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Micellar block copolymer templated galvanic displacement for epitaxialnanowire device integration†

Gregory S. Doerk,‡a Charles Dhong,xa Christine Politi,xb Ian Laboriante,a Carlo Carraroa

and Roya Maboudian*a

Received 16th February 2011, Accepted 13th April 2011

DOI: 10.1039/c1jm10693g

New strategies for catalyst nanoparticle placement in arbitrary patterns and non-planar geometries will

likely accelerate the large scale device integration of epitaxial semiconductor nanowires (NWs) grown

through the vapor-liquid-solid (VLS) process. Herein we report a technique for rational metal catalyst

nanoparticle deposition based on galvanic displacement onto semiconductor surfaces through block

copolymer micelle templates. Nanoparticle volumes and areal densities are controlled by the time in the

plating solution and by mixing homopolymer with the micelle suspension, respectively. Above

a minimal nanoparticle diameter, the mean diameters of epitaxial VLS-grown Si NWs scale directly

with mean sizes of the template deposited nanoparticle seeds from which they are grown. The substrate

selectivity and conformality of galvanic displacement makes possible two-level micro/nano-patterning

in a variety of geometries by applying micellar templates over photolithographically patterned masks;

the growth of single sub-50 nm diameter Si NWs in 600–700 nm diameter wells demonstrates a feature

size reduction greater than one order of magnitude. Two-point electrical measurements across single

NWs or a few NWs epitaxially bridging silicon-on-insulator electrodes after ex situ doping demonstrate

the viability of this approach for epitaxial NW device integration.

1 Introduction

The high crystal quality, extensive range of synthetically avail-

able materials,1 and capacity for controlled formation of axial,2

core–shell,3 and branched4,5 semiconductor NW heterostructures

made possible through VLS synthesis portends great promise for

its use in future technologies. Indeed single NW devices have

been demonstrated in numerous application areas including

nanoelectronics,6 photonics,7 and sensing.8 However, the full

potential of NW-based devices cannot be realized without the

development of large scale, high yield device integration strate-

gies with accurate spatial registry. For NWs synthesized by the

VLS method in general, devices are fabricated either on NWs

that have been transferred from the growth substrate to

aDepartment of Chemical and Biomolecular Engineering, University ofCalifornia, Berkeley, California, 94720, USA. E-mail: maboudia@berkeley.edu; Fax: +1 510 642 4778bDepartment of Chemical and Biological Engineering, University ofCol-orado, Boulder, Colorado, 80309, USA

† Electronic supplementary information (ESI) available: Auger electronspectra of substrates at different points in the templating process,a plot showing mean Au NP sizes versus galvanic displacementimmersion time for low and high areal densities, and a XPS spectrumof a plain Si substrate. See DOI: 10.1039/c1jm10693g

‡ Present address: IBMAlmaden Research Center, 650 Harry Road, SanJose, CA 95120-6099, USA.

x These authors contributed equally to this work.

This journal is ª The Royal Society of Chemistry 2011

a separate device substrate, or from NWs that have been grown

epitaxially in the location and geometry desired for the device, as

directed by the relation between the orientation of the substrate

surface and the NW preferred growth orientation. While post-

growth assembly techniques have advanced impressively towards

achieving large scale NW device integration,9,10 this integration

problem remains acute for devices composed of epitaxially

grown semiconductor NWs.

This severely hinders the development of technology utilizing

the unique advantages of epitaxial VLS-grown NWs. For

instance, highly aligned vertical NW arrays can be fabricated by

growing the NWs epitaxially on flat substrates if the crystallo-

graphic orientation orthogonal to the surface coincides with the

preferred NW crystal growth orientation. Such arrays may serve

as the basis for cellular biomolecular delivery platforms,11

photovoltaics,12 and three-dimensional integrated circuits.13,14 In

ideal cases, the position and size of each vertical nanowire is

determined by the same characteristics in the predeposited

metallic nanoparticle (NP) that serves as the growth catalyst.

Electron beam lithography may be used to control catalyst

particle placement for epitaxial NW growth,15,16 but its high cost

and serial nature make it unsuitable for large scale processing.

Colloidal metal NPs offer a possible lower cost solution, and

a number of techniques for positioning Au colloids (and hence

the resulting NWs) have emerged,17–20 even to the extent of single

NP resolution.19 However, the range of particle densities per unit

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area may be limited depending on the technique, and accurate

placement of a small number of catalyst NPs (1–3) still requires

high resolution nanolithography (< �200 nm) that will likely

drive up costs. Moreover, these patterning strategies may not be

easily applied to substrates with varied topography (i.e.

nonplanar). This is critical in case of fabricating epitaxially

bridging VLS-grown NWs in microtrenches, where features such

as rigid mechanical clamping21 and ultra-low electrical contact

resistances22 promote their use in potential applications such as

rapid microfluidic biosensors,23 directly grown logic gates and

photovoltaic devices,24 horizontal wrap-gate transistors,25 and

nanomechanical resonant NW mass sensors.26 Nevertheless, to

date the placement of Au colloids on trench sidewalls is still

essentially random.

Specialized metal deposition and templating chemistry may

help bridge the gap between the resolution achieved at low cost

with photolithography and the high resolution required to ach-

ieve devices consisting of single or only a few NWs, while

potentially reducing process complexity. One simple yet attrac-

tive way to deposit metals with nanoscale control is through

galvanic displacement, in which metal ions dissolved in (usually

aqueous) solutions are spontaneously reduced to elemental form

by electrons provided through oxidation of the underlying

substrate. Typically metals with redox potentials above hydrogen

may be deposited, while substrates must be able to oxidize

readily; these criteria still leave available a wide variety of

important metals and semiconductor substrates.27 The require-

ment for substrate oxidation makes the deposition inherently

selective, permitting the use of common masking materials such

as SiO2, SiNx, or some polymers. Furthermore, deposition is

conformal and thickness may be controlled by immersion time

and metal ion concentration. No seed layers are necessary,

though in most cases an additional solute like hydrofluoric acid

must be added to remove the oxidized substrate as it develops.

Gao et al. employed galvanic displacement from water-in-oil

microemulsions with micelles containing an aqueous solution of

a Au salt and hydrofluoric acid to deposit Au NPs selectively

onto Si microtrench sidewalls and flat Si substrates with

controlled nanoscale diameters to seed the subsequent growth of

epitaxial Si nanowires.28 Aizawa and Buriak demonstrated the

use of diblock copolymer micelles to template the deposition of

various size-controlled noble metal NPs on Si, Ge, InP and

GaAs.29,30 Unfortunately, both methods give rise to high NP

areal densities that may only be viable for the growth of dense

NW arrays. However, an intriguing aspect to surface patterning

with block copolymer micellar templates is the fact that the

micelles are often trapped in a non-equilibrium state; they are

then amenable to density modulation by the addition of inert

micelles,31 empty micelles, or homopolymers.32

In this report we demonstrate a general strategy for the

epitaxial integration of semiconductor NWs based on the

controlled deposition of catalyst NPs via galvanic displacement

through micellar diblock copolymer templates. The chemically

rate-controlled Au deposition affords direct control over catalyst

NP size, while the addition of polystyrene homopolymer enables

rational reduction of NP areal densities. Thus, NPs with a range

of sizes may be deposited using a single micellar suspension. The

mean diameters of epitaxial Si nanowires follow the mean sizes of

the as-deposited NPs from which they are grown. Using dip

8808 | J. Mater. Chem., 2011, 21, 8807–8815

coating, we have applied this templating route to both planar and

microtrench substrates patterned with Si ‘‘windows’’ in SiO2

where Au NPs can be deposited selectively. Single Si NWs with

diameters of approximately �50 nm may be grown in windows >

600 nm in diameter, indicative of an order of magnitude reduc-

tion in feature size. Electrical measurements across two Si

electrodes connected by single or few-numbered epitaxially

bridging Si NWs exhibit ohmic responses and specific contact

resistances that compare favorably with lowest values for other

contact methods reported in the literature. The catalyst place-

ment strategy presented here may be a promising step towards

achieving high spatial registry at low cost for the integration of

epitaxial NWs into future devices.

2 Experimental

2.1 Materials

Si(111) dice 10 mm� 11 mm (p-type, r¼ 5–10 mU cm) were used

for basic spin- and dip-coating experiments, as well as for

unpatterned nanowire (NW) growth studies. Two types of

patterned substrates were used. Si(111) substrates covered by an

SiO2 film approximately 400 nm thick and patterned with wells of

different shapes and sizes ranging from < �0.5 mm2 to > 20 mm2

that exposed the underlying Si were fabricated for vertical growth

studies. For epitaxial bridging NW growth, silicon-on-insulator

substrates with an approximately 500 nm thick buried oxide

layer and an approximately 1 mm thick highly boron doped

(> 1018 cm�3) device layer were used. Details about the fabrication

and patterning of these samples has been provided in a previous

publication.33 In all cases, before block copolymer templating,

samples were cleaned by sequential sonication in acetone,

isopropyl alcohol, and deionized (DI) water for 10 min. each,

followedby treatmentwith aUVozone cleaner for 10min. (Jelight

UVO-Cleaner Model No. 42). Two polymers were used in

templating. The block copolymer was polystyrene-block-

poly(2-vinylpyridine) (PS-P2VP, fP2VP ¼ 0.28; total

Mn¼ 46 kg mol�1, Mw/Mn¼ 1.18), and the homopolymer was PS

(Mn ¼ 6 kg mol�1, Mw/Mn ¼ 1.5). Both were synthesized using

anionic polymerization, following standard methods.34

2.2 Block copolymer templating and gold deposition

Micelles were formed by dissolving the PS-P2VP in toluene at

a mass concentration of 5 mg mL�1, and heating the solution

to > 70 �C for more than 10 min., during which time micelles

formed. The solution was then cooled to room temperature.

Homopolymer solutions were made by dissolving PS in toluene

at room temperature with a mass concentration of 5 mg mL�1,

and block copolymer/homopolymer mixtures were prepared with

the desired mass ratio by mixing the PS-P2VP micelle suspension

with the PS solution in the same volume ratio. All mixtures were

stored under refrigeration, but allowed to warm to room

temperature before use, and could be reused repeatedly over the

course of months.

Block copolymer templates were applied to substrates by spin-

or dip-coating. Before use, the liquid mixture was stirred using

a magnetic stir-bar for at least 30 min. For spin coating, 15 mL of

a polymer mixture was placed on a sample, which was then

spin-coated in air at 4000 rpm. Dip-coating was accomplished by

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pulling the sample out of the polymer mixture at a rate of

80–90 mm/min.

Metal deposition was accomplished via galvanic displacement.

Plating solutions were made by adding 100 mL of a 10 mM

aqueous solution of potassium tetrachloroaurate (KAuCl4)

along with 10 mL of 48–50% (w/w) hydrofluoric acid (HF) to 10

mL of DI water. Particles of different sizes were deposited

through the micelle template by immersing samples in the plating

solution for 10 min. or less, with negligible effect on the particle

density.† After metal deposition, samples were sequentially

immersed in at least 2 beakers of clean DI water followed by

drying under a nitrogen stream.

The polymer template was removed by immersing samples into

neat toluene heated to temperatures above greater than 70 �C for

at least 10 min. Unpatterned samples were then sonicated in the

same solution for 10 min. In patterned samples sonication

sometimes resulted in a significant loss in the number of Au

nanoparticles (NPs) and so it was not always done. Auger elec-

tron spectra confirm the removal of most of the polymer film

with sonication in heated toluene.†

2.3 Synthesis and boron doping of silicon nanowires

Si NWs were synthesized via vapor-liquid-solid (VLS) growth

from the template-deposited Au NPs. Samples were placed in

a quartz tube, hot-wall, atmospheric pressure chemical vapor

deposition reactor, with NW growth occurring at 835 �C for 5–

10 min. SiCl4 was transported into the reactor by bubbling

10 sccm of carrier gas (10%H2 in Ar) through liquid SiCl4 held at

0 �C, while 100 sccm of the carrier gas flowed directly to the

reactor. For the growth of Si NWs using Au catalyst NPs with

diameters above �20–30 nm, the thermodynamically preferred

growth direction is the <111> direction; thus a large proportion

of Si NWs grown from Au epitaxially from Si(111) surfaces will

stand perpendicular to the surface.5

Ex situ boron doping was performed using BBr3 in the same

reactor used in Si NW growth but with a different quartz tube.

The process has been described in a previous publication,35 but

here we briefly summarize it with details pertinent to this work.

Prior to doping, the sample was immersed in an iodine/potassium

iodide Au etchant solution (4 : 1 : 40 KI:I2:H2O) for more than

30 min. to remove the Au catalyst material, followed by rinsing

with isopropyl alcohol and drying under a nitrogen stream. The

sample was then cleaned by UV ozone treatment for 30 min. and

inserted into the reactor for doping, which was performed in two

steps. In the first step boron was introduced to the Si NWs at

700 �C by bubbling 6 sccm of the carrier gas (again 10%H2 in Ar)

through liquid BBr3 held at 0 �C for 1 min., while 270 sccm of the

carrier gas flowed directly to the reactor. At this stage, BBr3 is

believed to be reduced by H2 to deposit boron on the Si NW

surfaces. The BBr3 flow was then cut off and the carrier gas was

switched to Ar gas only. The sample was annealed under this Ar

carrier gas at the same temperature for 60 min. to drive-in and

activate the boron as a dopant.

2.4 Characterization

Scanning electron microscope (SEM) images were obtained using

either an Agilent 8500 field-emission SEM or a Leo 1550

This journal is ª The Royal Society of Chemistry 2011

Schottky field-emission SEM. Atomic force microscopy (AFM)

measurements were performed in air with a Digital Instruments

Nanoscope IIIa system in tapping mode. Mean NP volumes were

determined by flooding analysis using WSxM Scanning Probe

Microscopy Software,36 and the error bars presented in the

report represent the standard deviation in the mean NP volume

over several images. For aesthetic presentation, Gwyddion

scanning probe microscopy software was used to create topo-

graphic images. Au NP counts and areal densities were deter-

mined from SEM and AFM images. Error bars in such cases

represent standard deviations in the mean values obtained from

at least three images. All images used for measurements were

taken from the same samples, though repeated experimental

conditions resulted in highly consistent results. Surface cleanli-

ness and composition were characterized by Auger electron

spectroscopy (AES; PHI-Perkin-Elmer model 10-155) and X-ray

photoelectron spectroscopy (XPS; Omicrometer analyzer,

EA 125).

Two-point electrical measurements were performed using

a probe station with a HP4145B Semiconductor Parameter

Analyzer. Since a number of electrodes are interconnected by

epitaxially bridging Si NWs, and these in turn occasionally

bridge to a third ‘‘gate’’ electrode nearby, a small degree of

leakage current was observed in some measurements. We cor-

rected for this by measuring not just the resistance across the

NWs under test (‘‘source to drain’’, RSD) but also through the

other connected conduction paths (‘‘source to gate’’, RSG, and

‘‘gate to drain’’, RGD). The actual NW resistance (RNWs) is

obtained by the following equation:

RNWs ¼ RSDðRSG þ RGDÞRSG þ RGD � RSD

(1)

The measurements of RSG and RGD most likely underestimate

their true values due to the parallel conduction path through the

wire; therefore this measurement may overestimate the NW

resistance. The maximum negative error is included in Fig. 6(d);

however, a large potential barrier to electrical conduction by the

leakage pathsmakes this negative error small inmost cases for the

voltage range over which resistance is measured (�1 V to 1 V).

3 Results and discussion

A schematic for the micellar templating strategy is presented in

Fig. 1. The block copolymer used was polystyrene-block-

poly(2-vinylpyridine) (PS-P2VP; Mn ¼ 46 kg mol�1, Mw/Mn ¼1.18). In bulk form, the P2VP volume fraction (fP2VP ¼ 0.28)

implies that the polymer should microphase separate to cylin-

drical P2VP domains in a PS matrix;37 when dissolved in toluene

(5 mg mL�1) and heated to 70 �C the preferential interactions

between toluene and polystyrene result in the formation of

spherical micelles with P2VP cores and PS coronas (Fig. 1(a))

that are preserved when the solution is cooled to room temper-

ature.29 This micellar suspension is then mixed with a solution of

PS homopolymer (Mn ¼ 6 kg mol�1, Mw/Mn ¼ 1.5) in toluene at

the samemass concentration to obtain a newmixture with a mass

ratio of PS homopolymer to block copolymer that is equivalent

to the volume ratio of the two mixtures (Fig. 1(b)). A micellar

template film is applied to a silicon substrate by either spin or dip

coating. After the sample is visibly dry, templated samples are

J. Mater. Chem., 2011, 21, 8807–8815 | 8809

Fig. 1 Block copolymer templating scheme. (a) Micellization of

PS-P2VP occurs in toluene above room temperature, with P2VP cores

(red) and PS coronas (blue). (b) After being cooled to room temperature,

the micelle suspension is mixed with a solution of PS in toluene having the

same mass concentration in the desired mass proportion. (c) The micelle

template is applied to a semiconductor surface by spin- or dip-coating,

where the PS serves as a spacer between micelles. The sample is then

exposed to the aqueous galvanic displacement bath, where the P2VP

micelle cores allow selective access to the surface for metal deposition. (d)

The polymer template is removed, leaving an array of metal NPs.

Fig. 2 Results of the templating process with no added PS. SEM images

of (a) the initial micellar template, (b) after Au galvanic displacement for

10 min, and (c) after polymer removal. Scale bar ¼ 1 mm for all images.

(d) Au region of the XPS spectrum from the sample shown in (c).

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immersed in a Au plating solution consisting of potassium

tetrachloroaurate (KAuCl4) and excess hydrofluoric acid (HF) in

deionized (DI) water for variable times, during which Au is

deposited via galvanic displacement only through the P2VP cores

that give aqueous solutions access to the underlying Si surface, as

depicted in Fig. 1(c). The PS homopolymer acts as a spacer on

the surface, decreasing the surface areal density of the micelles.

Finally, the polymer is removed through immersion in hot

toluene (> 70 �C), followed by sonication (which may be skipped

if particle adhesion is poor) leaving an array of Au NPs with

controlled size and areal density (Fig. 1(d)). The effectiveness of

the polymer removal is verified using Auger electron

spectroscopy.†

The results of this templating approach (by spin coating) with

no PS homopolymer added are displayed in Fig. 2. Fig. 2(a)

shows a scanning electron microscope (SEM) image of the initial

polymer template, while Fig. 2(b) and (c) show SEM images after

10 min of galvanic displacement and subsequent polymer

removal, respectively. There are fewer particles (� 40–50 mm�2)

after galvanic displacement than there are apparent micelles

(> 60 mm�2), though the final number of particles after polymer

removal is approximately the same after galvanic displacement.

8810 | J. Mater. Chem., 2011, 21, 8807–8815

The imperfect formation of particles from the micellar template

may be attributed to the presence of a PS homopolymer layer

directly above (or underneath) particular micelles that prevent

direct access of the aqueous solution from the substrate. The Au

region of an X-ray photoelectron spectrum (Fig. 2(d)) taken

from the sample shown in Fig. 2(c) indicates that the deposited

material is indeed elemental Au. The small hump at �90 eV is

from the Si substrate itself.†

A primary reason for using block-copolymer micellar

templating is the capability to tune the particle areal density over

a wide range. We do this by adding PS homopolymer, the results

of which are summarized in Fig. 3. SEM images in Fig. 3(a–c)

depict as-deposited Au NPs after polymer removal for three

different mass fractions of PS-P2VP block-copolymer with PS

homopolymer as the other binary component (the total mass

concentration of the mixture was held constant). As the block-

copolymer mass fraction is increased from�0.03 (a) to�0.33 (b),

and finally up to 1.0 (c; undiluted block copolymer), the Au NP

areal density increases proportionally. Since the areal density (ra)

is proportional to D�2, where D is the domain size, reduction in

areal density is also possible by using higher molecular weight

block copolymers. In the case of PS-P2VP which may be

considered amphiphilic, or at least strongly segregated (at the

molecular weight used in this report or above), D may be expected

to scale with the degree of polymerization N as D f N2/3�1.38

As a result, polymers used to achieve this reduction in ra must

possess molecular weights at least �5–10 times larger, and the

weak scaling of micelle center-to-center distance with degree of

polymerization for metallic NPs deposited through PS-P2VP

micellar templates in the previous work by Aizawa and Buriak

corresponds more closely to the strongly segregated limit (i.e.

requiring molecular weights at least 10 times larger).30 On the

other hand, as demonstrated here extensive areal density tuning

may be achieved in a rational way through the mixture of only

a single block copolymer and a single homopolymer. Though the

Au NP size is controlled separately by galvanic displacement

This journal is ª The Royal Society of Chemistry 2011

Fig. 3 NP areal density control by homopolymer addition. (a–c) SEM

images of samples produced with block copolymer mass fractions of (a)

0.03 (1 min galvanic displacement), (b) 0.33 (5 min galvanic displace-

ment), and (c) 1.0 (10 min galvanic displacement). All images were taken

after polymer removal, and the scale bar ¼ 1 mm for each image.

Immersion times for each image were chosen to provide optimal NP

clarity and contrast, due to the decreased average Au NP deposition rate

with increased micelle areal density.† (d) NP areal density as a function of

the block copolymer mass fraction for dip-coated samples. The dashed

line depicts expected areal densities assuming it is proportional to the

volume fraction of block copolymer in the mixture. The error bars

represent standard deviations in mean values obtained over several

images.

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plating time (discussed later in this report), the surface micelle

areal density does have an notable impact. Specifically, NP

growth rate is reduced as ra is increased, a fact that we attribute

to increased competition for reactant. For this reason, lower

immersion times were necessary to achieve optimal NP clarity

and contrast in the SEM images depicted in Fig. 3(a–c). A graph

comparing Au NP growth at > 40 mm�2 and < 2 mm�2 is con-

tained in the ESI.†

Fig. 3(d) shows a plot of particle areal density on dip-coated

samples versus the block-copolymer mass fraction in the mixture.

Undiluted block copolymer results in particle areal densities

�42 mm�2, whereas particle areal densities less than 2 mm�2 are

readily obtained for block copolymer mass fractions

< 0.05 (homopolymer to block-copolymer mass ratios greater

than 20 to 1). Similar results were obtained for spin-coating. If

one assumes that the volume fractions of block copolymer and

homopolymer in the thin film micellar template are approxi-

mately identical with their volume fractions in the toluene

mixture, then the areal density of NPs may be considered directly

proportional to the block copolymer volume fraction, fBC, rep-

resented by a dashed blue line in Fig. 3 (The maximum NP areal

density is set to the experimentally measured value and volume

This journal is ª The Royal Society of Chemistry 2011

fraction and mass fraction of block copolymer are nearly iden-

tical, given the equivalent molar volume39 and nearly equivalent

molecular weights for the PS and P2VP monomer units). This

simple model provides a fair guide for expected NP areal densi-

ties, though there is a noticeable discrepancy at a mass fraction of

�0.09. This may be a consequence of the polymer deposition

method since no further steps that are commonly used in block

copolymer template processing like thermal or solvent annealing

have been applied. It may also be due to higher NP deposition in

the micellar template. The connections between processing

conditions and areal density will be investigated more thoroughly

in future work.

Various proposed applications for epitaxially VLS-grown

semiconductor NWs demand that their diameters be precisely

controlled. Since VLS-grown NW diameters are determined

primarily by the size of the metal NP used as a catalyst, this is the

most critical parameter to control. In the method proposed here,

catalyst NP size is regulated by the galvanic deposition rate

(encompassing chemical kinetic and mass transport factors) and

the time of deposition. In order to link NP size with NW diam-

eter, both were measured from the same sample before and after

growth. Furthermore, in order to minimize particle agglomera-

tion due to the high mobility of Au on Si at the elevated

temperatures of our growth reactor (> 800 �C)5,40,41 that wouldstrongly affect the nanowire diameter distribution, we used

samples with low areal densities (< 2 mm�2; template applied by

spin-coating). Fig. 4(a) and (b) show atomic force microscopy

(AFM) topography images of Au NPs on Si(111) deposited

through the method described in this report for immersion times

of 1 min. and 10 min., respectively, and Fig. 4(c) and (d) show

SEM images of the resulting Si NWs grown from the samples in

(a) and (b), demonstrating the clear increase in mean diameter

for NWs grown from the larger particles. The epitaxial nature of

the as-grown Si NWs is evidenced by the high degree of vertical

growth.5

If one assumes that the Au droplet on top of the NW is

approximately hemispherical and the NW radius is nearly

equivalent to the droplet radius, then the NW diameter is

expected to be proportional to the cube root of the Au NP

volume from which it is grown. The exact proportionality is

determined by the eutectic droplet contact angle on Si and the

equilibrium Si mole fraction in the droplet at the growth

temperature, as well as the Si supersaturation during growth and

the ellipticity of the droplet. We ignore the portion of the AuNPs

recessed in the Si substrate though this is offset to some extent by

the convolution of tip size with the NP size measurement.

Therefore, we focus on the proportional scaling between NW

diameter with Au NP size as determined by immersion time.

Fig. 4(e) shows the cube root of the mean Au NP volume as well

as the mean NW diameter versus immersion time. Both increase

monotonically with increasing immersion time, with a much

faster growth rate characterizing the initial NP nucleation stage

(< �1 min.). For times above one minute the mean NW diam-

eters scale directly with the NP size, as expected. A larger

difference between the mean NW diameter and the cube root of

the NP volume is seen for the 30 s. sample. In this case, fewer

NWs grow from smaller Au clusters, a fact that is commonly

attributed to the Gibbs-Thomson effect in which the supersatu-

ration in the droplet is energetically limited by the increased

J. Mater. Chem., 2011, 21, 8807–8815 | 8811

Fig. 4 Control over NP size by immersion time. (a–b) AFM topography

images of Au NPs deposited on Si(111) by galvanic displacement through

�2 mm�2 micellar templates for (a) 1 min. and (b) 10 min. The squares are

both 5 mm � 5 mm and the vertical scale is the same for both images

(500 nm). (c–d) Si NWs epitaxially grown from the Au NPs shown in (a)

and (b), respectively. The scale bars are 5 mm in both images. (e) Cubic

root of the mean NP volume and mean diameters of Si NWs grown from

these particles versus immersion time in the aqueous galvanic displace-

ment solution. For 1 min. and above the NW diameters and NP volume

cube roots scale with each other, while for less than 1 min. NWnucleation

barriers skew the measured mean NW diameter to a higher value. Con-

necting Lines are guides to the eye and the error bars represent standard

deviations.

Fig. 5 Block copolymer templating and Si NW growth on Si(111)

samples with patterned wells in a SiO2 hard mask. The PS:PS-P2VP mass

ratio is 10 : 1. (a) Si NWs grown in �100 mm2 wells (scale bar ¼ 10 mm).

(b) Si NWs grown in circular wells 600–700 nm in diameter (scale bar ¼10 mm). (c) Higher magnification of individual Si NWs grown in the wells

shown in (b) (scale bar¼ 1 mm). (d) Number of AuNPs per well area. The

solid line represents expected values based on measurements of areal

densities on unpatterned samples.

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droplet surface energy relative to chemical energy in the droplet

volume as the droplet size is reduced.42 Since the VLS NW

growth rate is proportional to supersaturation, smaller NWs

grow much more slowly or not at all; the overall NW density is

then reduced and the mean NW diameter is skewed to a higher

value than expected based on the geometric argument given

above. Similar Au NP growth behavior is observed for higher

areal density micellar templates, though the overall NP growth

rate is reduced.†

A significant advantage to galvanic displacement is its

substrate selectivity, which permits highly controlled conformal

metal deposition in patterns without additional lift-off steps.

Therefore, we applied a block-copolymer template (PS:PS-P2VP

ratio¼ 10 : 1) to Si(111) substrates with a patterned silicon oxide

hard mask by dip coating, and deposited Au NPs via galvanic

displacement for 1 min. The selective deposition of Au NPs

means that Si NWs can only be grown in the wells exposing the Si

8812 | J. Mater. Chem., 2011, 21, 8807–8815

substrate, as shown in Fig. 5(a)–(c). Critically, in Fig. 5(c) it is

demonstrated that individual NWs less than �50 nm in diameter

may be grown in wells larger than 600 nm in diameter, achieving

a reduction in feature size greater than an order of magnitude

with the same approximate pattern registry. For uniform

coverage of the micellar template, the number of NPs and

subsequently grown NWs in a well will be proportional to the

well area, which is consistent with experimental data shown in

Fig. 5(d). Note that the solid line is not a linear fit to the data, but

rather the expected number of particles based on the AuNP areal

densities measured on unpatterned samples. Subsequently grown

NWs track the Au NP density well.

The conformality of galvanic displacement suggests that this

templating technique may also be applied to appropriately

patterned substrates with trenches having the correct orientation

of sidewalls (namely Si(111) for the Au/Si epitaxial NW growth

system studied in this report5,28) so that epitaxial Si NWs most

often grow perpendicularly from them for the fabrication of

devices based on in-plane epitaxial NWs. We primarily focus on

NWs grown from the 30 : 1 template (ra < 2 mm�2) since this is

the most likely template to enable the fabrication of devices

based on single or few NWs using only micron-scale lithography.

Au NPs were deposited using a 1 min. immersion time in the Au

deposition bath. Fig. 6(a) shows Si NWs grown in a �2 mm�2

wide trench to epitaxially bridge two electrically isolated single-

crystal Si electrodes on a buried silicon oxide layer. The top of the

electrodes were also covered with silicon oxide, preventing Au

cluster deposition on the top surface as well. Selective growth of

Si NWs can be enhanced if the trench sidewalls are also covered

with an oxide mask and a second lithography step is used to

This journal is ª The Royal Society of Chemistry 2011

Fig. 6 Epitaxial Si NW growth from templated Au NP seeds (ra < 2 mm�2) on silicon-on-insulator trench sidewalls. (a) SEM image of Si NWs grown in

a 2 mmwide trench (scale bar¼ 2 mm) (b) SEM image (bottom) of a single Si NW grown in a photolithographically defined Si window depicted in the top

schematic diagram (scale bar ¼ 2 mm). (c) Current–voltage response for the Si NW shown in (b) after ex situ boron doping. (d) Resistance versus the

physical scaling parameter inside parentheses in eqn (2) for 11 single Si NW or few-NW epitaxially bridging device structures on the same sample. The

red line is a linear fit to the data.

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define micron-scale windows of exposed Si, as for example in

Fig. 6(b). In this case a single Si NW epitaxially bridges the two

highly-doped single-crystal Si electrodes near a third electrode

that may be used for further device fabrication if desired (the

capacitive coupling with the NW is too small to yield any gate

dependent response in the current configuration). After Si NW

growth, the sample was doped with boron through an ex situ

method reported previously35 to achieve an active dopant

concentration greater than 10�19 cm�3. After doping, the top

layer of silicon oxide covering the Si electrodes was removed and

two-point current–voltage measurements were performed on

various single-NW or few-numbered NW device structures.

Fig. 6(c) depicts the measured current–voltage curve for the

single Si NW in (b) evidencing an ohmic response that was

typical of these device structures. The two-point current–voltage

responses for 11 such structures on the same die were measured.

The high doping level minimizes the surface carrier depletion

width so that we may assume that the full NW diameters are

involved in electronic conduction. For a given two point

measurement, the NWs which conduct current may possess

different diameters (D), but are all approximately the same

length (L). In addition, if we assume that the electrode and probe

tip resistances are negligible, and that the doping is uniform

This journal is ª The Royal Society of Chemistry 2011

across the substrate, the resistance of the NWs (RNWs) is given

by:

RNW s ¼ 2RC þ r

4L

pPN

i D2i

!(2)

where RC is the contact resistance and r is the resistivity of the

NWs. Plotting measured resistance versus the size parameter

contained in parentheses in eqn (2) (Fig. 6(d)), we derive a resis-

tivity value of 10� 5 mU cm and an average contact resistance of

7 � 3 kU. Based on the average cross sectional area of bridging

NW device structures, the estimated specific contact resistance is

(1.0� 0.5)� 10�6 U cm2. This value is superior to specific contact

resistances reported for Si NWs using deposited metal films,43,44

and is comparable or lower than previous values for epitaxial

NW contacts,22,45 though NiSi contacts to Si NWs have been

shown to exhibit lower values.46 Nonetheless, the electrical data

presented here reveal the great potential for fabricating devices

based on epitaxially integrated NWs with catalyst deposition

achieved via micellar block-copolymer templated galvanic

displacement.

Since the micellar template is disordered, the registry for

particle deposition with the larger photolithographic pattern is

J. Mater. Chem., 2011, 21, 8807–8815 | 8813

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statistical and thus imperfect; hexagonal ordering might signifi-

cantly improve upon this. Using polymers with more uniform

chain lengths would likely enhance the hexagonal packing to

a small degree. We avoided thermal annealing since the micellar

template is a trapped morphology that is not favored in thin film

or bulk phases and would likely transfer to an in-plane cylin-

drical morphology (for the given P2VP volume fraction in the

PS-P2VP) if thermally annealed.47 On the other hand, solvent

annealing typically induces non-equilibrium morphologies and

has been shown to improve hexagonal ordering of spherical

micelles.48,49 It may be possible to apply solvent annealing to

block-copolymer micelle/homopolymer systems like the one

described here, though we are not aware of any preceding work

on solvent annealing of block copolymer micelles with low areal

densities or homopolymer molecular spacers. Besides micelles,

equilibrated block copolymer thin films may exhibit excellent

ordering and uniformity with the help of chemical epitaxy50 or

graphoepitaxy,51 though significant effort has been focused on

pattern density multiplication,50,51 opposing the effort of this

work. Nevertheless, Stuen et al. have recently shown that the

domain spacing of polystyrene-block-poly(methyl methacrylate)

(PS-PMMA) block copolymer templates consisting of perpen-

dicular cylindrical minority domains (PMMA) may be adjusted

by a factor of 3 by concurrent PS and PMMA homopolymer

addition, though a substantial reduction in domain size unifor-

mity was observed and the pattern areal densities were still well

above 100 mm�2 (assuming hexagonal packing).52 Papalia et al.

have also recently demonstrated areal density tuning of spherical

minority polyisoprene domains in polystyrene-block-poly-

isoprene (PS-PI) over a very broad range by the addition of PS

homopolymer and controlled film thickness, though hexagonal

ordering was not apparent and the minimal areal density

reported was > �20 mm�2.53 In any case, we conjecture that the

evolving developments in ordering block copolymer thin film or

micellar templates may be leveraged with galvanic displacement

for the epitaxial integration of NW devices via precise substrate

and pattern registry in various geometries.

4 Conclusions

In conclusion, we have demonstrated a facile technique to inte-

grate epitaxially grown semiconductor NWs into device struc-

tures based on galvanic displacement of catalyst seed NPs

through PS-P2VP block copolymer micelle templates. Areal

density tuning from > �40 mm�2 to < �2 mm�2 is achieved

through mixture with PS homopolymer, and Au NP size is

directly controlled by immersion time in the aqueous plating

solution. The inherent substrate selectivity and conformality of

galvanic displacement enables epitaxial NW registry with much

larger photolithographically defined patterns in both horizontal

and vertical geometries, and the growth of single NWs with

diameters of�50 nm in wells with diameters larger than�600 nm

indicate pattern size reductions greater than one order of

magnitude. Electrical testing of ex situ doped epitaxially bridging

Si NW device structures further establishes the viability of this

approach.

In this report, we have concentrated on the galvanic

displacement of Au through NP templates as this is the most

commonly used catalyst metal for VLS NW growth. However,

8814 | J. Mater. Chem., 2011, 21, 8807–8815

the likelihood of internal54 or surface40 Au contamination that

may severely impair the performance of devices based on VLS-

grown Si NWs has generated considerable interest in other

catalysts that may be more CMOS compatible.55 Fortunately, as

noted earlier in this report galvanic displacement may be used to

deposit a number of different metals on various semiconductor

surfaces.27 In principle, all these metals may be directed to

deposit as NPs through the templating method described here.

Beyond NW integration, the capability to deposit noble metal

NPs selectively on semiconductor surfaces with controlled

diameters, areal densities and pattern registry may prove valu-

able in plasmonic applications such as enhancing light absorp-

tion in thin film photovoltaics56 or for plasmonic chemical or

biological sensing.57

5 Acknowledgements

We acknowledge the support of the National Science Founda-

tion, Grants No. EEC-0832819 (through the Center of Inte-

grated Nanomechanical Systems) and DMR-0804646. We also

thank Marta Fernandez, Dr Alvaro San Paulo, and Dr Noel

Arellano for the use of patterned substrates, as well as Megan

Hoarfrost and Prof. Rachel Segalman for providing the

polymers.

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