gaas nanowires grown by movpe

Phys. Status Solidi B 247, No. 6, 1294–1309 (2010) / DOI 10.1002/pssb.200945495 p s sb

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eature Article

basic solid state physics

GaAs nanowires grown by MOVPE F

Jens Bauer*,1,2, Hendrik Paetzelt1,2, Volker Gottschalch1, and Gerald Wagner3

1University of Leipzig, Institute of Inorganic Chemistry, Semiconductor Chemistry Group, Johannisallee 29, 04103 Leipzig, Germany2University of Leipzig, Institute of Mineralogy, Crystallography and Materials Science, Scharnhorststrasse 20, 04275 Leipzig,

Germany3Leibniz Institute of Surface Modification, Permoserstrasse 15, 04318 Leipzig, Germany

Received 26 October 2009, revised 2 February 2010, accepted 19 February 2010

Published online 24 March 2010

PACS 61.46.Km, 62.23.Hj, 64.70.kg, 64.75.Jk, 78.55.Cr, 81.15.Gh

*Corresponding author: e-mail [email protected], Phone: þ49 341 2353311, Fax: þ49 341 2353400

GaAs nanowire (NW) growth was studied by metal-organic

vapour phase epitaxy (MOVPE). The vapour–liquid–solid

(VLS) mechanism with gold-based alloy particles and the

selective-area growth (SAG) mechanism on electron beam

lithographically prepared SiNx/GaAs mask structures were

applied. A special focus is set on thermodynamic aspects of the

VLS process. The alloy particle formation and the influence of

MOVPE growth parameters on the growth rate and the GaAs

NWmorphology are examined. Furthermore, the improvement

of the real structure with particular interest on the twin

formation is studied. Besides the commonly used continuous

VLS growth mode also a pulsed VLS growth mode with

alternating precursor supply is reported. Based on photo-

luminescence measurements the effect of strain in core/shell

NW structures is confirmed. For the SAG mechanism the

MOVPE growth parameters are determined and the real

structure is described.

� 2010 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

1 Introduction Nanostructures have remarkable pro-perties and a high potential to realize new nanoelectronic andnanooptoelectronic devices. Presently much research inter-est is concentrated on essential properties such as crystalstructure and crystal perfection, morphology, definiteheterojunctions, composition and doping. The most usedgrowth methods are the molecular-beam epitaxy (MBE) [1]and the metal organic vapour phase epitaxy (MOVPE) [2].Different MOVPE growth techniques were applied tofabricate various nano- and microstructures.

A novel class are nano- and microtubes. These nano-structures can be formed by releasing a strained multilayersystem from the substrate by selective etching of a sacrificiallayer. The multilayer system can be manufactured by meansofMBE [3] orMOVPE [4, 5]. Hence, III–V nanotubes with aspecific geometry can be realized. The defined strain state of

incorporated quantum-well structures offers new possibili-ties to tailor the electronic properties.

The preparation of freestanding III–V nanowires (NW)with diameters in the range from a few atomic distances tosome hundreds of nanometres is typically based on thevapour-liquid-solid (VLS) growth mechanism or the selec-tive-area growth (SAG) mechanism. Both techniques utilizethe epitaxy with lateral size limitation.

In the SAGmechanism limited regions (50 nm–1mm) onthe substrate surface are provided by nanolithography. Withappropriate growth conditions selective growth of a specificcrystal plane is favoured in MOVPE. This uniaxial growthenables the fabrication of NW from the limited substrateregions. Additionally, the SAG technique enables thefabrication of well-ordered NW arrays without the use of acatalyst material [6–8].


Phys. Status Solidi B 247, No. 6 (2010) 1295

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Figure 1 SchematicdiagramofVLSGaAsNWgrowthusingPGMandCGM.ThePGMdeposition cycle consisted of separate pulses ofTMG and arsine. Each TMG and arsine pulse was followed by a 2 sand 5 s H2 pulse, respectively. In CGM constant flows of TMG andarsine were simultaneously supplied instead.

The VLS growth involves the usage of initiators orcatalysts such as Au or other metals. The lateral sizelimitation (<10 nm–1mm) is realized by the wetting of aliquid nanodroplet on top of the substrate [9]. The dropletacts as a reaction vessel. In particular, the gaseous reactantsare absorbed at the droplet surface. After diffusion throughthe liquid agent the nucleation and crystal growth occurs atthe interface to the underlying substrate. The applied growthconditions are usually beyond the optimal MOVPE layergrowth region, i.e. the growth temperature is about 250 8Cbelow. Hence, the crystal growth from the nanodropletresults in the formation of an NW while the liquid is alwayssituated on the top. We have investigated a great variety ofmaterials for the fabrication of NW using the VLS growthmode, e.g. ZnO [10], GaN [11], GaAs [9] and InAs [12, 13].

Defined positioning [6, 9] as well as the realization ofradial and axial heterostructures [6, 14] have been demon-strated with both growth techniques.

The technological effort of the VLS mechanism is muchlower than of the SAGmechanism. Otherwise, the use of theliquid nanodroplet considerably increases the complexity ofthe growth system. Presently, the general understanding ofthe basic growth processes has not yet been solved for bothgrowth mechanisms.

In this paper, we report on the MOVPE growth of GaAsNW using the VLS and the SAG technique. We focus onfundamental aspects of the VLS mechanism. In particular,the droplet formation process (reaction of Au with GaAssubstrate and precursormaterial), theGaAsNWmorphologyand the crystal structurewith respect to differentVLSgrowthmodes are presented. We determined the optimal conditionsfor GaAs NW growth via the SAG mechanism to achieveuniform NW arrays. Additionally, we depict a growth routetowards complex nanostructures using a special combinationof SAG and VLS growth.

We analysed the morphology and structure of GaAs NWusing X-ray diffraction (XRD) methods, scanning electronmicroscopy (SEM) and high-resolution transmission elec-tron microscopy (HRTEM). The structural properties andcomposition of the gold droplets and the epitaxial relation-ships of Au droplet/GaAs NW/GaAs substrate weredetermined by a Philips X’Pert X-ray diffractometeroperating with copper Ka radiation. HRTEM provided theinformation about crystal structure and lattice defects. Inparticular, the formation of twins, and segments of sphaleritetype and wurtzite type structure were determined withrespect to several growth parameters.

2 The VLS growth mechanism The growth ofsemiconductor one-dimensional structures via the VLSmechanism was first observed in 1964 by Wagner and Ellis[15]. A fundamental semi-empiric description of the basicprocesseswas developed byGivargizov [16]. In this view thenanodroplet was assumed to be completely in liquid state andthe precursor gases were absorbed at the droplet’s surface.TheNWgrowth ratewas determined by the kinetic processesat the droplet/substrate interface. In recent years new

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experimental results were obtained suggesting a differentview. Seifert et al. [17] investigated the NW growth rate inMOVPE experiments. Based on their results the transport ofthe precursor molecules to the nanodroplet is suggested todetermine the growth rate. However, a consistent viewrespecting thermodynamics and kinetics inside the liquiddroplet has not yet been proposed.

In the following sectionswe focus on theVLSGaAsNWgrowth via the continuous growth mode (CGM) and thepulsed growth mode (PGM).We applied low-pressure metalorganic vapour phase epitaxy (LP-MOVPE) in a commercialAIX200 reactor with a total pressure of 5 kPa, a total gas flowof 7000 sccm and H2 atmosphere. We used trimethylgallium(TMG) and AsH3 (arsine) or tertiarybutylarsine (tBA).

The standard CGM procedure was as follows (seeFig. 1): Bymeans of thermal evaporation a gold film of a fewnanometre thickness (typical 6 nm) was deposited on (-1-1-1)AsGaAs. The sample was heated to 400–700 8C in AsH3/H2

atmosphere. After a period of 2min (stabilization of thegrowth temperature) TMG was turned on for further 2min(GaAs growth step). Usually V/III ratios of 10–40 (variationof the arsine supply) and a TMG partial pressure of 0.36 Pawere used. Finally the sample was cooled down maintainingthe arsine supply.

Prior to the growth via the PGM, we deposited a thingold film with a thickness in the range from 1 to 20 nm on a(-1-1-1)As GaAs substrate. During the PGM the TMG(pTMG¼ 0.36 Pa) and arsine (pAsH3¼ 7.1 Pa) were suppliedalternately into the reactor at a growth temperature of 480 8C(Fig. 1). The growth was started after a temperaturestabilization step of 2min. Then the formation of the liquidAu–Ga droplets was initiated by a 20 s TMG step. In thefollowing pulse sequence TMG and arsine were alternatelyswitched every 10 s. Arsine was provided during coolingdown.


1296 J. Bauer et al.: GaAs nanowires grown by MOVPEp

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Figure 2 (online colour at: www.pss-b.com) SEM images (b–d)(-110) cross-section, (a) tilted 208 towards the sample normal:(a, b) morphology of a gold film on (-1-1-1)As GaAs substrate afterthermal treatment (2min, 600 8C, AsH3/H2 atmosphere); (c, d)morphology of GaAs NW (CGM, 2min, 480 8C, V/III¼ 10,pTMG¼ 0.36 Pa) on (-1-1-1)As GaAs. After the thermal treatment/NWgrowth the coolingdownprocedurewasvaried: (a, c) arsinewasprovided during cooling down; (b) arsine was provided from 600 to480 8Cand then turnedoff; (d)noarsine supplyduringcoolingdown.V/2Q-scans of the samples (a–d). For comparison an untreated goldfilm on (-1-1-1)As GaAs substrate is also shown.

2.1 Droplet formation and Au–Ga–As reactionThe Au–Ga–As alloy droplet is essential for VLS NWgrowth.We analysed the droplet formation during heating upto growth temperature in AsH3/H2 atmosphere. In particularthe transient effects in the liquid Au–Ga–As alloy prior toand after theGaAs growth stepwere focussed. The gold filmsand their interactions with both the GaAs substrate as well asthe precursormaterialswere characterized by SEMandXRDtechniques (V/2Q-diffraction, F-scans).

The V/2Q-XRD scan of an evaporated gold film on(-1-1-1)As GaAs substrate is shown in Fig. 2. All peaks couldbe indexed by gold (JCPDS file Nr. 04-0784) and GaAs(JCPDS file Nr. 80-0003). A broad 111 Au reflection(FWHM� 1.48) was observed, which indicates a preferred[111] out-of-plane orientation of slightly misaligned goldgrains. During heating up gold islands are formed from thegold film (Fig. 2a and b). Without arsine supply the islandsmelt at 525� 25 8C [18, 19]. After the thermal treatment at600 8C with arsine supply (Fig. 2a) the 111 Au reflectionnarrows to FWHM� 0.18. A similar effect is obtained for thealloy droplets of GaAs NW when the arsine supply is alsoprovided after the NW growth step during cooling down toRT (Fig. 2c). As will be shown later, the alloy droplet is incomplete liquid state duringNWgrowth at V/III< 20. Basedon the similar XRD results (Fig. 2a and c) a liquid alloyformation between the gold islands and the GaAs substratematerial is proposed at 600 8C. This view is supported by thesimilar polycrystalline alloy particle morphology with arough surface structure, which is formed in both cases duringcooling down with arsine supply. Another situation wasobserved when the arsine supply was turned off duringcooling down (Fig. 2b and d). The particle surface is smoothwith well-defined crystal facets in that case. Some islandswere single crystalline. The phase analysis by V/2Q-XRD(Fig. 2b and d) showed the formation of the Au7Ga2 phase(b-Au–Ga phase) instead of pure Au. This experimentalresult is a proof for the Au–GaAs alloy formation. Thecrystallization of a liquid Au–Ga alloy has been previouslyinvestigated [19, 20]. Below 500 8C the formation of thea-Au–Ga phase (fcc gold with solved gallium via substi-tution) is reported, which is transformed peritectically intob-Au–Ga (Au7Ga2) at 390–400 8C.

To get informations about the azimuthal alignmentasymmetric V/2Q- and F-scans were taken from the goldfilm prior to and after the GaAs growth. Figure 3 representsthe F-scans of a 20 nm thick gold film deposited on(-1-1-1)As GaAs.

The three asymmetric GaAs substrate peaks of the 331family occur each 1208 and are directed towards theazimuthal h112i directions. Hence, they reflect the threefoldin-plane symmetry of the (-1-1-1)AsGaAsplane. TheF-scansof the Au film were taken from the 331 and 113 reflectionsshowing the sixfold symmetry of (111) Au. A 308 rotationwas observed between the 331 peaks of GaAs and Au. As aresult we obtained the following epitaxial relationship:

� 20

ð111ÞAu k ð-1-1-1ÞAs GaAs; h110iAu k h211iGaAs:

10 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

The same sample was treated for 2min at T¼ 650 8C inAsH3/H2 atmosphere (Fig. 4).

The additional 331 GaAs peaks (marked by the dots)indicate the interface reaction of Auwith GaAs at 650 8C. Asa result, we suggest the formation of a liquid Au–Ga alloyand the recrystallization to GaAs. Since the 331 GaAs peaksappear, twin formation is expected during recrystallizationof GaAs. The residual gold particles showed a different in-plane orientation after recrystallization (marked by thearrows):

ð111ÞAu k ð-1-1-1ÞAsGaAs; h110iAu k h110iGaAs:

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Figure 3 F-Scans of a gold film deposited on (-1-1-1)As GaAs.

Figure 5 V/2Q-Scans of 6 nm Au/(0001) Al2O3 specimens afterthe TMG treatment at 480 8C for different durations.

This orientation relationship has also been measured forgold on GaAs NW samples. The 113 Au reflections showedfurther XRD features which were symmetrically turned byabout 108 from the main orientation. Based on this result wesuggest the existence of the Au7Ga2 phase at the Au/GaAsinterface [21]. This result is in good agreement with theappearance of the Au7Ga2 phase in theV/2Q-scan (Fig. 2b).

We suggest the formation of an Au–Ga alloy layer at theAu/GaAs interface during the heating step to 600–650 8C.Simultaneously the arsine supply results in a separation ofthe liquid alloy into solid Au and GaAs.

Figure 5 shows XRD V/2Q-scans (ex situ after coolingdown) of the reaction of a gold film (6 nm thickness) on(0001) Al2O3 with TMG at 480 8C. The TMG partialpressure was 0.36 Pa. After 20 s TMG supply the 111 Aupeak at 2Q¼ 38.28 disappeared and polycrystalline Au–Gaalloys with lowmelting points (390–420 8C)were generated.Based on these experiments we suggest the formation of anAu–Ga melt from the gold film after sufficient exposure toTMG.

Figure 4 F-Scans of a gold film on (-1-1-1)As GaAs after thermaltreatment at 650 8C in AsH3/H2.

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2.2 The continuous VLS growth mode Themorphology of GaAs NW grown by the CGM can becontrolled by the MOVPE parameters and the growthsequence. In the growth temperature range from 400 to530 8C the GaAs NW exhibited a columnar morphology(Fig. 6a). At this temperature efficient MOVPE layer growthis not expected [22]. The situation changes for higher growthtemperatures. For T> 530 8C an additional MOVPE layergrowth is observed at the NW side facets. As a result, atapered GaAs NW morphology was found (Fig. 6b).

A more detailed investigation of the columnar NWmorphology was performed using the CGM standard growthprocedure at T¼ 480 8C. After GaAs NW growth character-istic parts were observed (sketch in Fig. 7a): (1) TheNW footregion is formed during initial NW growth. In this period thedroplet composition and the Au/GaAs interface geometrychange. (2) Then the composition and thewetting of the alloydroplet reaches a (kinetically stabilized) quasi-stationaryequilibrium state resulting in the column-like NW bodyregion. (3) After turning off the TMG supplymaintaining thearsine flow the laterally extendedNWneck region is formed.



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Figure 6 (online colour at: www.pss-b.com) Effect of the growthtemperatureon theGaAsNWmorphology. (a)T¼ 480 8C:VLSNWgrowth resulting in a columnar morphology and (b) T¼ 600 8C:additional gas-phase epitaxy on the side facets of the growing NWresults in a tapered NW morphology.

Figure 7 GaAs NW using the CGM standard growth procedure atT¼ 480 8Cwith a 6 nm thick goldfilmon (-1-1-1)AsGaAs substrate.(a) SEM image ((-110) view) after cooling down with AsH3/H2

atmosphere. Sketch: Habitus of a single NW responding to themajority of NW after the CGM standard growth procedure. Com-pletelyseparateddroplets (leftNW)formveryrarelyafter theappliedimmediate cooling down. (b) SEM image ((-110) view) of the samesample after a further annealing for 2min at 480 8C in AsH3/H2

atmosphere. After annealing the arsine supply was turned off andthuscoolingdownwasperformed inH2atmosphereonly.Partsof theneck regionwere resolvedby thedroplet. (c)TEMbright-field image((111) view) of a typical cross-section of GaAs NW showing athreefold symmetry.

In this growth step the necessary amount of gallium forGaAsepitaxy originates from the reservoir in the Au–Ga droplet.When TMGand arsine are turned off simultaneously no neckregion appears.

The alloy particle is located always on top of the GaAsNW. Applying a slow cooling down step with further arsinesupply the Au–Ga alloy is completely separated into a-Au–Ga phase and the GaAs neck region. The resulting alloyparticle showswell-defined facets (see left NW in Fig. 7a). Incontrast, an immediate cooling down results mostly in apolycrystalline solidification of the liquid Au–Ga melt inAu–Ga alloy phases and GaAs. The particle morphologyshows no crystalline facets in that case (see right NW inFig. 7a). Because of the polycrystalline solidification theepitaxial neck region is much less extended after immediatecooling down (see the different Dz in Fig. 7a).

In Fig. 7b the sample was heated up again with arsinesupply and cooled down without arsine. The neck region isalmost completely solved by the alloy particle. From theextension of the neck region (left NW in Fig. 7a) the alloycomposition during GaAs NW growth was estimated.Assuming a residual particle of pure gold the galliumfraction during NW growth was about xGa¼ 26 at.% (V/III¼ 10). Since the residual particle possibly contains someamount of gallium, this value can be even higher. However,this composition is in liquid state at growth temperature.

Based on the correlation of the GaAs NW morphologyparts with the VLS process we developed a growth sequence


to create in situ markers. In particular, we applied thetransient effects of a growth interruption. In Fig. 8 the TMGsupply was interrupted three times for 5min each. Duringturning off the TMG flow the situation was similar to theformation of the neck region accompanied by the increase ofthe NW diameter. When the TMG was turned on again thequasi-stationary growth state was reached after a shorttransient growth.As a result, a local diameter increase (in situmarker) was left by the TMG interruption. With increasinginterruption period from 5 s to 5min the diameter increase isextended.

The experiments further showed a maximal markerextension after a TMG interruption period of 5–30min. As aresult we suggest a complete macroscopic separation of thealloy droplet. This time scale is much higher than the alloyformation at the beginning of the GaAs NW growth at thesame arsine partial pressure. In particular, experiments

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Figure 8 (onlinecolour at:www.pss-b.com)GaAsNWwith in situmarkers at three defined positions (indicated by the arrows). In situmarkers were realized by the application of specific growth inter-ruptions. Inset: When the TMG supply is turned off the diameterincreases similar to the growth of the neck region. Then the TMGsupply is turned on again. The quasi-stationary equilibrium growthconditions are reaching and hence the diameter decreases again.

considering the temporal evolution of the GaAs NW lengthat V/III¼ 10 showed that the initial GaAs NW growth,i.e. the alloy formation, is completed after about 50 s [9].With increasing arsine partial pressure (V/III ratio) theinitial growth period increases and the separation perioddecreases.

2.3 The pulsed VLS growth mode The PGM wasstarted with a 20 s TMG pulse. As a result liquid (l) AumGandroplets (see also Fig. 5) are formed from the gold layer. Thisalloy formation can be formally described by the followingreaction:

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mAus þ nGaðCH3Þv3 þ 3nH ! AumGaln þ 3nCHv

4

(2.1)

After the TMG pulse an arsine pulse followed in thesecond step.We expect the complete arsine decomposition atthe droplet surface. The liquid Au–Ga alloy gets super-saturated and as a result GaAs crystallizes. The reaction ofthe Au–Ga alloy with arsine can be formally described by

Figure 9 SEM image ((-110) cross-section) of GaAs NW on(-1-1-1)As GaAs using PGM. The individual pulse steps are visibleby contrast-rich lamella (indicated by markers). The gold filmthickness was 3 nm.

AumGaln þ AsHv

3 ! GaAss þ AumGaln�1 þ

3

2H2;

(2.2)

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and finally

AumGaln þ nAsHv

3 ! nGaAss þ mAus þ 3n

2H2: (2.3)

The starting composition of the droplet is defined by thegold layer thickness, the TMG partial pressure and the pulselength. During the arsine pulse the composition and thesupersaturation change. As a result of the GaAs crystal-lization the gallium fraction in the Au–Ga melt decreases.After the alternating pulse sequence the growth process isfinished with an arsine treatment resulting in the solidifica-tion of the droplet (recrystallization of gold). Then thesamples were cooled in an AsH3/H2 atmosphere.

Starting from a 20 nm gold layer the NW diameters were15–100 nm. The mean vertical growth rate (in the arsenicpulses) ranged between 1 and 5 nm/s. The GaAs NW wascovered with well-faceted gold droplets at the top (Fig. 9).Diffraction patterns obtained after NW growth at 480 8Cshowed strong 111 GaAs and 111 Au peaks indicating thatpure gold or goldwith a small concentration ofGa (a-Au–Gaalloy) crystallized epitaxially. Additionally, F-scans of the331 GaAs and 311 Au diffraction peaks suggest thefollowing orientation relationship:

ð-1-1-1ÞAsGaAs k ð111ÞAu; h110iGaAs k h110iAu:

2.4 Thermodynamic and kinetic aspects of theVLS mechanism Givargizov [16] used a semi-empirictheory to describe the VLS mechanism. In particular, thegrowth process is divided into four partial steps: (1)precursor molecule transport in the gas phase to the surfaceof the liquid droplet; (2) condensation at the droplet’ssurface; (3) diffusion through the droplet; (4) nucleation andcrystal growth at the liquid/solid interface. The steps 1 and 4determine the GaAs NW growth rate in MOVPE.



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In particular, the growth rate of GaAs NW has beenalready investigated with respect to important growthparameters, i.e. the growth temperature, the V/III ratio andthe growth time [9]. Thus, the growth rate is limited bykinetics in the VLS growth and shows a strong dependenceon the V/III ratio. At low V/III ratios (V/III< 20 in ourstandard procedure at 480 8C) the slowest kinetic step is thenucleation and crystal growth at the liquid/solid interface.Since growth kinetics depends on the supersaturation of thenanodroplets, the Gibbs–Thomson effect has to be taken intoaccount. In contrast at high V/III ratios (V/III> 20) thesource molecule transport to the droplet is the growth ratelimiting step. In particular, the surface diffusion of TMGdecomposition products on the side facets of the alreadygrown GaAs NW is the dominating limitation process. Thegrowth rate of GaAs NW is about 7–11 nm/s at V/III¼ 10–40. As a result of the influence of the Gibbs–Thomson effectat low V/III ratios and the surface diffusion at high V/IIIratios the growth rate usually shows a strong dependence onthe NW diameter. However, at an intermediate V/III ratioof 20 we achieved a diameter-independent GaAs NWgrowth rate [9].

The VLS growth of GaAs NW and the LPE growth ofGaAs layers show similarities: (1) The growth systemconsists of a gas phase, a liquid melt and the growing crystal.(2) At least one precursor material can be provided from thegas phase. Themelt saturation and the LPEgrowth behaviourhave been studied for GaN [23], N-doped GaP [24], GaP andGaAs [25]. (3) Since the composition ratio Ga/As in the meltis much bigger than unity the supersaturation during crystalgrowth is determined by arsenic [26, 27]. The meltcomposition changes corresponding to the Au–Ga–As phasediagram. (4) Similar to the present droplet material for theVLSmechanism also different solvents have been examinedfor LPE growth to realize low growth temperatures, ahigher group-Vmaterial solubility and sharp heterojunctions[23, 27].

However, because of the nanoscopic dimensions of theliquid phase duringVLSNWgrowth the LPE view cannot becompletely applied to describe the VLS growth process.There are two main differences between both growthtechniques: (1) LPE is usually performed with a largevolume of a gallium-rich melt. Arsenic can be supplied fromthe gas phase. The amount of gallium necessary for crystalgrowth is taken from the gallium reservoir in the melt. Thissituation is similar to (a) the VLS growth of the neck regionin the CGM and (b) the arsenic pulse in the PGM. In bothcases the TMG supply is turned off. In contrast, duringcontinuous NW growth both elements, gallium and arsenic,need to be supplied from the gas phase. Otherwise the meltreservoir of galliumwas rapidly consumed. As a result of thesupply of both precursor materials, a quasi-stationary meltcomposition is reached during NW growth via the CGM. (2)In LPE, the arsenic diffusion through the macroscopic meltlayer is the slowest kinetic step and determines the GaAslayer growth rate. The situation inVLS growth is remarkablydifferent. Assuming a GaAs NW growth rate of 10 nm/s [9]


the growth of one GaAs monolayer in [-1-1-1]As direction(d¼ 0.32639 nm) takes about 30ms. The effect of diffusioncan be roughly estimated in a one-dimensional view. Thediffusion length z ¼

ffiffiffiffiffiffiffiffi2Dt

pdepends on the diffusion

coefficient D and diffusion time t. Diffusion data of arsenicin an Au–Ga–As melt are lacking, but usually the diffusioncoefficient is D¼ 10�4–10�5 cm2/s in metallic melts (Ref.[28], p. 229). A typical droplet for GaAs NW growth has adiameter of 100 nm and below. The one-dimensionaldiffusion over a distance of z¼ 100 nm with D¼ 10�5 cm2/s takes about 5ms. With a higher diffusion coefficient thistime gets even shorter. As a result we state, that the liquid/solid interface kinetics determines the NW growth rate.Diffusion can be neglected in the nanoscopic droplets.Hence, the composition and the temperature within the meltare spatially constant during VLSNW growth. Based on thisestimation, the vapour/liquid interface rather than the liquid/solid interface is expected to be situated near the thermo-dynamic equilibrium. In contrast, in LPE the liquid/solidinterface is usually situated near thermodynamic equilibriumconditions.

The thermodynamic equilibrium at the vapour/liquidinterface implies that the melt composition varies accordingto the vapour/liquid distribution coefficient. Because of thenanoscale dimension, additionally the Gibbs–Thomsoneffect has to be considered. Givargizov [16] presented asemi-empiric model to describe the NW growth rate withrespect to the NW diameter. This model is completelyconsistent with our understanding of VLS growth at V/III< 20. After Givargizov the NW growth rate follows theequation:

R ¼ b4svlv

kBT

� �21

dc� 1

d

� �2

; (2.4)

where b is a kinetic factor. The critical diameter dc, thevapour/liquid interface tension svl and the arsenic molarvolume v depend on the particular growth conditions. T isthe absolute growth temperature. The appearance of thecritical diameter is a result of the Gibbs–Thomson effect.For NWwith d< dc no NWgrowth can be obtained, becausethe vapour pressure over the nanodroplet is higher than thesupplied precursor partial pressure. At V/III ratios< 20 theGaAs NW growth rate follows the Gibbs–Thomson effect[9] and can be expressed by Eq. (2.4) (Fig. 10). In contrast toother III–V NW growth results [29, 30] we obtained anexcellent agreement between experiment and the Gibbs–Thomson relation in the whole diameter range (20–350 nm).For the particular growth conditions in Fig. 10 we obtained acritical diameter of about 14 nm. With svl¼ 1.02 J/m2,v¼ 11.1 cm3/mol and the slope in Fig. 10 the kinetic factorwas estimated: b¼(2.4� 0.1)� 10�6 cm/s. This value forMOVPE is about four orders of magnitude smaller than forHVPE conditions [16].

A similar result has been obtained for the PGM, too(Fig. 11). The growth rate was estimated from theNW lengthand the total duration of arsine supply (sum of all pulses). In

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Figure 10 (onlinecolourat:www.pss-b.com)GrowthrateofGaAsNW versus the NWdiameter in the CGM standard growth at 480 8Cand V/III¼ 10 (pTMG¼ 0.36 Pa). The characteristic follows theGibbs–Thomson relation. The line shows the fit of the experimentaldata after Eq. (2.4).

Figure 12 Estimated isotherms (squares 480 8C, stars 600 8C) inthe VLS growth temperature range. The eutectic line (dotted curve)limits the GaAs deposition in the system AuxGayAsz.

agreement with the increased supersaturation (increasedarsine partial pressure of 7.1 Pa) the critical diameter (about9 nm) was decreased compared to the CGM growth in Fig. 9.To consider the dynamic changes of the growth rate(originating from the cyclic composition and hence super-saturation changes) more detailed investigations will benecessary in future.

To account for the effect of thermodynamics in theGaAsNW growth process, the Au–Ga–As phase diagram has to beconsidered. Based on the phase diagram the dropletcomposition range can be defined and the growth conditionsrequired for successful NW growth can be discussed.Unfortunately, exact thermodynamic data for the tempera-ture range of 350–600 8C are lacking. Hence, we used DTAdata reported by Panish [31] and Leonhardt/Kuhn [32] toestimate the phase diagram for temperatures up to 600 8C.Our calculations are based on the following values:DSFGaAs ¼ 16.64 cal/mol/K, TF

GaAs ¼ 1511K, aGaAs: 5160–

Figure 11 (online colour at:www.pss-b.com)Meangrowth rate ofGaAsNW versus theNWdiameter in thePGMgrowth at 480 8CandpTMG¼ 0.36 Pa, pAsH3¼ 7.1 Pa from a 20 nm gold film. The char-acteristic follows theGibbs–Thomsonrelation.The lineshows thefitof the experimental data after Eq. (2.4).

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9.16 T (cal/mol), aAu–As: 4000–2.5 T (cal/mol), aGa–Au:�12 540 cal/mol.

Figure 12 shows the eutectic line (circles) in the primarydeposition region of GaAs. Furthermore, the estimatedisotherms at 480 8C (squares) and 600 8C (stars) are included.For both temperatures the GaAs solubility in the liquid Au–Ga melt increases with increasing gold fraction (decreasingabsolute amount of gallium at a constant absolute amount ofgold). During GaAs NW growth the melt composition issituated beyond the thermodynamic liquid/solid equilibriumconditions, i.e. the melt is supersaturated.

The quasi-stationary equilibrium composition isreached, when the same numbers of gallium and arsenicatoms are absorbed at the droplet surface as needed for theGaAs crystallization. In contrast, during the initial stage ofVLS growth (foot region) more gallium atoms are suppliedthan necessary for GaAs crystallization with a gold-richdroplet composition. This excess of gallium results in achange of the droplet composition towards a higher galliumfraction. Since the arsenic supersaturation increases withincreasing gallium fraction more GaAs crystallizes. As aresult, the quasi-stationary equilibrium state is reachedaccompanied by an increase of the growth rate. This processis not only affected by the external precursor supply, but alsoby the thermodynamic equilibrium conditions at the vapour/liquid interface, i.e. the vapour/liquid distribution coeffi-cient, and the Gibbs–Thomson effect. Note the distributioncoefficient depends on the melt and the vapour compositionsand the Gibbs–Thomson effect on the droplet size.Furthermore, the increasing droplet volume with increasinggallium content has to be considered. Hence, the Gibbs–Thomson effect is reduced and the precursor materialabsorbing area is increased with increasing liquid galliumfraction. It is not yet clear if besides the influence of theGibbs–Thomson effect the growth rate is additionallyenhanced by an increased supersaturation at higher goldfractions under quasi-stationary growth conditions.



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During GaAs NW growth via the CGM TMG and arsineare continuously supplied to the system. Corresponding tothe applied gas-phase ratio (V/III ratio) themelt compositionvaries. With increasing V/III ratio the gold fraction and thearsenic fraction are expected to increase following the trendof the isotherm (Fig. 12). At high arsine partial pressures themelt composition is limited by the eutectic line on the rightside in Fig. 12. However, GaAs NW growth was alsoobtained at high V/III ratios> 20 [9]. Another growth modelhas to be found to describe the different experimental resultsat high V/III ratios where the melt composition is fixed to theeutectic composition. Droplets with solid gold parts have tobe considered in this case [33].

2.5 Real structure The GaAs NW have sphaleritetype structure [9]. The side facets are mainly {11-2} polarplanes (Fig. 7c). The B-planes are by a factor of 1.6–2.0larger than the A-planes resulting in a threefold symmetry ofthe NW cross-section. In [-1-1-1]As growth directionrotational twin sections are formed during NW growth. Aphenomenological model for the twin formation waspresented previously [9].

Characteristic notches determine the surface structure oftwinned NW (Fig. 13). Because of the strong geometricrules, which depend on the crystallography of the sphaleritestructure, the variation of the NW geometry parametersduring twin growth can be determined. We denoted thewidths of the two polar side planes {-1-12}Ga and {11-2}Asby a and b, respectively (see Fig. 7c). Each twin sectionstarts with the formation of an inclined {-1-11}Ga plane at the{-1-12}Ga facets (region II in Fig. 13).Width a increases and

Figure 13 (online colour at: www.pss-b.com) Reconstruction ofthe twin structure of a single GaAs NW using the geometric model.The cross-section area is shown for the twin region marked by thedashed lines.


width b decreases in region II during growth (lower diagramin Fig. 13). The formation of a twin boundary is accompaniedby a change in the stacking along the [-1-1-1]As axis from‘AaBbCbAaBbCc’ to ‘AaCcBbAaCcBb’. As a result, thepolar planes of a twin section (region III in Fig. 13) have anopposite polarity compared to the matrix structure. Note thecorresponding switch of the widths a and b at each twinboundary in the diagram in Fig. 13. Width a increases andwidth b decreases also during growth in region III (lowerdiagram in Fig. 13). The side facets consist of T{11-1}As andT{-1-11}Ga planes in the twin region. After the second twinboundary in region IV the side facets reach their primaryvalues of region II. In region IV {001} and {11-1}As planesare formed.

The upper diagram in Fig. 13 shows the development ofthe NW cross-section area A ¼ ð

ffiffiffi3

p=4Þða2 þ b2 þ 4abÞ

during the different twin regions. Based on our theoreticdescription we can state that the NW cross-section areareaches a minimum at the twin boundaries. Since the NWcross-section area has a remarkable influence on the dropletwetting on top of the NW, this result further indicates thattwin formation is related to wetting instabilities. Inagreement we previously discussed a perturbation of themechanical equilibrium of the droplet surface to be theinitiator for twin formation [9].

We developed a modified growth sequence by which theincorporation of twin sections during NW growth could besignificantly reduced. The growth of theGaAsNW in Fig. 14was started with the MOVPE standard procedure at 480 8Cand V/III¼ 10. After 5min the arsine partial pressure wasabruptly increased. Then the GaAs NW growth wasperformed with V/III¼ 40 for further 5min. The SEMexperiments in Fig. 14 show a complex surface structure ofthe GaAs NW in their lower half. Hence, the formation ofclose twin sections is suggested at V/III¼ 10. In contrast, theupper NW halves consist of flat side facets. By means ofTEM investigations we verified that in the upper NW partonly very isolated thin twin sections are incorporated.

Recently, Joyce et al. [34] reported on the reduction oftwin formation at high V/III ratios. GaAs NW growth withour standard procedure always shows a high frequency oftwin sections, also at high V/III ratios> 40. Instead of thatthe use of the two-step procedure, i.e. starting growth at a lowV/III ratio and switching to the higher V/III ratio sub-sequently, allowed an almost twin-free GaAs NW growth.Another procedure with a special temperature profile toreach twin-free GaAs NW growth has been described [35].However, we were not able to reproduce these results.

The GaAs NW grown by the CGM have sphalerite typestructure in MOVPE. In contrast, a characteristic feature ofthe MBE NW growth is the formation of the hexagonalwurtzite (w) structure. Additionally, the structural propertiesof the PGM NW were determined with HRTEM. Figure 15shows the HRTEM cross-section image of a GaAs NWgrown by the PGM. This image and the selected areadiffraction pattern revealed the existence of sphalerite (c)and wurtzite (w) type structure in the GaAs NW.

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Figure 14 (online colour at: www.pss-b.com) SEM ((-110) view):Effect of the V/III ratio on the twin formation in GaAs NW. DuringNWgrowth theV/III ratiowas abruptly increased from10 to40.Asaresult, the twin density in the upperNWhalf (grown atV/III¼ 40) issignificantly reduced.

Figure 15 HRTEM image ((-110) view) of aGaAsNWconsistingof regions with sphalerite-type (c) and wurtzite-type (w) structures.The gold film thickness was 20 nm. Twins (T) exist within thesphalerite-type regions. The inset shows the extended wurtziteregionand the smoothsphalerite/wurtzite transitionofanNWgrownwith a higher galliumcontent in the droplet after initial alloyingwithTMG. The thickness of the gold filmwas 3 nm. Both NWdiametersare equal.

The foot region was from pure sphalerite type. Allsegments of sphalerite type structure contained stackingfaults and twins. After the stabilization of the Au–Ga droplet(20 s TMG supply) and the following arsine supply the NWgrowth was abruptly switched to wurtzite type. Theoccurrence of the wurtzite type structure is usually discussedas a result of the nucleation process. The two-dimensionalnucleation occurs preferentially at the VLS phase line and ata high As supersaturation in the Ga-rich Au–Ga droplet[1, 36]. With a higher gallium content in the startingcomposition of the droplet an extended w-region wasobserved (Fig. 15, inset, thickness of the gold layer was3 nm). After the last PGM switching sequence the NW neckregion was formed under arsine flow. The observed NWdiameter decrease probably results from the partial galliumconsumption from the Au–Ga droplet (see Eq. (2.3)). To thebest of our knowledge this is the first report on the formationof GaAs NW with wurtzite type structures grown byMOVPE. With a systematic investigation of the growthparameters on the crystal structure we are looking forward toachieve controllable sphalerite/wurtzite superlattice struc-tures via the VLS mechanism in PGM in future.

2.6 Nanowire heterostructures A main reason forthe investigations of the semiconductorNWstructure and theNW growth process is the fabrication of novel basicstructures for modern electronic and optoelectronic buildingblocks. For this purpose it is necessary to combine materialswith diverse electronic properties. NW offer two types ofheterostructures: (1) Axial NW heterostructures via the VLSmechanism emerge from switching of different precursor

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materials at optimized growth conditions. The particularNWgeometry seems to be advantageous for axial heterostruc-tures. For sufficient thin NW the combination of materialswith high lattice mismatch is expected to be achievablewithout dislocation formation [37]. (2) Lateral NW hetero-structures (‘core/shell’ structures) are established by con-ventionalMOVPE layer growth on the NW side facets. Suchcoating layers are useful to passivate the NW surface and forthe confinement of charge carriers inside the NW.

While the growth of axial heterojunctions with group-Vmaterial change has been successfully demonstrated [38],the growth of the desirable III-As heterostructures isproblematic. Frequently, axial heterostructure growth isaccompanied by characteristic phenomena like kinking,increased twin density, the formation of diffuse heterojunc-tions with alloy crystal gradients instead of abrupt interfaces[39, 40] and growth along a side facet of the already grownNW [41]. However, sharp axial InAs/GaAs heterojunctionshave been reported [42], which motivated a more detailedinvestigation of the material system [14].

We investigated the growth of (InGa)As sectionsembedded into GaAs NW. The NW growth was started withpure GaAs. Then TMI was spontaneously supplied inaddition. The partial pressure ratio was TMG/TMI¼ 2.After a certain time the TMI supply was stopped to grow the



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Figure 16 TEMbright-field image ((-110) view) of an axialGaAs/(InGa)As/GaAs double-heterostructure NW. The indium fractionalong the NW is illustrated by the EDX line scan. The (InGa)Assegmentmeasuresabout100 nm.Aconsiderableamountof indiumisfound in thealloyparticle, too.Hence, indiumis apparent in thealloydroplet during growth of the second GaAs segment.

Figure 17 (a) SEM image ((-110) view) of pure GaAs NW. (b)TEMbright-field image ((-110) view) ofGaAs/(AlGa)As core/shellNW. As a result of the (AlGa)As coating layer the NW diameterincreasesand themorphologyof theNWcross-section ((111)viewinthe insets) changes. (c) TEM bright-field image ((-110) view) of aGaAs/GaN core/shell NW. The polycrystalline GaN layer contin-uously surrounds the GaAs NW.

second GaAs NW section. The double-heterostructure NWhad a straight morphology without kinking. EDX mappingexperiments verified the existence of the (InGa)As section ineach NW [14]. A more quantitative examination was madebyEDX line scans at single heterostructureNW(Fig. 16). Animprovement of the sharpness of the heterointerfaces and ahigh incorporation of indium was achieved by the use ofgrowth interruptions prior to and after the (InGa)As growthstep [14].

Hiruma et al. [43] showed that the photoluminescenceefficiency of GaAs NW can be increased due to surfacepassivation. Thus, the optically inactive region at the GaAssurface (about 10 nm) can be avoided. By using asemiconducting coating with a higher bandgap an additionalcharge carrier confinement favours the photoluminescenceyield [44]. We investigated the effect of (AlGa)As and GaNas coating material. In both cases the coating layer wasdeposited directly after NW growth at the specific MOVPElayer growth conditions. Tapered GaAs NW were grownwith TMG(2.3 Pa) and tBA (3.6 Pa) at 550 8C.The (AlGa)Asgrowth was realized with trimethylaluminium (TMA,0.15 Pa), TMG (0.27 Pa) and arsine (35.6 Pa) for 10min at700 8C. The polycrystalline GaN layer was deposited withTMG (0.17 Pa) and 1,1-dimethylhydrazine (DMHy, 291 Pa)for 10min at 610 8C.

The core/shell samples were investigated by SEM andTEM experiments. GaAs NW without coating layer showed{-211} side facets. The NW diameters were about 10–15%less than the corresponding, crystallized gold particles(Fig. 17a). In contrast, GaAs NW with (AlGa)As coatinglayer showed a significantly higher diameter with respect tothe gold particle size (Fig. 17b). Furthermore, the NW cross-section changed to a hexagonal shape with {-110} sidefacets. Since the material contrast between (AlGa)As andGaAs is poor in TEM, the heterojunction could not bevisualized. However, the different side-facet geometries andthe increased NW diameters indicate on the successfulformation of the (AlGa)As coating layer.

GaAs/GaN NW are shown in Fig. 17c. The GaN coatinghad a thickness of 10–20 nm. The GaN deposition tempera-ture was considerably low. As a result, the layer was verydefect-rich with sphalerite/wurtzite transitions. However,the GaN layer was homogeneous and continuously


distributed along the NW. Hence, the layer was suitable aspassive coating material.

We used an Arþ-LASER (514.5 nm) with an outputpower of 60mW for the photoluminescence experiments.The LASER excitation spot size was about 100mm. Thesamples were mounted in a cryostat for low temperaturemeasurements (T¼ 2K). For detection a 1m doublemonochromator with a Peltier-cooled GaAs-photomultiplierwas used. The investigated GaAs NW were either grown on(111) Si or on (-1-1-1)As GaAs, whereby in the latter case theNWwere scratched on a carbon pad to get independent of thesubstrate luminescence.

The observed spectra (Fig. 18a and b) are dominated bybroad donor–acceptor pair transitions (DAP) from a neutraldonor (D0) to the neutral C-acceptor and in the case of Sisubstrate also to the neutral Si-acceptor. The C-doping isexpected to result from the uncomplete decomposition ofTMG at the low growth temperature and V/III ratio. Theincorporation of Si originates from the initial alloying of thegold droplet with the substrate material.

We investigated the photoluminescence of columnar aswell as tapered GaAs NW and distinguished the effect of

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Figure 18 Low-temperature photoluminescence investigation at(a) GaAs/(AlGa)As core/shell NW grown on (111) Si substrate and(b)GaAs/GaNcore/shellNWgrownon(111)AsGaAs.Dependingonthe shell material the GaAs NW is tensile ((AlGa)As-shell) orcompressively (GaN-shell) strained. As a result, the PL band showsa red shift or a blue shift, respectively.

both coatings. Experiments concerning the NW shapewithout any coating showed that tapered GaAs NW have amuch higher luminescence yield than columnar NW. Weconsider three reasons for explanation: (1) The tapered NWwere grown at higher temperatures, so that less lattice defectsmight be incorporated. (2) The precursor decomposition isnot completed in the growth window of columnar NW. Inparticular, TMG and tBA decompose at about 300–500 8C(see Ref. [45], p. 267). Hence, C-doping possibly occursduring columnar NW growth. (3) Between GaAs and air is afairly high diffraction index difference. Hence, columnarGaAs NW act as a waveguide for the generated light in thedirection of the NW axis. A great amount of the light isabsorbed by the substrate in that way. In taperedNW the lightcan leave the NW much easier.

Point (3) is supported by the PL experiments at core/shellNW. The luminescence yield of columnar NW wasconsiderably increased in the core/shell structures. Incontrast, the high luminescence yield of tapered NW is

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reduced by the coating layers (Fig. 18a and b). We suggest asimultaneous excitation of the NW and the coating materialwhich lowers the GaAs band luminescence intensity.However, besides the intensity also the energetic positionof the emission band changes. In particular, the GaAs PLband is red shifted if the NW are coated with (AlGa)As(Fig. 18a). This effect was previously observed by Seifertet al. [17]. The reason ismentioned to be the latticemismatchbetween NW and the cover material. The pseudomorphicdeposition of the wider lattice spaced (AlGa)As coatingresults in a tensile strain in the NW’s length direction. Indeedwe verified the opposite behaviour for the GaN coating,which has lower lattice spacing than GaAs (Fig. 18b).

3 The SAG mechanism The SAG mechanism pro-vides amethod for realizingwell-orderedNWarrayswithoutthe use of a catalyst material [46–48]. This method uses thehigh difference of growth rates of particular crystal planesat specific growth conditions for the one-dimensionalgrowth. We investigated the SAG growth of GaAs andInAs NW on various substrates and realized GaAs/InAsheterostructures [6].

3.1 GaAs nanowires The nanowires were grown on(-1-1-1)As GaAs and (111) Ge substrates. We used plasmaenhanced chemical vapour deposition (PECVD) to deposit athin SiNx-layer (thickness �15 nm) on the substrate [49].The SiNx was partially removed by electron beam litho-graphy (EBL) definition of well-ordered circular openings ina electron sensitive resist and wet chemical etching using adiluted ammonium fluoride etching mixture (AF 87.5–12.5VLSI Selectipur Merck). The radius of the openings wasvaried in dependence of the electron dose during the EBLbetween 100 and 300 nm. The triangular latticewith the pitcha between the openings was investigated from 500 to4000 nm. These SiNx templates were used to realizenanostructures of different size and alignment.

SAG was carried out using LP-MOVPE with a totalpressure of 5 kPa, a total gas flow of 7000 sccm and H2

atmosphere. TMG and TMI were used as group-III sourcematerials and arsine was used as group-V source material.We investigated the GaAs NW growth within a temperaturerange of 600–850 8C and a V/III ratio of 130–520. The TMGpartial pressure was kept constant at 0.14 Pa. At hightemperatures (T¼ 750–850 8C) and a highV/III ratio of 260–520 uniform arrays of vertically aligned GaAs NW wererealized on (-1-1-1)As GaAs [6].

For the growth of InAs NW the optimal growthconditions were found to be at a growth temperature of600 8C, a TMI partial pressure of 0.13 Pa and a V/III ratio of550.

Figure 19 shows a GaAs NW array grown on (-1-1-1)AsGaAs substrate. The growth time was 20min and the arraypitch between the circular openings in the 15 nm thick SiNx

layer was 1000 nm. The vertically aligned NW had ahexagonal cross-section formed by six {1-10} side-wallfacets and a top plane parallel to the substrate surface. The



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Figure 19 SEM image (titled by 458 to the surface normal) of ahexagonal arrangement ofGaAsNWgrownby theSAGmechanismwith an array pitch of 1mm (inset: top view).

Figure 21 HRTEM image ((-110) view) of a GaAs NW grown on(-1-1-1)As GaAs by the SAGmechanism. The GaAs NW consist ofsphalerite type structure with rotational twins.

NW had an average radius of r¼ 130 nm and a height ofh¼ 1600 nm.By changing the radius of the circular openingsin the SiNx layer the resulting NW radius can be controlled.

Figure 20 summarizes the investigations of the radius-dependent growth of SAG NW for T¼ 800 8C andV/III¼ 260 with a TMG flow of 3.75ml/min and a growthtime of 20min. The plotted reference curve represents therelationship between the growth rate Rg¼ h/t¼Cg/rþC0,where Cg is the rate coefficient dependent on r and theconstant C0 is related to the layer growth of GaAs on the(-1-1-1)As surface (for more details see Refs. [6, 50]).

These NW arrays were characterized by high-resolutionXRD in collaboration with the University of Siegen andEuropean Synchrotron Radiation Facility [51, 52]. With anX-ray beam focussed down to a spot size of 220� 600 nm2

Figure 20 NWheight h in dependence of theNWradius r for SAG(array pitch a) grownGaAs NW at V/III¼ 260. The SAG data werefitted using standard growth models (see Refs. [6, 45]).


even high-resolution XRD patterns of individual NW werepossible [53]. TEM investigations at GaAs NW showed thatthe real structure contains rotational twins in the sphaleritetype lattice [6] (see Fig. 21).

3.2 GaAs nanotree structures NW heterostruc-tures in axial and radial directions were fabricated from avariety of materials [11, 38, 54]. The production of complexthree-dimensional nanotree-like structures, based on sequen-tial seeding of multiple generation of epitaxial NW usingmetal seed particles has been reported recently [55, 56].

We investigated the combination of SAG and VLSmechanism for GaAs nanotree growth using a two-stepepitaxial growth. We produced well-ordered arrays ofcomplex, three-dimensional nanostructures build fromGaAs NW taking the advantages of both methods – theaccurate positioning possibility of the SAG mode and thehigh growth rate of the VLS mechanism.

The NW array in Fig. 19 forms the basis for the secondepitaxial VLS-grownNW. Therefore the SAG samples wereevaporated with a thin Au layer. The evaporation directionwas tilted by 458 from the NW length axis towards a {1-10}sidewall facet. To reach a homogeneous deposition on allside facets two further evaporation steps were performedwith a 1208 rotation around the NW axis in each step (seeFig. 22). The layer thickness was about 1 nm.

Table 1 summarizes the MOVPE parameters for bothgrowthmethods. In the SAGmechanismV/III ratios of about260 are necessary. TheVLS growthwas carried out at 480 8Cand V/III¼ 20. The growth time was 2min. Note the low V/III ratio of 20. The VLS-grown GaAs NW nucleate from the{1-10} sidewalls of the SAG NW with h-1-1-1iAs growthdirections (Fig. 23). The obtained NW exhibited diametersfrom 20 to 50 nm and lengths up to 1mm.

Figure 23 shows the GaAs nanotree structures. EachSAGNW is surrounded by a number of VLSNW (branches)which grow in the corresponding h-1-1-1iAs directions. In thetop-view image in Fig. 21 (right) the six-branch directions

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Table 1 MOVPE growth parameters for GaAs NW via the SAGand the VLS mechanism.

T (8C) TMGa(ml/min)

AsH3

(ml/min)V/IIIratio

t(min:s)

selective-area growth 800 3.75 50 260 20:00VLS growth 480 10 10 20 2:00

Temperature T, TMGa and arsine flow, V/III ratio, and growth time t.

Figure 22 Schematic illustration of the three-step thermal evapo-ration process.

are visible. In dependence of the array pitch and theadjustment of the circular openings (defined by EBL) adocking of neighbouring SAGNWvia theVLSNWseems tobecome possible in future.

4 Conclusion TheVLS and the SAGmechanismwereused to fabricate GaAs NW in MOVPE. Despite thefundamental concept of each mechanism is different, withboth of which GaAs NW are obtainable with controllablediameter and growth rate. Furthermore, positioning on thesubstrate surface and fabrication of axial and lateralheterostructures are possible. Themost promising advantageof the SAG mechanism over the VLS method is the absence

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of an additionalmaterial component such as gold. It is not yetclear, how much gold is incorporated into the GaAs NWduring VLS growth. However, the VLS mechanism allowsthe GaAs NW growth at a 250 8C reduced growthtemperature with up to 10 times increased growth rate. TheNWdiameter for the SAGmethod is presently limited by thenecessary substrate mask (>50 nm). With the VLS methodalso GaAs NW with diameters below 10 nm have beenreached.

Basic aspects of the VLS mechanism have beenfocussed. XRD measurements indicate a reaction of theevaporated gold layer with the underlying GaAs substrate.Furthermore, the alloy formation of a gold layer with TMGwas verified on inert Al2O3 substrate.We estimated a dropletcomposition of about Au0.74Ga0.26 during VLS growth(V/III¼ 10), which is in liquid state at growth temperature.Thermodynamic and kinetic aspects of the VLS NW growthwere considered. The Gibbs–Thomson effect was found todetermine the GaAs NW growth rate at low V/III ratios(<20).With increasingV/III ratio the gallium fractionwithinthe Au–Ga–As droplet decreases until the eutectic compo-sition is reached. At this point (V/III¼ 20) the conventionalview of the VLS mechanism is not more applicable.

Based on the knowledge of the specific GaAs NWmorphology a growth sequence with TMG interruptions hasbeen developed to position in situmarkers during GaAs NWgrowth. With a special two-step procedure we showed apromising route to increase the structural quality of VLSgrown GaAs NW by decreasing the twin density.Furthermore, photoluminescence experiments showed agood optical quality of GaAs NW as well as GaAs/(AlGa)As and GaAs/GaN core/shell NW. The misfit strainin the core/shell structure allows the adjustment of the GaAsbandgap.

Besides the commonly used continuous VLS growthmode we suggested a pulsed VLS growth mode withalternating TMG and arsine supply for the fabrication ofGaAs NW. Since CGM grown GaAs NW always exhibitsphalerite type crystal structure, PGM grown GaAs NWshowed regular, alternating sphalerite/wurtzite sections.

TheMOVPEparameters for SAGgrownGaAsNWweredetermined. The NW growth rate showed the expecteddependencies on the radius [9, 43]. The excellent two-

Figure 23 SEM images of a GaAs nanotreearray. Left: tilted by 458 to the surface normal;right: top-view.



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dimensional arrangement of SAG grown GaAs NW wasapplied as a starting point for a secondVLS-based GaAsNWgrowth step in preparation for the fabrication of tailored NWnetworks.

Acknowledgements The authors would like to thankK. Mergenthaler, M. Shirnow and A. Vogel for several growthexperiments and U. Pietsch, A. Davydok, A. Biermann andJ. Grenzer for special X-ray investigations. We thank G. Benndorffor support in the PL experiments. The work was supported byDeutsche Forschungsgemeinschaft (research group 522, DFGcontract GO 629 7/3).

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