GaN(0001) thin films are grown using radio frequency plasma assisted molecular beam epitaxy. By changing the growth temperature, anisotropic growth rate behavior is observed in both the step-flow growth mode and the 2D island growth mode. Tunneling scanning microscopy reveals, in the step-flow growth mode, strong influences from the growth anisotropy on the shape of the terrace edges, resulting in striking differences between hexagonal and cubic films. In the 2D nucleation growth mode, triangularly shaped islands are formed. The significance of growth anisotropy to growing high quality GaN films is discussed.published_or_final_versio

Seutter, SM

Tong, SY

Wu, H

Xie, MH

Zheng, LX

Zhu, WK

Physical Review Letters

HKU Scholars Hub

Title Anisotropic step-flow growth and island growth of GaN(0001) bymolecular beam epitaxyAuthor(s) Xie, MH; Seutter, SM; Zhu, WK; Zheng, LX; Wu, H; Tong, SYCitation Physical Review Letters, 1999, v. 82 n. 13, p. 2749-2752Issued Date 1999URL http://hdl.handle.net/10722/42453Rights Creative Commons: Attribution 3.0 Hong Kong LicenseVOLUME 82, NUMBER 13 P HY S I CA L REV I EW LE T T ER S 29 MARCH 1999Anisotropic Step-Flow Growth and Island Growth of GaN(0001) by Molecular Beam EpitaxyM.H. Xie, S.M. Seutter, W.K. Zhu, L.X. Zheng, Huasheng Wu, and S. Y. TongDepartment of Physics, The University of Hong Kong, Pokfulam Road, Hong Kong(Received 30 October 1998)Anisotropic growth is observed for GaN(0001) during molecular beam epitaxy for both the step-flow growth mode and two-dimensional (2D) nucleation growth mode. Using scanning tunnelingmicroscopy, we find that in the step-flow growth mode, growth anisotropy strongly influences theshape of terrace edges, making them strikingly different between hexagonal and cubic films. Inthe 2D nucleation growth mode, anisotropic growth results in triangularly shaped islands. Theimportance of understanding growth anisotropy to achieve better grown GaN films is discussed.[S0031-9007(99)08731-1]PACS numbers: 68.55.Jk, 61.16.Ch, 68.35.Bs, 68.35.FxThe growth of GaN thin films has attracted intenserecent interest, because important high-performance GaN-based devices have been fabricated [1–3]. For devicefabrication, the thin films are grown mostly by metal-organic vapor phase epitaxy (MOVPE). On the otherhand, the use of molecular beam epitaxy (MBE) to studyGaN growth has become increasingly popular because ofthe ease to incorporate in situ ultrahigh vacuum (UHV)characterization tools [4,5]. Up to now, there has beenlittle known concerning the microscopic aspects of growthmechanism of GaN films.We report here the observation of anisotropic growthbehavior of GaN(0001) thin films using radio frequency(rf) plasma assisted MBE. By changing the growthtemperature, we observe anisotropic growth behavior inboth the step-flow growth mode and 2D island growthmode. The speed in which the edge of step or islandgrows is mainly controlled by two factors: (i) the ease inwhich an arriving species (adatom) binds to a site at theedge [6] and (ii) the subsurface influenced diffusion rateof an adatom along an edge [7]. The former is mainlya local bonding interaction between an adatom and edgeatoms. The latter is the influence of subsurface atoms onthe mobility of an adatom. We shall use a microscopicpicture to explain the experimentally observed growthanisotropy and indicate which of the above two processesis dominant in the GaN(0001) case.The MBE system consists of a deposition chamberequipped with reflection high-energy electron diffraction(RHEED) and connected under UHV conditions to threeother chambers equipped with various characterizationtools. These include a variable-temperature scanningtunneling microscopy (STM), atomic force microscopy(AFM), low-energy electron diffraction (LEED), andx-ray photoemission spectroscopy (XPS). A Ga flux issupplied by a conventional Knudsen cell, and a N fluxis provided by an Oxford plasma unit. Typical pressurein the MBE chamber during deposition is 5 3 1025 torrat a N2 flow rate of 0.13 standard cubic centimetersper minute. The 6H-SiCs0001d substrate (Cree) is dicedinto 11 3 3 mm2 rectangular pieces. It is degreased0031-9007y99y82(13)y2749(4)$15.00in acetone and alcohol before being loaded into thevacuum. The surface is prepared in the UHV at 1100 –Cunder a flux of Si. This procedure results in a sp3 3p3 dR30– reconstruction on its Si face. STM picturesshow atomically flat terraces of more than 1000 Å wide[8]. On top of such a surface, a thick s,1 mmd GaNbuffer layer is grown at 650 –C. A GayN flux ratio ofapproximately 2 is maintained during growth, where theeffective fluxes of Ga and N are calibrated by growthrates of GaN in the Ga-limited and N-limited regimes,respectively. The growth rates are measured by RHEEDintensity oscillations and cross checked by total thicknessmeasurements. The typical growth rate is 0.26 Åysec.The atomic order of the substrate surface on which thegrowth starts and the GayN flux ratio are importantparameters to achieve a flat surface morphology. Previousstudies using different growth conditions have reported3D columnar growth [9].Further growth on such a thick GaN buffer is studiedboth at low temperature (500 –C) and high temperature(650 –C). Low temperature growth is shown to result in2D islands, while high temperature leads to the step-flowgrowth mode. To observe surface morphology by STM,the growing surface is thermally quenched at the sametime when the growth is stopped. To stop the growth, theGa and N shutters are closed simultaneously together withthe switching off of the plasma generator’s power supply.Sample thermal quenching is achieved by switching offthe dc power to the substrate. Since sample heating isby direct current, the sample cools down rapidly oncethe current is stopped. From RHEED, the surface of thefilm shows a “1 3 1” pattern. The film is determined tohave a Ga polarity [i.e., (0001) face], because we havevaried the Ga coverage on such samples and reproducedthe reconstruction structures reported by Smith et al. onthis face [10].In the following, we shall describe first the surfacemorphology of the GaN film under step-flow growth.Generally, steps act as sinks of adatoms. When thediffusion length of an adatom is long compared to theterrace width, a step-flow growth mode is normally© 1999 The American Physical Society 2749VOLUME 82, NUMBER 13 P HY S I CA L REV I EW LE T T ER S 29 MARCH 1999achieved where the growth is seen by the advance of stepsand no island is nucleated. An STM image of a GaNsurface grown in the step-flow growth mode at 650 –C isshown in Fig. 1. A number of terraces, some as wideat 1000 Å, are clearly visible. A most striking featurein this picture is that as one descends the terraces alonga symmetry crystallographic direction, e.g., f1 0 1 0g, thestep edges alternate between jagged and smooth. On agiven terrace, a smooth edge becomes a jagged edge bychanging a 60– direction.To explain this rather unusual step-edge anisotropy, werefer to the atomic configuration of a (111) bilayer lat-tice shown in Fig. 2(a). We label the symmetry crystal-lographic directions in both the fcc(111) (diamond or zincblende) and hcp(0001) (wurtzite) conventions. Consider-ing only the dangling bonds in the bilayer, we see thatthere are two types of step edges: Type A edge, normalto f1 1 2g or equivalent directions, is characterized by twodangling bonds per edge atom and type B edge, normalto f2 1 1g or equivalent directions, is characterized by onedangling bond per edge atom. It is reasonable to assume,and we shall show later that this assumption is correct,that a type A edge will grow faster than a type B edgebecause of the greater tendency for an arriving adatom tobind to a type A edge. If the bilayer is an island on asurface, the faster growth speed in the direction normalto a type A edge will ultimately result in a triangularlyshaped island whose vertices point along the f1 1 2g andequivalent directions. In the step-flow growth mode, thefaster growth speed in the direction normal to a type Aedge will result in a jagged shaped edge. Each protru-sion of the jagged edge is bounded by type B edges. Onthe other hand, the slower growth speed type B edge willFIG. 1. STM picture of step-flow growth mode of GaNat substrate temperature ­ 650 –C, sample bias 22.5 V, andtunneling current 0.1 nA.2750be smooth. To explain the alternating smooth-jagged se-quence of the terrace edges, we refer to Fig. 2(b), wherewe show that a hcp film has alternating type A and Bedges along a given crystallographic symmetry direction.However, at 660– or 180– away, the order of the type Aand B edges is reversed, exactly corresponding to the ob-servations of the STM picture. A corollary to this descrip-tion is that a smooth-jagged edge ordering is evidence thatthe terraces are separated by a single (or odd multiples)bilayer height. STM measurements of the terrace heightsin Fig. 1 confirm that all the terraces are separated by asingle bilayer height.The growth anisotropy will lead to step bunching in ahcp film, because the fast-growing steps will ultimatelycatch up with the slow ones as growth proceeds. Thisprocess will lead to qualitatively different edge appear-ances between a cubic and hexagonal film. Referring toFig. 2(b), for a fcc(111) film, the type A edge points al-ways along the f1 1 2g and equivalent directions on allterraces while the type B edge points always along thef2 1 1g and equivalent directions. Thus, in a step-flowgrowth mode for a fcc(111) film, step bunching does notoccur, and the appearance of smooth and jagged edgeswill remain on each terrace [see Fig. 3(a)]. On the otherhand, for a hcp(0001) film, all terrace edges will eventu-ally become smooth when the fast-growing type A edgescatch up with the slow-growing type B edges. Thus, ifone descends the terraces along the f1 0 1 0g direction,FIG. 2. (a) Plan view of a bilayer, showing type A, B edgesand fcc, hcp directions. For a hcp direction, the edge may beeither A or B, depending on the stacking sequence. (b) Cross-sectional view of hcp, fcc films on a hcp substrate.VOLUME 82, NUMBER 13 P HY S I CA L REV I EW LE T T ER S 29 MARCH 1999FIG. 3. Terrace step shapes due to step-flow growthanisotropy for (a) a fcc film and (b) a hcp film.one will encounter double height steps. On the otherhand, if one descends the terraces along the f1 1 0 0g di-rection and after descending each terrace, alternately turn-ing 660–, one will encounter only single bilayer stepheights [see Fig. 3(b)]. The strikingly different appear-ance in the terrace edge shapes between a fcc(111) filmand a hcp(0001) film is a direct consequence of growthanisotropy. The step bunching and smooth terrace edgesare indeed observed experimentally, as is shown in Fig. 4.Double bilayer steps have emerged at the expense ofsingle bilayer ones. Some step edges are extremelysmooth, with dimensions as long as 600 Å.Lowering the temperature leads to a growth modechange from step-flow to 2D nucleation. This is due to thereduction of diffusion length of adatoms on the surface.As explained before, the anisotropic growth speeds oftype A and B edges will result in triangularly shapedislands. Figure 5 shows a STM picture of GaN(0001)grown at 500 –C on top of the 650 –C buffer. The growthis monitored by RHEED intensity oscillations [8], andis stopped at the trough of the 5th oscillation. Theobservation of RHEED intensity oscillations indicatesFIG. 4. STM picture of GaN step-flow growth mode at 650 –Cshowing step bunching. Other conditions are the same asFig. 1.island nucleation growth, which is confirmed by the STMimage of Fig. 5 showing islands on terraces. As expected,the islands having the triangular shape are observed.Measurement of height differences by STM shows thatthe terraces have single bilayer height differences andthe triangular islands have single bilayer heights on theterraces. However, we also observe triangular islands ofthe same height on the same terrace which are oriented180– from each other. For single bilayer height hcpislands on terraces separated by a bilayer step height,their orientations should be 180– rotated, as is clear fromthe atomic model shown in Fig. 2(b). In order for 180–-rotated triangular islands to appear on the same terrace,they must be either: (i) hcp islands with a bilayer heightdifference, or (ii) the coexistence of hcp and fcc islandsof the same height. From the STM picture, we canrule out case (i) and conclude that the island growthmode has resulted in a mixture of hcp and fcc orientedislands. We have checked this with cross-sectional highresolution transmission electron microscopy (HRTEM)and the TEM image confirms the presence of cubic phaseregions embedded in the bulk hexagonal phase.We now comment on the assumption that the growthspeed of a type A edge is greater. For a hcp(0001) film,because of the alternating stacking sequence, the obser-vation of jagged step edges or triangular island verticespointing in the f1 0 1 0g direction does not prove that atype A edge grows faster. However, if triangular islandsare observed on a fcc(111) film and if the vertex direc-tion is identified, then the observation will prove whichstep edge grows faster. Avery et al. [11] have recentlyshown STM pictures which include that of island growthon a GaAs(111) film. Although these authors did not com-ment on the island shape, it is clear from their data thatapart from some roughness, most islands are indeed tri-angularly shaped with vertices pointing in the f1 1 2g andFIG. 5. STM picture of island nucleation of GaN at 500 –C.Other conditions are the same as Fig. 1.2751VOLUME 82, NUMBER 13 P HY S I CA L REV I EW LE T T ER S 29 MARCH 1999equivalent directions. Their data confirm that the assump-tion made here that type A edges grow faster is correct.Furthermore, as expected from our model for a cubic film,the islands imaged by Avery et al. [11] on different ter-races of the GaAs(111) film all have the same orientation.To address the issue of whether the edge-growth anisotropyis due to adatom/edge-atom bonding or subsurface layerinfluence on the adatom mobility, we observe that althougha hcp island and fcc island on a given terrace have differ-ent subsurface stacking geometries, the islands neverthe-less have similar sizes and qualitatively similar triangularshapes (but 180– misoriented). This observation suggeststhat the growth anisotropy in hcp and fcc GaN films isdominated by anisotropic edge accommodation, similar tothat found byMo et al. [6] and others [12] for homoepitaxyon Si(001). The same mechanism has been attributed tocausing step bunching in polytype SiC(0001) growth [13].It seems we now have evidence that growth anisotropy insemiconductors is mainly due to anisotropic adatom ac-commodation at an edge, a process different from what isfound in metal homoepitaxy, where the dominant effectis due to the anisotropic subsurface induced diffusion rate[7,14].Instead of a conclusion, we offer a discussion ofthe implications of the above results. In the 2D islandgrowth mode, we have observed the coexistence of bothhcp and fcc islands on the same terrace. Eventually,the islands will coalesce. Coalesced islands can easilybe seen in Fig. 5. When this happens, the coalescedislands will inevitably create line or plane defects due totheir boundaries’ mismatch. The resulting film will beinferior in structural order. If, on the other hand, onegrows in the step-flow mode, formation of the mixedphases can be suppressed by the atomic structure at thestep. The STM pictures in Figs. 1 and 4 show thatthe entire 1000 3 1000 Å2 area is of the hcp phase.Adatoms, once incorporated, will follow this structure,which will lead to a pure phase and a more uniformfilm. Lattice mismatch between substrate and film createsmisfit dislocations with edge components which runparallel to the interface. These dislocations are usuallyconnected to threading arms with vertical components.Most threading arms cross the entire thickness of thefilm and are harmful to the quality of the film. If athreading dislocation is of the screw type with Burgersvector b ­ c ­ 5.18 Å, it can be an effective nucleationcenter which causes double spiral growth. The edgesof the spirals form single bilayer steps as shown inFig. 1. Thus, a strategy to reduce the number of screwnucleation centers is to maintain the step flow growthfrom a vicinal substrate. With a large enough diffusionlength, atom incorporation at step edges dominates andnucleation centers on the terraces themselves are greatlysuppressed. The step flow growth may also lengthenparallel-running edge dislocation lines, resulting in areduction in the number of threading dislocations with2752vertical components. Thus, the quality of the film farfrom the interface is improved [15]. Another long-distance vertical defect is inversion domain boundaries(IDBs) [16]. To minimize IDBs, the starting substrateshould have atomically smooth terraces with the sameelemental termination. Because 4H or 6H SiC(0001) isnonisomorphic with 2H-GaNs0001d, stacking mismatchboundaries (SMBs) or double positioning boundaries(DPBs) [17,18] will occur at step edges between thesubstrate and the film. However, SMBs (or DPBs)are short-distance vertical defects because they can becured by stacking faults in the nitride film; thus, awayfrom the interface, the film quality is good. Details ofmorphological studies of GaN films grown in the step flowmode on vicinal SiC(0001) substrates with atomicallysmooth terraces will be reported elsewhere [15].This work is supported in part by HK RGC GrantsNo. HKU 7118/98P, No. 7117/98P, No. HKU 260/95P,U.S. DOE Grant No. DE-FG02-84ER 45076, and NSFGrant No. DMR-9214054.[1] S. Strite and H. Morkoc, J. Vac. Sci. Technol. B 10, 1237(1992).[2] S. Nakamura, M. Senoh, S.-I. Nagahama, N. Iwasa,T. Yamada, T. Matsushita, Y. Sugimoto, and H. Kiyoku,Appl. Phys. Lett. 70, 1471 (1997).[3] S. Nakamura, Jpn. J. Appl. Phys. 30, L1705 (1991).[4] A. R. Smith, R.M. Feenstra, D.W. Greve, J. Neugebauer,and J. E. Northrup, Phys. Rev. Lett. 79, 3934 (1997).[5] A. Pavlovska, E. Bauer, V.M. Torres, J. L. Edwards,R. B. Doak, I. S. T. Tsong, V. Ramachandran, and R.M.Feenstra, J. Cryst. Growth 189/190, 310 (1998).[6] Y.-W. Mo, B. S. Swartzentruber, R. Kariotis, M. B. Webb,and M.G. Lagally, Phys. Rev. Lett. 63, 239 (1989).[7] T. Michely, M. Hogage, M. Bott, and G. Comsa, Phys.Rev. Lett. 70, 3943 (1993).[8] M.H. Xie, S.M. Seutter, W.K. Zhu, L. X. Zheng,Huasheng Wu, and S. Y. Tong (to be published).[9] Y. Nakada and H. Okumura, J. Cryst. Growth 189/190,370 (1998).[10] A. R. Smith, R.M. Feenstra, D.W. Greve, M. S. Shin,M. Skowronski, J. Neugebaner, and J. E. Northrup, Appl.Phys. Lett. 72, 2114 (1998).[11] A. R. Avery, H. T. Dobbs, D.M. Holmes, B. A. Joyce, andD.D. Vvedensky, Phys. Rev. Lett. 79, 3938 (1997).[12] J. Y. Tsao, E. Chason, U. Koehler, and R. Hamers, Phys.Rev. B 40, 11 951 (1989).[13] T. Kimoto, A. Itoh, H. Matsunami, and T. Okano, J. Appl.Phys. 81, 3494 (1997).[14] R. T. Tung and W.R. Graham, Surf. Sci. 97, 73 (1980).[15] M.H. Xie, L.X. Zheng, C. H. Pang, Huasheng Wu,N. Ohtani, and S. Y. Tong (to be published).[16] J. E. Northrup, J. Neugebauer, and L. T. Romano, Phys.Rev. Lett. 77, 103 (1996).[17] B.N. Sverdlov, G.A. Martin, H. Morkoc, and D. J. Smith,Appl. Phys. Lett. 67, 2063 (1995).[18] S. Tanaka, R. S. Kern, and R. F. Davis, Appl. Phys. Lett.66, 37 (1995).

Anisotropic step-flow growth and island growth of GaN(0001) by molecular beam epitaxy

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Anisotropic step-flow growth and island growth of GaN(0001) by molecular beam epitaxy

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