Dealloying and Formation of Nanoporosity in Noble-Metal Alloys · ii Dealloying and Formation of...

Dealloying and Formation of Nanoporosity

in Noble-Metal Alloys

by

Mariusz Albert Bryk

A thesis submitted in conformity with the requirements

for the degree of Master of Applied Science

Graduate Department of Chemical Engineering and Applied Chemistry

University of Toronto

© Copyright by Mariusz Albert Bryk (2012)

ii

Dealloying and Formation of Nanoporosity

in Noble-Metal Alloys

Mariusz Albert Bryk

Master of Applied Sciences

Department of Chemical Engineering and Applied Chemistry

University of Toronto

2012

Abstract

Nanoporosity formation by selective dissolution of Ag-5 at. pct Au in perchloric acid has

been investigated with regards to the mechanism of stress-corrosion cracking (SCC),

film-induced cleavage in particular. It has been proven that dealloying of silver-gold

systems containing low concentration of gold leads to the formation of a three

dimensional nanoporous layer and that it can be carried out in a broad range of potentials

and concentrations of a dealloying solution. Therefore, stress-corrosion cracking

observed in these alloys may be caused, initiated or at least accompanied by the

formation of nanoporosity resulting from dealloying. These results will have impact on

the fabrication of cheaper nanomaterials where there is required large surface to volume

ratio with gold as the outermost layer. Understanding the role of dealloying will also help

us to design new materials of higher resistance against stress-corrosion cracking.

iii

Acknowledgements

I would like to express my sincere gratitude to Professor Roger Newman for the

supervision of the whole project and giving me this incredible opportunity to become a

member of his research group. His suggestions and guidance have been invaluable and

helped me to become a better scientist. I would also like to thank Professor Roger

Newman for accepting me as a PhD student.

Special thanks towards Dorota Artymowicz for the SEM sessions that we were doing

together and Anatolie Carcea for explaining to me how one may quickly organize and

effectively operate with computer softwares designed for studying corrosion. In some

cases doing my experiments was solely possible owing to his personal engagement.

Finally, I would like to thank all my colleagues for their assistance and help that I have

received from them during the project.

iv

Table of Contents

Introduction ...................................................................................................................................... 1

1 Dealloying ................................................................................................................................ 2

2 The dealloying critical potential ............................................................................................... 5

3 Nanoporous metals ................................................................................................................... 9

4 Stress-corrosion cracking (SCC) ............................................................................................ 12

4.1 The surface mobility SCC mechanism ......................................................................... 14

4.2 The film-induced cleavage SCC mechanism ................................................................ 16

4.3 Short summary of the two SCC mechanisms ............................................................... 21

5 Objectives ............................................................................................................................... 22

5.1 Single-shot fracture experiments .................................................................................. 22

5.2 Electrochemistry ........................................................................................................... 23

5.3 Digital simulations ........................................................................................................ 23

6 Experimental .......................................................................................................................... 24

6.1 Materials ....................................................................................................................... 24

6.2 Samples preparation ...................................................................................................... 24

6.3 Electrochemical cell ...................................................................................................... 25

6.4 Mounting samples ......................................................................................................... 26

7 Results and Discussion ........................................................................................................... 28

7.1 Formation of nanoporosity by dealloying of Ag-5 at. pct Au ....................................... 28

7.2 The composition of a nanoporous layer formed on Ag-5 at. pct Au ............................ 29

7.3 The growth of nanoporous layers on Ag-5 at. pct Au ................................................... 30

7.4 The anodic behaviour of Ag-5 at. pct in perchloric acid .............................................. 38

7.5 Fracture experiments..................................................................................................... 41

7.6 Premature cracks and coarsening of a nanoporous layer .............................................. 41

7.7 Fracture tests ................................................................................................................. 43

8 Conclusions ............................................................................................................................ 65

9 Further work ........................................................................................................................... 67

10 References ......................................................................................................................... 68

11 Appendix ........................................................................................................................... 70

v

List of Figures

Figure 1.1. Illustration of the anodic behaviour of binary alloys by Pickering (2). ............ 2

Figure 1.2. A schematic showing a cross-section of ligaments formed on the surface (2). 4

Figure 2.1. Possible misinterpretation of the critical potential by extrapolation. ................ 5

Figure 3.1. Transmission electron micrograph of Ag50 Au50 exposed to 50 % HNO3 (1).10

Figure 3.2. The growth of a gold island by the surface disordering (1). ........................... 10

Figure 3.3. Porosity evolution during selective oxidation proposed by Erlebacher (3).

Reproduced by permission of The Electrochemical Society. ............................................ 11

Figure 4.1. Crack propagation in the presence of an ionic contaminant by Galvele (19). 14

Figure 4.2. An illustration of the distances which atoms have to travel by Galvele (20).. 15

Figure 4.3. An illustration of film-induced cleavage generated during dealloying.

(Adapted from Ref.18.)...................................................................................................... 17

Figure 4.4. A cross-sectional view of a sample that was bent while maintaining the

dealloying potential (24). ................................................................................................... 18

Figure 4.5. Fracture surface after application of strain: (a) brittle and (b) ductile (25)..... 19

Figure 4.6. Crack front striations on the fracture surface of a copper single crystal (26). 20

Figure 4.7. Correlation of acoustic events with current for TG cracks in α-brass (28). .... 21

Figure 6.1. Electrochemical cell used in experiments: WE – working electrode, RE –

reference electrode, CE – counter electrode. ..................................................................... 25

Figure 6.2. Mounting a sample for a single-shot fracture test. .......................................... 26

Figure 6.3. An illustration of the electrochemical cell used in single-shot fracture tests:

WE – working electrode, RE - reference electrode, CE – counter electrode. ................... 27

vi

Figure 7.1. Nanoporous layer obtained by dealloying of Ag95Au5 dealloyed at E = 170

mV in 0.1 M HClO4 at room temperature. Charge passed Q = 3 C/cm2, i = 100 A/cm

2, t

= 35,000 s. ......................................................................................................................... 28

Figure 7.2. Fracture surface of Ag95Au5 dealloyed at E = 300 mV in 0.1 M HClO4 at

room temperature. Charge passed Q = 3 C/cm2, i = 5 mA/cm

2, t = 628 s. Sample fractured

in air. Marker 50 µm. ......................................................................................................... 29

Figure 7.3. EDX performed on the fracture surface of Ag95Au5 dealloyed at E = 300 mV

in 0.1 M HClO4 at room temperature. Charge passed Q = 3 C/cm2, i = 5 mA/cm

2, t = 628

s. Sample fractured in air. Marker 50 µm. ......................................................................... 29

Figure 7.4. Nanoporous layer obtained by dealloying of Ag95Au5 at E = 300 mV in 0.1 M

HClO4 at room temperature. Charge passed Q = 5 C/cm2, i = 6.48 mA/cm

2, t = 772 s. .... 30



2, t = 532 s. .... 31



2, t = 410 s. ...... 31



2, t = 296 s. ...... 32



2, t = 128 s. ...... 32

Figure 7.9. The observed and theoretical thickness of nanoporous layers obtained by

dealloying of Ag95 Au5 in 0.1M HClO4 at E = 300 mV. ................................................... 33

vii

Figure 7.10. Changes in capacitance versus charge passed during dealloying of Ag95 Au5

in ten consecutive steps at different potentials in 0.1 M HClO4 at room temperature (33).

........................................................................................................................................... 34

Figure 7.11. The development of relative area versus charge passed during dealloying of

Ag95 Au5 in ten consecutive steps at different potentials in 0.1 M HClO4 at room

temperature. (33)................................................................................................................ 35


in ten consecutive steps at different potentials in 0.1 M HClO4 ........................................ 35


Ag95 Au5 in ten consecutive steps at different potentials in 0.1 M HClO4 ........................ 36

Figure 7.14. Potentials used for potentiostatic dealloying of Ag95Au5 and Ag77Au23

performed in ten consecutive steps in 0.1 M HClO4 at room temperature. ....................... 37

Figure 7.15. Changes in capacitance versus charge passed during dealloying of Ag95Au5

and Ag77Au23 in ten consecutive steps at different potentials in 0.1 M HClO4. ................ 37


Ag95Au5 and Ag77Au23 in ten consecutive steps at different potentials in 0.1 M HClO4. . 38

Figure 7.17. Current densities observed during dealloying Ag95Au5 and Ag77Au23.......... 38

Figure 7.18. Polarization curves performed on Ag95 Au5 and Ag77 Au23 in 0.1 M HClO4 at

room temperature. Scan rate 0.5 mV/s. IR drop corrected. ............................................... 39

Figure 7.19. Polarization curves performed on Ag95 Au5 and Ag77 Au23 in 1.0 M HClO4 at


Figure 7.20. Potentiodynamic curves performed on Ag95 Au5 in HClO4 of different

aeration degree at room temperature. Scan rate 0.5 mV/s. IR drop corrected. .................. 40

viii

Figure 7.21. Potentiodynamic curves performed on Ag95Au5 in HClO4 of different

concentration at room temperature. Scan rate 0.5 mV/s. IR drop corrected. .................... 40

Figure 7.22. Premature cracks on the surface after dealloying of Ag95Au5 at room

temperature in 0.1 M HClO4 at E = 300 mV vs. MSE. Marker 200 µm. .......................... 41

Figure 7.23.Ag95Au5 surface free from premature cracks after dealloying at room

temperature in 0.1 M HClO4 at E = 170 mV vs. MSE. Marker 20 µm. ............................ 42

Figure 7.24. Nanoporous layer after dealloying of Ag95Au5 in 0.1 M HClO4 for t = 125 s

at E = 170 mV. Charge passed Q = 50 mC/cm2. ............................................................... 42

Figure 7.25. Nanoporous layer after dealloying of Ag95Au5 in 0.1 M HClO4 for t = 2528 s

at E = 170 mV. Charge passed Q = 0.5 C/cm2. ................................................................. 43

Figure 7.26. Crack depth versus thickness of a nanoporous layer for Ag95Au5 dealloyed in

0.1M perchloric acid at 300 mV. Samples fractured beyond the dealloying solution....... 44

Figure 7.27 Fracture surface of Ag95Au5 dealloyed at E = 170 mV in 0.1 M HClO4 at

room temperature. Charge passed Q = 3 C/cm2, i = 200 A/cm

2, t = 15,000 s. Sample

fractured in the dealloying solution at E = -300 mV vs. MSE. ......................................... 45

Figure 7.28. Magnification of the fracture surface of Ag95Au5 dealloyed at E = 170 mV in

0.1 M HClO4 at room temperature. Charge passed Q = 3 C/cm2, i = 200 A/cm

2, t =

15,000 s. Sample fractured in the dealloying solution at E = -300 mV vs. MSE. ............. 45

Figure 7.29. Grain boundary corrosion on the fracture surface of Ag95Au5 dealloyed at E

= 170 mV in 0.1 M HClO4 at room temperature. Charge passed Q = 3 C/cm2, i = 200

A/cm2, t = 15,000 s. Sample fractured in the dealloying solution at E = -300 mV. ........ 46

ix



2, t = 35,000 s. Sample

fractured in the dealloying solution at the positive potential after dealloying. ................. 46



2, t = 35,000 s. Sample

fractured in the dealloying solution at the positive potential after dealloying. ................. 47

Figure 7.32. Spot 1 - Fracture surface of Ag95Au5 dealloyed at E = 170 mV in 0.1 M

HClO4 at room temperature. Charge passed Q = 3 C/cm2, i = 100 A/cm

2, t = 35,000 s.

Sample fractured in the dealloying solution at the positive potential after dealloying. .... 47


room temperature. Charge passed Q = 5 C/cm2, i = 6.5 mA/cm

2, t = 772 s. Sample

transferred and fractured in deionised water right after dealloying. .................................. 48

Figure 7.34. Close-up of the fracture surface of Ag95Au5 dealloyed at E = 300 mV in 0.1

M HClO4 at room temperature. Charge passed Q = 5 C/cm2, i = 6.5 mA/cm

2, t = 772 s.

Sample transferred and fractured in deionised water right after dealloying. ..................... 48



2, t = 772 s. ..... 49



2, t = 772 s. ..... 49



2, t = 772 s ...... 50



2, t = 772 s ...... 50

x




transferred and fractured in deionised water right after dealloying. .................................. 51



2, t = 410 s. ..... 51

Figure 7.41. Spot 2 - fracture surface of Ag95Au5 dealloyed at E = 300 mV in 0.1 M


2, t = 410 s. ..... 52

Figure 7.42. Spot 3 - fracture surface of Ag95Au5 dealloyed at E = 300 mV in 0.1 M


2, t = 410 s. ..... 52



2, t = 410 s. ..................... 53



2, t = 532 s. Fractured in

air. ...................................................................................................................................... 53



2, t = 532 s. Fractured in

air. ...................................................................................................................................... 54

Figure 7.46. Close-up of the fracture surface of Ag95Au5 dealloyed at E = 300 mV in 0.1

M HClO4 at room temperature. Charge passed Q = 4 C/cm2, i = 7.2 mA/cm

2, t = 532 s.

Fractured in air. .................................................................................................................. 54



2, t = 532 s.

Fractured in air. .................................................................................................................. 55

xi



2, t = 532 s.

Fractured in air. .................................................................................................................. 55



2, t = 532 s.

Fractured in air. .................................................................................................................. 56



2, t = 532 s.

Fractured in air. .................................................................................................................. 56



2, t = 532 s.

Fractured in air. .................................................................................................................. 57



2, t = 128 s. After

dealloying sample rinsed in deionised water and fractured in air. .................................... 57



2, t = 128 s. After




2, t = 128 s. After


xii



2, t = 128 s. After




2, t = 128 s. After




2, t = 128 s. After


Figure 7.58. Side surface of Ag95Au5 dealloyed at E = 170 mV in 0.1 M HClO4 at RT.

Charge passed Q = 50 mC/cm2, i = 360 A/cm

2, t = 125 s. After dealloying the sample

was rinsed in deionised water, ethanol and dried in air. Fractured in air. ......................... 61

Figure 7.59. Side surface of Ag95Au5 dealloyed at E = 170 mV in 0.1 M HClO4 at RT.

Charge passed Q = 50 mC/cm2, i = 360 A/cm

2, t = 125 s. After dealloying the sample

was rinsed in deionised water, ethanol and dried in air. Fractured in air. ......................... 61

Figure 7.60. Spot 1 - Nanoporous layer on the side surface of Ag95Au5 dealloyed at E =

170 mV in 0.1 M HClO4 at RT. Q = 50 mC/cm2, i = 360 A/cm

2, t = 125 s. ................... 62

Figure 7.61. Spot 2 – Grain boundary surface of Ag95Au5 dealloyed at E = 170 mV in 0.1

M HClO4 at RT. Charge passed Q = 50 mC/cm2, i = 360 A/cm

2, t = 125 s.................... 62

Figure 7.62. Potentiodynamic curves of Ag77Au22Pt1 and Ag77Au21Pt2 in 0.1 M HClO4.

Scan rate 0.5 mV/s. IR drop corrected. Ag77Au23 plays a role of a reference. .................. 63

xiii

Figure 7.63. Fracture surface. Nanoporous layer obtained from dealloying of Ag77Au23 at

60oC in 0.1 M HClO4 at E = 520 mV. Q = 1.2 C/cm

2, i = 10 mA/cm

2, t = 120 sec. Sample

fractured in air. Fracture surface. ....................................................................................... 64

Figure 7.64. Fracture surface. Nanoporous layer obtained from dealloying of Ag77Au22Pt1

at 60oC in 0.77 M HClO4 at E = 520 mV. Q = 2 C/cm

2, i = 8 mA/cm

2, t = 240 sec.

Sample fractured in the dealloying solution. Fracture surface. ......................................... 64

Figure 11.1. Polarization curves performed on Ag95 Au5 in 0.1 M and 1.0 M HClO4 at



in ten consecutive steps at different potentials in 0.1 M HClO4 at room temperature. One

side of the sample exposed to the dealloying solution. ..................................................... 71


in ten consecutive steps at different potentials in 0.1 M HClO4 at room temperature. Both

sides of the sample exposed to the dealloying solution. .................................................... 71

Figure 11.4. Fracturing device: A – electrochemical cell, B – mobile part made of

syringes. ............................................................................................................................. 72

Figure 11.5. Mobile part: C - the syringe tube, D – piston, E – small syringe with the

sample mounted in. ............................................................................................................ 73

Figure 11.6. Changes in the imaginary part of impedance after dealloying in ten

consecutive steps at E = 120 mV for Ag95 Au5 in 0.1M HClO4. Charge passed in each step

Q = 600 mC/cm2. ............................................................................................................... 73

xiv



Q = 600 mC/cm2. ............................................................................................................... 74



Q = 600 mC/cm2. ............................................................................................................... 74



Q = 600 mC/cm2. ............................................................................................................... 75



Q = 600 mC/cm2. ............................................................................................................... 75


consecutive steps at E = 550 mV for Ag77 Au23 in 0.1M HClO4. Charge passed in each

step Q = 0.6 C/cm2. ............................................................................................................ 76


consecutive steps at E = 650 mV for Ag77 Au23 in 0.1M HClO4. Charge passed in each

step Q = 0.6 C/cm2. ............................................................................................................ 76

1

Introduction

There is growing interest in nanoporous metals for which we are finding nowadays more and

more useful applications e.g., production of sensors or filtration purposes. Nanoporous

structures may be produced in different methods. The simplest way to do that is free corrosion

in a highly concentrated oxidizing agent or the application of electrochemical techniques (1, 2).

The properties of nanoporous metals fabricated by dealloying are remarkable, for instance,

nanoporous gold obtained from dealloying of silver–gold alloys possesses impressive surface to

volume ratio (3). Moreover, after dealloying the size of ligaments can be tuned by annealing.

Since in case of silver-gold alloys the ligaments are covered with gold, such a nanoporous

material can also find various applications in catalysis. Until now, however, not much attention

has been devoted to silver-gold alloys containing low percentage of gold, for example, Ag-5 at.

pct Au which is investigated in this research. This is because it was believed that these systems

would not undergo dealloying. Apart from obvious benefits of using silver-gold alloys

containing low percentage of gold, there is another important aspect correlated to the formation

of nanoporosity. There is a general agreement that dealloying processes can be the cause of

stress-corrosion cracking (SCC) (4, 5). Thus, the ability of an alloy to undergo dealloying may

be the key characteristic for assessing whether or not one may expect stress-corrosion cracking

to occur in a given alloy. Although silver-gold alloys are not typical construction materials used

in the industry but due to their chemical and physical properties they constitute a perfect object

for studying the mechanism of stress-corrosion cracking, the understanding of which will

broaden our knowledge on corrosion and help us to limit or prevent this corrosion attack in

construction materials.

2

1 Dealloying

Selective dissolution of one constituent from an alloy has been practised for centuries. First

applications of this process date back to the Indians of pre-Columbian Central America who

developed the method of depletion gilding (1). The mechanism of dealloying became recently

particularly interesting due to the fact that selective dissolution of the more active metal of an

alloy leads not only to the enrichment of the alloy surface in the more noble metal but also,

under certain conditions, to the formation of a nanoporous layer on the surface.

A good introduction to the selective dissolution phenomenon is works carried out by Pickering

(2, 6) who investigated the electrochemical behaviour of binary alloys of the general formula

AnB(1-n), where A refers to the less noble metal. Pickering investigated binary alloys using

potentiodynamic polarization and noted that the polarization curves for different alloys

exhibited certain similarities. He singled out their common features and marked them as

particular regions, as shown in Figure 1.1.

Figure 1.1. Illustration of the anodic behaviour of binary alloys by Pickering (2).

Region (a) refers to the part of a curve which exhibits passivation-like character, similarly to

the passive region of metals that undergo electrochemical passivation on the surface.

3

The passivation-like state in region (a) is due to the surface enrichment in the more noble

metal. This prevents further dissolution of the less noble metal and visibly restricts the

dissolution current. Because of that, this part does not noticeably depend on potential. Region

(b) is a potential-dependent increasing current region. The sudden change in current results

from the selective dissolution of the less noble metal. Similarly to region (b), region (c) is also

potential-dependent with this difference that only the initial dissolution is selective in character.

The two metals dissolve in this region. Region (d) is solely due to the constant dissolution of

the less noble metal. The accumulation of the more noble metal on the surface is either not

sufficient or a nanoporous layer formed is non-adherent. Region (d) is typical for pure metals or

alloys which are very rich in the more active metal. Analyzing the difference between the

standard reduction potentials of metals and the mole fraction of the more noble metal (NB),

Pickering noted a strong dependence between the two and the anodic behaviour of an alloy. He

divided the process of selective dissolution into two different types: Type I dissolution -

containing regions (a) and (b), and Type II dissolution - made up of regions (a) and (c).

Pickering concluded that the tendency increases for the Type I dissolution along with

increasing difference between the standard reduction potentials of the two metals. The tendency

for the Type II dissolution increases with the increase in the mole fraction of the more noble

metal. Thus, there must be the highest concentration for the more noble metal in a given binary

alloy, where such a transition between the two types of dissolution behaviour occurs. Pickering

denoted this threshold as (NB*). The presence of dealloying thresholds was confirmed

experimentally by Sieradzki et al. (7) who proposed explanations for the lack of agreement

between the conventional percolation theory and the observed dealloying thresholds in alloys,

including silver-gold systems (8). Recently, an attempt was made to explain the experimental

value of parting limit in silver-gold alloys using high density percolation theory (9).

4

Not much was devoted by Pickering towards the determination of a potential at which the

transition between regions of the passivation-like state and the selective dissolution takes place.

It was, however, proposed that a potential at which the transition occurs was the critical

potential (Ec) and that further increase in polarization would result in the generation of porosity

on the surface. In case of an alloy containing gold and a less noble metal, for instance, silver, a

nanoporous structure formed will consist of ligaments made up from two different phases: a

layer of pure gold on the outside and a gold-enriched alloy inside the ligaments, as shown in

Figure 1.2.

Figure 1.2. A schematic showing a cross-section of ligaments formed on the surface (2).

Since the formation of nanoporosity leads inherently to the development of surface area,

nanoporous materials obtained by selective dissolution can find various applications, for

instance, as catalysts or high surface area electrodes for advanced applications. Another reason

for which the understanding of dealloying processes has become important is that these

processes play an important role in corrosion and may be partially or even solely responsible

for the destruction of construction materials.

5

2 The dealloying critical potential

It has been generally accepted that different reactions leading to corrosion of alloys can be

initiated or even entirely governed by dealloying processes e.g., as in the case of stress-

corrosion cracking (SCC) in Cu-Zn or Cu-Au alloys (10-12). For obvious reasons, the

determination of a potential at which dealloying occurs has attracted a great deal of attention.

As we mentioned in the first part, a selective dissolution of the less noble metal gives sudden

rise to the anodic current preceded by the passivation-like state. The dissolution of the less

noble metal will continue with the increase in polarization as long as the less noble metal is

“available” in the alloy. The determination of a dealloying critical potential may be done in

different ways. Using the method of potentiodynamic polarization one needs to locate the onset

of the rapidly growing anodic current and extrapolate it to the baseline current (Figure 2.1).

Figure 2.1. Possible misinterpretation of the critical potential by extrapolation.

At this point, one may want to ask whether the method of dynamic polarization is the best

approach for the determination of Ec? It appears not to be. Figure 2.1 shows that reading off the

value corresponding to the critical potential may be quite ambiguous due to the fact that the

6

change between regions (a) and (b) is never sharp and that, as often occurs in transition points,

the onset of the change may depend on several parameters. These will be discussed below.

In potentiodynamic polarization, scan rate is one of those fundamental parameters that must be

set before each experiment and is normally expressed in millivolts per second. Sieradzki et al.

(13) examined several silver-gold alloys in perchloric acid investigating how sweep rate and

the composition of a solution may affect the appearance of a critical potential. Starting with the

Gibbs-Thomson effect, which relates the effect of curvature on chemical potential, they derived

an equation for a critical potential of a given binary alloy of the composition ApBg (g = 1 - p),

where (B) is the more noble metal. Sieradzki et al. proved the critical potential to be a function

of both the composition of an alloy and the composition of a dealloying solution.

]))2(

(2

ln[2

)2(

2)ln( 2

2

2

g

g

aJ

ND

nqnq

k

g

g

aa

nq

TkEE

o

SSB

A

BoH

crit n Eq. 1

where:

Eo - the standard reduction potential [V], kB - Boltzmann’s constant [J/K], q – unit charge [C],

n – number of electrons, a – lattice parameter [m], Ns – the number of atoms per unit area, Ds –

surface diffusivity [m2/s], δ – dissolution potential of an alloy [V], Ω - atomic volume [m

3], γ –

isotropic solid/electrolyte interfacial free energy [J/m2], Jo – exchange current density [A/m

2],

T – temperature [K].

Sieradzki et al. (13) started with the premise that, similarly to other kinetically controlled

processes such like the brittle-ductile transition in solids, the change from the passive-like state

to the selective oxidation of the less noble metal was rate dependent. They experimentally

demonstrated that the critical potential was a function of sweep rate.

7

Basically, there are a limited number of methods available by means of which one could

precisely determine the value of a critical potential for a given alloy in a given solution. In the

past, the simplest approach was to take this potential value as the critical one, at which the

current density is i = 1 mA/cm2 (13). Unfortunately, both the older method and the method of

potentiodynamic polarization have been proven to be an overestimation.

A great deal of work was done by Dursun et al. (14) who investigated dealloying of silver-gold

systems in 0.1M perchloric acid. On the contrary to what had been believed in the past, they

observed the formation of porosity well below the empirical critical potential. The novelty of

their approach was the application of potential hold experiments in the vicinity of the critical

potential. It was noted that the dealloying current eventually became nearly constant for the

samples potentiostatically polarized above (Ec), whereas samples polarized below the

dealloying potential displayed current continually decaying with time. An important conclusion

coming from this work is that the determination of the dealloying critical potential by its

extrapolation from a potentiodynamic curve may contain a significant error. The dealloying

critical potential can be overestimated even by a hundred millivolts. In addition, the method of

potentiostatic polarization, owing to its nature, eliminates the influence from sweep rate which

affects the shape of a polarization curve. In the light of what has been said, determining a

critical potential by potentiostatic polarization appears to be more precise than doing it in the

potentiodynamic fashion, because the necessity of extrapolation may itself be very subjective

and thus, apart from other issues, produce a large error. One should notice, however, that

potentiostatic experiments can also pose a serious problem since, as it is with every experiment,

this method requires time and time is a very important parameter in the kinetically controlled

processes. The influence from time on current fluctuations in potentiostatic experiments were

further investigated by Dursun et al. (15) who proved that short-time experiments can produce

8

very illusive effects because the observed current may give an impression of its decaying

character, whilst if the same experiment lasted longer, the investigator could notice the

development of a steady-state current and conclude that we are above the critical potential.

Apart from such parameters like scan rate or the nature of a given method (15), precise

determination of a critical potential may be problematic in the presence of particular species in

a dealloying environment. Dursun et al. (16) also examined the influence of halides and

demonstrated that the dealloying critical potential for silver-gold alloys was sensitive to their

presence. They proved that the addition of halides could affect the size of ligaments and

volume of pores in the resultant nanoporous layer, for instance, its porosity in the solution

containing iodide was almost ten times larger than in the absence of this ion.

As we may see, the presence of a critical potential (Ec) is the key characteristic of various

dealloying systems. If the concentration of the more noble component in an alloy is NB ≤ NB*,

a rapid growth in current density will mark the transition from a passive-like state to the region

of selective dissolution. The location of a critical potential may depend on the following

parameters: the difference between the standard reduction potentials of the constituents (ΔEo),

the mole fraction of the more noble metal (NB), sweep rate in potentiodynamic polarization,

time in potentiostatic experiments, the composition of an alloy and the content of a dealloying

environment.

9

3 Nanoporous metals

The application of high-porosity materials is now rapidly gaining interest due to their

functionality and unique chemical and physical properties (3, 17). Nanoporous structures may

be created in different forms, for instance, bulk or thin films. Although there are other ways in

which nanoporous materials can be produced, dealloying is still most commonly used

technique. Nanomaterials obtained by dealloying exhibit large surface to volume ratio and

remain uniform on the macroscopic scale. Moreover, the size of ligaments and pores can be

controlled and tailored specifically for a given application, for instance, ligaments can be

coarsen by annealing to the required size ranging from nanometers to micrometers. There are

four basic features that must be met for an alloy to be turned into a nanoporous material: 1)

significantly large difference in the standard reduction potentials between or among (in case of

ternary alloys) the components, 2) the mole fraction of the more noble constituent cannot be too

high, 3) the material must have homogenous structure prior to dealloying, 4) surface diffusivity

of the more noble metal must be sufficiently high (3).

Generally, binary systems used for the generation of nanoporous structures should possess

similar physical properties between the constituents such as: good solubility of the metals

across a broad range of concentrations used in the production of a given alloy, similar thermal

expansion or homogeneity. Because of that, most commonly used for this purpose alloys are

silver-gold, copper-gold or other similar systems that meet the requirements mentioned above.

It should also be noted that the number of available well understood and easily applicable

dealloying techniques greatly contributes to the popularity of a particular alloy.

Historically, Forty was probably one of the first scientists, if not the first one, who proposed the

surface diffusion mechanism for the formation of porosity caused by the selective oxidation of

10

the less noble metal. Forty used in his experiments high purity single crystals of silver-gold

containing equal concentrations of the two elements. The samples were exposed to 50 % nitric

acid for 30 and 60 seconds of free corrosion (1).

Figure 3.1. Transmission electron micrograph of Ag50 Au50 exposed to 50 % HNO3 (1).

Forty noted that selective dissolution of silver led to the generation of gold-rich regions on the

surface. Annealing these samples at 720 K caused the islands to merge and grow. On the basis

of his experimentally made observations he proposed the mechanism of porosity formation

during dealloying.

Figure 3.2. The growth of a gold island by the surface disordering (1).

Based on the kinetic Monte Carlo algorithm (KMC) (8), Erlebacher performed digital

simulations of dealloying (MESOSIM) and proposed that the formation of a nanoporous

structure by selective dealloying occurs in particular steps (3), as shown in Figure 3.3.

0.1 µm

11

Figure 3.3. Porosity evolution during selective oxidation proposed by Erlebacher (3).

Reproduced by permission of The Electrochemical Society.

According to this kinetic model, the evolution of porosity on the surface exposed to a

dealloying environment begins with stripping of atoms of the less noble metal (a). The

formation and coarsening of islands of the more noble metal (b) occurs as the second step. This

process is followed by stripping atoms of the less noble metal from the next layer (c) and the

accumulation of the more noble atoms at the hills (d). Undercutting of the hills gives the typical

shape to the ligaments (e). The nucleation of new noble clusters ends the process (f).

Coarsening will continue while the electrolyte penetrates the surface.

Although the simulations proposed by Erlebacher are based on certain simplifications, for

instance, the model does not include reduction processes such as redeposition of the less noble

metal on the anode, it has been accepted that the simulations mimic very well several aspects of

nanoporosity formation in real systems. Moreover, owing to the nature of a digital simulation,

one can obtain good insight in what happens on the surface during dealloying which, for

obvious reasons, cannot be observed in situ.

Apart from the digital simulations, which aim was to visualize the growth of nanoporosity,

Erlebacher performed simulations of potentiodynamic and potentiostatic curves. Comparison of

the simulated curves with other observations revealed that: 1) the empirical critical potential

12

does not separate the transition between passivation-like region from the formation of

nanoporosity, 2) the intrinsic critical potential does not depend on the sweep rate, 3) both the

intrinsic critical potential and empirical critical potential follow the same trend with the

composition of an alloy, 4) the intrinsic critical potential falls before the empirical critical

potential, 5) surface diffusion plays several roles during dealloying (3).

4 Stress-corrosion cracking (SCC)

Stress-corrosion cracking is one the most dangerous corrosion attacks that leads to a sudden

destruction of construction materials. Such problems are commonly known and occur in a

variety of industries: shipbuilding, aircraft, mining or facilities such as underground

transmissions pipelines. The prevention against SCC requires both the appropriate design of a

construction and its constant monitoring. During this corrosion process the material, which a

given construction is made of, may not show any visible features on the surface and at the same

time it can be filled with microscopic cracks inside. Thus, a sudden catastrophic failure may

occur “without warning” and the damage identification may not be easy for a casual inspection.

In order for a material to undergo stress-corrosion cracking, a combination of three different

factors must take place: 1) susceptible material, 2) presence of stress, 3) corrosive environment.

It is often misunderstood that only metals, alloys in particular, can undergo this type of

corrosion. Stress-corrosion cracking may also occur in polymers when exposed to corrosive

chemicals. Similarly to metals, the attack is confined to particular types of polymers in

particular environments. Thus, the chemical environment causing stress-corrosion cracking in a

given material may not have this adverse effect on other materials exposed to the same

environment.

13

Although the phenomenon of SCC has been studied for many decades, certain aspects still

remain unclear, for instance, the exact mechanism of stress-corrosion cracking. Generally, we

can distinguish five types of stress-corrosion cracking (18):

Type A – occurs by intergranular slip-dissolution

Type B – passive systems; cracking is induced by Cl-

Type C – surface dealloying

Type D – caused by films other than oxides or dealloyed layers

Type E – hydrogen embrittlement

For a given occurrence of SCC, a number of different models might be proposed as plausible

explanations. Some of them could be ruled out on thermodynamic grounds and thus they

cannot be perceived and applied as a universal model that would account for all intergranular

and transgranular examples of cracking in construction materials. As it was mentioned earlier,

stress (both residual and applied) is one of the three main factors contributing to this corrosion

attack. A metal subjected to a tensile strain may release a part of this stress by the motion of

dislocations (glide, creep) in the bulk material. The internal tensions caused by an external

force can also be relaxed by the generation of cracks inside the material.

Several mechanisms of stress-corrosion cracking have been proposed over the years but two of

them deserve special attention, namely the surface mobility model (SM) proposed by Galvele

(19, 20) and the film-induced cleavage (FIC) model of Sieradzki and Newman (4, 5), partially

originating from early works of Edeleanu and Forty (21). The two models will be briefly

discussed in the next two chapters.

14

4.1 The surface mobility SCC mechanism

The surface mobility SCC mechanism has been proposed by the late José R. Galvele and is

based on the assumption that a high stress present at the tip of a crack may locally create a

vacancy deficient region (20). According to the SM mechanism, the driving force for the crack

propagation will be the surface vacancy movement towards the tip of a crack, where vacancies

are captured or, equivalently, by the flow of adatoms in the opposite direction. The movement

of surface vacancies is possible due to the action from a corrosive environment which is

thought to be the cause for the formation of low melting point species on the walls of a crack. It

is suggested that such contaminants are not only responsible for increasing the surface mobility

but they may also contribute to the appearance of surface vacancies required in this SCC

mechanism. The movement of adatoms occurs in the first atomic layers and only these are

assumed to be taking part in the propagation of a crack (20), as demonstrated in the figure

below.

Figure 4.1. Crack propagation in the presence of an ionic contaminant by Galvele (19).

Following the above assumptions, a crack propagates an atomic distance after the stressed

lattice has captured a vacancy. Thus, the diffusion of vacancies (or adatoms) on the surface will

15

be the rate controlling process. Taking this into account, Galvele proposed the following

equation for the crack propagation rate (19).

]1)[exp(...3

kT

a

L

Drpc S Eq. 2

where:

c.p.r. - crack propagation rate [m/s], Ds - surface diffusion coefficient [m2/s], L – diffusion

distance of vacancies [m], σ – elastic surface stress at the tip of a crack [N/m2], a – the atomic

size [m], k – Boltzmann’s constant [J/K], T- temperature [K].

Figure 4.2. An illustration of the distances which atoms have to travel by Galvele (20).

In general, Galvele forms the following postulates for the SM mechanism (20):

sufficiently high surface mobility is created by the environment

SCC occurs at the temperature lower than 0.5 Tm

only elastic stress is relevant

surface vacancies are captured by the tip of a crack

16

4.2 The film-induced cleavage SCC mechanism

In the great majority of instances, corrosion reactions between a metal and a corrosive

environment occur at the interface between the two phases or in the vicinity of the interface.

Corrosion products might exist in different chemical forms such like oxides, salts etc. It has

been proven in the past that, under certain conditions, corrosion products may exhibit very

specific mechanical properties (22, 23). Various examples of stress-corrosion cracking failures

have shown over years the presence of common features in the materials destroyed this way,

for instance, the appearance of brittleness (5). It has been suggested that these failures could

result from cleavage-like events introduced by a thin film formed on the surface of a metal. It

should be noted here that this brittle-like character of a fracture has also been observed in

materials that exhibit inherent ductility (5). By contrast to other explanations of SCC failures

(18), Sieradzki and Newman developed a concept (4) according to which the mechanical

properties of a nanoporous structure could be the cause of SCC in metal/solution systems.

Following this assumption, a nanoporous layer formed on the surface in the process of selective

dissolution can inject micro-cracks into the “healthy” part of the material untouched by

corrosion. It was suggested that a minimal thickness of a nanoporous layer is required for this

type of fracture to occur (4, 5). On the contrary to the surface mobility SCC concept, which

assumes a continuous character of a fracture, Sieradzki and Newman have proved that a

cleavage event may occur in a series of consecutive steps (4) and will penetrate the material

until its arrest. The arrest might be due to several reasons such like crack bifurcation at a triple

point, exhaustion etc. For the reasons that were discussed previously, selective dealloying will

function here as the driving force because it can be repeated after each cleavage event. If the

conditions remain unchanged, that is, the stress and corrosive environment are still present, the

whole process of selective dissolution and crack propagation may happen again and will

17

proceed this way until the crack has entirely penetrated the material. An illustration of this

process is shown below.

Figure 4.3. An illustration of film-induced cleavage generated during dealloying. (Adapted

from Ref.18.)

Several different experiments have been carried out in the past when investigating a variety of

systems susceptible to stress-corrosion cracking. In the film-induced cleavage model, the most

important aspect lies in its basic assumption, that is, corrosion reactions leading to the

generation of nanoporosity can be separated from the action of stress. Following this

assumption, after a nanoporous layer has been formed on the metal surface, a sudden

application of stress can “trigger” the nanoporous layer to inject a crack into the uncorroded

part of the substrate. By doing it with the proper strain rate and at the appropriate potential,

other SCC mechanisms could simply be ruled out owing to the fact that the reactions, which are

thought to be the cause of these mechanisms, would need more time to occur.

The above aspects were investigated by Kelly et al. (24) who carried out experiments on silver-

gold alloy (Ag-20 at. pct Au) in 1M HClO4 applying stress under different conditions to the

pre-dealloyed samples. The purpose of this work was to confirm the ability of nanoporous

layers to generate a brittle cleavage into the “healthy” part of material. The single-shot fracture

experiments were performed in and out of the dealloying solution and it has been successfully

stress-corrosion crack

nanoporous layer brittle crack

plastic blunting

crack-arrest mark (striation)

18

proven that only under particular conditions the mechanical properties of a nanoporous layer

allowed a crack to reach a sufficient velocity and penetrate the uncorroded substrate.

Figure 4.4. A cross-sectional view of a sample that was bent while maintaining the dealloying

potential (24).

The tendency for a brittle fracture disappeared in the case of samples fractured in air. This

suggests that nanoporous layers are very susceptible to ageing which can reverse the

embrittlement. Kelly et al. (24) reported that both intergranular (IG) and transgranular (TG)

cracks were observed in the investigated samples. This stays in accordance to the observations

previously made by Newman et al. who did similar experiments on alpha-brass (Cu-35 wt.%

Zn) in cuprous ammonia solution (25) prior to the experiments on Ag-20 at. pct Au. Apart from

obvious differences regarding the composition of the alloys and dealloying environments, the

significant distinction between these two investigations is that the tests performed on alpha-

brass were performed without application of a potential, that is, by free corrosion of the metal

for up to 100 minutes. After the exposure, several procedures were undertaken before the

samples were strained. Brittle fractures were observed only in two cases; in the samples

strained right after dealloying and these ones which were transferred to liquid nitrogen and

strained to fail right after, as shown in Figure 4.5.

19

Figure 4.5. Fracture surface after application of strain: (a) brittle and (b) ductile (25).

Identically to the experiments on the silver-gold alloy, rapid ageing of nanoporous layers was

observed in alpha-brass. Tests with shorter immersion (10 minutes) produced similar effects

with this difference that the brittleness observed was not this pronounced as it was in the case

of long immersion tests (100 minutes).

The presence of transgranular SCC events in different copper-zinc systems was very well

documented over the past years. However, one may ask here a question whether it would be

possible for SCC to occur in pure copper. Sieradzki et al. (26) examined the behaviour of pure

copper single crystals exposed to 1 M NaNO2 and demonstrated a detailed fractographic study

on that. The novelty of this study was the combined application of single crystals and potential.

The samples were polarized (E = 0.0 V vs. SCE) during slow strain rate tests.

The striations seen on the fracture surface of a copper single crystal, as tiny horizontal lines

(Figure 4.6), suggest that the crack growth could be cleavage-like in its character because the

striations are parallel to the crack front. This brittle and cleavage-like appearance of cracks is

similar to the previously observed in Cu-Zn alloys (25, 27) and indicates that the crack

propagation may have the same or a similar fracture mechanism.

20

Figure 4.6. Crack front striations on the fracture surface of a copper single crystal (26).

Apart from various single-shot fracture experiments on pre-dealloyed samples followed by

microscopic assessment, transgranular stress-corrosion cracking in Cu-Zn systems was also

investigated by the analysis of acoustic emission that may accompany SCC events. The main

idea behind these experiments is that if stress-corrosion cracking propagates in a given material

as a series of short cleavage-like events, it is possible that a single event produces enough of

acoustic emission which can be amplified, recorded and correlated with the corresponding

increase in the anodic current. Similarly to the works done but by Pugh et al., Newman and

Sieradzki (28) carried out an investigation whose purpose was to demonstrate the discontinuous

character of transgranular events in α-brass. The tests were performed in 1M sodium nitrate

with potential held at E = 0 V vs. SCE. The test electrodes were fractured using a tensile

machine. The results of the acoustic emission were compared with the current variations

proving the discontinuous character of transgranular SCC in α-brass, as shown in Figure 4.7.

Scanning electron microscopy performed on the investigated samples also confirmed the

intermittent progress of cracks in this alloy (28).

21

Figure 4.7. Correlation of acoustic events with current for TG cracks in α-brass (28).

4.3 Short summary of the two SCC mechanisms

In the light of what has been discussed above, one may notice that there are certain similarities

and differences between the two SCC mechanisms. Both models assume that the environment

plays a fundamental role, that is, it initiates the process and increases the dissolution current. As

regards the film-induced cleavage model, this mechanism requires the composition of an alloy

to be in the appropriate ratio. A nanoporous layer can only be formed on the surface, if the

percentage of the more noble component is high enough but does not exceed its maximum

either (NB*). If it does, the two constituents will be oxidized (at higher potentials) without the

formation of nanoporosity. On the other hand, nanoporosity will not be generated when the

concentration of the more noble component is too low, that is, NB < 3% (2). Note that its

presence in the minimum concentration is strongly required for covering the ligaments during

dealloying. By contrast, the SM model assumes no correlation between SCC and dealloying,

thus the range of compositions which an alloy may stay in is much broader. Following the SM

model assumptions, the propagation of cracks is continuous, whereas in the film-induced

cleavage model, the crack progress is assumed to be a series of short cleavage-like steps.

22

5 Objectives

The main objective of this research is to investigate the mechanism of stress-corrosion cracking

in metals, film-induced cleavage in particular. Following film-induced cleavage assumptions,

the action of stress and anodic dissolution are thought to be discrete aspects that can be

separated in time (29, 30). Because film-induced cleavage is correlated with dealloying, the

formation of nanoporosity will constitute an integral part of this research. By contrast to what

has been suggested by Galvele (31, 32), we believe that formation of nanoporosity in silver-

gold alloys containing low concentration of gold is possible and therefore dealloying may

accompany or be the cause of stress-corrosion cracking observed in these alloys. For the

purpose of this investigation, binary and ternary noble-metal alloys such as silver-gold and

silver-gold-platinum have been chosen. The investigation is divided into three different parts:

1) electrochemistry, 2) single-shot fracture tests and 3) digital modelling. The latter issue,

although not performed by the present researcher, will constitute an attempt to correlate the

dealloying behaviour of the alloys in real conditions with those modelled digitally (33).

5.1 Single-shot fracture experiments

The purpose of single-shot fracture tests is to find the conditions under which film-induced

cleavage occurs. These experiments will be performed with help of a lab-constructed device by

means of which it will be possible to fracture a sample at any time of experiment, that is, when

the physical properties of a nanoporous layer formed have met the assumed requirements. The

device will allow on fracturing samples by the application of a sudden strain, in or out of a

dealloying solution, at room or elevated temperature. Samples fractured in this way will be

examined with scanning electron microscopy (SEM) and the following aspects will be

investigated: crack depth, thickness of a nanoporous layer, intergranular stress-corrosion

23

cracking (IGSCC), transgranular stress-corrosion cracking (TGSCC) and grain boundary

corrosion

5.2 Electrochemistry

The aim of this part is to answer how electrochemical conditions e.g., a potential applied,

dealloying time, temperature or the concentration of a dealloying solution, can affect the

following features of a nanoporous layer: 1) the size of ligaments, 2) development of surface

area versus charge passed, 3) thickness of a nanoporous layer versus charge passed.

Since the formation of nanoporosity occurs in the region of selective dissolution, the

experiments will need to be carried out above the critical potential. This requires its value to be

established in the first instance.

5.3 Digital simulations

In this particular part of the research, the present student will contribute by performing required

electrochemical experiments (33). The digital modelling will be carried out using the modified

version of the software MESOSIM. We will attempt to correlate the dealloying behaviour of a

particular noble alloy with Kinetic Monte Carlo simulations of the dealloying process.

Specifically, the issue of retained residual material in ligaments and the development of surface

area during dealloying will be under consideration. The results will then be applied in further

quantifications.

24

6 Experimental

The anodic behaviour of Ag-Au alloys has been investigated in perchloric acid of different

concentrations. The present researcher did also an attempt to examine Ag-Au-Pt alloys, but due

to the amount of work which these alloys required, only a few significant results were obtained.

6.1 Materials

High purity binary alloys (99.99 %) of the compositions Ag-5 at. pct Au and Ag-23 at. pct Au

were provided by Goodfellow Metals in the form of thin foil 100 µm thick. Ternary alloys

(99.95 %) of the following compositions: Ag77Au22Pt1, Ag77Au21Pt2, Ag77Au19Pt4 were

supplied by Ames Laboratory in the form of rolled strips approximately 3 cm x 30 cm. Their

thickness varied from 100 µm to 300 µm.

6.2 Samples preparation

Prior to use in the electrochemical cell, the materials were cut into rectangular pieces, 10 mm x

3 mm (for single-shot fracture tests) and 5mm x 6mm (other purposes), abraded manually with

Buehler 2500 grit sand paper, polished with a diamond paste (1 m) and thoroughly rinsed with

deionised water and acetone. Samples prepared this way were ultrasonically cleaned up in

ethanol and annealed in two different ways: at 900oC for one hour in air and for two hours in

argon with 2.5% of hydrogen. After annealing, the samples annealed in air were furnace cooled

and water quenched after the temperature had dropped to approximately 400oC. The samples

annealed in the gas were furnace cooled until the temperature had dropped below 180oC and

water quenched. Care was taken so as not to apply any cold work to the annealed samples

during the preparations.

25

6.3 Electrochemical cell

The three electrode electrochemical cell consisted of a 500 mL glass container with the average

volume of a solution 200 mL.

Figure 6.1. Electrochemical cell used in experiments: WE – working electrode, RE – reference

electrode, CE – counter electrode.

In all electrochemical investigations, Gamry 600 and Gamry Femtostat potentiostats were

employed. Single-shot fracture tests were carried out in a lab-made device that has been

designed by the present researcher. All the solutions of perchloric acid were prepared from

stock chemicals and deionised water of the specific resistivity 18 MΩ-cm obtained from a

Barnstead Nanopure system. Experiments were carried out at room temperature. A saturated

mercury/mercurous sulphate electrode (MSE) was used as a reference electrode (E = 640 mV

versus the standard hydrogen electrode SHE). In order to limit any interactions with the

sulphate ion, during all experiments the reference electrode was always residing in a double

junction. All the potentials in this work are given versus the MSE electrode. A five centimetres

long platinum wire was used as a counter electrode. Before each experiment the platinum wire

was cleaned up electrochemically for 5 minutes in 0.25 M HNO3. In some cases a dealloying

solution was deaerated by purging it for 30 minutes with nitrogen.

WE RE N2 CE

26

6.4 Mounting samples

Depending on the type of an experiment, samples were mounted in three different ways.

Samples destined for potentiodynamic experiments were covered with lacquer (Microshield) so

that only one surface area was exposed to the solution. This was done to eliminate interference

from the edges which normally tend to corrode at higher rates than flat areas. Another type of

mounting was employed for samples that were to be fractured after dealloying (Figure 6.2).

These were soldered to a copper wire and covered with lacquer in such a way that only the

central part of the sample was exposed to the dealloying solution. Copper wire was used to

form a small loop attached to the sample, the purpose of which was to block the sample during

fracturing it by a short and rapid pull-up. The loop (0.1 - 0.12 g) was also covered with lacquer.

Figure 6.2. Mounting a sample for a single-shot fracture test.

Figure 6.3 illustrates the way samples were fractured in all single-shot fracture tests. Since each

experiment was carried out on a thin (100 µm) and very soft material, fracturing samples

required doing it in a sufficiently reliable way, that is, without bending the sample prior to the

fracture. A 60 mL syringe was adapted for this purpose. Holes were drilled in the syringe tube

in order to provide a free flow of the solution. A deep hole made in the centre of the piston

allowed on embedding samples in an easy way. An insulated piece of a metal drill was led

through the loop each time after loading the sample in the tube. The piston was kept

immobilized over each experiment until the final pull-up breaking a sample in two, after which

loop sample joint coated copper wire

27

a sample was transferred to deionised water. Since straining samples was always done

manually, the average time required for fracturing a sample was assumed as t = 0.5 s.

Figure 6.3. An illustration of the electrochemical cell used in single-shot fracture tests: WE –

working electrode, RE - reference electrode, CE – counter electrode.

Occasionally, samples used for any other purposes than the ones described above, were

prepared just by soldering the sample to a copper wire and insulating the joint with lacquer.

The wire was coated with heat-shrink tubing providing sufficient protection.

RE CE N2

WE

28

7 Results and Discussion

7.1 Formation of nanoporosity by dealloying of Ag-5 at. pct Au

According to percolation theory (33), in order to form a connected random structure, the critical

volume fraction should be at least 16%. Following this assumption, and taking into account the

composition of the investigated alloy, the formation of a 3D structure will only be possible if

the nanoporous layer contains the residual silver, which has not been removed during

dealloying. In other words, if the whole silver is dissolved, a nanoporous layer will not be

formed due to the insufficient concentration of gold in the alloy. Figure 7.1 presents a

nanoporous layer obtained by dealloying of Ag-5 at. pct Au in 0.1M HClO4.

Figure 7.1. Nanoporous layer obtained by dealloying of Ag95Au5 dealloyed at E = 170 mV in

0.1 M HClO4 at room temperature. Charge passed Q = 3 C/cm2, i = 100 A/cm

2, t = 35,000 s.

The above result has significant implications. The formation of nanoporosity by dealloying of

alloys containing low percentage of the more noble metal can be achieved. Thus, a dealloying

based mechanism of stress-corrosion cracking is possible in Ag-5 at. pct Au. This is in

contradiction to what was suggested by Galvele (31, 32).

Side surface

Fracture surface

Thickness

29

7.2 The composition of a nanoporous layer formed on Ag-5 at. pct Au

The composition of a nanoporous layer has been determined with SEM/EDX performed on the

fracture surface, as shown in Figure 7.2.

Figure 7.2. Fracture surface of Ag95Au5 dealloyed at E = 300 mV in 0.1 M HClO4 at room

temperature. Charge passed Q = 3 C/cm2, i = 5 mA/cm

2, t = 628 s. Sample fractured in air.

Marker 50 µm.

Figure 7.3. EDX performed on the fracture surface of Ag95Au5 dealloyed at E = 300 mV in 0.1

M HClO4 at room temperature. Charge passed Q = 3 C/cm2, i = 5 mA/cm


fractured in air. Marker 50 µm.

Ag 83%

Ag 95%

30

The above result confirms the assumption regarding the expected high concentration of the

residual silver in ligaments. This fact can be very important from a synthetic point of view

because nanoporous layers fabricated from silver-gold alloys containing low percentage of gold

will be cheaper than those obtained by dealloying of Ag77Au23. Moreover, rich in silver

ligaments covered with a thin layer of gold may display similar properties as gold rich Ag-Au

alloys and could successfully be used in catalytic processes.

7.3 The growth of nanoporous layers on Ag-5 at. pct Au

In order to assess how the thickness of a nanoporous layer depends on charge passed, there was

performed a series of potentiostatic experiments at E = 300 mV during which different amounts

of charge were passed. The results of these experiments are presented in the next five figures.

All the current densities presented under the figures have been averaged.



2, t = 772 s.

31



2, t = 532 s.

Figure 7.6 and Figure 7.7 show the continual reduction in the observed thickness of a

nanoporous layer along with smaller amounts of charge passed during dealloying.



2, t = 410 s.

32



2, t = 296 s.



2, t = 128 s.

The conclusion coming from the above presented figures is that the thickness of a nanoporous

layer depends on the charge passed. We need to bear in mind, however, that the samples were

transferred from one environment to another during the preparations. This could lead to certain

physical changes. Thus, it was necessary to illustrate a theoretical thickness of a nanoporous

33

layer and compare it with the observed thickness (Figure 7.9). A theoretical thickness was

calculated by making two different assumptions: 1) the percentage of residual silver in the

ligaments is Ag = 83 % (as determined with EDX), 2) there is no residual silver in the

ligaments (the nanoporous layer consists of gold). The calculations have been done using the

following equation (33).

][AgFd

qMH

Ag

Ag Eq. 3

where:

q – charge passed [C/cm2], MAg – molar mass of silver, dAg – density of silver, F – Faraday’s

constant, [Ag] – the original content of silver in the alloy, - the fraction of available silver

removed during dealloying.

Figure 7.9. The observed and theoretical thickness of nanoporous layers obtained by dealloying

of Ag95 Au5 in 0.1M HClO4 at E = 300 mV.

34

The aim of further experiments was to investigate the development of surface area in

nanoporous layers during their formation. An assumption was made that dealloying at different

potentials would take different amount of time for passing the same charge. Since nanoporous

layers undergo physical changes such as coarsening or collapse, it seemed to be possible that

these changes could be observed by indirect measurement of capacitance. Thus, for the purpose

of these experiments, there were selected five potentials: 120 mV, 170 mV, 220 mV, 300 mV

and 400mV. Electrochemical impedance spectroscopy (EIS) was employed because this

method has been proven to be capable of measuring surface area of nanoporous layers (34).

Every potentiostatic experiment consisted of the same sequence: dealloying and EIS

measurement. During each step, equal amount of charge Q = 0.6 C/cm2 was passed. Impedance

was measured at constant potential E = 70 mV. The results of all the potentiostatic experiments

are shown in Figure 7.10 and Figure 7.11. See the Appendix for the raw impedance data.

Figure 7.10. Changes in capacitance versus charge passed during dealloying of Ag95 Au5 in ten

consecutive steps at different potentials in 0.1 M HClO4 at room temperature (33).

7.55 E-4 [F/cm2]

35

Figure 7.11. The development of relative area versus charge passed during dealloying of Ag95

Au5 in ten consecutive steps at different potentials in 0.1 M HClO4 at room temperature. (33)

Figure 7.12 and Figure 7.13 present the above results in a logarithmic scale.


consecutive steps at different potentials in 0.1 M HClO4

Area increased 20 times

36

Figure 7.13. The development of relative area versus charge passed during dealloying of Ag95

Au5 in ten consecutive steps at different potentials in 0.1 M HClO4

The results demonstrated in Figures 7.10 and Figure 7.11 show that there is the direct

correlation between the changes in capacitance and surface area. In the case of the highest

dealloying potential, the developed surface area was 1700 larger than the initial one. Regardless

of the applied potentials, no decrease in capacitance was observed. This proves that, under the

applied conditions, dealloying of Ag-5 at. pct Au results in the continual development of

surface area and that the application of electrochemical impedance spectroscopy is possible for

the characterization of nanoporous layers during their formation. The reproducibility of the

above experiments is displayed in Figure 11.2 and Figure 11.3 in the Appendix. One needs to

keep in mind that in case of potentiostatic experiments performed at lower potentials, long

exposure of a nanoporous layer to the dealloying solution will lead to coarsening of ligaments.

This could be minimized by carrying out potentiostatic experiments in a dealloying solution of

higher concentrations. As it will be shown in the next chapter (Figure 7.21), current density

increases along with the concentration of a dealloying solution.

37

The development of surface area during dealloying of Ag-5 at. pct Au in 0.1M perchloric acid,

has been compared with Ag-23 at. pct Au. Since the empirical critical potentials for Ag-5 at.

pct Au and Ag-23 at. pct Au are different (Figure 7.18), the assumption was made that in order

to quantify the two alloys, one must stay in the same distance from the critical potential of Ag-5

at. pct Au and Ag-23 at. pct Au. The below table presents the potentials applied.

Alloy The empirical critical

potential vs. MSE

Potential applied I Potential applied II

Ag-5 at. pct Au Ec = 50 mV E = 300 mV E = 400mV

Ag-23 at. pct Au Ec = 300 mV E = 550 mV E = 650mV

Figure 7.14. Potentials used for potentiostatic dealloying of Ag95Au5 and Ag77Au23 performed

in ten consecutive steps in 0.1 M HClO4 at room temperature.

The results of these experiments are shown in Figure 7.15 and Figure 7.16.

Figure 7.15. Changes in capacitance versus charge passed during dealloying of Ag95Au5 and

Ag77Au23 in ten consecutive steps at different potentials in 0.1 M HClO4.

38


Ag95Au5 and Ag77Au23 in ten consecutive steps at different potentials in 0.1 M HClO4.

Changes in capacitance observed during dealloying of Ag-23 at. pct Au (Figure 7.15) indicate

that the development of surface area is significantly different compared to Ag-5 at. pct Au, as

shown in Figure 7.16. The table below shows the average current density vs. potential applied.

Ag-5 at. pct Au E = 300 mV i = 9.0 mA/cm2

Ag-5 at. pct Au E = 400 mV i = 20 mA/cm2



Figure 7.17. Current densities observed during dealloying Ag95Au5 and Ag77Au23.

7.4 The anodic behaviour of Ag-5 at. pct in perchloric acid

The anodic behaviour of Ag-5 at. pct Au has been investigated in perchloric acid. The obtained

potentiodynamic curves, Figure 7.18 and Figure 7.19, confirm the typical shift in a critical

potential in relation to the composition of a binary alloy. This was reported in the past (2, 15).

39

Figure 7.18. Polarization curves performed on Ag95 Au5 and Ag77 Au23 in 0.1 M HClO4 at room

temperature. Scan rate 0.5 mV/s. IR drop corrected.

Figure 7.19. Polarization curves performed on Ag95 Au5 and Ag77 Au23 in 1.0 M HClO4 at room


The empirical critical potential for Ag-5 at. pct Au was further investigated in 0.1M perchloric

acid of different degree of aeration and concentration, as shown in Figure 7.20 and Figure 7.21.

40

Figure 7.20. Potentiodynamic curves performed on Ag95 Au5 in HClO4 of different aeration

degree at room temperature. Scan rate 0.5 mV/s. IR drop corrected.

Figure 7.21. Potentiodynamic curves performed on Ag95Au5 in HClO4 of different

concentration at room temperature. Scan rate 0.5 mV/s. IR drop corrected.

All the potentiodynamic experiments displayed excellent reproducibility. Similar curves

performed on Ag-5 at. pct Au in 1.0 M and 0.1 M HClO4 are presented in the Appendix.

41

7.5 Fracture experiments

Prior to single-shot fracture tests, there was performed a series of experiments whose aim was

to examine physical properties of nanoporous layers. Following the assumptions of film-

induced cleavage, a nanoporous layer formed on the surface must be of the appropriate

thickness in order to trigger a brittle crack. The presence of premature cracks in the nanoporous

layer or coarsening of the layer could have adverse effects in these experiments. Due to the

above, only potentials giving sufficiently high rate of dealloying were used.

7.6 Premature cracks and coarsening of a nanoporous layer

Figure 7.22 and Figure 7.23 show the presence of premature cracks on the side surface and

their absence after dealloying performed at two different potentials. This fact was taken into

account during fracture tests. Figure 7.22 also confirms the observations that were reported in

the past (35, 36).

Figure 7.22. Premature cracks on the surface after dealloying of Ag95Au5 at room temperature

in 0.1 M HClO4 at E = 300 mV vs. MSE. Marker 200 µm.

42

Figure 7.23.Ag95Au5 surface free from premature cracks after dealloying at room temperature

in 0.1 M HClO4 at E = 170 mV vs. MSE. Marker 20 µm.

Closer examination of nanoporous layers revealed coarsening of ligaments and its strong

dependence on time, as shown in Figure 7.24 and Figure 7.25.

Figure 7.24. Nanoporous layer after dealloying of Ag95Au5 in 0.1 M HClO4 for t = 125 s at E =

170 mV. Charge passed Q = 50 mC/cm2.

Note that the ligaments seen in Figure 7.24 are almost three times smaller than the ligaments

shown in Figure 7.25. This is due to the shorter time of dealloying not smaller charge passed.

43

Figure 7.25. Nanoporous layer after dealloying of Ag95Au5 in 0.1 M HClO4 for t = 2528 s at E

= 170 mV. Charge passed Q = 0.5 C/cm2.

7.7 Fracture tests

The main goal of single-shot fracture tests was to find the conditions under which transgranular

or intergranular film-induced cleavage could be generated. For this purpose, there was

performed a series of experiments in which both charge passed and potential applied were

controlled. Generally, each such an experiment consisted of the following steps:

dealloying a sample at a given potential

fracturing the sample after dealloying

SEM imaging

Fracturing samples after dealloying was performed in three different environments:

1) in the dealloying solution (at the positive or negative potential), 2) in deionised water and

3) in air.

There has been identified a linear relation between the thickness of a dealloyed layer and crack

depth for Ag-5 at. pct Au at E = 300 mV, as shown in Figure 7.26.

44

Figure 7.26. Crack depth versus thickness of a nanoporous layer for Ag95Au5 dealloyed in 0.1M

perchloric acid at 300 mV. Samples fractured beyond the dealloying solution.

Since dealloying at higher potentials had produced premature cracks (Figure 7.22), it was

assumed that single-shot fracture tests should be performed on the material pre-dealloyed at

lower potentials, neglecting the fact that dealloying at lower potentials inherently extends time

of the process and introduces coarsening. Based on the results presented in Figures 7.4 to 7.8,

the amount of charge passed, Q = 3 C/cm2, was assumed to be high enough for generating

sufficiently thick nanoporous layer that could trigger a brittle crack. Several fracture tests were

performed during which samples were fractured at the negative potential (E = -300 mV vs.

MSE). This was done in order to eliminate or limit faradaic reactions that could result in

superficial dealloying of the freshly revealed fracture surface. The negative polarization was

usually done from the dealloying potential in one step, after which short time (5-7 s) was given

for the cathodic current to saturate at the constant level. Similar fracture experiments to these

performed at E = -300 mV were also performed at E = -150 mV. Identical results were

observed in both cases. An example of such experiments is shown in Figures 7.27 to 7.29.

45

Figure 7.27 Fracture surface of Ag95Au5 dealloyed at E = 170 mV in 0.1 M HClO4 at room

temperature. Charge passed Q = 3 C/cm2, i = 200 A/cm

2, t = 15,000 s. Sample fractured in the

dealloying solution at E = -300 mV vs. MSE.

Figure 7.28 shows a magnification of the area marked with the rectangle in Figure 7.27.

Figure 7.28. Magnification of the fracture surface of Ag95Au5 dealloyed at E = 170 mV in 0.1

M HClO4 at room temperature. Charge passed Q = 3 C/cm2, i = 200 A/cm

2, t = 15,000 s.

Sample fractured in the dealloying solution at E = -300 mV vs. MSE.

One may have an impression that the fracture surface seen in Figure 7.27 has a brittle character

with only a tiny ductile region visible in the centre of the sample. Higher magnification of the

46

brittle part, however, revealed the presence of features that could be a deposition or superficial

dealloying on the fracture surface, as shown in Figure 7.29.

Figure 7.29. Grain boundary corrosion on the fracture surface of Ag95Au5 dealloyed at E = 170

mV in 0.1 M HClO4 at room temperature. Charge passed Q = 3 C/cm2, i = 200 A/cm

2, t =

15,000 s. Sample fractured in the dealloying solution at E = -300 mV.

Since the negative polarization could cause re-deposition of silver, fracture tests were also

performed at the positive potential. The results are shown in Figures 7.30 to 7.32.




dealloying solution at the positive potential after dealloying.

47

Figure 7.31 shows the magnification of the area marked in Figure 7.30.




dealloying solution at the positive potential after dealloying.

Figure 7.32 presents the magnification of the spot 1 marked in Figure 7.31.

Figure 7.32. Spot 1 - Fracture surface of Ag95Au5 dealloyed at E = 170 mV in 0.1 M HClO4 at


2, t = 35,000 s. Sample fractured

in the dealloying solution at the positive potential after dealloying.

Because fracture tests performed in a dealloying environment inherently involve sudden

exposure of fresh areas of the tested material to the dealloying solution, other single-shot

1

48

fracture tests on pre-dealloyed samples were performed beyond the dealloying solution. This

was done to eliminate the presence of solution that could cause corrosion after the fracture.


temperature. Charge passed Q = 5 C/cm2, i = 6.5 mA/cm

2, t = 772 s. Sample transferred and

fractured in deionised water right after dealloying.

Figure 7.34 shows the deepest crack found on the sample demonstrated in the figure above. The

fracture surface was examined with SEM at four different spots, as shown below.

Figure 7.34. Close-up of the fracture surface of Ag95Au5 dealloyed at E = 300 mV in 0.1 M



transferred and fractured in deionised water right after dealloying.

1

2

3 4

49

The following four figures will present magnifications of the four spots marked in Figure 7.34.

Note that the depth of pits on the surface decreases along with the distance from the dealloyed

layer on the surface. This suggests that corrosion along grain boundaries has a gradual

character. Figure 7.38 presents corrosion-free surface found on the investigated fracture

surface.



2, t = 772 s.



2, t = 772 s.

50



2, t = 772 s

Figure 7.38 shows the magnification of the spot 4 marked in Figure 7.34.



2, t = 772 s

Similarly to the results presented above, Figure 7.39 shows a fracture test performed on a pre-

dealloyed sample. The only difference is smaller amount of charge passed (Q = 3 C/cm2). The

nanoporous layer obtained by passing the above charge was assumed to be thick enough

(Figure 7.6) to trigger a brittle crack.

51



2, t = 410 s. Sample transferred and

fractured in deionised water right after dealloying.

The next three figures will demonstrate magnifications of the spots marked in Figure 7.39.

Figure 7.40 shows superficial dealloying on the brittle part of the fracture surface. Since this

area of the sample was exposed to the dealloying solution for longer time than the deeper part,

one may expect more intense coarsening of ligaments.



2, t = 410 s.

1

2

3

52

Figure 7.41. Spot 2 - fracture surface of Ag95Au5 dealloyed at E = 300 mV in 0.1 M HClO4 at


2, t = 410 s.

Figure 7.42. Spot 3 - fracture surface of Ag95Au5 dealloyed at E = 300 mV in 0.1 M HClO4 at


2, t = 410 s.

Figure 7.43 demonstrates magnification of the rectangular area marked in Figure 7.42.

Similarly to Figure 7.38, the fracture surface is free from corrosion. This indicates that the

penetration of the material by dealloying processes has reached its maximum depth.

53



2, t = 410 s.

On the basis of the results presented so far, one can make a conclusion that in some cases

corrosion seen on the fracture surface could result from reactions caused by the dealloying

solution transferred to DI water together with the sample. The next set of experiments was

performed to eliminate this possibility (Figure 7.44 to 7.51). Before fracturing in air, the

samples were rinsed with DI water after dealloying and then fractured.



2, t = 532 s. Fractured in air.

54

Figure 7.45 demonstrates the deepest crack found on the sample presented above.




The next figures will show magnifications of the five spots marked in Figure 7.46. Note the

preferential attack of dealloying along the grain boundary.

Figure 7.46. Close-up of the fracture surface of Ag95Au5 dealloyed at E = 300 mV in 0.1 M


2, t = 532 s. Fractured

in air.

1

2

3

4 5

Uniform progress

of dealloying

Preferential progress

of dealloying

55

Closer examination of the spots marked in Figure 7.46 also showed surface dealloying. The

degree of grain boundary corrosion varied along with the distance from the side surface.







In some instances corrosion pits were found to be significantly deeper at the ductile part of the

fracture surface. This is seen in Figure 7.47 and Figure 7.50. The difference in the depth of pits

56

and degree of dealloying may be due to the limited access of the solution and different time

during which a given part of the sample was exposed to the solution during dealloying.







Also in this case, corrosion-free surface area was found, as shown in Figure 7.51. This means

that the material has been dealloyed along the grain boundaries only to some depth.

57




Similarly to the results demonstrated above, Figure 7.52 to Figure 7.57 will show a fracture test

performed on pre-dealloyed samples fractured in air, after dealloying performed at E =300 mV

during which smaller amount of charge was passed. This was done to limit the penetration of

grain boundaries.



2, t = 128 s. After dealloying sample

rinsed in deionised water and fractured in air.

58

Figure 7.53 shows the magnification of the rectangle marked in Figure 7.52 and four marked

spots which magnifications are presented in Figures 7.54 to 7.57.



2, t = 128 s. After dealloying sample

rinsed in deionised water and fractured in air.



2, t = 128 s. After dealloying

sample rinsed in deionised water and fractured in air.

1

2

3

4

59









60





The results demonstrated until now showed three different environments in which samples

were fractured, that is, in the dealloying solution, in DI water and in air. Since the smallest

amount of charge passed was Q = 1 C/cm2, a conclusion was made that only very shallow

dealloying could limit the propagation of corrosion processes along grain boundaries. For this

purpose, there was performed an experiment in which the sample was dealloyed at the potential

which never caused premature cracks. Compared with all the presented results, the amount of

charge passed in this experiment was significantly smaller (Q = 50 mC/cm2). After dealloying,

the sample was transferred from the dealloying solution to deionised water, rinsed with ethanol

and dried in air (Figure 7.58 to Figure 7.61). Although the fracture surface exhibited only

ductile features, the applied strain revealed grain boundaries which were examined. In this case

crack depth to the thickness of the nanoporous layer was beyond expectations.

61

Figure 7.58. Side surface of Ag95Au5 dealloyed at E = 170 mV in 0.1 M HClO4 at RT. Charge

passed Q = 50 mC/cm2, i = 360 A/cm

2, t = 125 s. After dealloying the sample was rinsed in

deionised water, ethanol and dried in air. Fractured in air.

Figure 7.59 shows the magnification of the rectangle marked in Figure 7.58.

Figure 7.59. Side surface of Ag95Au5 dealloyed at E = 170 mV in 0.1 M HClO4 at RT. Charge

passed Q = 50 mC/cm2, i = 360 A/cm

2, t = 125 s. After dealloying the sample was rinsed in

deionised water, ethanol and dried in air. Fractured in air.

The next two figures show magnifications of the spots 1 and 2, marked in Figure 7.59.

1

2

62

Figure 7.60. Spot 1 - Nanoporous layer on the side surface of Ag95Au5 dealloyed at E = 170

mV in 0.1 M HClO4 at RT. Q = 50 mC/cm2, i = 360 A/cm

2, t = 125 s.

As it was in any other case, examining the side of the grain revealed rough surface with pits.

Figure 7.61. Spot 2 – Grain boundary surface of Ag95Au5 dealloyed at E = 170 mV in 0.1 M

HClO4 at RT. Charge passed Q = 50 mC/cm2, i = 360 A/cm

2, t = 125 s.

Figure 7.61 ends the results of all the fracture tests performed on Ag-5 at. pct Au in perchloric

acid. There has not been found any particular relation between the ways in which samples were

fractured and the depth of grain boundary penetration. This differs from the results reported in

the past on the behaviour of silver-gold alloys containing higher concentration of gold (24, 29).

63

The next three figures will present initial results of the experiments done on the ternary alloys

that also constitute a part of this research. Figure 7.62 shows potentiodynamic curves

performed on Ag77Au22Pt1 and Ag77Au21Pt2 in 0.1 M HClO4. The material containing 4% of

platinum has not yet been well examined. The main purpose of this investigation is to test

whether small additions of platinum introduced to Ag77Au23 at the expense of gold, may have

an impact on stress-corrosion cracking. This is based on the assumption that lower surface

diffusivity of platinum may slow down the diffusion of gold during dealloying and affect the

formation of nanoporosity.

Figure 7.62. Potentiodynamic curves of Ag77Au22Pt1 and Ag77Au21Pt2 in 0.1 M HClO4. Scan

rate 0.5 mV/s. IR drop corrected. Ag77Au23 plays a role of a reference.

The reproducibility of the results obtained from potentiodynamic curves performed on the

ternary alloys left a bit to desire and will be investigated in the future. Long annealing, up to 50

hours, did not improve the reproducibility. Its lack was probably due to the inhomogeneity of

the material. Figure 7.64 presents the very initial results of potentiostatic dealloying of the

64

ternary alloys and show the difference in the size of ligaments between Ag77Au23 and

Ag77Au22Pt1.

Figure 7.63. Fracture surface. Nanoporous layer obtained from dealloying of Ag77Au23 at 60oC

in 0.1 M HClO4 at E = 520 mV. Q = 1.2 C/cm2, i = 10 mA/cm

2, t = 120 sec. Sample fractured

in air. Fracture surface.

Figure 7.64. Fracture surface. Nanoporous layer obtained from dealloying of Ag77Au22Pt1 at

60oC in 0.77 M HClO4 at E = 520 mV. Q = 2 C/cm

2, i = 8 mA/cm

2, t = 240 sec. Sample

fractured in the dealloying solution. Fracture surface.

65

8 Conclusions

The results obtained from dealloying of Ag-5 at. pct Au in perchloric acid have proven that it is

possible to generate a three dimensional nanoporous layer by dealloying of silver-gold systems

containing low concentration of gold (Figure 7.4 to Figure 7.8). Dealloying of these systems

can be carried out in a broad range of potentials and concentrations of a dealloying solution, as

shown in Figure 7.18 and Figure 7.19. Therefore, stress-corrosion cracking observed in these

alloys may be caused, initiated or at least accompanied by the formation of nanoporosity

resulting from dealloying. This stays in contradiction to Galvele’s suggestions (31, 32).

Moreover, since the formation of 3D structures is possible in Ag-5 at. pct Au, it might be

possible that other silver-gold alloys of lower gold concentration than 5 %, could also be

successfully dealloyed.

The thickness of a nanoporous layer formed depends on the charge passed (Figures 7.4 to 7.8).

The composition of nanoporous layers, determined with EDX, revealed surprisingly large

concentration of the residual silver in ligaments, as shown in Figure 7.3. The observed

thickness of nanoporous layers is smaller than it could be expected from the calculations

(Figure 7.9). The difference may result from the collapse of the layer but it also seems to be

possible that a part of charge passed may be consumed in other dealloying processes that are

not taken into account, for instance, dealloying may penetrate the bulk material in narrow

channels via grain boundaries.

During dealloying nanoporous layers formed on Ag-5 at. pct Au undergo physical changes

leading to coarsening (Figure 7.24 and Figure 7.25). Precise estimation of the degree of

coarsening may be difficult since, after dealloying, samples were transferred several times from

one environment to another during the preparations (DI water, and ethanol etc). It should also

66

be noted that SEM imaging requires samples to reside in vacuum where surface diffusion may

be different to those in liquid solutions and air.

The influence of oxygen dissolved in a dealloying solution (Figure 7.20) is negligible and does

not affect the appearance of the empirical critical potential of Ag-5 at. pct Au in 0.1M HClO4.

The concentration of a dealloying solution, however, does affect the anodic current which

increases along with the increase in the concentration of a solution (Figure 7.21). The empirical

critical potential can be determined for Ag-5 at. pct Au (Figure 7.18 and Figure 7.19) and

changes along with the concentration of the more noble metal, as reported in the past (2, 15).

Although the conditions under which film-induced cleavage occurs have not yet been found, a

great deal of work was done on studying grain boundary corrosion and, specifically, on finding

the answer how to eliminate or limit this process as it was found to be significantly interfering.

In order to do it, two different ways of fracturing samples have been examined, that is,

fracturing in the dealloying solution (at positive or negative potential), and fracturing beyond

the dealloying solution (DI water and air). No relation has been found between the crack depth

and the ways in which Ag-5 at. pct Au was fractured. The depth of a crack depended only on

the charge passed. Since all the samples were fractured manually, the rate of the applied strain

(approximately t = 0.5 s) could be too low and will have to be reconsidered in further single-

shot fracture experiments.

The presence of premature cracks on the surface has been confirmed (Figure 7.22). Premature

cracks may also negatively affect the generation of film-induced cleavage. Their formation and

influence will further be investigated.

There has been observed an excellent reproducibility of all potentiodynamic and potentiostatic

experiments performed on Ag-5 at. pct Au in the range of potentials of interest.

67

The fact that silver rich Ag-Au alloys undergo dealloying will have a significant impact on the

way nanoporous layer may be fabricated and used, for instance, as cheaper catalysts than

nanoporous gold fabricated from gold rich Ag-Au alloys.

9 Further work

Since experiments on polycrystalline materials inherently involve grain boundaries, further

investigation of film-induced cleavage in silver-gold systems will be divided into two different

parts during which one type of film-induced cleavage will be investigated at a time: 1)

transgranular film-induced cleavage (TG-FIC), 2) intergranular film-induced cleavage (IG-

FIC). In order to study TG-FIC, further single-shot fracture tests will be carried out on

monocrystals or samples of sufficiently large grains because it will eliminate grain boundaries.

IG-FIC will be studied on a given grain boundary using bicrystals. Before fracture tests, EBSD

will be performed in order to determine the type of a grain boundary. Both monocrystals and

bicrystals will be produced by the application of small amounts of cold work (up to 5%) and

long annealing at 900oC in argon plus 2.5% H2. These conditions have been already checked

and proved to produce excellent results. Prior to any of the above experiments, for a given

alloy, the highest dealloying potential causing no premature cracks will be determined. Because

the generation of film-induced cleavage requires a nanoporous layer to be sufficiently thick and

formed as quickly as possible, new conditions for a dealloying solution will be employed such

as higher concentration of perchloric acid and elevated temperatures.

68

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and Interpretation, 1188, 94 (1993).

35. N. A. Senior and R. C. Newman, Nanotechnology, 17, 2311 (2006).

36. R. C. Newman and N. A. Senior, Corrosion Science, 52, 1541 (2010).

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11 Appendix

Figure 11.1 shows the reproducibility of potentiodynamic curves performed on Ag-5 at. pct Au

in 0.1 M and 1.0 M HClO4 at room temperature. Compared with the potentiodynamic curves

presented in Figure 7.21, the below curves were obtained by potentiodynamic polarization of

the samples exposed to the dealloying solution from each side. This was done on purpose to

confirm that the changes in current density will be the same, regardless of the way samples are

exposed, that is, both sides or one side only.

Figure 11.1. Polarization curves performed on Ag95 Au5 in 0.1 M and 1.0 M HClO4 at room


Figure 11.2 and Figure 11.3 show the reproducibility of the potentiostatic experiments

performed on Ag-5 at. pct Au at E = 400 mV and E = 300 mV in 0.1M HClO4 for the samples

exposed from both sides and one side only.

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consecutive steps at different potentials in 0.1 M HClO4 at room temperature. One side of the

sample exposed to the dealloying solution.


consecutive steps at different potentials in 0.1 M HClO4 at room temperature. Both sides of the

sample exposed to the dealloying solution.

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The below image demonstrates the lab-made device designed by the present student. It was

used in all single-shot fracture experiments. The device consists of two main parts:

electrochemical cell and a mobile part. The cell was an adapted glasslock container normally

used for storage of sandwiches. Four holes were drilled in the lid of the container, one for the

mobile part and the three smaller ones for reference electrode (kept in a double junction),

counter electrode and the glass tube supplying nitrogen. The lid is separated from the glass with

a gasket. The gasket seals off the cell very well and prevents nitrogen escaping during

deaeration of a dealloying solution. The whole provided a very convenient and quick way of

loading samples at will and immediate start of an experiment without any delay. It was

especially required during experiments at elevated temperatures. The copper wire sticking out

from the mobile part connects the sample with the potentiostat.

Figure 11.4. Fracturing device: A – electrochemical cell, B – mobile part made of syringes.

The mobile part consists of two syringes of different sizes (Figure 11.5). The main part was

made of a largish 60 mL syringe. Holes were drilled in the syringe tube in order to provide an

undisturbed exchange of a dealloying solution. The hole seen in the piston has the same

diameter as the diameter of the little syringe. A piece of wood led through the syringe piston

A

B

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blocks the piston against moving during loading the mobile part into the container. It also

protects the sample against bending or any accidental application of cold work.

Figure 11.5. Mobile part: C - the syringe tube, D – piston, E – small syringe with the sample

mounted in.

The set of figures below shows changes in the imaginary part of impedance versus frequency

for the potentiostatic experiments performed on Ag-5 at. pct Au and Ag-23 at. pct Au.

Figure 11.6. Changes in the imaginary part of impedance after dealloying in ten consecutive

steps at E = 120 mV for Ag95 Au5 in 0.1M HClO4. Charge passed in each step Q = 600 mC/cm2.

C D E

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75





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steps at E = 550 mV for Ag77 Au23 in 0.1M HClO4. Charge passed in each step Q = 0.6 C/cm2.


steps at E = 650 mV for Ag77 Au23 in 0.1M HClO4. Charge passed in each step Q = 0.6 C/cm2.

Dealloying and Formation of Nanoporosity in Noble-Metal Alloys · ii Dealloying and Formation of...

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