Shape, size evolution and nucleation mechanisms of GaAs

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Shape, size evolution and nucleation mechanisms of GaAs nanoislands grown on (111)Si by low temperature metalorganic vapor phase epitaxy Ilio Miccoli, Paola Prete, and Nico Lovergine Cryst. Growth Des., Just Accepted Manuscript • DOI: 10.1021/acs.cgd.9b00225 • Publication Date (Web): 13 Aug 2019 Downloaded from pubs.acs.org on August 20, 2019

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Crystal Growth & Design

Shape, size evolution and nucleation mechanisms of GaAs nanoislands grown on (111)Si by low temperature metalorganic vapor phase epitaxy Ilio Miccoli,‡,* Paola Prete,,† and Nico Lovergine ‡,§ ‡ Dipartimento

di Ingegneria dell’Innovazione, Università del Salento, Via Monteroni, I-73100 Lecce, Italy.

 Istituto

per la Microelettronica e Microsistemi, Consiglio Nazionale delle Ricerche, SS Lecce, Via Monteroni, I-73100 Lecce, Italy.

Abstract. The shape, size evolution and nucleation mechanisms of GaAs nanoislands grown at 400°C on As-stabilized (111)Si by metalorganic vapor phase epitaxy are reported for the first time. GaAs crystallizes in the zincblend phase since the very early nucleation stages until the

*

Present address: Aixtron SE, Dornkaulstraße 2, 52134 Herzogenrath (Germany).

†

Corresponding author (Paola Prete). e-mail: [email protected]; phone: +39 0832 297 250;

fax: +39 0832 297 249; https://orcid.org/0000-0002-4948-4718 §

https://orcid.org/0000-0003-0190-4899

ACS Paragon Plus Environment

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formation of a continuous epilayer. GaAs nanoislands grow (111)-oriented on Si as truncated hexagonal pyramids, bound by six equivalent {120} side facets and a (111) facet at the top. Their diameter and height appear to increase linearly with the deposition time, yielding a constant aspect ratio of 1/4. The nanoisland density (before coalescence) stays constant with time at 21010 cm-2, suggesting that their nucleation occurs at specific Si surface sites (defects) during very early growth stages, rather than being due to the continuous formation of new nuclei. To understand the molecular-level mechanisms driving the low-temperature MOVPE growth of GaAs on Si, we applied a deposition-diffusion-aggregation (DDA) nucleation model, which predicts a linear evolution of the overall GaAs growth rate with surface coverage, in good agreement with experimental observations, under the assumption that direct impingement of trimethylgallium (Me3Ga) molecules onto the nanoislands surface dominates the material nucleation and growth rate; the contribution of Me3Ga adsorbed onto the As-stabilized (111)Si shows negligible, pointing out the reduced reactivity of Si surface (As-passivation). Our DDA model allows to estimate the effective reactive sticking coefficient of Me3Ga onto GaAs, which turns equal to 2.8210-5: the small value is compatible with the Me3Ga large steric hindrance and the competitive role of methyl radicals to surface adsorption at low temperature.

Keywords: GaAs heteroepitaxy, Si substrate, nanoislands, MOVPE, nucleation, depositiondiffusion-aggregation model, reactive sticking coefficients.


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Introduction The hetero-epitaxy of III-V compound semiconductors (in particular, GaAs) on Si substrates has been in the focus of intense research during the last four decades, as its exploitation allows the direct monolithic integration of highly efficient optoelectronic devices based on IIIVs with Si-based microelectronics. To this purpose, studies in the field have primarily dealt with the hetero-epitaxy on (100)Si substrates,1,2,3,4 as almost solely (100)-oriented Si wafers are used in microelectronics. However, in recent years the growth of III-Vs onto (111)Si substrates has been attracting an increasing number of studies, as (111)-oriented III-V layers can be employed for the subsequent growth of III-V nanowires.5,6,7,8,9 Free-standing III-V nanowires are more desirable for device integration on Si substrates than planar epilayers, as they permit (i) a facile relaxation of lattice-mismatch constraints and avoidance of mismatch-induced structural defects within the active volume of the device,10 and (ii) fabrication of novel and more (with respect to planar counterparts) efficient nanodevices, such as nanowire-based field-effect transistors,11 light-emitting diodes,12 laser diodes,13 photodetectors,14 and solar cells.15,16 Nanowire nucleation by most common selfassembly methods16,17,18 occurs along the semiconductor [111]-direction, which means that the preliminary deposition of a thin III-V continuous epilayer onto a (111)Si substrate can ensure their vertical alignment with high yield.6,7 Also, the use of a thin III-V buffer layer may provide the necessary chemical separation between the Si substrate, the nanowires and the metal catalyst (if any) employed for their self-assembly during early nucleation stages.



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The epitaxy of high quality III-V epilayers onto (111)Si remains however, as challenging as that on (100)Si due to (i) large lattice and thermal mismatches (4% and 55% respectively, for GaAs/Si), and (ii) the well-known polarity mismatch between III-Vs and Si crystals.1 As result, large densities of mismatch dislocations can be found at III-V/Si hetero-interfaces,1,19,20 and a high degree of mosaicity; while the latter is associated with the formation of anti-phase domains (APDs) at single atomic steps on the (100)-oriented Si surface,21,22 a large amount of rotational (twinned) crystal domains (i.e., domains rotated by 180° with respect to each other around the epilayer 111 lattice directions)7,23 are observed in (111)-oriented epilayers but no APD nucleation, since only bilayer-height steps exist on the (111)Si surface.7 As matter of fact, the crystallinity of III-V epilayers grown on (111)Si turns up better than that for growth of (100)-oriented Si substrates.2 The actual amount of defects and surface roughness of a III-V epilayer on Si strictly depend on materials interaction and nucleation mechanisms: for GaAs epitaxy, the large lattice mismatch with Si inevitably leads to three-dimensional (the so-called Volmer-Weber) nucleation3 and, as the growth proceeds further, islands coalescence. Clearly, the formation of a continuous GaAs layer on Si is desirable before structural defects are introduced by plastic relaxation; however, the coalescence of coherent islands brings along further defects, which may be eliminated/reduced by additional process steps, such as annealing treatments and/or the growth of thicker epilayers.1 Furthermore, the shape, size and density of coalescing islands largely determines the surface roughness of the as-obtained continuous layer. In this respect, best GaAs epilayer crystallinity and surface roughness has been achieved so far by the soACS Paragon Plus Environment

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called two-step growth method, consisting of a first low-temperature (400450°C) growth of a thin nucleation layer, followed by high temperature annealing and the growth of a second layer at 600650°C,19,24 the latter temperatures being typical for growth of device-quality GaAs by molecular beam (MBE) or metalorganic vapor phase (MOVPE) epitaxy. In the twostep growth approach, the low-temperature nucleation step is pivotal to the formation of a dense array of (nearly) coherent nano-islands which afterwards coalesce to form a continuous, thin and relatively smooth GaAs epilayer. It is thus of the utmost importance to investigate the specific mechanisms driving these phenomena. Early nucleation stages and the coalescence of GaAs nanoislands on (100)Si substrates have been previously studied in the literature,3,25 while only a few reports have been published on initial deposition onto (111)-oriented Si substrates.19 In this work, the crystalline phase, shape, size evolution and nucleation mechanisms of GaAs nanoislands on in-situ As-stabilized (111)Si by low-temperature (400°C) MOVPE up to the formation of a continuous thin layer are reported in details. Furthermore, a deposition-diffusion-aggregation (DDA) nucleation model, sensitive to the different reactivity of metalorganic precursors to GaAs islands and Si, is applied to explain the deposit growth rate change with substrate surface coverage, in good agreement with experimental observations.

Experimental



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GaAs epilayers were grown on (111)Si substrates by low (50 mbar) pressure MOVPE in an Aixtron 200RD reactor, using trimethylgallium (Me3Ga) and tertiarybutylarsine (tBuAsH2) as Ga and As precursors, respectively. To this purpose, p-type (111)-oriented Si wafers were cleaned in iso-propanol vapors for 1 h, etched at room temperature in diluted (5% by volume in d.i. H2O) HF for 2 min to remove the native surface oxide layer, thoroughly washed in d.i. H2O, and blown-dry in 6.0N pure N2. Immediately after this treatment, the substrates were loaded into the reactor chamber and kept under pure H2 flow until the process was initiated. The substrates were then heat-cleaned at 700°C for 30 min under H2, and further annealed in a H2+tBuAsH2 flow for 10 min. Thermal treatment in As-vapor has a twofold purpose: (i) it lowers the minimal temperature for decomposition of residual oxides left on the Si surface well below that (950°C) required for annealing under pure H2,2 ensuring almost complete deoxidation of the substrate; and (ii) it provides an As-stabilized Si surface,26,27 which leads to the nucleation of (111)B-oriented GaAs.7 To keep the Si surface under these As-stabilized conditions until growth start the substrate was cooled down to the final growth temperature under continuous H2+tBuAsH2 flow. GaAs was subsequently nucleated at 400°C on as-treated substrates using a tBuAsH2:Me3Ga molar flow ratio of 10:1. The growth time was varied between 5 min and 1 h, so to obtain various degrees of coverage of the Si surface by the GaAs deposit, i.e. from sparse nanoislands up to a continuous thin layer. The morphology and size of GaAs nanoislands were determined by a combination of field emission scanning electron microscopy (FE-SEM) and contact-mode atomic force microscopy (AFM) observations; FE-SEM observations were performed using a Zeiss microscope model ACS Paragon Plus Environment

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Sigma VP equipped with a Gemini electron column, whereas AFM topographic images were recorded by means of a VEECO Caliber microscope. Quantitative analyses of the FE-SEM and AFM images allowed to determine the size (diameter and height) distribution of GaAs nanoislands, their 3-dimensional shape, and surface coverage of as-deposited Si substrates. The nanoisland-substrate epitaxial relationship was assessed by X-ray diffraction (XRD) patterns recorded in the Bragg (-2) geometry using a Rigaku D-Max/Ultima+ diffractometer equipped with a MPA2000 thin-film attachment stage and a Cu-anode X-ray tube.

Results and discussion Figure 1 shows the surface morphology of GaAs-deposited samples grown at 400°C on (111)-oriented Si substrates for different growth times. As expected, the early nucleation of GaAs onto Si follows a 3-dimensional (Volmer-Weber) growth mode: indeed, Fig. 1 shows the formation of a dense array of relatively flat nanoislands. After 5 min growth, the nanoisland density reaches a value of 21010 cm-2, which remains almost constant upon increasing the growth time up to about 11 min, while the nanoisland size keeps increasing. After 15 min the nanoislands begin to coalesce, until an almost complete coverage of the Si substrate surface is obtained for about 1 h growth. The epitaxial relationship between the GaAs nanoislands and the (111)Si substrate is evidenced by the high-angle -2 XRD spectra reported in Fig. 2 for some of the samples discussed above: the XRD patterns show only the (111) and (333) diffraction peaks of



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zincblend (ZB) GaAs, along with the (111), (222), and (333) diffraction peaks of Si, demonstrating that the GaAs nanoislands are epitaxially well aligned along the substrate [111] growth direction. Yasuda et al.28 previously reported the low temperature (200°C) MBE nucleation of stable zincblend (ZB) phase GaAs onto (111)Si for layer thickness below 5 nm, but with increasing thickness the crystal phase changed to wurtzite (WZ) by formation of orderly-arranged stacking faults. However, no hints of WZ phase was observed in the XRD spectra of our samples, demonstrating that the ZB structure is the stable phase for MOVPEgrown GaAs at present temperatures.

Figure 1. (a-f) FE-SEM micrographs (30,000 magnification, 45°-tilt view) of the GaAs deposit on (111)Si for several growth times (indicated in the images). White markers in (d) represent 250 nm for micrographs (a-c), and 500 nm for micrographs (d-f).


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Figure 2. -2 XRD spectra of GaAs/(111) Si samples for increasing growth time, namely (a) 15 min, (b) 30 min, and (c) 60 min.



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Figure 3. (a) An AFM topographic image of GaAs nanoislands on (111)Si substrate after 15 min of growth; scan size is 11 μm2. (b) A typical AFM line scan, taken across the topographic image in (a). Noteworthy is the appreciable broadening of the GaAs diffraction peaks with respect to Si ones: although we cannot completely rule out the contribution of extended defects and mosaicity, most of the broadening shall be ascribed to the nanometric size (namely the Debye-Scherrer effect) of GaAs crystals. Indeed, the GaAs diffraction peaks become narrower for longer deposition times (i.e., larger nanoislands). Figure 3(a) reports an atomic force microscopy (AFM) topographic picture of the 15 min grown GaAs nanoislands [same sample of Fig. 1(d)], for which the nanoisland coalescence has not yet set in. The AFM topography allows a quantitative evaluation of the nanoisland heights [estimated from the analysis of several AFM line profiles, such as the one shown in Figure 3(b)]. The count histogram of Fig. 4(d) shows the results of such analysis. It appears that the nanoislands have a relatively narrow height distribution, which can be best-fitted by a Gaussian profile with an average nanoisland height have=(11.0±0.2) nm, and a standard deviation of around σh= (2.3±0.3) nm. Plan-view field emission scanning electron microscope (FE-SEM) micrographs of the same samples


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Figure 4. (a,b) Plan-view and cross-sectional FE-SEM micrographs of a GaAs nanoisland on (111)Si after 15 min growth. (c) Reconstructed 3-dimensional shape of a GaAs nanoisland; the crystallographic planes of the facets bounding the nanoisland and its orientation with respect to the Si in-plane 〈112〉 directions are indicated. Count histograms of the GaAs nanoislands heights (d) and base diameters (e); the solid curve in each diagram represents the Gaussian function best-fitting the experimental data, with average height have=(11.0±0.2) nm and base diameter Dave=(47.2±8.8) nm. demonstrate that GaAs nanoislands have a hexagonal in-plane section (Fig. 4(a)). By comparing the directions of the nanoisland hexagonal edges with those of the substrate cleaved edges (and assuming a parallel epitaxy relationship between the two materials), it



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turns out that each couple of nanoisland parallel edges is normal to one of the substrate three equivalent in-plane crystallographic directions. Clearly, these edges correspond to welldefined crystallographic side facets, laterally bounding the nanocrystal. In order to determine the crystallographic planes corresponding to these side facets, the contact angle that the GaAs nanoislands form with the Si substrate surface was measured through cross-sectional FE-SEM observations of cleaved samples. Fig. 4(b) shows a cross-sectional FE-SEM micrograph of a GaAs nanoisland grown for 15 min on (111)Si substrate. In this case, the FE-SEM primary electron beam is directed along the sample direction. The nanoisland has a contact angle of ∼35° along the direction; thus the angle between the Si(111) surface and each of the six nanoisland side facets is ∼39°, which is very close to the angular distance between {111} and {120} lattice planes (39.23°). On the basis of these results, it is possible to reconstruct the 3-dimensional shape of a typical GaAs nanoisland (Fig. 4 (c)), which consists of a truncated hexagonal pyramid bound by the {111} basal and top facets, and six {120} side facets. In this respect, we point out that our GaAs nano-islands do not show clear edges, but appear more rounded than in the schematics of Fig. 4(c): this is not due to a limited spatial resolution of our FE-SEM images but rather to the rounding of GaAs nano-facets, most likely due to a kinetic-roughening effect under the high vapour/surface supersaturation conditions associated with the low growth temperature of present samples. Systematic plan-view FE-SEM observations performed on all samples allowed a quantitative evaluation of the nanoisland base diameter D (taken as the diagonal length of the nanoisland hexagonal base). Fig. 4(e)


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reports a count histogram of the base diameter values, obtained from the analysis of several FE-SEM plan-view micrographs of the 15 min grown sample (Fig. 4(d)). The nanoisland diameter distribution can be again best-fitted by a Gaussian profile, with an average diameter Dave=(47.2±0.2) nm and a standard deviation σD=(8.8±0.5) nm. An estimate of the GaAs nanoisland aspect-ratio have/Dave= 0.23±0.01∼1/4 can be thus obtained. Noteworthy, the equilibrium crystal shape of a GaAs nanoisland on (100)Si was shown to approach h/D∼1/2 (Ref. 3).

Thus, the GaAs nanoislands grow flatter on (111)Si than on (100)Si, and this

guarantees a faster coverage of the substrate, along with a relatively smoother surface morphology of the final epilayer. Figure 5 shows the surface coverage θ of the Si substrate by GaAs (i.e. the fraction of the substrate surface covered by the GaAs nanoislands) as function of growth time. The data points in the Figure were calculated by digital image analysis of a series of plan-view FE-SEM micrographs of as-grown samples, while the associated error bars were estimated by taking into account the uncertainty in determining the nanoisland edges from FE-SEM images (corresponding to ±2 nm, the spatial resolution limit of the FE-SEM microscope). Since the surface of the 30 min grown sample showed slight fluctuations in θ, the error bar in this case is given by the maximum semi-dispersion of all estimated values. The solid line of Fig. 5, which is a simple guide for the eyes, points out that θ increases super-linearly up to about 15 min deposition time, after which a change in the growth regime occurs, clearly ascribable to the coalescence of GaAs nanoislands. The inset diagram in Fig. 5 demonstrates that the



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nanoisland average diameter increases linearly with time up to 15 minutes; as already mentioned, the nanoisland density remains instead constant at 2×1010 cm-2, suggesting that their formation occurs through the heterogeneous nucleation of GaAs nuclei at pre-existing metallic contaminations and/or defects on the Si surface, whose amounts in turn define the density of island nucleation sites. Taken together these findings suggest that, after an initial nanoisland nucleation and before their coalescence, the substrate coverage increasing is essentially accounted for by the increase of nanoisland dimensions. Therefore, assuming a Gaussian distribution for the nanoisland diameters around an average value D0 (Fig. 4(e)), the expected surface coverage θ can be estimated through the following relationship: 

 exp  D0     A(D)P(D, D0 )dD

(1)

0

where A(D)  ( 3D 2 / 2) is the hexagonal nanoisland area as function of its base diameter D,

P ( D , D0 )  e

2  D  D0    2    

/   2 is the Gaussian diameter distribution with mean value D0, δ is the

nanoisland surface density and σ is the average value of the diameter standard deviations [estimated from the samples FEG-SEM analyses of GaAs nanoislands before their coalescence]. A linear best-fit of nanoisland average diameter values, reported in Fig. 5(b), shows that the average diameter varies with the deposition time t as

D0  3.1  0.1 nm sec  t


(2)

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pointing out a GaAs nanoisland growth rate of 1.55 nm/sec normal to the in-plane substrate directions (i.e., normal to the nanoisland edges). Based on Eqs. (1) and (2) the expected surface coverage θ(t) can be finally calculated as function of growth time, and it is represented by the dashed curve in Fig. 5. The good agreement between the as-evaluated surface coverage and experimental values before nanoisland coalescence confirms that the coverage of Si by GaAs is entirely ascribable to the monotonic enlargement of a fixed number of nanocrystals nucleated during the very early stages of materials growth (i.e., during the first few minutes) rather than to the continuous nucleation of new nanoislands. After around 15 min the process enters the nanoisland coalescence regime, during which a decrease of the island density is observed, while FE-SEM analysis of these samples indicate a higher proclivity of GaAs to growth on top of coalesced nano-islands. These phenomena are in turn responsible for the deviation of the measured coverage from the calculated dashed curve in Fig. 5.



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Figure 5. (a) The fraction (θ) of Si surface covered by GaAs nanoislands as function of growth time. The solid line is only a guide for the eyes, while the dashed curve represents the materials coverage before nanoisland coalescence (t15 min) calculated according to Eqs. (1) and (2) [see text]. (b) Inset: average diameter of GaAs nanoislands D0 as function of growth time. The solid line is the linear best-fit of experimental data (see Eq. (2)). To gain more insights of the actual molecular-level deposition mechanisms of GaAs on Si, a modified version of the so-called deposition-diffusion-aggregation (DDA) nucleation theory,29 usually applied for modeling physical vapor deposition, such as MBE,29 and chemical vapor deposition processes,30,31 is here employed. According to the original DDA model, atomic/molecular species randomly adsorbed onto the growth surface (ad-species) can diffuse and then either aggregate/react with other ad-species to create stable crystal nuclei


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(nanoislands) or are captured by already existing nanoislands, thus contributing to their further growth. Although a MOVPE process includes several chemical reactions, both in the gas-phase and onto the substrate surface, the physical assumptions of the DDA model may be sufficient to identify which mechanisms drive the materials nucleation process and describe analytically its deposition rate (gr), whenever species adsorption/re-evaporation and their different reactive sticking probabilities at the growth surface play a major role, as it is the case in low temperature MOVPE. Under As-rich vapor and low temperature conditions the MOVPE growth of GaAs is limited by the supply and decomposition of the Me3Ga precursor; however, at 400°C only 5-7% of Me3Ga molecules decompose,32 likely through dissociative surface adsorption, leading to the formation of adsorbed di- methylgallium (Me2Ga) and mono-methylgallium (MeGa) intermediate species by sequential loss of methyl radicals (Me).33,34 These molecules may either directly adsorb onto existing GaAs nanoislands, sticking irreversibly with a certain probability (a process called direct impingement), or fall onto the (As-stabilized) Si surface. In this second case and based on the observed lack of continuous nucleation of new nanoislands, one can suppose that only precursor species adsorbed onto Si within a distance λ from a given GaAs nanoisland (where λ is the ad-species surface diffusion length before reevaporation) can reach the nanoisland and be incorporated, otherwise they will desorb (as schematically shown in Fig. 6 for Me3Ga). Indeed, if one takes into account exclusively the contribution of direct impingement, the growth rate of GaAs can be expressed as



g r' 

dh  f i Ai  v F f   v GaAs F1 i GaAs 1 dt S Sub

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(3)

Figure 6. Schematic of elementary surface processes taking place during the low temperature MOVPE deposition of GaAs onto Si, according to the deposition-diffusion-aggregation (DDA) nucleation model. Me3Ga stands for tri-methylgallium molecules. where h is the thickness of an equivalent GaAs compact layer, vGaAs is the GaAs molecular volume (cm3/molecule), F is the molecular collision frequency (molecules/cm2⋅sec) according to the kinetic theory of gases (see Eq. (8) below), η1 is the effective reactive sticking coefficient for direct impingement on GaAs nanoislands (namely, the probability of incorporating the film-forming species in each surface collision), while fi is the shape factor of the i-th GaAs nanoislands (see further below), whereas SSub and Ai  ( 3Di2 / 2) are respectively, the total substrate surface and, as before, the i-th hexagonal nanoisland in-plane areas. The last term on the right of Eq. (3) is written after assuming that all nanoislands have the same shape ACS Paragon Plus Environment

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factor (fi=f), so that one can write ∑ifiAi/STot=f θ where θ is the substrate surface coverage by the GaAs deposit. Noteworthy for θ =1 one obtains the growth rate of a continuous GaAs layer, that is gr’( θ=1)= νGaAsF η1f. Similarly, the contribution to the growth rate that comes from group-III species adsorbing onto the (As-stabilized) Si substrate within a diffusion length λ from each GaAs nanoisland, can be the expressed as

g  v GaAs F 2 '' r



A  i

'

i

(4)

S Tot



where Ai  3 Di     Di2 / 2 is the area of the hexagonal surface annulus around each '

2

GaAs nanoisland (i.e., the dark-grey shaded area in Fig. 6) and 2 is the reactive sticking coefficient for direct impingement on Si. The total GaAs growth rate can then be written as

g r  v GaAs F1 f   v GaAs F 2

A  i

S Tot

'

i

.

(5)

It is possible to distinguish two extreme cases: (i) if the direct impingement term dominates the deposition process (1»2), then the growth rate turns out directly proportional to the surface coverage θ ; (ii) if group-III species adsorbing nearby GaAs nanoislands appreciably contributes, gr will instead, more rapidly (i.e., for θ «1) approach the growth rate of a continuous film, namely ν GaAsF η 1f, which corresponds to a non-linear dependence on the surface coverage θ (see inset of Fig. 7).31 Therefore, the main mechanism controlling the GaAs growth dynamics onto Si can be clearly identified by evaluating the experimental growth rate gr as function of surface coverage.



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The GaAs growth rates of present samples were evaluated as the following. Before nanoisland coalescence, the GaAs equivalent thickness (hGaAs) can be estimated based on the sample nanoisland density and their diameter distribution (and assuming the same reconstructed 3dimensional shape of Fig. 4(c) for all nanoislands); it follows that 

hGaAs    V ( D) P( D)dD

(6)

0





where V ( D)  k  D 3 1  1  2k cot  2  1  2k cot   / 3 is the volume of a truncated hexagonal pyramid with a contact angle with the substrate α =35° (see discussion above on the nanoisland shape), P(D) is the Gaussian diameter distribution, and k = 3 h ave /Dave  / 2 = 3/8 . After the nanoisland coalescence has set in, i.e. for deposition times beyond 15 min, Eq. (5) can no longer be applied; still, the equivalent thickness of GaAs can be estimated from the Debye-Scherrer broadening of the (333)GaAs K α 1 and K α 2 XRD peaks (which are well separated and distinct from the substrate (333)Si peaks - Fig. 2(b),(c)). Assuming a negligible contribution from extended defects and epilayer mosaicity, the angular broadening WDS of the (333)GaAs peaks is related to the mean dimension (d) of as-grown samples in the direction normal to the scattering (111)-planes via the usual Scherrer equation35 d

 KB

WDS cos( B )

(7)

The XRD (333)GaAs peaks were best-fitted through the superposition of two Gaussian peaks, associated with the K α 1 and K α 2 fluorescence lines of the Cu anode X-ray source. Hence, an estimate of WDS was obtained from the full width at half maximum (FWHM) values of bestfitted peaks after subtracting to it the contribution of the instrumental and natural line ACS Paragon Plus Environment

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broadening. The equivalent thickness of the GaAs layer was then evaluated by multiplying the result of Eq. (7) by the measured surface coverage θ of the sample hGaAs  d   .

(8)

This last correction is necessary as no contribution to the GaAs peak broadening comes from voids in the GaAs deposit. The GaAs growth rate can be thus calculated as gr=hGaAs/t for both cases above (i.e. before and after coalescence of GaAs nanoislands).

Figure 7. Estimated GaAs growth rate (gr) values of deposited samples as function of the measured surface coverage θ . Full symbols were calculated via Eq. (5). Open symbols were estimated via the Debye-Scherrer broadenings of the (333) XRD peak of GaAs along with Eqs. (7) and (8). The dashed line represents the growth rate function gr(θ) of Eq. (3) best-fitting the



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data. The Inset diagram in the Figure shows the predicted (qualitative) trends for gr vs θ according to Eqs. (3) and (5). Fig. 7 shows the results of such analysis, where gr is reported as function of the surface coverage θ. Noteworthy, the growth rate increases linearly with the surface coverage, resulting in a maximum growth rate for the continuous ( θ =1) film of (5.89±0.15)×10-3 nm/sec. The linear dependence of experimental data observed in Fig. 7 demonstrates that the second term in Eq. (5) is negligible, indicating that group-III species that fall onto the Si substrate rapidly desorb before reaching the GaAs nanoislands ( η 2« η 1).36 The reason for this behavior should most likely be ascribed to the reduced chemical reactivity of the As-stabilized (111)Si surface.2 Therefore, under present conditions the direct impingement of group-III species onto the nanoisland surface dominates the nucleation dynamics of GaAs, this being the sole mechanism responsible for the increasing GaAs coverage of the Si substrate up to the formation of a continuous layer. On As-rich GaAs surfaces the low-temperature heterogeneous decomposition of Me3Ga is limited by its dissociative chemisorption,32-34 leading to the sequential loss of methyl radicals, and to Me, Me2Ga and MeGa molecules bound to the growth surface. While the Me2Ga molecule either readily loses one of its methyls or desorb, MeGa sticks more on the surface and its further dissociation (usually mediated by free As or As-H groups) gives rise to Ga incorporation into the crystal. Noteworthy, at temperatures close to 400°C, only 1/5 of available surface adsorption sites are occupied by Me3Ga molecules,37 a result of the precursor steric hindrance and the competitive effect of Me adsorption. According to this ACS Paragon Plus Environment

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view, the direct impingement (adsorption) of Me3Ga molecules onto the surface of GaAs nanoislands is supposed to be the driving mechanism of the growth process (Fig. 6) and in our DDA model its reactive sticking coefficient η1 for can be thus estimated from Eq. (3) 1  g r   1 /  GaAs Ff   2.82  10 5

(9)

where we have taken νGaAs=4.52×10-23 cm3/molecule as the GaAs molecular volume, the shape factor of GaAs nanoislands as f=ANanoislands/Ainterface =1.142 [i.e. the ratio of the GaAs nanoisland surface area to the nanoisland/substrate interface one, assuming the 3dimensional nanoisland shape in Fig. 4(c)] and F is the molecular collision frequency for Me3Ga given by





F  3.51 10 22 

PMe3Ga M Me3Ga T

 3.99 1017 molecules / cm 2  sec

in which MMe3Ga=114.8 gr/mol (molar weight of Me3Ga), T=673K (growth temperature), and PMe3Ga=3.1610-3 torr (nominal Me3Ga partial pressure in the MOVPE reactor chamber). The as-calculated low value of 1 agrees qualitatively with the reduced sticking probability expected for Me3Ga, as result of its steric hindrance and Me competition to surface adsorption at low temperature. In any case, the value of 1 falls within the order of magnitude range estimated by the DDA model for reactive sticking coefficients of other molecules in different chemical vapor deposition processes.30,31

Conclusions



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We reported in details on the crystal phase, morphology (shape), size distribution and nucleation dynamics of GaAs nanoislands grown by low-temperature (400°C) MOVPE on insitu As-stabilized (111)Si substrates. GaAs growth occurs in the stable ZB phase since the very initial stages and until the formation of a continuous epilayer. The nanoislands grow (111)oriented on the Si substrate in the form of truncated hexagonal pyramids, bound by six equivalent {120}-plane side facets and one (111)-plane facet at the top; the nanoislands aspect ratio approach ¼, which makes them relatively flat and guarantees a fast coverage of the substrate, along with a relatively smooth surface morphology of the final epilayer. Their diameter and height appear narrowly distributed with a Gaussian profile, the average values increasing linearly with deposition time. The nanoisland density (before coalescence) remains instead constant at 21010 cm-2, and we tentatively ascribe their nucleation as due to specific Si surface sites (defects, metal contaminations) during the very early growth stages (the first few minutes) rather than being due to the continuous formation of new nuclei. In this respect, a thorough understanding of the nature of those defects/contaminations and their control through more specific treatments of the As-passivated Si surface is necessary to optimize the crystalline structure and morphology (roughness) of the final GaAs epilayers. To understand the molecular-level mechanisms driving the 3D growth of GaAs on Si, we applied a DDA nucleation model sensitive to the different reactivity of the group-III precursor molecules onto the GaAs nanoislands and Si surfaces. The model predicts a linear evolution of GaAs growth rate with the deposit surface coverage (since the early stages of nanoisland growth and coalescence until the formation of a continuous layer) in good agreement with ACS Paragon Plus Environment

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experimental observations, under the assumption that the direct impingement of a specific group-III species (likely Me3Ga) onto the GaAs nanoisland surface dominates the nucleation and overall growth rate; the contribution of Me3Ga adsorption onto the As-passivated (111)Si surface shows instead, negligible. The letter hypothesis agrees with the observed lack of continuous nanoisland nucleation after the first few minutes. Finally, our DDA model allowed to estimate the effective reactive sticking coefficient of Me3Ga molecules for their direct impingement onto GaAs at 400°C, which turned out equal to

1=2.8210-5, its small value being compatible with the large precursor steric hindrance and competitive Me surface adsorption expected at the low temperature of present growth experiments.

Acknowledgments The authors wish to acknowledge R. Mucciato and A. Pedio for AFM and FE-SEM observations respectively, and F. Marzo for assistance to the growth experiments and artwork preparation.



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For Table of Contents Use Only

PAPER TITLE: Shape, size evolution and nucleation mechanisms of GaAs nanoislands grown on (111)Si by low temperature metalorganic vapor phase epitaxy AUTHORS: Ilio Miccoli, Paola Prete, and Nico Lovergine TOC Graphics:

SYNOPSIS: The shape, size and nucleation mechanisms of GaAs nano-islands grown by low temperature metalorganic vapor phase epitaxy on As-stabilized (111)Si are studied by combining experimental FE-SEM, AFM and XRD measurements of the nano-island morphology/size evolution with analysis by a deposition-diffusion-aggregation nucleation model: it is demonstrated that, after an initial heterogeneous nucleation stage, the direct impingement of trimethylgallium (Me3Ga) molecules onto the nano-islands surface dominates their further growth and coalescence, the contribution of molecules adsorbed onto the As-stabilized (111)Si


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free-surface remaining instead negligible, due to the reduced reactivity of the As-passivated Si surface.


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(30) Kajikawa, Y.; Noda, S. Growth mode during initial stage of chemical vapor deposition. Appl. Surf. Science 2005, 245, 281-289. (31) Kajikawa, Y.; Tsuchiya, T.; Noda, S.; Komiyama H. Incubation time during chemical vapor deposition of Si onto SiO2 from silane. Chem. Vap. Deposition 2004, 10, 128-133. (32) Stringfellow, G.B.; Organometallic Vapor-Phase Epitaxy: Theory and Practice. Academic Press (San Diego, London), 1999. (33) Den Baars, S.P.; Maa, B.Y.; Dapkus, P.D.; Danner, A.D.; Lee, H.C. Homogeneous and heterogeneous thermal decomposition rates of trimethylgallium and arsine and their relevance to the growth of GaAs by MOCVD. J. Cryst. Growth 1986, 77, 188-193. (34) Weyers, M.; Sato, M. A comparison of the growth of GaAs and GaP from trimethylgallium. Jpn. J. Appl. Phys. 1995, 34, 434-441. (35) Scherrer P. Bestimmung der Grosse und der Inneren Struktur von Kolloidteilchen Mittels Rontgenstrahlen. Nachr. Ges. Wiss. Göttingen, Math-Phys. Kl. 2 1918, 96-100. (36) Noteworthy is that the short residence time ads associated to a rapid desorption of group-III species from the Si surface also implies a reduced surface diffusion length, as

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(37) Creighton, J.R. Chemisorption and decomposition of trimethylgallium on GaAs(100). Surf. Sci. 1990, 234, 287-307.


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Shape, size evolution and nucleation mechanisms of GaAs

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