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Applied Surface Science 256 (2010) 4745–4756

Contents lists available at ScienceDirect

Applied Surface Science

journa l homepage: www.e lsev ier .com/ locate /apsusc

urface treatments toward obtaining clean GaN(0 0 0 1) from commercial hydrideapor phase epitaxy and metal-organic chemical vapor deposition substrates inltrahigh vacuum

zusa N. Hattori a,∗, Katsuyoshi Endoa, Ken Hattorib, Hiroshi Daimonb

Research Center for Ultra-Precision Science and Technology, Graduate School of Engineering, Osaka University, Yamada-oka 2-1, Suita, Osaka 565-0871, JapanGraduate School of Materials Science, Nara Institute of Science and Technology, Takayama 8916-5, Ikoma, Nara 630-0192, Japan

r t i c l e i n f o

rticle history:eceived 5 October 2009ccepted 1 March 2010vailable online 10 March 2010

ACS:8.37.Ef8.47.Fg8.43.Vx8.55.ag

eywords:aNcanning tunneling microscopyow-energy electron diffractioneflection high-energy electron diffractionurface structure

a b s t r a c t

We studied processes of cleaning GaN(0 0 0 1) surfaces on four different types of wafers: two types werehydride vapor phase epitaxy (HVPE) free-standing substrates and two types were metal-organic chemicalvapor deposition (MOCVD) films grown on these HVPE substrates and prepared by annealing and/or Ar ionsputtering in ultra high vacuum. We observed the surfaces through treatments using in situ low-energyelectron diffraction (LEED), reflection high-energy electron diffraction (RHEED), scanning tunnelingmicroscopy (STM), and Auger electron spectroscopy, and also using ex situ temperature programmeddesorption, X-ray photoelectron spectroscopy, X-ray diffraction, and secondary ion mass spectrometry.For HVPE samples, we obtained relatively clean surfaces under optimized three-step annealing con-ditions (200 ◦C for 12 h + 400 ◦C for 1 h + 500 ◦C for 5 min) without sputtering, after which the surfacecontamination of oxide and carbide was reduced to ∼20% of that before annealing. Clear GaN(0 0 0 1)1 × 1patterns were obtained by LEED and RHEED. STM images showed flat terraces of ∼10 nm size and stepsof ∼0.5 nm height. Upon annealing the HVPE-GaN samples at a much higher temperature (> 550 ◦C),three-dimensional (3D) islands with facets were formed and the surface stoichiometry was broken downwith the desorption of nitrogen in the form of ammonia, since the samples include hydrogen as an impu-rity. Ar+ sputtering was effective for removing surface contamination, however, postannealing could not

urface preparation recover the surface roughness but promoted the formation of 3D islands on the surface. For MOCVD/HVPEhomoepitaxial samples, the surfaces are terminated by hydrogen and the as-introduced samples showeda clear 1 × 1 structure. Upon annealing at 500–600 ◦C, the surface hydrogen was removed and a 3 × 3reconstruction structure partially appeared, although a 1 × 1 structure was dominant. We summarize thestructure differences among the samples under the same treatment and clarify the effect of crystal qual-ity, such as dislocations, the concentration of hydrogen impurities, and the residual reactant molecules

ce str

in GaN films, on the surfa

. Introduction

Gallium nitride (GaN) has excellent properties such as a directnd wide band-gap energy of 3.4 eV at room temperature (RT) in theexagonal phase (wurtzite, a = 0.319 nm and c = 0.519 nm) [1,2].

t also has high electron mobility [3], and has thus attracted muchttention for its potential use in a wide range of electronic devices.or device fabrication, various materials have been integrated onaN substrate surfaces, where the initial GaN surface itself often

∗ Corresponding author. Present address: Nanoscience and Nanotechnology Cen-er, The Institute of Scientific and Industrial Research, Osaka University, 8-1,

ihoga-oka, Ibaraki, Osaka 567-0047, Japan. Tel.: +81 6 6879 4281;ax: +81 6 6879 4283.

E-mail address: a-hattori@sanken.osaka-u.ac.jp (A.N. Hattori).

169-4332/$ – see front matter © 2010 Elsevier B.V. All rights reserved.oi:10.1016/j.apsusc.2010.03.001

ucture.© 2010 Elsevier B.V. All rights reserved.

plays a crucial role in the device operation. It is essential to expli-cate and control the properties of the surface of the GaN substrate.A number of groups have investigated procedures for cleaning GaNsurfaces for device fabrication [4,5]. There is currently, however,no standard method of preparing clean GaN substrates either bythe use of chemical solutions or by treatments in ultra high vac-uum (UHV), and the obtained results have not been reproduced bycommercially supplied samples.

The structures of clean surfaces are of particular importancesince knowledge of the structures is the first step in understandingthe fundamental issues of contact formation, chemical reactivity,

growth processes, and so on. The surface structures of as-grownsamples of Ga-polar GaN(0 0 0 1) grown by molecular beam epi-taxy (MBE) and metal-organic chemical vapor deposition (MOCVD)on substrates of Si-polar SiC(0 0 0 1), sapphire(0 0 0 1), Si(1 1 1), andso on [6–20] have been studied in situ. For the MBE growth of a

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746 A.N. Hattori et al. / Applied Sur

aN(0 0 0 1) surface under a Ga-rich condition, 2 × 2, 2 × 3, 3 × 2,× 4, 4 × 4, 5 × 5, 10 × 10, 5

√3 × 2

√13, 5 × 2.5,

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3 ×3 reconstruction structures have been reported [6–10]. In gen-

ral, the unintentional presence of arsenic on the surface leads tohese MBE reconstruction structures [11–13,16]. On samples growny MOCVD, an ordered 1 × 1 surface with a nearly one-to-onetoichiometric composition has been reported [19,20]. Some theo-etical studies have predicted that a GaN(0 0 0 1) surface exhibits

2 × 2 reconstruction structure under a Ga-rich condition [21]nd the minimum-energy configuration of GaN(0 0 0 1)2 × 2 haseen proposed [22,23]. Indeed, GaN(0 0 0 1)2 × 2 reconstructionurfaces have been reported experimentally on MBE samples underitrogen-rich conditions [24–27]. These studies were performed

n situ; films were fabricated by controlling the stoichiometry inHV, where the surfaces were subsequently observed. It is essen-

ial for device technology to study and develop surface-cleaningrocedures for ex situ prepared GaN substrate samples.

In general, commercially available GaN sample surfaces areurface-lapped (mechanically polished) and finished by chemicalechanical polishing (CMP), dry etching, and so on, after fabrica-

ion and subsequent exposure to air. For commercially availableamples, the main fabrication method is hydride vapor phasepitaxy (HVPE). We can also obtain MOCVD GaN film samplesy homoepitaxial growth on GaN substrates and heteroepitaxialrowth on other substrates. To obtain atomically clean surfaces, it isecessary to understand the phenomena involved in the processesf cleaning surfaces through chemical, thermal, gaseous reactionsdry etching), and so on, and to effectively combine such processes.

GaN crystal quality has been improved in recent years ands expected to improve further. GaN crystallinity (such as defectensity) and/or properties strongly depend on the growth condi-ions. In general, HVPE GaN has a threading dislocation densityf about 106–107 cm−2[28,29], while that of MOCVD heteroepi-axial GaN is 107–108 cm−2[30]. Currently, HVPE growth [31] is aery popular method of fabricating both crystalline substrates andomoepitaxial layers because of the high growth rates of around00 �m/h. Also many methods of stably fabricating larger andigher-quality GaN crystals at a low price by HVPE growth haveeen patented. Of course, dislocations severely influence deviceerformance [32–34], though GaN crystals have been realized asroducts in the laser market. To achieve electronic devices withxcellent properties, for instance, carrier transport, we shouldmprove the crystal quality, and surface-finishing methods, andnderstand the local (microscopic) surface structure that becomeshe interface under electrode materials. A dislocation density of lesshan 103 cm−2 is required for laser diode applications to obtain aigher yield and cost-effectiveness. Many researchers have tried to

abricate GaN crystals with lower dislocation densities. Ishii et al.uggested a guideline for the epitaxial growth of GaN(0 0 0 1) from aheoretical viewpoint [35,36]; the key is to introduce gallium atomsefore supplying nitrogen because of the unusual nitrogen kinetics

n the truncated GaN(0 0 0 1) [37].GaN surface structures are affected by impurities in crystals,

n particular, the presence of hydrogen as an impurity has beeneported [27,38–40]. It was shown that H in GaN exhibits uniqueeatures which were not observed in more traditional semicon-uctors such as Si or GaAs [40], for example, the formation of ancceptor–hydrogen complex, especially in GaN:Mg, with p-typeonductivity [41,42]. Also, the presence of ambient H changeshe formation energy of the GaN surface [43] and causes theormation of reconstruction structures [27,38,39]. The hydrogen

mpurity sometimes arises from an impurity component of theource gases, especially ammonia. The hydrogen concentration of019–1023 cm−3 was measured for MBE-GaN by nuclear reactionnalysis, and a strong correlation between hydrogen concentrationnd crystal strain (dislocations) was observed [44]. This correla-

cience 256 (2010) 4745–4756

tion leads to a higher density of the hydrogen impurity in MOCVDfilms than that in HVPE substrates, affecting the dislocation density,when the hydrogen impurity is supplied from the source gases.

Among the cleaning processes, annealing in UHV is one ofthe simplest treatments. However, there have not yet been anysystematic or reproducible studies on the surface structures of free-standing HVPE GaN(0 0 0 1) substrates. Ar ion sputtering is one ofthe general treatments used to remove surface contamination. Itwas reported that Ar+ sputtering is not suitable for obtaining a cleanGaN(0 0 0 1) surface because of the increase of surface roughnessand/or defects.

In this paper, we present the surface structures, chemical states,crystallinities, and morphologies of GaN substrates upon anneal-ing and also upon Ar+ sputtering (in particular, sputtering andpostannealing) in UHV, mainly through evaluation by low-energyelectron diffraction (LEED), reflection high-energy electron diffrac-tion (RHEED), scanning tunneling microscopy (STM), temperatureprogrammed desorption (TPD), and X-ray photoelectron spec-troscopy (XPS). We studied the dependence of sample quality infour different types of wafers: two HVPE freestanding substratesfrom different companies and two homoepitaxial MOCVD GaNfilms deposited on the former substrates. We found that even afterthe same thermal treatment, different structures appeared in eachsample. In the case of Ar+ sputtering, the surface roughness of thesamples was not recovered by postannealing, and instead, three-dimensional (3D) islands with facets appeared depending on theannealing temperature and duration. MOCVD films on HVPE sub-strates showed 3 × 3 structures, probably due to domain inversioncaused by crystal dislocations.

2. Experimental

Four different types of samples were used. Type-1 samples werecut from a wafer of freestanding and mirror-polished GaN(0 0 0 1)grown by HVPE (company A, commercial). Type-2 samples werecut from a type-1 wafer on which a GaN epitaxial film of ∼2 �mthickness had been grown by MOCVD without any polishingtreatment after the epilayer growth (company A, commercial).Type-3 samples were cut from a wafer of HVPE-grown freestandingGaN(0 0 0 1) that had been subjected to mirror-polishing (companyB, commercial). Type-4 samples were cut from a type-3 wafer onwhich a GaN epitaxial film of ∼4 �m thickness had been grown byMOCVD without any polishing treatment after the growth.

The samples were cut to a size of 2 × 7 mm2 from wafers of∼300 �m thickness. The orientation flat directions of the waferswere used as a reference to determine the azimuth directions ofthe sample surfaces. All samples were degassed at 200 ◦C for ∼12 hand annealed in UHV (base pressure: < 1 × 10−8 Pa) at 200–800 ◦Cfor 5 s–12 h, except in the case of measurements on as-introducedsamples. The sample temperature was monitored by two pyrom-eters [45] with emissivity � = 0.93 [46]. Some type-1 and type-3samples were cleaned by Ar ion sputtering with Ar gas at a pres-sure of ∼2 × 10−3 Pa (99.9995%) for 1–30 min at an incident energyof 0.5 keV, an incident angle of ∼50 ◦, and a beam current densityof 1.0–1.5 �A/cm2, with subsequent annealing. In situ LEED, STM,RHEED, and Auger electron spectroscopy (AES) measurements, andex situ XPS, TPD, X-ray diffraction (XRD), and secondary ion massspectrometry (SIMS) were performed. Three different UHV systemswere used for the in situ measurements. The main system [47] wasequipped with an Ar ion sputtering gun (PSP Technology, ISIS3000),

an RT-STM (JEOL, JSTM4500-XT), and LEED optics (WA Technology,MicroLEED). RHEED and LEED/AES measurements were performedin the two other UHV systems [48,49]. The base pressure in XPS(ULVAC-PHI, Quantum2000) and TPD (ESCO, EMD-WA1000S/W)was lower than 1 × 10−7 Pa. �–2� XRD (Cu K˛1) measurements were

A.N. Hattori et al. / Applied Surface Science 256 (2010) 4745–4756 4747

Fig. 1. LEED and RHEED patterns of as-introduced samples: (a) and (e) for type-1 sample (HVPE), (b) and (f) for type-2 sample (MOCVD), (c) and (g) for type-3 sample (HVPE)and (d) and (h) for type-4 sample (MOCVD), respectively. Note that the 00 beam in LEED was not at the screen center. The incident electron direction was [1 1̄ 0 0] in RHEED.E dicatr st of tt r in th

pm

aiweas1mpastwtce

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3

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p in LEED was 95 eV. In RHEED, Laue zones (L0 and L1) and transmission spots are ineciprocal lattice units corresponding to a and c, respectively. Note that the contrahe intensities of (e) and (g) are weak. (For interpretation of the references to colou

erformed for some type-3 and type-4 samples, and SIMS measure-ents were performed for some type-1 and type-3 samples.In STM measurements, topographic images were observed at

sample bias voltage Vs between −5 V and +5 V and tunnel-ng current It = 0.1–0.5 nA. The STM tip was a 0.3 mm� tungsten

ire sharpened by electrolytic etching and cleaned by subsequentlectron-bombardment heating. RHEED patterns were observedt the incident energy of the primary electron, Ep = 15 keV. XPSpectra were obtained using monochromatic Al K˛ X-rays (h� =486.6 eV). The XPS system equipped with Ar+ sputtering equip-ent (0.5 kV, Ar pressure: ∼5 × 10−3 Pa) enabled us to obtain depth

rofiles of XPS spectra. The core-level spectra were obtained withpass energy of 40 eV, an energy resolution of 0.8 eV, and an X-ray

pot size of ∼5 �m2. The binding energy (BE) of the spectra wasaken relative to the valence band maximum, which was measuredith the Al K˛ photon line. The background was subtracted using

he Shirley method [50]. Peaks in each spectrum were fitted by aonvolution of Gaussian and Lorentzian curves. All measurementsxcept TPD were performed at RT.

. Results and discussion

.1. Characterization of as-introduced GaN surfaces

Fig. 1 shows LEED and RHEED patterns of the four differents-introduced samples. While LEED patterns of the type-1 sam-le (Fig. 1(a)) and type-3 sample (Fig. 1(c)) show almost no specificeatures, the type-2 sample (Fig. 1(b)) and type-4 sample (Fig. 1(d))how clear 1 × 1 spots; the homoepitaxial MOCVD surfaces have× 1 surface ordering even after exposure to air, while the HVPE

urfaces do not exhibit such spots. Since it has been reported thatn MOCVD-grown GaN(0 0 0 1) surface is terminated by hydrogen27,51], we consider that the survival of the 1 × 1 ordering on the

OCVD samples is caused by an inert H-termination cap, similar tohe case of Si(0 0 1)–H surfaces [52]. Note that LEED in this Ep rangeeflects the ordering of a few surface layers. The slightly higher

ackground seen in Fig. 1(b) and (d) implies slight pollution of theurface by exposure to air. In contrast, the commercial HVPE GaNamples show no specific LEED spots. For the commercial HVPEaN substrates, surface-finishing treatments (CMP, etc.) were per-

ormed after lapping to obtain mirror-polished smooth surfaces,

ed by pink dashed arcs and cyan circles, respectively. The scales of a∗ and c∗ are thehe RHEED patterns (e) and (g) is enhanced compared with that of (f) and (h) sinceis figure legend, the reader is referred to the web version of the article.)

where the presence of a surface oxide layer was reported [53]. Thus,surface crystallinity and ordering were clearly destroyed in at leasta few surface layers by the surface-finishing treatments.

In the RHEED patterns for type-1 (Fig. 1(e)) and type-3 (Fig. 1(g))samples, where the electron mean free path at this Ep is ∼20 nm[54], we observed weak transmission diffraction spots (cyan circles)in addition to clear 1 × 1 surface spots (on dashed arcs, correspond-ing to Laue zones). The transmission spots indicate the existence of3D islands of GaN(0 0 0 1) on the as-introduced GaN(0 0 0 1) sur-faces. The main purpose of the surface-finishing treatments is toremove damage by lapping and to achieve surface flattening with-out damage. However, no effective polishing method or etchanthas yet been confirmed for GaN substrates, although macroscopicsurface roughness has been improved [55,56]. Hanser et al. mea-sured the defect density of their original HVPE-GaN crystals treatedby CMP using cathode-luminescence imaging, and reported that nodefects related to polishing were observed [56]. In contrast, Murataevaluated the thickness of damage (depth profile) in commercialHVPE GaN substrates treated by CMP by measuring photocurrentdensity. They concluded that the commercial HVPE GaN substratehas a damaged layer with a thickness of a few hundred nm andmany scratches exist on the surface [57].

On the other hand, in the RHEED patterns of the MOCVD film sur-faces (Fig. 1(f) and (h)), 1 × 1 surface-ordering patterns and sharpKikuchi lines without any transmission spots can be observed. Theintensity of Kikuchi lines strongly depends on the surface crys-tallinity and morphology; sharper Kikuchi lines are observed fromcrystals with perfect surfaces and perfect bulk lattices [54]. Fromthis viewpoint, the surface quality and crystallinity of type-2 andtype-4 samples were higher than those of type-1 and type-3 sam-ples. This is consistent with the surface-sensitive LEED results.We emphasize that evaluation by electron diffraction, especiallyRHEED, is useful and convenient for confirming and optimizingsurface-finishing treatments to obtain samples with high crys-tallinity as well as for the evaluation of surface roughness bymicroscopy.

Fig. 2(a) and (c) show the core-level spectra of Ga 3d5/2,3/2 and(b) and (d) show those of N 1s for (a) and (b) an as-introducedtype-1 sample (HVPE), and (c) and (d) an as-introduced type-2sample (MOCVD) obtained by XPS. The Ga 3d spectra are asym-metric and can be fitted with four symmetric Voigt components

4748 A.N. Hattori et al. / Applied Surface Science 256 (2010) 4745–4756

Fig. 2. Core-level XPS spectra of (a) and (c) gallium 3d, and (b) and (d) nitrogen 1s: (a) and (b) for an as-introduced type-1 (HVPE) sample, and (c) and (d) for an as-introducedt gt fun( es indb ns oft e.)

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ype-2 (MOCVD) sample. Ga 3d and N 1s spectra (black curves) are fitted with VoiNGa, NH2 , NH3 , and two Ga LMM Auger satellites), respectively. The red dashed linending between the type-1 and type-2 samples results in the different peak positioo colour in this figure legend, the reader is referred to the web version of the articl

ith different chemical shifts corresponding to Ga bonding to nitro-en: GaN (BE = 19.2–20.3 eV) [58], Ga bonding to oxygen: GaOx

BE = 19.6–21.0 eV) [59], Ga bonding to N–H: GaH−N (BE = 19.21 eV)60], and metal Ga (BE = 18.4–18.49 eV) [61]. It is known that hydro-en is unintentionally doped in GaN and hydrogen behavior in GaNlms has been investigated by Fourier transform infrared (FTIR)bsorption [60], electron-energy-loss spectroscopy [62], and son. Kong et al. confirmed the stretch mode of the N–H bond byTIR and denoted the hydrogen-related complex as Ga· · · H–N [60].hey also assigned a Ga· · · H–N peak in the Ga 3d spectra in XPS60]. Note that the energy difference between Ga 3d3/2 and Gad5/2 is less than 0.1 eV [63], negligible compared with our res-lution. Comparing the Ga 3d spectra for the MOCVD (Fig. 2(c)) and

VPE (Fig. 2(a)) specimens, markedly different features, i.e., a domi-ant Ga· · · H–N peak and negligible GaOx and metal-Ga peaks, werebserved on the MOCVD sample. The high ratio of the Ga· · · H–Nomponent on the MOCVD sample should be caused by the higherensity of the hydrogen impurity in the MOCVD sample than in the

able 1tomic percentages of surface species, Ga, N, O, C, and Cl, for the different sample types (1–4y XPS intensities and sensitivity factors for Ga 3d, N 1s, O 1s, C 1s, and Cl 2p.

Treatment Type n Ga (%)

GaN GaH−N GaOx Metal

As-introduced 1 3 16 9 13 2Sputtering 1 2 26 20 – –Annealing 1 1 18 15 8 1As-introduced 2 1 6 31 1 <1As-introduced 3 1 13 19 3 2As-introduced 4 1 8 25 1 <1

ere, n denotes the number of measured samples. As shown in Fig. 2, we can estimate thehose of NGa, NH2 , and NH3 from N 1s spectra. Here the sum of Ga, N, O, C, and Cl atomic%nnealing (200 ◦ C for 12 h + 400 ◦ C for 1 h + 500 ◦ C for 5 min). The Ar+ sputtering time is 2fter Ar+ sputtering. Note that all values have an error of about 20%.

ctions of four components (GaN, GaOx , GaH−N, and metal-Ga) and five componentsicate the sum of the component peaks. Note that the slight difference of the bandGa LMM Auger peaks in the BE representation. (For interpretation of the references

HVPE sample [44], in addition to hydrogen termination [27,51], asdescribed before. We suppose that the reduction of GaOx in thetype-2 sample is caused by the termination of surface hydrogen,which prevents surface oxidation even in air. For the enhancementof the GaOx and metal-Ga peaks in the type-1 sample, we considerthat CMP treatment after film fabrication destroys the surface ter-mination, promotes oxidation, and induces the formation of Gametal impurities and/or segregation. The surface oxidation [64] andthe existence of metal impurities [65] after CMP have been pointedout to be general phenomena. Upon comparing type-3 and type-4 samples, we found a small GaOx peak and negligible metal-Gapeaks in the type-4 sample. These results also indicate that CMPtreatment has a severe effect on the surface structures and condi-

tion. Note that the existence of Ga oxide layers with a thickness of∼1.5 nm was reported for HPVE samples [53].

The N 1s spectra can be fitted with five components [66,60]: N 1score electrons bonding to hydrogen: NH3 (BE = 405.6–406.2 eV) [63]and NH2 (BE = 397.7–399.72 eV) [66], an N 1s core-electron bond-

) and treatments (as-introduced, Ar+ sputtering, and 600 ◦ C annealing), determined

N (%) O C Cl

Sum NGa NH2 NH3 Sum (%) (%) (%)

40 15 22 6 43 5 11 < 146 22 25 4 51 2 – < 142 21 14 3 38 11 8 < 139 17 22 4 44 6 11 –37 20 20 3 43 11 8 < 134 25 19 4 48 8 10 –

component intensities of GaN, GaOx , GaH−N, and metal-Ga from Ga 3d spectra, andis fixed to 100%. The annealing is performed at 600 ◦ C for 10 min after three-step0 min, as shown in Fig. 8(a). No carbon was detected in the XPS spectra of samples

face Science 256 (2010) 4745–4756 4749

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Fig. 3. Schematic phase diagrams for (a) HVPE samples (type-1 and type-3) treatedby annealing, (b) HVPE samples (type-1 and type-3) treated by Ar+ sputtering and

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A.N. Hattori et al. / Applied Sur

ng to gallium: NGa (BE = 396.2–397.86 eV) [66,59], and two Ga LMMuger electrons with satellites [61]. There is no significant differ-nce of N 1s spectra between the HVPE and MOCVD samples. Thempurity components of O 1s, C 1s, and Cl 2p were observed byPS. From the intensities of these components, we can estimate

heir atomic percentages (atomic%) in the surface region using thestablished photoelectron cross-sections and mean free paths [63].able 1 summarizes the atomic% of each component for the differ-nt sample types and treatments. Here, the sum of Ga, N, O, C, andl atomic% is fixed to 100%. We can confirm that the as-introducedype-1 samples include non-negligible amounts of O, C, and Cl asmpurities.

.2. Schematic phase diagrams

Fig. 3 shows schematic phase diagrams of (a) HVPE substratestype-1 and type-3) after annealing, (b) HVPE substrates (type-1nd type-3) after Ar+ sputtering and subsequent annealing, and (c)omo epitaxial MOCVD films on HVPE substrates (type-2 and type-) after annealing, determined by LEED, RHEED, XPS, AES, TPD, andTM measurements. Relatively clean surfaces were obtained bynnealing without Ar ion sputtering for both HVPE and MOCVDamples (Fig. 3(a) and (c)); 1 × 1 structures exist on the HVPEurfaces and 1 × 1 and 3 × 3 structures coexist on homoepitax-

al MOCVD surfaces after annealing at approximately 500–600 ◦C.n contrast, on Ar+-sputtered surfaces, the surface morphologyould not be recovered by postannealing, although surface car-on contamination and oxide layers were almost removed. For allreatments (Fig. 3(a)–(d)) at a higher annealing temperature (≥

ig. 4. TPD curves of major mass numbers detected for (a) type-1 (HVPE) and (c) type-2 (Ma) and (c), respectively, considering the cracking elements and ratio. The temperature were similar to those for type-1 and type-2 samples, respectively.

subsequent annealing in UHV, and (c) homoepitaxial MOCVD samples (type-2 andtype-4) treated by annealing. Nitrogen sublimates in two forms: as ammonia around300–700 ◦C and as molecular nitrogen above 800 ◦C. The phases were observed atRT after annealing.

500–600 ◦C), nitrogen sublimation in the form of ammonia startedto occur and 3D islands with facets formed. Finally, samples weredamaged after annealing above 800 ◦C. In the following sections, wediscuss the relation between the structure, treatment, and sampletype.

3.3. Desorption from HVPE and MOCVD samples

Fig. 4(a) and (c) show the typical major TPD curves ofas-introduced type-1 and type-2 samples, respectively. The tem-

OCVD) samples. (b) and (d) show TPD spectra of desorption species estimated fromas ramped up at a rate of 0.5 K/s. The TPD curves from type-3 and type-4 samples

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erature was ramped up at a rate of 0.5 K/s. Similar TPD curveso those in Fig. 4(a) and (c) were observed for type-3 and type-4amples, respectively. TPD intensities were normalized by sampleeight; the intensities in Fig. 4(a) and (c) are comparable [67]. As

hown in the figure, the major detected values of m/e are 2, 12, 14,6, 17, 18, 28, 35, 36, 69, and 104 in atomic units (AMU), whichan be assigned to H2

+, C+, N+/CH2+, O+/NH2

+, OH+/NH3+, H2O+,

2+/CO+, Cl+, HCl+, Ga+, and GaCl+, respectively. Thus, the main

esorption species are considered to be H2, CH4, NH3, H2O, N2, CO,H)Cl, Ga, and GaCl. Here, we discuss the desorption species of H2,H4, H2O, N2, Ga, GaCl, Cl, and NH3, considering the cracking ratio.

Fig. 4(b) and (d) show the desorption curves of these speciesstimated from Fig. 4(a) and (c), respectively. H2O peaks below00 ◦C, N2 peaks above 800 ◦C, and broad H2 peaks from 400 to00 ◦ C can be seen for both samples. The origin of the H2O desorp-ion should be physical adsorption on the surface from humidity inhe air. There are distinctive differences in the NH3, CH4, Ga, andl peaks between the HVPE (type-1) and homoepitaxial MOCVDtype-2) samples. For the HVPE type-1 samples, m/e = 35 (Cl+), 36HCl+), and 104 (GaCl+) peaks were clearly observed, whereas theirntensities were negligible for the homoepitaxial MOCVD type-2amples. In the HVPE method, GaCl and NH3 react to produce GaN,ith H2 and HCl gases as by-products: GaCl + NH3 → GaN + H2 + HCl

31]. The desorption of GaCl, (H)Cl, NH3, and H2 in type-1 sam-les strongly indicates that they originated from unreacted sourceaterials and/or residual product materials. This implies that theVPE growth conditions are not yet optimized. Moreover, weonfirmed by XPS that type-1 and type-3 samples (HVPE GaN) con-ained Cl element as an impurity, while type-2 and type-4 sampleshomoepitaxial MOCVD GaN) did not show Cl peaks (Table 1). In thePD of type-2 and type-4 samples, no GaCl desorption and littlel desorption were detected. In the homoepitaxial MOCVD sam-les, we consider that the small amount of Cl desorption shouldriginate from the HVPE substrate under the MOCVD film. In theOCVD process, the gallium and nitrogen sources are usuallya(CH3)3 and NH3, respectively, and the products are GaN and CH4:a(CH3)3 + NH3 → GaN + 3CH4[68]. Indeed, we can see intense CH4esorption from type-2 and type-4 samples rather than type-1 andype-3 samples; the CH4 species is one of the by-product elementsn the MOCVD film and/or from surface carbon contamination. No

ethyl gallium (Ga(CH3)n+, n = 1–3) species were detected from

ny samples.For both sample types, there are broad H2 peaks in the tempera-

ure range of 400–800 ◦C (Fig. 4(b) and (d)). Hydrogen is a commonmpurity in many semiconductors, and it was reported that GaNlso contains hydrogen as an impurity [40,69]. Indeed, we con-rmed the existence of hydrogen in type-1 and type-3 samplessing SIMS. From the TPD results for type-2 samples (Fig. 4(d)),e point out the possibility of hydrogen desorption from the HVPE

ubstrate under the MOCVD film, similarly to the Cl desorption.lthough the hydrogen density is higher in the MOCVD film than

hat in the HVPE substrate, the total amount of hydrogen in the des-rption region should be smaller in the MOCVD film that that in theVPE substrate, as shown in the H2 intensity difference betweenig. 4(a) and (d).

The TPD results (Fig. 4(b) and (d)) show that NH3 desorptiontarted at ∼400 and ∼200 ◦C for the HVPE and homoepitaxialOCVD samples, respectively. A small amount of Ga desorption

lso occurs at ∼500 ◦ C for both samples. When the type-1 sam-le was subjected to annealing at 600 ◦C, according to the resultsf XPS the atomic weight ratio between Ga and N was 1.1 ± 0.2,

hile this value for an as-introduced type-1 sample was 0.9 ± 0.2

Table 1); the annealing results in a slightly Ga-rich or N-defectiveurface mainly due to NH3 desorption. There are two candidatesor the origin of ammonia desorption: one is unreacted ammoniatself, and the other is residual hydrogen leading to the production

cience 256 (2010) 4745–4756

of ammonia. The former would result in isolated NH3 desorption,which would correspond to the peak at ∼200 ◦ C obtained from theMOCVD samples in Fig. 4(d)), and the latter would result in thecoincident desorption of NH3 and H2.

Fig. 4(b) for a type-1 (HVPE) sample indicates the coincidentdesorption of NH3 and H2 at above 350 ◦C. This should originatefrom the reaction between GaN bulk and residual hydrogen. Theweak coincidence of NH3 and H2 desorption at the same temper-ature for the type-2 (MOCVD) sample suggests the same reaction,but with hydrogen produced from the HVPE substrate under theMOCVD film, as mentioned. Monatomic interstitial hydrogen hasbeen predicted to exist in two charge states in GaN [40,69–71]:H+ and H−, whereas the H0 state is unstable. H+ preferentiallyexists at the N antibonding site; the calculated N–H bond lengthis 0.102–0.104 nm, which is close to that in NH3 (0.104 nm [72]),while H− and H0 preferentially exist at the Ga antibonding site [69].The diffusion barrier for H+ is ∼0.7 eV while that for H− is ∼3.4 eV,implying mobile H+ slightly above RT and immobile H−[40]. ThisN–H+ configuration is a possible precursor state for NH3 dissocia-tion during the GaN fabrication and in NH3 desorption from GaN ata higher substrate temperature. Actually, Table 1 shows a decreaseof the N 1s component bonded to hydrogen after annealing at 600 ◦

C for HVPE (type-1) samples.Differences in the crystal quality would induce differences of the

residual NH3 desorption intensity at ∼200 ◦ C between type-1 andtype-3 (HVPE) samples and between type-2 and type-4 (homoepi-taxial MOCVD) samples. In general, HVPE GaN crystals have athreading dislocation density of about 106–107 cm−2[28,29], whichis less than that in MOCVD heteroepitaxial GaN, 107–108 cm−2[30],as mentioned before. Though we have no independent proof ofthreading dislocation density differences between HVPE GaN andMOCVD GaN, in the �–2� XRD spectra for type-3 and type-4 sam-ples, the (0 0 0 2) peak showed a wider full width at half maximum(FWHM) in MOCVD GaN than that in HVPE GaN (540 vs 288 arcsec)for 2�, corresponding to 60 and 110 nm, respectively, in crystallinedomain size. This suggests that the NH3 desorption at ∼200 ◦ C fromthe homoepitaxial MOCVD type-2 sample is caused by the lowercrystallinity (larger threading dislocation density), in addition tothe existence of the residual NH3 in the MOCVD films. We shouldnote here that the HVPE samples showed excellent crystallinity inthe bulk according to the results of XRD but lower atomic order-ing in a few surface layers compared with the MOCVD samplesaccording to the results of LEED.

The phase diagrams (Fig. 3) show the formation of 3D islandswith facets in the end stage of NH3 desorption on HVPE and MOCVDsamples upon annealing above 550–600 ◦C. Above 800 ◦C, crystaldisaggregation started; N2 sublimation occurred (Fig. 4) and thesamples broke into bits. Therefore, to obtain flat and stoichiomet-ric GaN surfaces, the annealing temperature should be below 550 ◦

C for both HVPE and MOCVD samples. However, 550 ◦ C is not suffi-ciently high to remove surface contamination or to recover surfaceroughness after sputtering, as shown in the following subsections.

3.4. Annealing of HVPE samples

Fig. 5(a) shows an LEED pattern and a typical STM image of anHVPE type-1 sample annealed (degassed) at 200 ◦ C for 12 h. LEEDpatterns showed a diffused 1 × 1 structure and STM indicated thatmany clusters remained on the surface. These clusters are ascribedto the formation of oxide/carbide layers [59]. The peak-to-valley(PV) height and root-mean-square (RMS) roughness were 2.5 and

0.86 nm, respectively. We found that three-step annealing resultsin relatively clean and flat surfaces. The established conditions foreach step were (1) 200 ◦C for 12 h, (2) 400 ◦C for 1 h, and (3) 500 ◦Cfor 5 min. After three-step annealing, the sharpest 1 × 1 LEED pat-terns appeared and STM images showed wide and flat terraces of

A.N. Hattori et al. / Applied Surface Science 256 (2010) 4745–4756 4751

Fig. 5. LEED patterns, STM images, and AES spectra of type-1 (HVPE) samples. (a) Typical STM image and LEED pattern (inset) after degassing at 200 ◦ C for 12 h. (b) Cross-sectional profile through the dashed line in (a); PV and RMS were 2.5 and 0.86 nm, respectively. (c) and (e) Typical STM images and LEED pattern (in inset of (c)) afterthree-step annealing (200 ◦ C for 12 h, 400 ◦ C for 1 h, and 500 ◦ C for 10 min). (d) and (f) Cross-sectional profiles through the dashed lines in (c) and (e), respectively. PVand RMS in (d) were 1.7 and 0.39 nm, respectively. The observed conditions were (a) Ep = 85 eV, Vs = −4.8 V, It = 0.3 nA; (c) Ep = 85 eV, Vs = −5.0 V, It = 0.1 nA; and (e)V nealin( rcles),( d 600( erred

∼1naisapatoketoati(O

sftamosao

s = +4.2 V, It = 0.5 nA. (g) Intensity change of (10) LEED spots as a function of an277 eV, open blue squares), N (392 eV, filled violet triangles), O (525 eV, open red cig) annealing was performed at 250 ◦ C for 12 h, 400 ◦ C for 1 h, 500 ◦ C for 10 min, anFor interpretation of the references to colour in this figure legend, the reader is ref

10 nm size and steps, as shown in Fig. 5(c). PV and RMS were.7 and 0.39 nm, respectively. When the duration of degassing wasot sufficiently long, the vacuum was too low in the followingnnealing process to obtain a sharp 1 × 1 LEED pattern and STMmages with wide flat terraces. We observed flat wide terraces withteps of height ∼0.5 nm on the type-1 GaN surfaces after three-stepnnealing, as shown in Fig. 5(e). This height is close to the latticearameter c, but no ordered atomic arrangements were observednywhere. We confirmed that the spot size and/or peak intensity inhe LEED patterns did not change from 5 min to 12 h in the final stepf annealing at 500 ◦C. We emphasize that the pressure should beept less than ∼5 × 10−7 Pa during 400 and 500 ◦C annealing, oth-rwise diffuse LEED patterns and many protrusions appear on theerraces in STM corresponding to oxide and/or carbide. When wemitted the second annealing (400 ◦C), it was impossible to obtainrelatively clean and flat surface. AES signals (Fig. 5(g)) showed

he reduction of oxide and carbide from the surface upon anneal-ng at ∼ 200 and 500 ◦C, which is consistent with the TPD resultsFig. 4(b)). On relatively clean surfaces, the AES intensities of C and

decreased to ∼20% of those for the as-introduced surfaces.Facets start to appear on the GaN surface above 550 ◦C. Fig. 6(a)

hows typical RHEED pattern of a type-1 sample annealed at 600 ◦Cor 20 min after the three-step annealing. We can see characteris-ic “chevrons” in the RHEED patterns, corresponding to facets, inddition to transmission spots from 3D GaN islands. These trans-

ission spot intensities were weak, indicating a small number

f epitaxial islands. The chevron patterns were broad, indicatingmall-size domains. Although the chevron directions are 〈1 1 2̄ 4〉t the [1 1̄ 0 0] incidence (Fig. 6(a)), we consider that these are partf 〈1 0 1 2̄〉 reciprocal rods, which correspond to {1 0 1 2̄} facets, from

g temperature (open black inverted triangles). In addition, the AES intensity of Cand Ga (1098 eV, filled green circles) are plotted against annealing temperature. In

◦ C for 10 min. The incident Ep for LEED and AES were 95 eV and 2 keV, respectively.to the web version of the article.)

the LEED results of Fig. 6(b). Fig. 6(b) and (c) respectively show aLEED pattern and typical STM image of a type-1 sample annealedat 600 ◦C for 5 min after the three-step annealing. We can see char-acteristic spots in the LEED pattern (some of them are indicated byarrows); the motion of these spots with changing Ep was differentfrom that of the fundamental spots, reflecting the inclined recipro-cal rods of the facets. The motion of the facet spots was restrictedin the directions from the 00 spots to the 10, 01, 1̄1, 1̄0, 01̄, 11̄spots, implying that the inclined reciprocal rods are oriented in the〈1 0 1̄ n〉 directions. Since the projection of the RHEED chevrons withthe 〈1 1 2̄ 4〉 direction (Fig. 6(a)) in the 〈1 0 1̄ n〉 direction gives the〈1 0 1̄ 2〉 direction, we consider the existence of six fold facets of{1 0 1̄ 2}. This facet plane is the so-called R-plane [73]. There havebeen many reports about facets formed on GaN surfaces. Most ofthem were formed during chemical etching [4,74] and at the begin-ning of crystal fabrication [75], where facets with a size of fewtens of �m, for example, {1 0 1̄ 1̄} facets, were formed. The {1 0 1̄ 2̄}orientation of the facet planes on GaN(0 0 0 1) films has also beenreported [76]. For MOCVD GaN(0 0 0 1) samples annealed at around800–900 ◦C in UHV, facet LEED patterns were shown in some Refs.[51,77,78].

According to STM images, many hillocks (a few-10 nm in diam-eter and a few nm in height) appeared on the surface, as shownin Fig. 6(c). Although the shapes of the facet islands were notclear in the STM images, these hillocks should be the origin of the

facets patterns in the LEED and RHEED patterns. From the resultsof XPS, we found that the facet surface was Ga-rich (Table 1), corre-sponding to the higher desorption of NH3 than Ga and GaCl belowapproximately 600◦ in TPD (Fig. 4(b)). At present, although themechanism of facet formation by annealing in UHV is not clear,

4752 A.N. Hattori et al. / Applied Surface Science 256 (2010) 4745–4756

Fig. 6. (a) Typical RHEED pattern of type-1 (HVPE) sample annealed at 600 ◦ C for 20 min after three-step annealing. The incident 15 keV beam was in the [1 1̄ 0 0] direction.T ) samm re indf obseri

ftdKos[

oisoSfoiSnpOtbit

FT

he chevrons indicate {1 0 1̄ 2} facets. (b) LEED pattern (Ep = 68 eV) of type-1 (HVPEarked by red circles. The other spots correspond to facets, some of which them a

acet surface. Numerous protrusions ascribed to facets with a size of a few nm weres referred to the web version of the article.)

acets are thought to grow mainly by the sublimation of nitrogen;he enhancement of grooves on the surfaces in the sublimation pro-uces isolated 3D islands with facets. For the reaction of GaN withOH solution [4], the selectivity of Ga to etching has been pointedut to induce facet formation. Instead of GaN facets, Blant et al.uggested that segregated Ga droplets comprised the protrusions79].

In type-3 HVPE samples, similar dependence of the behaviorf surface structures on the annealing temperature as that shownn the type-1 HVPE sample was confirmed. However, the micro-copic structures observed by STM were quite different from thosef the type-1 sample. Fig. 7 shows a RHEED pattern and typicalTM image of a type-3 sample after three-step annealing (200 ◦Cor 12 h + 400 ◦C for 1 h + 500 ◦C for 5 min). Sharp 1× 1 spots with-ut any transmission spots are seen in the RHEED pattern, as shownn Fig. 7(a). The LEED pattern also showed sharp 1 × 1 spots. TheTM image indicated a flat surface with numerous grains of a fewm in size as shown in Fig. 7(b). Some grains gathered to formetal-like shapes (indicated by a diamond in the inset in Fig. 7(b)).

n less than a few percent of the surface area, the petal-like struc-

ures coalesced with each other (indicated by diamonds in Fig. 7(c)),ut these structures are scattered. The side length of the diamonds

s ∼2 nm; thus, the area of one petal-like structure correspondso ≈ 6a × 6a. Note that the directions of the side of the diamonds

ig. 7. (a) RHEED pattern and (b) and (c) typical STM images of type-3 (HVPE) samples ahe conditions of observation are (b) Vs = +4.0 V, It = 0.5 nA and (c) Vs = +3.5 V, It = 0.7

ple annealed at 600 ◦ C for 5 min after three-step annealing. Fundamental spots areicated by arrows. (c) Typical topographic STM image (Vs = +2.0 V, It = 0.2 nA) of aved. (For interpretation of the references to colour in this figure legend, the reader

are parallel to the 〈1 1 2̄ 0〉 directions. These petal-like structuresresemble those in the STM image in Fig. 2(b) in Ref. [16], wherea 12×12-reconstructed structure which was any seen previouslyon neither the (0 0 0 1) surface nor the (0 0 0 1̄) surface [25]. Feen-stra et al. [16] associate this novel structure with the “1 × 1” Gaadlayers on GaN(0 0 0 1) layers and imply the existence of an inver-sion domain immediately below the novel reconstruction area ofthe surface. On type-3 sample surfaces, petal-like structures werereproducibly observed but neither pits nor holes were observed.They were present in at most a few % in the whole observed area.The type-3 sample was grown by the multi-step lateral epitaxialover growth and there is the defects distribution inside this crystal[80]. Thus, we associate the petal-like structure on the surface withthe high-density of defects area immediately below the inside ofthe crystal, and conclude that the petal-like structure in this papercorrelate with the novel reconstruction in the prior work [16], with-out any independent proof of the existence of an inversion domainthere. The petal-like structures observed on type-3 samples werenot observed on type-1 samples.

3.5. Ar ion sputtering and subsequent annealing of HVPE samples

Ar ion sputtering is commonly used to prepare clean surfacesand has been used for Ga compound III–V semiconductor surfaces

fter three-step annealing (200 ◦ C for 12 h, 400 ◦ C for 1 h, and 500 ◦ C for 10 min).nA. The incident electron direction in the RHEED pattern was [1 1̄ 0 0].

A.N. Hattori et al. / Applied Surface Science 256 (2010) 4745–4756 4753

Fig. 8. (a) Normalized XPS intensities of components of C 1s (purple triangles), O 1s (green inversed triangles), Ga 3d(GaN) (red solid squares), Ga 3d(GaOx ) (pink open squares),N 1s(N ) (blue solid circles), and N 1s(N ) (cyan open circles), as a function of Ar+ sputtering time for type-1 (HVPE) samples. Here, each intensity was normalized to the

a d (c) Rs d 500t on of

swXfrsstpvtc

irawwatoeoswwitribtNtsisccca

j

Ga H2s-introduced (no-treated) value. The Ar+ sputtering rate was ∼0.1 nm/min. (b) anputtering and subsequent three-step annealing (200 ◦ C for 12 h, 400 ◦ C for 1 h anhe references to colour in this figure legend, the reader is referred to the web versi

uch as GaP [81], GaAs [82], GaSb [83], and GaN [84]. For GaN, itas reported that nitrogen ions are more effective than Ar ande ions in removing C and O, and subsequently the annealed sur-

aces exhibited greater ordering [84,85]. In the previous studies theeduction of surface contamination and the appearance of 1 × 1urface diffraction patterns were revealed. Here, we present theurface structures and compositions obtained through Ar+ sput-ering and subsequent annealing, and discuss their microscopichenomena. We found that Ar+ sputtering of the GaN surface isery effective for removing oxide and carbide from the surface, buthe enhancement of surface roughness caused by the sputteringould not be recovered by postannealing.

The as-introduced GaN surfaces have carbon and oxygen asmpurities (Table 1 and Fig. 5(g)). Ar ion sputtering is effective foremoving C and O surface contamination. Table 1 and Fig. 8(a) showsignificant reduction of C 1s, O 1s, and Ga 3d (GaOx ) componentsith sputtering time. After 20 min Ar+ sputtering, the C 1s peakas not detectable and the surface oxygen component was onlyfew atomic%. As previously mentioned, HVPE GaN (type-1 and

ype-3) samples include Cl as an impurity. In XPS spectra, small butbvious Cl 2s, 2p, and Auger peaks (< 1 atomic%) always appeared,ven after Ar+ sputtering. This implies the existence of Cl in the bulkf the HVPE type-1 samples, as mentioned in Section 3.3. Fig. 8(a)hows a higher rate of increase of Ga 3d (GaN) than that of N 1s (NGa)ith sputtering time; the intensity of GaN increases over tenfoldhile that of NGa increases around threefold after 20 min sputter-

ng. The GaN and NGa intensities against all elements observed onhe surface increased from 16 to 26 atomic% and 15 to 22 atomic%,espectively, as shown in Table 1. The higher rate of increase of GaNs thought to be caused by the difference in sputtering efficiencyetween Ga and N. A molecular dynamics simulation [86] showedhat Ga atoms are always sputtered with N atoms in pairs while

atoms are mostly sputtered alone. This simulation also showedhat in Ar+ sputtering at an incident energy of 500 eV, the nitrogenputtering yield is about five times higher than that of gallium. Thisndicates that the Ar+-sputtered GaN surface does not maintain itstoichiometry and becomes Ga-rich at the near-surface. Under ouronditions (20 min sputtering), however, the stoichiometry did not

hange greatly (Table 1), despite the increase of the GaH−N and NH2omponents. The hydrogen included in samples is expected to playrole in stabilizing the stoichiometry against sputtering.

Fig. 8(b) shows a typical RHEED pattern of a type-1 sampleust after Ar+ sputtering for 20 min. Compared with Fig. 1(e), the

HEED patterns of type-1 samples: (b) after Ar+ sputtering for 20 min, (c) after Ar+

◦ C for 10 min). The incident electron direction was [1 1̄ 0 0]. (For interpretation ofthe article.)

1 × 1 spots are more diffused and the background level is higher.Faint transmission spots are seen. After the three-step annealing(200 ◦C for 12 h + 400 ◦C for 1 h + 500 ◦C for 5 min), the intensitiesof the transmission spots (cyan circles) increased (Fig. 8(c)). Thetransmission spots indicate the existence of 3D islands on thesurface. Indeed, we observed many islands in STM images of thethree-step annealed type-1 samples. In the RHEED patterns somefaint transmission diffraction patterns and chevron patterns corre-sponding to facets were also observed at different incident azimuthangles. The chevron patterns, implying the presence of facets, wereobserved after 600 ◦C annealing for both 10 min and 12 h. The exis-tence of epitaxial 3D GaN islands implies that the recovery froman Ar+-sputtered surface is not sufficient to form a flat surface;postannealing cannot produce a flat surface. In contrast, withoutAr+ sputtering, as mentioned in Section 3.4, the intensity of thetransmission spots is strong, indicating a large number of epitax-ial islands. For HVPE and MOCVD homoepitaxial surfaces withoutAr+ sputtering, 3D islands with facets also formed by annealingabove 550 and 600 ◦C, respectively. For Ar+-sputtered HVPE type-1 samples, the formation of 3D islands started at a slightly lowerannealing temperature (500 ◦C). This would imply that the sput-tering induced surface damage leads to easier sublimation, thusproducing 3D islands. Metallic Ga (˛ [87], ˇ [88], � [89], ı [90], �-Ga, Ga(II), or Ga(III) [91]) in the liquid phase shows halo patternsin diffraction, however, we could not observe any halo patterns inRHEED. According to the STM observation, many small clusters andgrains appeared to form.

3.6. Annealing of MOCVD samples

As shown in Fig. 1(b), (d), (f), and (h), as-grown MOCVDGaN(0 0 0 1) surfaces have a 1 × 1 structure even after exposureto air. In Section 3.1 we suggested the inert property of MOCVD-grown GaN(0 0 0 1) surfaces terminated by hydrogen. Similarly tothe HVPE samples, the MOCVD samples displayed transmissionspots from 3D GaN(0 0 0 1) islands in RHEED images after anneal-ing above 600 ◦C (550 ◦C for the HVPE samples), and were damagedafter annealing at ∼800 ◦C (Fig. 3(c)). However, the behavior of the

MOCVD samples below 600 ◦C was different from that of the HVPEsamples.

After annealing below 500 ◦C, LEED and RHEED patterns showed1 × 1 spots, which became sharper with narrower peak width withannealing, where H2O, carbide, and oxides were desorbed from

4754 A.N. Hattori et al. / Applied Surface Science 256 (2010) 4745–4756

F ype-4a ace re

tsscsniicntbssbttdltrattrtttlsdd

4

itttam1

ig. 9. Typical (a) LEED (Ep = 68 eV) and (b) RHEED ([1 1̄ 0 0] incidence) images of tt 550 ◦ C for 5 min. Clear 3 × 3 spots in the LEED pattern and faint and streaky surf

he surface (Fig. 4(d)). After annealing at 550–600 ◦C, weak 3 × 3uperspots appeared in the LEED and RHEED patterns for type-4amples, while 1 × 1 spots were dominant. Fig. 9(a) shows a typi-al LEED pattern of a type-4 sample. For the type-2 sample, the 3 × 3treaky spots in the RHEED pattern were very faint (Fig. 9(b)), ando 3 × 3 spots could be observed in the LEED pattern. These results

mply that the 3 × 3 domains on type-2 surfaces are small and fewn number. In STM we could not observe the 1 × 1 or 3 × 3 atomiconfiguration on any MOCVD homoepitaxial samples since the tun-eling current was not stabilized due to insulating character ofhese sample surfaces. It was reported a 3 × 3 reconstruction is notelong to a Ga-polar GaN(0 0 0 1) surface, but a N-polar GaN(0 0 0 1̄)urface [25,92]. Bermudez et al. reported a 3 × 3 reconstructiontructure observed on an MOCVD GaN(0 0 0 1) film on ˛-sapphirey LEED after N ion bombardment and annealing, and concludedhat the 3 × 3 structure was caused by a N vacancy on the defec-ive surface [93]. As mentioned in Section 3.3, the NH3 started toesorb from MOCVD GaN at the lower temperature due to their

ower crystallinity. Although we confirmed directly by XPS, theype-2 and type-4 samples surfaces are expected to become Ga-ich/N-defective after annealing at 550–600 ◦C, similar to the 600 ◦Cnnealed type-1 sample (Table 1). The observed 3 × 3, which seemso be different 3 × 3 reconstruction on GaN(0 0 0 1̄) surface, is ten-atively denoted as “defective-3 × 3” here. Further observations areequired to conclusively identify the “defective-3 × 3” reconstruc-ion, but, the intensities and FWHM of their RHEED spots indicatehat the domain number and size on type-4 samples are largerhan those on type-2 samples. Thus, we can conclude that it is ateast not unlikely that the formation of the “defective-3 × 3” recon-truction depend on the sample crystalline quality (the dislocationensity) and these facts suggest that type-4 samples have a higherislocation density.

. Conclusion

Using LEED, RHEED, STM, TPD, XPS, AES, SIMS, and XRD, we stud-ed the characteristics of four types of GaN sample surfaces: twoypes of HVPE GaN(0 0 0 1) surfaces and two types of homoepi-

axial MOCVD GaN(0 0 0 1) surfaces. The samples were subjectedo different treatments: the annealing of as-introduced samplesnd the annealing of Ar+-sputtered samples in UHV. Under opti-ized three-step annealing conditions (200 ◦C for 12 h + 400 ◦C forh + 500 ◦C for 5 min) without the sputtering, the surface contam-

and type-2 (MOCVD) samples, respectively, after annealing at 200 ◦ C for 12 h andconstruction spots in the RHEED pattern (arrows in (b)) are seen.

ination of oxide and carbide was reduced to ∼20% of that beforeannealing. Clear GaN(0 0 0 1)1 × 1 patterns were obtained by LEEDand RHEED for all types of samples. For the HVPE samples (type-1and type-3), 500 eV Ar+ sputtering was effective in removing C andO surface contamination, but the increased surface roughness as aresult of Ar+ sputtering could not be recovered by postannealing,and the formation of 3D islands with facets was enhanced on thesputtered surfaces. The homoepitaxial MOCVD (type-2 and type-4)surfaces showed 3 × 3 structures in the LEED and RHEED patternsafter annealing at 500–600 ◦C, while the 1 × 1 structure was dom-inant. No 2 × 2 reconstructed structures, expected to appear onin situ MBE surfaces and predicted in theoretical studies, wereobtained by annealing and Ar+ sputtering. For all sample types,3D islands with facets formed upon annealing above 550–600 ◦C,and the surfaces became rough with unbalanced stoichiometry(Ga-rich/N-defective). Therefore, we suggest that the surface tem-perature should be kept below 550 ◦C during GaN homoepitaxialgrowth to obtain high interface integrity and high crystallinity.

We also found that surface-lapping and finishing treatmentsseverely damaged the GaN surface properties. Although almost allsemiconductor substrates include a CMP step in their manufactur-ing processes, CMP techniques are commercially confidential andtheir phenomena are not well understood scientifically. Of course,CMP techniques with high performance are well established, espe-cially for Si wafers. However, CMP treatments for GaN wafers havenot yet been particularly successful [57,94,95]. We hope that a sci-entifically modified CMP process will be established in the nearfuture as a result of a technological breakthrough. At present, it isdifficult to commercially obtain well-ordered GaN, such as with a2 × 2 reconstruction structure, because the crystal and/or surface-finishing process is not of an acceptable quality. Diffraction is usefulfor evaluating the properties of wafers after surface-finishing treat-ments, i.e., not only their flatness but also their crystal ordering, aswe have done here using LEED and RHEED. To achieve electronicdevices with excellent properties, for instance, carrier transport, weshould improve the crystal quality, and surface-finishing methods,and understand the local (microscopic) surface structure becauseit becomes the interface under electrode materials.

Acknowledgements

This study was partly supported by a grant of the Global COE Pro-gram, “Center of Excellence for Atomically Controlled Fabrication

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echnology from the Ministry of Education, Culture, Sports, Sciencend Technology (MEXT) Japan and Kyoto-Advanced Nanotechnol-gy Network under sponsorship from the Ministry of Education,ulture, Sports, Science and Technology (MEXT), Japan. One of theuthors (ANH) was also supported in part by Special Coordinationunds for Promoting Science and Technology from the Ministry ofducation, Culture, Sports, Science and Technology for the Osakaniversity Program for the Support of Networking among Presentnd Future Women Researchers.

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