Copper Distributions in Aluminium Alloys

COPPER DISTRIBUTIONS IN ALUMINIUM ALLOYS

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COPPER DISTRIBUTIONS IN ALUMINIUM ALLOYS

T. H. MUSTER A. E. HUGHES AND

G. E. THOMPSON

Nova Science Publishers, Inc. New York

Copyright © 2009 by Nova Science Publishers, Inc.

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CONTENTS Preface

vii

Chapter 1

Introduction

1

Chapcter 2

Alloy Manufacture

5

Chapter 3

Alloy Microstructure

9

Chapter 4

Electrochemistry

25

Chapter 5

Corrosion

31

Chapter 6

Chemically Pretreated Surfaces

49

Conclusions

83

Acknowledgements

85

References

87

Index

99

PREFACE Aluminium alloys are used extensively throughout the world, in items such as decorative architectural applications through fasteners to high strength structural applications. Such a diverse range of application areas has a similarly diverse range of requirements for materials properties and performance. The mechanical properties are achieved through alloying aluminium with a wide range of elements. Copper, which is one of the major alloying additions, is added in varying amounts to many of the different aluminium alloy series, with the lowest levels in the purest wrought aluminium alloys (AA1xxx series) and the highest levels in the high strength AA2xxx series. The distribution of copper in aluminium alloys varies from copper atoms dispersed in solid solution through the formation of clusters of copper atoms and then onto to a range of intermetallic compositions and particle sizes. The presence of copper in all these forms, particularly in the AA2xxx series, has a significant impact on the chemistry and electrochemistry of the surface of the alloy and, hence, on the susceptibility to corrosion and approaches to metal finishing. This chapter examines how the copper distributions change, as a result of corrosion reactions, and explores the influence of these changes on continued corrosion. The influence of the distribution of copper in aluminium alloys on metal finishing processes and the redistribution of copper as a result of metal finishing is also examined.

Chapter 1

INTRODUCTION Aluminium alloys were first developed for commercial use in the mid 19th century in France [Polmear (1989)]. Since that time there has been considerable alloy development to produce the vast range of cast and wrought alloys that are available today. Aluminium alloys are used extensively throughout many industries. For example, an examination of the categories of the Aluminium Surface Science and Technology Conference proceedings from Bonn 2003 indicates applications in architecture, packaging , electronics, transport, lithography, capacitor foils and heat exchangers [ASST proceedings (2004)]. Other sources indicate the extensive usage of aluminium alloys in transport (30%), packaging (18%), building and construction (21%), mechanical engineering (8%), electrical engineering (9%), household articles (8%) with 8% assigned to miscellaneous uses [Riotinto website]. The high strength to weight ratio of alloyed aluminium makes it an excellent candidate for the transport industry where it is used in train, automotive, shipping and the aircraft industry sectors. A large number of alloys have been developed over the years to meet the requirements of different application areas, and new alloys and heat treatments continue to be developed as new needs arise. One example is in the automotive industry, where light weight aluminium casting alloys are replacing ferrous-based materials for engine heads and blocks. Specific alloys vary from producer to producer, but generally heads are made from AA319 and its variants, and blocks are beginning to be made from AA380 and its close variants. A further example is in the aerospace industry where Aluminium-Lithium-Copper alloys were developed to replace high strength Aluminium-Copper alloys, which have been used for nearly eighty years in aircraft manufacture [Bovard (2006)] and possibly

2

T. H. Muster, A. E. Hughes and G. E. Thompson

as long as one hundred years if airships and aluminium alloys of engines in the first aircraft are included. Some of these developments are successful whereas others are not immediately taken up. Aluminium-Lithium-Copper alloys, for example, have been available commercially since the 1950’s [Polmear (1989)], but there has been an ongoing reluctance to take up these alloys until there is an improvement in corrosion performance and fracture toughness [Bovard (2006)]. However, there are exceptions, and the alloy AA1420 (Al-5Mg-2Li-0.5Mn) has a high corrosion resistance and has been used successfully on at least one advanced military aircraft produced in the former Soviet Union [Polmear (1989)] The aircraft industry uses medium to high strength AA2xxx and AA7xxx alloys which typically have higher copper contents up to approximately 6 and 3 % respectively. Recent trends in the aerospace industry are aiming at increased operational lifetimes [Schmitt (1998), Brown (1992)] with corrosion issues associated with the airframe becoming a high priority. Older AA2xxx series alloys such as AA2024-T3 alloy and AA7xxx series alloys such as AA7075-T6 are the, so-called, legacy alloys [Bovard (2006)]. The corrosion issues related to an airframe increase markedly with operational lifetime beyond around twenty five years. Table 1 lists some of the legacy alloys [Bovard (2006)] which, not too surprisingly, are also some of the most studied alloys in terms of corrosion and metal finishing. AA2024-T3 is one of the most corrosion prone alloys because of the high levels of copper, and AA7075-T6 is particularly prone to intergranular attack [Davis (1999), Hatch (1984)]. The relationship of temper to corrosion performance should not be overlooked since a change in heat treatment can cause a significant change in the distribution of alloying compounds and intermetallics. For example, AA2024-T4 has a very high susceptibility to various forms of corrosion attack, which can change from intergranular to pitting by changing quench conditions (which influences the amount of precipitation of solute from solid solution [Hatch (1984)]. Similarly problems with stress corrosion in AA7xxx-series alloys can be overcome by heat treating to an over-aged condition (T7x, where a lower value of x means a greater degree of over-aging). Table 1. Legacy Alloys UNS Number AA2024-T3 AA7075-T6 AA7075-T73

Introduction 1935 1945 1960

Aircraft Douglas DC3 Boeing B29 Douglas DC9

Introduction

3

New alloy design has led to the development of alloys that supercede the more corrosion prone alloys. Some of the new alloys include AA7x5x-T77 alloy which has been used on the Grumman A-6, AA7058 on the Boeing 777, and AA7085 on the Airbus A380 [Bovard (2006)]. Improvements in temper have also led to improved resistance to intergranular attack in the AA7xxx alloys, as the T77 condition suffers only a minimal reduction in mechanical properties compared to the T6 temper. In the AA2xxx series, AA2190-T8 alloy has been introduced into the fuselage sheet. Such a broad range of applications comes with an equally broad range of service conditions which, in turn, require a wide range of approaches for aluminium finishing. Some common finishing processes include anodizing and conversion coating for corrosion protection, adhesive bonding and painting finishes varying from architectural facades to adhesive bonding in aircraft applications, to coating for casings for electronic hardware. [Laevers et al. (1993), Arai et al. (1984), Dunn et al. (1971)]. This chapter considers the influence of copper in corrosion and metal finishing of aluminium alloys. Copper takes a special place in the role of both corrosion of aluminium alloys and its metal finishing, because of its electrochemical properties. Copper is one of the most noble alloying elements used in aluminium alloy manufacture; hence, it exhibits distinctly different electrochemical characteristics from the aluminium matrix. Copper-containing phases on the surface tend to be cathodically protected since they exhibit a net cathodic reduction reaction: O2 + 2H2O + 4e- → 4OH- (neutral media)

...1

O2 + 4H+ + 4e- → 2H2O (acidic media)

...2

2H+ + 2e- → H2 ↑ (acidic media)

...3

whilst the aluminium matrix has a net anodic reaction which leads anodic dissolution via: Al → Al3+ + 3e-

…4

The presence of Cu accelerates the rate of surface electrochemical reactions in equations 1 to 4. Copper also tends to accumulate on the surface during both corrosion reactions, and in metal finishing, which further complicates subsequent surface processing and reaction with the external environment.

Chapcter 2

ALLOY MANUFACTURE In order to address the materials requirements of the broad range of applications where aluminium is used, it is alloyed with a range of different elements. On the basis of the predominant alloying metal, aluminium alloys have been divided into different series, designated AA1xxx series to the AA8xxx series. The alloy designations for wrought alloys is defined by a four-number system known as the International Alloy Designation System (IADS) [Starke and Staley (1996)]. The first number designates the major alloying addition (Table 2), the second number refers to modifications of the original alloy or to impurity levels, and the last two numbers indicate the specific alloy. Table 2 lists the wrought aluminium alloys according to their series and the predominant alloying elements with an example of a common alloy for each series and the concentration range of copper addition for the designated alloy. While the last two numbers may specify the individual alloy, the designation simply gives the compositional bounds for the alloy, and individual alloys made under the same alloy designation and the same temper can have slightly different compositions. The microstructure of the alloy is, in part, determined by the heat treatment that the alloy undergoes. Most heat treatments involve solution treating, where the alloy is taken to a sufficiently high temperature to generate the equilibrium single, solid phase for the alloy concerned but below the solidus temperature. For example, the solution treatment temperature in the Al-Zn-Mg-Cu system varies widely with zinc content and temperature of incipient melting. The complexity of the phase diagram as a function of zinc and magnesium content for this system, at the solution treatment temperature, is evident in Figure 1. When the alloy is quenched, the solute elements are retained in the single phase aluminium as a supersaturated solid solution, with the exception of insoluble

6

T. H. Muster, A. E. Hughes and G. E. Thompson

intermetallic particles of phases such as those provided in Figure 1. The solid solution then decomposes during the ageing treatment to produce precipitation hardening. Large constituent phase such as η (MgZn2) can also form on grain boundaries during the ageing process. Similarly the Al-Cu-Mg system is complex and solution treatment needs careful control since it has to be performed a few degrees below the solidus temperature. Table 2. Series Designation for Wrought Aluminium Alloys Series 1xxx

Alloying Metal -

Example Alloy AA1100 AA1200

2xxx 3xxx

Cu, Mg Mn

5xxx

Mg

AA2024 AA3003 AA3104 AA5005 AA5052

6xxx

7xxx 8xxx 1

Si, Mg

Zn. Mg, Cu Other Including Li and Ni

AA6060/

Common Usage

Cu (wt%)

aluminium foil food handling / packaging containers high strength beverage cans beverage cans structural alloys / automotive and decorative automotive structural alloys /general purpose/ automotive

0.05 – 0.201

Heat Treatable No

0.052

No

3.8 – 4.91 0.05 – 0.201 0.152

Yes No No

< 0.22

No

AA6061 AA7075

high strength

AA8081

automotive

< 0.12 < 0.12

Yes

0.15 – 0.41 1.2 – 2.01

Yes Yes

0.7 - 1.31

No

Polmear (1989), 2 Kaufman (2004).

Figure 2 shows the ternary phase diagram for the Al-Cu-Mg system. AA2024-T3 typically contains 1.2 to 1.8 wt% magnesium and 3.8 to 4.9 wt% copper which means, if AA2024 alloy was a ternary alloy, S and θ phases should be present in the α-aluminium matrix. Commercial alloys have additional elements which change the composition of precipitates so that in AA2024-T3 alloy there is a range of Al-Cu-Mn-Fe containing intermetallics as well as S and θ phases.

7

Alloy Manufacture

10 α+T

8

α+S+T

%M g

6

α+T+M+S

4

α+S

2

α+S+M

α

2

4

8

6

α+M α+T+M

10

12

%Zn Figure 1. Section of the Al-Zn-Mg-Cu phase diagram (1.5% Copper at 460°C): S= Al2CuMg; T=Al6CuMg4 + Al32(MgZn)49; M= MgZn2 + AlCuMg (after Polmear, 1989).

3

α+θ+S

α+S

2

1

% Cu

α+S

α+ θ

α+S+T α+S

α 1

2

3

% Mg Figure 2. Section of ternary Al- Mg-Cu phase diagram (1.5%Copper at 460°C) S= Al2CuMg, θ =Al2Cu, α = solid solution (after Polmear, 1989). The shaded section of the diagram is at 460°C whereas the other phases are for 190°C.

8

T. H. Muster, A. E. Hughes and G. E. Thompson Table 3. Temper Designations

Alloy Type Non Heat Treatable

Heat Treatable

Basic Treatment

Secondary Treatment

H

1

only cold worked

“ “

2 3

T

1

“

2

“

3

“

4

“

5

“

6

“

7

“

8

“

9

cold worked and partly annealed cold worked and stabilised partial solution treatment followed by natural ageing Annealed cast products only partial solution treatment followed by natural ageing solution treatment followed by natural ageing artificially aged solution treatment followed by artificial ageing solution treatment and stabilisation solution treatment, cold work, followed by artificial ageing solution treatment followed by artificial ageing and cold worked

Description of Secondary Treatment

Solution treatment is followed by quenching, usually into cold water, to achieve the maximum supersaturation of the alloying components. Quenching does, however, introduce residual stresses into the product (particularly sheet product) which can be reduced by stretching or roller levelling. From this point, wrought aluminium alloys can be divided into heat treatable alloys, which improve their mechanical properties with subsequent heat treatments, and nonheat treatable alloys, which develop their strength through strain hardening. This latter group includes the Al-Mg, Al-Mn and Al-Mg-Mn alloys. The mechanical properties of heat-treatable wrought alloys can be manipulated using artificial aging treatments, as described above. Artificial aging can also influence the corrosion performance of a particular alloy, and its surface response to metal finishing treatments. The temper designation for both classes of alloys are listed in Table 3. Slow cooling can also be used, but this depends on the individual alloy since there may be a propensity for the formation of large undesirable intermetallic particles.

Chapter 3

ALLOY MICROSTRUCTURE BULK MICROSTRUCTURE To understand how copper promotes corrosion in aluminium alloys and why it accumulates at the surface of alloys during metal finishing processes it is instructive to examine the microstructure of aluminium alloys, particularly those with higher copper contents. Traditionally, the “microstructure” has, by default, referred to the bulk microstructure, which is the focus of this section, but equally important is the surface microstructure (see following section), since this is the interface where reactions of relevance to metal finishing or corrosion commence. As noted previously, there are two broad classifications for wrought aluminium alloys; non-heat treatable and heat treatable [Hatch, (1984)]. Non-heat treatable alloys, which obtain most of their strength through solid solution hardening and strain hardening, contain major additions of chromium, iron, magnesium, manganese, silicon and zinc, whilst only minor additions of copper are permitted (i.e. 0.12 – 0.15 wt% in can stock alloys AA3003 and AA3104, AA1100, and 1 wt% in AA8280 and AA8081 alloys). Heat treatable aluminium alloys can contain increased levels of copper (up to 6.3 wt%) [Hatch (1984)]. Ultimately, these alloying elements are present in either solid solution in the matrix, intermetallic particles, or both. Copper, together with magnesium, zinc and silicon, are appreciably soluble at high temperature and considerably less soluble at low temperature. This results in the precipitation of various phases during solidification of the alloy [Hatch (1984)]. For the precipitate-hardened alloy, the mechanical properties are improved by the precipitation of alloying components from solid solution. These fine precipitates often start as clustering of alloying components called Guinier-

10

T. H. Muster, A. E. Hughes and G. E. Thompson

Preston (GP) zones that grow into a range of precipitates with increasing temperature and time. Initially, these precipitates are coherent with the aluminium lattice, which is desirable, but continued ageing will take the precipitates through degrees of coherency until they are so large that they become incoherent with the lattice. These phases are usually identified using greek letters and the degree of coherency is identified using “primed (‘)” symbols. For example, in the case of Al-Cu binary alloys where the θ-phase forms (Al2Cu), the stages are GP → θ” → θ’ → θ with each step indicating an incremental loss of coherency with the primary matrix. Table 4 indicates some of the typical hardening precipitates. A second class of precipitates are the dispersoids that form by solid state reaction during preheating of the ingot; they represent an important part of the alloy microstructure since they control grain growth. They are formed through precipitation with either chromium, manganese, titanium or zirconium and form dispersoids particles such as Al12CrMg2, Al20Cu2Mn3, Al12Mn3Si, Al3Ti and Al3Zr. These particles are usually a few nanometers up to 200 nm in size [Starke and Staley (1996)]. Table 4. Compositions of Hardening Precipitates Alloy Type Al-Cu Al-Mg Al-Si Al-Cu-Mg Al-Mg-Si Al-Zn-Mg Al-Li-Mg Al-Li-Cu

Precipitate Composition Al2Cu Al8Mg5 Si Al2CuMg Mg2Si MgZn2 (Al,(Cu,Zn))49Mg32 Al3Li Al2LiMg Al3Li Al2CuLi

Precipitate Nomenclature θ β S β η T δ δ T1

A third class of particles are the constituent particles, most of which are formed during solidification of the initial ingot. As stated above, at low levels, copper is present in solid solution in the matrix (α-Al). The AA2xxx, AA7xxx and AA8xxx series alloys also have copper in solid solution as well as being incorporated into a range of intermetallic phases called constituent particles. As noted previously, the microstructure of these alloys is complex and depends on thermal and ageing treatments. Common constituent particles and the alloys they appear in are listed in Table 5 [Stake and Staley (1996)].

Alloy Microstructure

11

Table 5. Some Typical Constituent Particles found in Wrought Al-Alloys Alloy 2X24 2X19 6013 7X75 7X50 7055 2090 2091 2095 8090

Constituent Particles Al7Cu2Fe, Al12(Fe,Mn)3Si, Al2CuMg Al2Cu , Al6(Cu,Fe) Al7Cu2Fe, Al12(Fe,Mn)3Si, Al2Cu) Al12(Fe,Mn)3Si Al7Cu2Fe, Al6(Fe,Mn), Al12(Fe,Mn)3Si, Mg2Si Al7Cu2Fe, Al2CuMg, Mg2Si Al7Cu2Fe, Mg2Si Al7Cu2Fe Al7Cu2Fe, Al3Fe, Al12Fe3Si Al7Cu2Fe , Al2CuLi, Al6CuLi3 Al3Fe

A simple summary of the microstructure of aluminium alloys is not possible since it varies considerably from series to series, alloy to alloy and even from temper to temper. General summaries of the alloy microstructure has been given elsewhere [Hatch (1984), Vander Voort (2004)] and are not repeated here except for consideration of the distribution of copper in the respective alloys.

AA1xxx Series Alloys The AA1xxx series include high and super purity aluminium and generally only contains impurity levels up to 1% of iron and silicon as major impurities. The types of intermetallic particles that form include Al3Fe and silicon particles. AA1xxx series alloys are commonly used as cladding for AA2024-T3 alloy in the aircraft industry to provide protection for the more corrosion prone AA2024-T3 alloy.

AA2xxx Series Alloys The AA2xxx series, containing copper and magnesium, are high strength aluminium alloys and are therefore often used in applications which require such strength i.e., aircraft manufacture. They have a high damage tolerance and fatigue resistance. The mechanical properties of the AA2xxx series alloys are determined by the ternary Al-Cu-Mg phase diagram. In Al-Cu-Mg ternary systems that fall in the α + S phase region (Figure 2), the precipitation of Al-Cu-Mg particles occurs in the

12

T. H. Muster, A. E. Hughes and G. E. Thompson

following sequence; Copper in solid solution with aluminium (α-Al), first precipitates out as clusters of solute elements before GP zones form as fine rods in the <001>α directions. Much later in the precipitation sequence, or more usually, where cold work is applied, S-phase precipitates form as laths within the microstructure on {012}α planes in <001>α directions. The development of these types of phases for AA2024-T3 alloy can be followed using Cu63 NMR as shown in Figure 2 (right), where copper in solid solution is clearly distinguished from the precipitated forms of copper such as the S-phase [Bastow (2003, 2005, 2006), Nairn (2006)]. The development of different forms of precipitate can also be followed, and quantified, as shown in Figure 2 (right) for the binary Al - 4 wt% Cu. S

α

GPZ α

θ''

4000

3000

2000 ppm

1000

6000

4000

2000 ppm

0

Figure 2. 63Cu NMR spectra. Left - of Al-4 wt% Cu solution treated, quenched and aged at 130oC for two hours showing the development of θ’-precipitates within the alloy. Right AA2024 solution treated and quenched and aged at 177oC for two hours showing the development of S-phase (Courtesy of T. Bastow).

Equally, electron microscopy can be used to follow the formation of precipitates in aluminium alloys. θ’ precipitates, in the quaternary Al-Cu-Mg-Ag alloy aged to the T6 condition, can be seen, viewed edge-on, along the [001]α plane in Figure 3(a), using transmission electron microscopy. The dominant Ω phase provides the mottled appearance, since it is inclined to the plane in which the image of the θ’ precipitates has been taken. The backscattered image in Figure 3(b) shows a triple point junction between three grains in an Al-Cu binary alloy.

Alloy Microstructure

13

The θ’ precipitates can be seen in the individual grains as well as the larger θ precipitates in the grain boundaries. The precipitation of the θ precipitates within the grain boundary has led to copper depletion in the vicinity of the grain boundary. As will be shown later these small changes in copper distribution have significance for corrosion performance. 100 nm

(a)

200 nm

(b)

Figure 3. (a) Transmission micrograph of Al-Cu-Mg-Ag alloy aged to the T6 condition showing θ’ precipitates as plates viewed edge on. A small amount of the minor phase (Σ) is observed as cuboids, and the dominant Ω-phase is present on [111]α incline planes. Imaged along [001]α (courtesy of Dr. R Lumley). (b) Backscattered scanning electron micrograph of a Al-Cu binary alloy showing θ’ precipitates in the matrix and θ precipitates in Cudepleted grain boundaries (provided by Professor G. Thompson).

In corrosion and metal finishing, the most widely studied, copper-containing alloy in the AA2xxx series alloys is sheet AA2024-T3, although AA2014-T3 alloy is becoming more prominent in aircraft manufacture, and therefore a subject of research, [Starke and Staley (1996), Henon (2006) Smith (1993)]. For AA2024T3 sheet, total constituent particle number densities have been reported from 300,000/cm2 [Chen et al. (1997)] to 530,000/cm2 [Juffs (2003), Hughes et al. (2006)] for polished surfaces and as high as 11,700,000/cm2 for the rolled surface [Juffs (2003), Hughes et al. (2006)]. However, for rolled and polished surfaces, the surface area occupied by intermetallic particles was similar, suggesting that rolling leads to significant breakup of intermetallic particles. This is also reflected in the average particle size which was much smaller for the rolled surface than the polished surfaces. Particle size distributions for the intermetallic particles in AA2024-T3 have been reported by Jakab et al. (2005) and Hughes et al. (2006) for polished surfaces, with slightly different distributions revealed, but with similar volume fractions of intermetallic particles (Table 5).

14

T. H. Muster, A. E. Hughes and G. E. Thompson Table 5. Intermetallic Particles Distributions in AA2024-T3 Alloy Average Particle Size (μm)

Source

% Surface Area

Polished Hughes et al. [2006), Juffs (2003) Jakab et al. (2005)

6.6 3.1

2.89 2.18

Rolled Hughes et al. [2006), Juffs (2003)

2.0

2.82

It is not clear whether the difference in the particle number density between 300,000 and 530,000/cm2 represents a significant variation. Certainly there will be batch variation and probable processing effects; further the particle population densities will depend on the resolution of the techniques used for the counting statistics. Another possibility for the differences in the published figures is the processing history of the alloy. Specifically, for sheet alloy, the gauge (or thickness) reflects the number of rolling passes that the alloy undergoes. Clearly, at each pass the potential exists for further breakup of intermetallic particles and changes in the intermetallic size and spatial distributions and grain refinement (see section on Surface Microstructure). Examination of cross sections of AA2024-T3 alloy revealed that the distribution of intermetallic particle density across the sheet can change significantly, as depicted in Figure 4 [Juffs (2003)]. The variation in particle density is accompanied by an increase in particle size towards the centre of the sheet; this is reflected in the larger particle size on the polished surfaces (toward the sheet centre) versus the rolled surface. The characteristics of intermetallic particle distributions is an area which warrants further investigation since second phase particles are often sites of corrosion initiation and the influence of clustering of these particles is largely unknown. Focusing on the larger intermetallic particles, Buchheit et al. (1997) reported that roughly 60% of the constituent particles of particle diameter exceeding 0.2 μm were Al2CuMg (S-phase). The remaining 40% of intermetallics comprised a range of Al-Cu-Fe-Mn containing phases. The composition of Al-Cu-Fe-Mn phases has been suggested to take various forms. Gao et al. (1998) suggest compositions based upon (Al,Cu)x(Fe,Mn)ySi such as modified forms of Al8Fe2Si or Al10Fe2Si type intermetallics, although in low silicon–containing AA2xxx series alloys, these compositions are different. For example, Buchheit et al. (1997) reported that of the remaining 40% of intermetallic particles, the most notable included Al7CuFe2, Al6MnFe2, (Al,Cu)6Mn, and a number of undetermined compositions in the class Al6(Cu,Fe,Mn) where the Cu:Fe:Mn ratios were

15

Alloy Microstructure

approximately 2:1:1. The Al-Cu-Fe-Mn particles consistently exhibit crosssectional diameters in the range 10 to 50 μm, possess a high hardness, and are generally irregular in shape [Liao and Wei (1999)]. Further Scholes et al. (2006) reported that this class of intermetallic particles underwent fracture during milling, whereas the S-phase particles remained largely intact.

100

Particle Count

90 80 70 60 50 2

4

6

8

10

12

14

16

Position Across Sheet Figure 4. Intermetallic Particle Count taken on frames across a section of AA2024-T3 with a thickness of 1.2 mm. The sample was mounted in bakelite and polished down to 1 μm (after Juffs, 2003).

There is emerging interest in the spatial relationship of intermetallic particles in surfaces. In an extensive study of clustering, Juffs (2003) examined several methods of determining clustering of intermetallic particles in aluminium alloys. One of the most sensitive methods was the pair correlation function in which the average number of nearest neighbours is determined as a function of distance from the average particle. Juffs (2003) observed clustering in AA2024-T3 alloy for both polished and rolled surfaces; indeed the number of nearest neighbours was more than double that expected on a polished surface with a random distribution of intermetallic particles, and a little under twice as many for the rolled surface. On the other hand Jakab et al. (2005) found no significant clustering in AA2024-T3 alloy. The statistical sampling between the two studies may explain the differences. In the former study, Juffs (2003) counted several thousand particles, whereas, Jakab et al. (2005) did not indicate the number of

16

T. H. Muster, A. E. Hughes and G. E. Thompson

particles counted, although it appeared to be less than one hundred. As will be further detailed in the section on corrosion, clustering may prove to be an important issue for pit initiation and is a promising area for further research. At the submicron scale of the alloy microstructure, there is an even distribution of Al20Cu2Mn3 dispersoids. Guillaumin and Mankowski (1999) have reported that the coarse S-phase intermetallic particles are surrounded by a dispersoid free zone. However, Buchheit et al. (1997) suggested that only those particles that precipitate after secondary solution heat treatment will have a precipitate free zone surrounding them. At the finest scales there are lenticular particles around 100 nm in length which comprise the Al2CuMg hardening precipitates. More generally in AA2xxx alloys, the hardening precipitates of θphase (Al2Cu) and S-phase (Al2CuMg), depend on the copper to magnesium ratio [Hatch (1984)].

AA3xxx Series Alloys The AA3xxx series is formed by alloying with manganese at levels between 1 and 1.25 % and is dispersion-hardened through the presence of Mn-containing particles (Al6Mn). According to Hatch (1984) AA3003 alloy is the only AA3xxx series of commercial interest. The larger constituent particles can generally be divided into two classes which vary in their relative number according to the alloy composition (including impurities such as iron and silicon which are typically at levels of 0.7 and 0.3 to 0.6, respectively). These impurity levels are also higher than that for the many of the AA1xxx series alloys. Afseth et al. (2001), for example, reported that 60% of the constituent particles in AA3005 alloy were Al6(Fe,Mn) and the remaining particles were α- Al12(Fe,Mn)3Si. The AA3xxx series of alloys has gained some attention in recent years due to its enhanced filiform corrosion susceptibility which is largely related to the near-surface microstructure. This is discussed in the next section.

AA4xxx Series Alloys The AA4xxx series is formed by alloying with silicon at levels up to 13 wt% and may contain low levels of copper (typically 0.3%). These alloys are often used in brazing allplications [Hatch, 1984]. AA4047, for example is used as a cladding for AA3005 but according to Hatch (1984) has no copper. Because of impirty iron in the alloy β-AlFeSi can form as well as silicon particles.

Alloy Microstructure

17

AA5xxx Series Alloys These alloys, based on alloying with magnesium, have some of the best corrosion resistance of all aluminium alloys as well as relatively good strength [Polmer (1989), Vander Voort (2004)]. The alloys are often used for welding applications, and are usually manufactured as plate [Polmear (1989)]. Because of their corrosion resistance, they are often employed in marine manufacture such as small craft or ship superstuctures. Magnesium contents range from as little as 0.8% and up to 5% for wrought alloys. Apart from the typical impurity phases such as those related to silicon and iron, the main phases are Al3Mg2 and β-phase (Al8Mg5). With significant levels of either silicon, copper or zinc, hardening precipitates such as Mg2Si, Al2CuMg and Al2Mg3Zn3 can also form. Very small levels of other elements are added to AA5xxx series alloys, such as chromium which forms Al18Mg3Cr2 dispersoids. [Vander Voort (2004)].

AA6xxx Series Alloys The 6xxx series alloys are alloyed with both magnesium and silicon which are usually added in a ratio whereby Mg2Si form by precipitation from solid solution or silicon is added in excess. These alloys gain their strength by heat treatment and precipitation of the Mg2Si phase. The most common form of production of AA6xxx series alloys is as extrusions where Si is added between 0.8 and 1.2% [Polmear (1989)] and quenching immediately from the extrusion die means the alloy only require subsequent low temperature ageing (e.g. 180°C) to improve mechanical properties. As with other alloy classes, copper can be added ; the medium strength alloys AA6013, AA6056 and AA6111, for example, have up to about 1% copper. This is to enhance precipitation hardening, reported to be a result of Q-phase formation. Alloy AA6111 finds major use as an automotive bodysheet alloy, having a good combination of formability and strength. The AA6xxx series alloys containing copper, do, however, have inferior corrosion properties to copper-free AA6xxx series alloys.

AA7xxx Series Alloys The AA7xxx series alloys system is based on the ternary Al-Zn-Mg system but has copper included to alleviate severe problems with stress corrosion

18

T. H. Muster, A. E. Hughes and G. E. Thompson

cracking. The quaternary alloys also display an improved response to agehardening. Some AA7xxx series alloys are considered to be weldable and are used extensively in transport. These alloys typically have 3 to 7 % zinc and 0.8 to 3.0% magnesium. Chromium, manganese and zirconium are also added for grain size control. The two phases that form in wrought Al-Zn-Mg alloys, through eutectic decomposition, are Mg2Zn and Al2Mg3Zn3 [ASM (2004)]; the dispersoid phase is Al12CrMg2. The AA7xxx series alloys are well known for their susceptibility to stress corrosion, which is due to grain boundary precipitation of η-phase (Zn2Mg) and depletion of the adjacent grains [Pickens and Langan (1987)]; it will not be dealt with in this chapter. The addition of copper can reduce the susceptibility to stress corrosion cracking. The Al-Zn-Mg-Cu alloys possess some of the highest tensile strengths and other mechanical properties of all aluminium alloys due to their age hardening properties [Polmear, 1989]. Because of their excellent mechanical properties, the alloys have been used extensively in aircraft manufacture and AA7075 series, particularly in the T6 or T73 condition, have been used most extensively (Table 1). T73 is a duplex aging that creates an overaged microstructure. AA7050 alloy is a further important structural alloy and it has increased levels of copper to assist age hardening. Birbilis and Buchheit (2005) include the following intermetallics for AA7075 in the T6 condition: Mg2Si, MgZn2, Al12Mn3Si, Al7Cu2Fe, Al2Cu, Al2CuMg, Al3Fe, Al12CrMg2, Al20Cu2Mn3, Al6Mn, Al3Ti, Al6Zr, Al3Mg2, Al32Zn49, Mg(Al,Cu). (While the most common intermetallic particles are listed in Table 5, a more detailed analysis of the intermetallic particles in a complex alloy like AA7075-T6 will reveal a much more extensive list of particles.) Clustering of constituent intermetallic particles in 7075-T6 on the as-received surface was also studied by Juffs (2003) using the radial distribution function. The highest degree of clustering occurred in this alloy when compared to other alloys (including AA2024-T3), with nearly 4 neighbours expected within a radius of 5 μm.

SURFACE MICROSTRUCTURE The nature of surface microstructure has historically emerged from tribological studies [Fishkis and Lin (1997), Schey (1983)] and, in recent years, has increasingly been addressed as part of filiform corrosion susceptibility [Asfeth (2001), Leth-Olsen (1997), Mol et al. (2002)]. The surface microstructure is often more complex than the corresponding bulk microstructure. At the most

Alloy Microstructure

19

fundamental level, magnesium, lithium and silicon typically segregate to the external surface during heat treatment of both cast and wrought alloys. Textor and Amstutz (1994) reported that magnesium and, particularly lithium, are enriched in the surface oxide by 2-4 orders of magnitude compared to their bulk concentration. Internal segregation to grain boundary interfaces and the development of internal depletion zones was considered briefly in the previous section. Segregation to the external surface occurs via the two routes of (i) bulk diffusion and (ii) grain boundary segregation. In general, the surface enriched elements have a high negative free energy for oxide formation and high diffusion coefficient in aluminium metal [Textor and Amstutz, (1994)]. However, the tribology of forming processes such as rolling and extrusion add an extra and complex dimension to the nature of surface layers [Fishkis and Lyn (1997)]. The study of segregation phenomena in metallic alloys goes back many years and there have been extensive studies on the theory of segregation [Seah (1980), Hondros and Seah (1977), Du Plessis and Tagauer (1992), Luckman (1988), Darken (1967), Hofmann (1987), Du Plessis and van Wyk (1988) and Guttmann (1975)] These theories deal extensively with binary alloys, with some consideration of ternary alloys, but do not include the influence of phase precipitation or precipitation in stacking faults. Nevertheless, they provide a basic understanding of the kinetic and thermodynamic driving forces for segregation to surfaces. Carney et al. (1990), for example, suggested that magnesium enrichment in binary alloy powders is in accordance with the diffusion rate of magnesium in molten aluminium. However, mass loss due to evaporation at high temperatures is higher than the rate of accumulation due to segregation and magnesium depletion can occur. Lea and Molinari (1984) and Viswanadham et al. (1980) showed that the maximum segregation occurred at around 475°C. Even on polished surfaces, the surface oxide on Al-Cu-Fe-Mn intermetallic particles apparently forms by preferential oxidation of aluminium in the intermetallic phase [Roberts et al. (2002)]. In addition to heat treatments providing a driving force for segregation, mechanical processing can significantly change the surface microstructure of the alloy. In the last ten years there has been renewed interest in the reasons for increased filiform corrosion in north western Europe. The increase in filiform susceptibility was due to the incorporation of lower etch rate processes in the metal finishing of architectural alloys. The lower etch rates failed to remove the surface modified layer produced by rolling and related heat treatments, which may be rendered extremely electrochemically active. Multiple-pass rolling causes considerable modification to the surface of aluminium alloys creating new surface layers with a very fine grain structure and

20

T. H. Muster, A. E. Hughes and G. E. Thompson

incorporated oxide (Figure 5) [Fishkis and Lin (1997), Afseth et al. (2001)]. This type of transformation is called a Grain Refined Surface Layer (GRSL) [LethOlsen (1998)]. The surface layers are characterised by a high porosity, very fine grain structure and large oxide content. Oxides that have been detected in the surface include γ-Al2O3, MgO and the spinel phase MgAl2O4. The latter is only formed above 350°C [Fishkis and Lin (1997), Scamans and Butler (1975) Pronko et al. (1988), M. Pijolat et al., (1988), C. Lea and J. Ball (1984), S.K. Toh et al., (2003)], although Lumley et al. (1999) reported its formation at 275°C. The mechanism of modification proposed by Fishkis and Lin (1997) was a three-step process involving: 1. Formation of surface cavities by plowing, adhesive wear, delamination wear or transverse surface cracking, 2. Filling of the cavities with wear debris, including oxide, metal fines and lubricants, 3. Covering the cavities with thin metal layers during continuing rolling, leading to a shingled surface appearance.

Oxide Fragments Surface Oxide

GRSL Intermetallic Particles

Bulk Metal

Figure 5. Schematic diagram of the restructured layer as a result of mechanical work such as rolling including the incorporation of oxides particulates and a recrystallised zone. The recrystallised layer is called the Grain Refined Surface Layer (GRSL).

Alloy Microstructure

21

There are various reports on the depth of the modified surface region. Fishkis and Lin (1997) reported that the recrystallised surface layer changed from a depth of 8 μm after a first pass roll to 3 to 5 μm on subsequent rolling. Afseth et al. (2001) examined a rolled sheet of AA3005 alloy and revealed a deformation layer about 1μm thick. Both studies suggest that recrystallisation results in a very fine surface grain structure of dimensions down to 40 nm at the outer surface and up to 400 nm elsewhere in the recrystallised zone. Similar results were reported by Leth-Olsen et al. (1998) for AA8006, AA3005 and AA1xxx series alloy. Subsequent heat treatment of the AA3005 can result in precipitation of manganese-containing particles which renders the surface layer extremely surface active [Afseth (2001), Scamans et al. (2003)]. Milling also produces changes in surface structure as reported by Scholes et al. (2006), where crushing of Al-CuFe-Mn type intermetallics to a depth of around 4 μm from the surface of milled AA2024-T3 alloy was observed. There was also significant folding of the matrix alloy creating subsurface crevices up to 5μm long and a few microns depth. In both these instances, the GRSL is likely to be in a metastable state with respect to ageing and changes in the oxide composition. The surface oxide can change as a result of exposure to the environment. Viswanadham et al. (1980) studied changes to the magnesium enriched surface oxide of an Al - 5.5 wt% Zn2.5 wt% Mg alloy. They observed an increase in aluminium oxides on the surface on the magnesium oxide after storage in various moist environments at 50°C. They attributed this enrichment to aluminium diffusion from the underlying alloy through grain boundaries in the surface magnesium oxide onto the external surface. A further impact of rolling on the surface microstructure relates to the intermetallic particles phases. Lunder and Nisancioglu (1987) reported that during mechanical processing, i.e., rolling, the large intermetallics in the surface are covered with a layer of the matrix aluminium alloy. Rolling also has a mechanical impact by breaking up the intermetallic particles as shown in Table 5. Particle counting of the two main classes of intermetallic particles in AA2024-T3 (S-phase (Al2CuMg) and Al-Cu-Fe-Mn-containing intermetallics) indicated that the number density on the rolled surface was twice as high and the average size was about one third (Table 5), suggesting intermetallic breakup [Hughes et al. (2006)]. Milling was also observed to result in the fracture of intermetallic particles and covering of the intermetallics particles with the aluminium alloy matrix [Scholes et al., (2006)]. Although in this case it was reported that the S-Phase particles remained largely intact. From studies of a range of aluminium alloys, Lunder and Nisancioglu (1987) have also observed that the alloy matrix can cover the intermetallic particles in rolled surfaces.

22

(a)

T. H. Muster, A. E. Hughes and G. E. Thompson

(b)

Figure 6. EELS Maps from AA7475-T6 alloy: (a) aluminium (blue), magnesium (green) and oxygen (red); (b) copper (red), zinc (green) and iron (Blue). Color rules: R + G = Yellow, R + B = Magneta, G + B = Cyan, R + G + B = White.

(a)

(b)

Figure 7. EELS Maps from AA2024-T3: (a) oxygen/aluminium/copper and (b) oxygen/aluminium/magnesium. Color rules: R + G = Yellow, R + B = Magneta, G + B = Cyan, R + G + B = White. See text overlay on figure for colour assignments.

These studies indicate that, generally, while the surface microstructure and composition may be understood, individual treatments may result in considerable variation in the surface layers. Examples of the surface oxide are given in threecolour maps in Figure 6 for AA7475-T6 where it is revealed that the surface layer

Alloy Microstructure

23

varies considerably in thickness and incorporates the matrix metal (bottom left) folded into an oxide which varies in thickness from 250 to 500 nm. The threecolour map of the copper, zinc and iron elemental maps shows that ironcontaining particles are incorporated into the surface oxide. The chattering associated with mechanical damage due to ultramicrotomy, changes in moving from the bulk to the surface but these changes are deeper than the oxide coating and indicate the depth of the GRSL. Figure 7 shows a thin continuous layer for the surface oxide for AA2024-T3 alloy with considerable incorporation of magnesium oxide. Copper oxide particles are also present in the surface oxide layer. These particles are of similar size to the Al20Cu2Mn3 dispersoid phase, but appear to be copper oxide since neither aluminiumn nor manganese was detected in them; the origin of these particles is not clear. This example demonstrates the complex nature of the surface oxide. In summary, the surface of aluminium alloys may have a complex structure and composition that depends on the processing history and the storage environment. It is the objective of metal finishing processes to remove such layers to produce a surface which has a well defined, reproducible structure and chemistry, thus minimizing the history of the alloy on its subsequent performance during coating processing and its performance in-service.

Chapter 4

ELECTROCHEMISTRY This section focuses on the reactivity of aluminium alloy surfaces, and the influence of copper with regards to electrochemical phenomena, particularly corrosion. Aluminium is thermodynamically unstable, appearing highly negative on the electrochemical series (standard reduction potential = -1.42 VSCE). However, aluminium owes its, often exemplary, corrosion resistance to its ability to form a passivating surface oxide layer. Under most conditions, this surface oxide is able to form and, if damaged, can easily reform [Davis (1999), Hatch (1984)]. The presence of the passivating film allows aluminium to achieve potentials in the proximity of -0.75 VSCE in aqueous solution (Table 6). The oxide formed at the surface of pure aluminium is composed of two layers. A thin (usually < 5 nm) compact amorphous passivating film is formed adjacent to the metal. This film has variously been reported as either γ-Al2O3 [Pryor, (1971)], γ’Al2O3 [Wilsdorf (1951)] or glassy [Fehlner and Mott, (1970)]. A thicker, hydrated, and porous oxide forms the outer surface layer. This thicker layer is generally designated as Al2O3.nH2O, and may be composed of aluminium hydroxides or oxyhydroxides, depending on the formation conditions [Alwitt (19740, Vedder and Vermilyea (1970), Davis (1999)]. When aluminium or its alloys are exposed to aqueous solutions, the protective oxide film may breakdown, resulting in corrosive attack of the underlying metal. At near-neutral pH the solubility of aluminium oxides is low (solubility constant <10-32); however, the stability of aluminium oxides is highly dependent upon the pH of the environment (Hemingway et al., 1991). Thermodynamically, aluminium oxides are stable between pH 4 – 9 (Pourbaix, 1966) and increasingly unstable at high and low pH. Copper oxides tend to be stable at high pH, and become increasingly soluble at low pH (Figure 8).

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T. H. Muster, A. E. Hughes and G. E. Thompson

Log (soluble species) [M]

5 Boehmite Al(OH)3

4

CuO Cu(OH)2

3

2

1

0 0

2

4

6

8

10

12

14

pH Figure 8. Solubility of Al2O3. H2O (Boehmite), amorphous Al(OH)3, CuO and hydrated CuO in distilled water as a function of pH (data from Pourbaix, (1966)).

In electrochemical systems, copper exhibits increased nobility in comparison with aluminium (standard reduction potential of +0.24 VSCE). The nobility of a metal is a direct reflection of its Gibbs free energy, which can be measured in terms of the work function (the energy required to remove an electron from the bulk of the metal into a vacuum). These are 4.17 ± 0.09 and 4.76 ± 0.23 eV for aluminium and copper, respectively, (Lide, 2001). When copper is alloyed with aluminium a common Fermi level is achieved, and the additional energy required to remove an electron from copper is maintained as a difference in the surface potential (in air) and a solution potential when immersed in an electrolyte (Muster and Hughes, 2006). As a result, the introduction of copper into aluminium alloys generally increases the solution potential in comparison with pure aluminium (Table 6). However, adding active metals, such as magnesium and zinc, tends to decrease the nobility of aluminium alloys. Figure 9 provides a summary of the effects of major copper and magnesium additions on the solution potentials of aluminium alloys. Up to a concentration of 4 wt%, the addition of copper to aluminium increases the solution potential by approximately 37 mV per percent copper. Further increasing the copper content (from 4 to 7 wt%) has limited effect on the solution potential (Figure 9, Table 6).

27

Electrochemistry

Potential (V)

- 0.70 Cu

- 0.78

- 0.86

Mg 2

4

6

Alloying element (wt %) Figure 9. Effects of copper and magnesium alloying elements on electrolytic solution potential of aluminium. Potentials (0.1N Calomel scale) are for high-purity binary alloys solution heat treated and quenched, then measured in a solution of 53 g/l NaCl plus 3 g/l H2O2, after Davis (1993).

Table 6. Solution potentials of relevant metals, alloys and secondary phases determined according to ASTM G69 (Davis, 1999) Metal Aluminium (99.999%) Copper (99.999%) Magnesium Zinc 2024-T3 6061-T6 7075-T6 7475-T7 8090-T3

Solution potential (VSCE) -0.75 +0.00 -1.64 -0.98 -0.60 -0.74 -0.74 -0.75 -0.70

% Cu

3.8 - 4.9 < 0.1 1.2 - 2.0 1.2 -1.9 1.0 -1.6

In addition to determining electrochemical potentials, the Gibbs free energy of alloying elements also controls the enrichment of elements at the alloy surface and in the surface oxides during corrosion processes [Skeldon et al. (1995), Zhou et al. (1999)]. Copper and other more noble elements (i.e. gold) have high Gibbs free energies of oxide formation per equivalent (ΔGº/n) relative to that of alumina

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T. H. Muster, A. E. Hughes and G. E. Thompson

and therefore show extensive enrichment at the metal/oxide interface. In contrast, magnesium, for example, has a lower ΔGº/n value than aluminium and, therefore, is more likely to appear in the oxide following corrosion processes [Zhou et al., 1999]. In instances where copper ions are released from the metal surface during electrochemical processes, the proportion of copper in the oxide, relative to its bulk alloy concentration, is decreased, and under most conditions oxidized copper is undetectable at the oxide/electrolyte interface [Zhou et al., 1999]. It follows that copper ions have an increased mobility through the surface oxides in comparison with aluminium ions, and that the Cu2+-O bond energy is lower than that of Al3+O. This allows copper ions to migrate more rapidly from the metal/oxide interface, through the oxide and into the electrolyte [Zhou et al., 1999]. Where copper is present at the surface of aluminium alloys, it shows vastly different behaviour to aluminium. The surface of copper is particularly efficient at supporting cathodic reactions (e.g., oxygen and water reduction). Limiting cathodic currents measured for pure copper and pure iron are reported to be in the vicinity of 1.5 mA cm-2, whereas limiting currents on pure aluminium are three orders of magnitude lower (0.5 – 1 μA cm-2) [Seegmiller et al., (2004)]. Note that the reference to iron is included since its electrochemical behaviour is somewhat complimentary to that of copper in aluminium alloys. In a stagnant and naturally aerated aqueous solution the oxygen reduction at the surface of copper is reported to occur at 20-30 μA cm-2 [Vukmirovic et al., 2003]. It has also been shown that copper oxide films do not significantly impede oxygen reduction, presumably due to their high electronic conductivity [Vukmirovic et al., 2003]. The high conductivity is related to the semiconducting nature of copper oxides, which is defined by the low bandgaps of 2.1 eV and 1.3 eV for Cu2O and CuO, respectively [Richardson et al., 2001]. As discussed earlier, the oxide that forms on most commercial alloys is predominantly a mixed aluminium and magnesium oxide (often with other alloying elements such as silicon) and is likely to possess a complex microstructure. The heterogeneities introduced by varying oxide coverages, particularly over intermetallic phases, are thought to lead to the onset of corrosion due to localised defects or poor oxide stability [Guillaumin et al., 2001; Vukmirovic et al., 2002]. Copper-containing aluminium alloys generally become more susceptible to corrosion with increasing copper concentration [Davis, 1999]. Copper containing alloys such as AA2xxx and AA7xxx series alloys are susceptible to pitting and intergranular corrosion, particularly in commonly encountered, chloride-containing environments [Leblanc and Frankel, 2002]. In addition to the altered stability of surface oxides, the precipitation of intermetallic phases results in a heterogeneous surface structure that exhibits areas of varying

Electrochemistry

29

solution potential in the presence of an electrolyte. This creates localized electrochemical cells between microstructural features at the surface. The electrochemical properties of the intermetallic phases found within aluminium alloys have been characterized by various authors, using bulk synthesized intermetallics [Buchheit, 1995 ; Ilevbare et al., 2004], microelectrode studies [Suter and Alkaire, 2001] and combinations of the two [Birbilis and Buchheit, 2005]. It is commonly accepted that all copper-containing intermetallics except S-phase are net cathodes in comparison with the surrounding aluminium matrix [Buchheit, 1995; Ilevbare et al., 2004]. S-phase generally possesses potentials more negative than -0.8 VSCE in a wide range of electrolytes and it is, at least initially, the site of preferred anodic reactions [Ilevbare et al., 2004]. Coppercontaining intermetallics have bulk electrochemical potentials that increase in nobility in the order: Al2CuMg < Al7Cu2Fe < Al6(Fe,Mn) < Al2Cu < Al20Cu2(Fe,Mn)3 [Birbilis and Buchheit, 2005; Ilevbare et al., 2004; Muster et al., 2007].

Figure 10. Net cathodic currnet densities for copper and intermetallic phases at the corrosion potential of AA7075-T651 alloy in naturally aerated 0.1 M NaCl (Data from Birbilis and Buchheit, 2005).

An understanding of the efficiency of individual intermetallics to function as cathodes can be obtained from the data of Birbilis and Buchheit (2005), who reported the average currents at bulk intermetallic phases at the corrosion potential of AA7075-T651 in 0.1 M NaCl (i.e. -965 mVSCE). Copper containing intermetallic phases support cathodic currents exceeding 310 μA cm-2 and 470 μA

30

T. H. Muster, A. E. Hughes and G. E. Thompson

cm-2 for Al7Cu2Fe and Al2Cu phases, respectively. Intermetallic phases containing manganese are shown in Figure 10 to support decreased cathodic reaction rates (less than 120 μA cm-2). Asfeth et al. (2002) demonstrated that the incorporation of increasing amounts of manganese into Al6(Fe,Mn) intermetallic structure improved the filiform corrosion resistance of AA3xxx alloys, a result that was attributed to solution potentials approaching that of the aluminium matrix. Schneider et al. (2004) showed that the incorporation of iron into bulk Al20Cu2(Fe,Mn)3 increased cathodic reaction efficiency to achieve similar levels to Al2Cu. They also showed the magnitude of cathodic currents determined for Al2Cu and Al20Cu2(Fe,Mn)3 to be three to ten times greater than those measured on AA2024-T3 alloy surfaces. Of the non-copper-containing intermetallic compounds present in aluminium alloys, iron-containing intermetallics, Al3Fe and α-Al12(Fe,Mn)3Si, associated with AA1xxx and AA3xxx series alloys, are known to act as cathodic sites [Asfeth et al., 2002 ; Davoodi et al., 2005]. Bulk Al3Fe has been shown to have a noble corrosion potential (similar to that of Al2Cu) and to support cathodic currents as high as 200 μA cm-2 [Park et al., 1999]. However, the superior ability of coppercontaining intermetallics to host cathodic reactions was noted by Birbilis and Buchheit (2005); they found that whilst Al3Fe was more noble than Al2Cu (and therefore exhibits the larger driving force for cathodic reactions), Al2Cu sustained higher cathodic rates than Al3Fe. Non-copper containing electrochemically active phases in aluminium alloys include Al32Zn49, MgZn2, Mg2Si, Al8Mg5 [Davis, 1999] and Al-Si-Mg particles in AA6xxx series [Guillaumin and Mankowski, 2000]. The previous considerations have revealed that intermetallic phases containing copper may be divided into two types:- (1) those with a cathodic potential with respect to the matrix and (2) the active Al2CuMg S-phase. The latter, as will become clear, has unique properties that influence aluminium alloy corrosion and surface finishing performance. The following sections detail the corrosion phenomena associated with copper present in (1) solid solution, (2) in cathodic intermetallic phases, and (3) S-phase intermetallics.

Chapter 5

CORROSION THE INFLUENCE OF COPPER IN SOLID SOLUTION Since copper has a more noble formation potential than aluminium, enrichment of copper can occur at the surface of the solid solution phase [Habazaki et al. (1997)]. As discussed in the following sections, the majority of the research concerning matrix enrichment has been carried out on aluminium alloys exposed to chemical polishing, electropolishing, alkaline etching and anodisinig. However, many of the findings appear to be systematic and are likely to hold for corrosion processes, whilst other behaviours appear to be dependent upon the electrochemical energy applied to the system [Jung et al. (1985)]. In a range of corrosive environments, aluminium has been shown to preferentially oxidise, resulting in the build-up of copper within a layer approximately 2 - 5 nm thick at the alloy surface [Jung et al. (1985), Habazaki et al. (1997)]. Strehblow et al. (1978) reported relatively thick enrichment zones (up to 65 nm) for anodized Al-Cu binary alloys; however, as stated by the authors, this apparent increased thickness of the copper-enriched zone may result from surface roughening. In environments where aluminium alloys continually experience anodic dissolution, then it has been suggested that even alloys with relatively little copper (0.06 to 26 at%) may display copper enrichment at the metal-oxide interface [Jung et al., (1985), Garcia-Vergara et al. (2004), Korovela et al. (1999), Blanc et al. (1997)]. Once a certain level of enrichment occurs, copper atoms (and other noble alloying elements) are thought to arrange themselves into clusters through surface diffusion and eventually protrude from the alloy surface with high curvature due to undermining of the surrounding aluminium matrix [Sieradzki (1993), Habazaki et al. (1997)]. The copper clusters may be released into the oxide layer by one of

32

T. H. Muster, A. E. Hughes and G. E. Thompson

two mechanisms: (1) they may be undermined and released as elemental copperrich nanoparticles (these particles are likely to have short lifetime due to their increased free energy, which increases significantly for particles with a small radius of curvature [Brinker and Scherer (1990)], or (2) copper ions may be oxidized directly from the protruding clusters, which is also promoted through increased surface curvature that moves the potential of copper to more negative values [Newman and Sieradzki (1994)]. Based upon experimental evidence, the surface of all aluminium alloys containing trace copper upwards will become enriched in copper over sufficient periods of time. Zhou et al. (1999) have also demonstrated that the level of copper enrichment is also influenced by grain orientation. As referred to previously, the process of copper enrichment at the aluminium alloy surface can vary with applied (over)potential, which suggests that observations from accelerated electrochemical processes may not be directly transferable to the progression of damage occurring at the corrosion potential. Jung et al. (1985) investigated the copper enrichment at the surface of Al-Cu binary alloys (containing up to 2 at% copper) after polarizing samples at either 500 mVSCE or -100 mVSCE in sulphuric acid. They found at -0.5 VSCE, where copper was cathodically protected, that the surface anodic film contained very little copper, but that copper particles (4 to 25 nm) were present on the surface. Further copper enrichment was present beneath the oxide film. These particles were thought to develop from an enriched layer beneath the oxide. Some of the particle dimensions were larger than the film thickness and, consequently, had no oxide covering. At a more aggressive potential of -0.1 VSCE less enrichment was achieved at the alloy/oxide interface, which was attributed to the active dissolution of copper. Under these more aggressive conditions there was evidence of copper (or Cu-Al Alloy) fragments incorporated into the oxide. In terms of general corrosion performance, the enrichment of copper at the alloy surface is also likely to increase the number of flaws that exist in the aluminium oxide. The level of copper and particularly its distribution at the grain boundaries can also have a profound effect on intergranular corrosion performance. Problems occur most commonly in high-copper wrought alloys (i.e. AA2xxx and AA7xxx), particularly when slow quenching leads to precipitation from solid solution and subsequent intermetallic particle growth, particularly along grain boundaries. Slow quenching or artificial aging (at 190 ºC) allows the formation of copper-rich precipitate phases close to the grain boundaries (these differ from coarse intermetallics that are developed prior to quenching or heat treatment). Copper (and other alloying elements) diffuse from the matrix into the particle, leaving a copper-depleted region surrounding the intermetallic [Zhang and Frankel (2003)].

Corrosion

33

In addition to grain boundaries, precipitate-free zones also develop around coarse secondary intermetallics such as Al2CuMg [Guillaumin and Mankowski (1999)]. Decreased copper in these zones leads to a more active potential (up to 120 mV have been reported between the AA2024-T3 alloy matrix and Cu-depleted zones [Davis (1999)] and, as detailed by Galvele and co-workers [Galvele and de Micheli (1970) Muller and Galvele (1977)], a more active breakdown potential at the depleted zones. The breakdown (or pitting) potential is an electrochemically measured potential that is associated with the breakdown of the oxide (or metallic phase), above which the oxide can no longer re-passivate the surface. Two breakdown potentials have systematically been reported for high Cu-Al alloys [Galvele and de Micheli (1970), Guillaumin and Mankowski (1999), Zhang and Frankel (2003)]. Galvele and Micheli (1970) attributed the first breakdown potential to the dissolution of copper-depleted zones, and, the second, to pitting of the grain bodies. For AA2024-T3 alloy, Guillaumin and Mankowski (1999) suggested that the first breakdown potential was related to the preferential dissolution of Al2CuMg particles and the second to matrix and grain boundary breakdown. More recent work by Zhang and Frankel (2003) on AA2024-T3 alloy agreed that Al2CuMg dissolution leads to the first breakdown potential, but that the second breakdown results from the initiation and growth of intergranular corrosion. Breakdown potentials provide a good indication of whether pitting or intergranular corrosion is likely to occur in particular environments but their values are somewhat reliant upon the exact experimental details (i.e. scan rate), and contributing localized corrosion processes cannot always be isolated. Nevertheless, experimental evidence has shown that the addition of copper initially shifts the breakdown potential of aluminium to a more noble value. However, during ageing of the alloy, the breakdown potential gradually decreases as the copper-containing alloys become increasingly susceptible to intergranular corrosion. In addition, the susceptibility of aluminium alloys to intergranular corrosion is not strictly reserved for alloys with higher copper concentrations. Svenningsen et al. (2004) studied model AA6xxx series extrusions, finding that alloys possessing low copper (approx. 0.02 wt%) were resistant to localized corrosion whilst copper additions of approximately 0.17 wt% were prone to intergranular attack. Electron microscopy showed the presence of coppercontaining precipitates (Al5Cu2Mg8Si6) at the grain boundaries of alloys susceptible to intergranular corrosion. The Al5Cu2Mg8Si6 phase was demonstrated to possess a noble potential in comparison to the matrix and accelerated matrix dissolution.

34

T. H. Muster, A. E. Hughes and G. E. Thompson

THE INFLUENCE OF COPPER IN CATHODIC INTERMETALLICS Cathodic constituent particles are present in AA2xxx, AA3xxx, AA6xxx and AA7xxx alloys. Copper-containing intermetallic phases (particularly phases with higher levels of copper in the Al-Cu-Fe-Mn family, such as Al6(Cu,Fe,Mn), Al2Cu, Al7Cu2Fe and Al20Cu2(Fe,Mn)3 phases), support high cathodic reaction rates under most solution conditions, and assist in the initiation of localized pitting corrosion [Scully et al., (1993), Chen et al., (1996), Schneider et al. (2004), Birbilis and Buchheit (2005)]. It is noted that cathodic intermetallic phases that do not contain copper (i.e. Al3Fe, α-Al(Fe,Mn,Si)) can initiate localized corrosion; however, initiation rates are much reduced compared to alloys with coppercontaining intermetallics [Nisancioglu (1990), Blanc and Mankowski (1997), Birbilis and Buchheit (2005)]. The ability of Al-Cu-Fe-Mn (or similar) phases to support high cathodic reaction rates commonly results in pitting initiation where breakdown of the oxide occurs adjacent to cathodic particles [Chen et al. (1996), Liao et al. (1998), Alodan and Smyrl (1998), Guillaumin and Mankowski (1999), Leclere and Newman (2002), Vukmirovic et al. (2002), Ilevbare et al. (2004)]. Figure 11 shows scanning electron micorgraphs of a section through two pits developed on AA2024-T3 alloy which had been immersed in 0.5M NaCl for 24 hours. The top surface and sections of the alloy can be seen in tilted micrographs (Figures 11 (a) and (b) where the top surface is rotated 30° from the normal allowing a view of the sectioned face), and sections of the S-phase pit can be seen in Figures 11(c) and (d). The large pit has two Al-Cu-Fe-Mn intermetallic particle remnants and displays excessive trenching with etch patterns. It is likely that most of the intermetallic particles have been removed, presumably by dissolution. Slow dissolution of Al-Cu-Fe-Mn phases have been observed, pitting has also been seen on the Al-Cu-Fe-Mn particles themselves [Ilevbare et al. (2004)], which is likely to result from the heterogeneous structure of the intermetallics, as shown by SEM and AFM imaging (Figure 12). The small pit to the left of the large pit in Figure 11a is associated with an S-phase particle and has quite different characteristics, which will be discussed in the following section. Whilst the appearance of trenches around Al-Cu-Fe-Mn phases have been observed by many, there is some uncertainly with regards to the mechanism of trenching initiation and development. Even less is understood about the propagation of pits from existing trenches. Trenching of the matrix surrounding cathodic intermetallics may be interpreted as a galvanic corrosion between the particle and the matrix [Buchheit et al. (1997), Ilevbare et al. (2004)], but is also

35

Corrosion

attributed to etching under the high pH conditions generated by cathodic reactions (oxygen and/or water reduction). Several studies have reported pH values around 9.5 at the edge of cathodic intermetallics [Park et al. (1999), Vukmirovic et al. (2002)].

(b)

(a) Cu-Fe-Mn-Al S-Phase

(d)

(c)

Figure 11. AA2024-T3 which has been corroded in 0.5 M NaCl for 24 hours. (a) secondary and (b) backscatter images of top surface and section of two pits. The large pit has Al-Cu-Fe-Mn intermetallic remnants and displays trenching where the small pit to the left (white square) has Cu particulates and is assumed to result from S-phase etchout, (c) secondary and (d) backscatter images of section of S-phase pit.

The high pH conditions are suggested to enable “cathodic” corrosion of the aluminium adjacent to the intermetallics, assisted by chemical reaction with the surface oxide, producing aluminate ions and hydrogen gas [Moon and Pyun (1997, 1999)]: +

−

Oxidation via oxygen vacancies: 2 Al + 3H 2O → Al2O3 + 6 H + 6e …5a −

Direct formation of hydroxide film: Al + 3OH → Al (OH )3 + 3e

−

…5b

36

T. H. Muster, A. E. Hughes and G. E. Thompson −

−

Dissolution of alumina: Al2O3 + 2OH + 3H 2O → 2 Al (OH ) 4 −

…6a −

Dissolution of aluminium hydroxide: Al (OH )3 + OH → Al (OH ) 4 …6b Combined −

oxidation − 4

Al + 4OH → Al (OH ) + 3e

(5)

and

dissolution

−

(6): …7

−

Water reduction : 2 H 2O + 2e → H 2 + 2OH

−

…8

Combined water reduction (8) and oxidation and dissolution (7):

Al + 3H 2O + OH − → Al (OH ) 4− +

3 H2 2

…9.

The rate at which equation 9 occurs is increased at high pH, and has been shown to be independent of the current density; the reaction rate is almost as high when the alloy is under open-circuit conditions as when cathodically polarized [Moon and Pyun, 1997 ; Buchler et al. (2000), Leclere and Newman (2002)]. Increased buffering of the solution decreases the rate of pitting around cathodic, intermetallic particles; this has been attributed to the inability of high pH conditions to be sustained across the surface of the surrounding matrix [Park et al. (1999), Vukmirovic et al. (2002)]. In order to establish the high pH conditions at cathodic intermetallics an equal number of anodic reactions are required at some distance from the particles [Buchler et al. (2000)]. These anodic reactions are thought to occur on the matrix aluminium, in the dispersoid-free zone surrounding the intermetallic, or, if present, on nearby Al2CuMg particles [Ilevbare et al. (2004)]. Leclere and Newman (2002) suggest that local galvanic cells were unlikely to exist over the small distances involved in trenching in a conducting electrolyte; however, the formation of colloidal alumina gels between anodic and cathodic regions on the surface may aid in maintaining potential differences and stabilizing pit growth [Park et al. (1999), Ilevbare et al. (2004)]. Furthermore, Ilevbare et al. (2004) reported that, as the pH rises at the site of cathodic intermetallics, and decreased pH conditions develop at anodic sites, increased galvanic potential differences are created (by comparison with potential differences at the bulk electrolyte pH). Schneider et al. (2004) suggested that a simple pH-induced description might not be adequate to explain all trenching

Corrosion

37

phenomena. For instance, increased chloride ion concentration led to an earlier onset of trenching but had little influence on oxygen reduction kinetics. Also, experimentally observed trench widths of 1 – 2 μm are significantly narrower than expected given that high pH zones extended up to 50 μm. Schneider et al. (2004) suggested that trenching could in fact be described by a model based upon a galvanic-couple assisted passive breakdown of the dispersoid-free zone, where the trenching may be viewed as an unstable pitting process. Such a model could describe the conversion of trenches to form stable pits containing acidic anolyte solution, if sufficiently high anodic currents could be generated. In contrast, the mechanism of a model, where trenches formed at high pH and later switching to an acid pitting mechanism, is not clear despite some attempted explanation [Park et al. (1999), Leclere and Newman (2002)].

Figure 12. Atomic force microscopy topography image of AA2024-T3 exposed to nonchromate deoxidation (Turco SmutGo) for 5 minutes. The large Al-Cu-Fe-Mn particle shows a heterogeneous microstructure. Trenching is observed around the Al-Cu-Fe-Mn particle and smaller pits to the bottom left were the site of Al2CuMg intermetallics (courtesy Dr T. Muster).

38

T. H. Muster, A. E. Hughes and G. E. Thompson

The mechanism of pitting in aluminium alloys at anodic sites has been the subject of corrosion research for many years. Local anodic reactions, occurring at the site of the oxide breakdown can result in pits development. Chloride (and other halide) anions assist in the breakdown of oxides and pit initiation [Davis (1993)]. Early stages in pit growth are thought to be autocatalytic due to the generation of a low pH environment at the base of the pit, which catalyses further pit growth [Hatch (1984)]. The presence of a cathode external to the pit, enables the anodic production of aluminium ions (eqn 10), which react with water forming aluminium hydroxides and releasing protons (eqn 11). Protons are reduced at cathodic sites within the pit to release hydrogen gas (eqn 12). The presence of various salts can alter the exact composition of corrosion products described in equation 11, depending on complexation chemistry. For instance, chloride ions complex with Al3+ ions at anodic sites and prevent the formation of aluminium hydroxides.

Al → Al(3aq+ ) + 3e −

…10

Al(3aq+ ) + 3H 2O → Al (OH )3( s ) + 3H (+aq )

…11

2 H + + 2e − → H 2( g )

…12

After the initiation of a pit, it may either passivate or continue with stable growth [Szklarska-Smialowska (1999)]. Whilst the autocatalytic mechanism may explain pit initiation and early stages of growth, it is suggested that cathodic reactions involving oxygen reduction are still required for stable pitting in most instances [Pride et al. (1994)]. For the continued growth of pits, the solution potential must exceed a critical pitting (breakdown) potential. Further, it has been shown that the ratio of current to pit radius must exceed approximately10-2A/cm for stable pit growth [Pride et al. (1994)]. The addition of copper to aluminium alloys increases the pitting potential; however, the corrosion potential is also increased and, in general, the resistance against pitting usually decreases with increasing copper content [Davis (1999)]. The surface of metallic copper is highly efficient at reducing oxygen, and therefore, copper-rich sites allow oxygen and proton reduction reactions to occur with an enhanced efficiency, thus increasing the probability of stable pit growth [Davis, 1999]. The high cathodic currents supported by copper-containing intermetallics are thought to rely upon mixed reaction control [Buchheit et al. (1999), Jakab et al. (2005)]. Jakab et al. (2005)

Corrosion

39

investigated the oxygen reduction reaction (equation 1) occurring on the surface of AA2024-T3 alloy. The amount of copper on the surface was varied (through exposure to an alkaline etch) and the corresponding changes in the oxygen reduction reaction (ORR) modeled. The charge-transfer controlled ORR was shown to increase linearly with copper coverage whilst the mass-transfer controlled ORR increased as a complex function of copper coverage and the inverse of the boundary layer thickness. The efficiency of copper containing intermetallics to act as cathodes also enables them to support copper plating reactions whereby soluble copper ions are reduced back to copper metal. Chen et al. (1996) studied the corrosion processes on polished AA2024 alloy, noting that Cu-deposition as nodules was observed on the surface of Al-Cu-Fe-Mn particles. This indicates that Cu is dissolved from either intermetallics particles or the matrix and undergoes reduction on the cathodic particles. Deposition of Cu is particularly a problem in cooling systems which comprise mixed metals, and may contain brass/copper and aluminium. Blackwood and Chong (1998) studied the deposition of Cu onto AA6063 plate from CuSO4 solutions with and without Cl- ions. They noted that the deposition of copper ions onto copper metal was more favoured than copper ions onto aluminium metal; however, in the presence of chloride ions, they proposed that the formation of CuCl+ lowered the energy barrier for deposition of copper onto aluminium allowing more deposition of copper onto the aluminium. It is worth noting here that the surface of aluminium alloys containing low amounts of copper can be altered by the presence of soluble copper ions [Bjørgum et al. (1995)]. Where heavy metals are deposited on the surface and become soluble (i.e. at low pH), they are able to reduce back to their metallic state at cathodic sites whilst oxidizing aluminium [Szklarska-Smialowska (1999)]. Copper is the greatest concern in this respect, as only a low concentration (i.e. 0.02 - 0.05 ppm) is needed to initiate pitting at neutral or acidic pH. Even alloys with greater than 99.5 % purity have cathodic regions that are active enough to plate copper onto their surface [Bjørgum et al. (1995)], and therefore the availability of copper ions from the environment can have severe consequences for all aluminium alloys. Several studies have shown the importance of cathodic Cu-containing intermetallics in controlling pit initiation and growth. Studies carried out on areas less than 0.01 mm2 on polished AA2024-T3 alloy have demonstrated that Al-CuFe-Mn phases are critical for pit initiation in chloride-containing environments [Leblanc and Frankel (2002)]. Small areas of the alloy, exposed to 0.5 M NaCl without the presence of Al-Cu-Fe-Mn phases, failed to display pitting over periods of two hours, even where active phases such as Al2CuMg phases were present. Liao et al. (1998) identified constituent particles as initiators of pitting

40

T. H. Muster, A. E. Hughes and G. E. Thompson

and of pit growth, however, they suggested that pits were shallow unless there was a cluster of constituent particles, which would promote deeper pitting. For instance, Figure 12 shows the subsurface attack on AA2024 during corrosion along grain boundaries and around intermetallic phases. The brighter contrast between corroded intermetallics and uncorroded intermetallics suggests copper enrichment of the remnant particles. Subsurface corrosion is thought to be promoted by the presence and linking of copper-rich intermetallic phases.

Figure 12. Left: Backscattered scanning electron micrograph of corroded AA2024 showing subsurface attack, Right: Conceptual model for particle-induced subsurface attack after Liao et al. (1998).

Given that stable pit growth requires a significant exchange current to continue growth [Pride et al. (1994)], it appears logical that the area fraction of the various phases, and their proximity to each other, could play a controlling role in pit formation. The work of Boag et al. (2005) and Ilevbare et al. (2004) also support the views of Liao et al. (1998), that the proximity and distribution of intermetallic phases may control initiation kinetics and sustained pit growth. Also, it is presently unknown whether colloidal copper particles and copper metal plated on the surface through adsorption and reduction of dissolved copper in solution, which may also deposit on the walls of a pit, play any significant role in stabilizing pit growth by providing efficient local cathodes. Further understanding of pit growth is likely to delineate the importance of intermetallic distribution profiles, colloidal copper particles and re-plated copper.

Corrosion

41

THE INFLUENCE OF COPPER IN S-PHASE INTERMETALLICS Al-Cu-Mg ternary alloy systems are a special case due to the presence of Al2CuMg (S-phase) particles. S-phase particles are of interest for two reasons; firstly, they make up approximately 64 % of the intermetallic surface area, and secondly, they have a corrosion potential significantly less noble than the common AA2024-T3 and AA7075-T6 alloys, and are subject to selective dissolution under most corrosive conditions [Buchheit et al. (1997)]. In AA2xxx series aluminium alloys, Al2CuMg phases are known to be the most active, and pits are commonly generated on, or around, these particles. In recent years, significant work has been performed to detail the mechanisms of damage associated with S-phase, and secondly, in developing methodologies to minimize its influence on corrosion and surface finishing processes [Juffs et al. (2001), McGovern et al. (2000), Yoon and Buchheit (2006)]. Several studies have noted that S-phase is anodic to the aluminium matrix and therefore preferentially attacked [Alodan and Smyrl (1998), Chen et al. (1996)]. In contrast, other studies have reported that the matrix surrounding S-phase particles has been attacked (analogous to trenching discussed above), which can eventually lead to undermining and ‘drop-out’ of the particle [Guillaumin and Mankowski, 1999]. Figure 13 shows a 2D profile across an Al2CuMg particle showed that has been etched around the side of the particle. These two observations appear at first glance to be conflicting; however, several detailed analyses of S-phase properties have led to consensus on most of the mechanistic details [Buchheit et al. (2000)]. It summary, the consensus is that the S-phase is initially anodic and then, upon the dealloying of magnesium and aluminium, a copper-rich phase is left [Chen et al. (1996)]. The enrichment in copper results in S-phase being transformed from being a net anode to a net cathode. For example, Figure 14 shows a plume of hydrated aluminium-magnesium oxide emanating from the surface of AA2024-T3 alloy. In Figure 14 (a) and (b) the plume of oxide is formed on the surface after treatment in cerium chloride solution; it demonstrates a morphology that is typical of S-phase dissolution. Figures 14 (c) and (d) show another plume of corrosion product from an S-phase particle. This plume also protrudes from the surface and bright copper-particulates are incorporated into the corrosion plume. These copper particulates are release from the copper-rich phase following the selective dissolution of aluminium and magnesium. Previously, in Figure 11, a small pit was identified as the site of an S-phase remnant. The corrosion product is full of heavy atomic number particulate material (bright spots in Figure 11 (d)), which were identified as copper. Higher magnification images (Figure 15), revealed that the individual bright spots in

42

T. H. Muster, A. E. Hughes and G. E. Thompson

Figure 11(d) themselves comprise clusters of nanoparticles 50 to 200 nm in size. This smaller pit shows the characteristics described by Buchheit et al. (2000) for the mechanical redistribution of Cu redistribution via transport in the corrosion product as will be discussed below.

a. Polished ~ 0.3 μm

b. Polarised

Figure 13. 2D profiles of Al2CuMg after (a) polishing and (b) 5 min of polarization at above the breakdown potential of S-phase (-750 mVSCE), after Guillaumin and Mankowski (1999).

Buchheit et al. (1999) demonstrated that potential scans, conducted on bulk S-phase, provided indirect evidence of the dealloying of S-phase. Anodic polarization scans, starting at 50 mV cathodic to the steady state open-circuit potential (OCP) in 0.5 M NaCl, showed a positive shift (up to 150 mV) in the OCP determined during the scan. These positive shifts in the OCP were attributed to cathodic corrosion, where increased alkalinity leads to the preferential dissolution of aluminium from the matrix. Interestingly, after the anodic scan, when one would expect to find the enrichment of elemental copper at the surface, the OCP remains at similar values to previously unexposed S-phase. Buchheit et al. (2000) suggested that electrical isolation of deposited colloidal copper means that it does not alter the OCP by any significant amount.

43

Corrosion

(a)

(b)

(c)

(d) Cu particulates

Figure 14. (a) Secoondary and (b) backscattered electron images of AA2024-T3 which has been exposed to acidified CeCl3 solution with 3% H2O2 and rinsed in ethanol (from Gorman, 1998). Milled surface of AA2024-T3 alloy which has been exposed to neutral salt spray for 8 days.

S-phase is an active anode with dissolution rates reported to be equivalent to 20 mA cm-2 (from visual estimation) during corrosion processes on an AA2024 alloy in 0.1 M sulphate and 0.005 M chloride solution [Ilevbare et al. (2004)]. Aldykewicz et al. (1995) reported decreased anodic currents, in the range of 30 to 60 μA cm-2, using scanning vibrating probe analysis over an AA2024 alloy in 0.012 M NaCl solution. There are a number of reasons for the large difference in the magnitude of estimated current exchanges, including electrolyte concentration, pH and non-Faradaic dissolution of magnesium and aluminium into solution. Buchheit et al. (1999) showed that in addition to the S-phase being an active anode, it supports high current densities for cathodic reactions (exceeding 1 mA cm-2). The efficiency of S-phase to support high cathodic reaction rates has been attributed to the formation of large surface area ‘sponges’ of copper produced by

44

T. H. Muster, A. E. Hughes and G. E. Thompson

dealloying [Buchheit et al., 1999]. (A visual example of a copper-rich sponge can be found in Figure 21 in the section on deoxidizing/desmutting.) Numerous studies have detailed the existence of Cu sponges [Buchheit et al. (1997), Buchheit et al. (1999), Chen et al. (1996), Guillaumin and Mankowski (1999), Kolics et al. (2001)]. The cathodic properties of the copper-rich phase lead to dissolution of the matrix surrounding the original particle. Also, the matrix is thought to be depleted in dispersoid particles in the areas surrounding S-phase particles [Guillaumin and Mankowski (1999)], although some analysis shows that this may not always be the case [Buchheit et al. (1997)]. Guillaumin and Manikowski (1999) reported what appeared to be copper deposits on the matrix directly adjacent to the copper-rich S-phase remnants. Buchheit et al. (2000) have also observed pitting around the periphery of S-phase particles and suggested the initiation of secondary galvanic cells. Numerous studies have reported copper enrichment at the surface of aluminum alloys associated with S-phase intermetallics [(Buchheit et al. (1997), Chen et al. (1996)]. Movement of copper particulates of 10 to 100 nm in dimension have been observed to originate from the site of Al2CuMg particles and to migrate to the outermost surface [Buchheit et al. (1997), Chen et al. (1996), Buchheit et al. (2000)]. Several potential mechanisms of copper enrichment have been suggested; including the direct oxidation of copper into solution and the nonFaradaic release of copper nanoparticles from copper-rich sponges [Buchheit et al. (2000)].

Cu-particles

(a)

Amorphous Oxide

(b)

Figure 15. (a) Secondary and (b) Backscattered electron images of Cu particles embedded in an amorphous oxide. These Cu particles are thought to be the corrosion product form the dissolution of S-phase depicted in Figure 10(c) and (d).

Studies have shown the presence of copper ions in solution during electrochemical testing of S-phase and AA2024 alloys. Vukmirovic et al. (2002)

Corrosion

45

found that copper ions were released from AA2024 alloy. Buchheit et al. (2000) used stripping voltammograms to confirm the release of both copper ions and elemental copper into solution during the dissolution of Al2CuMg into aqueous chloride solutions. Significant peaks were observed relating to the oxidation of Cu → Cu+ and Cu+ → Cu2+ for Al2CuMg when held at its open-circuit potential and, additionally, when polarized ± 50 mV with respect to the open-circuit potential. Therefore, evidence suggests that it is possible to remove copper from the surface of aluminium alloys even though bulk copper is thermodynamically stable at the corrosion potentials achieved in solution. As a consequence, it is likely that copper redistribution cannot be completely suppressed, even using electrochemical techniques [Buchheit et al. (2000)]. The importance of S-phase dissolution in the generation of surface copper was recognized in the work of Vukmirovic et al. (2002), who compared the amount of Cu released into solution from model and commercial AA2024 specimens. Commercial AA2024 alloy showed much higher levels of solubilised copper, which was attributed to its ability to form copper-rich sponges. The model sample, consisting of a lowcopper alloy aluminium covered with physical vapour deposited copper islands, produced much lower amounts of solubilised copper. In order for copper ions to directly oxidize from the alloy in solution, the potential of the copper phase must be decreased. Buchheit et al. (2000) suggested two mechanisms by which the open-circuit potential of copper can approach the more negative values achieved in Al-Cu-Mg aluminium alloys. Firstly, the complexing of copper ions with chlorides can alter the activity of copper ions in solution and decrease open-circuit potentials. This effect is only possible at pH values exceeding approximately 5.5 and becomes increasingly significant at high pH. Secondly, the radius of curvature of materials has a large effect on their free energy and reduces the open-circuit potential, particularly where curvatures approach the nanometer scale. This scenario is possible given that preferential leaching of magnesium and aluminium can lead to the presence of high surface area regions of copper. In order to reduce surface energy, the structure coarsens through coalescence mechanisms based upon differential solubilities [Brinker and Scherer (1990), Buchheit et al. (1997)]. Buchheit et al. (1997) suggested that particle remnants of metallic copper become detached from the surface, presumably through the mechanical action of pit growth and corrosion product formation or through solution movement, as depicted in Figure 16 [Buchheit et al. (1997)]. The argument that Cu-particles are released by mechanical action is supported by the work of Dimitrov et al. (1999), who immersed AA2024-T3 into both stirred (500 rpm) and unstirred 0.5 M NaCl solutions for periods up to 35 minutes.

46

T. H. Muster, A. E. Hughes and G. E. Thompson

Upon the removal of the AA2024-T3 samples, stirred samples routinely possessed 1.8 times more copper than that of the unstirred sample. Dimitrov et al. (1999) and later Vukmirovic et al. (2002) suggested that the rate of redeposition of copper fragments, previously removed from the surface via non-faradaic processes, was increased due to the added convection during stirring.

Figure 16. Schematic diagram of a mechanism for redistribution of copper by dissolution of large Al2CuMg and Al2Cu intermetallic particles after Buchheit et al. (2000). Dealloying of the particle is caused by anodic polarization of the particle by the surrounding matrix, resulting in the formation of a copper-rich network (which coarsens with age) that under hydrodynamic flow or stresses arising through corrosion product formation, releases small copper-rich clusters. The electrically isolated copper clusters oxidize at their OCP, and may precipitate as copper-oxides or be reduced back to elemental Cu at cathodic sites and serve as efficient local cathodes that stimulate secondary pitting.

S-phase shows some evidence of passivation under high pH, possibly as a result of the stability of magnesium oxide in alkaline environments [Ilevbare et al., 2004]. However, even at high pH, any passivity is overcome by the significant anodic polarization that exists when coupled to an aluminium matrix phase [Ilevbare et al., 2004]. The oxide covering the S-phase varies from that of the matrix oxide [Guillaumin et al., 2001] and is thought to have weak protective properties [Vukmirovic et al., 2002]. Therefore, the weakened oxide and anodic

Corrosion

47

polarization experienced by the S-phase particles results in rapid attack in corrosive environments.

Figure 17. Schematic diagram of the advanced stages of corrosion induced by dealloyed Sphase particles. Trenching is observed around the particles, however, the nature of the reactions within the trench is an area of controversy. It is thought that the trenching is initiated by alkaline conditions due to the cathodic activity of the Cu-rich S-phase remnant. However, in the latter stages of corrosion, acidity might also develop at the base of the trench setting up a differential aeration cell. Secondary pitting is also observed due reduction of Cu ions on the matrix which serve as local cathodic sites that stimulate anodic reactions on the adjacent matrix.

A significant body of work has detailed the peculiar behaviour of S-phase intermetallics, however, there appears to be some areas that could benefit from further study. The exact structure of dealloyed S-phase is not well understood. Studies by Sieradzki and co-workers have made some progress in this area of noble metal enrichment, describing the dealloyed structures in terms of percolation theory [Sieradzki (1993), Newman and Sieradzki (1994), Vukmirovic et al. (2002)]. The dealloyed structure is predicted to be determined by a combination of the dissolution rate and copper content. High dissolution rates do not allow surface diffusion and relaxation processes to coarsen the structure, leading to more open networks. Also, copper contents cannot be too high, otherwise dealloying will not lead to an isolation of the percolation network, which is required for the mechanical release of copper clusters. With a copper concentration of 25 at% and an orthorhombic structure, S-phase is considered to

48

T. H. Muster, A. E. Hughes and G. E. Thompson

be borderline in its ability to release copper clusters [Vukmirovic et al. (2002)]. Hydrodynamic forces would certainly assist in fragmentation. Finally, the influence of S-phase and its corrosion by-products on corrosion is not well documented or understood. Figure 17 shows a schematic of the types of corrosion processes that may occur around an S-phase after longer time periods (compared with the dealloying mechanism presented in Figure 16). Trenching is initiated around the dealloyed particle as it switches from behaving as a net anode to a net cathode. Some trenched areas may be converted to acidic pits due to the separation of anodic and cathodic regions due to the establishment of differential aeration cells, which are created via the precipitation of colloidal alumina gels (corrosion products) within the pit. Pitting itself may also be enhanced through copper plating at cathodic sites on pit walls, and the deposition of elemental copper on the aluminium matrix may result in the formation of secondary pits. The role of oxidized copper particles is not clear in terms of advanced corrosion, although their presence would ensure a constant supply of soluble copper, which is expected to decrease overall corrosion performance.

Chapter 6

CHEMICALLY PRETREATED SURFACES FINISHING PROCESSES Metal finishing is a term that encompasses a range of chemical processes for taking aluminum and its alloys from the as-received state to a state where the surface has been prepared (in a reproducible way) for further processing, such as the application of an organic coating or adhesive [Critchlow and Brevis (1996)]. Many of the metal finishing processes involve a series of treatment steps that require alternate immersion in acid and/or alkali treatments designed to modify the surface chemistry prior to anodizing or conversion coating. A range of mechanical treatments, such as grit blasting with graded alumina or silica, also exist for aluminium surface preparation [Critchlow and Brevis (1996)], but there is little or no information on how such mechanical treatments affect the copper distribution. During the last twenty years, there has been considerable pressure to make metal finishing processes more environmentally friendly and less of a health risk. This has led to research into a range of new processes. Much of this work has been performed in research environments where, often, a disjunct occurs between the outcomes of the research and the requirements of the metal finishing industry and its endusers. Such differences in perspectives between the research community and the metal finishers may be demonstrated by the following example. For many years, chromate/HNO3/HF deoxidisers have been used for the treatment of coppercontaining, sheet aluminium products. This deoxidizer effectively removes Sphase particles and essentially dissolves the Al-Cu-Fe-Mn intermetallic particles in proportion to their composition without preferential dissolution that leads to

50

T. H. Muster, A. E. Hughes and G. E. Thompson

copper enrichment of the intermetallic remnant [Hughes et al. (2003)] as might be observed with other deoxidizer combinations [Hughes et al. 2001]. Further, the etch rate is not so high as to remove sufficient aluminium matrix to continually expose new intermetallic constituent particles. Thus, it is possible to obtain a surface which has a significantly reduced population of intermetallic particles, but it does develop a copper-enriched layer beneath the surface oxide. Recently, Moffitt et al. [2001] who, on the observation of this thin copper enrichment layer beneath the surface oxide on AA2024-T3 alloy after treatment in a nitric acid based deoxidizer, suggested: “The surface … maintains a copper enriched surface despite the apparently long held claim, that this eliminates surface copper enrichments”. The issue here is not whether Moffitt is correct or those he quotes are incorrect, but the perspective from which they approach the concept of cleanliness. In metal finishing, nitric acid solutions remove copper-containing smut, giving a visually bright appearance which is all the processor needs to know for quality control; however, if the surface is examined with sophisticated techniques such as TEM, Auger or XPS, then copper will be detected in enriched layers and intermetallic remnants. Moffitt et al. [2001] partly condede this point in their next sentence, citing that the original claim was made in 1958 well before surface science techniques were commonly available. With this example in mind, one aim of this section is to highlight the differences between the requirements for metal finishing, and the approach of surface science, both of which have a part to play in the understanding of the current processes, and the development of new ones. The metrics used by each community can be used to further elaborate on the difference between the metal finishing industry and the research community. Metal finishers deal with etch rates, coating weights and processing bath maintenance to evaluate the operating state of coating lines. They use visual appearance as one of the major quality control tools, although more sophisticated tools are commonly available and used (X-Ray fluorescence, for example). The eveness of colour of the coating is a good example of a visual quality control tool [Schmidt-Hansberg and Schubach (2003)]. Metal finishers also use performance testing in standardized corrosion environments, such as Mil-C-81605 and Mil-C5541E in conjunction with standard operating procedures such as ASTM B117 to estimate performance “in-service”. Research laboratories, by contrast, often use sophisticated techniques to study small areas of samples. They often use fresh solutions rather than “aged” or “sweetened” solutions, which have some metal dissolved into it, or a surface condition, such as polished to 1 μm or less, as opposed to chemically cleaned. Hence, it is important to identify these differences when relating the surface chemistry, metallography and morphology to metal

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51

finishing parameters, when improving current processes and developing new ones. Even within the metal finishing industry, it is clear from a number of studies that changing the processing parameters, or introducing new processing products, can result in marked differences in the performance of the metal finishing process. Ketcham and Brown’s (1976) work on the performance of chromate conversion coatings (CrCC) pretreated with a variety of different commercial products, all qualified and commercially available (alkaline cleaners and deoxidation solutions), showed that considerable variation resulted between products in terms of the final corrosion resistance provided by the CrCC. In another study, Leggatt et al. (2002), found that CrCCs applied on AA2024, AA6061 and AA7075 alloys in the field, failed to pass corrosion resistance performance requirements despite the use of appropriate standards for application of the CrCC’s. These standards were designed to give the required performance. These studies show that etching and surface condition, including surface enrichment of alloying components play an important role in performance. From the point of view of the enrichment of alloying elements, particularly copper, there are two main considerations: 1. etch rate of the metal finishing solution 2. solubility of the alloying element in the metal finishing solution The etch rate determines the amount of buildup of material on the surface, but the chemistry of the treatment solution has a bearing on the whether that material is dissolved into solution. To this end sequestering agents are often added to commercial products to complex the metal ions to prevent them from depositing on the surface during the treatment step. From the perspective of etchants, a high etch rate (NaOH or fluoride) results in fast removal of aluminium, but leaves copper behind on the surface. For this reason, HNO3 is either used as a post treatment in the case of alkali treatments, or added to acidic treatment steps. In the following section the influence of individual treatment steps on aluminiuim-alloys will be examined to illustrate how they change the surface condition.

SOLVENT CLEANING The purpose of solvent cleaning, using either solution or vapour processes, is to remove oils and greases that may have been applied to the surface of the alloys for the purposes of corrosion protection during transport. Typical vapour

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T. H. Muster, A. E. Hughes and G. E. Thompson

degreasers included organo chlorines such as trichloroethane, trichloromethane and trichloroethylene. Solvent cleaning can be divided into the type of cleaning used in the metal finishing industry and those used in the research laboratory. In the former, solvents, such as those listed in Table 7, are often used, whereas in the laboratory, a different range of reagents such as acetone, ethanol or methanol are frequently used for sample preparation. In the metal finishing environment, solvent cleaning is disappearing. Table 7. Solvents Used for Wiping Workpieces Clean Metal Finishing Paraffin (kerosene) White Spirit Carbon tetrachloride Methylated Spirits

Research Environment Acetone Ethanol Methanol Trichloro ethane

Clearly the list on the left of Table 7 is different from that on the right and has some implications for drawing conclusions about metal finishing and corrosion processes from laboratory research. For example, Chidambaram and Halada (2001) explored acetone degreasing of polished AA2024-T3 alloy, and observed carbonyl groups on the surface; this led them to suggest the formation of surface acetates (both copper and aluminium). They further proposed the formation of acetic acid, and suggested that in the acidity produced from acetic acid, copper chloro-complexes are formed which assist corrosion. They also point to the role of photo-oxidation in modifying the speciation of the complexes. While these studies cannot be directly related to surfaces during metal finishing, they have some relevance in terms of artifacts produced in the research environment that are not likely to be present in the processing environment. In the metal finishing industry, issues associated with solvent wiping have also been identified. Trichloroethane, when used in vapour degreasing, can react with water to produce HCl which dramatically attacks the aluminium [King (1988)]. Commercial solvents have added stabilizers and inhibitors to prevent such reactions [King 1988]. Hughes et al. (1996) also noted the presence of surface chlorine, using XPS, on aluminium oxide after a solvent wipe with trichloroethane, supporting the model that where acidity develops on the surface, HCl may have an important role to play in subsequent corrosion events.

Chemically Pretreated Surfaces

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DETERGENT CLEANING Metal finishing has undergone considerable change in the last 20 years due to environmental legislation and concerns regarding occupational health. From the environmental point of view there has been a drive to remove volatile organic chemicals from industrial processes. This has impacted metal finishing processing due to the removal of vapour degreasers, and changes in the composition of paint systems from low-solids to high-solids based as well as a general trend away from organic-based to aqueous –based paint systems. Solvent cleaning using the vapour phase has been replaced by detergent cleaners, which generally have very low or non-existent etch rates. Under these circumstances, there are only minor changes from the original structure and chemistry of the surface, and possible adsorption of surfactants onto the oxide surface.

ALKALINE CLEANING Aqueous alkaline cleaners were traditionally used after vapour degreasers to further saponify any oils or grease remaining on the surface after vapour degreasing (or detergent cleaning) as well as give a mild etch to the surface [Wernick et al. (1987), King (1988)]. These solutions are usually carbonate-based and inhibited for non-etch cleaners, or caustic based with inhibitor for etch cleaners[Wernick et al., (1987)]. Inhibitors include silicates, chromates, phosphates, fluorides, silicofluorides or organics to reduce the rate of etching [Wernick et al., (1987)]. Modification to the surface during this process step usually involves the generation of basic zinc and magnesium oxides on the surface; these must be removed during deoxidizing or desmutting steps. Ketcham and Brown (1976) demonstrated that different formulations of these types of cleaners can have a significant impact on corrosion resistance of chromate conversion coatings, which is almost certainly related to residual species on the surface after cleaning and surface morphology. Build-up of transition metal alloying components during these processes is dependent on the etch rate. Etching of copper-containing alloys in NaOH (chemical milling, described below) leads to the development of a heavy copperrich smut on the surface. The loose copper-rich smut is readily removed in HNO3containing solution [Nelson et al. 2001, Dimitrov et al. (1999)]; however, it is not clear whether a copper-enriched layer is present at the Al/Al-oxide interface.

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T. H. Muster, A. E. Hughes and G. E. Thompson

Chemical cleaning in NaOH is moderately reactive, with reports of etch rates of approximately 0.18 and 0.28 μm/min for 0.1M [Liu (2005)] and 0.6M NaOH [Nelson (2001)] respectively. There have been a number of studies of the influence of 0.1M NaOH solutions on aluminium alloys. Liu et al. (2005) have observed copper enrichment and formation of copper nanoparticles in the surface oxide of a sputter deposited Al-30 at% Cu alloy after treatment in 0.1M NaOH. Based on work on anodized coatings and model binary Al-Cu alloys, they have estimated that the maximum buildup is around 1 x 1016 atoms/cm2 [Liu et al 2003]. This level of buildup occurs with the removal of only 5 nm of the alloy [Liu et al. 2005]. Continued etching of the aluminium matrix, resulting in the formation of aluminium oxide, leads to incorporation of copper particles which can be oxidized since they are no longer cathodically protected by the aluminium matrix. Cupric oxide is also a narrow band gap semi-conductor which can display n- or p-semicondiction [Siripala and Kumra, (1989); Di Quarto et al., (1985); Sutter et al. (1995); Millet et al. (1995)]. The electron donor capability of this oxide may mean that it could act as a cathode. Studies, using positron techniques 1 , of oxide/hydroxide formation on relatively high purity aluminium after alkaline cleaning, showed increased number of nano-sized defects, partly related to enrichment of small defects near the surface [(Wu et al. (1994)]. Nelson et al. (2001) examined the surfaces of AA2024-T3, AA6061-T6 and AA7075-T6 alloys after treatment in 0.6M NaOH with a prior treatment in HNO3/HF. They found considerable development of basic oxides on the surface incorporating alloying additions such as magnesium and zinc. The level of copper observed on AA6061-T6 alloy was only slightly lower than AA7075-T6 alloy, despite it having only one quarter as much copper in the alloy content. The level of copper enrichment for AA2024-T6 alloy was roughly three times that of AA6061-T6 alloy, even though it had ten times the bulk copper content, as determined by ICP analyses. However, as pointed out in the work of Roberts et al. (2002), the enrichment underneath the surface oxide may have been greater, but XPS probing was unable to penetrate through the surface oxide. Constituent particles undergo selective phase and elemental etching in alkaline solutions. For example, Lunder and Nisancioglu (1987) found enrichment of iron and manganese on a range of non copper-containing interemtallic 1

Positron techniques include Positron Annihilation Lifetime Spectroscopy (PALS) and Doppler Broadened Energy Spectroscopy (DBES). In both techniques, positrons impinge on a surface and annihilate with electrons within the material. If the positrons reside within a void in the material then they will have a longer lifetime than the rest of the material, which provides some information on the size of the void.

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constituent particles including Al3Fe, Al6(Mn, Fe) and Al6Mn. They also found that the alkaline treatment increased the susceptibility to neutral salt spray for a range of different alloys, with sheet AA5052 being the exception. Preferential removal of aluminium during alkaline treatment should lead to the formation of aluminium-rich oxides on the external surface of the constituent particles with probable enrichment of zinc and magnesium for particles containing these alloying elements. The silicated, carbonate based cleaners [King (1988)] have a much lower etch rate, but still modify the surface. Hughes et al. (1996) observed an increase in magnesium, zinc and silicon with increased immersion time in a silicated alkaline cleaner using XPS. It was suggested that the silicon was present in the form of silicate on the basis of the Si 2p binding energy. Moffitt et al. (2001) observed similar changes for AA2024-T3 alloy as well as AA7075-T6 alloy, although they did not analyse silicon on the surface. Both studies suggested that the levels of copper increased with alkaline treatment above that expected in the AA2024-T3 matrix. Conversely, Toh et al. (2004) did not observe any copper enrichment on AA7475-T7651 alloy after treatment in a mild alkaline cleaner, but did observe incorporation of silicon. These differences are likely to be due to differences in the etch rates of the alkaline cleaners, which are influenced by pH, operating temperature, solution chemistry and the composition of the intermetallic particles. All of these can vary with the product and even during processing leading to varied etch rates. It can be expected that, if there is a high etch rate of the aluminium, and low solubility of either the alloying elements in the etch solution, or compounds formed from the alloying elements in the etch solution, then there will be a build-up of the transition metals including iron, manganese and copper in the oxide coating the matrix as well as transition metals in intermetallic particles.

CHEMICAL MILLING Chemical milling is an alkaline etching process which operates at etch rates of 12 to 37 μm/min [Wernick et al., (1987)]. Chemical milling of aluminium alloys is achieved through the addition of 75 to 150 g/L (approx. 2 M to 4 M) of NaOH. The rate of milling can be modified via the addition of certain chemicals [Wernick et al. (1987)] of which silicate, phosphate and chromate are some of the most common. Milling copper-containing alloys leave a heavy smut on the surface which is removed using a nitric acid treatment [[Wernick et al., (1987), King (1988),

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Nelson et al. (2001), Hughes et al]. Liu et al. (2005) observed the development of copper nano-particles as a result of etching Al-30 at.% Cu alloy. Nanoparticles were formed as a result of enrichment of copper beneath the surface oxide, followed by clustering of copper atoms and occlusion by the oxide film. Other studies on more pure aluminium alloys show that the intermetallics in the surface tend to etch slower than the surrounding matrix [Lunder and Nisancioglu, (1987), Short and Shearsby (1969)]. Indeed, Montiero et al. (1991) found severe etching associated with the interemtallic particles for a commercial purity alloy. Even on pure aluminium, Wu and Herbert (1996) observed buildup of iron, copper and gallium after caustic etching 99.98 % pure Aluminium in 1M NaOH. It was suggested that copper particles were contained in a highly defective 10 nm oxide on the surface of the alloy.

WATER RINSING Relatively little work has been done on the influence of water rinsing, with the majority of the work undertaken some time ago. Clearly, for a chemically active aluminium surface, exposure to water at room temperature and trace ions can influence the amount of oxide and the adsorbed species. Most of the published data is related to the Forest Product’s Laboratory etch (FPL). The nascent oxide had a particular structure that was suitable for adhesive bonding or anodizing [Pocius (1981), Russell and Garnis (1976), Sun et al. (1978)]. Russell and Garnis (1976) demonstrated that the contact electrical resistance increased significantly with the length of the rinse time, with higher resistance generated on purer alloys. Studies of oxide growth in aqueous solution have largely been confined to processes at elevated temperature. At temperatures above 50°C to boiling, oxide growth is dominated by pseudo-boehmite (Al2O3.nH2O, n = 2 or 3) and boehmite (Al2O3.H2O) formation, whereas, at 40°C, bayerite eventually forms the dominant product. [Vedder and Vermilea (1970), Altenpohl (1966), Alwitt ((1974), Underhill and Rider (2005)]. Clearly, surface condition is important since it has been reported that as little as 100 ppm of Si in solution can retard oxide growth by absorption onto the aluminum surface [Vedder and Vermilea (1970), Altenpohl (1966)]. The adsorption of chromate ions on the surface can also significantly impede oxide formation at elevated temperature [Gorman et al. (2003)]. At room temperature, many species, like chromate and phosphate, adsorb irreversibly, thus having a major impact on oxide formation [Heine and Pryor (1967), Ergun et al. (1997), Sotoudeh et al. (1981), Böhni and Uhlig (1969), Badawy and Al-Kharafi (1997), Németh et al. (1998), Vermilyea and Vedder (1970)]. Additionally, oxide

Chemically Pretreated Surfaces

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growth is a thermally- activated process and the driving force for oxide formation will be much less at room temperature [Vedder and Vermilea (1970), Altenpohl (1966)]. Whether the presence of surface copper affects oxide growth in rinse water at room temperature is not well understood. At 40° to 50°C, the ultimate film thickness of oxide films developed in deionised water was similar for a range of alloys [Gorman et al. (2003), Underhill and Rider (2005)]. However, several studies suggest that the initiation of oxide growth may be more rapid with increased copper content [Gorman et al. (2003), Underhill and Rider (2005)]. The improved initiation of oxide growth may arise because the copper, which acts as a net cathode, promotes aluminium dissolution at a greater rate during the initial stages of immersion than on pure aluminium alloys.

DEOXIDATION/DESMUTTING Deoxidation or desmutting is a term used to describe removal of the oxides left after alkaline cleaning, and is intended to be a mild etching process rather than a milling process. The term “deoxidation” is preferred in North America, whereas “desmutting” is preferred in Europe. For the purposes of this chapter we will use the term “deoxidation” as the process has broader chemical implicatons than the simple removal of “smut”. As with any process that results in etching of aluminium alloy surfaces, preferential etching of aluminium results in a build up of alloying components leading to surface enrichment, as well as texturing of the surface. The first part of this section will deal with the types of deoxidisers that are available and their impact on surface enrichment, particularly copper enrichment, and the second part will deal with surface textures.

Surface Chemistry The composition of the surface and the residual chemical species after treatment in deoxidizing solutions varies according to the chemistry of the treatment solution and the etch rate. As mentioned previously, Ketcham and Brown (1976) observed that different deoxidisers had a significant influence on the corrosion resistance of the CrCC. Nelson et al. (2001) also observed a significant deterioration in performance in corrosion resistance of the CrCC on

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AA2024-T3 and AA7076-T6 alloys when a simple HNO3/HF treatment was used instead of a chromate-based deosidizer. This was attributed to a difference in the copper levels on the surface. Liu et al. (2004, 2005) observed that the CrCC coating weight decreased with the enrichment of copper at the surface due to anodic dissolution during coating. Therefore there is evidence that different surface treatments influence the performance and deposition kinetics of coatings such as CrCC. This is likely to be due to the surface chemistry in the form of the passivity of the surface oxides, absorbed species and surface morphology. None of these aspects is a metric used in the metal finishing industry so one of the aims of this section is to draw the connections between the characteristics measured using the sophisticated techniques in the research environment to the metrics used in the metal finishing industry. A review of the literature reveals that there are relatively few studies on deoxidisers in comparison with the numbers on conversion coating and anodizing. A number of studies are related to the Forest Products Laboratory (FPL) etch, which has been used as a preparative treatment for adhesive bonding. Since the surface texture related to this treatment is important for bonding, it is discussed in detail in the next section. Many of the studies relate to current, chromate based deoxidisers [Hughes et al. (1996), Moffitt (2001)2 , Gorman et al. (2003)] and chromate replacements for these deoxidisers which will be discussed below. Examination of the chromate/HNO3/HF deoxidizer indicates that it leaves a thin chromate-passivated oxide on the surface for a range of alloys [Gorman et al. (2003)]. There is also evidence of copper, after deoxidation, for a range of Alalloys, including AA1100-O, AA2024-T3, AA3004-H19, AA5005-O, AA6061T6 and AA7075-T6 alloys, despite some of these alloys having very low bulk copper content [Gorman et al. (2003)] and also some indication of copper enrichment with this deoxidizer on AA2024-T3 alloy [Moffitt et al. (2001), Gorman et al. (2003), Hughes et al. (1996)]. Figure 18 (a) shows a high magnification secondary electron image of the surface after deoxidation in a chromate/HNO3/HF for ten minutes. It is evident that there is a network structure across the surface which is decorated with larger nodules; the network and the nodules are assumed to be part of the oxide covering the surface. An XPS depth profile through this structure is presented in Figure 18 (b) where it can be seen that the copper signal increases with a corresponding decrease in the oxygen signal suggesting copper enrichment at the aluminium2

The Authors assume that Moffitt et al. used a chromate based deoxidiser since the product number they use (Parker Amchem Deoxidiser No 7) is a chromate based deoxidiser, although they do not explicitly note this in their experimental section.

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oxide/aluminium-metal interface [Hughes et al. (1996), Moffitt et al. (2001)]. The depth profile has been divided into three regions: Region I is the external surface which has a high chromium concentration, Region II is the oxide, and Region III is the interfacial region with the underlying metal. Given that the oxide ridges observed in Figure 18 (a) are the thickest part of the coating, and, therefore, the last part of the oxide to be removed during depth profiling, these results suggest that the copper enrichment is greatest at the interface between the oxide ridges and the underlying aluminium as depicted in Figure 19. The results do not indicate whether there is also a continuous film beneath the thinner parts of the surface oxide. The model presented in Figure 19 is not unlike models presented by CaicedoMartinez et al. (2003) for the development of surface textures during chemical polishing of aluminium alloys. However, in this case, the deoxidizer solution removes all alloying components and impurities, effectively leaving only copper enrichment on the surface. Species from the deoxidizer solution are also incorporated into the surface oxide, since it contains chromium and fluorine. Given that chromate in conversion coatings solutions react with the cathodic sites, which Liu et al. (2003) and Brown et al. (1992) have shown to be on the ridges of textured aluminium surfaces, then the chromate in the deoxidizer solution may be reacting with copper-enriched texture on the aluminium surface forming a network of mixed chromium, aluminium oxides or oxyfluorides. Aluminium metal ridges have been included beneath the oxide ridges in a model similar to that of Caicedo-Martinez et al. (2003) where the copper provides some cathodic protection to the metal beneath it and promotes anodic dissolution of the adjacent metal. These metal ridges have not been observed, as yet, experimentally. For the chromate-free deoxidizers, there are a range of compositions. Generally, if a high etch rate is required then the deoxidizer will contain HF. Thus, a simple deoxidizer formulation combines a mineral acid and HF, which is sufficient for low copper-containing aluminium alloys. High etch rate deoxidation can also be achieved with elevated temperature as discussed below. For the purposes of copper-removal, HNO3-based solutions are preferred. Chromate-free deoxidisers, with the ability to remove copper, include formulations based on HNO3/BrO3- [Bibber (1991,1993), Kloet et al. (2005), Toh et al. (2004)], rare earths [Hughes et al. (2001,2003), Kimpton et al. (2000), Campestrini et al. (2004)] and deoxidisers based on HNO3/HF [Montiero et al. (1991,1988), Nelson et al. (2001)], as well as HNO3/HF with added oxidants, i.e, oxidant/ HNO3/HF. A number of these deoxidisers have oxidants that are intended to perform a similar function to chromate, such as Fe3+ ions [Hughes et al.

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(2003,2004)]. The latter forms the basis of several, commercial, non-chromate deoxidizers. The presence of the chromium and fluoride containing oxide after deoxidation, leads into some often overlooked aspects of deoxidation and the relationship that it bears to the subsequent conversion coating step. In the last section, the possibility of oxide growth during water rinsing was mentioned. It is possible that the presence of this thin Cr-Al-F oxide layer, which shows some passivating properties [Hughes et al. (1996)], may prevent oxidation during rinsing after deoxidation, and also provides a nascent oxide structure that promotes rapid growth of the CrCC after immersion in the conversion coating bath. 60

(a) (a)

(b)

50

Al O Cu ( x5 ) Cr ( x5 ) F ( x5 )

Atomic %

40 30

I

II

III

20 10 0

500 500 nm nm

0

100

200

300

400

Sputter Time (sec)

Figure 18. (a) Secondary electron image of the surface of AA2024-T3 after deoxidation in Cr6+/HF/HNO3 and (b) aluminium, oxygen and copper (×5) XPS depth profiles through the surface in (a).

The etch rate of Fe3+/HNO3/HF deoxidizers are lower than that of the corresponding Cr6+/HNO3/HF deoxidizer by about two thirds, depending on the alloy [Hughes et al. (2003)], and result in a slower etching of the surface oxides that remain after alkaline cleaning [Hughes et al. (2004)]. Intermetallic removal is effective; S-phase are completely removed, and the Al-Cu-Fe-Mn intermetallic particles severely etched, although there in no data in the literature on whether there is preferential enrichment as a result of etching. The etch pattern on the matrix, typical of the Fe3+/HNO3/HF deoxidizer is shown in Figure 20, and is similar to the etch patterns on the surface of AA2024-T3 alloy after treatment in the chromate deoxidizer, but leaves iron-containing deposits on the surface

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[Bibber. 1993;, Hughes et al. (2004)] as well as copper enrichment, presumably at the aluminium/aluminium-oxide interface [Hughes et al. (2003)].

Aluminium Oxide Ridges Cu enrichment

Aluminium Oxide

Al alloy Aluminium Matrix

Figure 19. Proposed model of oxide ridges for the surface of AA2024-T3 after deoxidation in Cr6+/HF/HNO3 deoxidizer.

Cerium-based deoxidisers have also been investigated for cerium based conversion coatings [Hughes et al. (1999), Kimpton et al. (2000), Campestrini et al. (2004)]. One low etch rate version of this deoxidizer results in little more than preferential etching of Mg-oxides from the surface oxide. However, a high etch rate composition of this deoxidizer has the oxidant/ HNO3/HF formulation with Ce4+ being the oxidant, but it has an additional mineral acid in H2SO4. Both Hughes et al. (1999) and Campestrini et al. (2004) observed the retention of copper after immersion in a high etch rate Ce4+/H2SO4/HNO3/HF deoxidizer. Both studies observed the deposition of copper particles (200 nm in dimension) from the deoxidizer onto the matrix as well as onto the intermetallic particle remnants, which additional oxidants (H2O2, K2S2O8) assisted in removing. The selective dissolution of species (aluminium, magnesium, manganese, iron and silicon) can leave a range of structures during deoxidation. In deoxidizers containing HNO3 (and HF), there tends to be even etching of the intermetallic particles (Hughes et al. (1996, 2001, 2003), as well as a high rate of intermetallic removal from fallout due to undermining by etching in the presence of HF [Campestrini et al. (2004)]. Copper enrichment occurs when either other mineral

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T. H. Muster, A. E. Hughes and G. E. Thompson

acids are operated at elevated temperatures, or the deoxidizer composition does not include HNO3.

(a)

(b) Figure 20. a) Secondary electron image of the surface of AA2024-T3 alloy after deoxidation in Fe3+/HF/HNO3 and (b) surface reconstrtuction from strereopair of another part of the surface showing the scalloping (etching) pattern.

63

Chemically Pretreated Surfaces (a)

(b)

10 μm

20 μm

(d)

(c)

2 μm

5 μm

Figure 21. Electron micrographs of Cu-Fe-Mn-Al containing intermetallic particles after deoxidation at ambient temperature in various deoxidisers. (a) BrO3-/HNO3, (b) Fe3+/HNO3/HF, (c) Ce4+/H2SO4/HNO3/HF and (d) Na2S2O8/H2SO4/HF.

For example, Figure 21 shows some backscattered and secondary electron images of etched Al-Cu-Fe-Mn type intermetallic particles. In Figure 21 (a) and (b), where strong oxidants (i.e. BrO3- or Fe3+) are used in the deoxidizer formulation, there is even etching of the intermetallic particles, but no significant copper enrichment of the intermetallic particle remnants. In the case of the Ce4+/H2SO4/HNO3/HF deoxidizer, there has been selective removal of aluminium, iron and manganese leaving a copper sponge (Figure 21(c)). This sponge appears to have precipitated copper from solution on the surface. Where Na2S2O8/H2SO4/HF was used (Figure 21(d)), there is also selective removal of iron, manganese and aluminium, leading to copper-enrichment of the intermetallic particle, but no sponge was apparent in these particles, suggesting a different mechanism of copper enrichment. In this last case, there was no detectable copper in the deoxidizer solution indicating limited opportunity for deposition onto the intermetallic particles of dissolved copper from solution. In these latter two cases,

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the presence of SO32- appears to have a significant influence on the mechanism of dissolution, but the details have not yet been revealed. Where the copper becomes electrically isolated, presumably either through the oxidation of the underlying alloy, effectively electrically isolating the coppersponge from the aluminium matrix, or through undermining due to etching, the particle can dissolve to form Cu2+, which becomes available for deposition onto the surface. As stated previously, deposition (or plating out) has not been observed with HNO3-based deoxidisers, however, as seen above, it can occur with other formulations. There are a few studies that attempt to make the connection between etch rate and changes to the surface as a function of chemistry. Hughes et al. (2001,2003) examined the attack on AA2024-T3 alloy using H3PO4, H2SO4 and HNO3, alone and with additions of HF, and HF/Na2S2O8. In the mineral acids alone, at ambient temperature, the matrix was protected by the surface oxide, and the intermetallic particles were clearly protected by a layer of aluminium matrix, deposited during rolling [Lunder and Nisancioglu (1987)]. Etching of the matrix can be achieved at elevated temperatures (60°C) and a heavy copper smut can be developed on the surface [Campestrini et al. (2004)]. Heavy copper smut can also be developed at room temperature with the addition of an etchant, like HF, to the mineral acids in formulations, such as H2SO4/HF, H3PO4/HF, H2SO4/HF/Na2S2O8 and H3PO4/HF/Na2S2O8,. This occurs as a result of etching of the matrix without dissolution of the copper. It was also noted that the addition of Na2S2O8 as an oxidant had a minimal affect on reducing the copper levels [Hughes et al. (2001)]. In an endeavour to delineate the relationship between etch rate and the observations of various surface techniques, Hughes and co-workers (2001,2003) have proposed a qualitative model to describe the stages of attack by deoxidisers on a rolled aluminium alloy surface. The process proceeds in three stages: Stage 1:

Stage 2:

Upon immersion in the acidic deoxidiser solution, there is preferential dissolution of components of the oxide remaining after alkaline cleaning. These include the basic magnesium and zinc oxides and also silicon-containing phases. Low etch rate deoxidisers (Ce(IV), HNO3/BrO3- at 20°C or simple mineral acids) generally do not proceed beyond this Stage unless the operating temperature is elevated substantially. Deoxidisers with any significant etch rate, proceed beyond Stage 1 into Stages 2 and 3. During Stage 2, the surface oxide left after alkaline cleaning is completely removed and etching of the underlying alloy begins. This is a complex process since the alloy

Chemically Pretreated Surfaces

Stage 3:

65

contains a number of alloying additions in a variety of phases and each of these components must reach an equilibrium between accumulation and dissolution. Stage 2, therefore, is the intermediate stage in which the dissolution of aluminium leaves behind an accumulation of alloying components. Stage 2 is reached by deoxidiser combinations containing HF and also BrO3/HNO3 at 40°C or 60°C and probably H2SO4 deoxidisers operated at increased temperatures such as those used in adhesive bonding since these give the characteristic network of oxide ridges. The balance between surface oxide formation and dissolution moves towards thinning of the oxide and the build-up of alloying components, (particularly copper). It is not clear whether the rate of oxide thinning is delayed by the Cu(I)/Cu(0) redox couple, although data on low copper-containing alloys may help to resolve this issue. The presence of oxidants in the system may assist in copper dissolution (but this is likely to depend on the details of the chemistry) so that an equilibrium concentration might be obtained. In Stage 3, alloy component dissolution reaches an equilibrium between dissolution and accumulation. The Cr(VI) and Fe(III) deoxidisers, described earlier, (which contained HNO3, in combination with HF) and the BrO3-/HNO3 deoxidiser at 60°C reach Stage 3. Similar results were observed for rare-earth based deoxidisers when they contained fluoride.

The HNO3/BrO3- deoxidizer provides a good example of all three stages as a function of temperature and time. Figure 22 shows that at 20°C that there is very little attack on the matrix, but there is significant attack on the intermetallic particles; after 10 minutes the particles were nearly completely removed. Thus, the attack on the matrix oxide is confined to Stage 1. Attack on the matrix was evident at 40°C where, after 10 minutes, the deoxidizer had moved through Stage 2 and commenced Stage 3 since the characteristic etch patterns were observed. At 60°C the deoxidiser moves quickly into Stage 3, where the network of oxides ridges is apparent after deoxidation for only 5 minutes. To further illustrate the stages of deoxidation, the etch rates for a number of different aluminium alloys in chromate/HNO3/HF deoxidizer are shown in Figure 23. It is clear that the etch rates are higher initially for the AA7xxx series alloys and lower for AA2024-T3 alloy, but, subsequently, both alloys reveal similar values. The initial differences, during Stage 1, represent changes in the dissolution

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of the surface oxide. The 7475-T7651 alloy remains in Stage 1 and 2 the longest because this particular batch of alloy had a thick surface oxide [Toh et al. (2003)]. Eventually all the alloys display similar etch rates in Stage III which represent the dissolution of the substrate aluminium, where copper enrichment and intermetallic dissolution occur. A qualitative estimate of the three stages of dissolution is also included in Figure 23.

20ºC (a)

(b)

(c)

(d)

(e)

(f)

(g)

(h)

(i)

40ºC

60ºC

1 min

5 min

10 min

Figure 22. Scanning electron micrographs of AA2024-T3 alloy after deoxidation for various times and temperatures in BrO3-/ HNO3- based deoxidiser. (Scale marker for (c) = 1 µm, for all other images = 500 nm.). Reproduced with permission from Hughes et al. (2003).

A qualitative model, similar to that described, could be made more quantitative by combination of the knowledge of copper build-up under anodic coatings (see below), enabling prediction of the levels of copper-enrichment after the pretreatment process. The two main issues that need to be addressed here are: (i) the buildup of copper at the aluminium/aluminium-oxide interface, and (ii) the removal of intermetallic particles from the surface. Hence the copper-enrichment could be described by: Cuenrich = f(EM, EIM,Dcu)

…13

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Chemically Pretreated Surfaces

where EM is the etch rate of the matrix, EIM, is the etch rate of the intermetallic particles and Dcu is the amount of redeposited copper. To determine the influence of these terms on copper enrichment, it would be necessary to know the copper content of the matrix, as well as of individual particles. It would also be necessary to know the influence of the solution chemistry on etching of the intermetallic particles. If it is assumed that these three variables are independent, then equation 13 can be separated into the sum of three independent terms: Cuenrich = f(EM) + f( EIM) + f(Dcu)

….14

19 18

I

Etch Rate (μm/h)

17

II

III

16 15 14 13 12 11 0

5

10

15

20

Time (mins)

Figure 23. Etch rates for chromate-based deoxidizer three different aluminium alloys. ● = 7475-T7651, ■ = 2024-T3 and ▲= 7075-T6. I= stage I, II = Stage 2 and III = Stage 3.

Clearly, terms like f(EIM) and f(Dcu) would need to incorporate the surface area of the copper and well as the coverage of the aluminium. For a number of deoxidisers such as the chromate/HNO3/HF and HNO3/BrO3- (60°C) deoxidizers, where the intermetallics are virtually completely removed and there is no significant redeposition of copper then equation 14 is simply related to the buildup of copper under the surface oxide. For more complicated systems, specific models would need to be developed for the enrichment of copper due to intermetallic particle dissolution and copper redeposition. A simpler approach to

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T. H. Muster, A. E. Hughes and G. E. Thompson

populating the terms of equation 14 is through the use of a look-up table which gives a figure for a selected acid combination under specified conditions. 40 nm

5 nm 40 nm

Oxide

5 nm

Cu-enrichment

Aluminium

Figure 24. Model of surface oxide structure based on scanning electron microscopy stereo pairs after Venables et al. (1979). (Cu-enrichment has been added by the authors and did not appear in the original paper by Venables et al. (1979))

Figure 25. Copper enrichment levels are the electropolishing of aluminium-copper binary alloys as a function of copper content in the alloy (after Liu et al. (2003)).

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Surface Structure Many copper-containing aluminium alloys develop surface textures or structures comprising networks of ridges across the surface during electropolishing or chemical treatment (e.g. Figures 20 or 21). Liu et al. (2003) have shown that the individual ridges consist of a metal ridge covered by an oxide. A number of models have been put forward for the structure of the ridge network on the surface [Brown (1949), Hunter and Robinson (1973)]. At this point, it seems to be controlled by the distribution of alloying elements in the surface [Thompson et al. (1987), Caicedo-Martinez et al. (2003)]. These types of structures are extremely important for adhesive bonding applications, particularly in the aerospace industry. Work by Venables et al. (1979) suggested that a particular type of nodular oxide structure (Figure 24) promoted good bond strength but was susceptible to degradation during storage, particularly if fluoride was present. In early work, it was recognized that this structure was somehow related to copper in the processing bath or the copper content of the alloy [Pocius, (1981), Sun et al. (1978, 1980)], but precise details have not been determined. Based on the work of Caicedo-Martinez et al. (2003) the authors have proposed that there is copper enrichment beneath the oxide ridges which facilitate dissolution of the adjacent matrix; these have not, as yet, been observed experimentally.

ELECTROPOLISHING The electrochemical polishing of a metal is a process where a DC current is applied to oxidize the metal and remove it from the surface. Leveling occurs due to the field gradient across the surface, which results in higher dissolution rates at raised areas of the surface than in troughs. According to Thompson et al. (1987), an anodic oxide film forms on aluminium during electropolishing; the film develops at the metal/film interface through O2- ingress and with the outwardly mobile Al3+ ions ejected into the electrolyte under the influence of the electric field. In addition, a through-film dissolution of aluminium ions occurs, along with dissolution of the outer oxide regions, which are exposed to the reactive electrolyte. For aluminium-copper alloys, preferential oxidation of aluminium enables development of a copper enriched layer in the alloy immediately beneath the film developed in the electropolishing bath. Copper enrichment at the alloy surface during electropolishing has been observed to reach levels corresponding to

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40 at%, which are contained in a layer approximately 2 nm thick (Habazaki et al., 1995). The thickness of the enriched layer is thought to be independent of bulk copper concentration and electropolishing conditions. However, the amount of copper contained within this thin enriched layer can vary significantly. Liu et al. (2003) electropolished aluminium-copper binary alloys ranging in copper concentrations from 0.0025 at.% to 1.5 at.% at 20 V in perchloric acid/ethanol solution (20/80 v/v) at 278 K for various times. They reported metal removal rates of 40 nm s-1. Copper enrichments at the surface of each binary alloy were independent of time between one minute and fifteen minutes electropolishing, indicating that all enrichment occurred within the first minute. Figure 25 shows that the amount of copper enrichment was found to increase rapidly up to 0.4 at% copper, whereas alloys containing greater amounts of copper showed a relatively weak enrichment. Liu et al. (2003) introduced an enrichment factor, defined as the ratio of copper atoms in the enriched layer (atoms cm-2) compared to that of the bulk concentration (at%). Therefore, the enrichment factor decreases for alloys with increasing copper concentration (Figure 25). If significant copper enrichment is achieved, copper-rich clusters formed at the alloy surface may be released into the oxide film, which leads to smutting of the electropolished surface. As described previously, nitric acid etchants are generally used to remove copper smut. For instance, Liu et al. (2003) used a 50 wt% HNO3 solution at ambient temperature for 60 s to remove smut after electropolishing laboratory samples. It follows that decorative aluminium alloys used to present bright surface finishes usually have a low amount of copper. Typically, AA5xxx series alloys containing minimal copper are used, such as; AA5005, AA5050, AA5252, and AA5657 [Hatch (1984)].

ANODISING Anodizing is an electrochemical process that, in appropriate electrolytes, generates a cellular-structured, porous anodic oxide film on the surface of aluminium. Anodizing on copper-containing alloys is reported to show reduced film growth rates due to the preferential dissolution of intermetallic compounds [Takahashi et al. (2003)]. Unlike acid electrolytes, where porous anodic films are generally formed, Thompson et al. (1987), examined barrier-type anodic film formation on aluminium-copper alloys in near-neutral electrolytes. Anodic oxide growth proceeds at both the solution-film and alloy-film interfaces under the electric field

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[Habazaki et al. (1996)]. Habazaki et al. (1996) reported that about 40% of the film thickness was formed at the film/electrolyte interface due to the outward diffusion of Al3+ and the remainder is formed at the metal/oxide interface through the inward diffusion of O2-/OH- species. For binary solid solution alloys the efficiency of film growth is only marginally influenced by the development of a layer of copper-enrichment in the alloy, which occurs immediately adjacent to the alloy-film interface [Habazaki et al. (1995), Liu et al. (2004)]. During anodizing of such alloys, it has been demonstrated that the anodic film develops initially in the absence of incorporated copper species, due to preferential oxidation of aluminium. Such anodic oxidation allows copper to accumulate at the alloy/film interface until it reaches a concentration of about 40 at% in a layer of thickness of about 2 nm. Habazaki et al (1995) used RBS to demonstrate that the extent of copper enrichment at the interface was independent of the anodizing conditions and was similar to the enrichment developed by electropolishing of an alloy of equivalent copper content. At the critical level of enrichment, copper oxidizes at the alloy/film interface and enters the oxide in its alloy proportions in the continued presence of the enriched layer. The applied voltage was also found to influence the incorporation of copper into anodized films (Table 8), where increased voltages up to 300 V led to a linear increase in copper enrichment at the alloy surface. For the particular conditions used by Habazaki et al. (1996), 140 V was calculated as being a critical voltage, which, if exceeded, would lead to the incorporation of copper into the anodized film. Table 8. Compositional analysis determined by RBS of enriched alloy layer and anodized films formed under varying applied voltage at 5 mA cm-2 in 0.1 M ammonium pentaborate at 293 K (Habazaki et al., 1996) Anodizing voltage (V) 15 50 150 300 1

Copper in enriched alloy layer (at%)1 4.8 16.0 41.2 Not determined

Atomic aluminium:copper in anodized film (× 10-3) ≤ 0.2 ≤ 0.2 ≤ 0.2 4.4

The bulk alloy copper content was 0.9 at%.

Where a significant enrichment of copper has occurred at the metal/film interface, oxygen gas-filled voids develop in the barrier film through the semiconducting nature of the Cu(II)-O bond; the gas pressure achieves very high levels, with eventual rupture of the surrounding alumina. Interestingly, the enrichment and consequent gas generation appears to be alloy-grain orientation

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dependent [Zhou et al. (1999)]. In separate studies, where such gas generation was eliminated by anodizing of thin layers of the copper-containing alloy, the outward mobility of Cu2+ ions was shown to be 2-3 times that of Al3+ ions, which gives rise to their loss at the film/solution interface and generation of an anodic film of copper content less than that of the bulk alloy. For the bulk alloy, the consequence of gas generation and film rupture is to generate a heavily flawed anodic film. Concerning porous anodic film formation, Shimizu et al. (1997) examined anodizing of a binary Al-Cu alloy that had been artificially aged to develop a fine distribution of θ’ precipitates. Porous anodic film formation, proceeding exclusively at the alloy-film interface due to O2- ion ingress under the field, led to recession of the alloy/film interface. However, when θ’ precipitates of relatively high copper content relative to the adjacent matrix where encountered, immediate gas generation developed to sufficiently high pressures that ruptured the anodic film. Subsequently, film repair proceeded, removing the θ’ precipitate and generating a porous anodic film of enhanced porosity.

(a)

(b)

Figure 26. Transmission electron micrographs of (a) unclad 2014 T3 alloy and (b) clad and after anodizing in sulphuric acid, revealing the resultant anodic films of contrasting morphologies (courtesy of Prof. G. Thompson).

Figure 26 reveals transmission-electron micrographs of anodic films formed on clad and unclad AA2014-T3 alloy in sulphuric acid. A distinct affect of copper on the alloy is revealed in the anodic film which displays increased porosity compared with that on the comparatively pure clad material. This is related to

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copper enrichment in the alloy, gas generation and possible film rupture and film repair. Interestingly, additional mechanisms of anodic film formation are under consideration, which question the established mechanism of film growth at the aluminium/film interface that has been thought to proceed through dynamic equilibrium with field assisted dissolution at the pore base [Garcia-Vergara et al. (2006)]. That is, for anodic films formed in acid electrolytes (eg. sulphuric acid and phosphoric acid with significant amounts of incorporated electrolyte-derived anions) stress induced flow of the film material is being explored to explain pore formation. In other words, plastic flow of the anodic film material is thought to proceed from the barrier layer beneath the pore base to the adjacent cell material. For the film formed on the AA2014 alloy shown in Figure 26, it appears that the lateral porosity or layered film regions are associated with a cyclic oxidation of copper from the enriched layer, in addition to the consequences of oxide plasticity and flow under the field [Iglesias-Rubiane et al. (2006)].

CONVERSION COATING Conversion coating is a term used to describe a coating process that transforms (“converts”) the natural oxide on the surface of aluminium to an oxide with more desirable properties, such as improved paint adhesion and corrosion resistance. The most commonly used conversion coatings are chromate (and chromate-phosphate) conversion coatings (CrCC), which are a family of closely related, commercially-available processes. Thus, products for cleaning, deoxidation and conversion coating are available for a range of applications, but since this review is concerned with copper distributions, the focus is on the processes used for treating high-strength, copper-containing alloys, which are used extensively in the aerospace industry. Non-chromate coating formulations are also available [Buchheit and Hughes (2003), Nylund (2000)] with the Ti/Zr processes [Knudsen et al. (2003), Deck and Reichgott (1992), Lunder et al. (2004), Tomlinson (1997)] and a Co-based process [Schreiver (1992), Schreiver (1996), Roland (1998), Hughes et al. (2004)] being the most common; however, these do not perform well on Copper-rich alloys [Chalmers (1995)]. There are also a number of other processes such as permanganate processes [Bibber (1991), Hughes et al. (2006)], silane treatments, self-assembled monolayers and rare earth processes [Schmidt-Handsberg and Schubach (2003), Rivera et al. (2004)], the last of which particularly benefits from the presence of copper at the surface [Hughes et al. (1995,2004), Palomino et al. (2006), Campestrini et al. (2001)].

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Conversion coating is a multi-step process involving a number of the cleaning steps described previously, prior to the deposition of the conversion coating. The cleaning steps usually involve a mild alkaline cleaning followed by a rinse and then, for copper-containing alloys, treatment in a HNO3/HF deoxidizer usually contains chromate 3 . As stated previously, in aerospace applications, the CrCC process works best with a chromate-based deoxidizer prior to conversion coating [Ketcham and Brown, 1976]. Further, the use of nitric acid in the deoxidizer formulation for copper-containing alloys generally removes copper associated with the smut and the vast majority of the intermetallics from the surface [Hughes et al. (1996,2003)]. Thus, the only copper of any significance on the surface is localised in an enriched layer between the aluminium oxide and the aluminium metal [Moffitt et al. (2001), Hughes et al. (2001)]. Despite the fact that the Cr(VI)/HNO3/HF pretreatment removes most of the copper from the surface, the alloying content (particularly copper) of aluminium alloys still has a marked influence on the coating weight. As pointed out in several texts [Ketcham and Brown (1976), Buchheit and Hughes (2003), Eppensteiner and Jenkins (1995)], the amount of deposition of the CrCC is greater with purer alloys, and lower for highly alloyed aluminium. Thus, copper-containing alloys of commercial importance, such as AA7xxx and AA2xxx series, can be expected to have some of the lightest coating weights of the aluminium alloy family. Typical coating times are up to a few minutes for most alloys, with increased coating times leading to some unusual behaviour. Liu et al. (2005) observed that at longer immersion times (up to 20 minutes) the amount of deposited CrCC on AA2014-T6 and a range of magnetron sputtered binary Al-Cu alloys, decreased as a direct result of an increase in the surface copper levels. Trathen et al. (1993) reported that the corrosion resistance and adhesive failure of a paint system on the CrCC both showed unusual behaviour in the time to pitting in neutral salt fog testing as well as adhesive failure, across a range of alloys for coating times of 10 minutes. There have been numerous studies of the deposition of the CrCC onto aluminium-alloys generally, but fewer on the deposition onto copper-containing aluminium-alloys. At the most basic level, film growth proceeds through contact of the alloy with the coating solution where the fluoride attacks and thins the surface oxide. With sufficient thinning electron tunneling may proceed [Brown et al. (1993), Katzman et al. (1976)]. Fluorine has been detected at the CrCC/metal interface [Hughes et al. (1997), Vasquez et al. (2002), Abd Rabbo et al. (1978), 3

For Al alloys with low copper content, a mineral acid/HF mixture is commonly used such as H3PO4/HF.

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Treverton and Davis, (1977)], suggesting that fluorine-containing species such as AlF3 or AlOF have formed at the interface. Once the reaction starts then the mechanism of coating in its most basic form is described by [Katzman et al. (1979)]: Cr2O72- + 2Al0 +2H+ + 6HF → 2AlF3 + 2CrOOH + 3H2O

....14

The reaction scheme can be elaborated to include the effects of thinning of the surface oxide on the aluminium prior to the deposition reaction. As the reaction proceeds, more chromium species are deposited, however, a limiting thickness is eventually reached which depends on bath chemistry. It should be noted that there are two classes of CrCC processes: unaccelerated and accelerated. The unaccelerated processes contain a mixture of HF, chromic acid and nitric acid and are operated at a pH below 2. Accelerated processes, in addition to the usual components have an added accelerator, the most common of which is K3Fe(CN)6 [Xia and McCreery (1999)]. The coating weight [Arrowsmith et al. (1984)], coating thickness [Juffs et al. (2002), Katzman et al. (1979), Campestrini et al. (2001)] or intensity of characteristic lines from the conversion coating [Schram et al. (1998), Kendig et al. (1993), Drozda and Maleczki (1985), Xin and McCreey (1999), Treverton and Amor (1985)] increase with time. For unaccelerated coatings, at longer coating times, the coating weight tends to stabilize [Katzman et al. (1979)]. For accelerated coatings (those containing K3Fe(CN)6), the coating weight also becomes self limiting, but, at longer immersion times, the coating becomes loose and powdery with poor adhesion. Typical coating thicknesses for the non-accelerated and the accelerated formulations are shown in Figure 27. These data are derived from several sources as listed in the figure caption, and, generally, for five minutes immersion time in the coating solution. It is evident that the coatings from the accelerated formulations are about an order of magnitude thicker than those from the unaccelerated solutions. Coatings on copper-rich intermetallic particles are much thinner than the surrounding matrix. It should be noted that there is considerable variation in the coating thickness with the accelerated coating. For example, Treverton and Amor (1985) estimated a coating thickness of over 1 micron after only 3 minutes coating. Hagans and Hass (1994) found lower values, but than others, this may be due to the low pH (pH 1) of the coating solutions that were used. Osbourne (2001) has show that virtually no coating is developed at pH 1 whereas maximum coating weights are obtained around pH 2.

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T. H. Muster, A. E. Hughes and G. E. Thompson 1200 accelerated no acceleration

Coating Thickness (nm)

1000 800 600 400 200

b b

FeAl3

FeAl3

a

FeAl 3 FeAl3

CuMgAl 2 CuMgAl2

CuAl2

Cu2FeAl7

Cu2FeAl7

b c a c

CuAl CuAl2 2

a

CuAl2

--

e f

Matrix

d

--

b c

Matrix

Matrix

a

--

0

Figure 27. Coating thicknesses for a range of surfaces in unaccelerated and accelerated coating solutions. a = Juffs et al. (2002), b = Hagans and Haas (1994), c = Vasquez et al. (2002), d = Katzman et al. (1979), e = Sun et al. (2001), f = Treverton and Amor (1985). In the case of Vasques et al. (2002) the matrix value from Hughes et al. was used as the 100% value and the coating thickness over the IM phases was estimated accordingly.

The coating thickness on the intermetallic particles is similar for both the unaccelerated and the acceleterated processes as can be seen in Figure 27. From Figure 28, for the unaccelerated coating, it can be seen that the coating develops over the matrix more quickly than over the intermetallic particles after five seconds immersion, which agrees with the results of Hagans and Haas (1994) for shorter coating times. Clearly, the coating thickens over the intermetallic particles and develops to the same thickness as the matrix (~ 100 nm) over longer times as seen for five minute coatings in Figure 27. In the accelerated process, the coating over the matrix is nearly an order of magnitude thicker than the intermetallic particles, which have a similar thickness to the coating in the unaccelerated process. Hagans and Haas (1994) suggested that this was due to the formation of copper-ferrocyanides on the surface of the copper-containing intermetallic

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particles. Cyano species were also observed on iron-containing species, but the particular phase was not identified. This is discussed in more detail below.

Figure 28. Atomic force microscopy of the surface of AA2014 alloy after 5 seconds at ambient temperature in a chromate conversion coating solution (0.8gl-1 NaF , 4.0gl-1 CrO3 and 3.5gl-1 Na2Cr2O7). Reproduced with permission from Liu et al. (2001).

The role of the ferricyanide accelerator has not yet been fully elucidated although there are two models for its role. The first, proposed by Treverton (1981), suggests that ferricyanide preferentially adsorbs onto the surface of chromium oxide/hydroxide gel particles as they form at or near the surface. The adsorbed ferricyanide then blocks the sites for chromate absorption onto the gel particles; thus more Cr(VI) is available for reaction with the aluminium surface. The second model, proposed by Xia and McCreery (1999), considers the Fe(II)/Fe(III) couple to act as a catalyst (redox mediator in their terminology), with reduction of Fe(III) to Fe(II) during aluminium oxidation and being oxidized from Fe(II) to Fe(III) by Cr(VI) reduction. These reactions facilitate the generation of Cr(III) hydroxyoxide on the surface. Xia and McCreery (1999) also examined a number of other catalysts, such as IrCl62-, Fe3+ and V3+, which accelerated the deposition of the CrCC, but none was as effective as K3Fe(CN)6. Hagans and Haas (1994) examined the influence of ferri/ferro cyanide

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accelerators and confirmed that they resulted in coating thicknesses about an order of magnitude greater than the unaccelerated formulations on the matrix (Figure 27). However, the nature of the surface at which reactions occur are unclear in these models. The issue relating to the nature of the reacting surface has, to some extent, been clarified by Osborne (2001), who discussed the nature of chromate conversion coating in a sol gel context. He suggested that at low pH (< pH 1.2) individual sol particles are formed adjacent to the surface, whereas, at more moderate pH a gel may form on the surface. Thus, the reacting surface is probably either the external surface of the sol particles at low pH, or the surface of the gel coating at more moderate pHs. The other area of debate is the nature of the cyano complexes on the surface; again, while there is strong evidence for certain species e.g. Berlin green [Xia and McCreery (1999)], the isomerisation reactions between iron and chromium cations in hexacyano complexes are complex, and there are probably a number of species present at the surface [Basset Brown et al. (1968)]. The literature on the relationship of copper distributions on the deposition mechanism is complicated by the differing methods of preparation of the aluminium alloys. In an effort to understand the mechanism of deposition on wellcharacterised substrates, many studies have focused on polished aluminium alloy surfaces [Hagans and Haas (1994), Xia and McCreery (1999), Juffs et al. (2002), Vasquez et al. (2002), Liu et al. (2000)], as well as on polished bulk intermetallics [Juffs et al. (2001,2002), Vasquez et al. (2002), Lunder et al. (2005)]. There are relatively few studies that deal with the commercial method of preparation, including appropriate alkaline cleaning, deoxidation and conversion coating, as used on aerospace alloys [Hughes et al. (1997), Lyttle et al. (1995)]. Several studies use chemical pretreatment steps prior to deposition of a coating [Treverton and Davies (1994,1997), Schram et al. (1998), Arrowsmith et al. (1984), Katzman et al. (1979), Drodza and Maleczki (1985), Campestrini et al. (2001), Sun et al. (2001), Meng and Frenkel (2004), Kloet et al. (2005), Vasquez et al. (2002)]. Some of these works use a deoxidation step designed to increase the surface levels of copper. The difference between these various starting surfaces (polished or deoxidized/copper-enriched) may have major implications for the application of the findings to properly deoxidized surfaces. For example, Sun et al. (2001) showed that HF/H2SO4 pretreatment of polished surfaces led to a level of copperenrichment. After conversion coating this copper-enriched surface, copper was found on the external surface of the CrCC as well as the CrCC/metal interface. In the same study, no copper was found on the external surface of the CrCC on polished surfaces. Thus the copper left on the surface after HF/H2SO4 pretreatment had migrated onto the CrCC external surface during the conversion

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coating process. Copper-enrichment may occur via a mechanism similar to that described in corrosion [Zahavi et al. (1978)] and anodizing [Habazaki et al. (1975)] where small copper nanoparticles detach from the enriched layer and are incorporated into the developing coating during its formation. With regard to the work on polished surfaces, it is important to note the differences between them and the chemically pretreated surfaces. Polished surfaces have a much higher exposed area of intermetallic particles than the chemically treated surface used in practical applications, typically 3 to 6% for AA2024-T3 alloy [Hughes et al. (2006), Jakab et al. (2005), Buchheit et al. (1997)], but probably less for other alloys. The oxide covering the matrix is a very thin aluminium oxide [Nylund and Olefjord (1994)] and thus compositionally different to that produced on the rolled surface or the deoxidized surface, as described previously, with its GRSL (see section on surface microstructure). Thus, the polished surface is likely to be more active than the deoxidized surface where the intermetallic particles have been removed. The drive to understand the deposition process on these polished surfaces has led to a number of studies of deposition of chromate onto intermetallic phases and significant advances have been made in the understanding of deposition onto these types of phases [McGovern et al. (2000), Juffs et al. (2001,2002), Vasquez et al. (2002)]. Figure 27 shows coating thicknesses over intermetallic phases is generally thinner than the matrix in the accelerated case. McGovern et al. (2000) studied the deposition of CCC’s onto an Al-Cu-Mg ingot containing various AlCu-Mg phases. Raman spectroscopy showed that the intensity of the 860 cm-1 band, indicative of a mixed chromium oxide, decreased over phases with higher copper content. They also noted larger carbon and nitrogen peaks at the sites of copper-containing intermetallics using AES. Vasquez et al. (2002) examined deposition onto various intermetallic phases in AA2024-T3 alloy, as well as onto intermetallic phases manufactured by laser ablation, and observed the deposition of similar species onto all phases. Hagans and Haas (1994) used AES, XPS and ion-beam depth profiling to study CCC formation on AA2024-T3 for formation times of up to 3 min. The rate of film formation was faster over the matrix and slower over intermetallic phases. It was suggested that ferrocyanide in the conversion coating formulation interacted with copper-rich intermetallic phases to form compounds such as Cu4Fe(CN)6 and Cu2Fe(CN)6 on the surface. All these studies show that the oxide covering the intermetallic phases is much thinner. Copper appears to have a major role in limiting the thickness of the coating on these phases, with Al2Cu having the lowest coating thickness.

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In terms of the coating on the matrix, it was noted above that coppercontaining alloys tend to have a reduced coating weight. As stated previously, Sun et al. (2001) observed different copper distributions in the CrCC as a result of different pretreatment conditions, but the model of the coating for AA7075-T6 alloy was similar to that proposed by Hughes et al. (1997) for AA2024-T3 alloy and Treverton and Davies (1981) for pure aluminium. This model is slightly different to that proposed by Vasquez et al. (2002) for AA2024-T3 alloy; they observed some copper on the external surface of the CrCC due to the presence of copper-containing intermetallics. A polished surface was used without deoxidation and, therefore, the intermetallic phases had not been removed. More recently Meng and Frankel (2004) examined a number of AA7xxx series alloys in (i) polished and (ii) polished and acid etched conditions. They used an HF/H2SO4 acid etch to artificially increase the copper content at the surface prior to chromating. Hughes et al. (2001) found that this type of pretreatment leads to a loose copper-containing smut on the surface of AA2024-T3 alloy, which is not typical of copper buildup on the alloy during commercial processing in a HNO3 based deoxidiser. However, Meng and Frankel (2004) found two breakdown potentials in polarization curves of the polished AA7xxx surfaces. The first was associated with the dissolution of hardening precipitates while the second was combined intergranular and selective grain attack. The breakdown potential increased with copper content of the AA7xxx series alloy for the polished surface, indicating that the increased copper within the alloy improved the corrosion resistance for both the polished, and the polished and conversion coated alloys. The first breakdown was not observed in the conversion coated specimens, indicating that either the conversion coating treatment had removed these phases or that they had been passivated. A detailed examination of a number of these works seems to suggest that the CrCC process itself does not significantly change the prevailing copper distribution on the alloy surface. The kinetics of the coating process, however, is strongly influenced by the presence of copper. The work of Sun et al. (2001) best illustrates this point, since no enrichment of copper was observed during chromate conversion coating of polished surfaces of AA2024-T3 alloy, whereas copper, already present on the surface due to deoxidation, remained enriched at the metal surface beneath the CrCC, and some of this copper was present on the external surface of the CrCC. These results suggest that, while the fluoride ions in the coating solution (equation 14) are intended as an etchant for the underlying metal, they only thin the surface oxide so that electron tunneling through the oxide can promote deposition reactions, rather than having a strong etching effect on the underlying aluminium. The role of fluoride as an etchant for the oxide is

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reinforced by the work of Campestrini et al. (2004). They anodized aluminium prior to conversion coating, and demonstrated that deposition of the conversion coating only began after the anodized layer was removed by fluoride attack. Indeed, it is often overlooked that for the CrCC process, a chromate deoxidizer is preferable. The composition of the deoxidizer solution is similar to the CrCC, but more acidic, and the deoxidizer leaves a surface oxide rich in chromium containing species; these could well be a nascent CrCC, so that the subsequent immersion in the CrCC solution completes the coating process but at higher pH where deposition reactions occur. The fact that the CrCC may not change the copper distribution means that the role of a copper-enriched layer on the deposition process needs to be examined. One view [Vasquez et al. (2002)] is that the presence of copper on the surface simply reduces the available aluminium surface area for the reaction with the conversion coating solution. For this model to work, a surface coverage effect could only be obtained if the enriched copper layer at the alumnium/aluminiumoxide interface comprised thin islands (assuming appropriate deoxidation) on the surface, rather than a continuous layer. (In the latter case the whole surface would be inactive and no coating would precipitate.). There is no real evidence to date of the overall distribution to support either model. Copper-enrichment arises from four sources; the matrix, hardening precipitates, dispersoid particles and large intermetallic particles. Since the large intermetallic particles are removed during deoxidation, copper-enrichment is likely to be due to dissolution of the matrix, hardening precipitates and the dispersoid particles which may provide some granularity to the enriched layer thickness. However, the enrichment due to the solid solution component should be continuous. Hence, it seems more likely that the copper influences the reaction kinetics. It should be noted that commercial formulations contain sequestering agents to complex copper in solution to minimize the amount of plating out of dissolved copper. To understand how the copper-enriched layer has an influence on the kinetics of reaction of underlying alloy composition it is instructive to examine the reactions of the CrCC solution with the intermetallic phases. Ferricyanide has been observed to deposit onto the copper-containing intermetallic phases and form an insoluble compound, as suggested by Hagans and Haas (1994). Juffs et al. (2002) proposed, however, that, upon adsorption of the ferricyanide onto the copper, it is reduced to ferrocyanide due to cathodic activity. Concurrently, the chromate that is also adsorbed onto the copper-containing sites, is reduced to Cr(III), not through the Fe(II)/Fe(III) couple but through cathodic reduction at the intermetallic phase. Hence, the activity of the copper containing sites is lost, because, with Cr(III) and ferrocyanide on the surface, there is no mechanism for

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oxidation of the ferro/ferri couple, thus, removing the accelerating role of the Fe(II)/Fe(III) couple. Under either of these mechanisms (copper ferri/ferro cyanides or cathodic reduction of Fe(III) and Cr(VI)), the loss of the ferricyanide at the reaction interface, would mean that the coating solution will behave more like an unaccelerated solution giving thinner coatings over intermetallic phases as seen in Figure 27.

Other Conversion Coatings As mentioned earlier, other non-chromate coating formulations are also either under development or available commercially [Buchheit and Hughes (2003), Nylund (2000)]. Those that are available commercially include the Ti/Zr processes [Knudsen et al. (2003), Deck and Reichgott (1992), Lunder et al. (2004), Tomlinson (1997)], Co-based process [Schreiver (1992,1996), Roland (1998), Hughes et al. (2004)], permanganate processes [Bibber (1993), Hughes et al. (2006)], silane treatments, self-assembled monolayers [Schmidt-Handsberg and Schubach (2003)] and rare earth processes. Most of these processes do not perform adequately on copper-containing alloys [Chalmers (1995)]. This is probably because many of these formulations are specifically designed to interact with a hydrated aluminium oxide surface, rather than a copper-rich surface, to form bridging Al-O-M bonds from which a coating “superstructure” can be built. One obvious exception to this is the rare earth processes which benefit from the presence of copper-enrichment [Hughes et al. (2004), Scholes et al. (2006), Campestrini et al. (2004)], but these are currently limited to application to architectural alloys [Schmidt-Handsberg and Schubach (2003)]. This coating system is intrinsically different from the other coating systems in that it relies on the cathodic activity of the surface for the deposition reaction. Unlike CrCCs the coating weight increases with copper content of the alloy. Formulations contain H2O2 which is cathodically reduced to generate a local pH rise in which the rare earth (principally cerium species) is deposited onto the surface. A second exception is the cobalt-based process [Schreiver (1992), Schreiver (1996), Roland (1998)] which has a poor performance on copper-containing alloys compared with alloys of reduced copper content [Chalmers (1995)].

CONCLUSIONS The importance of microstructure and chemistry of aluminium alloys has been demonstrated for copper-containing aluminium alloys. Alloying elements such as copper are present in solid solution, intermetallic particles or both. The distribution of alloying elements within aluminium alloys depends on alloy composition and processing history. Copper is one of the most noble alloying elements that is added to aluminium alloys. Because of this property, the copper distribution has a major impact on corrosion processes. Generally, copper-containing intermetallic particles will facilitate cathodic reactions resulting in the anodic dissolution of the adjacent aluminium matrix leading to pitting and other degradation reactions. In a special case, S-phase (Al2CuMg) intermetallic particles, while initially behaving as a net anode (i.e. undergoing dissolution), eventually become copper–rich and transform into very efficient cathodes. Alloys containing these intermetallic particles such as some of the AA2xxx series alloys, can expect to have severe corrosion issues related to the presence of copper. In metal finishing, proper pre-treatment, such as deoxidation, should remove the majority of intermetallic particles including the copper-containing particles. However, copper enrichment occurs just beneath the surface oxide wherever etching processes are used. While nitric-acid containing processes remove copperrich smut from the surface, they cannot remove the copper enriched layer which forms as a result of etching. This layer gives rise to some typical coating variation of the different alloy series such as variation in coating weight during chromate conversion coating. The investigation and understanding of copper distributions in aluminium alloys is part of a larger topic of microstructure and process control, fabrication, forming and joining, surface modification and recycling to provide light alloys

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that meet current requirements and future challenges. The future challenges are to develop cheaper, stronger, more formable and lighter alloys to compete with composites.

ACKNOWLEDGEMENTS The authors will like to thank Mr Tim Harvey, Dr Scott Furman, Dr R. Lumley and Dr R. Taylor for critical reading of the text and for feedback. Mr Tim Harvey is also acknowledged for assistance with the references. The authors would also like to thank past and present colleagues and students in the Light alloys Group in Corrosion and Protection Centre at Manchester.

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INDEX # 2D, 41, 42

A absorption, 56, 77 accelerator, 75, 77 acetic acid, 52 acetone, 52 acid, 32, 37, 49, 50, 52, 55, 59, 61, 68, 70, 72, 73, 74, 75, 80, 83 acidic, 3, 37, 39, 48, 51, 64, 81 acidity, 47, 52 adhesion, 73, 75 adsorption, 40, 53, 56, 81 aerospace, 1, 2, 69, 73, 74, 78 AFM, 34 Africa, 98 Ag, 12, 13 age, 18, 46 ageing, 6, 8, 10, 17, 21, 33 agents, 51, 81 aging, 2, 8, 18, 32 aid, 36 air, 26 aircraft, 1, 2, 3, 11, 13, 18 alkali, 49, 51

alkaline, 31, 39, 46, 47, 51, 53, 54, 55, 57, 60, 64, 74, 78 alkalinity, 42 alloys, vii, 1, 2, 3, 5, 6, 8, 9, 10, 11, 12, 13, 14, 15, 16, 17, 18, 19, 21, 23, 25, 26, 27, 28, 29, 30, 31, 32, 33, 34, 38, 39, 41, 44, 45, 49, 51, 53, 54, 55, 56, 57, 58, 59, 65, 67, 68, 69, 70, 73, 74, 78, 79, 80, 82, 83, 85 aluminium, vii, 1, 3, 5, 6, 8, 9, 10, 11, 12, 15, 17, 18, 19, 21, 22, 23, 25, 26, 27, 28, 29, 30, 31, 32, 33, 35, 36, 38, 39, 41, 42, 43, 45, 46, 48, 49, 51, 52, 54, 55, 56, 57, 58, 59, 60, 61, 63, 64, 65, 66, 67, 68, 69, 70, 71, 73, 74, 75, 77, 78, 79, 80, 81, 82, 83 aluminium alloys, vii, 1, 2, 3, 5, 8, 9, 11, 12, 15, 17, 18, 19, 21, 23, 26, 28, 29, 30, 31, 33, 38, 39, 41, 45, 54, 55, 56, 57, 59, 65, 67, 69, 70, 74, 78, 83 aluminum, 44, 49, 56, 89, 90, 92, 93 aluminum surface, 56 ammonium, 71 amorphous, 25, 26, 44 anions, 38, 73 anode, 41, 43, 48, 83 application, vii, 1, 49, 51, 78, 82 aqueous solution, 25, 28, 56 aqueous solutions, 25 argument, 45 artificial, 8, 32 ASTM, 27, 50, 87, 89 Atlas, 95

100

Index

atoms, vii, 31, 54, 56, 70 attacks, 52, 74 attention, 16 Australasia, 94, 96 Australia, 94, 96 automotive, 1, 6, 17 availability, 39

B backscattered, 12, 43, 63 band gap, 54 barrier, 39, 70, 71, 73 behaviours, 31 benefits, 73 binding, 55 binding energy, 55 blocks, 1, 77 Boeing, 2, 3 boiling, 56 bonding, 3, 56, 58, 65, 69 bonds, 82 borderline, 48 bounds, 5 brass, 39 breakdown, 25, 33, 34, 37, 38, 42, 80 by-products, 48

C carbon, 52, 79 carbonyl groups, 52 cast, 1, 8, 19 casting, 1 catalyst, 77 catalysts, 77 cathode, 38, 41, 48, 54, 57 cations, 78 cavities, 20 CCC, 79 cell, 47, 73 cerium, 41, 61, 82 Chalmers, 73, 82, 89 chemical, 31, 35, 49, 53, 57, 59, 69, 78

chemicals, 53, 55 chemistry, vii, 23, 38, 49, 50, 51, 53, 55, 57, 64, 65, 67, 75, 83, 88 chloride, 28, 37, 38, 39, 41, 43, 45 Chloride, 38, 52 chromium, 9, 10, 17, 18, 59, 60, 75, 77, 78, 79, 81, 89 cladding, 11, 16 classes, 8, 16, 17, 21, 75 cleaning, 51, 52, 53, 54, 57, 60, 64, 73, 74, 78 clustering, 9, 14, 15, 18, 56 clusters, vii, 12, 31, 42, 46, 47, 70 Co, 73, 82 coatings, 51, 53, 54, 58, 59, 61, 66, 73, 75, 76, 82 cobalt, 82 commercial, 1, 16, 28, 45, 51, 56, 60, 74, 78, 80, 81 commodities, 95 communication, 91 community, 49, 50 complexity, 5 components, 8, 9, 51, 53, 57, 59, 64, 65, 75 composites, 84 composition, 6, 14, 16, 21, 22, 23, 38, 49, 53, 55, 57, 61, 62, 81, 83 compositions, vii, 5, 14, 59 compounds, 2, 30, 55, 70, 79 concentration, 5, 19, 26, 28, 37, 39, 43, 47, 59, 65, 70, 71 conductivity, 28 conductor, 54 Congress, iv, 98 consensus, 41 construction, 1 continuing, 20 control, 6, 10, 18, 38, 40, 50, 83 controlled, 39, 69 convection, 46 conversion, 3, 37, 49, 51, 53, 58, 59, 60, 61, 73, 74, 75, 77, 78, 79, 80, 81, 83 cooling, 8, 39 copper, vii, 2, 3, 5, 6, 9, 10, 11, 12, 13, 16, 17, 18, 22, 23, 25, 26, 27, 28, 29, 30, 31, 32, 33, 34, 38, 39, 40, 41, 42, 43, 44, 45, 46,

101

Index 47, 48, 49, 51, 52, 53, 54, 55, 57, 58, 59, 60, 61, 63, 64, 65, 66, 67, 68, 69, 70, 71, 72, 73, 74, 75, 76, 78, 79, 80, 81, 82, 83 copper oxide, 23, 28 correlation, 15 correlation function, 15 corrosion, vii, 2, 3, 8, 9, 11, 13, 14, 16, 17, 18, 19, 25, 27, 28, 29, 30, 31, 32, 33, 34, 35, 38, 39, 40, 41, 42, 43, 44, 45, 46, 47, 48, 50, 51, 52, 53, 57, 73, 74, 79, 80, 83 corrosive, 25, 31, 41, 47 coverage, 39, 67, 81 covering, 21, 32, 46, 58, 79 cracking, 18, 20 CRC, 93 cross-sectional, 15 cyanide, 77

D decomposition, 18 defects, 28, 54 deformation, 21 degradation, 69, 83 degree, 2, 10, 18 density, 14, 21, 36 deposition, 39, 48, 58, 61, 63, 64, 74, 75, 77, 78, 79, 80, 81, 82 deposits, 44, 60 diffusion, 19, 21, 31, 47, 71 dispersion, 16 distilled water, 26 distribution, vii, 2, 11, 13, 14, 15, 16, 18, 32, 40, 49, 69, 72, 80, 81, 83 distribution function, 18 donor, 54 Doppler, 54

E earth, 65, 73, 82 electric field, 69, 70 electrical, 1, 42, 56 electrical resistance, 56

electrochemical, 3, 25, 26, 27, 28, 29, 31, 32, 44, 69, 70 electrochemical reaction, 3 electrochemistry, vii electrolyte, 26, 28, 29, 36, 43, 69, 71, 73 electrolytes, 29, 70, 73 electron, 12, 13, 26, 34, 40, 43, 44, 54, 58, 60, 62, 63, 66, 68, 72, 74, 80 electron microscopy, 12, 68 electronic, iv, 3, 28 electronics, 1 electrons, 54 electrostatic, iv energy, 19, 26, 28, 31, 39, 45, 55 engineering, 1 engines, 2 environment, 3, 21, 23, 25, 38, 39, 52, 58 environmental, 53 equilibrium, 5, 65, 73 etching, 31, 35, 51, 53, 54, 55, 56, 57, 60, 61, 62, 63, 64, 67, 80, 83 ethane, 52 ethanol, 43, 52, 70 Ethanol, 52 Europe, 19, 57 evaporation, 19 evidence, 32, 33, 42, 45, 46, 58, 78, 81 expert, iv exposure, 21, 39, 56 external environment, 3 extrusion, 17, 19

F fabrication, 83 failure, 74 family, 34, 73, 74 fatigue, 11 faults, 19 feedback, 85 Fermi, 26 Fermi level, 26 filiform, 16, 18, 19, 30 film, 25, 32, 35, 56, 57, 59, 69, 70, 71, 72, 73, 74, 79

102

Index

film formation, 70, 72, 73, 79 film thickness, 32, 57, 71 films, 28, 57, 70, 71, 72 fines, 20 flow, 46, 73 fluorescence, 50 fluoride, 51, 60, 65, 69, 74, 80 fluoride ions, 80 fluorides, 53 fluorine, 59, 75 foils, 1 folding, 21 food, 6 fracture, 2, 15, 21 fragmentation, 48 France, 1 free energy, 19, 26, 27, 32, 45

H2, 3, 61, 63, 64, 65, 78, 80 handling, 6 hardness, 15 health, 49, 53 heat, 1, 2, 5, 8, 9, 16, 17, 19, 21, 27, 32 heavy metal, 39 heavy metals, 39 heterogeneous, 28, 34, 37 high pressure, 72 high temperature, 5, 9, 19 host, 30 House, 96 household, 1 hydrodynamic, 46 hydrogen, 35, 38 hydrogen gas, 35, 38 hydroxide, 35, 36, 54, 77 hydroxides, 25, 38

G I gallium, 56 gas, 35, 38, 71, 72, 73 gauge, 14 gel, 77, 78, 88 gels, 36, 48 generation, 38, 45, 53, 71, 72, 73, 77 Germany, 87 Gibbs, 26, 27 Gibbs free energy, 26, 27 gold, 27 grain, 6, 10, 13, 14, 18, 19, 21, 32, 33, 40, 71, 80 grain boundaries, 6, 13, 21, 32, 33, 40 grain refinement, 14 grains, 12, 18 greek, 10 groups, 52 growth, 10, 32, 33, 36, 38, 39, 40, 45, 56, 57, 60, 70, 73, 74 growth rate, 70

H H19, 58

images, 35, 41, 43, 44, 63, 66 imaging, 34 immersion, 49, 55, 57, 60, 61, 64, 74, 75, 76, 81 impurities, 11, 16, 59 inactive, 81 indication, 33, 58 industrial, 53 industry, 1, 2, 11, 49, 50, 51, 52, 58, 69, 73 inhibitor, 53 inhibitors, 52 initiation, 14, 16, 33, 34, 38, 39, 40, 44, 57 initiation rates, 34 injury, iv intensity, 75, 79 interface, 9, 28, 31, 32, 53, 59, 61, 66, 69, 71, 72, 73, 74, 78, 81, 82 intermetallic compounds, 30, 70 intermetallics, 2, 6, 14, 18, 21, 29, 30, 32, 34, 35, 36, 37, 38, 39, 44, 47, 56, 67, 74, 78, 79, 80 ions, 28, 32, 35, 38, 39, 44, 45, 47, 51, 56, 59, 69, 72, 80

103

Index iron, 9, 11, 16, 17, 22, 23, 28, 30, 54, 55, 56, 60, 61, 63, 77, 78 isolation, 42, 47 Italy, 88

J Jung, 31, 32, 92

K kerosene, 52 kinetics, 37, 40, 58, 80, 81 King, 52, 53, 55, 92

L laser, 79 laser ablation, 79 lattice, 10 leaching, 45 lead, 28, 41, 44, 45, 47, 55, 71 legislation, 53 lifetime, 2, 32, 54 linear, 71 literature, 58, 60, 78 lithium, 19 Lithium, 1, 2 lithography, 1 localised, 28, 74 London, 94, 96 lubricants, 20

M magnesium, 5, 6, 9, 11, 16, 17, 18, 19, 21, 22, 23, 26, 27, 28, 41, 43, 45, 46, 53, 54, 55, 61, 64 magnetic, iv magnetron, 74 maintenance, 50 manganese, 9, 10, 16, 18, 21, 23, 30, 54, 55, 61, 63

mass loss, 19 matrix, 3, 6, 9, 10, 13, 21, 23, 29, 30, 31, 32, 33, 34, 36, 39, 41, 42, 44, 46, 47, 48, 50, 54, 55, 56, 60, 61, 64, 65, 67, 69, 72, 75, 76, 78, 79, 80, 81, 83 mechanical, iv, vii, 1, 3, 8, 9, 11, 17, 18, 19, 20, 21, 23, 42, 45, 47, 49 mechanical properties, vii, 3, 8, 9, 11, 17, 18 mechanical treatments, 49 media, 3 melting, 5 metal ions, 51 metallography, 50 metals, 26, 27, 39, 55 methanol, 52 metric, 58 microelectrode, 29 microscopy, 12, 33, 37, 77 microstructure, 5, 9, 10, 11, 12, 16, 18, 19, 21, 22, 28, 37, 79, 83 military, 2 mobility, 28, 72 models, 59, 67, 69, 77 monolayers, 73, 82 Moon, 35, 36, 94 morphology, 41, 50, 53, 58 movement, 45

N NaCl, 27, 29, 34, 35, 39, 42, 43, 45 nanometer, 45 nanometer scale, 45 nanometers, 10 nanoparticles, 32, 42, 44, 54, 79 natural, 8, 73 network, 46, 47, 58, 59, 65, 69 New York, iii, iv Ni, 6 nitric acid, 50, 55, 70, 74, 75 nitrogen, 79 NMR, 12 nodules, 39, 58 normal, 34 North America, 57

104

Index O

observations, 32, 41, 64 occlusion, 56 occupational, 53 occupational health, 53 Ohio, 88 oil, 6 oils, 51, 53 organic, 49, 53 organic chemicals, 53 orientation, 32, 71 orthorhombic, 47 oxidants, 59, 61, 63, 65 oxidation, 19, 36, 44, 45, 52, 60, 64, 69, 71, 73, 77, 82 oxide, 19, 20, 21, 22, 23, 25, 27, 28, 31, 32, 33, 34, 35, 38, 41, 44, 46, 50, 52, 53, 54, 55, 56, 57, 58, 59, 60, 61, 64, 65, 66, 67, 68, 69, 70, 73, 74, 75, 77, 79, 80, 81, 82, 83 oxides, 20, 21, 25, 27, 28, 38, 46, 53, 54, 55, 57, 58, 59, 60, 61, 64, 65 oxygen, 22, 28, 35, 37, 38, 58, 60, 71 oxyhydroxides, 25

P packaging, 1, 6 paper, 68 particle density, 14 particles, 6, 8, 9, 10, 11, 13, 14, 15, 16, 18, 19, 21, 23, 30, 32, 33, 34, 36, 39, 40, 41, 44, 45, 46, 47, 48, 49, 54, 55, 56, 60, 61, 63, 64, 65, 66, 67, 75, 76, 77, 78, 79, 81, 83 passivation, 46 passive, 37 percolation, 47 percolation theory, 47 performance, vii, 2, 8, 13, 23, 30, 32, 48, 50, 51, 57, 82 pH, 25, 26, 35, 36, 38, 39, 43, 45, 46, 55, 75, 78, 81, 82 pH values, 35, 45 phase diagram, 5, 6, 7, 11

phosphate, 55, 56, 73 phosphates, 53 physics, 88 plastic, 73 plasticity, 73 play, 40, 50, 51, 52 polarization, 42, 46, 80 polarized, 36, 45 poor, 28, 75, 82 poor performance, 82 population, 14, 50 pore, 73 porosity, 20, 72, 73 porous, 25, 70, 72 positron, 54 positrons, 54 powders, 19 precipitation, 2, 6, 9, 10, 11, 13, 17, 18, 19, 21, 28, 32, 48 prediction, 66 preparation, iv, 49, 52, 78 pressure, 49, 71 private, 91 probability, 38 probe, 43 process control, 83 production, 17, 38 promote, 40, 80 propagation, 34 property, iv, 83 protection, 3, 11, 51, 59 protons, 38 pseudo, 56

Q quality control, 50

R radial distribution, 18 radius, 18, 32, 38, 45 Raman, 79 Raman spectroscopy, 79

Index random, 15 range, vii, 1, 3, 5, 6, 10, 14, 17, 21, 29, 31, 43, 49, 52, 54, 57, 58, 59, 61, 73, 74, 76 rare earth, 59, 73, 82 rare earths, 59 reaction rate, 30, 34, 36, 43 reactivity, 25 reading, 85 reagents, 52 recession, 72 recycling, 83 redistribution, vii, 42, 45, 46 redox, 65, 77 reduction, 3, 25, 26, 28, 35, 36, 37, 38, 39, 40, 47, 77, 81 reflection, 26 relationship, 2, 15, 60, 64, 78 relaxation, 47 relaxation process, 47 relaxation processes, 47 relevance, 9, 52 repair, 72, 73 research, 13, 16, 31, 38, 49, 50, 52, 58 resistance, 2, 3, 11, 17, 25, 30, 38, 51, 53, 56, 57, 73, 74, 80 resolution, 14 retention, 61 risk, 49 rods, 12 rolling, 13, 14, 19, 20, 21, 64 room temperature, 56, 57, 64

S S phase, 11 salt, 43, 55, 74 salts, 38 sample, 15, 45, 46, 52 sampling, 15 Scanning electron, 66 scanning electron microscopy (SEM), 34, 68 science, 50 search, 95 segregation, 19 self limiting, 75

105

separation, 48 series, vii, 2, 3, 5, 10, 11, 13, 14, 16, 17, 18, 21, 25, 28, 30, 33, 41, 49, 65, 70, 74, 80, 83 services, iv shape, 15 shipping, 1 silane, 73, 82 silica, 49 silicate, 55 silicates, 53 silicon, 9, 11, 14, 16, 17, 19, 28, 55, 61, 64 sites, 14, 30, 36, 38, 39, 46, 47, 48, 59, 77, 79, 81 sol-gel, 88 solid phase, 5 solid state, 10 solidification, 9, 10 solubility, 25, 51, 55 solutions, 39, 45, 50, 51, 53, 54, 57, 59, 75, 76 solvent, 51, 52 solvents, 52 South Africa, 98 Soviet Union, 2 Spain, 96 spatial, 14, 15 speciation, 52 species, 53, 56, 57, 61, 71, 75, 77, 78, 79, 81, 82 spectra, 12 spectroscopy, 79 S-phase, 12, 14, 16, 21, 29, 30, 34, 35, 41, 42, 43, 44, 46, 47, 48, 49, 60, 83 sponges, 44, 45 stability, 25, 28, 46 stabilize, 75 stabilizers, 52 stages, 10, 38, 47, 57, 64, 65 standard operating procedures, 50 standards, 51, 87 statistics, 14 steady state, 42 stock, 9 storage, 21, 23, 69 strain, 8, 9 strength, vii, 1, 2, 6, 8, 9, 11, 17, 69, 73

106

Index

stress, 2, 17, 18, 73 stretching, 8 students, 85 substrates, 78 sulphate, 43 summaries, 11 Sun, 56, 69, 76, 78, 80, 96, 97 supply, 48 surface area, 13, 41, 43, 45, 67, 81 surface chemistry, 49, 50, 58 surface diffusion, 31, 47 surface energy, 45 surface layer, 19, 21, 22, 25 surface modification, 83 surface region, 21 surface structure, 21, 28 surface treatment, 58 surfactants, 53 susceptibility, vii, 2, 16, 18, 19, 33, 55 switching, 37 symbols, 10 systematic, 31 systems, 11, 26, 39, 41, 53, 67, 82

T temperature, 5, 9, 10, 17, 55, 56, 57, 59, 63, 64, 65, 70, 77 tensile, 18 tensile strength, 18 theory, 19, 47 thermal, 10 thermodynamic, 19 Ti, 73, 82 time, 1, 10, 32, 48, 55, 56, 65, 70, 74, 75 time periods, 48 titanium, 10 tolerance, 11 toughness, 2 transfer, 39 transformation, 20 transition, 53, 55 transition metal, 53, 55

transmission, 12, 72 transmission electron microscopy (TEM), 12, 50 transport, 1, 18, 42, 51 treatable, 8, 9 trend, 53 tribological, 18 tribology, 19 trichloroethylene, 52 tunneling, 74, 80

V vacancies, 35 vacuum, 26 values, 32, 33, 35, 42, 45, 65, 75 variables, 67 variation, 14, 22, 51, 75, 83 visual, 43, 44, 50 voids, 71

W warrants, 14 water, 8, 26, 28, 35, 36, 38, 52, 56, 57, 60 wear, 20 weight ratio, 1 welding, 17 workers, 33, 47, 64

X XPS, 50, 52, 54, 55, 58, 60, 79

Z zinc, 5, 9, 17, 18, 22, 23, 26, 27, 53, 54, 55, 64 zinc oxide, 64 zirconium, 10, 18 Zn, 5, 6, 7, 10, 17, 18, 21