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Influence of buffer layer and 6H-SiC substrate polarity on the nucleation of AlN grown by the sublimation sandwich technique
Journal of Crystal Growth 233 (2001) 177–186
Influence of buffer layer and 6H-SiC substrate polarity on the nucleation of AlN grown by the sublimation sandwich technique Y. Shia, Z.Y. Xiea, L.H. Liua, B. Liua, J.H. Edgara,*, M. Kuballb a
Department of Chemical Engineering, Durland Hall, Kansas State University, Manhattan, KS 66506-5102, USA b H.H. Wills Physics Laboratory, University of Bristol, BS8 1TL, UK Received 23 July 2000; accepted 13 June 2001 Communicated by A.F. Witt
Abstract The initial nucleation stage of AlN crystal grown by the sublimation method on on-axis (0 0 0 1)Si and (0 0 0 1)C 6HSiC as well as 3.51 off-axis (0 0 0 1)Si 6H-SiC substrates was investigated. Sublimation growth from a pure sintered AlN source was carried out in a resistively heated furnace with a source temperature of about 18001C at a nitrogen pressure of 500 Torr. Direct growth on Si-terminated as-received SiC substrates was discontinuous, marked by sparse nucleation and slow lateral growth of nuclei. Several hexagonal sub-grains were usually obtained on the substrates after a long growth time (more than 3 h) due to the incomplete coalescence of the nuclei. In contrast, no sublimation growth occurred on the as-received C-terminated substrates. To enhance two-dimensional (2D) growth, an AlN epitaxial layer was first deposited on the substrates by MOCVD before sublimation growth. Continuous films could then be grown on all the substrates with AlN MOCVD buffer layers. The tensile stress of the AlN layer due to thermal expansion coefficient mismatch between SiC and AlN caused cracking across the AlN/SiC interface into the SiC substrates during the cooling process limiting the maximum of the thickness and lateral size of the AlN crystals. r 2001 Elsevier Science B.V. All rights reserved. Keywords: A1. Crystal morphology; A1. Nucleation; A2. Growth from vapor; A2. Seed crystals; B1. Nitrides
1. Introduction The GaN-based semiconductors have developed rapidly in blue and green light emitting diodes and blue emitting lasers [1]. The primary substrates for epitaxial growth of GaN are SiC
*Corresponding author. Tel.: +1-913-532-4320; fax: +1913-532-7372. E-mail address: [email protected] (J.H. Edgar).
and sapphire, which are poorly lattice and thermal expansion coefficient matched to GaN. Aluminum nitride, a wurtzite structure wide band gap (6.2 eV) semiconductor, is a good candidate substrate for GaN epitaxial films due to its relatively small lattice constant mismatch along the a-axis ( 3.5%), good thermal stability (melting point >25001C), high resistivity and similar coefficient of thermal expansion [2]. The use of AlN buffer layers in MOCVD significantly improves the quality of GaN epitaxy [3] also
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suggesting that AlN would be an ideal substrate for GaN epitaxy. The most successful method for producing bulk AlN crystal is by sublimation growth (vapor phase transport). In this technique, AlN undergoes the reaction AlN=Al+12N2 in a high temperature source zone and this reaction is driven in the reverse direction to form AlN single crystals in a lower temperature zone. This technique was most successfully developed by Slack and McNelly [4,5] in the mide-1970s. Large single crystals of AlN up to 1 cm long and 0.4 cm in diameter were grown by self-seeding in the sharp tip of sealed tungsten crucibles in an rf induction furnace. Recently, Rojo et al. [6,7] have further advanced this technique to produce single crystals up to 15 mm in diameter. Tanaka et al. [8] investigated the axis orientation of AlN freely nucleated on graphite crucible walls. Most relevant to this work, Balkas et al. [9] produced single crystal platelets of AlN (p1 mm thick) on 10 10 mm2 6H-SiC substrates in a resistively heated graphite furnace. However, the crystals contained small (B2 2 mm2) individual hexagonal sub-grains, which was attributed to the severe degradation of the SiC substrates resulting in isolated stable nucleation sites. A prior investigation by the author’s group [10] showed freely nucleated needles and platelets and exhibited much better crystal quality than the crystals grown on 6H-SiC substrates although by the latter method the crystal orientation and initial nucleation was usually better controlled. Up to now, no systemic investigation of the initial nucleation stage of AlN sublimation on 6H-SiC substrate has been reported, although proper nucleation critically effects the quality of the subsequently grown crystals. Nor has the buffer layer technique common to MOCVD growth been employed to prepare substrates for the sublimation growth. Herein for the first time the initial nucleation stage of AlN crystal grown by sublimation method on on-axis (0 0 0 1)Si and (0 0 0 1)C 6HSiC as well as off-axis (0 0 0 1)Si 6H-SiC substrates were investigated. The influence of the AlN buffer layer deposited on different SiC substrates by MOCVD on the subsequent AlN sublimation growth were also studied.
2. Experimental procedures 2.1. Crystal growth Sublimation growths were conducted in a resistively heated furnace using tungsten wire mesh heating elements. The seed and source were enclosed in two concentric tungsten crucibles to contain the Al vapor as it deteriorates the tungsten heating elements. Sintered AlN with approximately 1% oxygen as the main impurity analyzed by a standard inert gas fusion method (TC-136, Leco Co.) was used as source and seed holders. The as-received (0 0 0 1) 6H-SiC silicon terminated (on-axis and 3.51 off-axis) and carbon terminated wafers (only on-axis) were cut into about 1 cm2 substrates. All the substrates were ultrasonically cleaned in organic solvents (TCE, acetone and methanol), rinsed with DI water, and dried with high purity nitrogen gas before loaded into the furnace. The distance between the source and the substrate was kept at 2 mm. In order to reduce SiC dissociation from the backside of the seed, it was covered by an AlN cap. The growth ambient was 99.99993% pure nitrogen. Before growth, the furnace was evacuated to a pressure less than 10 4 Torr, then purged by nitrogen twice to remove residual gases. During the heating process the furnace was maintained at 800 Torr, as this higher nitrogen pressure suppressed both AlN sublimation and the decomposition of the SiC substrate which may roughen the substrate surface causing unfavorable AlN nucleation. As soon as the temperature reached the desired value, the background pressure was decreased to 500 Torr very quickly to start the sublimation growth. The growth pressure was maintained by an automatic throttle valve. At the end of the growth, the system pressure was increased to 800 Torr again during cooling to avoid unwanted AlN recondensation. The vertical temperature profile of the growth chamber was measured by an optical pyrometer focused on a moving target at different furnace output powers and pressure. To obtain comparable results, the growth temperature was kept constant for all growths by fixing the output power. The in situ growth temperature was
Y. Shi et al. / Journal of Crystal Growth 233 (2001) 177–186
measured by the same optical pyrometer focused on the top lid of the outside crucible. According to the pre-measured temperature profile, at our experimental condition the source temperature was about 18001C and the source-substrate temperature gradient was about 101C for 2 mm source-substrate distance. Several experiments carried at the temperature above 20001C showed that higher growth temperatures had no significant influences on the nucleation and morphology except the growth rate. Therefore the low temperature was employed to avoid the degradation of the furnace and the SiC substrates. 2.2. MOCVD technique The AlN buffer layer growth was performed in a vertical-type MOCVD reactor operated at low pressure (76 Torr). Before loading into the reactor, the 6H-SiC substrates were cleaned using the same procedure described above. Trimethylaluminum (TMA) and ammonia (NH3) were the Al and N sources, and Pd-cell purified H2 was the carrier gas. The substrates were preheated at 11001C for 10 min in 3 slm of H2 flow for surface cleaning. Then the AlN films were deposited at the same temperature under the H2, NH3 and TMA flow rates of 3, 3 slm and 30 sccm, respectively. An approximate 2 mm thick AlN buffer layer for the subsequent sublimation growth was obtained for the 4 h growth time. 2.3. Characterization The as-grown AlN crystals were first characterized by Nomarski differential interference contrast microscopy (NDIC). Macrostructural features were observed at magnifications of 5–40X. X-ray diffraction measurements were conducted using XDS 2000 diffractometer (Scintag Inc.) with the Cu Ka1 radiation. The surface morphology of the samples was examined by Digital Instruments Nanoscope E AFM in contact mode. MicroRaman spectra were obtained at Ultraviolet Renishaw with the 325 nm line of a HeCd laser as excitation source. The spot size and the spectral resolution were 1–2 mm and 3–4 cm 1, respectively.
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3. Results 3.1. Long time growth result of AlN on (0 0 0 1)Si 6H-SiC Initially, different growth temperature, pressure and substrate-seed distance were tried on the Sipolarity 6H-SiC substrate to optimize the growth condition. The pressure during the heating and cooling process was first kept consistent with the growth pressure. The typical sublimation time was usually about 10 h. In all experiments, crystalline AlN was deposited on the majority of the SiC substrates. Every AlN crystal thus formed contains several hexagonal sub-grains. Grain size varied slightly under different growth conditions with average about 1 mm2. Fig. 1 shows the typical morphology of AlN sublimation crystal; this sample was grown at 18001C, 500 Torr for 10 h with source–substrate distance 2 mm. The same result was reported by Balkas et al. [9] who speculated that the sub-grains were caused by separate, isolated nucleation, due to the thermal decomposition of the SiC substrate. To suppress this decomposition, a slightly higher (800 Torr) nitrogen pressure was employed during our initial heating to the growth temperature. The effectiveness of this procedure was tested by examining a SiC substrate subjected to only the heating and cooling steps under 800 Torr nitrogen pressure, without any AlN grow. No significant SiC decomposition was observed by optical microscopy. Nevertheless, AlN crystals grown by sublimation on SiC substrates using this technique still contained many sub-grains. Thus, it is not SiC decomposition, but the initial nucleation state which leads to the sub-grain growth mode. Therefore the initial AlN nucleation and subsequent growth was studied to provide insights into the sub-grain formation. 3.2. Nucleation and further growth on as-received 6H-SiC substrates Although the surface polarity of the SiC substrate and the misorientation has a pronounced impact on the morphology of the epitaxy IIInitrides films [11–14], Balkas et al. did not specify
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Fig. 1. Optical micrograph of AlN crystal grown on the on-axis (0 0 0 1)Si 6H-SiC substrate at 18001C, 500 Torr for 10 h with source–substrate distance 2 mm. The magnification is 120 .
which they employed in Ref. [9]. Therefore, three kinds of 6H-SiC wafers i.e. on-axis (0 0 0 1)Si, (0 0 0 1)C 6H-SiC and 3.51 off-axis (0 0 0 1)Si were used as the substrates in our paper. To investigate the initial nucleation and subsequent growth mode systematically, AlN was deposited on each kind of substrate under the above-described experimental condition for 15, 45 and 120 min, respectively. Their optical micrographs are shown in Fig. 2 denoted by (1), (2) and (3) in a sequence of increasing growth time. The mechanism of sub-grains formation were clearly seen from Fig. 2a (1–3) for the crystals grown on on-axis Si-terminated substrates. After 15 and 45 min growth, only individual nuclei were formed, with only slow lateral growth of the nuclei after 120 min. This growth mode inevitably led to the formation of hexagonal sub-grains due to the incomplete coalescence of the nuclei. The nuclea-
tion and growth modes on the 3.51 off-axis Siterminated substrate as shown in Fig. 2b(1–3) was very similar with those of on-axis Si-terminated substrate. Relatively large macrosteps were formed on the off-axis SiC substrates for the 15 and 45 grown samples. The 120 min growth sample showed larger grain size in Fig. 2b(3) than in Fig. 2a(3) probably because the stepped features enhanced the coalescence of the nuclei. Because the AlN deposited layers were discontinuous for 15 min growth samples on on-axis and off-axis Siterminated substrates, the (0 0 2) AlN peak could not be detected by the y–2y X-ray diffraction. Only a very weak and broad (0 0 2) peak could be detected for the 45 min samples. In contrast, the diffraction (0 0 2) peak from the 120 min growth sample was strong and sharp due to the formation of large size crystals. The growth on the on-axis C-terminated substrates was strikingly different. After sublimation growth, the majority of the substrate surfaces were very smooth and featureless. Some unusual areas where deposition formed irregular crystals are presented in Fig. 2c (1–3). The (0 0 2) AlN peak was not detected in the X-ray diffraction pattern of the as-grown surface even after 120 min of growth, but it was detected from the backside of this substrate. Apparently, AlN cannot be deposited on the C-terminated SiC substrate; instead the aluminum and nitrogen vapor recondensed on the backside of the substrate. A 14 h growth sample verified this hypothesis. As seen in the optical micrograph (Fig. 3), the as-grown surface only contained sparse hexagonal AlN crystals, which did not cover the whole substrate. In contrast, the sublimation growth on the Si-terminated substrate for only 10 h produced a thick AlN layer composed by several coalescence hexagons as shown in Fig. 1. After a long growth time, some regions of the C-terminated surface decomposed, so isolated AlN crystals could form. We conclude that bare C-terminated SiC wafer cannot be used as the substrate for AlN sublimation growth. Sasaki et al. [11,12] observed that the Si-terminated 6H-SiC surface produced smooth and featureless GaN by MOCVD technique, while the C-terminated surface yielded prominent hexagonal pyramids of GaN. The reason is generally
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Fig. 2. Optical micrographs of AlN crystals grown on the different as-received (0 0 0 1) 6H-SiC substrates with magnification 60 : (a) on-axis Si-terminated, (b) 3.51 off-axis Si-terminated and (c) on-axis C-terminated substrates; (1–3) denote the growth time sequence of 15, 45 and 120 min.
presumed that the surface mobility of adatoms is low on the N-terminated GaN surface grown on the C-terminated SiC substrate so step flow growth is hindered and statistical roughening proceeds [15]. Ren and Dow [16] proposed that the observed differences in GaN surface morphologies were due to large local microscopic lattice mismatch differences. The electrical polarity of the GaN/SiC interface created interfacial charges which locally altered the lattice mismatch. For our AlN films prepared by MOCVD, the different polarity substrates did not have any significant differences in morphology as observed by AFM, nor were there differences in the crystal quality as measured
by X-ray rocking curve. The reason why AlN could not be sublimated on C-terminated surface of SiC needs further investigation. 3.3. Nucleation and further growth on 6H-SiC substrates with AlN MOCVD buffer layer Since the AlN nucleation directly on the asreceived (0 0 0 1)Si 6H-SiC substrate was discontinuous which inevitably lead to sub-grain growth, the buffer layer technique successfully employed in GaN epitaxy growth by MOCVD [3] was introduced in the sublimation growth. Unlike the very thin buffer layer adopted in MOCVD growth, an
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Fig. 3. Optical micrograph of AlN crystal grown on the on-axis (0 0 0 1)C 6H-SiC substrate at 18001C, 500 Torr for 14 h with source– substrate distance 2 mm. The magnification is 120 .
Fig. 4. Optical micrographs of AlN crystals grown on the 3.51 off-axis (0 0 0 1)Si 6H-SiC substrates with magnification 60 : with AlN MOCVD buffer layers in the growth time sequence of 15, 45 and 120 min.
approximate 2 mm thick AlN buffer layer was grown on the SiC substrates since the AlN buffer layer may partially decompose at a high sublimation temperature before the nucleation begins. As was done for the as-received substrates without buffer layer, the nucleation and subsequent growth on the different substrates with buffer layer were studied for 15, 45, 120 min growth, respectively. The subsequent sublimation growth on the three kinds SiC substrates with AlN MOCVD buffer layer showed very similar morphologies. Therefore only the optical micrographs of off-axis series samples are presented in Fig. 4. The AlN sublimation growth on the MOCVD buffer layer exhibited continuous growth mode, but the density of cracks caused by the stress during cooling process
increased with the growth time. For the 15 min growth sample, only a few cracks were observed. The crack density increased much more for the 45 min grown sample. After 120 min growth the stress rings were clearly observed on the as-grown surface indicating a high level of stress, and some AlN pieces even peeled off from the substrates. Despite the crack, using the buffer layer single grain growth was achieved by the continuous nucleation on the buffer layer. The relative intensity of the AlN (0 0 0 2) X-ray diffraction peaks from the series of off-axis substrates increased with the growth time (Fig. 5). The X-ray rocking curve from the AlN crystals grown on the Si-terminated substrates had slightly better quality than those on the C-
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Table 1 FWHMs of double crystal X-ray rocking curve of (0 0 0 2) peak AlN films on different (0 0 0 1) 6H-SiC substrates with AlN MOCVD buffer layer Growth condition
Si-terminated on-axis substrates
Si-terminated 3.51 off-axis substrates
C-terminated on-axis substrates
No growth 15 min growth 45 min growth
511 359 369
581 293 367
560 491 647
Fig. 5. y–2y X-ray diffraction of the samples grown on the 3.51 off-axis (0 0 0 1)Si 6H-SiC substrates with AlN MOCVD buffer layers for 15, 45 and 120 min.
terminated substrates. As seen from the results summarized in Table 1, the AlN sublimation crystals for 15 min growth had better qualities than those of the AlN MOCVD buffer layer without growth sample, but the qualities deteriorated for 45 min growth because of the increase in the crack density. The X-ray rocking curve for the 15 min growth sample on the off-axis substrate with buffer layer is presented in Fig. 6. It has the narrowest full-width at half-maximum value of 293 arcsec. The atomic force microscopy (AFM) images of the off-axis series samples also showed that the morphology of the AlN deposited by MOCVD improved after the subsequent sublimation growth due to the formation of more faceted surface. Asreceived off-axis (0 0 0 1)Si 6H-SiC substrates typically contain polishing induced randomly oriented scratches as shown in Fig. 7(a). The surface of the AlN MOCVD buffer layer grown on the off-axis substrate had irregularly oriented island-like features, indicating a 3D growth mode in Fig. 7(b). After 15 min of subsequent sublimation growth, a zig-zag stepped surface structure was produced with smooth terraces as shown in Fig. 7(c). Step bunching was also clearly seen on the image. Most terrace widths and step heights varied in the range of 300–500 nm and 10–50 nm, respectively. Therefore, we conclude that the AlN sublimed on the
Fig. 6. Double crystal X-ray rocking curve around the (0 0 0 2) AlN reflection from the 15 min growth samples grown on the 3.51 off-axis (0 0 0 1)Si 6H-SiC substrates with AlN MOCVD buffer layers.
MOCVD buffer layer by the step-flow growth mode. Fig. 7(d) showed the image of subsequent 45 min sublimation growth sample. The zig-zag stepped surface was still observed but each terrace surface became much rougher. This is consistent with the results of X-ray rocking curve. According to the characterization of the threading dislocation structure in GaN films by X-ray rocking curve [17], the symmetric (0 0 0 2) rocking curves are only sensitive to the screw dislocations due to its Burgers vectors lying parallel to [0 0 0 1], causing the distortion of c-plane. Therefore, the more regular and smooth the surface is, the lower the screw dislocation density, and hence, the (0 0 0 2) peaks are narrower.
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Fig. 7. AFM images of off-axis series samples: (a) as-received off-axis wafer, (b) with AlN MOCVD buffer layer, no further growth, (c) 15 minutes subsequent sublimation growth and (d) 45 min subsequent sublimation growth. Z scale of AFM: 200 nm.
4. Discussion Although the continuous nucleation and step flow growth mode could be achieved for the AlN sublimation growth using the MOCVD buffer layer technique, cracks still appeared even after only 15 min growth. For AlN growth on 6H-SiC, compressive lattice mismatch stresses are relieved after a few nanometers of growth [18]. So the only remaining stress is due to the mismatch of thermal expansion coefficients when the composite AlN/ SiC structure is cooled to the room temperature. Since the thermal expansion of AlN is greater than that of SiC [19], the tensile stress is the inevitable result of the AlN growth layer on SiC substrates. For the AlN sublimation directly on the asreceived substrate, cracking is not a serious
problem because most stresses could be relaxed by the sub-grain boundaries. But for the single grain growth on the substrate with MOCVD buffer layer, the unrelaxed residual tensile stresses will bend the whole structure toward the AlN layer and cause serious cracking. As mentioned above, some pieces even peeled off from the substrate for the 120 min growth sample. Analysis of the front and back sides of these pieces by micro-Raman indicated that the cracks occurred across the AlN/SiC interface into the SiC substrates. As the Raman spectroscopy shows in Fig. 8, all the peaks of the as-grown surface can be assigned as AlN Raman modes, the peaks of the peeling-surface are due to the SiC Raman modes.
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Fig. 8. Raman spetra of the crystalline separated from SiC substrate: (a) as-grown surface, (b) back side of the peeling crystal.
Cracking has been a serious problem which limits the maximum of the thickness and lateral size of the AlN crystals sublimated on the SiC substrates with AlN MOCVD buffer layers. In contrast, a compressive thermal mismatch stress is introduced in the AlN layer grown on the sapphire substrate because the thermal expansion of AlN is smaller than that of sapphire. Since the cracking threshold for the compressive strength is generally much higher than that of the tensile strength for AlN ceramic, cracking may be reduced by growth on sapphire substrate. Therefore sapphire was also tried as a substrate for AlN sublimation growth with the MOCVD buffer layer first deposited. Unfortunately the sapphire decomposed at the AlN sublimation temperature (above 17001C). This decomposition was not suppressed even by a thick AlN MOCVD layer. Thus, sapphire is not a viable substrate for the AlN sublimation growth.
investigated in this paper. Direct growth on Sipolarity as-received substrates was discontinuous caused by the initial sparse nucleation and slow lateral growth of nuclei. Several hexagonal subgrains were usually obtained on the substrates due to the incomplete coalescence of the nuclei. In contrast, no sublimation growth occurred on the as-received C-terminated substrates. Continuous films could then be grown on all the substrates with AlN MOCVD buffer layers. The tensile stress due to thermal expansion coefficient mismatch in the AlN layer during cooling process caused cracking across the AlN/SiC interface into the SiC substrates which limits the maximum of the thickness and lateral size of the AlN crystals. The best way to overcome this problem is to use a homoepitaxial seed for further sandwich sublimation technique growth of AlN.
Acknowledgements 5. Conclusion The influence of the buffer layer and 6H-SiC substrate polarity on the nucleation of AlN grown by the sublimation sandwich technique have been
We are grateful for the support of this research from BMDO (Contract No. N00014-98C-0407) and ONR (Contract No. N0001499-1-0104).
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[9] C.M. Balkas, Z. Sitar, T. Zheleva, L. Bergman, R. Nemanich, R.F. Davis, J. Crystal Growth 179 (1997) 363. [10] J.H. Edgar, L.H. Robins, S.E. Coatney, L. Liu, J. Chaudhuri, K. Ignatiev, Z. Rek, Mater. Sci. Forum 338–342 (2000) 1599. [11] T. Sasaki, T. Matsuoka, J. Appl. Phys. 64 (1988) 4531. [12] T. Sasaki, T. Matsuoka, J. Appl. Phys. 61 (1987) 2533. [13] F.A. Ponce, M.A. O’Keefe, E.C. Nelson, Philos. Mag. A. 74 (1996) 777. [14] P. Vermaut, P. Ruterana, G. Nouet, Philos. Mag. A. 1215 (1997) 1215. [15] E.C. Piquette, P.M. Bridger, R.A. Beach, T.C. McGill, MRS Internet J. Nitride Semicond. Res. 4S1 (1999) G377. [16] S.Y. Ren, J.D. Dow, Appl. Phys. Lett. 69 (1996) 251. [17] M.J. Hordon, B.L. Averbach, Acta Metallogr. 9 (1961) 237. [18] K. Hiramatsu, T. Detchprohm, I. Akasaki, Jpn. J. Appl. Phys. 32 (1993) 1528. [19] K. Wang, R. Reeber, MRS Internet J. Nitride Semicond. Res. 4S1 (1999) G318.
Influence of buffer layer and 6H-SiC substrate polarity on the nucleation of AlN grown by the sublimation sandwich technique Y. Shia, Z.Y. Xiea, L.H. Liua, B. Liua, J.H. Edgara,*, M. Kuballb a
Department of Chemical Engineering, Durland Hall, Kansas State University, Manhattan, KS 66506-5102, USA b H.H. Wills Physics Laboratory, University of Bristol, BS8 1TL, UK Received 23 July 2000; accepted 13 June 2001 Communicated by A.F. Witt
Abstract The initial nucleation stage of AlN crystal grown by the sublimation method on on-axis (0 0 0 1)Si and (0 0 0 1)C 6HSiC as well as 3.51 off-axis (0 0 0 1)Si 6H-SiC substrates was investigated. Sublimation growth from a pure sintered AlN source was carried out in a resistively heated furnace with a source temperature of about 18001C at a nitrogen pressure of 500 Torr. Direct growth on Si-terminated as-received SiC substrates was discontinuous, marked by sparse nucleation and slow lateral growth of nuclei. Several hexagonal sub-grains were usually obtained on the substrates after a long growth time (more than 3 h) due to the incomplete coalescence of the nuclei. In contrast, no sublimation growth occurred on the as-received C-terminated substrates. To enhance two-dimensional (2D) growth, an AlN epitaxial layer was first deposited on the substrates by MOCVD before sublimation growth. Continuous films could then be grown on all the substrates with AlN MOCVD buffer layers. The tensile stress of the AlN layer due to thermal expansion coefficient mismatch between SiC and AlN caused cracking across the AlN/SiC interface into the SiC substrates during the cooling process limiting the maximum of the thickness and lateral size of the AlN crystals. r 2001 Elsevier Science B.V. All rights reserved. Keywords: A1. Crystal morphology; A1. Nucleation; A2. Growth from vapor; A2. Seed crystals; B1. Nitrides
1. Introduction The GaN-based semiconductors have developed rapidly in blue and green light emitting diodes and blue emitting lasers [1]. The primary substrates for epitaxial growth of GaN are SiC
*Corresponding author. Tel.: +1-913-532-4320; fax: +1913-532-7372. E-mail address: [email protected] (J.H. Edgar).
and sapphire, which are poorly lattice and thermal expansion coefficient matched to GaN. Aluminum nitride, a wurtzite structure wide band gap (6.2 eV) semiconductor, is a good candidate substrate for GaN epitaxial films due to its relatively small lattice constant mismatch along the a-axis ( 3.5%), good thermal stability (melting point >25001C), high resistivity and similar coefficient of thermal expansion [2]. The use of AlN buffer layers in MOCVD significantly improves the quality of GaN epitaxy [3] also
0022-0248/01/$ - see front matter r 2001 Elsevier Science B.V. All rights reserved. PII: S 0 0 2 2 - 0 2 4 8 ( 0 1 ) 0 1 5 6 0 - 3
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suggesting that AlN would be an ideal substrate for GaN epitaxy. The most successful method for producing bulk AlN crystal is by sublimation growth (vapor phase transport). In this technique, AlN undergoes the reaction AlN=Al+12N2 in a high temperature source zone and this reaction is driven in the reverse direction to form AlN single crystals in a lower temperature zone. This technique was most successfully developed by Slack and McNelly [4,5] in the mide-1970s. Large single crystals of AlN up to 1 cm long and 0.4 cm in diameter were grown by self-seeding in the sharp tip of sealed tungsten crucibles in an rf induction furnace. Recently, Rojo et al. [6,7] have further advanced this technique to produce single crystals up to 15 mm in diameter. Tanaka et al. [8] investigated the axis orientation of AlN freely nucleated on graphite crucible walls. Most relevant to this work, Balkas et al. [9] produced single crystal platelets of AlN (p1 mm thick) on 10 10 mm2 6H-SiC substrates in a resistively heated graphite furnace. However, the crystals contained small (B2 2 mm2) individual hexagonal sub-grains, which was attributed to the severe degradation of the SiC substrates resulting in isolated stable nucleation sites. A prior investigation by the author’s group [10] showed freely nucleated needles and platelets and exhibited much better crystal quality than the crystals grown on 6H-SiC substrates although by the latter method the crystal orientation and initial nucleation was usually better controlled. Up to now, no systemic investigation of the initial nucleation stage of AlN sublimation on 6H-SiC substrate has been reported, although proper nucleation critically effects the quality of the subsequently grown crystals. Nor has the buffer layer technique common to MOCVD growth been employed to prepare substrates for the sublimation growth. Herein for the first time the initial nucleation stage of AlN crystal grown by sublimation method on on-axis (0 0 0 1)Si and (0 0 0 1)C 6HSiC as well as off-axis (0 0 0 1)Si 6H-SiC substrates were investigated. The influence of the AlN buffer layer deposited on different SiC substrates by MOCVD on the subsequent AlN sublimation growth were also studied.
2. Experimental procedures 2.1. Crystal growth Sublimation growths were conducted in a resistively heated furnace using tungsten wire mesh heating elements. The seed and source were enclosed in two concentric tungsten crucibles to contain the Al vapor as it deteriorates the tungsten heating elements. Sintered AlN with approximately 1% oxygen as the main impurity analyzed by a standard inert gas fusion method (TC-136, Leco Co.) was used as source and seed holders. The as-received (0 0 0 1) 6H-SiC silicon terminated (on-axis and 3.51 off-axis) and carbon terminated wafers (only on-axis) were cut into about 1 cm2 substrates. All the substrates were ultrasonically cleaned in organic solvents (TCE, acetone and methanol), rinsed with DI water, and dried with high purity nitrogen gas before loaded into the furnace. The distance between the source and the substrate was kept at 2 mm. In order to reduce SiC dissociation from the backside of the seed, it was covered by an AlN cap. The growth ambient was 99.99993% pure nitrogen. Before growth, the furnace was evacuated to a pressure less than 10 4 Torr, then purged by nitrogen twice to remove residual gases. During the heating process the furnace was maintained at 800 Torr, as this higher nitrogen pressure suppressed both AlN sublimation and the decomposition of the SiC substrate which may roughen the substrate surface causing unfavorable AlN nucleation. As soon as the temperature reached the desired value, the background pressure was decreased to 500 Torr very quickly to start the sublimation growth. The growth pressure was maintained by an automatic throttle valve. At the end of the growth, the system pressure was increased to 800 Torr again during cooling to avoid unwanted AlN recondensation. The vertical temperature profile of the growth chamber was measured by an optical pyrometer focused on a moving target at different furnace output powers and pressure. To obtain comparable results, the growth temperature was kept constant for all growths by fixing the output power. The in situ growth temperature was
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measured by the same optical pyrometer focused on the top lid of the outside crucible. According to the pre-measured temperature profile, at our experimental condition the source temperature was about 18001C and the source-substrate temperature gradient was about 101C for 2 mm source-substrate distance. Several experiments carried at the temperature above 20001C showed that higher growth temperatures had no significant influences on the nucleation and morphology except the growth rate. Therefore the low temperature was employed to avoid the degradation of the furnace and the SiC substrates. 2.2. MOCVD technique The AlN buffer layer growth was performed in a vertical-type MOCVD reactor operated at low pressure (76 Torr). Before loading into the reactor, the 6H-SiC substrates were cleaned using the same procedure described above. Trimethylaluminum (TMA) and ammonia (NH3) were the Al and N sources, and Pd-cell purified H2 was the carrier gas. The substrates were preheated at 11001C for 10 min in 3 slm of H2 flow for surface cleaning. Then the AlN films were deposited at the same temperature under the H2, NH3 and TMA flow rates of 3, 3 slm and 30 sccm, respectively. An approximate 2 mm thick AlN buffer layer for the subsequent sublimation growth was obtained for the 4 h growth time. 2.3. Characterization The as-grown AlN crystals were first characterized by Nomarski differential interference contrast microscopy (NDIC). Macrostructural features were observed at magnifications of 5–40X. X-ray diffraction measurements were conducted using XDS 2000 diffractometer (Scintag Inc.) with the Cu Ka1 radiation. The surface morphology of the samples was examined by Digital Instruments Nanoscope E AFM in contact mode. MicroRaman spectra were obtained at Ultraviolet Renishaw with the 325 nm line of a HeCd laser as excitation source. The spot size and the spectral resolution were 1–2 mm and 3–4 cm 1, respectively.
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3. Results 3.1. Long time growth result of AlN on (0 0 0 1)Si 6H-SiC Initially, different growth temperature, pressure and substrate-seed distance were tried on the Sipolarity 6H-SiC substrate to optimize the growth condition. The pressure during the heating and cooling process was first kept consistent with the growth pressure. The typical sublimation time was usually about 10 h. In all experiments, crystalline AlN was deposited on the majority of the SiC substrates. Every AlN crystal thus formed contains several hexagonal sub-grains. Grain size varied slightly under different growth conditions with average about 1 mm2. Fig. 1 shows the typical morphology of AlN sublimation crystal; this sample was grown at 18001C, 500 Torr for 10 h with source–substrate distance 2 mm. The same result was reported by Balkas et al. [9] who speculated that the sub-grains were caused by separate, isolated nucleation, due to the thermal decomposition of the SiC substrate. To suppress this decomposition, a slightly higher (800 Torr) nitrogen pressure was employed during our initial heating to the growth temperature. The effectiveness of this procedure was tested by examining a SiC substrate subjected to only the heating and cooling steps under 800 Torr nitrogen pressure, without any AlN grow. No significant SiC decomposition was observed by optical microscopy. Nevertheless, AlN crystals grown by sublimation on SiC substrates using this technique still contained many sub-grains. Thus, it is not SiC decomposition, but the initial nucleation state which leads to the sub-grain growth mode. Therefore the initial AlN nucleation and subsequent growth was studied to provide insights into the sub-grain formation. 3.2. Nucleation and further growth on as-received 6H-SiC substrates Although the surface polarity of the SiC substrate and the misorientation has a pronounced impact on the morphology of the epitaxy IIInitrides films [11–14], Balkas et al. did not specify
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Fig. 1. Optical micrograph of AlN crystal grown on the on-axis (0 0 0 1)Si 6H-SiC substrate at 18001C, 500 Torr for 10 h with source–substrate distance 2 mm. The magnification is 120 .
which they employed in Ref. [9]. Therefore, three kinds of 6H-SiC wafers i.e. on-axis (0 0 0 1)Si, (0 0 0 1)C 6H-SiC and 3.51 off-axis (0 0 0 1)Si were used as the substrates in our paper. To investigate the initial nucleation and subsequent growth mode systematically, AlN was deposited on each kind of substrate under the above-described experimental condition for 15, 45 and 120 min, respectively. Their optical micrographs are shown in Fig. 2 denoted by (1), (2) and (3) in a sequence of increasing growth time. The mechanism of sub-grains formation were clearly seen from Fig. 2a (1–3) for the crystals grown on on-axis Si-terminated substrates. After 15 and 45 min growth, only individual nuclei were formed, with only slow lateral growth of the nuclei after 120 min. This growth mode inevitably led to the formation of hexagonal sub-grains due to the incomplete coalescence of the nuclei. The nuclea-
tion and growth modes on the 3.51 off-axis Siterminated substrate as shown in Fig. 2b(1–3) was very similar with those of on-axis Si-terminated substrate. Relatively large macrosteps were formed on the off-axis SiC substrates for the 15 and 45 grown samples. The 120 min growth sample showed larger grain size in Fig. 2b(3) than in Fig. 2a(3) probably because the stepped features enhanced the coalescence of the nuclei. Because the AlN deposited layers were discontinuous for 15 min growth samples on on-axis and off-axis Siterminated substrates, the (0 0 2) AlN peak could not be detected by the y–2y X-ray diffraction. Only a very weak and broad (0 0 2) peak could be detected for the 45 min samples. In contrast, the diffraction (0 0 2) peak from the 120 min growth sample was strong and sharp due to the formation of large size crystals. The growth on the on-axis C-terminated substrates was strikingly different. After sublimation growth, the majority of the substrate surfaces were very smooth and featureless. Some unusual areas where deposition formed irregular crystals are presented in Fig. 2c (1–3). The (0 0 2) AlN peak was not detected in the X-ray diffraction pattern of the as-grown surface even after 120 min of growth, but it was detected from the backside of this substrate. Apparently, AlN cannot be deposited on the C-terminated SiC substrate; instead the aluminum and nitrogen vapor recondensed on the backside of the substrate. A 14 h growth sample verified this hypothesis. As seen in the optical micrograph (Fig. 3), the as-grown surface only contained sparse hexagonal AlN crystals, which did not cover the whole substrate. In contrast, the sublimation growth on the Si-terminated substrate for only 10 h produced a thick AlN layer composed by several coalescence hexagons as shown in Fig. 1. After a long growth time, some regions of the C-terminated surface decomposed, so isolated AlN crystals could form. We conclude that bare C-terminated SiC wafer cannot be used as the substrate for AlN sublimation growth. Sasaki et al. [11,12] observed that the Si-terminated 6H-SiC surface produced smooth and featureless GaN by MOCVD technique, while the C-terminated surface yielded prominent hexagonal pyramids of GaN. The reason is generally
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Fig. 2. Optical micrographs of AlN crystals grown on the different as-received (0 0 0 1) 6H-SiC substrates with magnification 60 : (a) on-axis Si-terminated, (b) 3.51 off-axis Si-terminated and (c) on-axis C-terminated substrates; (1–3) denote the growth time sequence of 15, 45 and 120 min.
presumed that the surface mobility of adatoms is low on the N-terminated GaN surface grown on the C-terminated SiC substrate so step flow growth is hindered and statistical roughening proceeds [15]. Ren and Dow [16] proposed that the observed differences in GaN surface morphologies were due to large local microscopic lattice mismatch differences. The electrical polarity of the GaN/SiC interface created interfacial charges which locally altered the lattice mismatch. For our AlN films prepared by MOCVD, the different polarity substrates did not have any significant differences in morphology as observed by AFM, nor were there differences in the crystal quality as measured
by X-ray rocking curve. The reason why AlN could not be sublimated on C-terminated surface of SiC needs further investigation. 3.3. Nucleation and further growth on 6H-SiC substrates with AlN MOCVD buffer layer Since the AlN nucleation directly on the asreceived (0 0 0 1)Si 6H-SiC substrate was discontinuous which inevitably lead to sub-grain growth, the buffer layer technique successfully employed in GaN epitaxy growth by MOCVD [3] was introduced in the sublimation growth. Unlike the very thin buffer layer adopted in MOCVD growth, an
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Fig. 3. Optical micrograph of AlN crystal grown on the on-axis (0 0 0 1)C 6H-SiC substrate at 18001C, 500 Torr for 14 h with source– substrate distance 2 mm. The magnification is 120 .
Fig. 4. Optical micrographs of AlN crystals grown on the 3.51 off-axis (0 0 0 1)Si 6H-SiC substrates with magnification 60 : with AlN MOCVD buffer layers in the growth time sequence of 15, 45 and 120 min.
approximate 2 mm thick AlN buffer layer was grown on the SiC substrates since the AlN buffer layer may partially decompose at a high sublimation temperature before the nucleation begins. As was done for the as-received substrates without buffer layer, the nucleation and subsequent growth on the different substrates with buffer layer were studied for 15, 45, 120 min growth, respectively. The subsequent sublimation growth on the three kinds SiC substrates with AlN MOCVD buffer layer showed very similar morphologies. Therefore only the optical micrographs of off-axis series samples are presented in Fig. 4. The AlN sublimation growth on the MOCVD buffer layer exhibited continuous growth mode, but the density of cracks caused by the stress during cooling process
increased with the growth time. For the 15 min growth sample, only a few cracks were observed. The crack density increased much more for the 45 min grown sample. After 120 min growth the stress rings were clearly observed on the as-grown surface indicating a high level of stress, and some AlN pieces even peeled off from the substrates. Despite the crack, using the buffer layer single grain growth was achieved by the continuous nucleation on the buffer layer. The relative intensity of the AlN (0 0 0 2) X-ray diffraction peaks from the series of off-axis substrates increased with the growth time (Fig. 5). The X-ray rocking curve from the AlN crystals grown on the Si-terminated substrates had slightly better quality than those on the C-
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Table 1 FWHMs of double crystal X-ray rocking curve of (0 0 0 2) peak AlN films on different (0 0 0 1) 6H-SiC substrates with AlN MOCVD buffer layer Growth condition
Si-terminated on-axis substrates
Si-terminated 3.51 off-axis substrates
C-terminated on-axis substrates
No growth 15 min growth 45 min growth
511 359 369
581 293 367
560 491 647
Fig. 5. y–2y X-ray diffraction of the samples grown on the 3.51 off-axis (0 0 0 1)Si 6H-SiC substrates with AlN MOCVD buffer layers for 15, 45 and 120 min.
terminated substrates. As seen from the results summarized in Table 1, the AlN sublimation crystals for 15 min growth had better qualities than those of the AlN MOCVD buffer layer without growth sample, but the qualities deteriorated for 45 min growth because of the increase in the crack density. The X-ray rocking curve for the 15 min growth sample on the off-axis substrate with buffer layer is presented in Fig. 6. It has the narrowest full-width at half-maximum value of 293 arcsec. The atomic force microscopy (AFM) images of the off-axis series samples also showed that the morphology of the AlN deposited by MOCVD improved after the subsequent sublimation growth due to the formation of more faceted surface. Asreceived off-axis (0 0 0 1)Si 6H-SiC substrates typically contain polishing induced randomly oriented scratches as shown in Fig. 7(a). The surface of the AlN MOCVD buffer layer grown on the off-axis substrate had irregularly oriented island-like features, indicating a 3D growth mode in Fig. 7(b). After 15 min of subsequent sublimation growth, a zig-zag stepped surface structure was produced with smooth terraces as shown in Fig. 7(c). Step bunching was also clearly seen on the image. Most terrace widths and step heights varied in the range of 300–500 nm and 10–50 nm, respectively. Therefore, we conclude that the AlN sublimed on the
Fig. 6. Double crystal X-ray rocking curve around the (0 0 0 2) AlN reflection from the 15 min growth samples grown on the 3.51 off-axis (0 0 0 1)Si 6H-SiC substrates with AlN MOCVD buffer layers.
MOCVD buffer layer by the step-flow growth mode. Fig. 7(d) showed the image of subsequent 45 min sublimation growth sample. The zig-zag stepped surface was still observed but each terrace surface became much rougher. This is consistent with the results of X-ray rocking curve. According to the characterization of the threading dislocation structure in GaN films by X-ray rocking curve [17], the symmetric (0 0 0 2) rocking curves are only sensitive to the screw dislocations due to its Burgers vectors lying parallel to [0 0 0 1], causing the distortion of c-plane. Therefore, the more regular and smooth the surface is, the lower the screw dislocation density, and hence, the (0 0 0 2) peaks are narrower.
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Fig. 7. AFM images of off-axis series samples: (a) as-received off-axis wafer, (b) with AlN MOCVD buffer layer, no further growth, (c) 15 minutes subsequent sublimation growth and (d) 45 min subsequent sublimation growth. Z scale of AFM: 200 nm.
4. Discussion Although the continuous nucleation and step flow growth mode could be achieved for the AlN sublimation growth using the MOCVD buffer layer technique, cracks still appeared even after only 15 min growth. For AlN growth on 6H-SiC, compressive lattice mismatch stresses are relieved after a few nanometers of growth [18]. So the only remaining stress is due to the mismatch of thermal expansion coefficients when the composite AlN/ SiC structure is cooled to the room temperature. Since the thermal expansion of AlN is greater than that of SiC [19], the tensile stress is the inevitable result of the AlN growth layer on SiC substrates. For the AlN sublimation directly on the asreceived substrate, cracking is not a serious
problem because most stresses could be relaxed by the sub-grain boundaries. But for the single grain growth on the substrate with MOCVD buffer layer, the unrelaxed residual tensile stresses will bend the whole structure toward the AlN layer and cause serious cracking. As mentioned above, some pieces even peeled off from the substrate for the 120 min growth sample. Analysis of the front and back sides of these pieces by micro-Raman indicated that the cracks occurred across the AlN/SiC interface into the SiC substrates. As the Raman spectroscopy shows in Fig. 8, all the peaks of the as-grown surface can be assigned as AlN Raman modes, the peaks of the peeling-surface are due to the SiC Raman modes.
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Fig. 8. Raman spetra of the crystalline separated from SiC substrate: (a) as-grown surface, (b) back side of the peeling crystal.
Cracking has been a serious problem which limits the maximum of the thickness and lateral size of the AlN crystals sublimated on the SiC substrates with AlN MOCVD buffer layers. In contrast, a compressive thermal mismatch stress is introduced in the AlN layer grown on the sapphire substrate because the thermal expansion of AlN is smaller than that of sapphire. Since the cracking threshold for the compressive strength is generally much higher than that of the tensile strength for AlN ceramic, cracking may be reduced by growth on sapphire substrate. Therefore sapphire was also tried as a substrate for AlN sublimation growth with the MOCVD buffer layer first deposited. Unfortunately the sapphire decomposed at the AlN sublimation temperature (above 17001C). This decomposition was not suppressed even by a thick AlN MOCVD layer. Thus, sapphire is not a viable substrate for the AlN sublimation growth.
investigated in this paper. Direct growth on Sipolarity as-received substrates was discontinuous caused by the initial sparse nucleation and slow lateral growth of nuclei. Several hexagonal subgrains were usually obtained on the substrates due to the incomplete coalescence of the nuclei. In contrast, no sublimation growth occurred on the as-received C-terminated substrates. Continuous films could then be grown on all the substrates with AlN MOCVD buffer layers. The tensile stress due to thermal expansion coefficient mismatch in the AlN layer during cooling process caused cracking across the AlN/SiC interface into the SiC substrates which limits the maximum of the thickness and lateral size of the AlN crystals. The best way to overcome this problem is to use a homoepitaxial seed for further sandwich sublimation technique growth of AlN.
Acknowledgements 5. Conclusion The influence of the buffer layer and 6H-SiC substrate polarity on the nucleation of AlN grown by the sublimation sandwich technique have been
We are grateful for the support of this research from BMDO (Contract No. N00014-98C-0407) and ONR (Contract No. N0001499-1-0104).
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