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 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 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

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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

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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

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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.

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(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

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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