SCHOOL OF MATERIALS, THE UNIVERSITY OF MANCHESTER

A Comparison of nanowire production

methodologies Copper electrodeposition into etched ion-track polycarbonate and porous anodised aluminium

oxide templates

Matthew Webb

2013

Advanced Metallic Systems CDT

The University of Manchester

1

Abstract

Metal wires have been found to have increasing yield strengths with decreasing diameters

in the sub-micron regime. Previous work in Manchester has shown that there is a universal

scaling law for FCC metals. A lack of reliable method for the production of nanowires below

200nm in diameter has resulted in there being few studies of wires below this size.

Electrodeposition into a template is a promising method for producing wires of sufficiently

low diameter. Two templates were compared: commercially available etched ion-track

polycarbonate filters (EMD Millipore Corporation, Billerica, MA, USA), and porous anodised

aluminium oxide films grown by a two-stage anodising process.

Electrodeposition into the polycarbonate templates was successful, with a pore filling rate

of 90% achieved. However the templates were found to have a larger than specified pore

diameter and a large variation in the size of the pores (970nm ± 245nm for templates specified

as 800nm). This was due to overlapping pores formed during the random radiation production

method. Pore angles were also found to be non-perpendicular to the template surface. Thus, the

templates are not suitable for growing nanowires for mechanical testing. The alumina pores were

found to have a more suitable pore size distribution and pore angle; however issues with preparing

the templates for electrodeposition caused unsatisfactory deposition to occur.

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Acknowledgements

Many thanks to my supervisor Brian Derby for invaluable guidance and advice given

throughout this project. Thanks also to Rui Dou for experimental help and Joseph Ward for

his enthusiasm during his time on the project. Thanks to Michael Faulkner for his essential

help during SEM sessions.

A huge thank you to Claire Hinchliffe and Brad Wynne of the Advanced Metallic Systems

CDT for being supportive and understanding throughout my time at the CDT. Thanks also to

my course mates on the CDT for making my time there thoroughly enjoyable (and providing

a roof over my head!).

Finally, a massive thanks to my family, friends, band mates and other assorted loved ones

for their continuing support. I really appreciate it.

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Contents

1. Introduction .................................................................................................................................... 5

1.1. Aims and objectives ................................................................................................................ 7

2. Manufacture of nanowires ............................................................................................................. 8

2.1. FIB milling ................................................................................................................................ 8

2.1.1. Limitations and Drawbacks of FIB milling nanopillars .................................................. 11

2.2. Selective Etching of Eutectic Microstructures ...................................................................... 14

2.3. Electrodeposition into template ........................................................................................... 16

2.3.1. AAO templates .............................................................................................................. 16

2.3.2. Ion track templates ...................................................................................................... 22

2.4. Summary of production methods ......................................................................................... 28

3. AAO Film growth ........................................................................................................................... 30

3.1. Barrier Films .......................................................................................................................... 30

3.2. Porous Films .......................................................................................................................... 31

3.3. Pore regularity ...................................................................................................................... 34

3.4. Pore size ................................................................................................................................ 35

4. Deformation mechanisms and mechanical testing of nanowires ................................................ 36

4.1. Deformation mechanisms in sub-micron diameter nanowires ............................................ 36

4.1.1. Storage of geometrically necessary dislocations .............................................................. 37

4.1.2. Dislocation starvation and dislocation nucleation ............................................................ 38

4.2. Mechanical testing of nanowires .......................................................................................... 39

4.2.1. Testing in bending ......................................................................................................... 39

4.2.2. Nanoindentation ........................................................................................................... 42

5. Experimental methods .................................................................................................................. 48

5.1. Porous anodic aluminium oxide (AAO) templates ................................................................ 48

5.2. Etched ion track polycarbonate templates ........................................................................... 51

5.3. Electrodeposition of copper into templates ......................................................................... 51

6. Results and discussion .................................................................................................................. 52

6.1. Polycarbonate membrane pore analysis .............................................................................. 53

6.1.1. 800nm template analysis .............................................................................................. 53

6.1.2. 400nm template analysis .............................................................................................. 59

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6.1.3. 220nm template analysis .............................................................................................. 60

6.1.4. 100nm template analysis .............................................................................................. 62

6.1.5. Pore angle ..................................................................................................................... 62

6.1.6. Polycarbonate damage susceptibility ........................................................................... 63

6.1.7. Summary of etched ion track polycarbonate template analysis .................................. 66

6.2. Electrodeposition into polycarbonate membrane ................................................................ 67

6.2.1. Sample 1 ........................................................................................................................ 67

6.2.2. Sample 2 ........................................................................................................................ 69

6.2.3. A note on current density-time plots ............................................................................ 72

6.2.4. Sample 3 ........................................................................................................................ 72

6.2.5. Sample 4 ........................................................................................................................ 73

6.2.6. Summary of electrodeposition into polycarbonate templates ..................................... 76

6.3. Aluminium Oxide template growth ...................................................................................... 76

6.3.1. First stage anodisation .................................................................................................. 76

6.3.2. Ridges formed on aluminium ........................................................................................ 77

6.3.3. Second stage anodisation ............................................................................................. 78

6.4. Electrodeposition into AAO templates ................................................................................. 81

6.4.1. Sample 1 ........................................................................................................................ 81

6.4.2. Sample 2 ........................................................................................................................ 84

6.4.3. Sample 3 ........................................................................................................................ 87

6.4.4. Summary of electrodeposition into porous aluminium oxide templates ..................... 89

6.5. Etched ion track polycarbonate templates vs. AAO templates ............................................ 90

7. Conclusion ..................................................................................................................................... 92

8. Future works ................................................................................................................................. 94

References ............................................................................................................................................ 97

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1. Introduction

Size effects, in terms of dimensions of internal features, are well known to change a

material’s properties. An example of this is the Hall-Petch relationship in metals, which

states that the metal’s yield strength is inversely proportional to the square root of the

average grain diameter [1,2]. This relationship holds for grain diameters ranging from

millimetres down to tens of nanometres [3]. Sample size effects have been less thoroughly

investigated however, as classical mechanics traditionally states that a material’s properties

are independent of sample size.

Sample size effects have been observed, however, as early as 1951, where micrometre scale

metallic whiskers were shown to have an increasing yield strength in tensile testing with

decreased whisker diameter [4]. This has led to studies of the mechanical properties of

nanowires with sub-micron diameters.

Prior works have shown that this ‘smaller is stronger’ tenet is apparent in metal nanowires.

In the sub-micron regime a nanowire’s yield strength increases with reducing diameter [5,6].

A previous project in Manchester has proposed a relationship between the strength, ,

normalised by shear modulus, , and the wire diameter, , normalised by Burgers vector, ,

for face-centred cubic (FCC) nanowires (shown in Equation 1 and Figure 1) [7].

(

)

Equation 1

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Figure 1: Plot of Normalised shear stress against Normalised pillar diameter for a variety of FCC

nanowires showing linear fit [7].

This paper only includes one result with pillars of diameter less than 180nm. A contributing

factor to this is a lack of reliable methods for producing nanowires below approximately

200nm in diameter, with FIB milled pillars (with diameters greater than ~200nm) more

commonly investigated.

Above this diameter, both BCC and FCC nanowires have been produced via methods

including focussed ion beam (FIB) milling and by etching away the matrix of a eutectic

composite [5,8]. FCC metals have been investigated below this diameter using an

electrodeposition into anodised aluminium oxide (AAO) templates method, based on the

process by Masuda and Fukuda and developed by Dou and Derby [9][10][11]. More research

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on nanowires below that achievable via FIB milling is required to fully understand the

strengthening processed occurring.

1.1. Aims and objectives

The main aim of this project was to show that it is possible to produce copper nanowires via

electrodeposition in two different types of template: porous aluminium oxide and pre-

manufactured track-etched polycarbonate (EMD Millipore Corporation, Billerica, MA, USA).

A comparison of the two methods was also carried out. The steps taken to achieve this aim

were as follows:

A literature review was conducted on the various methods used to produce

nanowires, as well as the deformation mechanisms in the sub-micron regime and

mechanical testing mechanisms.

Training at techniques such as electropolishing, anodising and electrodeposition

required for the production of nanowires.

Training in analytical and imaging techniques including scanning electron microscopy

in order to view the templates and nanowires.

Literature review

Following is a review and summary of the different methods used to produce metal

nanowires, including: FIB milling, selective etching of a eutectic matrix and electrodeposition

into a variety of different templates. Different strengthening mechanisms that explain the

size dependent increase in strength have been suggested and these are discussed and

reviewed, as well as a variety of different mechanical testing methods used to characterize

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the materials. A section on the growth of porous aluminium oxide is also included due to its

relevance to the project as a template.

2. Manufacture of nanowires

Research into metal nanowires has been ongoing for a quarter of a century for a variety of

purposes, including in magnetic, electronic and optoelectronic devices [12]. Nanowires have

been produced using a variety of methods, each with benefits and limitations: size and

otherwise. The following section details the various methods of producing nanowires of

various metals reported in literature. Some of the nanowires reported were produced in

order to perform strength and deformation testing while some were produced for other

investigations.

2.1. FIB milling

FIB (Focused Ion Beam) milling of nanopillars was developed by Uchic et al. in 2002 and

reported in literature in 2004 [5,13]. A FIB machine can be used for either imaging or

sputtering/milling. For imaging applications a low beam current is used, whereas a high

beam current is used for sputtering and milling. For milling applications the ions in the beam

(usually but not exclusively Gallium+ ions) knock atoms out of their lattice sites. If sufficient

energy is transferred to the target atom it can create a collision cascade where the displaced

atom interacts with further atoms in the lattice [14]. The effects of this are discussed later in

this section.

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Figure 2: FIB image of FIB milled nanopillars with diameters of a) 860nm and b) 300nm [6].

In the Uchic et al. study nanopillars ranging from 500nm to 4000nm were machined from a

bulk sample crystal [5]. The location in the bulk was selected within a single grain so as to

produce single crystal nanopillars. Nickel was first used because the mechanical behavior of

bulk nickel is well known. An intermetallic alloy, Ni3-Al-Ta was also used to create

nanopillars as this has shown distinctly different flow mechanisms and so comparisons of

compression data could be made with the pure Nickel. The operating parameters of the FIB

machine, such as beam current, were not detailed. The diameters of the nanopillars

produced ranged from 4000nm to 500nm with aspect ratios of between 2:1 and 4:1 [5].

These aspect ratios were much lower than that of metal whiskers reported by Brenner in

1956 which had aspect ratios of around 1000:1 [4]. Mechanical testing of the nanopillars

was performed with a nanoindentor equipped with a flat-punch tip [5]. This is method is

discussed in Section 4.2. Mechanical testing.

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Figure 3: Schematic diagram of a FIB machine[14].

Uchic et al.’s method was developed by Greer et al. in 2005 and applied to gold for the first

time, and the minimum nanowire diameter was reduced to 400nm [6,15]. Gold nanopillars

with diameters of 400nm and 7450nm were produced via a FIB milling method with aspect

ratios of approximately 3:1. A single crystal <0 0 1> oriented gold disk was used as a

precursor. A gallium ion beam with beam current of 3000pA was used to machine a circular

crater with a 4000nm diameter pillar remaining in the center (see Figure 2). This pillar was

then machined into a pillar of the required diameter and aspect ratio by using finer beam

currents and by adjusting the tilt angle of the sample [6,15].

Further works on FIB milling of nanopillars have been carried out using a similar method to

that detailed above on a variety of materials, including gold, nickel, molybdenum and

niobium [5,6,15–26]. A summary of prior studies is shown in Table 1.

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Author Date Material Crystal structure Diameter

Uchic et al. [5] 2004 Nickel Ni3-Al-Ta

FCC 500nm – 4000nm

Greer & Nix [15] 2005 Gold FCC 400nm – 7450nm

Greer et al. [6] 2005 Gold FCC 400nm – 7450nm

Volkert & Lilleodden [16] 2006 Gold FCC 180nm – 8000nm

Greer & Nix. [25] 2006 Gold BCC 290nm – 7450nm

Kim & Greer. [23] 2008 Molybdenum BCC 200nm – 900nm

Brinckmann et al. [26] 2008 Gold Molybdenum

BCC 250nm – 946nm 205nm+

Shan et al. [17] 2008 Nickel FCC 150nm – 400nm

Frick et al. [27] 2008 Nickel FCC 165nm – 4500nm

Schneider et al. [19] 2009 Molybdenum Niobium

BCC

300nm – 3000nm

Schneider et al. [18] 2009 Molybdenum BCC 200nm – 5000nm

Kim & Greer. [22] 2009 Gold Molybdenum

FCC BCC

250nm – 1000nm

Kim et al. [21] 2009 Niobium BCC 200nm – 900nm

Kim et al. [20] 2010 Tungsten Molybdenum Tantalum Niobium

BCC

Han et al. [24] 2010 Vanadium BCC 183nm – 2790nm

Table 1: Summary of FIB milled nanowire studies.

2.1.1. Limitations and Drawbacks of FIB milling nanopillars

FIB milling of nanopillars is advantageous in that it can be applied to a wide range of sample

materials and it is relatively easy to produce a range of diameters of nanowire specimens to

study any size effect [16]. The method does however have a number of drawbacks, notably

contamination by the Ga ions in the beam implanting into the sample material and

consequent surface damage caused by irradiation [28]. Kiener et al. investigated a range of

aspects of the FIB process that could have an impact on mechanical properties. An

amorphous layer containing gallium ions of between 2nm and 50nm thickness, depending

on the beam current used, is formed on the surface of the sample. Gallium ions can also

penetrate and implant in the surface of the sample up 40 – 50nm, i.e. the same depth as the

amorphous layer. The mean surface roughness of a FIB milled sample was shown to be

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between 10.7nm and 16.4nm, an increase from a roughness of 10.6nm pre-milling. This

increase of surface roughness was independent of milling conditions such as incident angle

and beam current. FIB milling can also introduce defects into the sample material such as

point defects, dislocations and precipitates. These various effects were shown to have an

influence on the mechanical properties of the sample, with yield stress increases from

several MPa and up to GPa depending on the type of hardening occurring [29].

Figure 4: Schematic diagram showing possible changes caused to a nanowire by FIB milling. ta shows the amorphous layer at the surface of sample while td shows an area

with precipitates and dislocations [29].

A further study into the dislocations introduced by the FIB milling process was undertaken

by Bei et al. in 2007 [30]. This study suggests that the dislocations added by the FIB milling

13

decrease the strength of the nanopillars and that a FIB machined sample will not reach the

theoretical strength of the material whereas non-FIB machined samples have been shown

to achieve this [8,30].

In contrast to this a study by Shan et al. in 2008 argues that the FIB milling production

method actually causes the increase in strength demonstrated [17]. Two types of dislocation

are present in a FIB milled sample: loop-like dislocations that are small in size and larger line

defects that can extend through the entire pillar. It was suggested that the line defects were

pre-existing dislocations present in the bulk crystal pre-milling and that the small loop

dislocations were introduced by the FIB milling process. The study goes on to show that the

dislocations are removed via a process of mechanical annealing in the early stages of

deformation, and this dislocation-free state is suggested to be the cause of the high

mechanical strengths shown in FIB nanopillars [17]. This is also discussed in the Bei et al.

paper, which introduces the possibility that the FIB-damaged layer on the surface of the

sample has an increased hardness in comparison to the undamaged sample. As the samples

are reduced in diameter, the proportion of harder damaged material increases, increasing

the strength of the sample [8].

A further disadvantage of FIB milled pillars is the noticeable vertical taper in the pillars (see

Figure 5). This taper is more severe as pillars of lower diameter are produced [31]. Zhang et

al. reported that taper in the pillar results in a non-uniform strain distribution in the pillar

during mechanical testing resulting in an increased measured elastic modulus as well as

taper induced hardening [32].

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Figure 5: SEM images of Au nanopillars produced by electrodeposition [(a) to (h)], and FIB milling (i), showing noticeable taper in the FIB sample.

FIB milling has been shown to have an effect on the mechanical properties on nanowires, by

either strengthening or weakening mechanisms or a combination of both. The effects

discussed in this section intensify as the sample size is reduced, and this is perhaps the

reason that FIB nanopillars have not been produced with diameters of below 150nm. There

is therefore a need to look into alternate methods for the production of nano-scale samples.

2.2. Selective etching of eutectic microstructures

A novel method for producing nano-scale pillars for mechanical testing while avoiding the

issues introduced by the FIB milling process was developed by Bei et al. in 2007 [8]. Here,

Bei et al. furthered a method of growing a Nickel-Aluminium-Molybdenum matrix by

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directional solidification developed by the same group in 2005 [33]. An Ni-45.5Al-9Mo (at%)

alloy was melted and drop cast into a copper mold. An optical floating zone furnace was

used to directionally solidify the ingots. These rods were used as the feed material and

directionally solidified rods were used as seed material and were heated with a xenon arc

lamp. A homogeneous molten zone was ensured by rotating the seed and feed rods in

opposing directions. This process results in a rod-like eutectic microstructure with a NiAl

matrix and Molybdenum fibers as shown in Figure 6 [33]. In the 2007 Bei et al. study, the

molybdenum fibers were exposed by preferentially etching away the matrix. The diameter

of the exposed nanopillars was taken to be square root of the cross-sectional area of the

pillar due to their square cross-sections. Bei et al. produced nanopillars ranging from 360nm

to 1000nm using this method [8].

Figure 6: SEM images of NiAl-Mo eutectic alloy showing (a) transverse view, (b) longitudinal view and (c) high magnification transverse view, where the NiAl matrix has been preferentially etched [33].

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The preferential etching of a directionally solidified eutectic method of producing

nanopillars has advantages over the FIB method. However, there have been limited studies

using this method and the minimum diameter of nanopillars produced is 360nm. Whilst

eliminating the structural effect of the FIB process, this method does not reduce the lower

limit to diameters of nanowires. The method is also only viable with materials showing an

appropriate fibrous eutectic microstructure. Therefore a different method of nanowire

production is required to investigate the properties of nanowires of different metals below

the 150nm lower limit of FIB.

2.3. Electrodeposition into templates

A promising method of producing nanowires with sub-150nm diameters is by

electrodeposition into a porous template. This method is limited only by the size of pores

present in a template film. Templates, including lab-produced and commercially available

ion-track etched polymers and lab-produced porous aluminium oxide templates, are

discussed in this section.

2.3.1. AAO templates

Porous anodised aluminium oxide (AAO) templates have been shown as viable for the

electrodeposition of nanowires using a variety of materials. Aluminium oxide pores have

been shown to be self-organizing during anodizing under appropriate conditions [34,35].

Using an appropriate electrolyte a variety of different metals can be electrodeposited into

the pores. Further details on the growth of AAO templates and methods of increasing pore

regularity are included in section 3. AAO Film growth.

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A DC electrodeposition method was used by Dou & Derby in 2008 to produce gold

nanowires using AAO templates. A schematic diagram is show in Figure 7 and the following

labeled steps correspond to that diagram. An AAO template was produced using the

Masuda and Fukuda two-stage anodizing method (discussed in section 3.3) [9–11]. 99.999%

pure aluminium was first electropolished in a perchloric acid-ethanol mixture. One surface

of this substrate was anodised at 40V for 12 hours. To improve pore regularity the oxide film

formed was removed by etching in a 0.1M chromic acid and 0.1M phosphoric acid mixture

(a). The substrate was then anodised a second time for 2 hours (b). 0.1M phosphoric acid

was then used to etch the template and produce different pore diameters. PMMA was

coated onto the surface of the film (c) and the aluminium substrate was removed by etching

in 0.1M copper chloride solution (d). The oxide barrier layer was removed using copper

chloride solution (e). A gold layer was sputtered onto the template to act as a conductive

substrate (f). The PMMA layer was dissolved (g) and Gold nanowires were electrodeposited

into the template using an 0.02M KAu(CN)2 and 0.25M Na2CO3 electrolyte of pH13 (h). The

gold coated template was used as the working electrode and a platinum mesh was used as a

counter electrode with a voltage of -1.2V DC at an Ag/AgCl reference electrode. A 3M NaOH

solution was used to etch away the top of the AAO template, leaving the gold nanowires

exposed (i). Each nanowire was reported to be a single crystal with a random

crystallographic orientation resulting in no overall texture. Gold nanowires with diameters

30nm to 80nm were produced via this method. Different pore diameters were produced by

chemically etching the 30nm pores in 0.1M phosphoric acid for up to 60 minutes to thin the

cell walls and thus widen the pores to 55nm (30 minute etch) and 80nm (60 minute etch)

[11].

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Figure 7: Schematic diagram showing method of producing gold nanowires by Dou [36].

Dou also produced nickel and copper nanowires using the same method [36]. Nickel

nanowires were deposited using an aqueous solution of 0.5M NiSO4∙6H2O, 0.1M NiCl2∙6H2O

and 0.2M H3BO3. Aluminium oxide has been shown to dissolve at a pH of less 4.2 and so

35M Na2SO4 was added to adjust the pH value from pH4 to pH4.8 and 0.01M C7H4NNaO3S

was added to increase adhesion. This study reports a pore filling ratio of 90%, however the

electrodeposition was inhomogeneous and pores were filled at different rates. This causes

some pores to overflow with nickel and form balls on the surface of the template. A precise

19

ion beam polish system (PIPS) was used to create a smooth surface with nanowires of

uniform height [36].

Figure 8: SEM image of gold nanowires [36].

Gelves et al. reported in 2006 that the DC electrodeposition methods can be difficult due to

the need to remove the oxide barrier layer leaving fragile AAO films [37]. Gelves suggested

an AC electrodeposition method wherein the AAO barrier layer can be sufficiently thinned

to allow electrodeposition without the need to remove the aluminium substrate and the

oxide barrier layer formed between the aluminium and the pores. A two-step anodizing

process was performed in 0.3M sulfuric acid at 25V for 2 hours for the first stage and 4, 6 or

8 hours for the second stage. After completion of the second anodizing step the voltage was

gradually decreased from 25V to 9V over 16 minutes and held for 5 minutes to reduce the

thickness of the barrier layer (this process is known as current limitation. Barrier layer

20

thickness is proportional to anodizing voltage (see Section 3 for more details). Copper was

electrodeposited into the pores from a 0.5M CuSO4 and 0.285M H3BO3 aqueous solution

using a 200Hz sine wave at 10Vrms for 10-15 minutes. The aluminium substrate was used as

one electrode and a copper plate as the counter electrode. The proportion of pores filled

using this method is between 20% and 60% with the nanowires produced having a diameter

of 20nm. The nanowires were liberated and dispersed for SEM imaging by dissolving the

substrate and no mechanical testing was carried out [37].

An AC electrodeposition method was also reported by Borissov et al in 2009 [38]. Here the

author used a two stage anodizing process as mentioned above. After the second anodizing

stage the sample was etched in 2M sulfuric acid to reduce the barrier layer thickness and

widen the pore diameter. Iron was deposited into the pores using an alternating current at

300Hz and 35Vac, in an electrolyte containing 120g/L FeSO47H2O, 45g/L boric acid and 0.5

g/L ascorbic acid. Nanowires grown using this method had a mean pore diameter of 50nm ±

2nm. The proportion of pores filled was reported to be almost 100%, a feat attributed to the

thinning process of the barrier layer [38].

Figure 9: Schematic diagram showing advantages of AC electrodeposition over DC [37].

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A third method of pulsed electrodeposition (PED) has also been investigated. The process

was first reported by Nielsch et al. in 2000 and is similar to that of the AC electrodeposition

process [39]. A porous AAO template is produced using a two-step method, and the

aluminium oxide barrier layer is thinned by chemical etching or by current limitation [40].

Metal is electrodeposited into the pores using a millisecond voltage pulse followed by a gap

of around 1 second (see Figure 10). The negative pulse causes electrodeposition of the

metal in the pore bottom while the positive pulse discharges any capacitance in the barrier

layer as well as repairing any discontinuities in the oxide layer [40]. PED offers more control

over operation parameters than AC deposition such as ion concentration in the pores, as the

time delay allows for ion diffusion from the pore bottom to the pore opening [39]. Nielsch

reported a pore diameter of 45-55nm with a pore filling rate of almost 100%. This filling rate

is highly dependent on the amount of thinning of the barrier layer, and fewer pores are

filled in areas where the barrier layer is thicker [39].

Figure 10: Pulsed electrodeposition method signal plot [40].

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Electrodeposition into a porous alumina template can be performed using DC, AC or PED

methods. Both AC and PED methods have less production stages and as such are simpler.

Both however depend on the barrier layer being thin enough to allow ion transfer. DC

electrodeposition requires more production stages, including the removal of the original

aluminium substrate. This necessitates working with extremely delicate freestanding porous

alumina film. However, once the template is prepared for electrodeposition, the process is

more straightforward and deposition can be performed in the pores without the intervening

alumina barrier layer.

2.3.2. Ion track templates

An alternative to using porous AAO templates for the electrodeposition of nanowires is to

use etched ion track templates. These templates can be produced by chemically etching a

sample which has been exposed to ionizing radiation. This process was first demonstrated

by Possin in 1970 who produced tin, indium and zinc wires with diameters as small as 40nm

[41]. Radiation damage tracks in mica were etched in strong acid via a process

developed by Fleischer and Price [42]. Possin chose Mica as a template material due to its

large relative etching rates along the track axis to transverse to the axis. This results in

almost uniform pores throughout their 12μm length. The pores were filled by an

electroplating process and the produced wires have an almost uniform cross section with

diameters ranging from 40nm to 100nm [41].

A similar template method was used by Molares et al. to produce copper wires in 2001 [43].

Here the authors used polycarbonate foils between 30 and 40μm as a template. The foils

were irradiated with Au197 and Pb208 ions with sufficient energy per nucleon (11.4MeV) that

the penetration range was larger than the thickness of the foils and that the stopping power

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of the ions (dE/dx) was large enough to result in homogeneous etching of the pores. The

foils were then chemically etched with 6M NaOH solution at 50°C, with 10% methanol as

well as 1% surfactant. Different pore diameters of between 30nm and 200nm were created

by adjusting etching time. A 100nm gold layer was sputtered onto the membrane to serve

as a cathode for electrodeposition. A copper electrode was used as an anode in a two-

electrode electrochemical cell, with the sputtered gold substrate used as the cathode. The

anode used was cone shaped in order to ensure homogeneous deposition throughout the

sample. A high concentration copper sulfate pentahydrate solution (238g/L) was used as an

electrolyte to ensure a sufficient amount of copper ions were available for

electrodeposition. The electrical conductivity of the solution was increased by adding 21g/L

sulphuric acid, which also has the effect of lowering the cathode overvoltage. The wires

produced were shown to be either single crystals, or poly crystalline (depending on

electrodeposition conditions) by X-ray diffraction and TEM analysis. The authors conclude by

stating that polycarbonate ion track films are suitable for the production of nanowires up to

30μm in length and as low as 30nm in diameter.

Figure 11: SEM image of single (a) and poly (b) crystalline copper wires in pores as well as overgrown caps on the surface of a polycarbonate template [43].

24

A second paper published by Molares in 2001 provides a more in depth investigation into

the electrodeposition conditions that produce single or polycrystalline wires [44]. Single

crystals were produced when existing nuclei growth dominated the nucleation and

formation of new grains. Parameters thought to affect this, and thus investigated, were

current density (controlled by varying the electrodeposition cell voltage), temperature and

the type of electrolyte used. When a commercial electrolytic copper bath is used as the

electrolyte and a constant current density applied, increasing the temperature caused the

wires grown to be single crystals. The transition temperature between poly and single

crystals occurs at 60°C. Below this temperature, additions in the solution, including

brightening agents, are absorbs at the cathode’s surface, which increases the cathode

overvoltage and thus increases the electrode polarization. Increasing the temperature

decreases this cathode overvoltage, as well as increases both surface diffusion and efficient

transport of ions towards the electrode. A combination of these effects is attributed to

encourage growth of existing nuclei, and thus create single crystal wires. In a simple salt

electrolyte (the same 238g/L copper sulfate pentahydrate solution as used in the previous

paper), decreasing the current density was shown to favour have the same effect as

increasing the temperature in the previous setup and single crystal nanowires were grown

for current densities below 35mA/cm2 at both 50°C and 60°C.

The two Molares papers reveal four different regimes during electrodeposition (see Figure

12). A peak in current at the beginning of electrodeposition is attributed to the charge of the

double layer and the creation of the diffusion layer. A constant (albeit slightly decreasing)

current is then shown during growth of the nanowires in the pores. Once the pores fill, caps

begin forming on the film surface, increasing the electrodeposition area and thus increasing

25

current flow. Finally the caps combine to become a layer and current flow becomes

constant again (at a higher rate than during nanowire growth due to the increased area)

[43]. Increasing the current density (by increasing the voltage) causes the copper to fill the

pores more rapidly as expected. Current density was also shown to increase lineally with

applied voltage. This knowledge of current flow is important during electrodeposition as in

order to observe the nanowire grown it is preferable to stop the process at the end of the

2nd phase before caps are grown on the surface of the film. Performing the

electrodeposition at a lower voltage and increasing the time before caps form would offer

more control.

Figure 12: Schematic current time diagram (left) and current time curves for different applied voltages (right) during copper deposition into polycarbonate pores. [43,44]

Enculescu et al. used a similar method to grow copper nanowires in a polycarbonate ion

track template, however unlike Molares et al. who grew a ‘forest’ of wires, single nanowires

were produced [45]. Enculescu argued that single nanowires can be used to more accurately

measure electric properties. A 30μm polycarbonate film was irradiated by a single heavy ion

with kinetic energy of greater than 1GeV. A gold layer was sputtered onto the back of the

foil as support and for use as a cathode and the ion track damage trail was etched in a

26

solution containing sodium hydroxide and methanol. The composition of the etching

solution was show to have an effect on the pore shape: as it adjusts the etching velocity

along the track and the etching rate of the bulk at a different rate. In an etching solution

with no methanol pore diameter was measured to be 160nm at the pore opening and 40nm

close to the tip and the pore was described as almost cylindrical. Adding methanol to the

etching solution increases the etching rate across the bulk, , compared to the etching

rate along the track, , and thus increases the pore angle (see Equation 2). This was

thought to be an effect related to the rate of diffusion of the etchant in the pore.

( ) (

) for

Where is the pore radius at a distance along the track, is the etching time.

Equation 2

This is an important result when designing an experimental setup for creating nanowires by

a template and electrodeposition method as it shows that the wires grown adapt the same

form as the pores in the template.

Unlike the regularly ordered templates produced by a two-stage anodizing process on

aluminium, the ion track templates discussed above have either randomly spaced pores

(both Molares et al. papers) or produce single pillars (Enculescu et al.). Burek and Greer

claim that this is primarily due to limited control of pore size and placement in conventional

methods [31]. Here the authors propose an electron beam lithographic method that allows

full control over pore size and placement. This allows pores to be produced in sizes ranging

from 750nm to 25nm without the taper associated with etching of heavy ion track methods.

A 100nm gold layer is deposited onto a silicon substrate with a 20nm titanium layer

27

adhesion layer between. The gold here acts as a cathode during electrodeposition as in

other electrodeposition methods. A PMMA layer was spin coated onto the substrate at a

thickness closely matching the height of the nanopillar required and baked to set. An

electron beam lithography system was used to expose the PMMA in a computer generated

pattern. Precise isolation of nanopillars was possible and pillars were spaced up to 50μm

apart. The PMMA was developed immediately after exposure in a solution of

methylisobutylketone and isopropyl alcohol for 60s. This creates a template with attached

cathode layer akin to similar electrodeposition methods discussed in this section.

Figure 13: (a) SEM image of regularly ordered array, with both tension and compression samples, (b) to (i) FIB images of tension (with caps) and compression nanopillars.

Burek and Greer proceed to grow both gold and copper nanopillars. Gold was deposited in a

commercially available [Au(SO3)2]2- plating bath, and copper was deposited using a 125g/L

copper sulphate pentahydrate and 50g/L sulphuric acid solution. Both Galvanostatic (DC)

and reverse pulse (AC) conditions were used for electrodeposition. Deposition was stopped

when the pillars filled the pores for compression samples, and caps were allowed to form on

the surface of the template for to provide grip points for tension samples (see Figure 13).

28

Finally the PMMA template was removed in a bath of acetone. The gold nanopillars grown

were limited to nanocrystalline microstructure, with only small changes to grain size

possible. This was the case for both AC and DC deposition. This was attributed to the

electrolyte, specifically the instability of the gold sulfite salt in solution. For copper

deposition, DC current densities of below 20mA/cm2 yields single crystal pillars. This is a

lower value than the 35mA/cm2 reported by Molares et al. [44]. Deposition below

10mA/cm2 was not possible due to current being used by side reactions. AC deposition was

found to increase the percentage of single crystal pillars grown. It was found however, that

AC deposition can have effects on future mechanical testing by introducing voids and by

creating inhomogeneous pillars.

2.4. Summary of production methods

FIB milling has been used to produce nanopillars ranging from 150nm to 8000nm and in a

variety of different materials. Pillar size and aspect ratio can be selected and created with

minimal manufacturing limits. However, FIB milled pillars have a number of disadvantages.

Firstly, radiation effects from the ion beam can cause an amorphous layer up to 50nm on

the surface of the pillars, increase the surface roughness of the pillar, as well as introduce

point defects, precipitates and dislocations into the pillar. These can all have an effect on

the mechanical properties, with strengthening mechanisms, weakening mechanisms, or a

combination of the two. These effects are increased as pillar diameter is decreased,

imposing a lower limit on pillar diameter produced. FIB milled pillars also have a noticeable

vertical taper which causes non-uniform strain gradients during mechanical testing further

effecting mechanical properties.

29

Etching away the matrix of a directionally solidified eutectic results in nanopillars without

many of the disadvantages of FIB milled pillars. There is none of the radiation damage

effects associated with FIB milling and pillars are taper free. However, the method has only

been used to produce molybdenum pillars and selection of pillar diameter is limited. The

minimum diameter produced via this method is 360nm.

A final method of producing nanowires is by electrodeposition into a template. Templates

can be made with a variety of materials, with popular options being regularly ordered

porous aluminium oxide and ion track polymer templates. Nanowires grown have the same

form as the pores in the template, and as such the shape and size of wires grown is limited

only by the shape and size of the template available. Aluminium oxide templates have been

used to grow wires from 30nm up to 450nm, and pillars grown are in a regular hexagonal

array. Disadvantages of aluminium oxide templates is the complex production method

required to remove the barrier layer oxide formed between the porous oxide and the

aluminium substrate, however an AC electrodeposition process can simplify this. Polymer

templates have been used to produce wires from 25nm to 750nm. This upper value is not a

hard limit, and is only in place due to lack of relevance of studying wires above this

diameter. The pores in the templates can be produced by either etching radiation damage

tracks caused by the bombardment of heavy ions, or by electron beam lithography.

Nanowires of a variety of different metals can subsequently be grown in the templates by

electrodeposition. Single crystal or polycrystalline pillars can be grown depending on

electrodeposition conditions. The electrodeposition into a template method produces wires

that lack the radiation damage and tapering effects of FIB milled samples, and also the

method overcomes the size limitations imposed by both the FIB milling and etching of

30

eutectic methods. Electrodeposition into a template therefore provides an ideal method for

the production of nanowires for study.

3. AAO Film growth

As this project involves the production of porous aluminium oxide templates, this section is

included to discuss the theory of oxide growth and pore regularity.

Anodic alumina films can be grown in two distinct varieties: Barrier type films, which consist

of a uniformly thick amorphous layer of alumina that are often utilised as dielectrics, and

porous films, which consist of a regular series of alumina pores formed above a thin barrier

layer and have practical uses in corrosion and wear protection, as well as providing keying

layers in adhesive bonding and base layers for paints [35,46]. These films are formed by the

electro-chemical process of anodising (anodic oxidation), where the sample is used as the

anode in a circuit as direct current is passed through an aqueous electrolyte.

3.1. Barrier Films

Barrier type films are formed in electrolytes at a near-neutral pH. Al3+ ions are produced via

oxidation of aluminium atoms in the sample, and O2- ions are derived from the electrolyte.

The applied electric field causes the solid state migration of the ions towards the cathode

and anode (sample) respectively (see Figure 14). The film is formed at both the anode-film

interface and the film-electrolyte interface, with a growth ratio theoretically proportional to

the current carried by the ions: 40% at the anode-film interface, corresponding to the O2-

31

ions, and 60% at the film-electrolyte interface, corresponding to the Al3+ ions [47].

Figure 14: Schematic diagram of ion migration in the formation of a) barrier type films, and b) porous type films [48]

3.2. Porous Films

Porous films are formed in non-neutral pH electrolytes. In 1978 Thompson et al. showed

that during anodisation a non-uniform growth rate of the anodic film occurs, with thicker

film material forming above any ridges in the surface of the sample [35]. These ridges can be

formed during an electropolishing process performed on the sample during preparation, by

etching or mechanical polishing and scratching. This results in a cellular topographical

appearance on the film identical to that of the original aluminium. These ridges develop to

form roughly hexagonal cells. Pore formation occurs within these preferred cells and

continued anodisation causes pore diameter growth in order to maintain a uniform barrier

32

layer thickness, which is dependent on the voltage applied. Once the cells reach a diameter

where they interact, a steady-state porous morphology is formed.

This description was further developed by Shimzu et al. [49]. In a non-neutral pH electrolyte,

a barrier type film can only grow at the anode-film interface as outwardly migrating Al3+ ions

are lost to the electrolyte via field-assisted dissolution as well as direct injection, rather than

forming oxide. Once a certain thickness of barrier type film is formed, increased local

stresses caused by alumina being less dense than aluminium causes cracking. This critical

thickness is proportional to the applied voltage. Stress concentration factors are higher at

any protuberances in the anode, which causes any cracks that form to typically occur at

these locations. The cracks are rapidly healed at a high current density, which leads to the

thickening of the film at these locations (see Figure 15).

33

Figure 15: Schematic diagram showing the formation of thicker oxide film above ridges in the sample [49].

The persistent film cracking and repeating events cause the protuberances to grow thicker

and wider. This causes an increase in curvature in the film in the regions between the

protuberances, and an increase in the field in these areas. This increases the rate of field-

assisted dissolution of the Al3+ ions into the electrolyte until the rate of thinning becomes

equal to the rate of growth caused by the O2- ions. This results in the voltage dependent

34

uniform thickness of barrier layer observed by Thompson. At this time, barrier layer growth

is overtaken by pore growth.

Once the surface of the sample is covered by such pores, geometric constraints between the

pores force the pores to rearrange slightly into a regular hexagonal formation. This is

possibly due to a reduction in the diameters of both the cells and pores [50]. The result of

this procedure is a regularly ordered porous film above a barrier film with thickness

proportional to the applied voltage.

3.3. Pore regularity

To improve pore regularity a further process must be carried out before anodising. A

pretexturing process can be carried out where a master with hexagonal concaves is pressed

into the aluminium [51,52]. This results in a negative of the same topography on the surface

of the aluminium and serves to enhance the regularity of the ridges used for nucleation of

pores as discussed above. An alternative method for improving pore regularity is a 2-step

anodising process developed by Masuda et al. and used by Dou [10,36]. In this method a

long period anodising step is performed to produce a thick porous oxide on the surface. This

forms a hexagonal concave pattern on the surface of the aluminium as discussed in section

3.2. The oxide is then chemically etched off and the aluminium is anodised again, with the

textured ridges serving as nucleation points for the growth mechanisms discussed above.

Pore regularity is also affected by the choice of electrolyte, as well as the anodising voltage

[10][53][54].

35

3.4. Pore size

The use of different acid solutions during anodizing has an effect on the size of cell formed.

Larger cells allow larger pores to be produced as the pores must be smaller than the cell

size. Li et al. reported the pore distances (hence cell sizes) of three different acids: Sulfuric,

oxalic and phosphoric [12]. Each acid has an optimal anodizing voltage to produce regularly

ordered pores. Sulfuric acid produces 60nm cells at 25V; Oxalic acid produces 95nm cells at

40V; Sulfuric acid produces 420nm cells at 160V. Pore distance was found to increase

linearly with anodising voltage. These values largely correspond to that of Masuda et al.

who report sulfuric acid producing 65nm cells at 27V, oxalic acid producing 100nm cells at

40V, and phosphoric acid producing 500nm cells at 195V [10,53,54]. The cell size and

anodizing voltage for phosphoric acid for the two groups differ somewhat, however both

agree with an approximate 2.5nm/V proportionality constant between applied voltage and

cell size as reported by Masuda [53]. Li et al. themselves reported a proportionality

constant of 2.81nm/V [12].

36

Figure 16: SEM images of porous alumina formed in sulphuric acid at (a) 20V, (b) 23V, (c) 25V, and (d) 27V, showing increasing regularity [54].

4. Deformation mechanisms and mechanical testing of nanowires

4.1. Deformation mechanisms in sub-micron diameter nanowires

In order to design a successful experiment to investigate the deformation of nanowires it is

important to understand the deformation mechanisms occurring. It has been shown that in

37

a sub-micron regime, nanopillars strength increase with decreasing diameter [5,7]. The

precise cause of this effect is still unclear, however a number of mechanisms have been

suggested: the storage of geometrically necessary dislocations, dislocation starvation and

limited dislocation nucleation.

4.1.1. Storage of geometrically necessary dislocations

An indentation size effect is well known in the mechanical testing of metals [55–57]. Here

the hardness of the sample during indentation testing increases as the size of the indention

decreases. This effect is more pronounced in the sub-micron indentations. Fleck et al. report

that during indentation the strain directly below the tip of the indenter is large while the

strain far from the tip is zero [58]. This results in a strain gradient across the deformation.

This strain gradient necessitates a density, , of geometrically necessary dislocations

(GNDs). The density of these dislocations increases with the inverse of the length scale (see

Equation 3).

Where is the dislocation density, is the macroscopic plastic shear

stress, is the magnitude of the Burger’s vector and is the local

length scale of the deformation field.

Equation 3

Nix and Gao further this description by stating that the density of GNDs increase with the

inverse of the indentation depth (see Equation 4 and Figure 17) [59].

Where is the dislocation density, is the magnitude of the Burger’s

vector, is the indentation depth and is the angle of the indentation tip.

Equation 4

38

Figure 17: Geometrically necessary dislocations (GNDs) formed by indentation testing [59].

Nix and Gao proceed to show that an increase in the number of GNDs create a hardening

effect responsible for this indentation size effect: The square of the indentation hardness is

proportional to the reciprocal of the indentation depth.

4.1.2. Dislocation starvation and dislocation nucleation

Greer and Nix present an alternative strengthening mechanism in the absence of strain

gradients in small volumes [6,15]. During the deformation of a bulk material, dislocation

motion leads to the multiplication of dislocations through mechanisms such as double cross

slip and Frank-Read sources. This leads to an increased dislocation density and strain

hardening of the material, increasing its yield strength. In a sufficiently small sample such as

nanowires, dislocations can only travel a short distance before coming into contact with a

free surface, due to the large surface area to volume ratio, and annihilating. This reduces

both the number of dislocations in the sample, as well as the possibility of dislocation

39

multiplication, leading to a dislocation-starved state. For further deformation to occur new

dislocations must be nucleated. This requires high stresses, and thus leads to increased yield

strength.

Volkert and Lilleodden provide a further description of this dislocation hardening [16]. In a

bulk material the dislocation density increases during deformation as discussed above. The

increased dislocation density leads to large internal stresses which can activate dislocation

sources. In a dislocation starved state, the internal stresses of the sample will be lower.

Thus, while large samples can have high dislocation density and can more easily reach the

local stresses required for dislocation nucleation, smaller samples have a lower dislocation

density and lower internal stress and as such require higher applied stress to deform.

4.2. Mechanical testing of nanowires

As discussed above in the introduction, nanowires have frequently been shown to increase

in strength with decreasing diameter in the sub-micron regime. Proposed mechanisms for

this effect have been discussed above. Experimentally, various mechanical testing methods

have been used to demonstrate and investigate the effect.

4.2.1. Testing in bending

Wu et al. investigated gold nanowires with diameters ranging from 40nm to 250nm in

bending [60]. Wires were produced using an electrodeposition into porous alumina and

subsequently released from the original substrate and dispersed onto a new substrate

containing FIB milled trenches. Wires found to span the trenches were secured at both ends

by electron beam induced deposition of lines of platinum. An AFM (atomic force

microscope) cantilever was used to place the nanowire under lateral loading. Figure 18

40

shows a schematic of the experimental setup while Figure 19 shows a 200nm sample in

various states of deformation.

Figure 18: Experimental setup of nanowire bending using AFM [60]

Figure 19: AFM (c,d,e) and SEM (f) images of a 200nm gold nanowire before and after plastic deformation [60]

41

Wu reported that, while the young’s modulus, E, of nanowires remain independent of wire

diameter, the yield strength increases up to 100 times that of the bulk with decreasing wire

diameter and ultimate strengths approaching theoretical (E/10) for the smallest wires

tested.

Figure 20: Young’s modulus and Yield strength for various different nanowire diameters under bending [60]

42

The results of this paper agree with the general theory that nanowires strengths increase

with decreasing diameter. However, as the test was performed in bending, large strain

gradients were introduced, which increase with decreasing length scales i.e. wire diameter.

This makes it difficult to differentiate between strain gradient hardening methods and other

true length scale strengthening mechanisms such as dislocation starvation and nucleation.

Thus it is not possible to determine whether a decrease in sample size causes a fundamental

change in deformation mechanisms in the absence of strain gradients.

4.2.2. Nanoindentation

Uchic et al. proposed a mechanical testing method based on traditional uniaxial

compression testing in order to eliminate strain gradients [5]. Here, nickel and nickel alloy

nanopillars between 500nm and 40μm were produced via a FIB milling method. To perform

a compression test, a nanoindentation system was fitted with a flat-punch tip. Compression

tests were performed at displacement rates of between 1nm/s and 5nm/s and single

nanopillars were tested at a time. Uchic reported an increased strength of up to 25 times

the bulk value for the smallest pillars tested (see Figure 21 and Figure 22).

43

Figure 21: Engineering stress strain curves for bulk and FIB milled nickel tested in compression (A) and SEM image of a 20μm nanopillar pre (B)and post (C) deformation [5]

Figure 22: Engineering stress strain curves for bulk and FIB milled Ni3Al-T alloy tested in compression (A) and SEM images of 20μm nanopillars after deformation [5]

Greer and Nix used a similar uniaxial compression testing methodology on single FIB milled

gold nanopillars [15]. The diamond tip was etched off a standard Berkovich indenter leaving

a flat circle with a 9μm diameter. This was used to compress single nanopillars between

44

400nm and 7450nm in diameter at a strain rate of 0.004s-1. Pillars were located using an

optical microscope calibrated to the indenter. Unlike Uchic et al. pillars were subjected to

multiple loading/unloading cycles to ensure elastic unloading. The smallest pillars tested

achieved a yield strength of 550MPa, 18 times that of bulk. Elastic moduli of 43GPa was also

recorded, similar to that of bulk gold of the same crystal orientation.

Figure 23: Stress-strain curves for FIB produced nanopillars showing loading/unloading cycles and increased yield strength with decreased pillar diameter [15].

A similar method of using a flat-punch tip indenter has been used to test the mechanical

properties single FIB milled pillars by other groups [6,16–18,20,22,23,25,27,61,62]. The

method has also been used to test individual pillars produced by an etching of eutectic

method [8]. All of these investigations have shown evidence for the ‘smaller is stronger’

phenomenon. Performing uniaxial compression tests ideally produce no strain gradients in

the sample and are therefore useful to investigate other strengthening mechanisms in the

45

sub-micron regime. However, FIB milled pillars often show at least a small degree of taper in

the side walls of the pillar and so strain gradients are introduce through the length of the

pillar (as discussed in the FIB production method above). Nano-indentation compression

testing is therefore more suitable for nanowires produced in a taper free method.

All the reports discussed above compress only single pillars at a time. This can be time

consuming to test a statistical significant number of pillars. Also, only low aspect ratio pillars

can be tested without the risk of buckling (aspect ratios of between 2 and 3 have been

recommended [32]). There is also a limit on how small flat-punch tip can be created, and in

various methods of production the wires are too close together to be compressed

individually. Dou and Derby compressed a large number of gold and nickel nanowires grown

by electrodeposition into an alumina template in parallel [11,36]. A 10μm diameter

cylindrical diamond flat punch tip was fitted to a nanoindentor system. The diameter of this

was much greater that the diameter of the nanowires (30nm-80nm), and much larger than

the average cell size (that is the average size of template pore plus pore wall) of 100nm.

Therefore it is not possible to test individual wires without removing wires from the

substrate. A Precise ion polish system was used on electrodeposited wires in an alumina

template to create a uniform height for an area of wires. The wires were then compressed

in parallel and SEM images were used to determine the number of wires compressed. Dou

reported that the Young’s modulus of the gold nanowires to be 79.4GPa, very close to that

of polycrystalline gold of 80GPa. The Young’s modulus reported by Dou is comparable to

that of the polycrystalline bulk due to it being the average of many pillars of random

orientations. The wires showed increased yield strength with decreasing wire diameter. The

elastic unloading shows that no collapse or buckling of the wires takes place during testing.

46

This was thought to be due to the stabilising effect of the large density of nanowires. The

stress-strain plots recorded showed few features compared to other nano-indentation

compression plots (performed on single nanopillars). This was attributed to the stress data

being an average of multiple wires (N>1000).

Figure 24: Side view of the Diamond flat punch tip used by Dou and Derby[36]

47

Figure 25: Nanowire forest after compression testing showing residual indentation [36]

Figure 26: Engineering Stress-Strain curves for 30-70nm nanowires tested in compression [11]

48

To conclude, mechanical testing is possible on sub-micron scale samples. To avoid significant

strain gradients effects a uniaxial compression test has been developed and used on a

variety of different nanowires. A nanoindentation machine is fitted with a flat punch tip to

compress either single nanopillars or a forest of nanowires. FIB produced nanopillars can

still have strain gradients across the length of the pillar and so uniaxial compression tests

are best suited to nanowires produced in a taper free method, such as electrodeposition

into a template. All papers discussed above show an increase in yield strength with a

decreased wire diameter.

5. Experimental methods

This project used an electrodeposition into template method to produce nanowires. Porous

anodic aluminium oxide (AAO) films and prefabricated etched ion track polycarbonate were

used as templates.

5.1. Porous anodic aluminium oxide (AAO) templates

A 2 stage anodizing process based on that of Masuda and Fukuda and developed by Dou and

Derby was used to produce regularly ordered aluminium oxide templates [10,11]. The

process is summarised in Figure 27.

49

Figure 27: Schematic diagram showing production stages of AAO template and subsequent

nanowires. i) Aluminium substrate after 1st anodising stage and subsequent oxide removal, ii) 2nd anodising stage and pore widening, iii) Coating with PMMA, iv) Aluminium etched away, v) Oxide

barrier layer etched away, vi) Oxide coated with metal, vii) PMMA removed, viii) metal electrodeposited into pores, ix) Oxide template etched away.

50

Samples were cut from a high purity (99.999%) aluminium sheet to form the substrate

material. The substrates were prepared for electropolishing by degreasing in acetone. The

substrates were then electropolished at 20V for 5 minutes in a 1:4 by volume mixture of

perchloric acid and ethanol (VWR International, Lutterworth, UK; Fisher Scientific,

Loughborough, UK). Care was taken that the solution did not reach above 10°C by

performing the procedure in an ice bath.

The electropolished substrates were anodised in a 0.3M oxalic solution (Fisher Scientific,

Loughborough, UK) for 12 hours at 40V. This produces an aluminium oxide film on the

surface of the substrate as discussed above in section 3. The film microstructures were

characterised using scanning electron microscopy (SEM) (XL 30, Philips, Amsterdam,

Netherlands).

The aluminium oxide film formed by the first anodising stage was removed by chemical

etching in a mixed solution of 0.1M chromic acid and 0.1M phosphoric acid (Fisher Scientific,

Loughborough, UK). This was done at 60°C for 2 hours. SEM images of the ridges formed on

the aluminium were taken.

The second stage of anodising was carried out in conditions identical to the first stage for 2

hours. The anodising time is proportional to the final thickness of porous film. Again, film

microstructures were characterised using SEM.

2g poly methyl methacrylate (PMMA) was dissolved in 20ml acetone and a layer was applied

to the porous alumina using a cotton tip. Once dry the aluminium substrate was removed by

etching in 0.1M copper chloride solution (Fisher Scientific, Loughborough, UK) for 5 hours.

The exposed alumina barrier layer was removed by etching in 0.1M phosphoric acid for 1

51

hour at 30°C. A gold/palladium coating system (Model 682 PECS, Gatan Inc., Pleasanton, CA,

USA) was used to provide a new substrate to support the oxide film, as well as to act as the

cathode during electrodeposition. The thickness of the gold coating must be such that the

pores in the template are completely covered; we estimated a thickness of coating of

approximately equal to that of the pore diameter to be sufficient. Any further coating is

unnecessary, however a thicker layer results in the easier handling of the sample. Finally the

PMMA layer was removed by dissolution in acetone.

5.2. Etched ion track polycarbonate templates

Prefabricated etched ion track polycarbonate filters were purchased in a variety of pore

sizes (100nm, 200nm, 400nm and 800nm) (EMD Millipore Corporation, Billerica, MA, USA).

To prepare the samples for electrodeposition an identical gold/palladium coating process to

that used on the AAO templates was performed. This is the only pre-processing stage

required as the filters already have the pores at the specified diameter.

5.3. Electrodeposition of copper into templates

Once each sample type was gold coated (and the PMMA layer removed in the case of the

AAO template) the further preparation and electrodeposition stages were identical for each

type of template. The template was cut into sample sizes of approximately 0.25cm2. This

was purely for practical reasons as it meant that each template prepared could be used for 4

experiments. The samples were then mounted to a 3cm x 1cm copper sheet. An electrical

contact was made between the copper and the gold/palladium coating by using either silver

conductive paint or conductive SEM sample mount pads. Lacomit lacquer (Xylene, butanol,

Butyl acetone and acetone solution) was used to secure the sample to the copper mount

52

and to cover the rear side and edges of the copper mount to prevent current leakage

through additional electrodeposition occurring in areas other than in the template.

An aqueous solution of 250g/L copper sulphate pentahydrate (CuSO4.5H2O) was used for

electrodeposition. This was based off the finding of Molares et al. and Enculescu et al. who

both reported positive results [44][45]. This is below the solubility limit of copper sulphate

pentahydrate of 316g/L, whilst still providing a large concentration of copper ions for

deposition. Electrodeposition was performed using a potentiostatic method at room

temperature. The Au/Pd coating on the mounted sample served as the target cathode and a

copper anode was used. A calomel reference electrode was also used as a reference. DC

voltages between 1500mV and 200mV were used for time periods ranging from 10s to 300s.

6. Results and discussion

Scanning electron microscopy (SEM) was used to image the etched ion track polycarbonate

templates and an analysis of pore size was carried out. Images of the polycarbonate

templates after electrodeposition were also studied. SEM images at multiple stages of

anodised aluminium oxide template production were taken, as well as images after

electrodeposition. The two different types of templates were compared, as were the

nanowires grown in both templates.

53

6.1. Polycarbonate membrane pore analysis

6.1.1. 800nm template analysis

Four different pore size filters were purchased: 100nm, 220nm, 400nm and 800nm (EMD

Millipore Corporation, Billerica, MA, USA). A scanning electron microscope (SEM) was used

to acquire images of each different pore size. A thin layer of gold/palladium was coated

onto the templates to reduce electron charging within the SEM. This layer was

approximately 3nm thick and so does not significantly affect pore size measurements. A

relatively low accelerating voltage was used to further discourage charging. Figure 28 shows

an SEM image of a template with pores specified at 800nm diameter.

Figure 28: SEM image of 800nm pores in a polycarbonate template. Template has been coated in a

~3nm layer of Au/Pd to avoid electron charging.

54

The pores in the filter are clearly not in a regular order. This is due to the production

method used to produce the filters being similar to that of Molares et al. as discussed above

[43,44]. The polycarbonate films were exposed to radiation bombardment and the damage

tracks were etched to produce pores of the specified diameter. This results in a template

with no pore ordering; some areas of the template have a large concentration of pores and

some areas are devoid of pores. Areas with a large concentration of pores often have

overlapping pores, the effects of this on pore size are revealed once the image is put into

image analysis software.

ImageJ software (ImageJ, U. S. National Institutes of Health, Bethesda, MD, USA) was used

to perform a pore size analysis of the template [63]. The scale of the image was calibrated

using the scale bar included in the original SEM image (Figure 28). A threshold was applied

to convert the image to a black and white binary image (Figure 29). To improve the accuracy

of pore measurement partial pores at the edge of the image, as well as any dark areas of

template picked up by the threshold from the original image were removed by selecting all

particles under a critical diameter; 800nm pores have an area of 0.5μm2 so a value of 0.25

μm2 was selected as a critical diameter to ensure no complete pores were selected (see

Figure 30). Figure 31 shows that no complete pores were removed from the image.

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Figure 29: 800nm pores after threshold applied to convert image to binary.

Figure 30: 800nm pores after incomplete pores and areas of template were removed.

56

Figure 31: Incomplete pores and dark template removed from Figure 30 showing no complete pores

were removed.

Particle analysis was performed on the remaining black areas, i.e. pores. 528 particles were

analysed, with a mean diameter of 1034nm ± 281nm (standard deviation). Whilst the

quoted value of pore diameter of 800nm is within the standard deviation of the calculated

mean, it is still considerably higher. The standard deviation is 27% the value of the mean.

The reason for both the high mean and large standard deviation is related to the random

nature of pore placement. A large number of pores overlap each other, which ImageJ

interprets as a single large pore. In areas of high pore density multiple pores can overlap to

form a large area. The largest area taken by a single particle during analysis was 4.919μm,

requiring at least 10 overlapping pores at minimal overlap (800nm diameter pores have a

0.5 μm2 area). As no pores have an area a drastically reduced area less than that of a single

pore this pore overlap considerably increases both the mean pore diameter value and the

standard deviation of this value.

57

To accurately measure the size of pores formed from individual etched ion tracks (i.e. pores

where no overlap occurs) a filter is applied to the image. ImageJ was set up to only select

pores with a circularity of at least 0.95. This selects all individual pores and the result is

shown in Figure 32. The image still however includes some overlapping pores. This occurs

when the overlap is almost perfect, or when a number of almost perfect overlaps create a

pore without affecting its circularity. The mean pore diameter from this image is 871nm

with a standard deviation of 51nm. This mean is much closer to the specified value of

800nm, and the standard deviation shows a much lower spread of pore diameters.

Interestingly the specified value is not within the standard deviation of the calculated mean,

unlike in the initial calculation. A possible reason for the larger than specified mean pore

diameter is discussed in the later in this section.

Figure 32: Image of 800nm pores after removal of particles with circularity <0.95. Note small number

of overlapping pores remain.

58

The number of pores in the original sample image used for pore size analysis (Figure 30) was

528. The number of pores with little to no overlap (Figure 32) was 256. The result of this is

that only 48% of nanowires produced will have the desired cylindrical shape and required

diameter. This is extremely relevant when designing an experiment involving mechanical

testing of the nanowires where wire diameter is a critical value.

The porosity of the template was calculated using the image with all pores displayed,

including partial pores at the edge of the image (Figure 29), however dark areas of template

were not included (a filtering process similar to that of Figure 30 was used but with a smaller

limit imposed). A porosity of 20.8% was recorded. This is at the upper limit of the porosity as

quoted in the specifications (see Figure 33). This suggests that either: a) this particular area

of the template has a lower than average number of overlapping pores, as the area of two

individual pores for example would be greater than if there was any overlap between the

two; b) there are a number of pores with an elliptical appearance, and hence a larger area,

due to a non-normal pore angle; c) this particular area of template received a higher than

average radiation dose during manufacture; or else d) this area of template is fairly average

and the lower porosity quoted in the specifications is merely to ensure any filters supplied

with a lower than average radiation dose remain within the parameters.

800nm IsoporeTM Membrane Filters

Chemistry Polycarbonate

Pore size 0.8 μm

Filter diameter 25 mm

Thickness 7 - 22 μm

Porosity 5 – 20 %

Figure 33: Summary of specifications of 800nm pore diameter filters [64].

59

To further investigate if the sample area above is similar to that of other areas of the filter, a

larger number of images were analysed. The average pore diameter of all pores was found

to be 970nm with a standard deviation of 245nm. This is slightly lower than that of the area

analysed above (1034nm ± 281nm), however it is within the standard deviation of the

original value showing that the original area was not too far outlying. The value for pore

diameter when only counting individual (i.e. not overlapping) pores was found to be 839nm

with a standard deviation of 67nm. Again this is slightly, but not significantly, lower than

that of the original area (871nm ± 51nm). However it does move the specified value of

800nm to within the standard deviation of the mean value. This result also shows that there

are differences, albeit small, in pore diameter in different areas of the template.

The mean pore size values clearly show that whilst the quoted value of 800nm is accurate

for individual pores, the chance for pore overlap is high; over 50% of pores are formed of

multiple overlapping pores. This results in the average pore diameter being much higher

than that quoted, as well as a large standard deviation of pore sizes. Analysis was carried

out on a variety of other filters (100nm, 200nm and 400nm) to ascertain whether this was

true for all filters (a summary of all pore analysis results is shown in Figure 41).

6.1.2. 400nm template analysis

The 400nm pore size filter had an overall average pore diameter of 436nm ± 97nm. Once

only individual (not overlapping) pores were selected the average pore size was reduced to

394nm ± 45nm. The porosity of the filter was 12.9%, almost exactly between the quoted

value of 5 – 20%.

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6.1.3. 220nm template analysis

The 220nm filter had an overall average pore diameter of 343nm ± 138nm. Once only

individual (not overlapping) pores were selected the average pore size was reduced to

259nm ± 20nm. The porosity was measured to be 26% (13.8% quoted). The porosity value

for this filter is considerably higher than that quoted. This is at least partially due to the SEM

image analysed (see Figure 34). The image has a notable brightness gradient across the

vertical axis, from darker at the top to lighter at the bottom. As well as this there are areas

of filter that are as dark as the pores. Due to these two effects it was difficult selecting only

pores when applying the threshold converting the image to binary (Figure 35). To get an

acceptable image to analyse, a convolve filter was applied to a copy of the image to create a

blurred version. This image was used as a background to subtract from the original image

eliminating most of the brightness gradient. After this additional processing a threshold

image of all pores was obtained. However, this image still included some unavoidable areas

where part of the filter was selected as pores. This additional area slightly increases the

measured porosity of the filter as this extra area is included.

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Figure 34: SEM image of 220nm filter. Not brightness gradient from the top of the image (darker) to the bottom of the image (lighter).

Figure 35: SEM image of 220nm pores when attempting to apply a binary threshold, highlighting issues due to a vertical brightness gradient.

62

Figure 36: Threshold image after selecting all pores in 220nm template. Note some areas of filter remain selected, increasing average pore diameter and porosity.

6.1.4. 100nm template analysis

The 100nm pore size filter had an average pore diameter for all pores of 163nm ± 34nm.

Individual (not overlapping) pores had an average pore size of 151nm ± 14nm. The porosity

of the filter was 11.2%, more than double than the quoted value of 4.18%. The image used

had none of the difficulties as described for the 220nm filter and this value is not artificially

increased.

6.1.5. Pore angle

A number of pores in Figure 28 show a brightness gradient across the diameter of the pore,

as well as appearing to have an elliptical cross section. This is caused by the radiation

incident angle, and therefore pore angle, being not normal to the filter surface. The effect is

shown more clearly in Figure 37. The non-perpendicular pore angle can occur in both single

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pores as well as overlapping pores. Time restraints have limited further investigation (e.g.

side on images) into specific values of the pore angles (as discussed in future studies). This

effect also influences pore diameter and porosity measurements, as it results in an elliptical

pore. However, due to the selection process when measuring non overlapping pores (as

discussed above) being a circularity limit this has a minimal effect on average pore diameter

recorded for single pores.

Figure 37: SEM image showing non-perpendicular pore angle in 800nm polycarbonate template.

6.1.6. Polycarbonate damage susceptibility

It is worth discussing the damage susceptibility of the templates, as handling is required for

both preparing for, as well as during, electrodeposition. After the gold/palladium coating

the filters are cut into quarters (as discussed in experimental methods above). Figure 38

shows the damage caused by a cut. The damage is limited to an area of less than 20μm from

the cut. This shows that the template is suitable to be split into separate samples as the

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Lacomit lacquer securing the samples to their copper mount cover more than this affected

area.

Figure 38: SEM image of damage during cutting of polycarbonate templates. Damage penetrates up to 20μm from the edge.

The polycarbonate templates remain fragile during handling. Figure 39 shows the damage

caused by handling with a crocodile clip. This damage occurred during an early

electrodeposition run before samples were mounted to a copper sheet, and the damage

influenced the decision to mount the samples.

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Figure 39: SEM image of damage caused by a crocodile clip during handling.

It is also possible to damage the polycarbonate templates during SEM imaging. Figure 40

shows a dark area in the centre of the image. This is known as specimen contamination and

is caused by residual gas being struck by the electron probe. The dark area shows where an

increased magnification view was used for focussing. Upon reducing the magnification the

area around the previous selection showed significant damage. This is due to stresses

caused by the incoming electrons of the SEM. It is therefore not possible to put imaged

templates to further use.

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Figure 40: SEM image of damage to a 100nm polycarbonate template due to electron imaging. The dark area at the centre shows where an increased magnification was used for focussing purposes,

with damage around this area.

6.1.7. Summary of etched ion track polycarbonate template analysis

A summary of polycarbonate template properties is shown in Figure 41. Whilst the average

pore diameter for individual pores is close to the specified value, the overall average pore

diameter for each template size when including all pores is considerably greater than that

specified. The standard deviation for each samples is at least 20% that of the mean (100nm

template) and as much as 40% of the mean (220nm template). This large variance in pore

diameter, as well as the fact that the average pore diameter is much greater than that

specified, results in templates unsuitable for the growth of nanowires for many applications

(as discussed later).

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Specified pore

diameter

Average pore

diameter (only

individual pores) (nm)

Average pore

diameter (all pores)

(nm)

Specified porosity Measured

porosity

800nm 839 ± 67 970 ± 245 5 – 20% 20.8%

400nm 394 ± 45 436 ± 97 5 – 20% 12.9%

220nm 259 ± 20 343 ± 138 13.8% 26%*

100nm 151 ± 14 163 ± 34 4.18% 11.2%

Figure 41: Table showing summary of pore analysis measurements. *porosity value for 220nm

template artificially inflated (see above).

6.2. Electrodeposition into polycarbonate membrane

Electrodeposition into the etched ion-track polycarbonate templates was carried out of a

range of potential values and time scales. All electrodeposition was carried out using a

250g/L copper sulphate pentahydrate (CuSO4.5H2O) solution, with a copper anode, a

saturated calomel reference electrode and the sample as the cathode. Only the 800nm pore

size template was used due to time restrictions. Significant results are discussed below.

6.2.1. Sample 1

Figure 42 shows a sample after electrodeposition for 300 seconds at an offset potential of -

1500mV. The majority of pores are unfilled, and the few that are show growth only on the

side walls of the pores. It was thought that this was due to an insufficient thickness of the

gold/palladium coating. If this coating was too thin, the pores would not be completely

covered. The copper would therefore be begin to be deposited at the edge of the pore

bottom and proceed to grow vertically through the pore, resulting in the partially filled

pores observed. A magnified view of this effect is shown in Figure 44 (taken from a separate

sample). Energy-dispersive X-ray spectroscopy was used to analyse the elemental

composition of the sample surface (Figure 43). The spectrum showed low copper peaks,

proving that copper had been deposited, if only in limited amounts.

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Figure 42: SEM image of 800nm template after copper electrodeposition at an offset potential of -1500mV for 300 seconds. The majority of pores are unfilled.

Figure 43: Energy-dispersive X-ray spectroscopy spectrum for copper electrodeposited in

polycarbonate template (Figure 42). Small copper peaks are shown, due to limited amount of copper deposited.

69

Figure 44: SEM image partially filled pores in an 800nm template after copper electrodeposition at an offset potential of -1500mV for 300 seconds.

6.2.2. Sample 2

Following the results above, a thicker conductive gold/palladium layer was applied to

subsequent samples. Electrodeposition was repeated using the same parameters as above

(potential difference of -1500mV for 300s). The result of this is shown in Figure 45. This

image shows a range of different results. Some pores are partially filled like those discussed

above in the previous sample. Conversely, there are large deposits formed on the surface of

the template. These were confirmed to be copper by Energy-dispersive X-ray spectroscopy

and a reduced brightness view is shown in Figure 46. Caps are formed once a pore has been

completely filled with copper (as discussed above in the literature review, and by Molares et

al. [43,44]). It is energetically favourable to continue growth at a current nucleation point

than to nucleate new growth areas. Therefore, the earliest pores to begin filling will proceed

at a quicker rate than others. Once the pore is full copper growth is free of the size

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restrictions of the pore and so a large cap forms on the surface. The increased surface area

increases the overall current flow of the electrodeposition. This is in part the cause of the

erratic current-time plot for the deposition run (Figure 47). Large current peaks are

observed, possibly due to the increased current flow. A more likely reason for the current

peaks is the formation of copper tendrils at the edges of the copper mounting sheet. This

occurred at areas of the copper sheet where the Lacomit lacquer did not completely cover

the mount, and electrodeposition was therefore able to take place. As the copper sheet was

thin (<1mm) the growths proceeded outwards into the electrolyte. This has a similar effect

on the current flow due to an increase in conductive area as the caps discussed above.

However, unlike the caps these tendrils are not supported by the polycarbonate template

and easily break off due to the slight movement of the magnetic stirrer in the electrolyte.

The tendrils then begin to grow again and this is likely to cause the peaks in the current flow

observed in Figure 47.

Figure 45: SEM image of electrodeposition into an 800nm template at an offset potential of -1500mV for 300 seconds, including partially filled pores, completely filled pores and copper caps.

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Figure 46: SEM image of copper caps formed on a polycarbonate template.

Figure 47: Current-Time plot of copper electrodeposition into 800nm polycarbonate template at

-1500mV. Note: Current density values are for indication only due to the lack of accurately measuring electrodeposition area (especially during cap or tendril growth, See section 6.2.3).

72

6.2.3. A note on current density-time plots

The potentiostatic set up used recorded of current density with time rather than current. In

order to calculate a current density the software required a surface area value be inputted.

As the samples used were not a regular size, were secured by an area of lacquer, and the

templates had a porosity value, this deposition area was estimated. The area was

approximated to a square with sides of 0.5cm (the software was calibrated in centimeters),

and the porosity was assumed to be 20%. This resulted in a value of 0.05cm2. However, as

caps or tendrils grow on the sample the area increases, which increases current flow. As this

change in area is not taken into account by the software, and the increase in current is

displayed as an increase in current density. It is therefore advisable to interpret the current-

time plots as a representation of current flow, and not to use actual current density values.

6.2.4. Sample 3

The large caps grown on the surface of the previous sample suggest that the

electrodeposition time was too long. Ideally the electrodeposition process would end before

the caps start to form. Figure 48 shows the results of electrodeposition into a template for

60 seconds at the same potential as used previously (-1500mV). The reduction in

electrodeposition duration drastically reduces the number of caps grown, with the majority

of the template surface being cap free. The image shows a mixture of unfilled, partially filled

and full pores. A possible reason for the uneven pore filling rate in the sample is undesirable

reactions taking place in some of the pores. If the current supplied is too high it is possible

to deplete the solution inside the pores. This will cause hydrogen from the solution to

reduce in place of the copper. The hydrogen can build up and form bubbles, causing

reduced or no copper deposition to take place in that pore. Mechanical stirring of the

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solution can reduce the hydrogen buildup, however there is minimal flow at the bottom of

the pores. The pores where hydrogen reduction occurs will have a lower chance of copper

nucleation, whilst other pores may be fully filled.

Figure 48: SEM image of electrodeposition into an 800nm template at an offset potential of -1500mV for 60 seconds, showing full and partially filled pores.

6.2.5. Sample 4

To increase the number of filled pores the applied potential was reduced, thereby reducing

(and ideally eliminating) hydrogen evolution. Figure 49 shows a much increased ratio of

filled pores to empty pores (approximately 90% filled). This was achieved using an offset

potential of -1000mV for 60 seconds. Despite the reduced potential and hence lower

electrodeposition rate, the time period was not increased. This was to ensure that no caps

were grown on the surface of the template. The pores were shown to be comprised of

copper by energy-dispersive X-ray spectroscopy (Figure 50). The spectrum shows large

peaks corresponding to the Lα and Kα values for copper. Large sulphur peaks are also

present, due to the copper sulphate pentahydrate electrolyte. The brightness of each pore

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corresponds to the height of the nanowire grown within the pore. A large amount of pores

are the same height, approximately that of the thickness of the template. Time restraints

however meant no profile image of the wires could be taken to fully study wire height

Figure 49: SEM image of electrodeposition into an 800nm template at an offset potential of -1000mV for 60 seconds (for clarity: scale bar reads 20μm and magnification reads 1501x).

Figure 50: Energy-dispersive X-ray spectroscopy spectrum of polycarbonate template sample 4 (Figure 49). Large copper peaks are present confirming copper deposition.

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The current-time plot for this sample has none of the peaks associated with large increase

then decrease in electrodeposition area. This corresponds with the fact that tendrils were

not observed to grow on the sample. The increase in current throughout the deposition

suggests an increase in deposition area. This is supported by theory, which states that

nucleation is more likely to begin at the interface between the pores and the gold/palladium

coating at the bottom of the pores. The reduced surface area required here reduces the

energy cost of nucleation, causing preferential nucleation to occur. The growth of copper

from these points along the sides of the pore increases the electrodeposition area. Thus

complete pores are likely to be formed when multiple edge growths combine. This results in

polycrystalline wires, a result confirmed by comparing the caps formed with those produced

by Molares et al. shown in Figure 11, which are of a circular nature rather than the multi-

faceted caps associated with single-crystal wires [43]. Further crystallographic analysis could

be used to confirm the structure of the wires, but was not performed due to time restraints.

Figure 51: Current-Time plot of copper electrodeposition into 800nm polycarbonate template at

-1000mV. Note: Current density values are for indication only due to the lack of accuracy in measuring electrodeposition area.

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6.2.6. Summary of electrodeposition into polycarbonate templates

Electrodeposition into polycarbonate templates was performed at potentials of between -

1500mV and -1000mV for time periods ranging from 15s to 300s. The best results were

found when performed at -1000mV for 60 seconds. The wires produced filled the majority

of the pores (approx. 90%) and were of a fairly uniform height, as demonstrated by SEM

images. Increasing the current (by increasing the potential) results in a lower pore filling

ratio, most likely due to hydrogen reduction in some pores. Grown nanowires were shown

to be polycrystalline in structure, as shown by the circular caps formed on the surface of the

template under longer electrodeposition.

6.3. Aluminium Oxide template growth

A two-stage anodizing process by Masuda et al. and developed by Dou and Derby was used

to produce ordered anodised aluminium oxide (AAO) templates [10,11]. An analysis of the

templates created is presented following.

6.3.1. First stage anodisation

After electropolishing, the aluminium substrate is much less rough than that of the original

aluminium, however there remain features such as ridges on the surface (time restraints

prevented analytical methods such as AFM (atomic force microscopy) being used to

determine surface morphology). After a long-period anodizing (12h) in 0.3M oxalic acid at

40V an alumina oxide film is formed on the surface. Figure 52 shows the surface of the

sample after this first anodizing stage. The pores formed are not in any regular order,

despite the self-ordering nature of porous alumina due to space restraints [50]. The average

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pore size is 1640nm ± 420nm. This is larger than suggested by literature, and the

distribution in size is fairly large [10]. The reason for the lack of regularity and large pore size

distribution is the large thickness of the oxide layer causing various distortions [9].

Figure 52: SEM image of alumina film after 12h anodising in 0.3M oxalic acid at 40V.

6.3.2. Ridges formed on aluminium

The thick, distorted layer of oxide was removed by etching in a 0.1M Chromic acid and 0.1M

phosphoric acid solution. The anodizing process causes ridges to form on the aluminium

substrate, as shown in Figure 53. The ridges form cells and are highly ordered. ImageJ was

used to determine the pore size by a similar process to that described for the polycarbonate

templates. The pores were found to be 61.5nm ± 3.1nm, showing a low degree of size

distribution. The cell interval (pore center to pore center) distance was found to be 112.5nm

± 11.5nm. This result agrees with analysis performed on images taken by Dou, of 64.2nm ±

8.0nm and 105.4nm ± 8.6nm respectively [36]. Time restraints prevented the depth of the

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cells to be determined through a technique such as AFM (cells were found to be

approximately 5nm in depth by Dou [36]). These ridges form nucleation points during the

next anodizing stage.

Figure 53: SEM image of aluminium substrate after removal of oxide layer, showing ridges formed during anodisation.

6.3.3. Second stage anodisation

A second anodising process was carried out for 1 hour in conditions identical to that of the

first stage. Figure 54 shows the porous aluminium oxide film formed. The pores show a high

degree of regularity, and closely resemble the ridges formed during the first anodising stage

in terms of approximate size and regularity. Pore size analysis using ImageJ (Figure 55 and

Figure 56) shows an average pore diameter of 61.5nm ± 3.0nm. This agrees with the result

of Masuda and Fukuda in the paper first reporting the 2-stage method of 67nm ± 6nm [10].

The inter pore distance (cell size) was found to be 108.6nm ± 10.5nm. This matches the cell

size reported by a number of sources of 99nm ± 8nm, 112nm, 95nm ± 5nm and 100nm for

79

identical anodizing conditions [10][65][38][53]. The cell size to voltage proportionality

coefficient of 2.7nm/V also corresponds with that reported by Li et al. of 2.8nm/V [12].

These values confirm the oxide layer’s resemblance to the ridges formed in the previous

stage, with both parameters being similar (pore size of 61.5nm ± 3.0nm compared to pore

size of 61.5nm ± 3.1nm for the ridges, and cell size of 108.6nm ± 10.5nm compared to

112.5nm ± 11.5nm of the ridges). This provides evidence for the ridges being nucleation

points for the oxide formation on the surface of the aluminium, and also suggests that the

pores formed are straight and parallel. The appearance of the inside of the pores being fully

dark with no shadow supports this, as the inside of angled pores would be visible. Time

restraints prevented profile images of the templates to be taken so this could not be

experimentally verified. Profile images would also be used to determine the thickness of the

oxide layer and the depth of the pores. The porous aluminium oxide has been reported to

have a growth rate of 220nm/min, which suggests a thickness in this case of 13.2μm [36].

80

Figure 54: SEM image of porous aluminium oxide formed after second anodising process showing similar appearance to the ridges formed during the first anodising stage (Figure 53).

Figure 55: SEM image of porous aluminium oxide formed after second anodising process used for pore analysis (higher magnification results in an increase in accuracy).

Figure 56: Image of pores after ImageJ threshold and particle analysis applied to Figure 55.

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6.4. Electrodeposition into AAO templates

Prior to electrodeposition, the templates produced required further processing stages.

These are discussed in the experimental methods, but in summary are: coating with PMMA,

etching of aluminium, etching of oxide barrier layer, porous alumina coated with

gold/palladium, removal of the PMMA. Electrodeposition took place in a 250g/L copper

sulphate pentahydrate (CuSO4.5H2O) solution identical to that used for the polycarbonate

templates. A copper anode and saturated calomel reference electrode were used, with the

gold/palladium of the sample used as the cathode. As with the polycarbonate templates,

significant results are discussed below.

6.4.1. Sample 1

Following the positive results reported above for the polycarbonate template sample 4,

similar electrodeposition parameters were used to deposit into the porous alumina

template. A potential difference of -1000mV was applied. The time period was reduced to

30 seconds to prevent caps forming due to the template being thought to be thinner than

the polycarbonate (approx. 13.2μm compared to 20μm). Figure 57 shows the surface of the

template after electrodeposition. Large angular growths up to 20μm in length are evident

on the template surface. The angular nature of the growths suggests that single crystal

nanowires were grown beneath them (see Figure 11). The pores here are much narrower

than those in the polycarbonate template and so multiple nucleation points in each pore are

less likely, supporting the suggestion of single crystal growth. The growths were confirmed

to be copper using Energy-dispersive X-ray spectroscopy (Figure 58). The spectrum shows a

low copper peak however, indicating a low amount of copper deposited. Peaks representing

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gold and palladium are present due to a thin (~3nm) Au/Pd coating applied prior to SEM

imaging to prevent electron charging on the surface of the alumina.

Figure 57: SEM image of the surface of porous alumina template sample 1, showing angular copper growths.

Figure 58: Energy-dispersive X-ray spectroscopy spectrum of aluminium oxide sample 1 (Figure 57).

Small copper peaks indicate low amount of copper deposited.

83

An increased magnification image shows that in areas without the large growths the

majority of pores are empty (Figure 59). There are a number of possible reasons for this: a)

The electrodeposition setup is not conducive for copper deposition; b) the electrical contact

between the copper mount sheet and the gold/palladium coating of the sample was not

adequate; c) the oxide barrier layer was not fully removed during the production process; or

d) the PMMA coating on the pores was not fully removed. Whilst this experiment yielded no

nanowires, it confirmed that the alumina templates survived the processing stages between

their initial growth and the electrodeposition. The pore size has also remained unchanged at

60.8nm ± 3.0nm (61.5nm ± 3.0nm reported above for pore size before further processing).

This unchanging of pore size does suggest that the barrier layer was not fully removed, as

once that layer is removed the solution will begin to etch away at the insides of the pores

increasing their diameter. However this is not proof in itself as the etching takes place at the

bottom of the pores and the SEM image shows the pore tops, and pore widening may still

have occurred.

Figure 59: SEM image of the surface of porous alumina template sample 1 under increased magnification, showing predominantly empty pores.

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6.4.2. Sample 2

During electrodeposition of the previous sample, the electrodeposition parameters and set-

up were identical to that of the successful polycarbonate template deposition and so the

incorrect electrodeposition conditions suggestion can be discounted. The electrical contact

from the mount to the sample is sufficient enough to support the large copper growths and

so was not thought to be the cause of low yield of nanowires. Finally there was no evidence

of any remaining PMMA on the surface of the sample from the SEM images. During the

production of the subsequent templates a longer period phosphoric acid etch was used to

ensure the entirety of the barrier layer was removed. Figure 60 shows the result of

electrodeposition into a porous alumina oxide template using the same electrodeposition

parameters: -1000mV potential for 30 seconds. Here, large bulbous copper growths are

observed, in contrast to the angular growths observed in the previous sample. Referring to

Figure 11, these growths suggest a polycrystalline structure. The growths were again

confirmed to be copper by energy-dispersive X-ray spectroscopy (Figure 61). An explanation

for these growths could be that a number of nanowires have grown in close proximity to

one another, caps formed on the alumina surface and new copper growths have nucleated

on these caps. Hence it is difficult to ascertain whether the wires themselves are of single or

polycrystalline structure, and without removing the caps and performing further

crystallographic analysis it is not possible to tell. The current-time plot for this run shows a

continuous increase in current, reflecting the increase in electrodeposition area. Areas

without the large growths show an almost zero pore filling rate as in the previous sample

(Figure 63).

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Figure 60: SEM image of the surface of porous alumina template sample 2, showing large circular copper growths (scale bar 20μm).

Figure 61: Energy-dispersive X-ray spectroscopy spectrum of growths on aluminium oxide sample 2 (Figure 60).

86

Figure 62: Current-time plot for alumina oxide template sample 2 showing continuous increase in

current flow.

Figure 63: SEM image of the surface of porous alumina template sample 2 at higher magnification, showing empty pores.

87

6.4.3. Sample 3

The previous sample was etched for a longer time period, ensuring that the alumina oxide

barrier layer was completely removed. Whilst the sample showed large amounts of copper

deposited, no pores were observed (with the exception of those presumably under the

growths). The energy-dispersive X-ray spectroscopy spectrum for the previous sample

showed a slight carbon peak. This peak is clearer when viewing the spectrum for an area of

the template with fewer copper growths (Figure 64). This may indicate that some PMMA

may remain in the pores, preventing electrodeposition in those pores. In order to

investigate this, a new technique for removing the PMMA was used. Previously the PMMA

was removed by submersing the sample in a beaker of acetone. It is possible therefore

possible for the region near the sample surface to become fully saturated with PMMA whilst

there is still some PMMA in the pores. A new removal method was developed utilizing a

magnetic stirrer and refreshing the acetone during the process.

Figure 64: Energy-dispersive X-ray spectroscopy spectrum of an area with little copper growths on aluminium oxide sample 2, showing carbon peak.

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Electrodeposition was performed on the sample under similar conditions to the previous

samples (potential of -1000mV for 30s). SEM images of the sample surface following

deposition showed regions of empty pores along with regions with a high pore filling ratio

(Figure 65). The full pores were shown to contain copper by Energy-dispersive X-ray

spectroscopy (Figure 66). Interestingly the peaks for carbon are still present. The average

wire diameter of wires grown was calculated to be 58.0nm ± 5.4nm, comparable to that of

the pore diameter. It is unclear why certain areas were more successful than others in

electrodeposition and further investigation is required.

Figure 65: SEM image of the surface of porous alumina template sample 3, showing region of high pore filling surround by a region of empty pores.

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Figure 66: Energy-dispersive X-ray spectroscopy spectrum of an area of full pores, showing copper peaks.

6.4.4. Summary of electrodeposition into porous aluminium oxide templates

Electrodeposition was carried out at a potential of -1000mV for 30 seconds in a 250g/L

copper sulphate pentahydrate (CuSO4.5H2O) solution. Despite this set up being successful in

growing nanowires in etched ion track polycarbonate templates it was less efficacious when

used with porous alumina templates. A series of production stages were changed to

investigate possible reasons for the lack of deposition, including increased etching of the

alumina barrier layer and improved removal of the PMMA coating. Improving these two

methods results in a template with areas of high pore filling ratio, surrounded by areas with

empty pores. It is clear that more research into the preparation of alumina templates for

electrodeposition is required. Despite the lack of large areas of nanowires grown, the SEM

images show a positive result in that the porous templates formed after the second

anodising stage successfully survived the further preparation stages and the pore diameter

remains unchanged through these processes.

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6.5. Etched ion track polycarbonate templates vs. AAO templates

A comparison of both types of template and of the nanowires produced in the templates

follows. The requirement of the wires to be mechanically tested by uniaxial compression is

considered.

Etched ion-track polycarbonate templates have a number of disadvantages due to the way

they are produced. Pores are arranged randomly, with areas of high pore density as well as

areas devoid of pores. This makes compression testing of nanowires produced difficult

whilst the wires remain attached to the substrate. Compression of individual wires is

dependent on being able to identify and locate a wire, and on the removal of all surrounding

wires. Meanwhile the simultaneous testing of multiple wires may result in buckling due to

the lack of stabilising effect from a large density of wires. Nanowires grown in porous

alumina templates are of a highly ordered nature and as such are suitable for mechanical

testing by simultaneous compression as shown by Dou and Derby [11]. The removal of wires

from the substrate for individual testing is possible in both cases; however it is difficult to

perform uniaxial compression or tension tests with free wires.

The diameter of produced nanowires is obviously highly critical during testing of how the

mechanical properties vary with diameter. Thus it is important to have a large degree of

consistency in this parameter. The pores in the polycarbonate templates have a large

standard deviation, as much as 40% the mean value. Due to the random pore arrangement

many pores (over 50%) are in fact multiple overlapping pores, which increases the pore

diameter as well as giving the pore a non-regular shape. The shape can introduce stress

concentration areas in the nanowire which could change its mechanical properties. Contrary

to this, porous alumina templates have a low variation in pore size, with the standard

91

deviation being less than 5% of the mean. This makes the nanowires grown in the

polycarbonate templates unsuitable for mechanical testing, whereas nanowires grown in

alumina templates are valid.

Investigation of how mechanical properties vary with diameter requires a range of

diameters be available. Both templates presented are suitable for an etching process

widening the diameter, using phosphoric acid for the alumina templates and a sodium

hydroxide aqueous solution for polycarbonate [11,66]. Moreover multiple diameters of

polycarbonate filters are available without the requirement for further processing.

For accurate compression testing of the produced nanowires, the nanowires must be

parallel to each other and at a normal angle to the compression axis. A substantial number

of pores (approximately 25% in Figure 37) in the polycarbonate templates were observed in

SEM images to have a non-perpendicular pore angle. This would cause the produced

nanowires to have the same angle, affecting mechanical testing results. Pore angles for the

alumina templates appear in SEM images to be normal to the surface, a result supported by

the matching pore diameters of the ridges in the aluminium and the final measured pore

diameters, as well as the high pore density (which create geometric constraints on the pore

angle). Literature also supports this assumption [65]. However, further investigation is

required to confirm the pore angles for each template.

If the aspect ratio of the wires requires increasing, a thicker template is needed. Shorter

wires can be grown in larger templates by reducing the electrodeposition time. Alumina

templates can be as thick as required, with templates up to 160μm produced [65]. There

have been no investigations into the upper limit of this growth whilst keeping the pores

produced straight and parallel however. Contrary to this, the polycarbonate templates are

92

only available in a specific thicknesses (up to 30μm, depending on pore diameter) and so

this is a limiting factor in the length of nanowires grown [64].

Results from the electrodeposition experiments indicate that whilst porous alumina oxide

templates can be consistently produced through the 2-stage anodising process, there is

further research required into preparing the templates for electrodeposition. This is shown

by the unreliable electrodeposition of nanowires. The large number of processing stages

mean that it is difficult to ascertain which stage the complications arise. The polycarbonate

templates have a much simpler processing route, simply requiring a conductive coating, and

this is reflected in the more consistent electrodeposition results (i.e. pores were filled to

some degree in each experiment).

7. Conclusion

A literature review was performed concluding that a promising method of producing

nanowires without the taper and ion damage associated with focussed ion beam (FIB)

milling, and without the size restraints of an etching of a directionally solidified eutectic, is

electrodeposition into a template. Research into the production of highly ordered porous

anodised aluminium oxide templates was carried out. Finally research into mechanical

testing methods took place. Uniaxial testing was shown to important in testing nanowires

mechanical properties in the absence of strain gradients.

Characterisation of polycarbonate templates was carried out. Templates with a specified

800nm pore diameter were found to have an average pore diameter of 970nm with a

standard deviation of 245nm. This is considerably higher than the specified diameter and

the variation in pore diameter was large (25% the mean). Similar results were found for

93

400nm, 220nm and 100nm templates, with average pore sizes being much higher than the

specified mean with large standard deviations of up to 40% the mean (for 220nm

templates). The reason for the large averages and variation in pore diameters is the large

number of overlapping pores (over 50% of all pores). This also creates pores with a non-

circular cross section, possibly affecting mechanical properties. This result, combined with

the evidence of pores at a non-normal angle to the template surface show that the

polycarbonate templates are not suitable for uniaxial compression testing of nanowires.

Porous anodised aluminium oxide templates were produced through a two-step anodising

process. Pore analysis was performed showing pores to be 61.5nm ± 3.0nm. This matches

templates grown using the same anodising conditions reported in literature. There was a

much lower variation in pore diameters compared to the polycarbonate templates. SEM

images suggested that the pores were straight, parallel with no non-normal pore angles.

Literature further supports this; however this was not experimentally verified duo to time

constraints. The pores are also regularly ordered over a long scale, and as such the

nanowires grown using these templates would be more conducive for mechanical testing via

simultaneous uniaxial compression.

Potentiostatic electrodeposition of copper into the polycarbonate templates was carried out

in a 250g/L copper sulphate pentahydrate (CuSO4.5H2O) solution, with a copper anode, a

saturated calomel reference electrode and using the samples as a cathode. A pore filling

rate of approximately 90% was achieved when applying a potential difference of -1000mV

for 60 seconds. The nanowires grown were shown to be copper using energy-dispersive X-

ray spectroscopy. Using the same parameters when depositing into porous alumina

templates produced less satisfactory results. This was thought to be due to complications

94

arising during one of the production stages between initially producing the porous template

in during the second anodising stage and the electrodeposition. The porous structure was

consistently shown to be intact after electrodeposition with no change in pore size.

Nanowires grown by electrodeposition into etched ion-track polycarbonate templates are

not suitable for the investigation of mechanical properties. This is due to their inconsistent

pore diameter, unordered pore arrangement and varying pore angle. However,

electrodeposition into porous anodised aluminium oxide templates is a promising method

of growing nanowires for this purpose. Pores have a low distribution in diameter, are

regularly ordered and the pores are normal to the template surface. This makes mechanical

testing of the grown nanowires properties possible via simultaneous uniaxial compression

testing.

8. Future works

A key future investigation is into improving the pore filling ratio during electrodeposition.

The preparation stages of the aluminium oxide templates need refining, for example

ensuring all the barrier layer oxide and PMMA layer are removed. Improvements to the

electrodeposition experimental set up could also be investigated. A conical anode could be

used, as a flat anode caused preferential deposition in border areas of the cathode [44].

Additions in the electrolyte could also be considered. Multiple papers have reported using a

small (21g/L – 32g/L) sulfuric acid additions, which are reported to increase the conductivity

of the solution and lower the cathode overvoltage [44,45]. Changes to the electrodeposition

parameters such as potential could also be investigated. During deposition into the

polycarbonate templates, the potential was reduced to avoid hydrogen reduction at the

95

cathode. This could still be occurring at the lower potentials used. Copper has been

reported to have been deposited at potentials as low as -20mV, however time periods

should be increased to compensate for the lower deposition rate [44].

During this work a DC deposition methodology has been used. This was the simplest method

of deposition into the polycarbonate templates, and was used for the alumina templates to

enable comparison between the two. However, the DC method requires many additional

processing steps, including the removal of the aluminium substrate and removal of the

alumina barrier layer. An AC deposition method after the thinning of the barrier layer, such

as that used by Gelves et al. could eliminate these additional stages (see Figure 9) [37].

Gelves performed the anodizing in sulfuric acid however so further investigations into this

process using oxalic acid is required.

The pore angles were not fully characterized during this project. The angles of the pores

were investigated by top down SEM images, however further investigation via profile

images is necessary to fully describe the templates.

The investigation of the relationship between nanowire size and its yield strength requires

the production of nanowires with various diameters. This can be achieved using the alumina

templates through etching of the pores. Investigations into the etching rate of different

chemicals would be useful.

The objective of this project was to compare different methods of producing nanowires for

mechanical testing. An important future study would be these mechanical tests. Nanowires

grown in the porous alumina templates are suitable for simultaneous uniaxial compression

using a flat punch tip fitted nanoindentor. Copper nanowires grown by electrodeposition

96

have not been mechanically tested and so it is important to see if the scaling law suggested

by Dou and Derby applies [7]. Furthermore, the templates are suitable for the growth of a

large range of different metals, as long as the alumina is not reactive to the

electrodeposition solution. Therefore a large number of different metals could be tested.

The micromechanics of strengthening and deformation mechanisms of the nanowires could

be investigated by transmission electron microscopy (TEM).

An additional study could look at the effects of radiation on nanowires. Recent theoretical

work involving nanoporous metals, which have a very similar diameter length scale to metal

nanowires, has suggested that such metallic nanostructures should be highly resistant to

radiation damage [67]. There is a need for such radiation resistant metals for future

applications in nuclear fission and fusion power plants. Thus underpinning research into

deformation mechanisms and the influence of radiation damage on the mechanical

properties of a variety of nanowires, leading to studies of tungsten nanowires could be

performed.

Finally, an investigation to look at thermal effects on BCC nanowire deformation could be

performed. Unlike FCC structured metals, BCC metals do not have a close packed plane, and

for this reason their slip planes are thermally activated. Higher temperatures activate more

slip systems and screw glide mechanisms. Temperature dependant deformation behaviour

could be tested using an elevated temperature nanoindentor.

97

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