 precursor corresponded to ∼10-5 mol of In, yielding V/III ∼ 104. Figure 3g-i further demonstrates that our system operated in the N-rich regime by showing very little impact on nucleation density due to varying the already excessive reservoir of the group V resource. All samples were examined by XRD and identified as wurtzite InN. Representative XRD data is shown in Figure 4a. Dominant peaks indexed to wurtzite InN (002), (101), and Al2O3 (006) were observed in all samples and located at 31.2, 33.05, and 41.66, respectively. Although the relative intensities of the peaks varied between samples, the InN (002) and (101) and Al2O3 (006) peaks exhibited the strongest intensity. The average fwhm of the InN (002) peak was ∼700 arcsec, indicating some defect broadening. Smaller peaks (∼100 times weaker in intensity than the InN (002) peak) were observed and are indexed as K-β reflections for the following: /, InN (002); †, Al2O3 (006); and ‡, InN (004). The dominance of the InN (002) and (101) peaks indicated that there was clear texturing of the InN crystallites. InN rods exhibited a preference for orienting either their (002) or their (101) planes parallel to the substrate surface. The arrangements of the
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Figure 3. Matrix of micrographs showing evolution of morphology due to varying growth conditions. Nominal growth conditions are Tgr = 440 C, TIn= 280 C, and fNH3 = 100 sccm. Each row shows the effect of varying a single variable while keeping the others at the nominal values. (a-c) Increasing growth temperature reduces sidewall nucleation and increases c-axis alignment normal to the substrate. (d-f) Increasing In flux simply increases surface coverage. (g-i) Increasing NH3 flow has little effect due to a highly group V rich environment. The scale bar applies to all images.
Figure 5. (a) Evolution of preferred crystallite orientation showing increasing population of rods aligned normal to the substrate at increased temperature. (b and c) Schematic of possible facet intersections described in the text. Figure 4. (a) Representative XRD data identifying grown material as wurtzite InN and showing the predominance of (002) and (101) crystallite orientations illustrated in parts b and c, respectively. Small peaks are identified in the text.
hexagonal InN relative to the substrate for these two preferred orientations are illustrated in Figure 4b and c. Details regarding the evolution of this preferred alignment with growth conditions are discussed below. Discussion We have demonstrated growth of wurtzite InN and observed two preferred crystalline orientations with respect to the substrate. It is desirable to understand what factors influence crystallite arrangement during nucleation and growth and use this information to target the desired morphologies. Aligning
the nanorod axis normal to the growth substrate has a number of potential benefits, including polarization-dependent ensemble optical behavior, field emission applications, and a known epitaxial relationship between crystallite and substrate. To this end, preferential crystallite orientation was tracked by calculating the ratio of the InN (002) and (101) integrated XRD peak areas. This value essentially represents the fraction of crystallites oriented with their (002) planes parallel to the surface of the substrate18 or, in other words, the InN c-direction normal to the c-Al2O3 substrate, as illustrated in Figure 4b. These peak ratios, plotted as a function of growth temperature in Figure 5a, varied by nearly a factor of 20 and increased with increasing growth temperature. This is likely due to either a thermal destabilization of (101)-oriented crystals or thermally activated diffusion, enabling nucleating crystals to align with their c-axis normal to
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the substrate. The presence of (002)-oriented InN rods on the cAl2O3 substrate (Figure 4b) is not surprising, since both present hexagonally arranged atoms. However, the atoms on the InN (101) facet (Figure 4c) exhibit rectangular symmetry. Although the conventionally defined lattice mismatch is difficult to compute for this situation, the lack of atomic registry requires significant strain or interfacial defect structure and may cause destabilization at higher temperatures. The trend of preferred alignment of the InN c-axis normal to the substrate was accompanied by a clear morphological trend of increased nanorod alignment normal to the substrate, as seen in Figure 3a-c and Figure 2b. The correlation between the increased relative intensity of the (002) XRD peak and increased rod alignment normal to the substrate provides strong evidence that these rods grew along the c-direction. This assignment is consistent with InN rods grown by hydride vapor phase epitaxy.13,16 The fact that the InN Æ001æ direction lies along the rod axis suggests a possible crystallographic underpinning for the branched structures observed in Figure 3. In the present system, one may immediately suspect stacking faults occurring during nucleation. The cubic zinc-blende (ZB) and hexagonal wurtzite structures of InN differ only by a low-energy stacking fault of the (111) plane in the ZB system (this would correspond to the (001) plane in the hexagonal system). If a small region of ZB material formed during nucleation, an obtuse angle of 109.5 would exist between the (111) type planes in this system. If branches of hexagonal material nucleated at these (111) type planes and commenced growth along the Æ001æ hexagonal direction (this would be commensurate with the Æ111æ-type cubic direction and is indeed the growth direction observed for our hexagonal InN rods), these rods would form angles of 109.5 with respect to one another. This is the approximate angle we find between branches in our structures. It seems quite plausible that a combination of stacking faults and regions of cubic phase material was responsible for the observed branching. We note that no XRD peaks indexed to ZB InN were observed, indicating that these ZB regions, if they exist, either constitute a small fraction of the grown material or are not parallel to the growth surface. Verification of this speculated mechanism would require detailed TEM analysis and is beyond the scope of this manuscript. Further examination of the tip facets revealed that these planes formed a 60.9 ( 2.9 angle relative to the rod axis. Two types of planes are consistent with this value. The (112) and (101) planes would form angles of 58.4 and 62, respectively. One structural characteristic that may be used to infer crystallography of these facets is the line of intersection created between the tip and sidewall facets (Figure 5b and c). In our material, this line has a component along the rod axis. This situation is shown in the experimental results of Figures 2c and 3h and i and schematically in Figure 5b. This relative arrangement was also observed in the plan view image of Figure 2a, where the ridges formed by the tip facets were rotated 30 relative to ridges formed by the sidewall facets. This relative arrangement between the tip and sidewall facets allows for two possible planar assignments; tip and sidewall facets are, respectively, (i) (101) and (110) or (ii) (112) and (100). Although the current data cannot distinguish between these possibilities, dangling bond counts19 and first principles calculations for the analogous GaN system20 predict the (100) to be energetically favorable, supporting option (ii). Thermodynamic and kinetic factors combined to result in the trends revealed in Figure 3a-c and Figure 5a. Considerations such as dangling bond counts and dipole-dipole attractions
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Figure 6. (a) Micrograph of near optimal growth with a high degree of rod alignment (“A”) and a very small fraction of crystallites oriented at an inclined angle (“B”). (b) Replot of relative alignment data with optimized growth included showing enhanced c-axis alignment normal to the substrate at elevated growth temperature. The vertical dotted line indicates the temperature limit imposed by etching.
between vapor-phase III/nitride molecules and the polar InN crystal12 define the thermodynamically preferred faceting. Simple dangling bond counting predicts the (001) basal plane as the most stable surface, which should lead to basal plane platelets instead of structures with extended (100)- or (110)planes like those exhibited in this and other work.14,16 Dipoledipole interactions have been cited as the primary driving force for the anisotropic growth of InN along the c-direction,13 and in fact, fast-growing c-planes have been observed in planar GaN19,21 despite their low energy, as predicted by bondcounting. However, sufficient surface diffusion of adatoms or molecules is necessary to produce the thermodynamically preferred morphology. Our results indicated that lower temperatures imposed kinetic limitations which slowed adatom diffusion. Slow diffusion at low growth temperature caused adatoms to persist on the rod sidewalls and eventually nucleate new crystals, resulting in polycrystalline and branched structures. Additionally, adatoms and crystallites forming on the bare substrate are less likely to diffuse into the preferred epitaxial alignment and may instead nucleate at some nonoptimal orientation with respect to the substrate. These effects were most obvious in Figure 3a, where lower diffusion rates led to nucleation and growth of crystals on the low-energy sidewalls, resulting in a randomly oriented polycrystalline morphology with low-aspect-ratio branched structures. On the other hand, higher diffusion rates increased the probability of atoms or molecules finding the bonding sites available on the thermodynamically determined fast-growing surfaces. Thus, a higher growth temperature yielded more rods oriented normal to the substrate, larger aspect ratios, and a decrease in branching, as demonstrated in Figure 3a-c. Finally, the understanding of growth behavior developed here was used to target a desired InN morphology. Accordingly, two final growths were carried out at elevated growth tempera-
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tures to enhance InN rod alignment normal to the growth substrate and moderate In precursor flux so as to obtain wellseparated rods. The parameters used for these growths were as follows: TIn = 280 C, fNH3 = 100 sccm, and Tgr = 560 and 600 C. The resulting sample grown at 560 C is pictured in Figure 6a. This sample exhibited a high fraction of crystals oriented normal to the substrate (labeled “A”), with a small amount of material comprised of extremely short rods protruding from the substrate at an inclined angle (labeled “B”). The reduced length of these shorter rods suggested that growth either halted at some point or commenced much more slowly. The aligned rods had a slight inverse taper. We speculate that, early in the growth process, reactant availability is reduced due to the growth of both the aligned and inclined rods. Later, growth of inclined crystals stops and all reactants are incorporated into the aligned rods. This variation of available reactants could be responsible for the slight diameter modulation observed. Branched structures, such as that shown in Figure 6a, were present but occurred at a much lower frequency under these growth conditions. A high degree of epitaxial alignment is demonstrated in the XRD peak ratios plotted in Figure 6b. This plot includes the previous results of Figure 5a in order to highlight the improved alignment. An XRD peak associated with the (101) alignment was still visible and is likely associated with the shorter crystallites (“B”). There was virtually no growth at the highest temperature of 600 C, presumably due to increased etching of the InN, as predicted by ref 12. This etching sets an upper temperature limit to InN growth in this environment and is indicated with a vertical dotted line in Figure 6b. Conclusion InN nano- and microscale rods were successfully produced in a CVD reactor utilizing InCl3 as the group III source. This precursor bypasses several key reaction steps necessary when using metal-organic precursors and provides an effective In/ Cl ratio for nanorod formation. The growth product consisted of hexagonally faceted rods growing along the c-direction and exhibited two preferred alignments relative to the c-Al2O3 substrate. These orientations were (001) and (101) planes parallel to the substrate. The evolution of preferred crystal alignment with growth temperature was tracked with XRD peak analysis and showed that increased growth temperature resulted in preferred epitaxial alignment of InN(001)//Al2O3(001). This change in preferred orientation was accompanied by a change in morphology from randomly oriented low aspect ratio branched crystals to highly anisotropic single rods growing normal to the substrate. This behavior was, in part, a consequence of temperature-dependent surface diffusion. Slow diffusion at low temperatures reduced the role of dipole-dipole driven growth along the polar c-direction and increased that of nucleation and growth on low-energy sidewalls.
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Increased growth temperatures reduced sidewall nucleation and either increased epitaxial rod alignment normal to the growth substrate or destabilized misoriented crystallites. By understanding the growth parameters which most influence the crystal characteristics of interest, we achieved the desired growth of strongly aligned and well-separated InN rods. Acknowledgment. This work was funded by the Office of Naval Research. A.K., a student at Thomas Jefferson High School for Science and Technology in McClean Virginia, acknowledges support from the NRL Science and Engineering Program and thanks Dr. A. Epshteyn for his help with the XRD technique.
References (1) Tansley, T. L.; Foley, C. P. J. Appl. Phys. 1986, 59, 3241. (2) Bhuiyan, A. G.; Hashimoto, A.; Yamamoto, A. J. Appl. Phys. 2003, 94, 2779. (3) Davydov, V. Y.; Klochikhin, A. A.; Emtsev, V. V.; Kurdyukov, D. A.; Ivanov, S. V.; Vekshin, V. A.; Bechstedt, F.; Furthmuller, J.; Aderhold, J.; Graul, J.; Mudryi, A. V.; Harima, H.; Hashimoto, A.; Yamamoto, A.; Haller, E. E. Phys. Status Solidi B 2002, 234, 787. (4) Wu, J.; Walukiewicz, W.; Yu, K. M.; Ager, J. W., III; Haller, e. E.; Lu, H.; Schaff, W. J.; Saito, Y.; Nanishi, Y. Appl. Phys. Lett. 2002, 80, 3967. (5) Matsuoka, T.; Okamoto, H.; Nakao, M.; Harima, H.; Kurimoto, E. Appl. Phys. Lett. 2002, 81, 1246. (6) Chen, J.; Chen, G.; Stern, E.; Reed, M. A.; Avouris, P. Nano Lett. 2007, 7, 2276. (7) O’Leary, S. K.; Foutz, B. E.; Shur, M. S.; Eastman, L. F. Appl. Phys. Lett. 2005, 87, 222103. (8) Zhang, R.; Zhang, P.; Kang, T.; Fan, H.; Liu, X.; Yang, S.; Wei, H.; Zhu, Q.; Wang, Z. Appl. Phys. Lett. 2007, 91, 162104. (9) Liang, C. H.; Chen, L. C.; Hwang, J. S.; Chen, K. H.; Hung, Y. T.; Chen, Y. F. Appl. Phys. Lett. 2002, 81, 22. (10) Cai, Z.; Garzon, S.; Chandrashekhar, M. V. S.; Webb, R. A.; Koley, G. J. Electron. Mater. 2008, 37, 585. (11) Tang, T.; Han, S.; Jin, W.; Liu, X.; Li, C.; Zhang, D.; Zhoa, C.; Chen, B.; Hand, J.; Meyyapan, M. J. Mater. Res. 2004, 19, 423. (12) Won, Y. S.; Kim, Y. S.; Kryliouk, O.; Anderson, T. J. J. Cryst. Growth 2008, 310, 3735. (13) Kryliouk, O.; Park, H. J.; Won, Y. S.; Anderson, T. J.; Davydov, A.; Levin, I.; Kim, J. H.; Freitas, J. A., Jr. Nanotechnology 2007, 18, 135606. (14) Kuykendall, T.; Ulrich, P.; Aloni, S.; Yang, P. Nature Lett. 2007, 6, 951. (15) Mamutin, V. V.; Shubina, T. V.; Vekshin, V. A.; Ratnikov, V. V.; Toropov, A. A.; Ivanov, S. V.; Karlsteen, M.; S€ odervall, U.; Willander, M. Appl. Surf. Sci. 2000, 166, 89. (16) Shalish, I.; Seryogin, G.; Yi, W.; Bao, J. M.; Zimmler, M. A.; Likovich, E.; Bell, D. C.; Capasso, F.; Narayanamurti, V. Nanoscale Res. Lett. 2009, 4, 532. (17) Park, C.; Yeo, S.; Kim, J.; Yoon, D.; Anderson, T. J. Thin Solid Films 2005, 498, 94. (18) Birkholz, M.; Selle, B.; Conrad, E.; Lips, K.; Fuhs, W. J. Appl. Phys. 2000, 88, 4377. (19) Jindal, V.; Shahendipour-Sandvik, F. J. Appl. Phys. 2009, 106, 083115. (20) Northrup, J. E.; Neugebauer, J. Phys. Rev. B 1996, 53, 10477. (21) Hiramatsu, K.; Nishiyama, K.; Motogaito, A.; Miyake, H.; Iyechika, Y.; Maeda, T. Phys. Status Solidi A 1999, 176, 535.
