EFFECT OF CU CONTENT ON CORROSION BEHAVIOR AND CHROMATE CONVERSION COATING PROTECTION OF 7XXX SERIES AL ALLOYS

Transcription

1 EFFECT OF CU CONTENT ON CORROSION BEHAVIOR AND CHROMATE CONVERSION COATING PROTECTION OF 7XXX SERIES AL ALLOYS DISSERTATION Presented in Partial Fulfillment of the Requirements for the Degree Doctor of Philosophy in the Graduate School of The Ohio State University By Qingjiang Meng, M.S. * * * * * The Ohio State University 2003 Dissertation Committee: Dr. Gerald S. Frankel, Adviser Dr. Rudolph G. Buchheit Dr. Michael J. Mills Dr. Richard L. McCreery Approved by Adviser Dept. of Materials Science & Engineering

2 ABSTRACT The addition of Cu in Al-Zn-Mg alloys increases the mechanical strength and resistance to stress corrosion cracking of 7xxx series aluminum alloys (AA7xxx). The peak aged T6 temper provides the maximum mechanical strength by precipitation hardening. However, the presence of noble Cu makes AA7xxx-T6 more susceptible to localized corrosion, such as pitting, crevice and intergranular corrosion (IGC). In order to protect AA7xxx-T6 from localized corrosion, protective chromate conversion coatings (CCCs) must be used. Cu has been reported to affect the CCC protection performance. The exact roles of Cu content in corrosion behavior and CCC protection of AA7xxx-T6 are the focus of this study. Polarization and Electrochemical Impedance Spectroscopy (EIS) approaches were used in combination with materials characterization techniques, such as Focused Ion Beam (FIB), SEM, TEM, High Resolution TEM (HRTEM), Scanning TEM (STEM), and X-ray Photoelectron Spectrometry (XPS). Electrochemical tests on AA7xxx-T6 with various Cu content in deaerated chloride solution found that all alloys except for essentially Cu-free AA7004-T6 had two breakdown potentials, which increased logarithmically with increasing Cu content. Transient dissolution of the fine hardening precipitates and the surrounding solid solution in a thin surface layer was found in the Cu-containing alloys polarized at potentials ii

3 between the two breakdown potentials. Stable dissolution associated with combined IGC and selective grain attack was found above the second breakdown potential. EIS tests revealed that the overall influence of Cu on the corrosion behavior was detrimental due to Cu enrichment in aerated chloride solution. TEM and STEM analysis revealed that CCC was heterogeneous on the heterogeneous microstructure of AA7075-T6. The coatings formed on coarse intermetallic particles were much thinner than CCC formed on the matrix. It was found that the CCC formed on the matrix mainly consisted of a Cr III OOH backbone and chemisorbed HCr VI O - 4. A sol-gel model for CCC formation was supported by the observations in this study. Finally the Cu content can have different effects on CCC protection: Cu is beneficial to CCC protection for coatings formed on polished AA7xxx- T6, but Cu is detrimental if it is enriched on the surface prior to CCC formation. iii

4 TO MY WIFE LIU-DAN AND MY SON MICHAEL FANDI MENG iv

5 ACKNOWLEDGMENTS I would like to acknowledge my advisor, Dr. Jerry Frankel, for his advisory and mentorship during the course of my graduate study at Fontana Corrosion Center. Dr. Frankel has not only helped me to grow as a researcher in the fields of electrochemistry and corrosion, but also helped me in developing my professional skills and personality for my future career. It would have been impossible to finish my thesis without his incredible inspiration, encouragement, and advice. I wish to thank Dr. Rudy Buchheit for his helpful discussion and support during the course of my study and research. I am also grateful for a lot of good comments and suggestions given by Dr. Michael Mills and Dr. Richard McCreery, both on my advisory committee. I would like to acknowledge the Strategic Environmental Research and Development Program for financially supporting this project. During the past few years I have received a lot of kind help and support from all the former and current FCC group members. They are Dr. Eiji Akiyama, Dr. Jian Zhang, Dr. Donghui Lu, Dr. Thodla Ramgopal, Dr. Weilong Zhang, Dr. Wenping Zhang, Dr. Patrick Leblanc, Dr. Valerie Laget, Mr. Xiaodong Liu, and others. I would also like to thank Mr. Henk Colijn, Mr. Cameron Begg, Dr. Lisa Hommel, and Dr. Steve Goss, who taught me how to use the state-of-the-art electron microscopes, XPS, and SIMS. My special thanks go to Ms. Suqin Meng, who helped me to settle down in Columbus at the v

6 beginning of my study at Ohio State and made TEM samples for me. I would like to thank Ms. Cindy Flore and Ms. Dena Bruedigam for their office supply. I would like to acknowledge Mr. Gary Dodge, Mr. Steve Bright, Mr. Lloyd Barnhart, and Mr. Ken Kushner for helping and training me to use a variety of tools. I also wish to thank Dr. Suliman Dregia, Mr. Mark Cooper, Ms. Mei Wang, and Ms. Wendy Ranney for their kind help with my academic problems. My most sincere thanks go to my families. My parents live in a city thousands of miles away, but I can feel their support every day. Their love and encouragement has accompanied me from my birth to now and will forever be treasured. Finally, I want to thank my wife, Liu-Dan, and my lovely son, Michael, who had been with me through numerous cherishing days and nights. Without their support and encouragement, it is hard to imagine that I could have finished my studies at Ohio State. vi

7 VITA October 29, Born Jiamusi, P.R. China B. S. Materials Science and Engineering, Beijing University of Aeronautics & Astronautics, Beijing, P.R. China M.S. Materials Science and Engineering, Beijing University of Aeronautics & Astronautics, Beijing, P.R. China present...graduate Research Associate The Ohio State University M.S. Materials Science and Engineering, The Ohio State University PUBLICATIONS 1. Q. Meng, G. S. Frankel, H. O. Colijn, S. H. Goss, Characterization of the Region around MnS Inclusions in Stainless Steels, Nature, accepted for publication. 2. Q. Meng, T. Ramgopal, G.S. Frankel, The Influence of Inhibitor Ions on Dissolution Kinetics of Al and Mg Using the Artificial Crevice Technique, Electrochemical and Solid-State Letters, 5 (2), B1 (2002). 3. H. Xu, Q. Meng, C. Jiang, S. Gong, Aging Effect on the Structure and Shape Memory Property of Ti 51 Ni 13 Pd 36 High Temperature Shape Memory Alloy, Proceeding of the 3rd Pacific Rim International Conference on Advanced Materials and Processing (PRICM 3), ASM, 2039, July FIELDS OF STUDY Major Field: Materials Science and Engineering vii

8 TABLE OF CONTENTS Abstract... ii Dedication...iv Acknowledgments... v Vita... vii List of Tables... x List of Figures... xi Chapters: 1. Introduction Literature Review Metallurgy of 7xxx Series Al Alloys Localized Corrosion of Al Alloys Pitting Corrosion of Al Alloys Criteria for Evaluation of Pitting Corrosion of Al Alloys Intergranular Corrosion of Al Alloys Role of Alloying Addition and Intermetallics in Localized Corrosion of Al Alloys Role of Chromate in Corrosion Inhibition of Al Alloys Chromate in Solution Chromate Conversion Coatings Formation of CCC Composition and Structure of CCC Coatings Corrosion Protection Mechanism Chromate in Paint Cu Enrichment and Redistribution Research Objective Page viii

9 3. Effect of Cu Content on Corrosion Behavior of 7xxx Series Al Alloys Introduction Experimental Results Microstructure Polarization Curves and Types of Corrosion EIS Measurement under Free Corrosion Conditions Discussion Mechanism of Localized Corrosion of AA7xxx-T Role of Cu Content in Corrosion Resistance of AA7xxx-T Summary Characterization of Chromate Conversion Coating on AA7075-T Introduction Experimental Materials SEM and EDS TEM and STEM Electrochemistry Results Discussion Determination of CCC Composition CCC Coating Formation on Matrix Coating Formation on Intermetallics Summary Effect of Cu Content on the Protection of AA7xxx-T6 by Chromate Conversion Coatings Introduction Experimental Results CCC Breakdown EIS Measurement on CCCs Cu Enrichment by Acid Pretreatment Discussion Role of Cu Content in CCC Protection of AA7xxx-T Relevance to Salt Spray Testing Summary Conclusions and Future Work Conclusions Future Work Bibliography ix

10 LIST OF TABLES Table Page 3.1 Composition of 7xxx Al alloys measured by ICP-MS in wt% Coarse intermetallic particles identified by SEM/EDS Breakdown potentials for AA7xxx-T6 in deaerated 0.5 M NaCl, ph= Data obtained from XPS spectra. BE is binding energy, and FWHM is full width at half maximum height Ratios of O/Cr concentration for possible mixed Cr oxide, 77% Cr(III) and 23% Cr(VI) x

11 LIST OF FIGURES Figure Page 2.1 Variations in cross-sectional shape of pits Schematic diagrams for pit initiation models. (a) adsorption mechanism; (b) anion penetration and ion migration model; (c) film breakdown theory Metastable pit transients observed on 302 stainless steel polarized at 420 mv SCE in 0.1M NaCl solution Anodic and net current densities change as a function of potential for 100 nm Al film in 0.1 M NaCl solution Concentration of Al 3+, Al(OH) 2+, and H + as a function of the product of the depth x and the current density i in a unidirectional pit Schematic cyclic polarization showing E p and E R Effect of aging time on the corrosion behavior, at constant potential, of Al-3.33 Cu in de-aerated 1 M NaCl solution at 25 C. Aging temperature 240 C.ο: passive; : IGC; : pitting + IGC; : pitting Variations of pitting potential as a function of alloying element content of (a) Al-Cu, (b) Al-Zn, and (c) Al-Mg binary alloys Pitting potentials for freshly deposited samples, E p and aged samples E p a, along with repassivation potentials, E R, for pure Al and AlNb alloys Schematic of standard coating system applied to Al alloys Polarization curves of AA2024-T3 in an aerated, oxygen-stirred 1 M NaCl base solution containing different amounts of persulfate and dichromate ions xi

12 2.12 Time series showing the effect of CrO 4-2 inhibitor on the anodic current spikes associated with metastable pitting polarized at -0.5 V SCE Schematic diagram illustrating duplex mechanism based model for the interaction of chromate with aluminum surface Proposed mediation mechanism by ferricyanide. Arrows represent redox cross reactions Schematic illustrating CCC formation by a sol-gel mechanism Mud-crack morphology of CCC formed on AA2024-T Auger depth profiling showing that CCCs are mainly composed of Cr and O Schematic representing the structure of CCCs formed on AA Cr K edge XANES reference spectra from solutions of 0.1 M CrO 4 2- (top) and Cr(H 2 O) 6 3+ (bottom), and mixed Cr(VI)/Cr(III) Schematic illustration of the bipolar model for interfaces for Al. MeO 4 2- is a metal oxyanion such as MeO 4 2- and CrO Schematic illustration of self-healing process of CCCs Cr VI concentration determined from the absorbance at 339 nm as a function of time after immersion of a CCC in nanopure water. Curvers are labeled with the ratio of the geometric CCC area to solution volume Model for adsorption of Cr VI to solid Cr III hydroxide showing the Cr III -O-Cr VI covalent bonding (a) The optical picture and (b) schematic representing Cu enrichment and redistribution during localized corrosion in AA2024-T Static SIMS chemical maps showing relative Cr and Cu concentration of a CCC on AA2024-T3 alloy at the outermost surface and at a level approximately 20% through the CCC Corrosion potential as a function of Cu content for Al-Cu solid solution and intermetallic compounds. Solid solution phase E corr values were determined in 1 M NaCl + 3% H 2 O 2. For the intermetallic E corr values, open data denotes measurements made in aerated 0.5 M NaCl; closed data denotes measurements made in 1 M NaCl + 3% H 2 O xii

13 3.1 Microstructure of three orthogonal sections of (a) AA7004-T6, (b) AA7039-T6, (c) AA7029-T6, (d) AA7075-T6, and AA7050-T6 sheets Microhardness versus Cu content curve showing that Cu addition increases the hardness of AA7xxx-T TEM micrographs showing the grain boundary regions in (a) 7004, (b) 7039, (c) 7029, (d) 7075, and (e) 7050 alloys Composition of η phase precipitates on grain boundary as a function of alloy Cu content in AA7xxx-T Potentiodynamic polarization curves for AA7xxx-T6 in deaerated 0.5 M NaCl, ph=3.56 with a scan rate of 0.2 mv/s Correlation between the breakdown potentials in deaerated 0.5 M NaCl, ph=3.56 and alloy Cu content Potentiostatic polarization curves at various applied potentials for AA7075-T6 in deaerated 0.5 M NaCl, ph= Variation of charge density as a function of applied potentials in 1 h potentiostatic polarization on AA7075-T6 in deaerated 0.5 M NaCl, ph=3.56. The arrows show the value of E 2 for each alloy Optical micrographs of (a) 7004, (b) 7039, (c) 7029, (d) 7075, and (e) 7050 samples potentiodynamically polarized to the current density of 1 ma/cm 2 in deaerated 0.5 M NaCl, ph= Cyclic anodic polarization curves for the same AA7075-T6 sample in deaerated 0.5 M NaCl, ph=3.56 at a scan rate of 0.2 mv/s. The sample was polarized three times: the apex current densities for three scans were 3, 10, and 10 ma/cm 2, respectively. When each cyclic polarization was finished, the sample was at open circuit until the OCP was stable Current transient of AA7075 polarized at 725 mv SCE for 500 s in deaerated 0.5 M NaCl, ph= SEM (a) secondary electron and (b) backscattered electron images of the surface of AA7075 sample polarized at 725 mv SCE for 500 s in deaerated 0.5 M NaCl, ph= XPS spectra measured from samples of AA7075 after mechanical polishing and a subsequent polarization at 725 mv SCE for 500 s in deaerated 0.5 M NaCl, ph=3.56. (a) Al 2p, (b) Zn 2p, (c) Mg 2p, xiii

14 (d) Cu 2p, (e) O 1s, and Cl 2p TEM micrograph of the product layer formed on AA7075-T6 when polarized at 725 mv SCE for 500 s in deaerated 0.5 M NaCl, ph= SEM micrographs of the surface of (a) 7039 polarized at 870 mv SCE, and (b) 7050 polarized at 720 mv SCE in deaerated 0.5 M NaCl, ph=3.56 for 24 h Metallographical micrograph of (a) top surface, (b) as polished cross-section, and etched cross-section of AA7075-T6 polarized in deaerated 0.5 M NaCl, ph=3.56 at 680 mv SCE for 5h SEM micrograph of FIB cross-section of the same AA7075 sample As in Figure 3.6 polarized at 680 mv SCE in deaerated 0.5 M NaCl, PH=3.56 for 5 h evidencing microstructural pitting Time evolution of the open circuit potentials of AA7xxx-T6 within 168 h immersion in aerated 0.5 M NaCl Polarization resistance determined by EIS tests as a function of immersion time for AA7xxx-T6 in aerated 0.5 M NaCl Open circuit potential versus time for 7xxx-T6 in deaerated 0.5 M NaCl, ph= Comparison of the breakdown potentials and the OCPs in deaerated and aerated chloride solutions. E 1 and E 2 are denoted by circles and squares, respectively. The OCPs in deaerated solution, ph=3.56 after about 30 min prior to the polarization tests are denoted by triangles. The range of OCP in deaerated solution during this period is given by the solid double-ended arrows. The OCPs in aerated solution after 168 h are denoted by crosses and are connected by a line. The range of OCP in aerated solution during this period is given by the dashed double-ended arrows. The variations in potential for AA7075 and AA7050 in aerated solution were so small that they were not included (a) Time evolution of the potential applied to AA7004-T6 and (b) potentiodynamic polarization curve for AA7004-T6 after 2 min holding at 1.15 V SCE showing only one breakdown potential for AA7004-T6. The solution was deaerated 0.5 M NaCl, ph=3.56 and the scan rate was 0.2 mv/s Potentiostatic polarization curve for AA7004-T6 at 930 mv SCE for 1 h in deaerated 0.5 M NaCl, ph=3.56 under two different xiv

15 conditions before polarization: (a) under OCP for 30 min and (b) under potential control at 1150 mv SCE for 2 min SEM images of CCC coatings formed on (a) Al 2 CuMg, (b) Al 3 (Fe Cu), and (c) Mg 2 Si intermetallic particles in AA7075-T TEM micrograph of CCC on matrix of AA7075-T High resolution TEM micrograph of CCC on matrix of AA7075-T Nano-EDS line profile of 3 min CCC on matrix of AA Ratio of O/Cr and Zn/Al signals from raw EDS data for the linescan shown in Figure 4.4 and transmitted electron intensity or brightness in STEM HAADF image inset on Figure Cr 2 O 3 powder standard for determination of K OCr factor. (a) STEM image of Cr 2 O 3 powders, and (b) intensity O/Cr ratio from raw EDS data for the EDS linescan. The linescan was made along the red line on the fine powder in the STEM image TEM micrograph of coating formed on Al 2 CuMg particle in AA7075-T TEM micrograph of coating on Mg 2 Si particle in AA7075-T TEM micrograph of coating on Al 3 (Fe Cu) particle in AA7075-T Nano-EDS line profile of coating on Al 2 CuMg particle in AA7075. No Fe signal was detected in linescans Nano-EDS line profile of coating on Mg 2 Si particle in AA7075. No Fe signal was detected in linescans Nano-EDS line profile of coating on Al 3 (Fe Cu) particle in AA7075 at open circuit (~ -610 mv SCE) Backscattered electron image of Al 2 CuMg ingot analog containing A: Al-33.6Mg-28.6Cu, B: Al-33.7Mg-15.4Cu, and C: Al-36.2Mg-4.5Cu by atomic percentage Backscattered electron image of Mg 2 Si ingot analog consisting of Mg 2 Si and Si Open Circuit Potential (OCP) for AA7075-T6 in Alodine solution xv

16 4.16 Current transients for Al 2 CuMg and Mg 2 Si ingot analogs in Alodine solution at 600 mv SCE, which is the OCP of AA7075-T6 in Alodine solution CCC coating thickness following a logarithmic-linear growth kinetics Comparison of breakdown potentials for CCC and bare AA7xxx-T Anodic polarization curves for bare and CCC AA7075-T6 (1 µm finish) in 0.5 M NaCl at a scan rate of 0.2 mv/s. The solutions for the bare and CCC samples were deaerated by Ar gas and open to air, respectively Optical macrograph of CCC coated 7075 (1 µm finish) polarized to 635 mv SCE with the ending current density of A/cm 2 in aerated 0.5 M NaCl SEM images of CCC coated 7075 (1 µm finish) polarized to 635 mv SCE with the ending current density of A/cm 2 in aerated 0.5 M NaCl. (a) secondary electron image, (b) backscattered electron image Variation of OCP for CCC coated AA7xxx as a function of immersion time during immersion in aerated 0.5 M NaCl The equivalent circuit model used to fit EIS data for CCC. R s is solution resistance, R c is CCC coating resistance, and C c is CCC coating resistance Coating resistance of CCC on polished and acid pretreated AA7xxx-T6. Polarization resistance of bare polished AA7xxx-T6 was measured as control experiments. All EIS measurements were conducted at OCP after 48 h immersion in aerated 0.5 M NaCl Optical micrograph of acid pretreated surface of AA7075-T XPS spectra measured from as polished and acid pretreated AA7075-T6 samples Optical micrograph of CCC on (a) as-polished and (b) acid pretreated AA7075 samples Breakdown potential and OCP for polished AA7xxx-T6 coated with CCC in aerated 0.5 M NaCl during168 h immersion. The error bars for the OCPs represent the OCP variation during 168 h immersion xvi

17 CHAPTER 1 INTRODUCTION High strength 7xxx series Al-Zn-Mg-Cu alloys (AA7xxx) such as AA7075 and AA7178 have been widely used in aircraft structure applications because of their high strength to weight ratio. The addition of Cu as one of alloying elements in Al-Zn-Mg alloys greatly improves the mechanical strength of 7xxx series alloys by precipitation hardening. The peak aged T6 temper provides the maximum mechanical strength. However, Cu-containing AA7xxx in the T6 temper (AA7xxx-T6) is extremely susceptible to various forms of localized corrosion in chloride environments, such as pitting, crevice corrosion, intergranular corrosion (IGC), exfoliation corrosion, and stress corrosion cracking (SCC). Since sudden and unpredicted failures due to localized corrosion can lead to catastrophic disaster, the United States Air Force spends over one billion dollars every year for fighting corrosion of aging aircraft. In order to protect high strength Al alloys from localized corrosion and other forms of corrosion, protective chromate conversion coatings (CCCs) must be used. Chromates in solution are also extremely effective in inhibiting corrosion, even at dilute concentrations. However, the usage of chromates will be banned by the Environmental Protection Agency (EPA) 1

18 because chromates are toxic and carcinogenic. Until now, no environmentally friendly coating system has yet been found to be as effective as CCCs. This is partially attributed to inadequate understanding of mechanisms of CCC formation and protection. The objective of this work is to develop a better understanding of localized corrosion, CCC formation, and protection by CCCs in AA7xxx-T6. Since the alloy composition, especially Cu content, plays a critical role in corrosion and CCC protection, the effect of Cu content on corrosion behavior and CCC protection is mainly focused in this work. This dissertation consists of six chapters including this introduction. Chapter 2 is a brief literature review on several topics. Physical metallurgy of AA7xxx Al alloys is reviewed first. Aspects of pitting and intergranular corrosion of Al alloys, especially the mechanisms and the effect of alloying elements and intermetallics, are reviewed and discussed. Roles of chromate in the forms of aqueous solution, CCCs, and paints in corrosion inhibition of Al alloys are covered. The formation, composition, and structure of CCCs as well as corrosion protection mechanism by CCCs are reviewed in detail. Chapter 2 also reviews Cu enrichment and redistribution as they are important issues influencing corrosion resistance and CCC protection. Finally, key unresolved issues are proposed at the end of chapter 2. Chapter 3 through 5 cover the details of the technical findings of this dissertation. They are written as stand-alone papers, and will be submitted individually for publication. Chapter 3 describes studies on the effect of Cu content on corrosion behavior of AA7xxx-T6. Electrochemical methods, such as polarization and electrochemical impedance spectroscopy (EIS), in combination with many materials characterization 2

19 techniques, such as optical metallography, SEM/EDS, TEM, FIB, STEM/Nano-EDS, and XPS, were utilized in this study. The corrosion mechanism and the exact role of Cu content in corrosion of AA7xxx-T6 were presented and discussed. Chapter 4 describes the characterization of CCCs formed on AA7075-T6. SEM/EDS, FIB, TEM, HRTEM, and STEM/Nano-EDS line profiling were utilized in this study. The composition and structure of the coatings formed on the matrix and coarse intermetallic particles were presented. Based on coating thickness measurements by FIB and polarization tests on Mg 2 Si and Al 2 CuMg ingot analogs, a sol-gel model for the CCC formation on the heterogeneous microstructure of AA7075-T6 was supported. Chapter 5 describes studies on the effect of Cu content on the protection of AA7xxx-T6 by CCCs. Electrochemical methods such as polarization and EIS as well as examinations by optical metallography and SEM were utilized in this study. The effect of acid pretreatment on CCC protection was also presented. Finally, a twofold role of Cu on CCC protection was presented and discussed. Chapter 6 summarizes the findings of this work and states conclusions on the effect of Cu content on corrosion behavior, CCC formation, and protection by CCCs in AA7xxx-T6. Chapter 6 also makes suggestions for future work and some unresolved issues. 3

20 CHAPTER 2 LITERATURE REVIEW High strength aluminum alloys are widely used in aircraft structure applications. However, high strength aluminum alloys are extremely susceptible to various forms of localized corrosion in chloride environments, such as pitting, crevice corrosion, intergranular corrosion (IGC), exfoliation corrosion, and stress corrosion cracking (SCC). In order to protect high strength aluminum alloys from localized corrosion and other forms of corrosion, protective chromate conversion coatings (CCCs) must be used. Chromates in solution are also extremely effective in inhibiting corrosion, even at dilute concentrations. The purpose of this chapter is to provide a brief review of literature work on localized corrosion, CCC formation, chromate protection in aqueous, CCC, and paint forms, and the Cu enrichment effect. Key unresolved issues will be discussed at the end of this chapter. 2.1 Metallurgy of 7xxx Series Al Alloys 7xxx series Al alloys (Al-Zn-Mg and Al-Zn-Mg-Cu) are heat-treatable and are strengthened by precipitation hardening [1]. The precipitation hardening involves the 4

21 formation of a fine dispersion of second phase particles, which hinders the movement of dislocations and thus increases the strength [1]. The precipitation hardening sequence is: SSSS G. P. zone intermediate phase equilibrium phase SSSS is a supersaturated solid solution of Al-Zn-Mg-(Cu), which is formed by solution heat treatment and subsequent quenching. Supersaturated solutes and vacancies retained in solid solution result in precipitate formation during aging heat-treatment. Under artificial aging at elevated temperature, spherical Guinier-Preston (GP) zones precipitate coherently with the lattice structure of the Al matrix [1]. GP zones are ordered solute rich clusters of less than 1 nm in size, which can dramatically increase the strength. The strength can be maximized by the peak aged T6 temper. During tempering to the T6 condition, some GP zones grow to form intermediate phase precipitates [1]. These intermediate precipitates are much larger than GP zones and are partly coherent with the Al matrix. Overaging T7 temper (higher aging temperature or longer aging time) results in the formation of coarse equilibrium phase precipitates. The coherency with the matrix is completely lost, which causes a decrease in strength. The intermediate and equilibrium phases formed during aging treatment mainly depend on alloy composition. When the Zn/Mg ratio is greater than 2, the stable precipitate phase is MgZn 2 or Mg(Zn Cu Al) 2 (η phase) with hexagonal close packed structure for Al-Zn-Mg-(Cu) alloys [1]. When the Zn/Mg ratio is less than 2, the stable precipitate phase is Mg 32 (Al Zn) 49 or Mg 32 (Al Zn Cu) 49 (T phase) with cubic structure for Al-Zn-Mg-(Cu) alloys [2]. Other than the homogeneous precipitation in the matrix, coarse precipitates of about 50 nm in size heterogeneously form on grain boundaries during aging treatment. 5

22 The formation of grain boundary precipitates typically results in the formation of precipitate free zones (PFZ) along grain boundaries. The formation of PFZ has been attributed to depletion of vacancies and solutes around the grain boundary regions [3]. The width of PFZ depends on quenching rate after solution heat treatment. The slower quenching rate results in the wider PFZ [3]. In the microstructure of AA7xxx-T6 alloys, some coarse constituents or intermetallic particles, such as Al 3 Fe, Al 2 CuMg, and Al 7 Cu 2 Fe, are present [1]. These coarse particles are formed during the solidification process and are not changed in subsequent solution and aging treatments [1]. The details about these particles will be discussed later. 2.2 Localized Corrosion of Al Alloys Pitting, crevice corrosion, exfoliation, and intergranular corrosion are various forms of localized corrosion. Since pitting and crevice corrosion share the same growth mechanism, and exfoliation corrosion is one type of interganular corrosion in wrought Al alloys, pitting and intergranular corrosion, as typical forms of localized corrosion of Al alloys, will be reviewed in this chapter Pitting Corrosion of Al Alloys Pitting corrosion is defined as localized accelerated dissolution of metals that occurs as a result of a breakdown of the otherwise protective passive film on the metal/alloy surface [4]. In an aggressive environment, typically containing chloride, pits initiate and grow in an autocatalytic manner, where the local environment within the pits 6

23 becomes more aggressive because of decrease in ph and increase in chloride concentration, which further accelerates the pit growth. The pit growth usually takes a variety of shapes in cross-section (Figure 2.1) [5]. Pit shapes can be simply divided into isotropic and anisotropic groups. Shapes in Figure 2.1a-e are isotropic, while those in Figure 2.1f are anisotropic and are called microstructural orientated pitting. The variation in pit shape could mainly depend on the microstructure of metals or alloys such as alloy composition and aspect ratio of grains. Extensive research has been conducted to study the mechanisms of pitting corrosion of stainless steels and Al alloys. Even though there are some differences in pitting corrosion between stainless steels and Al alloys, e.g., hydrogen bubbles form at the active pit surface in Al alloys, both materials basically share the similar mechanism. In general, pitting corrosion involves three stages: pitting initiation, metastable pitting, and pitting growth. Pit Initiation As mentioned above, aggressive anions such as chloride are believed to cause passive film breakdown. However, the exact mechanism of the passive film breakdown is still unclear. A number of models have been proposed to explain passive film breakdown or pit initiation [6-15]. Those models can be roughly classified into three main categories: adsorption mechanism, penetration mechanism, and film breaking mechanism (Figure 2.2) [4]. These models have been reviewed in depth in the literature [4, 16-18]. Briefly, each of these mechanisms emphasizes certain important aspects of pitting initiation. For example, the adsorption theory emphasizes the importance of adsorption of aggressive 7

24 anions like chloride ions [6, 7, 15]. A competitive adsorption of chloride ions and oxygen finally may lead to film thinning. The penetration model emphasizes the importance of anion penetration and ion migration through the passive film [8]. A point defect model as a modified or related penetration model has been developed by MacDonald and coworkers [10-12]. The point defect model addresses the transport of cationic vacancies to the metal/oxide interface controlling pit initiation instead of anion penetration. The point defect model has been fitted to experimental data such as pitting potential and induction time for pitting corrosion of Al and Al alloys in halide. However this model can not explain metastable pitting, and some assumptions such as the electrode potential and vacancy migration in extremely high electric field (on the order of 10 6 to 10 7 V/cm) are suspicious. The film breaking model involves the breakdown and repair of the passive film simultaneously [9]. Mechanical stresses due to electrostriction and surface tension cause the passive film breakdown, which is repaired rapidly. According to this filmbreaking model, pits initiate as a result of the passive film breakdown only when stable pits grow afterward. In summary, these models address important aspects of pit initiation such as aggressive ion adsorption, ion penetration and migration, and stress-induced breakdown of passive film. Although these models obtained some experimental support, no comprehensive or universal model can account for pitting corrosion in all metal/environment systems. This indicates that pit initiation is rather complicated and a combination of these models could explain pitting for a certain metal/environment system. 8

25 Metastable Pitting Metastable pits are pits that survive for a very short lifetime on the order of seconds or less. They can initiate and grow to the micron size at potentials far below the pitting potential and also above the pitting potential during the induction time prior to the onset of stable pitting [4]. Figure 2.3 shows typical metastable pit current transients on stainless steels in chloride solution under an applied anodic potential. The current increases corresponding to growth of metastable pit followed by a sharp current decrease due to repassivation process. Since mestastable pits experience initiation, growth, and repassivation, a better understanding of these three stages for the stable pit can be gained through study of metastable pitting [4]. Metastable pitting phenomenon was first observed in stainless steel in the early 1970 s [19, 20]. Frankel and coworkers used the term of metastable pitting for the first time [21]. Over the past 30 years, metastable pitting has been systematically investigated by analyzing pit current density for individual metastable pit and stochastic approaches to groups of metastable pits [4, 20, 22, 23]. These detailed studies show that the early development of stable pits appears to be identical to that of metastable pits, and the probability of stable pitting is directly correlated to the intensity of metastable pitting events [4]. Frankel [4] suggested that metastable pit growth in stainless steels is limited by the ohmic drop associated with porous pit cover of passive film remnant. Metastable pits repassivate probably when the porous cover ruptures and the pit electrolyte is diluted [4]. In contrast to a huge amount of studies on stainless steels, few studies have focused on Al or Al alloys. Pride et al. [24] studied metastable pitting on pure Al. They found that the number of metastable pits and the current spikes increase with increasing applied 9

26 potential below pitting potential and the chloride concentration. A critical transition from metastable pitting to stable pitting in Al has been found in their study. More discussion regarding this critical transition will be made in pit stability section later. Pit Growth Above the pitting potential, stable pits grow at a rate depending on alloy composition, local pit environment and pit bottom potential. Due to the autocatalytic nature of pitting corrosion, the local pit environment and bottom potential is severe enough to prevent repassivation. Pit growth can be controlled by each or a combination of three factors: charge-transfer, ohmic control, or mass transport [25-29]. For a hemispherical pit, different rate controlling factors would lead to specific relationships between current I, current density i, pit radius or depth r, time t, and potential E. Under charge transfer control, Tafel s law describes i exp E. (true when E- Eeq > 50 mv) [4]. Under ohmic control (E- E bot is fixed), it can be derived I r and i I/r 2 1/r. From Faraday s law, i dr/dt, leading to r t 1/2 and thus I t 1/2 and i t -1/2. Ohm s law determines i E [4]. Under mass transport control, according to Fick s laws, i 1/r, thus i t -1/2. i is E independent [4]. The similar i-t relationship for ohmic control and mass transport control makes it difficult to distinguish. For a 3D bulk sample, the non-steady state nature of pit deepening and the problem with accurate measurements of pit current density complicate the clear identification of the i-e relationship [4]. In a conventional measurement of i-e 10

27 relationship, current may come from several pits with unknown active surface areas and presumably is evenly distributed on the pits. However, the assumption of even distribution is not possible since different pits initiated at different potentials grow at different rates. Artificial pit electrodes, formed by imbedding a wire in epoxy have been extensively used to study iron and stainless steel behavior [30, 31]. The artificial pit electrode geometry forms a single pit in which the whole electrode area is active, generates a natural pit environment, and provides an ideal one-dimensional transport condition. For Al and Al alloys, similar to artificial pit electrodes, artificial crevice electrodes have been used since large crevice area facilitates the escape of H 2 bubbles [32, 33]. The results indicate that pits can grow either in the active state without salt film precipitation or in a salt-film-covered state. The active state is dominated by ohmic control while a salt-film-covered state is dominated by mass transport control. Other single pit techniques include the exposure of small area, laser irradiation of a small spot, and implantation of an activating species at a small spot [34-37]. These studies suggested different viewpoints of either ohmic control or mass transport control. Besides the electrochemical methods, non-electrochemical techniques have been also used. Hunkeler and Bohni [38] measured the time for pit to penetrate Al foils of varying thickness to determine the pit growth rate. They found that at fixed applied potential, pit depth d and current density i were time dependent: d t 1/2 and i t -1/2. Pit growth on Al was ohmic controlled since the growth rate was correlated to the conductivity of the electrolyte. Detailed studies of 2D pit in Al and other types of thin films by Frankel and coworkers found that the high current density increased linearly 11

28 with potential and reached a limiting value at higher potentials (Figure 2.4). Therefore, the pit growth at the beginning is controlled by ohmic control and after some time controlled by the mass transport [4, 39-41]. Pitting Stability Local pit environment and chemistry are believed to be very important for pit growth and repassivation. Among the various species present within pits such as metal cations, metal hydroxide, Cl -, and H +, acidification within pits as a result of hydrolysis is generally recognized to be a critical factor. Galvele [42, 43] calculated the acidification in 1D pits, based on metal dissolution, hydrolysis, and mass transport. He found that a critical value of the product x i (x is pit depth and i is current density), was the critical acidification within pits to sustain pit growth (Figure 2.5). This critical product can be used to explain the pitting potential and repassivation potential, and determine the current density required to initiate pitting and to sustain pit growth at a defect of a given size in passive film such as crack. Although, for some metals, other factors like chloride concentration are more important than acidification, they will roughly scale with acidification. Thus the critical value x i (sometimes I pit /r pit used) can be used as a criteria for pitting stability. Williams et al. [22] correlated pit stabilization with metastable pitting. They suggested that I pit /r pit for metastable pits formed on steels must exceed A/cm for stable growth. For Al, Pride et al. [24] found that the transition from metastable pits to stable pits occurs when I pit /r pit is higher than 10-2 A/cm. At a higher current density during pit growth, a salt film may form on the pit surface due to saturation of ionic species. For Al pits in chloride solution, this salt film 12

29 was considered to be aluminum chloride (AlCl 3 ) or aluminum oxy-chlorides such as Al(OH) 2 Cl and Al(OH)Cl 2 according to measured ph and possible hydrolysis processes [36, 44-47]. Upon salt film precipitation, as described above, the pit growth is under mass transport control. A salt film can enhance pitting stability by acting as buffer of ionic species that can dissolve into pit to sustain a severe condition in the pit environment such as high acid concentration [4]. The potential distribution in pits is considered to be another important factor to stabilize pit growth. When the IR drop is less than a critical value, pit growth stops due to repassivation, if the alloy undergoes an active/passive transition in the pit environment. [48-50]. In fact, all of the factors above might be generalized to pit growth current density, since a pit must maintain a minimum current density for stabilize growth. However, the critical pit current density and effect of environment factors need to be investigated further Criteria for Evaluation of Pitting Corrosion in Al Alloys Many electrochemical studies of pitting corrosion have found that there exist characteristic potentials. Using cyclic polarization techniques, two characteristic potentials can be determined, which correspond to pit initiation and repassivation (Figure 2.6). One is pitting potential (E P ), sometimes called critical potential or breakdown potential (E B ), above which stable pits initiate and grow rapidly. The other is repassivation potential (E R ), sometimes called protection potential, below which growing pits repassivate and stop growing. It should be noted that the values of these two 13

30 characteristic potentials can depend somewhat upon the methods used and potential scan rate [16]. Moreover, since pitting corrosion is considered to be stochastic, stochastic approaches have been developed to handle the scatter of pitting potential [51]. Both E P and E R have been extensively used to evaluate the susceptibility to pitting corrosion of various materials in a given environment [16]. It is generally recognized that materials exhibiting higher E P and E R are more resistant to pitting corrosion. However, this criterion is not satisfactory for evaluating pitting susceptibility of Al and Al alloys. For instance, in deaerated chloride solution, AA2024-T3 alloy exhibits higher E p and E R than high purity Al (99.999%), but high purity Al is more resistant to pitting corrosion than AA2024-T3. This contradiction arises from local cathodic reactions associated with noble Cu or Fe containing intermetallic particles in AA2024-T3. It is the severe local cathodic reaction on intermetallics that drives the corrosion potential (E corr ) in aerated chloride solution slightly above the pitting potential. For evaluation of pitting, another factor should be taken into account: the difference between the pitting potential and corrosion potential in aerated chloride environments. Therefore, the criterion should be that the higher difference between E p and E corr and the higher E p and E R, the more resistant to pitting corrosion of Al alloys Intergranular Corrosion of Al Alloys Intergranular corrosion (IGC) is selective attack of grain boundaries or closely adjacent regions without appreciable attack of the grains [52]. It is well known that high strength Al alloys such as AA2024-T3 and AA7075-T6 are very susceptible to IGC. Roughly speaking, IGC susceptibility of Al alloys is attributed to the electrochemical 14

31 heterogeneity between the active grain boundary regions and noble grain interiors. However, the exact mechanism of IGC is still unclear. In general, there are three different theories to explain IGC susceptibility of Al alloys: (a) galvanic couple theory; (b) precipitate free zone breakdown model; (c) anodic dissolution of grain boundary precipitates. The galvanic couple theory, the most classic explanation, attributes IGC susceptibility to the difference in corrosion potential of the anodic grain boundaries and cathodic grain interiors [52, 53]. The galvanic couple theory was first proposed by Dix and coworkers who measured the corrosion potentials for mounted grain boundary regions and grain interiors on coarsely grained Al-4%Cu alloy in NaCl-H 2 O 2 [53]. Galvele et al. proposed a precipitate free zone (PFZ) breakdown model based on the IGC studies of Al-Cu alloys in chloride solution [54]. They correlated various tempers and microstructures to IGC susceptibility (Figure 2.7) and concluded that the difference in pitting potential rather than the difference in corrosion potential was the main factor leading to IGC in Al-4%Cu alloy in some aging conditions. In the peak-aged condition, a continuous Cu depleted zone or precipitate free zone (PFZ) is present along grain boundaries with a lower breakdown potential than grain interiors and discrete grain boundary precipitates (noble Al 2 Cu particle in Al-Cu alloys). When the corrosion potential of the Al-Cu alloy is higher than PFZ breakdown potential, the continuous PFZ breaks down and preferentially corrodes leading to IGC. In the over-aged condition, Al- Cu alloys are more resistant to IGC since the PFZ disappears. Studies of stress corrosion cracking (SCC) of Al-4%Cu alloy, which is IGC under static tensile stress condition, also revealed the similar mechanism [55]. However, the PFZ breakdown model encounters 15

32 difficulties in explaining IGC susceptibility of the Al alloys where active grain boundary precipitates are present. The third theory is based on anodic dissolution. This theory attributes the IGC susceptibility to the anodic dissolution of active grain boundary precipitates. In Al-Mg alloys (AA5xxx), the formation of continuous grain boundary Mg 5 Al 8 precipitates is believed to provide a preferential anodic path along grain boundaries since the Mg 5 Al 8 precipitate is anodic to the matrix in most electrolytes. Studies of the intergranular SCC (IGSCC) behavior of AA5083 also revealed that anodic dissolution of continuous grain boundary precipitates (β) resulted in IGSCC. Buchheit et al. studied IGC and SCC of Al- Li-Cu alloys [56-58]. It was found that selective anodic dissolution of T 1 (Al 2 LiCu) subgrain boundary precipitates developed an aggressive occluded environment with a low ph, which led to continuous boundary attack. As mentioned above, Al-Zn-Mg-Cu (7xxx series) Al alloys are susceptible to IGC and SCC. Overaging (T7 temper) and RRA treatment (retrogression and reaging) can greatly improve the IGC and SCC resistance of 7xxx Al alloys. It should be noted here that according to electrochemical theory, susceptibility to IGC is a prerequisite for susceptibility to SCC. Maitra and English [59] found that 7075 alloy plate in the T6 temper exhibited two breakdown potentials. Between the active and noble breakdown potentials, the alloy exhibited IGC, while the noble breakdown potential corresponded to pitting in the matrix. However, only one breakdown potential was found in the T7 temper, which corresponded to matrix pitting. The active breakdown potential for T6 temper was postulated to be due to the breakdown of PFZ, which was Zn and Mg enriched. Overaging to the T7 temper decreased the susceptibility by reduction of the 16

33 difference between the two breakdown potentials and by the reduction in amount of solute atoms segregated [59]. Obviously, the change in microchemistry of the grain boundary regions with different tempers could be responsible for the IGC behavior of 7xxx Alloys. The PFZ on the grain boundaries in 7xxx Al alloys could be the active path for IGC. Microchemical analyses of the PFZ in 7xxx alloys under different aging conditions have been conducted over the past 30 years. Monitoring the plasma loss of the electron beam in a TEM, Doig et al. [60-63] found that Mg solute segregated at the grain boundaries in as-quenched Al-Zn-Mg alloys, but for overaged alloys, Mg solute was depleted in PFZ due to the formation of η (MgZn 2 ) precipitates on the grain boundaries. Chen et al. [64] studied grain boundary regions of Al-Zn-Mg alloys using Auger Electron Spectroscopy (AES). It was found that Mg and Zn segregated to the grain boundaries in the as-quenched temper. For the overaged temper, only a fraction of the total Mg at the grain boundaries was incorporated into the MgZn 2 grain boundary precipitates, the remainder still segregated at the boundaries. However, the AES work is suspecious since the probe of electron beam was 1 µm, which is large relative to the PFZ of about 100 nm in size. Park and Ardell [65] used Analytical Electron Microscopy (AEM) to study solute composition of PFZ in the 7075 alloy in the T6, T7 and RRA tempers. They found that Zn was depleted in the PFZ at grain boundaries similarly for all three tempers. In contrast, the depletion of Cu in PFZ was far more sensitive to the aging condition, being larger in both the T7 and RRA tempers than in the T6 temper. It was suggested that the reduction in Cu concentration in the PFZ was responsible for the increased SCC resistance of 7075 alloys in the T7 or RRA tempers. 17

34 Recently, Ramgopal and Frankel [66-68] systematically investigated IGC of AA7150. They used AEM to analyze the composition of the matrix, PFZ, and grain boundary precipitates of AA7150 in the T6 and T7 tempers. They found that the composition of the PFZ and matrix was similar in the T6 and T7 tempers, whereas Cu content in Mg(Zn Cu Al) 2 grain boundary precipitates increased dramatically in the T7 temper. Based on their measured composition data, they prepared thin film analogs to the matrix, PFZ, and Mg(Zn Cu Al) 2 grain boundary precipitates for the T6 and T7 tempers by the flash evaporation method. Their electrochemical work on the thin films in deaerated 0.5M NaCl revealed that the PFZ and matrix in the T6 and T7 tempers behaved similar in terms of breakdown potentials. However, it was found that the breakdown potential for the grain boundary precipitate increased in the T7 temper due to dramatic increase in Cu concentration. They further studied the dissolution kinetics of the matrix and PFZ thin film analogs in Cu-ion containing solution using the artificial crevice electrode technique. They developed a new explanation for the effect of temper on the IGC of AA7150. Dissolution of grain boundary precipitates in the T7 temper develops a localized microchemistry environment containing a higher Cu ion concentration than in the T6 temper. The presence of high Cu ion concentration ennobles the dissolution of the PFZ, which makes the T7 temper more resistant to IGC than the T6 temper. However, the Cu ion concentration in the localized microchemistry was not measured in their study. A further investigation is still needed to understand the role of Cu content on IGC of AA7xxx alloys. 18

35 2.2.4 Role of Alloying Addition and Intermetallics in Localized Corrosion of Al Alloys Addition of alloying elements, especially Cu, can significantly increase the mechanical strength of Al alloys such as Al-Cu-Mg alloys (2xxx series) and Al-Zn-Mg- Cu alloys (7xxx series) by precipitation hardening. Due to the limited solubility of many elements in aluminum, alloying elements are often distributed not only in the Al solid solution, but also in fine precipitates and coarse intermetallic particles. As mentioned before, grain boundary precipitates play an important role in the IGC of Al alloys. Coarse intermetallic particles play a crucial role in the corrosion behavior of Al alloys. The micro galvanic coupling between the matrix and the intermetallic particles is generally believed to result in pitting corrosion and further develop IGC into the deep structure of Al alloys. In this section, the role of alloying elements in solid solution and intermetallic particles in pitting corrosion of Al alloys will be reviewed. Alloying Elements Muller and Galvele first studied the role of alloying elements in pitting corrosion of Al-Zn, Al-Mg, and Al-Cu binary alloys in dearated 1 M NaCl [54, 69]. Zn, Mg, and Cu as alloying elements have different effects on the pitting potential of Al alloys (Figure 2.8). Pitting potential decreased greatly with increasing Zn content up to 3wt% and remained the same with further increase in Zn content. There was no influence of Mg on pitting potential. Pitting potential increased dramatically with increasing Cu content up to 5wt%. Furthermore, they studied the corrosion morphology of these three binary alloys. It was found that tunnel-like pits formed on Al-3Zn, and crystallographically shaped pits 19

36 on Al-3Mg and Al-3Cu. Sato and Newman studied metastable pitting on Al-Zn alloys. They found that the rate of pit nucleation was potential dependent regardless of the alloying addition. It was suggested that Zn addition influenced the pit growth instead of the pit nucleation events. Since the mid 1980 s, many studies have been conducted on surface chemistry and corrosion properties of stainless Al alloys containing W, Ta, Mo, Nb, and Cr [39, 70-75]. These studies provide some clues to explain the role of alloying elements on pitting potential. Due to the low solubility of above alloying elements in aluminum, thin films of supersatuarated Al binary alloys have been prepared by non-equilibrium methods such as sputter deposition. The electrochemical studies revealed that the pitting potential of aluminum can be dramatically increased by the addition of these elements. One of explanations is that enrichment in the passive film plays an important role in improving pitting resistance. Moshier and coworkers using X-ray Photoelectron Spectroscopy (XPS) conducted surface analysis of the passive films formed on Al-Mo, Al-Ta, Al-Cr, and Al-W alloys [71-74]. They found significant incorporation of the alloying elements into the passive film. It was suggested that a more protective passive film enriched with the solute atoms was responsible for improved pitting resistance by impeding the ingress of chloride ion through the passive film. McCafferty and coworkers [76] have proposed that enrichment of the solute atoms decreased the ph of zero charge, ph pzc of the passive film, which repelled chloride ion from the electrode surface. In contradiction to the passive film effect viewpoint, Smialowska [77] suggested that the solute elements in the active pit surface play the critical role instead of solute in the passive film. She proposed that the low solubility of the solute oxide in the acidic pit 20

37 environment is responsible for improved pitting resistance. Another explanation has been proposed by Frankel and coworkers [39] based on their measurement of thin film pit growth kinetics for Al-Nb, Al-Mo, and Al-Cr thin films by sputter deposition. They found that stable pits initiated at potentials only about 30 mv higher than they repassivated (Figure 2.9). It was suggested that the addition of noble alloying elements increased the pitting and repassivation potential by ennobling the dissolution kinetics of pit growth rather than the passive film effect. However, the exact mechanism by which alloying elements alter the dissolution kinetics is still unclear. Regardless, this dissolution kinetics viewpoint provides a new insight to understand the role of alloying elements such as Zn, Mg, and Cu in Al alloys in pitting corrosion. In the light of the dissolution kinetics viewpoint, Ramgopal and Frankel [33] recently studied the dissolution kinetics of Al-Zn, Al-Mg and Al-Cu binary alloys using the artificial crevice electrode technique. It was found that Zn, Mg, and Cu addition had different effects on repassivation potential and the dissolution kinetics. The addition of Cu increased the repassivation potential and lowered the dissolution kinetics. The addition of Zn decreased the repassivation potential and enhanced the dissolution kinetics. The addition of Mg had little or no effect on the repassivation potential by changing the dissolution kinetics. They suggested that the role of alloying elements was to mainly change the surface overpotential and thus shifted the repassivation potentials. Intermetallic Particles Intermetallic particles (IMCs) can be grouped into coarse intermetallic particles and fine precipitates. In Al alloys, coarse intermetallic particles form during the 21

38 solidification process, while fine precipitates including hardening precipitates in the matrix and grain boundary precipitates form during the aging process. The type and composition of intermetallics mainly vary with Al alloy composition and heat treatment [1]. The primary coarse intermetallics found in Al-Cu-Mg alloy such as AA2024-T3 are Al 2 Cu (θ), Al 2 CuMg (S), and Al 20 Cu 2 (Fe Mn) 3 [1, 78, 79]. The coarse intermetallics Al 3 Fe, Al 7 Cu 2 Fe, Al 2 CuMg, and Mg 2 Si are found in Al-Zn-Mg-Cu alloys such as AA7075-T6 [1, 78-80]. The fine precipitates for AA2024-T3 and AA7075-T6 are Al 2 CuMg and Mg(Zn Cu Al) 2, respectively [1]. As mentioned earlier, intermetallic particles play a crucial role in localized corrosion of Al alloys. It should be noted that the role of fine precipitates, mainly grain boundary precipitates, in localized corrosion of Al alloys has been reviewed in the intergranular corrosion section above. The role of coarse intermetallic particles in pitting corrosion of Al alloys will be reviewed below. The coarse intermetallic particles mentioned above can be further divided into two groups: active and noble particles relative to the Al matrix. Al 2 Cu, Al 3 Fe, Al 7 Cu 2 Fe, and Al 20 Cu 2 (Fe Mn) 3 are found to be noble to the matrix, while Al 2 CuMg and Mg 2 Si are active to the matrix [1]. Buchheit [81] compiled the corrosion potentials of various intermetallic phases in Al alloys, showing that the intermetallics exhibit different electrochemical properties from the matrix. Pits are readily found at the periphery of noble particles in Al alloys during exposure to chloride solution. It is generally accepted that noble Fe- or Cu-containing intermetallic particles act as cathodes and support oxygen reduction. As a result, a high ph local environment is established at the noble particles, which causes grooving of the surrounding Al matrix by alkaline dissolution. The alkaline attack must then somehow 22

39 switch to acid attack to result in a stable pit, which requires an acid environment. Electrochemical studies have been conducted on Al 3 Fe and Al 2 Cu [82-84]. Nisancioglu [82] found that near the open circuit potential in NaOH solution, Al 3 Fe underwent a preferential dissolution of Al, which resulted in an Fe rich surface. It was suggested that Fe enrichment on the Al 3 Fe surface is detrimental to cathodic behavior due to the formation of a protective Fe oxide. The presence of Mn and Si in Al 3 Fe can reduce the effect of Fe on both anodic and cathodic rates. Mazurkiewicz and Piotrowski [83, 84] found that Al 2 Cu underwent dissolution to form Al and Cu ions at the open circuit potential and under anodic polarization in sulfate solutions. Cu ion release was also found in Rotating Ring-Disk Electrode (RRDE) experiments on Al 2 Cu and Al 7 Cu 2 Fe at the OCP and under anodic and cathodic polarization in chloride solution [85]. The corrosion potentials for Mg 2 Si (β) and Al 2 CuMg (S) particles in chloride solution are and V SCE, respectively [81]. Both Mg-containing phases are active to the matrix and act as anode. They are susceptible to active dissolution or Mg dealloying when exposed in acidic solution or chloride solution [85-90]. Mg 2 Si phase in AA6000 dealloyed in 0.1 M phosphoric acid and MgO was found on the Mg 2 Si particles [86]. Buchheit and coworkers [85, 87-90] studied the electrochemical behavior of Al 2 CuMg phase in the form of both synthesized bulk and real phases in AA2024-T3. They found that S phase supported rapid anodic and cathodic reaction kinetics and selective dissolution of Mg and Al readily occurred under anodic and cathodic polarization. Dealloying of active S phase left Cu-rich remnants, which was cathodic to the matrix and therefore caused grooving by alkaline dissolution and then pitting at the dealloyed S phase. They also proposed that decomposition of Cu-rich remnants of S 23

40 phase resulted in Cu release and redistribution, which further accelerated corrosion of the Al alloys. This hypothesis has been supported by RRDE experiments on S particles. The detail about Cu enrichment and redistribution will be reviewed below. In summary, alloying addition and various intermetallic particles play important roles in the corrosion properties of Al alloys. They also govern the corrosion protection of Al alloys by chromate. In 7xxx series Al alloys, the role of Cu content and Cucontaining particles in corrosion behavior and chromate protection is not at all clear. The next section is dedicated to the review of literature on the chromate protection in Al alloys. 2.3 Role of Chromate in Corrosion Inhibition of Al Alloys In order to protect high strength Al alloys from localized corrosion, a coating system is applied. A standard coating system uses a chromate conversion coating (CCC) covered by organic coatings including primer and topcoats (Figure 2.10). CCC can offer corrosion protection and enhance adhesion of organic coatings. Chromates in solution are also extremely effective in inhibiting corrosion, even at dilute concentrations [91, 92]. Chromates as pigment are incorporated into paint used for further protection of Al alloys. Many studies have focused on the corrosion inhibition of Al alloys by chromates and obtained a better understanding of the mechanism Chromate in Solution Hexavalent chromium oxo-species (denoted as Cr(VI) or Cr VI ) mainly include chromic acid (H 2 CrO 4 ), dichromate (Cr 2 O 2-7 ), bichromate (HCrO - 4 ), and chromate 24

41 (CrO 2-4 ). The speciation of Cr VI depends on ph and concentration [93]: dichromate is predominant at low ph (ph 2~6) and high concentration of Cr VI ; bichromate is predominant at low ph (ph 2~6) and low concentration of Cr VI ; chromate is predominant at high ph (alkaline); below ph 1, the main species is H 2 CrO 4. In this review, chromates is used to indicate all forms of Cr VI oxoanions (i.e. all except chromic acid), unless it is noted specifically. Chromate species in aqueous solution obey the following equilibria [94]. H 2 CrO 4 = HCrO H + K = 4.1 (2.1) HCrO - 4 = CrO H + K = (2.2) Cr 2 O H 2 O = 2 HCrO 4 - K = (2.3) 2CrO H + = Cr 2 O H 2 O K = (2.4) When added into a chloride containing solution, chromates can shift the pitting potential of pure aluminum in the noble direction for open surfaces [24, 95] and can reduce or even stop crevice corrosion in occluded sites [96]. Chromates can change the behavior of metastable pitting, and thus affect pitting corrosion [24]. Studies of pit growth and repassivation of pure Al using the foil penetration technique [38] and image analysis of 2D thin film pits [92] found that pit growth was affected only when the concentration ratio of chromate to chloride was 1:1 or more. There is a dispute whether chromate is an anodic or cathodic inhibitor for Al and Al alloys. As an anodic inhibitor, the chromate should affect the anodic dissolution kinetics of Al or Al alloys or form a protective film. Akiyama et al. [32] and Meng et al. [97] used the artificial crevice electrode technique and found that chromate did not affect the anodic dissolution kinetics. They suggested that chromate inhibition must be 25

42 something other than anodic inhibition. They also suggested that the local ph might not change with the addition of chromate during pit growth since the charge density ratio of H 2 evolution to anodic dissolution remained constant (about 15%) with the chromate concentration, which is against the hypothesis of chromate inhibition by local ph change. Sehgal et al. [98] found that a small amount of dichromate ions can effectively inhibit pitting at the open circuit potential. It was also found that the presence of dichromate ions greatly reduced the cathodic portion of the polarization curves of AA2024-T3 alloy in oxygen bubbled 1 M NaCl with strong oxidizing agents, whereas there was no effect on the anodic portion (Figure 2.11). These results support the notion that chromates act as cathodic inhibitors, particularly inhibiting the oxygen reduction reaction on the cathodic sites of Al and Al alloys. Hence, chromates protect Al and Al alloy from corrosion through inhibiting the cathodic reactions under the open circuit condition. Pride and Scully [24] studied metastable pitting of Al in the presence of chromates. It was found that a small amount of chromate on the order of 10-5 M profoundly reduced the metastable pitting nucleation rates, background current density, and transient peak current density (Figure 2.12). The authors concluded that chromates greatly reduce the probability for metastable pits to reach a critical I pit /r pit value of 10-2 A/cm, above which stable pit can survive and grow. There are four major mechanisms for chromate inhibition [96, ]: (a) ph pzc model (b) Competitive adsorption model (c) Chromate incorporation into passive film (d) Duplex bipolar/repassivation model 26

43 Kendig and coworkers [99] measured the zeta potential (related to surface charge) and ph of zero charge (ph pzc ) for an anodized aluminum oxide of as a function of ph and concentration of chromates. It was found that chromates significantly lowered the surface charge and reduced ph pzc on the surface of aluminum oxide without the chromate reduction to inert Cr III oxide. This decrease in ph pzc would repel aggressive chloride ions from the surface. Heine et al. [100] found that passive film formed on aluminum in chromate solution exhibited lower ionic resistance and higher electronic resistance than amorphous thermal films of the same thickness. It was suggested that the formation of some microcrystalline γ-al 2 O 3 at the film/solution interface might be responsible for the low ionic resistance, while the higher electronic resistance was attributed to the inclusion of H + ions in the film and removal of some Al 3+ ions in the presence of chromate. Bohni and Uhlig [95] found a linear relationship between the logarithm of the Cl - activity and the logarithm of the CrO 2-4 activity required to shift the pitting potential of Al to noble value of 0.8 V vs NHE (reversible potential of oxygen reduction). A competitive adsorption model was proposed by McCafferty [96]. This competitive adsorption model considers that aggressive ions (Cl - ) and inhibitive ions (CrO 2-4 ) compete for adsorption sites on the metal surface. Breakdown of the passive film occurs as pitting on an open surface or crevice corrosion on occluded sites, only if the ratio of the surface coverage of aggressive to inhibitive ions exceeds a certain critical value, θ crit, which was supported by XPS surface analysis [96]. This model successfully predicts the linear relationship between log(a Cl- ) and log(a Cr6+ ), which was observed by Bohni and Uhlig [95]. Also, this model is 27

44 in good agreement with the linear decrease in pitting potential with the logarithm of the Cl - activity at constant CrO 2-4 concentration. Edeleanu and Evans [101] first attributed chromate inhibition to a redox reaction between chromate and aluminum to form aluminum oxide and Cr III ions. XPS analysis [102] revealed that the passive film formed on aluminum in 0.2 M Na 2 CrO 4 solution contained hydrated Cr III and Al III oxide, and adsorbed Cr VI species. The structure of passive film in chromate solution was found to be composed of an underlying barrier layer of amorphous γ-al 2 O 3 and outer microcrystalline Cr III oxide layer. The reduction of Cr VI into Cr III at flaw sites within the passive film in chromate solution, blocking pores and defects, was confirmed in the studies of anodized aluminum in chromate solution [ ]. In 0.1 M chloride solution containing 0.2 M chromate, however, a large amount of chloride species was detected in the film, no matter whether chlorides were initially present in the solution or were added only after 18 h immersion of the samples in the chromate solution. These results indicate that the incorporation of chromate and reduction to Cr III is unable to prevent the Cl - ingress in the passive film, which is not in agreement with the competitive adsorption model described previously. The authors attributed the chromate inhibition effect to the increase in potential at which a critical Cl - concentration was established at the Al/oxide interface. Similar to the chromate studies, a study of the interaction of molybdate anions with the passive film on Al found that the reduced MoO 2 state had no significant role in protecting the Al substrate [111]. Instead, the molybdate layer selectively impeded the ingress of anions such as Cl -, O 2- and OH -, which restricted not only the growth of passive film and also the Cl - reaching the Al/film interface to inhibit pitting [111]. 28

45 Recently, Clayton and coworkers [110] proposed a duplex bipolar/repassivation model for the interaction with chromate with Al surface (Figure 2.13). They used XPS to study the anodic film formed on AA1060 exposed to chromate and chloride solution. They found a structure change in the hydrated Al anodic film by chromate and chloride ions, which was consistent with deprotonation. Chromate induces the deprotonation process of aluminum hydroxide to form protective aluminum oxide, while chloride induces deprotonation process of aluminum hydroxide to form aluminum oxyhydroxide, which has poor corrosion resistance. When Al is exposed to chromate prior to chloride, chromate adsorption on passive aluminum oxide acts as a bipolar layer to inhibit the chloride ingress by the bipolar model. When scratched Al is exposed to chloride and chromate solution, the chloride attack in the scratch is inhibited by the formation of a stable Al(III)-OH-Cr(VI) compound. In contrast to numerous studies of chromate inhibition on pure Al, only a few studies of chromate inhibition on Al alloys have been conducted since heterogeneous microstructure of Al alloys such as AA2024 complicates the chromate inhibition [108, 112, 113]. As described previously, intermetallic particles play an important role in localized corrosion of Al alloys. Dealloying of Al 2 CuMg in AA2024-T3 in chloride solution causes Cu enrichment and redistribution, which further accelerates the corrosion. In the presence of chromate, dealloying of Al 2 CuMg phase is strongly inhibited [108]. Kendig et al. [112] studied the potentiostatic current transients for Al 2 Cu and Al 2 CuMg model intermetallics at the OCP of AA2024-T3 in 0.01 M NaCl with and without 0.01 M chromate. It was found that chromate rapidly passivated cathodic Al 2 Cu and anodic Al 2 CuMg phases. The authors attributed rapid passivation of intermetallics to rapid 29

46 chromate adsorption on intermetallics. Clark et al. [113] investigated chromate effects on Cu and AA2024-T3 using a galvanic corrosion approach. They suggested that chromate inhibition of oxygen reduction on Cu-rich intermetallics was attributed to the formation of Cr(III) monolayer on Cu. In summary, numerous studies have been carried out to develop a better understanding of the mechanism of Al alloy corrosion inhibition by chromate. However, little attention has been paid to the mechanism of chromate protection of 7xxx Al alloys Chromate Conversion Coatings Chromate conversion coatings (CCC) are applied to Al alloys to improve corrosion resistance and to enhance subsequent adhesion of organic coatings. Basically CCC acting as a barrier layer can separate aggressive environments from the Al substrate underneath. Another unique function that CCC can also provide is self-healing once fresh Al substrate is exposed by a scratch or damage in the coating. Although CCC is extremely effective in corrosion protection, the usage of chromate will be banned by the EPA s regulations due to the toxic and carcinogenic nature of chromate. In recent years, considerable efforts have been put into developing environmentally friendly replacement coatings, but so far no coating has been found to be as effective as CCC. This is partially attributed to inadequate understanding of CCC formation and protection. Therefore, it is necessary to better understand the mechanism for CCC formation and protection in Al alloys, particularly 7xxx Al alloys such as 7178 and 7075 widely used in aircraft structure applications. In this section, several important aspects of CCC such as CCC formation and protection will be reviewed in detail. 30

47 Formation of CCC The formation mechanism of CCC is very complicated, because the whole coating process involves many complex factors, such as ph and composition of the chromating solution, the composition and microstructure of Al alloys, and the pretreatment prior to coating. Before the coating process, Al alloy samples are cleaned in alkaline solution, thoroughly rinsed in water, then cleaned in acidic solution, and thoroughly rinsed in water again. Cleaning in alkaline solution is to remove grease, oil and soils, while cleaning in acidic solution is to reduce the oxide layer and activate the metal surface. The pretreatment procedure has a significant effect on the subsequent chromating process and resulting properties of CCC since the pretreatment changes the sample surface and could cause Cu enrichment and redistribution [ ]. After the pretreatment, CCCs are applied to Al or Al alloys simply and efficiently by spraying or immersion in the chromating bath for a few minutes. Commercial Alodine 1200S bath with about ph 1.6 is widely used in industry. The composition of Alodine 1200S solution is listed in Table 1 [117]. The species in Alodine solution can be classified into three groups in terms of functions during the chromating process: chromates acting as a coating forming agent, fluoride as an activator, and ferricyanide as an accelerator. It has been generally accepted that CCCs form on Al or Al alloys via a redox reaction between chromates and Al. 2Al 2Al e - (2.5) Cr 2 O H + + 6e - 2Cr(OH) 3 + H 2 O (2.6) Overall reaction: Cr 2 O Al + 8H + 2Al Cr(OH) 3 + H 2 O (2.7) [4] 31

48 Other possible reactions from the literature are also listed here. 2HCrO Al + 8H + 2Al Cr(OH) 3 + 2H 2 O (2.8) [118] Cr 2 O Al + 2H + + H 2 O 2CrOOH + 2AlOOH (2.9) [119, 120] Without the presence of fluoride ions, the formation of CCC coatings is very slow. The role of F - ions in coating formation has been generally believed to be twofold [116, 119]. First, the fluoride ions at low ph attack the initial Al oxide film (Al 2 O 3 ) and activate the Al surface for the subsequent redox reaction and coating deposition. Second, the fluoride ions can dissolve AlOOH in the formed CCC and provide the bulk chromating solution access into the Al/CCC interface to prevent the passivation of Al surface and promote the further coating growth. The dissolution reactions by fluoride ions are listed as follows, Al 2 O HF 2AlF 3-6 (soluble) + 6H + + 3H 2 O (2.10) [116, 119] AlOOH + 6HF AlF 3-6 (soluble) + 3H + + 2H 2 O (2.11) [116, 119] Ferricyanide K 3 Fe(CN) 6 is an accelerator for CCC formation. When ferricyanide is added into the coating solution, coating corrosion resistance is greatly improved as well as both the coating weight/thickness and the coating growth rate are increased [115, 116, ]. The accelerating function of ferricyanide is far from clear. Treverton and Davies [122] proposed that a CrFe(CN) 6 layer adsorbed on the surface of growing hydrated Cr III oxide precipitates inhibited the chromate adsorption and as a result more chromates were available for further redox reaction. However, this mechanism has not been supported experimentally. Xia and McCreery [123, 124] using Fourier Transform Infrared (FTIR) and Raman Spectroscopy studied the chemical structure of accelerated CCCs and the roles of ferricyanide in the CCC formation on AA2024-T3. Their results 32

49 revealed that CCC growth is mediated by Fe(CN) 6 3-/4- (Figure 2.14). Without the presence of Fe(CN) 6 3-, direct redox reaction between Al and chromate proceeds slowly. However, Al reacts quickly with Fe(CN) 6 3- to produce ferrocyanide (Fe(CN) 6 4- ) and Al 3+, and Cr(VI) is quickly reduced to Cr(III) by Fe(CN) 4-6, which oxidizes back to Fe(CN) 3-6. As a result, Al is oxidized into Al 3+ while Cr VI is reduced into Cr III at a fast rate. The authors also suggested that in principle, any redox system similar to ferricyanide/ferrocyanide, which has a redox potential between that of Cr VI /Cr III and that of Al/Al 3+ and exhibits fast redox kinetics with Cr VI and Al, can act as a mediator through the mediation mechanism. Based on this idea, the authors tested IrCl 2-/3-6, V 3+/2+ and Fe 3+/2+ in the coating process. It was found that all of them accelerated the coating formation and the coatings had similar properties to ferricyanide accelerated CCC, although these redox systems accelerated less than the ferricyanide/ferrocyanide system. This result indirectly supports the ferricyanide mediation mechanism. CCC Nucleation and Growth Model Although CCC forms via a redox reaction as described above, the mechanism of CCC nucleation and growth on Al or Al alloys is still unclear. Heterogeneous microstructure of Al alloys further complicates the CCC growth kinetics. For CCC nucleation and growth on pure Al, several models below have been proposed. (a) Uniform growth model (b) Sideways growth model (c) Sol-gel growth model 33

50 Katzman et al. [119] proposed a uniform coating growth model based on Auger Electron Spectroscopy (AES) results of CCCs formed on commercial purity Al in CrO 3 + NaF chromating solution. They attributed the steady-state coating thickness to the equilibrium reached between the coating growth and the dissolution of surface deposited CrOOH by HF. The uniform growth model suggested that initial Al oxide was completely dissolved by HF and the bare Al reacted with chromate via a redox reaction to form an amorphous mixture of AlOOH and CrOOH, which precipitated uniformly on the Al surface to form the CCC. AlOOH in the coating was dissolved by HF leaving behind the less soluble CrOOH in the coating. When CCC continued to grow, anodic (oxidation of Al into Al 3+ ) and cathodic (reduction of Cr VI into CrOOH) reactions occurred separately at metal/ccc and CCC/solution interfaces, respectively. Obviously, this model has difficulty explaining the charge transfer through the wet coating during the CCC growth. Brown et al. [ ] used Transmission Electron Microscopy (TEM) to examine ultramicrotomed cross-sections of CCC formed on high purity Al (99.98% and %) in the same chromating solution as Katzman et al. [119]. The authors improved the uniform model for CCC nucleation and growth by taking the flaw sites in Al into account. The modified uniform growth model emphasizes electron tunneling on reduced Al oxide film in contrast to completely dissolving Al oxide film by HF in the Katzman s model. The thickness of the initial Al oxide film is reduced by HF until electron tunneling is possible. The reduction of Cr VI occurs either over the general surface covered with the reduced Al oxide film by electron tunneling or on a flaw by local conduction. For less purity Al (99.98%), CCC is formed preferentially on the 34

51 cathodic impurity segregated flaw sites, such as dislocations and grain boundaries. During the coating growth, cathodic reduction of Cr VI occurs at these flaw sites to form the hydrated Cr III oxide, while anodic oxidization of Al occurs over the general Al surface to reach a steady state between the dissolution of Al oxide film by HF and reformation of Al oxide film by oxidizing Al. The hydrated Cr III oxides grow laterally and perpendicularly to cover the whole Al surface with CCC eventually. For high purity Al ( %), the impurity segregation is negligible. Electron tunneling on the reduced Al oxide film plays a dominant role in the formation and growth of CCC. A uniform coating is formed on most of the Al surface and continues to grow except for some discrete holes. In these hole regions, the Al oxide films are thick and thus the coating formation is hindered since electron tunneling in thick oxides is difficult. In the cathodic regions, CCC deposits above the thinned alumina and prevents the solution ingress to the underlying alumina layer. In the anodic regions like holes, anodic dissolution of Al occurs. However, this model cannot explain how CCC at the cathodic regions further grows by electron tunneling through thick hydrated chromium oxide layers. SEM and AFM observation of the nucleation and early-stage growth of CCC on Al and Al alloys revealed that the structure of CCC exhibits spherical particles or nodules [132, 133]. Arrowsmith et al. [133] proposed a sideways growth model. Spherical particles nucleate from a strand-like solid probably due to the condensation of Cr(OH) 3, and grow by further polymerization at the surface. The spherical Cr(OH) 3 particles grow sideways and merge to form a CCC layer. The Al surface keeps in contact with the fresh solution through the gaps between the spheres, and the CCC continues to grow. This sideways growth model cannot explain the observation [133] that a new set of spherical 35

52 particles formed on top of the first layer after a long time of immersion. Nonetheless, this model might account for good corrosion protection and good paint adhesion of CCC since spherical particles have relatively high surface area, which can efficiently adsorb Cr VI and promote adhesion. Furthermore, this model implies that the sol-gel condensation and polymerization play an important role in CCC nucleation and growth. Recently, CCC formation has been described as a sol-gel process based on the fact that CCC has many characteristics found in dip or spin coated xerogel films [134, 135]. In contrast to the forced hydrolysis in sol-gel processing, colloidal Cr 3+ produced via a redox reaction undergos hydrolysis, condensation and polymerization, followed by drying to form CCC coatings (Figure 2.15). This viewpoint provides new insight to the study of CCC formation. The CCC nucleation and growth models described above are based on pure Al substrates. Al alloys, such as AA2024 and AA7075, have heterogeneous surfaces consisting of Al matrix and intermetallic particles, which complicate the nucleation and growth process of CCC. AFM studies of the nucleation and early growth of CCC on AA2024-T3 showed that Al-Cu-Fe-Mn intermetallic particles were covered faster with CCC relative to Al matrix, while Al-Cu-Mg exhibited lower coating coverage relative to Al matrix [132]. This is attributed to the electrochemical activity of the intermetallic particles. Al-Cu-Fe-Mn is cathodic to Al matrix, while Al 2 CuMg is anodic to Al matrix. A similar AFM study of nucleation and growth of CCCs formed on AA2024-T3 in combination with Field Enhanced Scanning Electron Microscopy (FE-SEM) and TEM was conducted by Brown and Kobayashi [131]. They concluded that CCC formation and growth on intermetallics strongly depended on the size, shape, and composition of 36

53 intermetallics. The exact mechanism for CCC nucleation and growth on Cu-containing Al alloys, particular in 7xxx Al alloys, needs further investigation Composition and Structure of CCC Coatings CCC is about 0.2~1µm thick (Alodine treatment for 1~3 min) [136] and amorphous in nature [137, 138]. CCC exhibits a mud-cracking morphology (Figure 2.16) due to the coating dehydration and shrinkage during aging. In the past 30 years, CCC formed on Al or Al alloys has been examined by many physical and chemical methods, especially surface analytic techniques, to investigate the composition and structure of CCC. It has been generally established that CCC formed on Al or Al alloys by immersion in commercial Alodine 1200S bath is mainly composed of Cr (hydrated Cr III hydroxide and adsorbed Cr VI ) and O (Figure 2.17) [136]. Little Al III is found in CCC. Several models for the CCC structure have been proposed based on the analyses by surface sensitive techniques and electron microscopy [118, 120, 123, 133, 136, 137, 139]. Among them, the model of CCC structure proposed by Hughes has been well accepted (Figure 2.18) [120]. At the Al/coating interface, there exists a thin barrier layer composed of mixed Cr and Al oxide containing some amount of F. The bulk CCC is polymeric hydrated Cr III hydroxide. Cr VI species are adsorbed in the outer layer of CCC. Ferricyanide is present through the CCC. Accurate measurement of Cr VI concentration in CCC has become a focus of research because adsorbed Cr VI is believed to be responsible for the self-healing property of CCC, which will be discussed later. Many measurements of the Cr VI concentration were conducted using XPS and about 9% Cr VI concentration was estimated in CCC 37

54 formed on high purity Al [140]. The accuracy of the XPS data on Cr VI concentration is suspect since it is now known that the photo-reduction of Cr VI into Cr III occurs during the XPS measurements [103, 141]. Another problem with XPS measurements is that ultrahigh vacuum could cause the CCC structure to change. Fortunately, X-ray Absorption Spectroscopy (XAS) can overcome the problems that XPS encounters. XAS is performed in ambient conditions employing higher energy X-rays, which greatly reduces photo-reduction of Cr VI into Cr III. XAS spectra involve two regions: extended X- ray absorption fine structure (EXAFS) and X-ray absorption near edge structure (XANES). EXAFS is fine oscillations at 30~1000 ev above an X-ray absorption edge, which can be used to study the coordination and bond length of Cr to its nearest O neighbour. The energy range in XANES is 0~30 ev above the X-ray absorption edge, which can be used to study the Cr concentration and oxidation state. In XANES, the edge height is proportional to the concentration of total Cr. The distinct pre-edge peak corresponds to Cr VI (Figure 2.19). The pre-edge height relative to the edge height is proportional to the Cr VI concentration. XANES of CCCs revealed approximate 20~30% of the total Cr as Cr VI [136, 142, 143]. A Cr VI concentration of 25% was obtained by ultraviolet absorption [123]. As mentioned above, EXAFS has been used to obtain information about the coordination of Cr with O. EXAFS analysis of CCC revealed that Cr VI is in tetrahedral coordination with a Cr-O distance of 1.71 Å, while Cr III is in octahedral coordination with a Cr-O bond distance of 1.99 Å [136]. Raman spectroscopy analysis by Xia et al. [123] revealed that the Cr-O stretch from chromium species in CCC is different than that in either CrO 2-4 or Cr 2 O 2-7. Cr-O bonding in CCC has been suggested to be the covalent 38

55 bonding between polymeric Cr III hydroxide and CrO /Cr 2 O 7 (figure 14) [123]. This covalent bonding is found to be reversible and allows subsequent release of Cr VI from CCC into the solution resulting in self-healing, which will be discussed later. Despite the huge amount of studies of CCC formed on Al or Al alloys, only few efforts have been put on studying the composition and structure of CCC formed on intermetallic particles in Al alloys. As described previously, intermetallics, especially Cucontaining particles, govern the corrosion properties of high strength Al alloys. These intermetallic particles have been believed to affect the properties and corrosion protection of CCC formed on Al alloys. However, the knowledge regarding the formation, composition, and structure of CCC formed on the intermetallics is much less than that of CCC formed on pure Al or Al alloy matrix. Part of the reason is the practical difficulty in characterizing the coating formed on the intermetallic particles on the order of 1~10 um in size and obtaining depth composition distribution due to the limited spatial resolution of many surface analytic techniques. In order to overcome the spatial limitation of surface analytical techniques and obtain composition and thickness information for CCCs on intermetallics, large-area intermetallic model samples in the form of bulk ingots and thin films have been studied. McGovern et al. [144] investigated the CCC formation on a specially cast Al-Cu-Mg ingot using Raman spectroscopy. The 860 cm -1 peak was used to monitor the Cr(VI)-O- Cr(III) bond in CCC. They found that CCC formation was suppressed on Al-Cu-Mg phases and was lower as the Cu content in the Al-Cu-Mg phase increased, as a result of passivation by adsorbed ferricyanide. Juffs et al. [145, 146] studied CCCs formed on the macroscopic couples. The intermetallic castings of Al 3 Fe, Al 7 Cu 2 Fe and Al 2 Cu (θ) were 39

56 separately coupled to AA1100. Analyses by surface analytical techniques revealed that coatings on the matrix were ten times thicker than over the intermetallic phases and increased linearly with immersion time. On the intermetallic phases, the decomposition of ferri/ferrocyanide and the fluoride attack were found. Vasquez et al. [147, 148] investigated CCCs formed on AA2024-T3 alloys and thin film analogs of Al 2 Cu (θ), Al 2 CuMg (S), Al 20 Cu 2 (Fe Mn) 3, and matrix (Al-4Cu) using a variety of surface analytical techniques. A refined view of CCCs formed on AA2024-T3 was proposed [148]. It has suggested that CCCs formed on AA2024-T3 are heterogeneous. The thickness of coatings on θ and S intermetallic particles was one tenth of that of coatings on the matrix. The coating on Al 20 Cu 2 (Fe Mn) 3 particles was rough and exhibited an island structure. The cyanide enriched oxide films covering intermetallics were different from those on the matrix, and Cr was depleted in oxide film formed on Cu rich intermetallics. The composition and structure of CCC formed on constituent intermetallic particles in Al alloys, particular 7xxx alloys is still unknown and further investigation using higher spatial resolution tools is needed Corrosion Protection Mechanisms Accelerated corrosion tests such as salt spray tests are usually used to evaluate the corrosion protection properties of inorganic or organic coatings. Alodine CCC coated Al or Al alloys can survive 168 h salt spray tests, even up to 1000 h in some cases without significant corrosion occurrence. CCCs can improve the pitting potential [149, 150]. The highly improved corrosion resistance of CCC coated Al or Al alloys has been attributed to the passive nature of CCC coating. 40

57 It is generally accepted that CCC provides a barrier insulation separating aggressive environments from the Al substrate and inhibits the cathodic reactions (in most cases oxygen reduction in neutral solution). Although CCCs exhibit a mud-cracking morphology, there could exist a barrier type film at the bottom of CCCs next to the metal/coating interface as described previously. In the case where Cr VI species leach out completely, the insoluble coating still remains the corrosion protection at some level [151], which suggests the barrier properties of CCC. In addition, as suggested by Clayton [110], CCCs could act as bipolar membranes according to Sato s bipolar model (Figure 2.20) [ ]. Sato pointed out that the metal oxide film formed on metal as a corrosion product is intrinsically anion selective. Anion selectivity means that only anions can pass through the film. The surface charge on the film would change to a negative value when CrO 2-4 or MoO 2-4 ions are adsorbed and the film becomes cation selective, which means only cation can adsorb and pass through the film. Sato s model predicts the best condition would be a bipolar membrane where the inner layer is anion selective while the outer layer is cation selective. The structure of CCCs matches this condition since the inner layer is Al III /Cr III oxide and the outer layer is enriched in Cr VI. This bipolar membrane would discourage the adsorption and ingress of aggressive Cl - ion, as well as the egress of metal cations. Kendig et al. [99] investigated the surface charge in the anodized Al by adsorbed Cr VI and confirmed that chromate significantly lowered the surface charge. This bipolar membrane function may further improve corrosion protection by CCCs. Another unique corrosion protection property of CCCs is self-healing. Selfhealing refers to the ability of CCCs to heal small chemical or mechanical defects such as 41

58 scratches or pits. It is well known that when scribed CCC Al alloy panels are salt spray tested, no corrosion is observed in the scribe after 168 h. The self-healing of CCCs has been attributed to soluble or leachable Cr VI stored in CCCs. The CCC self-healing process is thought to involve several steps (Figure 2.21). Leachable Cr VI species release from CCCs to the solution, migrate to the active damage sites such as scratches or pits, reduce to insoluble Cr III oxide or hydroxide, precipitate and block the exposed metal [91, 155]. Since the release of Cr VI is the first step of CCC self-healing process, the Cr VI release kinetics is important. Using radioactive tracer technique, Glass [156] first reported that the leaching rate of Cr VI species in NaCl solution was very rapid at the beginning and decreased gradually with immersion time until Cr VI concentration in the solution reached an equilibrium value. A study of the adsorption of Cr VI on the Al surface found that the adsorption rate at the beginning was rapid and decreased with immersion time until an equilibrium value was reached [91]. It is obvious that the adsorption and leaching of Cr VI are correlated to each other. Recently, Xia et al. [157] used ultravioletvisible spectroscopy to study the storage and release of soluble Cr VI species from CCCs. The Cr VI concentration in the solution over a CCC reached an equilibrium value, which depended on ph, ionic strength, and the ratio of the CCC surface area to the solution volume (A/V) (Figure 2.22). The adsorption of Cr VI by synthetic Cr III hydroxide formed a Cr III -O-Cr VI mixed oxide (Figure 2.23) and led to an equilibrium concentration of Cr VI in the solution. The authors proposed that the mechanism for storage and release of Cr VI was similar to Langmuirian adsorption-desorption equilibrium of Cr VI on a porous and insoluble Cr III hydroxide matrix as shown in the following reaction. 42

59 Cr III -OH (solid)+hcr VI O - 4 (aq)+h + (aq) Cr III -O-Cr VI O 3 H (solid) + H 2 O (2.12) [157] In this reversible equilibrium, adsorption is favored at low ph during CCC formation while desorption is favored in field conditions. The release kinetics might be controlled by diffusion of Cr VI in CCCs [158]. Zhao et al. [91] studied the migration of soluble Cr VI species from AA2024-T3 coated with a CCC to the uncoated alloy. The authors designed an artificial scratch cell in which a CCC coated sample and a bare sample separated by an 1.8 mm gap were exposed to 0.1 M NaCl for 96 h. Raman spectroscopy showed that Cr VI was released from the CCC, migrated across the gap, and protected the exposed AA2024-T3. The Cr transferred onto the untreated surface was made up of Cr(VI) and Cr(III) in a similar proportion to that found on Alodine treated surfaces. The Cr VI concentration in the solution was M. Acceleration by ferricyanide retained in CCCs was thought to be responsible for the rapid formation of deposits on the bare sample in the presence of so low concentration of released Cr VI. Jeffcoate et al. [155] used XANES to measure total Cr and Cr VI of CCCs in exposure to water and sulfate. It was found that the total Cr content of exposed CCC decreased by 13%, and amount of Cr VI decreased by 69%. A significant deposition of Cr III on the etched exposed Al close to the CCC was found. In summary, studies of the aspects of CCCs on Al or Al alloys suggest that CCCs have many unique characteristics such as barrier properties, the ability to store and release soluble Cr VI, and the self-healing function. These characteristics of CCCs can be used to guide the development of chromate-free coating systems. 43

60 2.3.3 Chromate in Paint As mentioned previously, a standard coating system uses CCC covered by paint. The main role of the paint is to separate the aggressive environment from the Al alloy substrate. Chromate (SrCrO 4 ) added into primer as pigment can improve the corrosion protection performance of the primer coating and the whole coating system [159, 160]. Corrosion inhibition offered by SrCrO 4 pigment has been generally believed to be associated with release of Cr VI from sparingly soluble SrCrO 4 salt (K sp = ) [161, 162]. Similar to CCCs, the primer containing SrCrO 4 pigment also has self-healing property, which was described previously. The release of Cr VI is governed by the chemical dissolution of SrCrO 4 in contrast to the equilibrium release of soluble Cr VI stored in CCCs. 2.4 Cu Enrichment and Redistribution The critical role of intermetallic particles in localized corrosion of Al alloys was described previously. Many studies revealed that Cu-containing intermetallic particles govern the corrosion and CCC protection of high strength Al alloys due to the noble nature of the Cu rich particles. In chloride environment, Cu-rich particles, particularly Al 2 CuMg (S) particles, often lead to Cu enrichment and redistribution, which in turn is detrimental to corrosion resistance and CCC protection. Buchheit and coworkers [87] attributed Cu enrichment in AA2024 to dealloying of S-phase, which accounts to about 60% of total intermetallic particles (Figure 2.24). S phase is susceptible to dealloying in acidic solution or chloride solution. Selective dissolution of Mg and Al leaves behind Cu rich sponge remnants. Buchheit et al. further 44

61 pointed out that Cu redistribution was attributed to the formation and redeposition of Cu ions although the corrosion potential of AA2024 is well below the reversible potential for Cu/Cu 2+. There are two possible explantions for this seemingly thermodynamic contradiction. Buchheit et al. [87] proposed that the Cu rich sponge remnants undergo physical coarsening, which results in nonfaradaical liberation of mechanically and electrically isolated metallic Cu clusters. Metallic Cu clusters suspended in the solution or isolated in the corrosion product can achieve a corrosion potential that is not controlled by the alloy potential. In an aerated solution, metallic Cu is oxidized into Cu ions. Dissolved Cu 2+ ions can drift around by solution convection and redeposit on the alloy surface, reducing back to metallic Cu. This then leads to the localized corrosion in other places. This mechanism is partially supported by rotating ring-disk electrode experiments conducted on the S phase at open circuit, anodic and cathodic polarization at potentials well below the reversible potential for Cu/Cu 2+ [85]. Sieradzki and coworkers [163, 164] proposed a different viewpoint that Cu ions are formed directly from Cu rich sponge remnants on the alloy surface. In this viewpoint, the curvature effect is thought to be responsible for the formation of Cu ions rather than the liberation of metallic Cu clusters. The curvature of Cu rich clusters on the surface shifts the reduction potential for Cu in the anodic direction (equation 2.13), dramatically when the radius r is very small. E Cu = E Cu 2γ CuΩ nfr Cu (2.13) [164] where ECu is the potential of the Cu rich remnant, E Cu is the reversible potential for Cu, n is 2 equiv/mol, F is Faraday s constant (96487 C/equiv), Ω Cu is the molar volume of Cu, r 45

62 is the radius of the surface curvature, and γ Cu is the surface energy of Cu. The formation of Cu ion is possible at the alloy corrosion potential when the radius is about 40 nm [164]. This direct Cu oxidation viewpoint also could explain the Cu ion formation in the deaerated solution, which Buchheit s viewpoint cannot explain. Besides the Cu enrichment and redistribution from the S phase, other arguments have been made to explain Cu surface enrichment, which do not require any Cu oxidation or long-range redistribution of Cu from the S phase. An argument is that Cu on the surface around the intermetallics comes from the surrounding matrix [163, 165]. According to this viewpoint, the active intermetallic particles such as the S phase rapidly dealloy leaving behind the porous Cu rich remnants, which act as local cathodic sites. The oxygen reduction reaction occurring at these cathodic sites increases local ph to alkaline. In local alkaline solution, adjacent Al matrix around intermetallics dissolves also leaving behind Cu, which is originally in the Al matrix. Cu enrichment from both S phase dealloying and matrix dealloying is possible. Recently, the rotating ring-disk electrode experiments [85] indicate that Cu enrichment and redistribution from the S phase is dominant when AA2024-T3 is immersed in chloride solution for a short time, whereas, matrix dealloying contributes to Cu enrichment an redistribution more than S phase delloying after a long time immersion. Cu enrichment and redistribution also were found during CCC formation [166]. Cu smut forms on the surface when Al alloy samples are immersed in alkaline degreasing solution in the CCC pretreatment process, which interferes with the formation of CCCs [167]. Hagans and Haas using XPS depth profiling found the enriched Cu layer formed on the surface of polished 2024-T3 alloys remained intact during the CCC formation 46

63 [116]. Sun et al. [168] used SEM, TEM and XPS to study the effect of acidic pretreatment on corrosion protection of CCCs on AA2024-T3. It was suggested that Cu enrichment on the surface was deleterious to CCC protection. Hughes et al. [120] found that CCCs formed on 2024-T3 alloys were heterogeneous and Cu rich regions were present at the alloy/coating interface. It is obvious that the heterogeneous microstructure of substrate alloys plays a great role in CCC composition and structure. Recently, SIMS element mapping [147] revealed that the CCC was heterogeneous, and Cu enrichment was found in the outermost layer of CCC on AA2024-T3 (Figure 2.25). The CCC only partially covered the intermetallic particles. The authors suggested that the redistribution of intermetallic dissolution products in the CCC might be attributed to the projection of the intermetallic composition through the CCC toward the CCC surface. Regions of the CCC above intermetallics were depleted in Cr, but enriched in Cu. In Cu enrichment regions of CCC, ferricyanide and fluoride were depleted. Heterogeneous CCC where Cu rich regions are present may provide paths for selective dissolution and lead to the breakdown of CCC coatings when exposed to the corrosion environment. In summary, Cu enrichment and redistribution are detrimental to the corrosion resistance of Cu-containing Al alloys. Cu enrichment and redistribution during the pretreatment and subsequent CCC formation would decrease the effectiveness of CCC protection. However, the exact mechanism for the role of Cu enrichment and redistribution on corrosion behavior and CCC protection of Al alloys, particularly AA7xxx, is unclear. Therefore, further investigation is necessary. 47

64 2.5 Research Objective The issues mentioned in above literature review are too numerous to be covered by a single thesis. Many issues have been found to be directly or indirectly related to Al alloy composition, especially Cu content. Noble element Cu is distributed in the Al matrix and intermetallic particles in high strength Al alloys. Cu in the matrix and intermetallic particles plays an important role in corrosion behavior, CCC formation and protection in Cu-containing Al alloys. Cu enrichment and redistribution during the corrosion process and the pretreatment process decrease the corrosion resistance and CCC protection, respectively. In addition, the roles of Cu mainly have been studied on Al-Cu-Mg alloys (2xxx series) such as AA2024-T3, but little attention has been paid to the roles of Cu on Al-Zn-Mg-Cu alloy (7xxx series) behavior. It is well known that AA7xxx is quite different from AA2xxx in terms of alloy composition and microstructure, which result in different corrosion behavior and CCC protection. The motivation of this thesis is to develop a better understanding of the effect of Cu content in corrosion behavior and CCC protection of AA7xxx-T6. Listed below are some of objectives that will be addressed in this thesis: (1) Effect of Cu content in localized corrosion Although Cu addition in Al-Zn-Mg alloys can increase the mechanical strength, it has been generally accepted that Cu is detrimental to the corrosion resistance of AA7xxx. On the other hand, Cu addition increases the pitting or breakdown potential of Al-Cu binary alloys [54, 169] and also increases the breakdown potential of Cu rich intermetallic particles (Figure 2.26) [169]. This indicates that the exact role of Cu in pitting corrosion of AA7xxx might be complex. Furthermore, the role of Cu content in 48

65 IGC is still unclear since the microchemistry of grain boundary regions on the order of about 100 nm is extremely difficult to characterize. SEM, TEM, Scanning TEM (STEM), XPS, and Focused Ion Beam (FIB) sectioning was employed in this work along with the electrochemical techniques such as polarization and EIS. (2) Formation and breakdown of CCC on heterogeneous microstructure CCC formation and protection have been found to depend on the composition of Al alloys, especially Cu content. Salt spray tests on CCC coated AA6061 (0.15~0.40 wt%), AA7075 (1.2~2.0 wt%), and AA2024 (3.8~4.9 wt%) panels found that survival time or corrosion protection decreased with increasing Cu content of Al alloys [170]. CCC formation process and CCC protection ability are related to each other. Although several models for CCC nucleation and growth have been proposed, there are still unexplained phenomena, such as how electrons transport through the formed CCC and whether or how the sol-gel mechanism applies to CCC nucleation and growth. In order to develop better understanding of CCC formation and protection, the composition, structure, and cross-sectional morphology of CCCs formed on Al matrix and intermetallic particles in AA7xxx are critical issues, which was investigated in this work by using SEM, FIB, AFM, TEM, and STEM/Nano-EDS profiling. (3) Effect of Cu enrichment and redistribution on corrosion and CCC protection Cu enrichment and redistribution are believed to be responsible for poor corrosion resistance of Cu-containing Al alloys. Pretreatment process prior to CCC also can cause Cu enrichment and redistribution on the surface, which further complicate the CCC formation and protection [168]. However, the knowledge about the effect of Cu 49

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76 FIGURES (a) (b) (c) (d) (e) (f) Figure 2.1. Variations in cross-sectional shape of pits [2]. 60

77 (a) Adsorption mechanism (b) Anion penetration and ion migration model (c) Film breakdown theory Figure 2.2. Schematic diagrams for pit initiation models. (a) adsorption mechanism; (b) anion penetration and ion migration model; (c) film breakdown theory [1]. 61

78 Figure 2.3. Metastable pit transients observed on 302 stainless steel polarized at 420 mv SCE in 0.1 M NaCl solution [18]. Figure 2.4. Anodic and net current densities change as a function of potential for 100 nm Al film in 0.1 M NaCl solution [37]. 62

79 Figure 2.5. Concentration of Al 3+, Al(OH) 2+, and H + as a function of the product of the depth x and the current density i in a unidirectional pit [39]. 63

80 Log current density Figure 2.6. Schematic cyclic polarization showing E p and E R [1]. Figure 2.7. Effect of aging time on the corrosion behavior, at constant potential, of Al Cu in de-aerated 1 M NaCl solution at 25 C. Aging temperature 240 C. ο: passive; : IGC; : pitting + IGC; : pitting [51]. 64

81 a b c Figure 2.8. Variations of pitting potential as a function of alloying element content of (a) Al-Cu, (b) Al-Zn, and (c) Al-Mg binary alloys [51, 66]. 65

82 Figure 2.9. Pitting potentials for freshly deposited samples, E p and aged samples E p a, along with repassivation potentials, E R, for pure Al and AlNb alloys [36]. topcoat primer CCC Al alloy substrate Figure Schematic of standard coating system applied to Al alloys [169]. 66

83 Figure Polarization curves of AA2024-T3 in an aerated, oxygen-stirred 1 M NaCl base solution containing different amounts of persulfate and dichromate ions [96]. Figure Time series showing the effect of CrO 4-2 inhibitor on the anodic current spikes associated with metastable pitting polarized at -0.5 V SCE [21]. 67

84 Figure Schematic diagram illustrating duplex mechanism based model for the interation of chromate with aluminum surface [108]. Figure Proposed mediation mechanism by ferricyanide. Arrows represent redox cross reactions [122]. 68

85 Figure Schematic illustrating CCC formation by a sol-gel mechanism [133]. 20 um Figure Mud-crack morphology of CCC formed on AA2024-T3 [134]. 69

86 Figure Auger depth profiling showing that CCCs are mainly composed of Cr and O [134] Figure Schematic representing the structure of CCCs formed on AA2024 [118]. 70

87 Figure Cr K edge XANES reference spectra from solutions of 0.1 M CrO 4 2- (top) and Cr(H 2 O) 6 3+ (bottom), and mixed Cr(VI)/Cr(III) [134]. 71

88 Figure Schematic illustration of the bipolar model for interfaces for Al. MeO 4 2- is a metal oxyanion such as MeO 4 2- and CrO 4 2- [108]. Figure Schematic illustration of self-healing process of CCCs [170]. 72

89 Figure Cr VI concentration determined from the absorbance at 339 nm as a function of time after immersion of a CCC in nanopure water. Curvers are labeled with the ratio of the geometric CCC area to solution volume [155]. Figure Model for adsorption of Cr VI to solid Cr III hydroxide showing the Cr III -O- Cr VI covalent bonding [155]. 73

90 (a) Solution flow corrosion product - hydrous gel Cu Cu e - Cu H 2 O Cu(OH) 2 + 2H + Cu clusters dealloyed region secondary pitting Al 2 CuMg Al + 3H 2 O Al(OH) 3 + 3H + + 3e - α-al matrix Mg + 2H 2 O Mg(OH) 2 + 2H + + 2e - (b) Figure (a) The optical picture and (b) schematic representing Cu enrichment and redistribution during localized corrosion in AA2024-T3 [171]. 74

91 Figure Static SIMS chemical maps showing relative Cr and Cu concentration of a CCC on AA2024-T3 alloy at the outermost surface and at a level approximately 20% through the CCC [145]. 75

92 Figure Corrosion potential as a function of Cu content for Al-Cu solid solution and intermetallic compounds. Solid solution phase E corr values were determined in 1 M NaCl + 3% H 2 O 2. For the intermetallic E corr values, open data denotes measurements made in aerated 0.5 M NaCl; closed data denotes measurements made in 1 M NaCl + 3% H 2 O 2 [167]. 76

93 CHAPTER 3 EFFECT OF CU CONTENT ON CORROSION BEHAVIOR OF 7XXX SERIES AL ALLOYS 3.1 Introduction High strength Al alloys such as AA7075 and AA7178 commonly have been used in aircraft structure applications. The addition of Cu as one of alloying elements in Al- Zn-Mg alloys greatly improves the mechanical strength of 7xxx series alloys (AA7xxx) by precipitation hardening. The peak aged T6 temper provides the maximum mechanical strength. However, Cu-containing AA7xxx alloys in the T6 temper (AA7xxx-T6) are very susceptible to various forms of localized corrosion in chloride environments, such as pitting, crevice, intergranular corrosion (IGC), and exfoliation corrosion. As pits or crevices grow, they often develop IGC, which can propagate rapidly and deeply into a structure. In addition, Cu enrichment and redistribution, which arise from dealloying of Cu rich intermetallics or matrix dissolution, have been observed in Cu containing Al alloys in chloride solution and further accelerate the corrosion process [1-3]. The microstructure of AA7xxx-T6 is heterogeneous, containing the Al matrix, coarse intermetallic compound particles (IMCs), and the grain boundary region, including 77

94 the precipitate free zone (PFZ) and grain boundary precipitates. Cu is distributed in the microstructure of AA7xxx-T6 in different forms. In the Al matrix, fine hardening precipitates on the order of nm in size, such as G. P. zones or η phase, contain most of Cu in the alloy [4]. Cu also exists in the coarse IMCs of about a few microns in size, such as Al 2 CuMg (S), Al 2 Cu (θ), or Al 7 Cu 2 Fe particles, and in the grain boundary precipitates of about nm size, such as Mg(Zn Cu Al) 2 [4-7]. The Cu distribution in the microstructure affects the susceptibility to localized corrosion. Pitting corrosion usually occurs in the Al matrix near Cu- or Fe-containing intermtallic particles owing to galvanic interaction with the Al matrix [8]. IGC is generally believed to be associated with Cucontaining grain boundary precipitates and the PFZ along grain boundaries [4, 9]. The grain boundary regions in AA7050-T6 and AA7150-T6 have been investigated using analytical transmission electron microscopy [9-11]. The PFZ was found to be about 30~70 nm on either side of the boundaries and the grain boundary precipitates were about 50~100 nm in size. Zn and Cu depletion were found in the PZF in both alloys. On the grain boundaries in AA7150, MgZn 2 precipitates with considerable solubility of Cu and Al were identified [9]. The Cu content in the precipitates was about 12 at%. It is very difficult to quantitatively measure the exact composition of the PFZ and grain boundary precipitates due to experimental limitations, such as the relatively large electron probe size in analytical transmission electron microscopy, the narrow PFZ regions and the small size of the precipitates on grain boundaries. Therefore, more sophisticated analytical tools are needed to characterize the grain boundary region. Electrochemical measurements have been conducted on AA7xxx, especially AA7075 and AA7150 in the T6 temper in deaerated NaCl solution [9, 12]. Two 78

95 breakdown potentials have been found in potentiodynamic polarization scans. Maitra and English [12] suggested that for AA7075-T6, the active breakdown potential was associated with IGC and the noble breakdown potential with pitting in the matrix. They attributed the IGC susceptibility of the T6 temper to Mg and Zn solute segregation or enrichment in the grain boundary regions. It was suggested that this enrichment caused the grain boundary region to have a more active breakdown potential than the matrix. However, this inference was contradicted by Park and Ardell, who used analytical transmission electron microscopy (TEM) to show that the grain boundary region in AA7075-T6 exhibited considerable Zn depletion relative to the matrix [11]. Ramgopal et al. characterized the grain boundary region in AA7150-T6 using analytical TEM and thin film analogs of the PFZ and grain boundary precipitates [9, 13]. The OCP and breakdown potential for grain boundary Mg(Zn Cu Al) 2 precipitates in deaerated 0.5M NaCl increased with Cu content. Based on artificial crevice polarization experiments, it was proposed that IGC in the T6 temper was caused by anodic dissolution of Mg(Zn Cu Al) 2 grain boundary precipitates, which led to the creation of aggressive occluded environments. This anodic dissolution mechanism for IGC was also found in Al-Li-Cu alloys [14-16]. Most of investigations reported in the literature have focused on high Cu content Al alloys such as AA7075 and AA7150. The relationship between the microstructure, especially Cu content, and corrosion behavior is still unclear. To investigate the role of Cu content, the corrosion behavior of several AA7xxx-T6 alloys, with Cu content ranging from essentially Cu-free up to 2 wt%, in deaerated and aerated NaCl solution was studied in this work. Metallography and Focused Ion Beam (FIB) sectioning of the corroded 79

96 alloys were employed to determine the corrosion forms. The grain boundary regions were characterized using state-of-the-art scanning TEM (STEM), which provided a better understanding of the correlation between the microstructure and corrosion behavior of AA7xxx-T Experimental Five types of AA7xxx in the T6 temper with various Cu content were used in this work. The compositions of these alloys, as determined by inductively coupled plasma mass spectroscopy (ICP-MS), are listed in Table 3.1. The AA7004 alloy contains only wt% Cu and can be considered to be essentially Cu-free. AA7039, AA7075, and AA7050 were commercial grade sheets purchased in the T6 temper. AA7004 and AA7039 were not commercially available, so they were cast by vacuum induction melting, homogenized at 500 C for 6 h, and hot-rolled to sheets. The rolled sheets were solution heat treated at 475 C for 30 min and quenched in water, then artificially aged at 120 C for 24 h (T6 temper). Samples were cut from these five alloy sheets. A nonaqueous polishing procedure was used to minimize corrosion during polishing. The samples were mechanically ground down to 1200 grit finish with ethanol, cleaned ultrasonically in ethanol, and finally dried by a cold air stream. Microhardness measurements on five types of Al alloys were conducted using a Buehler Micromet II digital microhardness tester. The load to make a pyramid indentation was 300 g and the load time was 20 s. The microstructures of the Al alloys were characterized using optical metallography, Scanning Electron Microscopy (SEM), Transmission Electron 80

97 Microscopy (TEM), and Scanning TEM (STEM). Optical metallography was performed on samples with three different orientations (L, LT, and ST) that were polished to 1 µm finish with alumina suspended in ethanol followed by etching in Keller s reagent for about 30 to 60 s. SEM characterization was performed on corroded samples using an FEI Sirion SEM with field emission gun (FEG-SEM) operating at 12 kv. TEM characterization was performed using a Philips CM200T TEM operating at 200 kv. To avoid possible composition change of the alloys during electropolishing and ensure the accuracy of the composition measurement, TEM membranes of the alloys were prepared by an FEI Strata Dual Beam 235M FIB using a 30 kev Ga ion beam and 5 kev electron beam. The membranes had an area of about 15 µm 5 µm. The membranes were thinned in the FIB to a thickness of about 100 nm for electron transparency. The membranes were then plucked out of the bulk under an optical microscope using a sharp pyrex needle of about 1 µm diameter, and placed on a 200 mesh Au specimen grid with a formvar/carbon support film for TEM and STEM analysis. STEM characterization of the grain boundary precipitates and PFZ was conducted with an FEI Tecnai TF20 STEM operating at 200 kv. The drift-corrected EDS line profiles were acquired and quantified using the FEI/Emispec TIA software. The probe size was less than 2 nm with sufficient brightness supplied by a field emission gun (FEG) and the step size was about 5 nm. The dwell time for each pixel point was 5 s. The EDS quantification used the Cliff-Lorimer method with theoretical k-factors [17]. Electrochemical polarization measurements were performed on samples ground to 1200 grit in 0.5 M NaCl. The ph of the solution was adjusted to 3.56 by the addition of HCl. The solution was deaerated with Ar gas to decrease the corrosion potential and 81

98 allow for clear observation of the breakdown potentials. Potentiodynamic scans were performed at a rate of 0.2 mv/s. Potentiostatic polarization tests were conducted at different potentials relative to the breakdown potentials for 1 h, 5 h, and 24 h. Polarization experiments were performed using a Gamry PC4/FAS1 potentiostat. A Pt counter electrode and saturated calomel reference electrode (SCE) were used. The corrosion morphologies of the surfaces and cross-sections were examined metallographically after polarization tests to determine the nature of attack. To ensure the validity of the metallography, same samples were cross-sectioned after potentiostatic polarization tests with FIB. Rough sectioning was performed first with the ion beam perpendicular to the corroded surface at a current of 20 na. Final fine sectioning was performed at a current of 1 na. The sample was then tilted to 30 in order to conduct SEM imaging of the cross-section in the crater. To study the change in composition of the sample surface after the polarization tests, X-ray Photoelectron Spectrometry (XPS) measurements were performed using an AXIS Ultra spectrometer controlled by a VISION data system. The entrance and exit slit widths for the hemispherical analyzer were set to 4 mm and a pass energy of 20 ev was used. A monochromatic Al Kα X-ray line with energy of ev and 150 W was used as the incident radiation source. The take-off angle was 90. The binding energies of the elements of interest were calibrated with respect to the C 1s line at ev. Electrochemical Impedance Spectroscopy (EIS) measurements were performed at the open circuit potential during 168 h immersion of the samples in 0.5 M NaCl open to the air. The impedance measurements were performed as a function of frequency between 10 khz and 10 mhz using a sinusoidal voltage modulation of 10 mv. The 82

99 experiments were conducted using a Princeton Applied Research (PAR) Model 273 potentiostat with a Solartron Model 1255 frequency response analyzer. The open circuit potential during the immersion was also recorded. 3.3 Results Microstructure Figure 3.1 shows orthogonal metallographic microstructures for the five AA7xxx- T6 sheets. All alloys have an elongated grain structure even though the aspect ratios are different. SEM/EDS was used to identify the coarse intermetallic particles. The type and composition of the coarse intermetallic particles are listed in Table 3.2. Figure 3.2 shows the microhardness data for the five AA7xxx alloys. The hardness was found to increase with increasing Cu content, which supports the notion that the addition of Cu into Al-Zn- Mg alloys improves the mechanical strength of AA7xxx by precipitation hardening. TEM and STEM were used to characterize the microstructure of all alloys on a microscopic scale. Figure 3.3 shows TEM bright field micrographs of the grain boundary regions of all five alloys. Fine hardening precipitates are evident in the matrix. PFZ and grain boundary precipitates are found in each alloy. The composition of the matrix, PFZ, and grain boundary precipitates was determined by STEM/Nano-EDS. The composition of the matrix including the hardening precipitates, which mirror the composition of the precipitates on the grain boundary, was found to be very close to the nominal alloy composition for each alloy. The composition of the PFZ was found to be almost identical for all five alloys: Al containing about 2 wt% Mg. The measured composition of the grain boundary precipitates deviated considerably from the expected stoichiometry of Mg(Zn 83

100 Cu Al) 2. This deviation occurred because the small grain boundary precipitates did not extend through the whole thickness of TEM membranes (about 100 nm thick). As a result, the X-ray excitation volume included both the precipitates and surrounding Al rich PFZ, and the measured Al concentration in the precipitates was erroneously high. To correct for this artifact and determine the exact composition of the grain boundary precipitates, the approach of Ramgopal et al. [9] was followed. The composition of the PFZ was assumed to be Al-2%Mg and the PFZ was considered to be the only material in the X-ray excitation volume when the probe was positioned in the center of the PFZ. It was further assumed that the stoichiometry of grain boundary precipitates was Mg(Zn Cu Al) 2 since η phase has considerable solubility for Cu and Al [4]. The Mg concentration in this Cu- and Al-containing η phase was fixed at at% regardless of the Cu and Al concentration. Based on these assumptions, the measured composition of the precipitate, (C i ) mp, is a sum of contribution from the PFZ and the precipitates [9]: (C i ) mp = V pfz (C i ) pfz + V p (C i ) p (3.1) where C i is the composition of the i-th element (Mg, Zn, Cu, Al) in at%; (C i ) mp is the measured composition of the precipitate in at%; (C i ) pfz is the composition of the PFZ, which is considered be Al-2wt%Mg (2.2 at% Mg); (C i ) p is the calculated actual composition of the precipitate in at%; V pfz and V p are the fractions of the X-ray excitation volume for the PFZ and precipitate, respectively, during analysis of the precipitate. The volume fractions are determined from the Mg analysis and by fixing the Mg content in the precipitate at at%. Since the probe and X-ray excitation volume were much smaller than the precipitates, the accuracy for composition calculation mostly depends 84

101 upon the accuracy for EDS quantification. It should be noted here that the EDS quantification was made using the Cliff-Lorimer method with theoretical k-factors. Figure 3.4 shows the calculated composition of η phases in the five 7xxx alloys converted to wt%. The Mg concentration was fixed at at%, but the wt% was found to vary slightly as the composition of the particles changed for each alloy. Al concentration in the grain boundary precipitate also remained nearly at the same level for the five alloys. However, with increasing alloy Cu content, the Cu concentration in the precipitate dramatically increased from zero in AA7004 up to 18.6 wt% in AA7050, and the Zn concentration in the precipitate decreased from 65.8 wt% in AA7004 down to 41.0 wt% in AA Polarization Curves and Types of Corrosion Figure 3.5 shows the anodic polarization curves for AA7xxx-T6 alloys with various Cu contents in deaerated 0.5 M NaCl solution with ph Two breakdown potentials were found for all except the Cu-free AA7004. The data for the breakdown potentials are listed in Table 3.3. For the alloys exhibiting two breakdown potentials, a sharp increase in current density occurred at the lower, more active breakdown potential referred to as the first breakdown potential (E 1 ). As the potential increased further, the current density increased until reaching a value of about 1 ma/cm 2, and then decreased. Above the more noble breakdown potential, referred to as the second breakdown potential (E 2 ), the current increased again. This indicates that the dissolution between E 1 and E 2 is not stable, and stable dissolution occurred only above E 2. Two breakdown potentials for Cu-containing AA7xxx-T6 have been reported in the literature [9, 12]. 85

102 Figure 3.6 shows the relationship between the breakdown potentials and alloy Cu content on a semi-logarithmic scale. For the Cu-containing alloys, both breakdown potentials increased logarithmically with increasing Cu content. The difference between the two breakdown potentials for Cu-containing alloys was found to be nearly constant, mv, as shown in Table 3.2 and Figure 3.6. For Cu-free AA7004, only the second breakdown potential was observed, and it was associated with stable dissolution. Extrapolation of the line fitting the E 1 values to the Cu content of AA7004 indicates that an E 1 value expected for this alloy, if it exists, would be about mv SCE. To determine the types of attack corresponding to each breakdown potential, constant potential tests were conducted at different potentials relative to the breakdown potentials in the same deaerated NaCl solution. All Cu-containing alloys exhibited similar potentiostatic polarization behavior. Figure 3.7 shows the current density transients for AA7075-T6 at various applied potentials from just above the first breakdown potential to well above the second breakdown potential (E 1 = -780 mv SCE and E 2 = -720 mv SCE). The current densities were on the order of 1 ma/cm 2 within first tens of seconds during polarization then decreased dramatically. At applied potentials between E 1 and E 2, the current densities decreased to about A/cm 2 after 1 h, which was close to the value of passive current density in the ph 3.56 solution. Above E 2, the current densities at 1 h were much higher, and A/cm 2 for potentials of -700 mv SCE and -680 mv SCE, respectively. The charge densities for the Cu-containing alloys during the 1 h potentiostatic polarization were calculated by integration of the current density. Figure 3.8 shows the variation of charge density as a function of applied potential for all five alloys. Cu-free 86

103 AA7004 with only one breakdown potential behaved differently than the other four Cucontaining alloys with two breakdown potentials. For AA7004, the charge density increased dramatically above the single breakdown potential. For the Cu-containing alloys, the charge densities were very small (about 0.45 C/cm 2 ) for potentials between E 1 and E 2. These small amounts of charges were mostly generated by the large current transient occurring at the very beginning of polarization. However, when the applied potential was above E 2, the charge density increased dramatically with further increasing applied potential. Therefore, the first breakdown potential E 1 was mostly due to transient dissolution on the surface. In contrast, stable or sustained dissolution occurred above the second breakdown potential E 2. To further determine the corrosion forms corresponding to transient dissolution and sustained dissolution at E 1 and E 2, respectively, the corrosion morphology was examined after potentiodynamic and potentiostat polarization tests. Figure 3.9 shows the surface morphologies of all five alloys, which were potentiodynamically polarized to the current density of 1 ma/cm 2. The final potential for each scan was between the two breakdown potentials for each alloy except for AA7004, where the final potential was higher than the single breakdown potential. Again, the behavior of AA7004 was different than that of the other alloys. Many pits were found on the surface of AA7004 after this treatment. In contrast, large dark corroded regions without obvious pits and some bright uncorroded regions were found on the originally shiny surface of the Cu-containing alloys. This dark corrosion morphology is discussed in more detail below. Since the Cu-containing alloys have similar polarization behavior, AA7075-T6 with intermediate Cu content was studied in more detail in this work. An AA7075-T6 87

104 sample was cyclically polarized in deaerated 0.5 M NaCl three times, Figure The apex current densities for the three scans were 3, 10, and 10 ma/cm 2, respectively. When each cycle was finished, the sample was left at the open circuit until the OCP of the sample was stable. As shown in Figure 3.10, two breakdown potentials (E 1 = -780 mv SCE and E 2 = -720 mv SCE) were evident in the first cycle. In the second and third cycles, the first breakdown in the first scan disappeared and only the second breakdown was seen. The sample surface changes during exposure at potentials between E 1 and E 2 as a result of the transient attack. In contrast, the second breakdown potential is reproducible in subsequent cyclic polarization and corresponds to stable forms of dissolution such as pitting and IGC. To further identify the cause of transient dissolution between the two breakdown potentials, another AA7075 sample was polarized at 725 mv SCE, which is between the two breakdown potentials, for 500 s in deaerated 0.5 M NaCl, ph=3.56. Figure 3.11 shows a current transient similar to the curves shown in Figure 3.7. The current density was high at the beginning (about 10-3 A/cm 2 ) and then decayed to A/cm 2 after about 300 s. The total charge during 500 s for 1 cm 2 exposure area of the sample was calculated to be C. During the polarization, it was observed by the unaided eye that several dark spots appeared and spread out on the originally shiny sample surface until the whole surface became dark. After the polarization test, the sample was rinsed gently with deionized water and dried rapidly in a desiccator. Nearly the whole surface was dark under an optical microscope, similar to the corrosion morphology shown in Figure 3.9d. Figure 3.12 shows secondary and backscattered electron images of the sample. The dark region observed by the unaided eye and the optical microscope was a thin layer, which 88

105 seemed to undergo uniform corrosion and could be corrosion product uniformly covering the surface. This thin layer is rippled, a feature probably caused during rinsing with deionized water. Also shown in Figure 3.12 is a region at the middle of the SEM images that was an uncorroded residual surface region. Since this thin uniformly corroded layer on the surface is estimated to be much less than 1 µm based on the SEM images as shown in Figure 3.12, SEM/EDS is not suitable to measure its composition. X-ray Photoelectron Spectroscopy (XPS) was performed on this treated AA7075 sample and on an as-polished sample surface for composition comparison purposes. On the as-polished sample surface, the naturally formed Al oxide is a few nm thick [4], which is less than the analysis depth for photoelectrons with a take-off angle of 90 [18]. Hence, the photoelectrons in the XPS measurements originated from both the Al oxide and underlying Al matrix in the aspolished sample. In the case of the polarized sample, the photoelectrons mainly were produced from the thin corroded layer or corrosion product. Figure 3.13 shows XPS spectra for Al, Zn, Mg, Cu, O, and Cl. The C 1s spectra for the as-polished and polarized samples were also obtained for energy calibration. The detailed data obtained from the XPS spectra are listed in Table 3.4. The oxidation states were determined using the XPS database [18]. It should be noted that the binding energy for the elements in the Al matrix deviated from the binding energy for the standard pure metals. The comparison of XPS spectra for the as-polished and polarized samples in Figure 3.13 provides important clues regarding the surface change after the polarization between the two breakdown potentials. Al 0 in the matrix is visible in the Al 2p spectrum for the as-polished sample, but not for the polarized sample. The Al and O spectra indicate the formation of Al(OH) 3 on the 89

106 polarized sample surface. In addition, Mg disappeared and the amount of Zn was reduced, while Cu was enriched on the polarized sample surface. Chloride was also found on the polarized sample surface, as expected. The change in chemical states and amounts of the elements indicates that preferential dissolution of Mg, Al, and Zn occurred during the polarization between the two breakdown potentials, in a fashion very similar to dealloying of Cu-containing intermetallic particles like Al 2 CuMg. A cross-sectional foil of the surface of this treated AA7075 sample was prepared by FIB and examined by TEM. Figure 3.14 shows the product layer and underlying Al substrate in cross-sectional view. The amorphous Al(OH) 3 layer was about 120 nm thick. The roughness of the interface between the layer and Al matrix suggests that the product layer was produced by uniform attack of the Al matrix. The potentiodynamic polarization experiments described above were relatively short in duration. The corrosion morphology was also examined after long-term potentiostatic polarization tests. AA7039-T6 and AA7050-T6 samples representing low and high Cu content, respectively, were polarized at potentials between their two breakdown potentials in deaerated 0.5 M NaCl, ph=3.56 for 24 h. The potentials applied to AA7039 and AA7050 were -870 mv SCE and 720 mv SCE, respectively. Similar to the curves in Figure 3.7, one large current transient was observed at the beginning of these two long-term polarization tests, followed by a low current density on the order of 10-5 A/cm 2. Figure 3.15 shows the corrosion morphologies of AA7039 and AA7050 after the 24 h polarization tests. Small pits and semi-continuous shallow attack were found on the grain boundaries in the low Cu content AA7039. The high Cu AA7050 exhibited continuous and relatively deep IGC attack as well as corrosion product on the attacked 90

107 grain boundaries. These morphologies indicate that IGC attack occurred in both alloys during 24 h exposure between E 1 and E 2. An AA7075-T6 sample was polarized at -680 mv SCE in deaerated 0.5 M NaCl, ph=3.56 for 5 h. This potential is above the second breakdown potential of AA7075 (E 2 = -720 mv SCE). Figure 3.16 shows the corrosion morphology of AA7075 in top view and metallographic cross-section view. The corrosion morphology in AA7075 reveals a combination of IGC and selective grain attack, sometimes referred to as microstructural pitting [19]. To ensure that the metallographic corrosion morphology is not an artifact due to grain fallout during mechanical polishing, the same AA7075 sample after polarization was cross-sectioned using FIB, which excludes this possibility. It is evident in Figure 3.17 that the FIB section also reveals a combination of IGC and selective grain attack. This mode of attack involves complete dissolution of certain grains. It is interesting that the attack ceases at the neighboring grain, instead of continuing to form a typical hemispherical isotropic pit. Intergranular attack can be seen at the neighboring grain face and in the underlying region in the FIB-etched section EIS Measurement under Free Corrosion Conditions All of the above electrochemical polarization measurements were performed in deaerated NaCl solution and were controlled by a potentiostat, which can supply as much current as is needed to support anodic dissolution of the alloys. In reality, corrosion occurs under free corrosion conditions, where local cathodes on the surface supply cathodic current to support anodic reactions at the local anodes. The primary cathodic reaction to consider in aerated neutral solutions is oxygen reduction. To investigate this 91

108 type of exposure, all five alloys were immersed in stagnant, air exposed 0.5 M NaCl solution for up to 168 h. Figure 3.18 shows the open circuit potential (OCP) as a function of immersion time. The OCP increased with Cu content in the alloys. The OCPs for Cufree AA7004 and AA7039, which has only wt% Cu, decreased after 48 h immersion, eventually reaching low values. In contrast, the OCP for the intermediate and high Cu-containing alloys remained constant, at values about equal to their respective E 2 potentials. EIS measurements were carried out at OCP during this immersion time. Figure 3.19 shows the variation of low frequency impedance or polarization resistance for all five alloys as a function of immersion time. The polarization resistance of high Cucontaining alloys was lower than that for low Cu-containing alloys under free corrosion conditions. It is known that Cu enrichment and redistribution occur in Al-Cu or Al-Cu- Mg alloys as a result of dealloying of the Cu-containing matrix and active Cu-containing intermetallic inclusions such as Al 2 CuMg particles and release of Cu particles, corrosion, and redeposition [1-3, 20-24]. Cu enrichment and redistribution in Al-Zn-Mg-Cu alloys during free corrosion could account for the variation of OCP and polarization resistance as a function of immersion time. For high-cu alloys such as AA7075 and AA7050, Cu remnants or clusters develop with time. These Cu-rich areas on the surface are cathodic sites that facilitate oxygen reduction, which drives the corrosion process. For these alloys, the OCP remains high and the corrosion rate (inversely related to the polarization resistance) increases with time. For AA7004 and AA7039, Cu enrichment and redistribution are negligible or less significant. Without Cu enrichment, there is insufficient cathodic reaction to sustain localized corrosion. As a result, the OCP drops, and the corrosion rate remains low and decreases with time in Al-Zn-Mg alloys. 92

109 3.4 Discussion Mechanism of Localized Corrosion of AA7xxx-T6 To develop a better understanding of the role of Cu content in corrosion behavior, it is necessary to understand the mechanism of localized corrosion of AA7xxx-T6. It is generally recognized that there should be a relationship between corrosion behavior and alloy composition and microstructure. By studying corrosion behavior and microstructure of a series of AA7xxx with various Cu content, it is possible to establish this relationship. One of key observations from the polarization curves described above is that all alloys have almost the same difference between their two breakdown potentials, about mv. This implies that the two types of dissolution or corrosion corresponding to the two breakdown potentials might be related each other, although the first breakdown potential corresponds to transient dissolution, while the stable dissolution occurs above the second breakdown potential. Potentiostatic polarization between the two breakdown potentials and potentiodynamic polarization to 1 ma/cm 2 with the final potential between the two breakdown potentials both led to the creation of dark corroded regions on the surface. The morphology of the dark regions observed using SEM is much like Al hydroxide corrosion product. Since the applied potentials were still below the second breakdown potentials, no large and deep pits were observed on the surface. This excludes the possibility that the corrosion product is produced from large discrete pits and deposited on the un-attacked surface. XPS surface analysis proves that the dark regions are those attacked uniformly and the thin layers in the dark regions are mainly Al(OH) 3 with Cu enriched and Mg depleted. TEM observation on the FIB prepared cross-section of the 93

110 layer and underlying Al matrix further proves that the Al matrix surface was subject to corrosion attack and the amorphous hydroxide layer was produced by direct attack of the Al matrix. This uniform corrosion on the surface could be attributed to dealloying of fine hardening precipitates containing Cu. Electrochemical studies revealed that Cucontaining intermetallic particles such as Al 2 CuMg and Mg(Zn Al Cu) 2 are susceptible to dealloying in chloride solution [13, 25]. These Cu-containing intermetallics are found to have their own breakdown potentials for dealloying [13, 25]. After the polarization of AA7075 at 725 mv SCE for 500 s, the charge for 1 cm 2 exposure area was C. Assuming a uniform dissolution process that was 100% current efficient, simplifying the Al matrix (91% Al) as 100% Al, and applying Faraday s law, the depth of Al matrix dissolved was 75 nm. This calculated value is on the same order of magnitude as the 120 nm thick uniform Al(OH) 3 layer observed by TEM. This further supports that the direct uniform attack of the Al matrix including the hardening particles corresponds to the first breakdown potential. The TEM micrographs in Figure 3.3 show the AA7xxx-T6 microstructure, where fine hardening precipitates of less than 1 nm in size are in the matrix, and relatively coarse precipitates of about 50 nm and PFZ are on the grain boundary. The area or volume fraction of fine hardening precipitates is much greater than that of grain boundary precipitates. It is impossible to directly determine the composition of the hardening particles because of their very small size. According to physical metallurgy of Al alloys, fine hardening precipitates and grain boundary precipitates should have similar composition in the T6 temper. In this work, the compositions of grain boundary 94

111 precipitates were determined using STEM/Nano-EDS. Figure 3.4 reveals that Cu content in grain boundary Mg(Zn Al Cu) 2 phase precipitates dramatically increased and Zn decreased with alloy Cu content. This variation of composition is probably also true for the fine hardening precipitates (GP zones and η ), which have similar composition to the grain boundary precipitates. Ramgopal and Frankel [13] deposited thin film analogs to Mg(Zn Cu Al) 2 precipitates and investigated electrochemical behavior of thin film analogs in deaerated 0.5 M NaCl. They found that breakdown potential of Mg(Zn Cu Al) 2 analogs increased with Cu content [13]. The breakdown of these thin film analogs appeared to be a localized dealloying phenomenon. The breakdown potentials they measured were lower than the first breakdown potentials of the alloys studied in this investigation. Also, they found that a considerable amount of Cu (17 at%) was needed to increase the breakdown potential beyond that of MgZn 2. Despite these differences, the trend of increasing breakdown potential of Mg(Zn Cu Al) 2 with increasing Cu content is useful for explaining the behavior of the alloys studied in this investigation. The Cu content of the Mg(Zn Cu Al) 2 grain boundary precipitates and fine hardening precipitates in the matrix increases with increasing alloy Cu content. The breakdown potential associated with dealloying for Mg(Zn Cu Al) 2 analogs was shown to increase with Cu content. These observations can be combined to explain the increasing trend of E 1 with alloy Cu content by considering that the first breakdown potential in Cucontaining AA7xxx-T6 is the result of preferential dissolution or dealloying of fine Mg(Zn Cu Al) 2 hardening precipitates in the matrix. Since these fine hardening precipitates are highly dispersed coherently or semicoherently in the Al solid solution containing Zn, Mg, and Cu, the preferential dissolution of these fine precipitates can also 95

112 cause the reaction of the Al solid solution in the matrix. This combined dissolution of fine hardening precipitates and surrounding Al solid solution results in the formation of an Al(OH) 3 product layer, which limits the depth of attack, making it a transient process. Attack of the grain boundary particles is also possible, but IGC does not develop until longer hold times or higher potentials for kinetic reasons, as discussed below. When the applied potential is above the second breakdown potential (E 2 ), stable pits form on the matrix by breakdown of the passive film or Al(OH) 3 formed between E 1 and E 2. The breakdown potential for stable pitting or dissolution is dependent upon the composition of 7xxx alloys, which are composed of mainly Al and a small amount of alloying elements including 4~6 wt% Zn, 1.6~3 wt% Mg, and 0~2 wt% Cu. Muller and Galvele have systematically investigated the pitting potential of high purity binary Al-Zn, Al-Mg, and Al-Cu solid solutions [26, 27]. Zn, Mg, and Cu as alloying elements were found to have different effects on the pitting potential of Al alloys. The pitting potential decreased greatly with increasing Zn content up to 3 wt% and then remained constant with further increase in Zn content. There was no influence of Mg on pitting potential. Pitting potential increased dramatically with increasing Cu content up to 5 wt%. In this study, the STEM/Nano-EDS composition measurements reveal that the composition of the Al matrix including the fine hardening particles is very close to the alloy composition. Since all five alloys contain 4~6 wt% Zn, the effect of Zn (> 3 wt%) on the breakdown potential of the matrix should be about the same. On the other hand, increasing Cu content in Al matrix from 0 to 2 wt% dramatically increases the breakdown potential of the matrix, E 2. Above E 2, selected grain attack or matrix dissolution is evident in the corrosion morphology as well as IGC. The pits are much larger than the size of the fine 96

113 hardening particles, so it is reasonable to consider the matrix with hardening particles to be a homogeneous phase regarding pit stability. Therefore, the increase in the second breakdown potential corresponding to stable dissolution in AA7xxx-T6 with increasing alloy Cu content can be explained by the increase in Cu content in the Al. Another interesting phenomenon is that shallow IGC formed and propagated on the surface during long-term polarization between E 1 and E 2 (Figure 3.15), or deeply into the alloys above E 2 along with selective grain attack (Figure 3.16 and 3.17). The moresevere attack observed for the higher Cu alloy at a controlled potential between E 1 and E 2 in deaerated solution is in contradiction to the observation of increasing breakdown potential with increasing Cu content. Furthermore, it is not consistent with the viewpoint of Ramgopal et al., who indicated that higher dissolved Cu concentration in intergranular crevices in T7 temper alloys decreases the severity of attack relative to the T6 temper [9]. It is worthwhile mentioning the effect of temper on the polarization behavior. Only one breakdown potential was found for AA7075-T7 [12] or AA7150-T7 [9] in deaerated chloride solution. The single breakdown potential for either of the alloys in the T7 temper was found to be equal to the first breakdown potential for the T6 temper. It is possible that the T7 temper decreases the critical current density for stable pitting, and thus decreases the breakdown potential for stable dissolution or pitting Role of Cu content in Corrosion Resistance of AA7xxx-T6 Electrochemical studies in this work reveal that corrosion behavior of AA7xxx-T6 is strongly dependent on Cu content. The breakdown potentials measured in deaerated chloride solution increase with alloy Cu content. It should be noted here that the second 97

114 breakdown potentials are more important than the first breakdown potentials since the second breakdown potential corresponds to the stable dissolution or localized corrosion, while the first breakdown potential corresponds to transient dissolution. From this point of view, increasing Cu content is beneficial to the localized corrosion resistance of AA7xxx because of its effect on increasing the breakdown potentials. It is interesting to consider that the increase in breakdown potential with Cu content might be related to the beneficial role of Cu in increasing resistance of AA7xxx to stress corrosion cracking [4]. However, the corrosion resistance of the alloys at open circuit in aerated solution decreases with increasing alloy Cu content. These contradictory effects of Cu content must be rationalized. One important consideration is the role of Cu on the OCP. Figure 3.18 presents the evolution of the OCP in aerated solutions over a period of 1 week. The OCP evolution in the deaerated solution during a shorter period, the stabilization period prior to potentiodynamic polarization tests, is given in Figure The relationship between the breakdown potentials and the OCPs in deaerated and aerated chloride solutions is given in Figure This figure combines the data of Figures 3.6, 3.18, and The OCPs in deaerated solution after about 30 min prior to the polarization tests are denoted by triangles. The range of OCP in deaerated solution during this period is given by the solid double-ended arrows. The OCPs in aerated solution after 168 h are denoted by crosses and are connected by a line. The range of OCP in aerated solution during this period is given by the dashed double-ended arrows. The variations in potential for AA7075 and AA7050 in aerated solution were so small that they were not included. In deaerated solution, the OCPs for all alloys except Cu-free AA7004 were below their 98

115 respective first breakdown potentials. The OCP for AA7004 in deaerated solution was mv SCE within about the first 250 s. This is just below its second breakdown potential (-951 mv SCE). Then the OCP decreased with the immersion time until a stable value, mv SCE, was reached. The exposure of AA7004 during the open circuit stabilization period at potentials above E 1 likely resulted in dissolution of the Cu-free η precipitates similar to the behavior of the Cu-containing alloys during potentiostatic exposure to potentials between their E 1 and E 2 values. As a result, the transient dissolution associated with E 1 was not observed on the subsequent polarization scan and only stable localized corrosion associated with E 2 was observed. This is similar to the behavior of AA7075-T6 during repeated cyclic polarization shown in Figure Experiments were performed in an attempt to evaluate this possibility by avoiding open circuit exposure at higher potentials. An AA7004 sample was immersed into deaerated 0.5 M NaCl at ph=3.56 under potential control at V SCE. To prevent possible overload of the potentiostat during immersion, an AA7039 sample was polarized at 1.15 V SCE prior to immersion of AA7004. The AA7004 sample was electrically shorted to the AA7039 sample, immersed in the solution, and then the AA7039 sample was disconnected electrically. The subsequent experiment was performed only on the AA7004 sample. The Gamry Virtual Front Panel (VFP) software was used in this experiment to avoid any open circuit exposure. After 2 min polarization at 1.15 V SCE, the potential was scanned upwards at a rate of 0.2 mv/s. In this fashion, the potentiodynamic polarization curve for AA7004 was obtained as shown in Figure 3.22b. Only one breakdown potential was found again and its value was the same as found previously. This sample was not exposed at a higher open circuit potential, and the 99

116 absence of a breakdown at the expected E 1 value suggests that AA7004 does not exhibit the phenomenon associated with the first breakdown potential. Another AA7004 sample was immersed for 2 min under potential control at V SCE in the same fashion, and then jumped to 930 mv SCE and held there for 1 h, Figure Figure 3.23 also shows the result for a sample held at 930 mv SCE after OCP exposure for 30 min in deaerated chloride solution, as was the procedure for all experiments shown in Figures 3.7 and 3.8. A large current transient was found during the first 400 s for the sample that was not first exposed at open circuit. After the transient, the current density was nearly same as the value for the AA7004 sample that was first exposed at OCP for 30 min. This result suggests that Cu-free AA7004 does in fact exhibit the transient dissolution phenomenon associated with E 1. Further work is required to resolve the contradictory data regarding AA7004. It is possible that the transient dissolution phenomenon occurs, but is difficult to observe at low potentials in AA7004. It is also possible that the transient dissolution phenomenon does not occur because some Cu content is needed. Dealloying of Mg and Zn from the η phase precipitates might occur at very low potentials and therefore be very difficult to sense electrochemically [13]. In a Cu-containing alloy, the η phase precipitates would contain some amount of Cu and dealloying would leave a Cu-rich remnant. In this view, E 1 might be associated with the breakdown of this Cu-rich remnant layer. In a Cu-free alloy, this layer would not be present, and the Cu-rich remnant η phase precipitates would just suffer the very low potential attack. At higher potentials, all alloys exhibit the matrix breakdown and IGC associated with E

117 The contradictory behavior in deaerated and aerated solutions can be understood by considering the OCP in aerated solution. As shown in Figures 3.19 and 3.21, the OCPs during 1 week in aerated solution are quite different for the different alloys. For Cu-free AA7004, the OCP started out just above its E 1, but then dropped down to very low values after several days. The OCP for AA7039, with extremely low Cu content, reached a value approximately equal to its E 2 after 2 days and then also decreased with immersion time to very low values. From the comparison of OCP and breakdown potentials for these alloys, it is expected that the transient dissolution associated with the fine hardening particles would occur during the early stages of exposure. Some IGC or selective grain attack could have taken place in the case of AA7039 during the period when its OCP of was close to its E 2. However, the long-term behavior indicates that sustained localized attack was not possible. This is supported by the high polarization resistances measured during the exposure to the aerated solution. In sharp contrast, the OCPs for other three alloys with intermediate and high Cu content remained at high values, which are a little higher than their own second breakdown potentials. As mentioned above, this is likely the result of Cu enrichment at the surface, which enhances the oxygen reduction reaction in aerated solutions. The potential is pinned close to the breakdown potential associated with stable attack (E 2 ) because of the relative non-polarizability of the localized corrosion reactions. Sustained localized corrosion is possible at these potentials, as is reflected by the decrease in polarization resistance with increasing Cu content. Hence, the role of Cu content is detrimental to localized corrosion resistance of AA7xxx in aerated solutions. 101

118 3.5 Summary In this study, the corrosion behavior of 7xxx alloys with various Cu content was investigated using a combination of SEM, TEM, STEM, and XPS analysis of the microstructure. The following conclusions are drawn: 1. Two breakdown potentials were observed for all alloys except for essentially Cufree AA7004. The two breakdown potentials increase logarithmically with alloy Cu content in deaerated chloride solution. 2. The first breakdown potential corresponds to transient dissolution associated with attack of the fine hardening particles and the surrounding solid solution in a thin surface layer. The Cu content of these particles likely mirrors that of the grain boundary particles, which increase in Cu content as the alloy Cu content increases. The Cu content in the hardening particles controls the first breakdown potential. 3. The second breakdown potential is associated with combined IGC and selective grain attack, and is controlled by the Cu content in the matrix, including the hardening particles. 4. IGC develops after long times at potentials between the two breakdown potentials. 5. Under free corrosion conditions in aerated chloride solutions, the corrosion potential increases and the polarization resistance decreases as the Cu content increases as a result of Cu enrichment on the surface, which facilitates the oxygen reduction reaction. Because of this effect, the overall influence of Cu on the 102

119 corrosion behavior is detrimental, despite the increase in breakdown potentials with Cu content. REFERENCES 1. R. G. Buchheit, M. A. Martinez, and L. P. Montes, J. Electrochem. Soc., 147, 119 (2000). 2. R. G. Buchheit and R. K. Boger, in Proceedings of the Research Topical Symposium on Localized Corrosion, Corrosion 2001, NACE, Houston, TX (2001). 3. N. Dimitrov, J. A. Mann, M. Vukirovic, and K. Sieradzki, J. Electrochem. Soc., 147, 3283 (2000). 4. J. E. Hatch, Aluminum: Properties and Physical Metallurgy, ASM, Metals Park, OH (1983). 5. R. P. Wei, C.-M. Liao, and M. Gao, Met. Mater. Trans. A, 29A, 1153 (1998). 6. M. Gao, C. R. Feng, and R. P. Wei, Met. Mater. Trans. A, 29A, 1145 (1998). 7. M. Puiggali, A. Zielinski, J. M. Olive, E. Renauld, D. Desjardins, and M. Cid, Corrosion Science, 40, 805 (1998). 8. E. Lunarska, E. Trela, and Z. Szklarska-Smialowska, Corrosion, 43, 219 (1987). 9. T. Ramgopal, P. I. Gouma, and G. S. Frankel, Corrosion, 58, 687 (2002). 10. J. K. Park and A. J. Ardell, Metallurgical Transaction, 15A, 1531 (1984). 11. J. K. Park and A. J. Ardell, Acta Metall. Mater., 39, 591 (1991). 12. S. Maitra and G. C. English, Metallurgical Transaction A, 12A, 535 (1981). 13. T. Ramgopal, P. Schmutz, and G. S. Frankel, J. Electrochem. Soc., 148, B348 (2001). 14. R. G. Buchheit, J. P. Moran, and G. E. Stoner, Corrosion, 46, 610 (1990). 15. R. G. Buchheit, J. P. Moran, and G. E. Stoner, Corrosion, 50, 120 (1994). 103

120 16. R. G. Buchheit, F. D. Wall, G. E. Stoner, and J. P. Moran, Corrosion, 51, 417 (1995). 17. D. B. Williams and C. B. Carter, Transmission Electron Microscopy: a Textbook for Materials Science, Plenum Press, New York (1996). 18. J. F. Moulder, W. F. Stickle, P. E. Sobol, and K. D. Bomben, Handbook of X-ray Photoelectron Spectroscopy, ed. J. Chastain, Perkin-Elmer Corporation, Eden Prairie, Minnesota (1992). 19. Annual Book of ASTM Standards, ASTM, Philodelphia, PA (1995). 20. R. G. Buchheit, R. P. Grant, P. F. Hlava, B. Mckenzie, and G. L. Zender, J. Electrochem. Soc., 144, 2621 (1997). 21. R. G. Buchheit, R. K. Boger, M. C. Carroll, R. M. Leard, C. Paglia, and J. L. Searles, JOM, 53, 29 (2001). 22. N. Dimtrov, J. A. Mann, and K. Sieradzki, J. Electrochem. Soc., 146, 98 (1999). 23. M. B. Vukmirovic, N. Dimitrov, and K. Sieradzki, J. Electrochem. Soc., 149, B428 (2002). 24. H. M. Obispo, L. E. Murr, R. M. Arrowood, and E. A. Trillo, J. Mater. Sci., 35, 3479 (2000). 25. R. G. Buchheit, L. P. Montes, M. A. Martinez, J. Micheal, and P. F. Hlava, J. Electrochem. Soc., 146, 4424 (1999). 26. I. L. Muller and J. R. Galvele, Corrosion Science, 17, 995 (1977). 27. I. L. Muller and J. R. Galvele, Corrosion Science, 17, 179 (1977). 104

121 TABLES AND FIGURES Alloy Zn Mg Cu Cr Fe Si Mn Al bal bal bal bal bal Table 3.1. Composition of 7xxx Al alloys measured by ICP-MS in wt%. alloy Coarse intermetallics 7004 Al 3 Fe 7039 Al 3 (Fe Si) 7029 Al 3 (Fe Cu Si) 7075 Al 2 CuMg, Al 3 (Fe Cu), Mg 2 Si 7050 Al 7 Cu 2 Fe Table 3.2. Coarse intermetallic particles identified by SEM/EDS. alloy E 1 (mv SCE) E 2 (mv SCE) E 2 -E 1 (mv) 7004 N/A -951 ± 3 N/A ± ± ± ± ± ± ± ± 1 52 Table 3.3. Breakdown potentials for AA7xxx-T6 in deaerated 0.5 M NaCl, ph=

122 Element Compound State BE (ev) FWHM (ev) Sample Al 2p Al in matrix Al polished Al 2p γ-al 2 O 3 /Al(OH) 3 Al polished Al 2p Al(OH) 3 Al polarized Zn 2p3/2 Zn in matrix Zn polished Zn 2p3/2 Zn in Al(OH) 3 Zn polarized Mg 2p Mg in matrix Mg polished Mg 2p polarized Cu 2p3/2 Cu in matrix Cu polished Cu 2p3/2 Cu in Al(OH) 3 Cu polarized O 1s γ-al 2 O 3 /Al(OH) 3 O 2- /OH polished O 1s Al(OH) 3 OH polarized Cl 2p3/2 Cloride Cl polarized Table 3.4. Data obtained from XPS spectra. BE is binding energy, and FWHM is full width at half maximum height. 106

123 107 Figure 3.1. Microstructure of three orthogonal sections of (a) AA7004-T6, (b) AA7039-T6, (c) AA7029- T6, (d) AA7075-T6, and AA7050-T6 sheets

124 200 Vickers Hardness (HV) Cu Content (wt%) Figure 3.2. Microhardness versus Cu content curve showing that Cu addition increases the hardness of AA7xxx-T6. 108

125 A B 200 nm 200 nm C D 200 nm 100 nm E 100 nm Figure 3.3. TEM micrographs showing the grain boundary regions in (a) 7004, (b) 7039, (c) 7029, (d) 7075, and (e) 7050 alloys. 109

126 80 Composition of Eta Phase (wt%) Zn Mg Al Cu Alloy Cu Content (wt%) Figure 3.4. Composition of η phase precipitates on grain boundary as a function of alloy Cu content in AA7xxx-T6. 110

127 Potential (V SCE) E 1 E Current Density (A/cm 2 ) Figure 3.5. Potentiodynamic polarization curves for AA7xxx-T6 in deaerated 0.5 M NaCl, ph=3.56 with a scan rate of 0.2 mv/s. 111

128 -0.6 Potential (V SCE) E 2 E Cu Content (wt%) Figure 3.6. Correlation between the breakdown potentials in deaerated 0.5 M NaCl, ph=3.56 and the alloy Cu content. 112

129 7075-T6 E 2 = -720 mv E 1 = -780 mv Current Density (A/cm 2 ) mv -700 mv -720 mv -722 mv -730 mv -740 mv -760 mv Time (Sec) Figure 3.7. Current transients at various applied potentials for AA7075-T6 in deaerated 0.5 M NaCl, ph=

130 3.6 Charge Density (Coulomb/cm 2 ) Potential (V SCE) Figure 3.8. Variation of charge density as a function of applied potential in 1 h potentiostatic polarization for AA7xxx-T6 in deaerated 0.5 M NaCl, ph=3.56. The arrows show the values of E 2 for each alloy. 114

131 115 Figure 3.9. Optical micrographs of (a) 7004, (b) 7039, (c) 7029, (d) 7075, and (e) 7050 samples potentiodynamically polarized to the current density of 1 ma/cm 2 in deaerated 0.5 M NaCl, ph=3.56.

132 Scan 1 Scan 2 Scan 3 Potential (V SCE) Current Density (A/cm 2 ) Figure Cyclic anodic polarization curves for the same AA7075-T6 sample in deaerated 0.5 M NaCl, ph=3.56 at a scan rate of 0.2 mv/s. The sample was polarized three times: the apex current densities for three scans were 3, 10, and 10 ma/cm 2, respectively. When each cycle was finished, the sample was left at open circuit until the OCP was stable. 116

133 10-2 Current Density (A/cm 2 ) Time (sec) Figure Current transient of AA7075 polarized at 725 mv SCE for 500 s in deaerated 0.5 M NaCl, ph=

134 (a) Secondary electron image (b) Backscattered electron image Figure SEM (a) secondary electron and (b) backscattered electron images of the surface of AA7075 sample polarized at 725 mv SCE for 500 s in deaerated 0.5 M NaCl, ph=

135 5000 counts/s Al 2p Intensity Polarized as-polished Binding Energy (ev) (a) Al 2p 2000 counts/s Zn 2p 1/2 Zn 2p 3/2 Intensity polarized as-polished Binding Energy (ev) (b) Zn 2p Figure XPS spectra measured from samples of AA7075 after mechanical polishing and a subsequent polarization at 725 mv SCE for 500 s in deaerated 0.5 M NaCl, ph=3.56. (a) Al 2p, (b) Zn 2p, (c) Mg 2p, (d) Cu 2p, (e) O 1s, and Cl 2p. 119 (Continued)

136 Figure 3.13: (Continued) Intensity 250 counts/s Mg 2p polarized as-polished Binding Energy (ev) (c) Mg 2p 10 3 counts/s Cu 2p 3/2 Cu 2p 1/2 Intensity polarized as-polished Binding Energy (ev) (d) Cu 2p (Continued) 120

137 Figure 3.13: (Continued) O 1s counts/s Intensity polarized as-polished Binding Energy (ev) 525 (e) O 1s 500 counts/s Cl 2p 3/2 Intensity polarized Binding Energy (ev) 190 (f) Cl 2p Figure XPS spectra measured from samples of AA7075 after mechanical polishing and a subsequent polarization at 725 mv SCE for 500 s in deaerated 0.5 M NaCl, ph=3.56. (a) Al 2p, (b) Zn 2p, (c) Mg 2p, (d) Cu 2p, (e) O 1s, and Cl 2p. 121

138 Pt overlayer product layer matrix Figure TEM micrograph of the product layer formed on AA7075-T6 when polarized at 725 mv SCE for 500 s in deaerated 0.5 M NaCl, ph=

139 (a) (b) Figure SEM micrographs of the surface of (a) AA7039-T6 polarized at 870 mv SCE, and (b) AA7050-T6 polarized at 720 mv SCE in deaerated 0.5 M NaCl, ph=3.56 for 24 h. 123

140 40 µ m (a) (b) (c) Figure Metallographical micrograph of (a) top surface, (b) as polished crosssection, and etched cross-section of AA7075-T6 polarized in deaerated 0.5 M NaCl, ph=3.56 at 680 mv SCE for 5 h. 124

141 10 µm FIB cross section Figure SEM micrograph of FIB cross-section of the same AA7075 sample as in Figure 3.16 polarized at 680 mv SCE in deaerated 0.5 M NaCl, ph=3.56 for 5 h, evidencing microstructural pitting. 125

142 -0.6 Open Circuit Potential (V SCE) Immersion Time (h) Figure Time evolution of the open circuit potentials of AA7xxx-T6 within 168 h immersion in aerated 0.5 M NaCl. 126

143 Polarization Resistance (ohm cm 2 ) Immersion Time (h) Figure Polarization resistance determined by EIS tests as a function of immersion time for AA7xxx-T6 in aerated 0.5 M NaCl. 127

144 Open Circuit Potential (V SCE) deaerated 0.5 M NaCl, ph Time (sec) Figure Open circuit potential versus time for 7xxx-T6 in deaerated 0.5 M NaCl, ph=

145 OCP in deaerated NaCl OCP in aerated NaCl after 168h E 2 E 1 Potential (V SCE) Cu Content (wt%) Figure Comparison of the breakdown potentials and the OCPs in deaerated and aerated chloride solutions. E 1 and E 2 are denoted by circles and squares, respectively. The OCPs in deaerated solution, ph=3.56 after about 30 min prior to the polarization tests are denoted by triangles. The range of OCP in deaerated solution during this period is given by the solid double-ended arrows. The OCPs in aerated solution after 168 h are denoted by crosses and are connected by a line. The range of OCP in aerated solution during this period is given by the dashed double-ended arrows. The variations in potential for AA7075 and AA7050 in aerated solution were so small that they were not included. 129

146 -0.8 Potential (V SCE) mv/s Time (sec) (a) Potential (V SCE) Current Density (A/cm 2 ) (b) Figure (a) Time evolution of the potential applied to AA7004-T6 and (b) potentiodynamic polarization curve for AA7004-T6 after 2 min holding at 1.15 V SCE showing only one breakdown potential for AA7004-T6. The solution was deaerated 0.5 M NaCl, ph=3.56 and the scan rate was 0.2 mv/s. 130

147 Current Density (A/cm 2 ) Under OCP for 30 min Under potential control at V SCE for 2 min E app = mv SCE Time (sec) Figure Potentiostatic polarization curve for AA7004-T6 at 930 mv SCE for 1 h in deaerated 0.5 M NaCl, ph=3.56 under two different conditions before polarization: (a) under OCP for 30 min and (b) under potential control at 1150 mv SCE for 2 min. 131

148 CHAPTER 4 CHARACTERIZATION OF CHROMATE CONVERSION COATING ON AA7075-T6 4.1 Introduction Chromate conversion coatings (CCCs) have been subjected to intensive study over the past decades due to the effective corrosion protection they confer to Al alloys and the desire to develop an effective and environment-friendly replacement. The composition and structure of CCCs formed on pure Al and Al alloys have been extensively investigated using a variety of analytical tools [1-12]. It is generally accepted that CCCs form via a redox reaction between Al in the alloy and Cr(VI) species in the chromating solution. CCCs are amorphous and mainly composed of hydrated mixed Cr(III)/Cr(VI) oxide (the ratio of Cr(III) to Cr(VI) is about 3:1) [5, 8]. Cr(VI) can be stored in CCCs and released as soluble chromate species when CCCs are exposed to solution [13]. This characteristic of CCCs is believed to lead to the important selfhealing ability of CCC-treated alloys. In order to better understand the unique characteristics of CCCs, an in-depth understanding of CCC formation and growth is necessary. However, the mechanism of 132

149 CCC formation and growth on Al and Al alloys has been debated. On pure Al, several possible models have been proposed. Katzman et al. [1] proposed a uniform coating growth model based on Auger Electron Spectroscopy (AES) depth profiling results of CCCs formed on commercial purity Al in CrO 3 + NaF chromating solution. This model suggests that the initial Al oxide is completely dissolved by fluoride in solution and the bare Al reacts with chromate via a redox reaction to form a uniform CCC, with anodic and cathodic reactions occurring separately at the Al/CCC and CCC/solution interfaces, respectively. This model has difficulty explaining charge transfer through the coating during CCC growth. Brown et al. [14, 15] used Transmission Electron Microscopy (TEM) to study ultramicrotomed CCC sections and concluded that electron tunneling through Al oxide is essential in the formation and growth of CCCs. In contrast to uniform growth models, Arrowsmith et al. [16] proposed that spherical Cr III hydroxide particles nucleate and grow sideways until a coating layer forms by merging of these spherical particles. The gaps between spheres allow contact between fresh solution and Al. Note that all the chromating solutions in the above studies were ferricyanide-free. In ferricyanide containing solution, such as the commercial bath Alodine 1200S, Fe(CN) 3-6 is an accelerator, which can greatly increase CCC coating weight. Xia and McCreery [8, 9] studied the chemical structure of accelerated CCCs and the roles of ferricyanide in CCC formation on AA2024 using Fourier Transform Infrared (FTIR) and Raman Spectroscopy. Their results suggest that CCC growth is mediated by Fe(CN) 3-/4-6. Without the presence of Fe(CN) 3-6, direct redox reaction between Al and chromate proceeds slowly. However, Al reacts quickly with Fe(CN) 6 3- to produce ferrocyanide (Fe(CN) 6 4- ) and Al 3+, and Cr(VI) is quickly reduced to Cr(III) by Fe(CN) 6 4-, which 133

150 oxidizes back to Fe(CN) 3-6. They also found that Cr(VI) reversibly chemisorbs on Cr(OH) 3, which is responsible for storage and release of Cr(VI) species from CCCs. Recently, CCC formation has been described as a sol-gel process based on the fact that CCC has many characteristics found in dip or spin coated xerogel films [17, 18]. In contrast to the forced hydrolysis in sol-gel processing, colloidal Cr 3+ produced via a redox reaction might undergo hydrolysis, condensation and polymerization, followed by drying to form CCC coatings. This viewpoint provides new insight to the study of CCC formation. CCC formation on real Al alloys such as AA2024 and AA7075 is complicated by their heterogeneous microstructure, which includes a matrix phase and a variety of intermetallic particles. These intermetallics play a critical role in the corrosion of Al alloys in chloride solution because they act as anodes or cathodes, and form galvanic couples to the Al matrix [19, 20]. The reactivity of the different intermetallics might affect CCC formation on the intermetallics, and thus influence the overall performance of the coating for corrosion protection of the Al alloy. Most investigations of CCCs on Al alloys have focused on AA2024, which contains Al 2 CuMg and Al-Fe-Cu-Mn coarse intermetallics with dimensions on the scale of microns to tens of microns. In practice, it is difficult to characterize the coating formed on particles of this size and to get depth distribution of the composition because of the limited spatial resolution limitation of surface sensitive techniques. Waldrop and Kendig [7] studied the formation of CCC on AA2024-T3 using Atomic Force Microscopy (AFM). The CCC on the Al matrix phase was found to nucleate and grow very fast in the form of nodules. Nucleation and growth of the coating was more rapid on Al-Fe-Cu-Mn particles than on Al-Cu-Mg. A similar 134

151 AFM study of nucleation and growth of CCCs formed on AA2024-T3 in combination with Field Enhanced Scanning Electron Microscopy (FE-SEM) and TEM was conducted by Brown and Kobayashi [21]. They concluded that CCC formation and growth on intermetallics strongly depended on the size, shape, and composition of intermetallics. It is obvious that AFM and SEM only provide topographic information of CCCs, and information about composition and thickness of CCCs formed on intermetallics is still lacking. In order to overcome the spatial limitation of surface analytical techniques and obtain composition and thickness information for CCCs on intermetallics, large-area intermetallic model samples in the form of bulk ingots and thin films have been studied. McGovern et al. [22] investigated CCC formation on a specially cast Al-Cu-Mg ingot using Raman spectroscopy. The 860 cm -1 peak was used to monitor the Cr(VI)-O-Cr(III) bond in CCC. They found that CCC formation was suppressed on Al-Cu-Mg phases and was lower as the Cu content in the Al-Cu-Mg phase increased as a result of passivation by adsorbed ferricyanide. Juffs et al. [23] studied CCCs formed on macroscopic couples. Castings of intermetallic Al 3 Fe, Al 7 Cu 2 Fe and Al 2 Cu (θ) were separately coupled to AA1100. Analyses by surface analytical techniques revealed that coatings on the matrix were ten times thicker than over the intermetallic phases and increased linearly with immersion time. On the intermetallic phases, the decomposition of ferri/ferrocyanide and fluoride attack were found. Vasquez et al. investigated CCCs formed on AA2024-T3 alloys and thin film analogs of Al 2 Cu (θ), Al 2 CuMg (S), Al 20 Cu 2 (FeMn) 3, and matrix (Al-4Cu) using a variety of surface analytical techniques [24, 25]. A refined view of CCCs formed on AA2024-T3 was proposed. It was suggested that CCCs formed on 135

152 AA2024-T3 are heterogeneous. The thickness of coatings on θ and S intermetallic particles was one tenth of that of coatings on the matrix. The coating on Al 20 Cu 2 (FeMn) 3 particles was rough and exhibited an island structure. The cyanide enriched oxide films covering intermetallics were different from those on the matrix, and Cr was depleted in oxide film formed on Cu rich intermetallics. Little attention has been given to CCCs on AA7075, another important aerospace alloy. Lytle et al. [5] investigated the structure and chemistry of CCCs on AA7075-T6 using X-ray Absorption Spectroscopy (XAS), X-ray Photoelectron Spectroscopy (XPS), FTIR, and AES techniques. They found that CCC thickness ranged from 200 nm to 1 um for coating time of 1~3 min. Fe(CN) 3-6 was found to be present throughout the top 2 nm of the CCC surface. Cr(VI) in tetrahedral coordination with O was found at a concentration of 23 ± 2%, comprising hydrated mixed Cr(III)/Cr(VI) oxide. It should be noted that Lytle et al. studied only the CCC formed on the matrix phase. From the literature, information about the morphology and composition of CCC layers on intermetallics in AA7075-T6 is still lacking. Similar to AA2024-T3, the structure of AA7075-T6 is heterogeneous. The matrix phase of AA7075-T6 is a solid solution of Al and several alloying elements containing fine hardening precipitates on the scale of nm, which are coherent or semi-coherent with the solid solution [26]. The intermetallic particles are categorized into two groups. The first group are coarse intermetallics of about 1~10 µm in size, such as Al-Fe-Cu particles, which are formed during the solidification process [27, 28]. The second group of intermetallics are grain boundary precipitates of about 100 nm in size, such as Mg 2 (Zn Cu Al) particles, which are formed preferentially on the grain boundary during the artificial aging process [29]. In chloride 136

153 containing solution, pitting corrosion usually occurs near coarse intermetallic particles due to galvanic coupling to the matrix [30]. Intergranular corrosion is generally believed to be associated with grain boundary precipitates and the precipitate free zone along grain boundaries [29, 31]. In this study, the CCC formed on AA7075-T6 was characterized using Scanning TEM (STEM). Focused Ion Beam (FIB) cross-sectioning was used to prepare TEM samples of CCCs. Preparation of CCC TEM samples by conventional methods is impossible because the CCC is very thin and can easily peal off when it is mechanically polished. The ultramicrotomy method has been used to prepare cross-section TEM samples of CCCs on Al alloys [4, 15]. The encapsulated samples are trimmed with a glass knife followed by sectioning with a diamond knife to about 10 nm thick. The TEM image resolution can be degraded by mechanical deformation during ultramicrotomy sectioning. In contrast to the ultramicrotomy method, FIB sectioning offers significant advantages. First, FIB sectioning by ion sputtering generates TEM specimens with considerably less mechanical deformation, which improves resolution. Furthermore, a dual beam SEM/FIB tool equipped with EDS can identify the intermetallic particles on the sample surface and allow site-specific sectioning at the particles. 4.2 Experimental Materials AA7075-T6 samples ( mm) were cut from a commercial sheet. The alloy composition was determined by inductively coupled plasma mass spectroscopy (in weight percent, wt%): Zn 5.37, Mg 2.41, Cu 1.36, Fe 0.24, Si 0.04, Cr 0.18, Mn

154 and balance Al. A bulk analog of the Mg 2 Si phase was synthesized by sintering compressed pellets of Mg 2 Si powder at 1000 C for 2 h in an oven that was evacuated and then back-filled with Ar gas. The bulk analog sample of Al 2 CuMg phase used by Buchheit et al. [32] was also studied in this work. Metallographic sections of both intermetallic bulk analogs were prepared. All sample surfaces were polished using a nonaqueous polishing procedure to limit corrosion during polishing. The samples were mechanically ground with successively finer SiC papers through 1200 grit in ethanol, polished down to 0.05 µm with a suspension of alumina in ethanol, cleaned ultrasonically in ethanol, and finally dried in a cold air stream. The polished and air dried (24 h) AA7075-T6 specimens were immersed in commercial Alodine 1200S solution at room temperature for 3 minutes, then rinsed thoroughly with deionized water (resistivity > 18 MΩ) before air drying for 24 h SEM and EDS CCC coated (3 min Alodine immersion time) AA7075-T6 and bulk intermetallic samples were examined using an FEI Sirion FEG-SEM operating at 5 kv and 10 kv, respectively. All phases were analyzed by EDS for elemental composition TEM and STEM Cross-sectional TEM samples were prepared from CCC-treated specimens by an FEI Strata Dual Beam 235M SEM/FIB tool using a 30 kev Ga ion beam and 5 kev electron beam. Prior to sectioning, specific sites underneath the coating such as the matrix, and Al-Cu-Mg, Al-Fe-Cu particles were located and identified by SEM/EDS in 138

155 the FIB, and then covered by deposition of 1.5 µm thick Pt layer to protect the CCC coating during the FIB sectioning. The membrane had an area of 15 µm 5 µm, and was thinned in the FIB to a thickness of about 100 nm for electron transparency. The membrane was plucked out of the bulk sample under an optical microscope using a sharp pyrex needle of about 1 µm in diameter, and placed on a 200 mesh Au TEM grid with a formvar/carbon support film for TEM/STEM analysis. For high resolution TEM (HRTEM) analysis, a 5 wedge of membrane was thinned with FIB to a thickness of less than 50 nm, plucked, placed on a 200 mesh Au TEM grid with carbon support film containing an array of holes. The HRTEM observation was made in the region of the membrane above a hole to remove the influence of the carbon film. A TEM grid with support film was dipped into a suspension of 99% pure Cr 2 O 3 standard powder in ethanol and dried in air. This sample was used for determination of the sensitivity factors for O and Cr. TEM characterization of FIB-sectioned CCC membranes was conducted with both an FEI CM200T TEM operating at 200 kv and an FEI CM300 TEM operating at 300 kv. EDS line profiles were acquired in the FEI Tecnai TF20 STEM operating at 200 kv and quantified using the FEI/Emispec TIA software. The probe size was less than 2 nm and the step size was about 10 nm (4 nm in some measurements). The EDS quantification used the Cliff-Lorimer method with theoretical k-factors. To study the CCC growth kinetics, CCCs on samples exposed for various coating times up to 10 min were sectioned with FIB and the coating thickness was measured directly in the FIB. 139

156 4.2.3 Electrochemistry Open circuit potential (OCP) measurements were performed on AA7075-T6 samples in the Alodine solution. This OCP value was then used to polarize Mg 2 Si and Al 2 CuMg intermetallic analogs immersed in Alodine solution for determination of anodic or cathodic polarity of coarse Mg 2 Si and Al 2 CuMg intermetallic particles present in AA7075-T6 during CCC coating formation. Electrochemical measurements were carried out using a Gamry PC4/FAS1 potentiostat. A three-electrode configuration was used with a saturated calomel reference electrode (SCE) and a Pt mesh counter electrode. To avoid reference electrode contamination, two cells were used and solution bridged with luggin probe with a membrane at one end. 4.3 Results The varying composition and reactivity of different parts of the heterogeneous microstructure of AA7075-T6 are expected to influence the formation of CCC across the alloy. It is indeed possible to identify the location of coarse intermetallic particles during SEM observation of a polished and CCC-treated surface as shown in Figure 4.1. The CCC on the matrix exhibits a typical mud-crack morphology. It should be noted that exposure to the vacuum in the SEM dehydrated the CCCs rather rapidly. The typical mud-cracking morphology associated with the CCC does not extend over the particles. Figure 4.1 shows three different types of particles, which were identified by EDS to be Al-Cu-Mg, Al-Fe-Cu, and Mg-Si phases. The EDS analysis is only qualitative due to the small size of the particles relative to the excitation volume, but it allows approximate identification even in the presence of the CCC, which is relatively 140

157 thin over these phases. The EDS data indicate that the stochiometric compositions of these particles are likely Al 2 CuMg (S), Al 3 (Fe Cu), and Mg 2 Si (β) phases. The SEM image shows steps in the coating at the Al-Cu-Mg and Mg-Si particles, suggesting that the thickness of the coating is less on top of these particles. In contrast, no step is observed at the Al-Fe-Cu particle; the level of the coated surface is almost the same as that of the matrix immediately surrounding it. However, the pattern of mud-crack in the immediate vicinity to this particle is different than the coarse cracking seen further away, suggesting that the CCC is thinner in this region. Other than the mud cracks, the CCC on the matrix was macroscopically uniform. Figure 4.2 is a TEM image of CCC coating on the AA7075 matrix. Coating thickness in the range of about 200~400 nm thick has been observed by inspection of several samples. Details of the CCC structure are evident in Figure 4.2. A cell-like structure is evident with a cell dimension about equal to 5 nm. The alternating light and dark regions are apparently the result of density fluctuations. A high resolution TEM image of a CCC coating on the AA7075 matrix is given in Figure 4.3. This image was taken on a different sample and with a different microscope. The 5 nm features seen in Figure 4.2 are visible as large patches, but a finer structure on the order of 0.8 nm is also evident. It is possible that the structure in Figures 4.2 and 4.3 represents the Cr(III) backbone and interpenetrating channels comprising the CCC. The nano-eds line profiling was conducted on CCCs formed on the matrix. Figure 4.4 shows nano-eds line profiles of the CCC coating on the matrix from the Pt to the matrix along the line on the inset STEM micrograph. The micrograph shows bands within the CCC and grain or sub-grain boundaries in the matrix. The coating is composed 141

158 of Cr and O, and large variation in X-ray intensity for the Cr and O signals is found in the coating. The Cr and O signals have the same intensity variation pattern. Figure 4.5 shows the ratio of the O/Cr raw EDS signals in the CCC region and the ratio of Zn/Al signals in the Al matrix for the linescan shown in Figure 4.4 as well as transmitted electron intensity or brightness along the same line in the STEM high-angle angular dark field (HAADF) image inset in Figure 4.4. Raw X-ray intensity ratios are proportional to material composition. In the CCC region, the O/Cr ratio is nearly constant even though the transmitted electron intensity or brightness profile fluctuates. This suggests that the brightness fluctuations are the result of CCC density fluctuations. In the Al matrix region, the peaks of Zn/Al ratio shown in Figure 4.5 are aligned with the electron density fluctuations associated with grain boundaries or subgrain boundaries features. It is likely that the fluctuations in X-ray intensity are the result of compositional differences at the boundaries. Furthermore, as shown in Figure 4.4, no Al was detected anywhere in the CCC, whereas a low amount of Fe was found in outer part of the CCC. It is believed that Fe in the form of ferri/ferro cyanide was incorporated during CCC formation from the Alodine solution. It is interesting that peaks associated with Mg and Zn, and a small peak associated with Cu, are found at the outer surface of the CCC, just below the Pt layer. This indicates the outer layer of CCC on the matrix could be enriched in Mg, Zn, and Cu. However, this enrichment may be an artifact or noise caused by X-ray counting transients when electron probe scanned from Pt overlayer with high atomic number to the nearby CCC. It should be noted that the data in the line profiles shown in Figure 4.4 are raw integrated intensity values, which are proportional to the concentration for the elements. 142

159 Sensitivity factors (K-factors) are required for the conversion from raw intensity to concentration. Although the EDS quantification can be made by the Emispec TIA software using theoretical K-factors, these standardless K-factors would make the concentration much more scattered. Another more precise approach for quantification is to experimentally determine K-factors using the TEM. Since O and Cr are of main interest, the K-factor for O and Cr was determined with % purity Cr 2 O 3 standard powder. Figure 4.6a is a TEM image of fine Cr 2 O 3 powders, which was analyzed by STEM/Nano-EDS line profiling. Figure 4.6b shows raw integrated intensity ratio of O/Cr, along the line drawn in Figure 4.6a on the fine powder. The Cliff-Lorimer method (Equation 4.1) was used to determine the K-factor for O and Cr. C O /C Cr = K OCr (I O /I Cr ) (4.1) where C O /C Cr denotes the concentration ratio of O to Cr; I O /I Cr denotes the raw integrated intensity ratio of O to Cr; K OCr is a sensitivity factor for O and Cr. By averaging the data points in Figure 4.6b, I O /I Cr for the Cr 2 O 3 standard was determined to be 0.353± Given Cr 2 O 3 where C O /C Cr is 1.5, K OCr was determined to be 4.25±0.23. The application of this factor is discussed below. Figure 4.7 is a TEM image of the coating on a coarse Al 2 CuMg intermetallic particle with diameter of about 2 µm. This particle is covered with a uniform coating layer about 30 nm thick, and it is clearly much thinner than the coating on the surrounding matrix. The sharp step in the coating thickness at the interface between the Al 2 CuMg particle and the matrix indicates that large gradients in coating formation rate can occur locally during immersion in Alodine solution, and that CCC formed by precipitation over the matrix does not migrate laterally to cover the particle where the rate 143

160 of formation is lower. Figure 4.8 is a TEM image of the coating on a coarse Mg 2 Si intermetallic particle about 1.5 µm in size. Due to the similar atomic masses of Mg 2 Si and the Al matrix, the contrast between the particle and matrix is less clear on the TEM image. Similar to the coating on the Al 2 CuMg particle, the layer on the Mg 2 Si particle is uniform and thin, about 20 nm in thickness. This coating is also much thinner than the coating on the surrounding matrix, and a step in thickness is also observed at the boundary of the two phases. Figure 4.9 is a TEM image of the coating on a coarse Al 3 (Fe Cu) intermetallic particle about 5 µm in size. The coating on this particle has a variable thickness and might even be discontinuous in spots. This result is consistent with the observation by Vasquez et al. for the CCC on AlCuFeMn particles in AA2024-T3 [24]. The average thickness of the coating on the Al 3 (Fe Cu) particle is about 50 nm, which is also much less than the average coating thickness on the matrix phase. However, the thickness of the particle coating is nearly the same as that of the coating on the surrounding matrix so no clear change in thickness is evident. The thickness of the coating on the three types of intermetallics and surrounding matrix observed in TEM is consistent with the SEM observations described above. Again, the nano-eds line profiling was conducted on the coatings formed on intermetallic particles. Figure 4.10 shows the x-ray intensity for Cr, O, Al, Cu, and Mg detected along a line profile of the coating on a coarse Al 2 CuMg particle with the scan location and direction indicated by the arrow on the inset STEM micrograph. In the first 60 nm of the line profile, the beam traversed the region of the Pt overlayer, and the apparent intensity of the other elements was attributable noise. The last 80 nm of the profile is in the bulk of the Al 2 CuMg particle. Between the bulk particle and coating, 144

161 there is a 40 nm region where the Mg signal is lower than that in the bulk particle, but the Cu signal is higher than that in the bulk particle. This indicates that Mg dealloying occurred during coating formation on the Al 2 CuMg particle. Between the Pt overlayer and particle is the region associated with the CCC, and the elements Al, Mg, Cr, and O were detected, which indicates that an Al/Mg/Cr mixed oxide or hydroxide covered the Al 2 CuMg surface. Compared to the Mg and Al signals in the coating, the Cr signal was rather small, suggesting that only a small amount of Cr species was incorporated in the coating on the Al 2 CuMg particle. Figure 4.11 shows the line profile of the coating on a coarse Mg 2 Si particle along the line indicated by the arrow on the inset STEM micrograph. The first 65 nm of the scan corresponds to the Pt overlay. In this Pt overlay region, the intensity for all signals is abnormally high for an unknown reason. The intensity level for signals can be regarded as a noise level in the Pt overlay region. Similar to the profile of the coating on Al 2 CuMg particle, the Mg signal in the region between the bulk Mg 2 Si particle and coating is lower than that in the bulk particle and the Si signal is higher, which indicates that Mg dealloying occurred during coating formation on the Mg 2 Si particle. A 50 nm dealloyed layer is evident in the profile in Figure In the region of the coating, Mg, O, and a small amount of Cr were detected, indicating that the Mg 2 Si intermetallic particle was covered by mixed Mg/Cr oxide during immersion in Alodine solution. Figure 4.12 shows the line profile of the coating on a coarse Al 3 (Fe Cu) particle. This coating is composed primarily of Cr and O, which is similar to the CCC formed on the matrix. It should be noted that Fe is also detected, and it is present at a much higher level than in the film formed on the matrix. Another line scan (not shown) was made 145

162 across the thinner region of the coating formed on the same Al 3 (Fe Cu) particle. The coating in this region, though much thinner, had the same composition as the thicker coating formed on other parts of the particle. Furthermore, there was no observable difference in the composition of the Al 3 (Fe Cu) particle under thin and thick regions of the coating. The results from characterization of coatings on intermetallics and matrix suggest that coating formation greatly depends on the local electrochemical reactivity of intermetallics and surrounding matrix. It is well known that Al 2 CuMg and Mg 2 Si intermetallics are active relative to the Al matrix and that Mg dealloying usually occurs in acidic solution, which might help account for thin coatings formed on Al 2 CuMg and Mg 2 Si intermetallics. However, the electrochemical behavior of Al 2 CuMg and Mg 2 Si intermetallics present in Al alloys has not yet been studied in Alodine solution. Figure 4.13 and 4.14 are backscattered electron images of the Al 2 CuMg and Mg 2 Si ingot analogs, respectively. In the Al 2 CuMg analog, EDS analysis indicates that three Al-Cu- Mg phases with various Cu content are present. In the Mg 2 Si analog, Mg 2 Si and Si phases exist. Before testing the electrochemical behavior of Al 2 CuMg and Mg 2 Si intermetallics, an AA7075-T6 sample was immersed in Alodine solution for 250 sec and the open circuit potential (OCP) was recorded. The variation of OCP as a function of time is shown in Figure The alloy OCP is found to be around -600 mv SCE with little variation. The epoxy mounted Al 2 CuMg and Mg 2 Si analogs were polarized at -600 mv SCE in Alodine solution. The current transients for the Al 2 CuMg and Mg 2 Si analogs are shown in Figure Both analogs exhibit anodic currents, which are large initially, 146

163 about A/cm 2, and decrease with time. The anodic current implies that anodic dissolution is predominant during coating formation on Al 2 CuMg and Mg 2 Si phases in AA7075-T6. It is likely that Mg dealloying occurs during immersion in Alodine solution. This result is consistent with TEM characterization of the coatings on Al 2 CuMg and Mg 2 Si particles. Finally, in order to investigate CCC formation kinetics on the matrix, AA7075 samples polished to a 0.05 µm finish were coated by immersion in Alodine solution for various periods of time. The CCC coatings were simply sectioned with the FIB and the SEM capability of the dual beam SEM/FIB instrument was used to determine the thickness of the CCCs on the matrix. Figure 4.17 shows the CCC thickness on the matrix as a function of coating time from 15 sec up to 10 min. The data points represent averages of thickness of three different coating samples for the same coating time. The coating growth on the matrix is found to follow a log-linear kinetics. The line in Figure 4.17 is a fit to the data and has the form: T = log(t) t (4.2) where T represents the coating thickness in nm and t is coating time in sec, with R= It should be noted here that the thickness of the coatings was measured in vacuum, which dehydrated and thus shrank the CCCs. If we assume that the relative amount of coating shrinkage rate was almost same for various coating times, the CCC thickness would be larger than that shown in Figure 4.17, but would follow the same trend. Campestrini et al. using Rutherfold Backscattering Spectrometry (RBS) also found similar trend for the thickness and weight of CCCs formed on an Alclad 2024 alloy as a function of the immersion time in Alodine solution [11]. 147

164 4.4 Discussion The results of this study show that the CCC on AA7075-T6 exhibits different morphology and composition on the matrix and coarse intermetallics in AA7075-T6. Plane-view and cross-sectional images reveal that the coating is much thicker on the matrix than on the coarse intermetallics. This key observation implies that CCC formation is locally dependent on the electrochemical reactivity of the exposed regions, such as the matrix, and three types of coarse intermetallic particles. As a result, the mechanisms of coating formation on matrix and intermetallics might be different Determination of CCC Composition Nano-EDS line profiling of the CCC as shown in Figure 4.4 indicates that the CCC formed on the matrix of AA7075-T6 is composed of Cr and O as well as a small amount of ferri or ferrocyanide. This result for the CCC is consistent with previous XPS, AES, XANES Raman, and FTIR results [5, 6, 8]. XANES, Raman and FTIR results in the literature also show that Cr oxide in the CCC is mixed Cr(III) and Cr(VI) oxide [5, 8]. Xia et al. [8] used UV-Vis to determine the ratio of Cr(VI) to Cr(III) in CCCs formed on AA2024 to be 3 ± 0.6. Lytle et al. [5] used XANES to measure the composition of CCCs formed on AA2024 and AA7075. They found that CCCs contain 23 ± 2 % Cr(VI) with the balance Cr(III). Although about 3:1 Cr(III)/Cr(VI) ratio has been established, the exact formulas of Cr(III) and Cr(VI) oxide in CCCs are still unclear. The Cr(III) oxide could be CrOOH, Cr 2 O 3 or Cr(OH) 3, while the Cr(VI) oxide could be HCrO - 4, Cr 2 O 2-7, or CrO 2-4. At about ph 1.6 of Alodine solution, CrO 2-4 is not stable and would transform itself to HCrO - 4 or Cr 2 O 2-7. Hence, the possibility of CrO 2-4 as the Cr(VI) oxide present in 148

165 CCCs is ruled out. There exist six possible combinations of Cr(III) and Cr(VI) oxides, as listed in Table 4.1. By taking the composition of 0.77Cr(III) Cr(VI) measured by Lytle et al. using XANES [5], the theoretical values of O/Cr atomic concentration ratios for the six possible Cr mixed oxide formulas were calculated and listed in Table 4.1. Although STEM/EDS used in this work cannot sense the oxidation states of Cr, the concentration of total Cr in a mixed Cr oxide can be measured by EDS. Figure 4.5 shows the ratio of O/Cr signals (I O /I Cr ) from raw EDS data for the linescan shown in Figure 4.4. By averaging the data points in the bulk CCC region, I O /I Cr for the CCC was determined to be ± According to Equation 4.1, the atomic concentration O/Cr ratio, C O /C Cr, is calculated to be 2.52 ± 0.35, given the experimentally determined K- factor (K OCr = 4.25 ± 0.23). By comparison with the theoretical values of C O /C Cr for possible Cr mixed oxide formulas listed in Table 4.1, the C O /C Cr value of 2.46 for the combination of CrOOH and HCrO 4 - best matches the experimental value of Therefore, the Cr mixed oxide in the CCC has a composition that is consistent with a mixture of Cr III OOH and HCr VI O - 4. The CCC is, of course, not a simple mixture of oxides; the structure is likely a Cr III OOH backbone with chemisorbed HCr VI O - 4 [13]. It should be noted that CCCs should contain some amount of water. Since water evaporated rapidly from CCCs under ultra high vacuum in TEM, there is no influence of water on above measurement and calculation CCC Coating Formation on Matrix In the literature, most investigations have been focused on CCC formation on the matrix. It is generally believed that when the alloy matrix is immersed in Alodine 149

166 solution, the redox reaction between Al and Cr(VI) species (dominant as HCrO - 4 ) occurs, which is given by Equation 4.3 [33]. Al + HCrO H + = Al 3+ + Cr H 2 O (4.3) Hydrolysis of Cr(III) species forms the CCC backbone, onto which Cr(VI) species are incorporated by chemisorption [8, 13]. In the CCC formation on the matrix, ph seems to play an important role. Forced hydrolysis of Cr(III) salts such as Cr(NO 3 ) 3 9H 2 O by addition of hydroxide reveals that ph 3~4 is necessary to form Cr(III) hydroxide [8, 34]. Deprotonation during redox reaction would increase the local ph in the solution close to the matrix surface leading to a ph gradient from the surface to the bulk solution of ph 1.6. Once ph in the near-surface region increases to 3~4, Cr 3+ cations would hydrolyze and form a hydrated Cr (III) hydroxide. A sol-gel model for CCC formation on matrix, as shown in Figure 2.15, has been proposed [8, 9, 18]. This model involves at least three important steps. First, in the presence of activiator F - and accelerator Fe(CN) 3+ 6 in Alodine solution, the redox reaction forms Cr(III) cations rapidly and intensively, which hydrolyze into monomers of Cr(III) hydroxide in the near surface region with ph of about 3~4. Second, Cr hydroxide monomers bind to each other and form dimers or trimers. Cr(VI) species in the solution can be adsorbed chemically on Cr(III) hydroxide monomers, dimers, or trimers up to a saturation value of about 1:3 Cr(VI):Cr(III). Third, colloidal Cr(III) hydroxides condense and precipitate on the matrix surface followed by polymerization in the solution resulting in sol formation on the matrix surface. When CCC coated alloy is exposed to air, the condensation or polymerization process in CCC gel continues accompanied by dehydration during air aging or drying. 150

167 This sol-gel model for CCC formation on the matrix is supported by the TEM and HRTEM micrographs of CCC coating on the matrix shown in Figure 4.2 and 4.3. The cell-like structure of CCC coatings with a finer structure on the order of 0.8 nm clearly suggest condensation or polymerization of dimers and trimers of Cr(III) hydroxide. This CCC structure represents the Cr(III) backbone and interpenetrating channels comprising the CCC. The channels in the CCC provide a pathway for transport of Cr(VI) from adsorption sites along the backbone to the bulk solution. This transport has been measured and modeled in recent work for CCCs on AA2024-T3 [35]. The diffusivity of Cr(VI) in the CCC structure was found to be on the order of cm 2 s -1, which is much less than typical diffusivities for ions in electrolytes. It was suggested that this low diffusivity resulted from the binding of Cr(VI) to immoble Cr(III) and the physical barrier of the convoluted pathways in the CCC [35]. The observed structure would provide this type of tortuous path for Cr(VI) transport. Another feature of the CCC that is evident in Figure 4.2 is a layering of the structure in the top 20 nm of the film, which is quite different than the random cell-like structure in the rest of the film. It is possible that this layering is an artifact developed during exposure to the vacuum in the FIB chamber prior to Pt deposition. As mentioned above, exposure to vacuum accelerated the formation of mudcracking. It is also possible that the Pt deposition process (ion beam decomposition of an organo-metallic gas) altered the surface structure because of local heating. However, this layering could represent the real structure of the top layer of the CCC. The sol-gel model is also supported by the observation of CCC growth kinetics on the matrix. In spite of shrinkage of CCC coating in high vacuum and ex-situ thickness 151

168 measurement, it is found that CCC coating on the matrix follows a logarithmic-linear growth kinetics. At the beginning of CCC formation, especially within first 30 seconds, Cr(III) hydroxide monomers are produced rapidly in the presence of fluoride and ferricyanide. The production rate of monomers decreases with immersion time in Alodine solution since the precipitation of CCC sol blocks and reduces the exposed area on the matrix. After 30 seconds, precipitation and condensation of CCC sol are dominant and proceed at a linear rate after 30 seconds. The interplay of production of monomers and precipitation of CCC sol results in the logarithmic-linear growth kinetics Coating Formation on Intermetallics Nano-EDS compositional analyses showed that the coating on Mg-containing intermetallics was mixed Al/Mg/Cr oxide for Al 2 CuMg and mixed Mg/Cr oxide for Mg 2 Si, and the coating on Al 3 (Fe Cu) was Cr oxide or CCC. From TEM micrographs, it is evident that the coatings formed on Mg-containing intermetallics were about 20~30 nm thick, much thinner than the coatings formed on the Al matrix. The coatings formed on Al 3 (Fe Cu) intermetallic particles is discontinuous but nearly as thick as the coating on the matrix. The difference in coating thickness and composition for various intermetallic particles should be related to the electrochemical reactivity of the intermetallics. The four types of coarse intermetallic particles studied in this work can be electrochemically categorized into two groups, depending upon whether they are anodic or cathodic relative to the Al matrix in acidic Alodine solution. The Mg-containing intermetallic particles such as Al 2 CuMg and Mg 2 Si are anodic, and Al 3 (Fe Cu) intermetallic particles could be cathodic. Buchheit compiled a list of corrosion potentials 152

169 for a range of intermetallic phases in aluminum alloys [36]. The galvanic series on the noble to active order is: Al 3 Fe > Al matrix > Al 2 CuMg > Mg 2 Si [36]. Although these galvanic relationships were obtained in NaCl or NaCl + H 2 O 2 solution, they provide a qualitative guidance to the galvanic series for the four intermetallic particles in AA7075- T6 in acidic Alodine solution. Anodic dissolution of Al 2 CuMg and Mg 2 Si analogs at the alloy OCP in Alodine solution partially supports these galvanic relationships. It is well known that Al 2 CuMg and Mg 2 Si are very susceptible to dealloying in acidic or chloride containing solution [32, 37, 38]. In 0.1 M NaCl solution with ph of 3, the Al 2 CuMg phase in AA2024-T3 dealloyed with preferential Mg and Al dissolution leaving behind spongy Cu rich residue, which results in Cu redistribution and pitting corrosion in AA2024-T3 [37]. Similar to Al 2 CuMg, Mg 2 Si phase in AA6000 dealloyed in 0.1M phosphoric acid with ph of 1.6 and MgO was found on the Mg 2 Si particles [38]. The anodic current transients for Al 2 CuMg and Mg 2 Si analogs polarized at the alloy OCP in Alodine solution indicates that the anodic dissolution dominated hydrogen evolution or chromate reduction. For both Mg-containing intermetallics, the anodic reaction consists mainly of anodic dissolution of Mg or dealloying during coating formation. In the case of Al 2 CuMg, Al would be corroded too. The line profiles also support this dealloying process. Only a small amount of Cr species was detected in the coatings on both Mg containing intermetallics. It is possible that the local ph remained low over the anodic Mg-containing intermetallics, which inhibited the condensation and polymerization of Cr(III) hydorixde monomers. In subsequent rinsing in deionized water, the unpolymerized monomers were rinsed off the coating surface of Mg-containing intermetallics, leaving sharp steps in coating thickness. During the immersion of the alloy 153

170 and two Mg-containing intermetallic analogs in Alodine solution, no hydrogen evolution was observed by unaided eyes, which indicates that hydrogen evolution is less significant than chromate reduction. Once the thin mixed oxide coating formed on the Mgcontaining intermetallics, their electrochemical reactivity was apparently reduced as indicated by the lack of recession of the intermetallics below the matrix/ccc interface, and the decreasing current densities on the order of 50~100 µa/cm 2 during polarization at the alloy OCP in Alodine solution. Strong passivation of the active particles is evident. In contrast to the active Mg containing intermetallic particles, the Al 3 (Fe Cu) intermetallic particles are cathodic relative to the Al matrix. As mentioned, the primary cathodic reaction is thought to be chromate reduction, which increases the local ph over Al 3 (Fe Cu) particles and results in the formation of standard CCC, or a Cr hydroxide, on the particles. Since the CCC also forms on surrounding matrix, it is likely that the cathodic reaction was not localized to these particles. HF in the coating bath affects the oxide surface layer and makes the Al matrix surface catalytic to both dissolution and chromate reduction. That can account for the absence of an obvious step in coating thickness between the cathodic particles and surrounding matrix. However, TEM observation showed that the CCC on the Al 3 (Fe Cu) particles was about 100 nm thick, which is thinner than the CCCs formed on the matrix. Much more Fe was found in the coating on Al 3 (Fe Cu) particles than in the CCC on the matrix. It is likely that the source of this Fe was ferri/ferro cyanide from the bath, which has been shown to passivate noble, Cu rich phases and reduce the amount of film formation [22]. It is not known why the film on the Al 3 (Fe Cu) particles is discontinuous; there was no correlation to variation in 154

171 the particle composition. However, there might be a variation in cyanide passivation across the surface. 4.5 Summary In this study, CCC formed on AA7075 was characterized using SEM, TEM, STEM, and FIB. The following conclusions are drawn: 1. The Cr mixed oxide in CCCs formed on AA7075 matrix has a composition that is consistent with a CrOOH backbone and chemisorbed HCrO The CCC formed on AA7075 is heterogeneous. The coating on the matrix is much thicker than coating on coarse Al 2 CuMg, Mg 2 Si, and Al 3 (Fe Cu) intermetallics. 3. The coating formation on the intermetallics is dependent on electrochemical reactivity of the intermetallics, local ph change, and interaction of the intermetallics with HF and ferricyanide. 4. A sol-gel model for CCC formation is supported by the observations in this study. REFERENCES 1. H. A. Katzman, G. M. Malout, R. Bauer, and G. W. Stupian, Applied Surface Science, 2, 416 (1979). 2. J. A. Treverton and N. C. Davies, Surface and Interface Analysis, 3, 194 (1981). 3. Z. Yu, H. Ni, G. Zhang, and Y. Wang, Applied Surface Science, 62, 217 (1992). 4. G. M. Brown, K. Shimizu, K. Kobayashi, G. E. Thompson, and G. C. Wood, Corrosion Science, 33, 1371 (1992). 155

172 5. F. W. Lytle, R. B. Greegor, G. L. Bibbins, K. Y. Blohowiak, R. E. Smith, and G. D. Tuss, Corrosion Science, 37, 349 (1995). 6. A. E. Hughes, R. J. Taylor, and B. R. W. Hinton, Surface and interface analysis, 25, 223 (1997). 7. J. R. Waldrop and M. W. Kendig, J. Electrochem. Soc., 145, L11 (1998). 8. L. Xia and R. C. McCreery, J. Electrchem. Soc., 145, 3083 (1998). 9. L. Xia and R. L. McCreery, J. Electrochem. Soc., 146, 3696 (1999). 10. M. W. Kendig, A. J. Davenport, and H. S. Isaacs, Corrosion Science, 34, 41 (1993). 11. P. Campestrini, E. P. M. v. Westing, and J. H. W. d. Wit, Electrochimica Acta, 46, 2553 (2001). 12. M. J. Vasquez, G. P. Halada, and C. R. Clayton, Electrochimica Acta, 47, 3105 (2002). 13. L. Xia, E. Akiyama, G. Frankel, and R. McCreery, J. Electrochem. Soc., 147, 2556 (2000). 14. G. M. Brown, K. Shimizu, K. Kobayashi, G. E. Thompson, and G. C. Wood, Corrosion Science, 35, 253 (1993). 15. G. M. Brown, K. Shimizu, K. Kobayashi, G. E. Thompson, and G. C. Wood, Corrosion Science, 34, 1045 (1993). 16. D. J. Arrowsmith, J. K. Dennis, and P. R. Sliwinski, Transactions of the Institute of Metal Finishing, 62, 117 (1984). 17. J. H. Osborne, Progress in organic coatings, 41, 280 (2001). 18. G. S. Frankel and R. L. McCreery, Interface, winter, 34 (2001). 19. T. G. Dunford and B. E. Wilde, in Field Metallurgy, Failure Analysis, and Metallography, ASM International, Metals Park, OH (1987). 20. G. S. Chen, M. Gao, and R. P. Wei, Corrosion, 52, 8 (1996). 21. G. M. Brown and K. Kobayashi, J. Electrochem. Soc., 148, B457 (2001). 156

173 22. W. R. McGovern, P. Schmutz, R. G. Buchheit, and R. L. McCreery, J. Electrochem. Soc., 147, 4494 (2000). 23. L. Juffs, A. E. Hughes, S. Furman, and P. J. K. Paterson, Corrosion Science, 44, 1755 (2002). 24. M. J. Vasquez, G. P. Halada, C. R. Clayton, and J. P. Longtin, Surf. Interface Anal., 33, 607 (2002). 25. M. J. Vasquez, J. R. Kearns, G. P. Halada, and C. R. Clayton, Surf. Interface Anal., 33, 796 (2002). 26. J. K. Park and A. J. Ardell, Metallurgical Transactions A, 14A, 1957 (1983). 27. M. Gao, C. R. Feng, and R. P. Wei, Met. Mater. Trans. A, 29A, 1145 (1998). 28. W. K. Johnson, British Corrosion Journal, 6, 200 (1971). 29. J. E. Hatch, Aluminum: Properties and Physical Metallurgy, ASM, Metals Park, OH (1983). 30. E. Lunarska, E. Trela, and Z. Szklarska-Smialowska, Corrosion, 43, 219 (1987). 31. T. Ramgopal, P. I. Gouma, and G. S. Frankel, Corrosion, 58, 687 (2002). 32. R. G. Buchheit, L. P. Montes, M. A. Martinez, J. Micheal, and P. F. Hlava, J. Electrochem. Soc., 146, 4424 (1999). 33. P. L. Hagans and C. M. Haas, Chromate Conversion Coatings, in ASM Handbook, Vol 5, Surface Engineering. 1996, ASM International: Metal Park, OH. 34. E. Matijevic, Acc. Chem. Res., 14, 22 (1981). 35. E. Akiyama, A. J. Markworth, J. K. McCoy, G. S. Frankel, L. Xia, and R. L. McCreery, J. Electrochem. Soc., 150, B83 (2003). 36. R. G. Buchheit, J. Electrochem. Soc., 142, 3994 (1995). 37. R. G. Buchheit, R. K. Boger, M. C. Carroll, R. M. Leard, C. Paglia, and J. L. Searles, JOM, 53, 29 (2001). 38. K. Mizuno, A. Nylund, and I. Olefjord, Corrosion Science, 43, 381 (2001). 157

174 TABLES AND FIGURES Possible mixed Cr oxide formula CrOOH HCrO CrOOH Cr 2 O Cr 2 O HCrO Cr 2 O Cr 2 O Cr(OH) HCrO Cr(OH) Cr 2 O 7 C O /C Cr Table 4.1. Ratios of O/Cr concentration for possible mixed Cr oxide, 77% Cr(III) and 23% Cr(VI). 158

175 (a) (b) (c) Figure 4.1. SEM images of CCC coatings formed on (a) Al 2 CuMg, (b) Al 3 (Fe Cu), and (c) Mg 2 Si intermetallic particles in AA7075-T6. 159

176 Pt layer CCC 100 nm matrix Figure 4.2. TEM micrograph of CCC on matrix of AA7075-T6. 160

177 Figure 4.3. High resolution TEM micrograph of CCC on matrix of AA7075-T6. 161

178 Cr O Al Zn Mg Cu Fe Pt CCC 200 nm matrix matrix Intensity (counts) Pt layer CCC Position (nm) Figure 4.4. Nano-EDS line profile of 3 min CCC on matrix of AA

179 1.0 CCC region Al matrix X-ray Intensity Ratio Pt layer O/Cr ratio Zn/Al ratio Transmitted Electron Intensity (counts) Position (nm) Figure 4.5. Ratio of O/Cr and Zn/Al signals from raw EDS data for the linescan shown in Figure 4.4 and transmitted electron intensity or brightness in STEM HAADF image inset on Figure

180 (a) 0.5 Intensity O/Cr Ratio Position (nm) (b) Figure 4.6. Cr 2 O 3 powder standard for determination of K OCr factor. (a) STEM image of Cr 2 O 3 powders, and (b) intensity O/Cr ratio from raw EDS data for the EDS linescan. The linescan was made along the red line on the fine powder in the STEM image. 164

181 CCC Pt layer Al 2 CuMg 500 nm matrix Figure 4.7. TEM micrograph of coating formed on Al 2 CuMg particle in AA7075-T6. 165

182 Pt layer Mg 2 Si 1 µm matrix CCC Figure 4.8. TEM micrograph of coating on Mg 2 Si particle in AA7075-T6. 166

183 Pt layer CCC Al-Fe-Cu phase matrix 1 µm Figure 4.9. TEM micrograph of coating on Al 3 (Fe Cu) particle in AA7075-T6. 167

184 Al 2 CuMg dealloyed layer particle Intensity (counts) Pt 200 nm Cr O Al Cu Mg coating 400 Pt layer Position (nm) Figure Nano-EDS line profile of coating on Al 2 CuMg particle in AA7075. No Fe signal was detected in linescans. 168

185 Mg 2 Si matrix Pt Cr O Mg Si nm dealloyed layer particle Intensity (counts) Pt layer coating Position (nm) Figure Nano-EDS line profile of coating on Mg 2 Si particle in AA7075. No Fe signal was detected in linescans. 169

186 800 Pt Al 3 (Fe Cu) particle nm Intensity (counts) CCC Cr O Al Fe Cu 200 Pt layer Position (nm) Figure Nano-EDS line profile of the thicker region of the coating on Al 3 (Fe Cu) particle in AA

187 A B C Figure Backscattered electron image of Al 2 CuMg ingot analog containing A: Al- 33.6Mg-28.6Cu, B: Al-33.7Mg-15.4Cu, and C: Al-36.2Mg-4.5Cu by atomic percentage. 171

188 Mg 2 Si Si Figure Backscattered electron image of Mg 2 Si ingot analog consisting of Mg 2 Si and Si. 172

189 -0.4 Open Circuit Potential (V SCE) Immersion Time (sec) Figure Open Circuit Potential (OCP) for AA7075-T6 in Alodine solution. 173

190 300 Current Density (ua/cm 2 ) Al 2 CuMg analog Mg 2 Si analog Immersion Time (sec) Figure Current transients for Al 2 CuMg and Mg 2 Si ingot analogs in Alodine solution at 600 mv SCE, which is the OCP of AA7075-T6 in Alodine solution. 174

191 Coating Thickness (nm) Coating Time (sec) Figure CCC coating thickness following a logarithmic-linear growth kinetics. 175

192 CHAPTER 5 EFFECT OF CU CONTENT ON THE PROTECTION OF AA7XXX-T6 BY CHROMATE CONVERSION COATINGS 5.1 Introduction CCCs effectively protect Al and Al alloys from corrosion because of some unique characteristics. First, CCCs provide a barrier layer that separates aggressive environments from the Al substrate. Second, CCCs might act as a bipolar membrane [1], which inhibits the adsorption and ingress of aggressive Cl - ion, as well as the egress of metal cations. Third, CCC has a unique self-healing property, which has been attributed to soluble or leachable Cr VI that is stored in the CCC. The CCC self-healing process is thought to involve several steps. Leachable Cr VI species release from the CCC to the solution [2-4], migrate to the active damage sites such as scratches or pits [5], and reduce to insoluble Cr III oxide or hydroxide. This Cr III film blocks the exposed metal [5, 6] and inhibits cathodic reactions [7]. The corrosion protection performance of CCCs can be influenced by many factors such as the composition, microstructure, and surface pretreatment of the Al alloy substrate. Ketcham [8] conducted salt spray tests on CCC treated panels of AA6061-T6 176

193 (0.15~0.40 wt% Cu), AA7075-T6 (1.2~2.0 wt% Cu), and AA2024-T3 (3.8~4.9 wt% Cu). None of the CCC treated AA6061 samples exhibited corrosion even after 960 h exposure. All CCC treated AA7075 samples passed 480 h exposure with no sign of corrosion but most were corroded after 960 h exposure. All CCC treated AA2024 samples passed 240 h exposure without corrosion but most were corroded after 480 h exposure. Ketcham suggested that corrosion protection by CCC decreases with increasing Cu content of Al alloys. Recent studies revealed that salt spray testing eroded CCCs and leached Cr VI species stored in CCCs [9, 10]. As described in previous chapters, CCC formation is heterogeneous on Al alloys with heterogeneous microstructures. The coatings formed on coarse intermetallic particles are much thinner than CCC formed on the Al matrix. Corrosion of CCC coated Al alloys is likely to inititate at these intermetallic particles covered by thin coating layers, particularly after long-term exposure to the corrosive environments. Cu enrichment during pretreatment and CCC processing plays a critical role in CCC formation and performance. Cu smut forms on the surface when Al alloy samples are immersed in alkaline degreasing solution in the CCC pretreatment process. This layer interferes with the CCC formation [11]. Using XPS depth profiling Hagans and Haas found the enriched Cu layer formed on the surface of polished 2024-T3 alloys remained intact during the CCC formation [12]. Sun et al. [13] used SEM, TEM and XPS to study the effect of acidic pretreatment on corrosion protection of CCCs on AA2024-T3. It was suggested that Cu enrichment on the surface was deleterious to CCC protection. Hughes et al. [14] found that CCC formed on AA2024-T3 was heterogeneous and Cu rich regions 177

194 were present at the alloy/coating interface. Recently, SIMS element mapping revealed that Cu enrichment was found in the outermost layer of CCC on AA2024-T3 [15]. Most investigations reported in the literature have focused on CCC formation and effects of CCCs on AA2024, whereas little attention has been paid to CCCs on AA7xxx. The purpose of this chapter is to develop a better understanding of the effect of Cu content of AA7xxx on protection by CCCs and the effects of acid pretreatment. 5.2 Experimental The five AA7xxx-T6 alloys described in chapter 3 were also studied in this part of the investigation. The AA7xxx-T6 samples were polished to 1200 grit (about 6 µm finish) in ethanol with SiC paper. Some AA7075-T6 samples were further polished to 1 µm finish with a suspension of alumina in ethanol. Some of the samples with 1200 grit finish were given an acid pretreatment by immersion in a solution composed of 1.5 ml HF, 10 ml H 2 SO 4, and 90 ml H 2 O for 30 s followed by rinsing thoroughly in deionized water and drying by a cold air stream. Both as-polished and acid pretreated samples were coated with CCC by immersion in commercial Alodine 1200S solution at room temperature for 3 min, and then rinsed thoroughly with deionized water before air drying for 24 h. Potentiodynamic polarization measurements were performed on CCC AA7xxx- T6 samples in aerated 0.5 M NaCl at a scan rate of 0.2 mv/s. For comparison, polarization measurements were also performed on bare AA7xxx-T6 samples in deaerated 0.5 M NaCl with ph 3.56, which were described in chapter 3. Polarization experiments were performed using a Gamry PC4/FAS1 potentiostat. A Pt counter 178

195 electrode and saturated calomel reference electrode (SCE) were used. CCC treated samples were immersed in the chloride solution for 24 h prior to the polarization measurements. To study CCC breakdown, a CCC treated AA7075-T6 sample was potentiodynamically polarized to a current density of A/cm 2. The surface of the CCC treated AA7075-T6 sample after the polarization test was examined using optical microscopy and scanning electron microscopy (SEM). The SEM examination was performed using an FEI Sirion FEG-SEM operating at 10 kv. Electrochemical Impedance Spectroscopy (EIS) measurements were conducted at open circuit in aerated 0.5 M NaCl for 168 h on as-polished and acid pretreated samples coated with CCC. The impedance measurements were performed as a function of frequency between 10 khz and 10 mhz using a sinusoidal voltage modulation of 10 mv. The experiments were conducted using a Princeton Applied Research (PAR) Model 273 potentiostat with a Solartron Model 1255 frequency response analyzer. Open circuit potentials (OCP) for CCC treated samples were also recorded. To study the change in morphology and composition of the sample surface after the acid pretreatment, X-ray Photoelectron Spectrometry (XPS) measurements as well as metallographical examinations were performed on AA7075-T6 samples with and without the acid pretreatment. The details about the XPS instruments and measurements were described in chapter

196 5.3 Results CCC Breakdown In contrast to polished samples, every CCC treated AA7xxx-T6 alloy exhibited only one breakdown potential in 0.5 M NaCl, as shown in Figure 5.1. Breakdown potentials for as-polished alloys measured in deaerated 0.5 M NaCl are also shown in Figure 5.1. As was found for the bare alloys, the breakdown potential of CCC treated samples also increased logarithmically with increasing Cu content of substrate alloys. The breakdown potential for CCC was about 10 mv above the second breakdown potential for each of the bare substrate alloys. This result implies that the breakdown of CCC is associated with the second breakdown of the substrate alloys. As described in chapter 3, matrix dissolution by selective grain attack occurs above the second breakdown potential in the case of bare alloys. To better observe CCC breakdown sites, AA7075-T6 samples were polished to a 1 µm finish. AA7075 samples with finer finish with and without CCC were potentiodynamically polarized in aerated and dearated 0.5 M NaCl, respectively. It should be noted that the deaerated solution was intentionally acidified to ph 3.56, while the aerated solution was just open to the air without ph control. Slight acidification of the deaerated solution was found to significantly lower the OCP for better determination of the breakdown potentials. There is little effect of ph on breakdown potentials in this region. Figure 5.2 shows the polarization curves for both samples. Similar to the polarization curve for bare AA7075-T6 with 1200 grit finish, two breakdown potentials were also observed for the bare AA7075 with 1 µm finish. However, the values of the two breakdown potentials for the 1 µm finish (E 1 = -757 mv SCE and E 2 = -698 mv 180

197 SCE) were both 22 mv higher than the values for 1200 grit or 6 µm finish. Others have found that breakdown potential increases as surface roughness decreases [16]. The polarization curve for CCC shown in Figure 5.2 reveals that the CCC treated sample had only one breakdown potential, which was right above the second breakdown potential for the bare substrate alloy. There might be some effect of aeration on the breakdown potential, but deaeration is typically considered to only decrease the OCP by eliminating oxygen reduction, with little or no effect on the breakdown potential. Below the breakdown potential for CCC, small current transients can be seen in the passivation region. Such transients are typically associated with metastable pitting. Metastable pitting transients were not observed for the bare AA7075. However, the larger passive current density and the peak associated with E 1 for the untreated alloy resulted in a large background current, so small current transients would have been impossible to see. Another 1 µm finish AA7075 sample coated with CCC was potentiodynamically polarized to -635 mv SCE, above the breakdown potential, at which point the current density was A/cm 2. Figure 5.3 is an optical macrograph of the CCC surface. The round region with a light color in the macrograph is the exposed area of 1 cm 2. The light color of the exposed region of the CCC indicates that Cr VI species leached out into the solution during the 24 h immersion period prior to the potentiodynamic polarization. After the polarization scan, several small pits of a few hundred microns in size were found on the CCC surface. Figure 5.4 shows SEM secondary and backscattered electron images of one of the small pits observed in Figure 5.3. Breakdown apparently occurred on the CCC coated matrix. The extent of damage makes it difficult to know whether the initial breakdown was associated with coarse intermetallic particles. 181

198 5.3.2 EIS Measurement on CCCs Figure 5.5 shows the open circuit potentials (OCPs) for CCC treated samples as a function of immersion time. There was no significant variation of the OCPs up to 168 h immersion for each alloy. The OCPs of CCC treated alloys with high Cu content, AA7075 and AA7050, were about 200 mv higher than those of CCC coated alloys with low Cu content, AA7004, AA7039, and AA7029. EIS measurements were performed at the OCP on CCC coated AA7xxx-T6 during the 168 h immersion in aerated 0.5 M NaCl. All EIS data were well fitted by the simple equivalent circuit model shown in Figure 5.6. Low frequency impedance or coating resistance of the CCCs, R c, was extracted. There was no significant variation for coating resistance of each of the CCC coated alloys during the 168 h immersion period. Figure 5.7 shows the coating resistance of aspolished AA7xxx samples coated with CCC after 48 h immersion. There was dramatic increase in coating resistance from Cu-free AA7004 to low Cu-containing AA7039 (0.077 wt% Cu) and the coating resistance continued to increase as the Cu content increased above wt%. The low frequency impedance or polarization resistance for the uncoated alloys follows the opposite trend with Cu content, as is also shown in Figure 5.7. As described in chapter 3, Cu enrichment and redistribution during corrosion of the bare alloys enhances the oxygen reduction reaction in aerated environments, and decreases the polarization resistance of the high Cu content alloys relative to the low Cu content alloys. In contrast, higher Cu content improves the corrosion resistance of CCC treated AA7xxx-T6 alloys, even in aerated chloride solutions. EIS measurements were also performed after 48 h immersion in aerated 0.5 M NaCl on CCC treated AA7xxx-T6 samples that were pretreated by immersion in HF + 182

199 H 2 SO 4 acid prior to the CCC treatment, as shown in Figure 5.7. The coating resistance for CCC on the acid pretreated alloys first increased with increasing Cu content in the substrate alloy similar to the behavior of polished and coated samples. However, the coating resistance then decreased when Cu content was greater than 0.69 wt% (AA7029) Cu Enrichment by Acid Pretreatment An optical micrograph of the surface of an acid pretreated AA7075-T6 sample shows that the alloy was etched during the acid pretreatment, Figure 5.8. Etched grains and etched coarse intermetallic particles in the alloy microstructure are evident in Figure 5.8. X-ray Photoelectron Spectroscopy measurements (XPS) were performed on large areas of the surfaces of both as-polished and acid pretreated AA7075-T6 samples. Figure 5.9 shows Al, Zn, Mg, Cu, and O XPS spectra. The C 1s spectrum was also obtained for energy calibration. The chemical states for all the elements were identified; the XPS data for the as-polished sample were listed in Table 3.4 in chapter 3. Relative to the XPS spectra for the as-polished sample, there was significant change in the XPS spectra for the acid pretreated sample. The Al 0, Zn 0 and Mg 0 peaks disappeared. The height of the Al 3+ and O 2- peaks corresponding to Al(OH) 3 were dramatically reduced, which indicates that Al(OH) 3 on the surface was dissolved in the presence of HF to produce soluble AlF However, the Cu 0 peak height greatly increased. The significant change in element peaks indicates that the surface of the acid pretreated sample was enriched in Cu and depleted in Al, Zn, and Mg. Optical micrographs of CCC treated samples are shown in Figure The CCC formed uniformly on the as-polished sample. No mud cracks were observed in the CCC 183

200 since the dehydration of the CCC was slight after 24 h aging in air. In contrast, the CCC formed on the acid pretreated sample was thinner and non-uniform. The etched grain structure underneath the CCC is still visible in Figure 5.10b. The color of the CCC on the acid pretreated sample was light yellow as opposed to golden color of the CCC on the aspolished sample. 5.4 Discussion Role of Cu Content in CCC Protection of AA7xxx-T6 Low frequency impedance or coating resistance of CCC on the as-polished and acid pretreated AA7xxx-T6, as determined by EIS, was shown in Figure 5.7. Coating resistance of CCC on as-polished AA7xxx-T6 was found to increase with Cu content. In contrast, coating resistance of CCC on acid pretreated AA7xxx-T6 was found to increase with increasing Cu content, and then decrease when Cu content is greater than 0.69 wt% (7029). For the Cu-containing alloys, coating resistance of CCC on either as-polished or acid pretreated alloys is much higher than that for bare polished alloys in control experiments. These results above indicate that CCC indeed increases corrosion resistance of AA7xxx-T6 and the role of Cu content in CCC protection is complex, which greatly depends on surface treatment. The CCC formed on as-polished AA7075-T6 was characterized in detail in chapter 4. For as-polished samples, CCC formed on the matrix was found to be much thicker than coatings formed on coarse intermetallic particles. No Cu enrichment on the as-polished samples was found during the CCC formation. The increase in coating resistance in aerated solution with Cu content is related to the effect of Cu on the 184

201 breakdown potential measured in deaerated solution. Figure 5.11 is a combination of Figures 5.1 and 5.5. It allows a comparison of the CCC breakdown potentials and the OCPs of CCC treated samples in aerated solution. The OCPs at the end of the 1 week immersion and the variations of the OCPs during the immersion are represented by cross symbols and error bars, respectively. The breakdown potential for CCC on polished AA7xxx-T6 increased with substrate Cu content. For all CCC treated alloys, the OCP remained below the respective breakdown potentials. The OCP of AA7004 was only about 30 mv below its breakdown potential, and this amount was decreased by the 10 mv magnitude AC signal applied during EIS measurements. This alloy exhibited the lowest coating resistance and pits were indeed found on CCC on the AA7004 samples after only 24 h of immersion. For every other alloy, the OCP was at least 60 mv below the breakdown potential during the whole 1 week exposure. As mentioned above, Cu enrichment occurs on uncoated Cu-containing samples during immersion in aerated chloride solution. For CCC treated samples Cu enrichment does not occur because the CCC is protective. As a result, the OCP stays below the breakdown potentials, and the alloys with higher Cu content exhibit higher resistance to localized corrosion because of the resistance of the underlying alloy to breakdown. In contrast to the as-polished AA7xxx-T6 samples, the acid pretreatment or etching dramatically altered the surface of the samples. The acid pretreatment etched the surface and enriched it in metallic Cu. Similar observations have been reported by Sun et al., who studied CCCs on AA2024-T3 [13]. They found that significant Cu enrichment at the oxide-alloy interface after 30 s acid pretreatment in the same solution. 185

202 To explain the poor protection performance of CCCs on the acid pretreated samples, the role of ferricyanide in CCC formation on Cu enriched layers must be taken into account. Ferricyanide is an accelerator that mediates the Al dissolution and chromate reduction reactions [17]. However, ferricyanide was found to inhibit CCC formation on Cu or Cu rich intermetallic particles [18]. Evidence for the ferricyanide inhibition of CCC formation on the acid pretreated sample is found in the optical micrographs of CCCs formed on the as-polished and acid pretreated samples. The CCC formed on the aspolished AA7075-T6 sample was golden in color, whereas the color of the CCC formed on the acid pretreated sample was light yellow. Furthermore, the CCC on the acid pretreated sample was so thin that the microstructure underneath is visible. A larger amount of Cr VI was likely stored in the thicker CCC formed on the as-polished sample than on the acid pretreated sample. For AA7xxx with higher Cu content, the Cu enrichment after the acid pretreatment is more significant and thinner and less protective films form. Based on the above discussion, Cu can have different effects on CCC protection: Cu is beneficial to CCC protection on as-polished AA7xxx-T6, but it is detrimental if enriched on the surface prior to CCC formation Relevance to Salt Spray Testing In industry, the standard method to evaluate CCC protection is salt spray testing (SST). SST is simple and reproducible, but the testing time is very long and the evaluation by simple inspection after testing is qualitative. Polarization and EIS are alternative and quantitative methods, although these methods have not been adopted as 186

203 industrial standards. Nonetheless, there is some relevance of the results obtained by polarization and EIS to the results obtained by SST. As mentioned above, a report on SST of CCC coated AA6061, AA7075, and AA2024 concluded that CCC protection performance decreases with increasing Cu content in the substrate alloys [8]. However, the opposite conclusion was drawn by EIS measurements in this work. This discrepancy might arise from two reasons. It is possible that pretreatment of the sample prior to CCC treatment in the earlier work resulted in Cu enrichment on the surface. The typical pretreatment procedure involves degreasing by immersion in alkaline solution and deoxidizing or desmutting by immersion in acid solution. Although this pretreatment is not so aggressive as the HF + H 2 SO 4 acid pretreatment used in this work, the surface of the samples is altered. During the pretreatment, Cu enrichment would be significant in high Cu content alloys such as AA7075 and AA2024. Therefore, CCC protection on high Cu content AA7075 and AA2024 could be dramatically decreased by the Cu enrichment introduced in the pretreatment. Another possible reason is erosion of the CCC by long-term SST. It is known that a long-term SST erodes the CCC and leaches out stored Cr VI species [9, 10]. The self-healing function of CCC is gradually eliminated by erosion during SST. It is possible that Cu enrichment and OCP ennoblement occurs after long term SST when the CCC is severely degraded. In this case, the higher Cu content alloys would be more susceptible, as for untreated samples. 187

204 5.5 Summary The effect of Cu content on CCC protection of AA7xxx-T6 was investigated by polarization, EIS, and XPS as well as metallography and SEM examination. The following conclusions are drawn: 1. For each of AA7xxx-T6 alloys, CCC have only one breakdown potential, which is slightly higher than the second breakdown potential for the bare alloy. CCC breakdown occurs in the matrix region. 2. Alloy Cu can have different effects on CCC protection: Cu is beneficial to CCC protection for coatings formed on polished AA7xxx-T6, but Cu is detrimental if it is enriched on the surface prior to CCC formation. REFERENCES 1. D. Chidambaram, C. R. Clayton, and G. P. Halada, J. Electrochem. Soc., 150, B224 (2003). 2. A. L. Glass, Materials Protection, 7, 26 (1968). 3. L. Xia, E. Akiyama, G. Frankel, and R. McCreery, J. Electrochem. Soc., 147, 2556 (2000). 4. E. Akiyama, A. J. Markworth, J. K. McCoy, G. S. Frankel, L. Xia, and R. L. McCreery, J. Electrochem. Soc., 150, B83 (2003). 5. J. Zhao, G. S. Frankel, and R. L. McCreery, J. Electrochem. Soc., 145, 2258 (1998). 6. C. S. Jeffcoate, H. S. Isaacs, A. J. Aldykiewicz, and M. P. Ryan, J. Electrochem. Soc., 147, 540 (2000). 7. W. J. Clark and R. L. McCreery, J. Electrochem. Soc., 149, B379 (2002). 8. S. J. Ketcham, Properties of Chemical Films on Al Alloys, in The Finishing of Aluminum, G. H. Kissin, Reinhold Publishing Corp., New York, NY (1963). 188

205 9. F. W. Lytle, R. B. Greegor, G. L. Bibbins, K. Y. Blohowiak, R. E. Smith, and G. D. Tuss, Corrosion Science, 37, 349 (1995). 10. G. O. Ilevbare, J. R. Scully, J. Yuan, and R. G. Kelly, Corrosion, 56, 227 (2000). 11. N. Fin, H. Dodiuk, A. E. Yaniv, and L. Drori, Applied Surface Science, 28, 11 (1987). 12. P. L. Hagans and C. M. Haas, Surface and Interface Analysis, 21, 65 (1994). 13. X. Sun, R. Li, K. C. Wong, and K. A. R. Mitchell, Journal of Materials Science, 36, 3215 (2001). 14. A. E. Hughes, R. J. Taylor, and B. R. W. Hinton, Surface and interface analysis, 25, 223 (1997). 15. M. J. Vasquez, J. R. Kearns, G. P. Halada, and C. R. Clayton, Surf. Interface Anal., 33, 796 (2002). 16. N. J. Laycock and R. C. Newman, Corrosion Science, 39, 1771 (1997). 17. L. Xia and R. L. McCreery, J. Electrochem. Soc., 146, 3696 (1999). 18. W. R. McGovern, P. Schmutz, R. G. Buchheit, and R. L. McCreery, J. Electrochem. Soc., 147, 4494 (2000). 189

206 FIGURES -0.6 Potential (V SCE) E (CCC) b E (bare) 2 E (bare) Cu Content (wt%) Figure 5.1. Comparison of breakdown potentials for CCC and bare AA7xxx-T6. 190

207 E b E 2 Potential (V SCE) -0.8 CCC 7075 E bare um finish surface Current Density (A/cm 2 ) Figure 5.2. Anodic polarization curves for bare and CCC AA7075-T6 (1 µm finish) in 0.5 M NaCl at a scan rate of 0.2 mv/s. The solutions for the bare and CCC samples were deaerated by Ar gas and open to air, respectively. The ph of the deaerated solution was acidified to

208 5 mm Figure 5.3 Optical macrograph of CCC coated 7075 (1 µm finish) polarized to 635 mv SCE with the ending current density of A/cm 2 in aerated 0.5 M NaCl. 192

209 (a) (b) Figure 5.4. SEM images of CCC coated 7075 (1 µm finish) polarized to 635 mv SCE with the ending current density of A/cm 2 in aerated 0.5 M NaCl. (a) secondary electron image, (b) backscattered electron image. 193

210 -0.6 Open Circuit Potential (V SCE) min CCC Immersion Time (h) Figure 5.5 Variation of OCP for CCC coated AA7xxx as a function of immersion time during immersion in aerated 0.5 M NaCl. 194

211 C c R s R c Figure 5.6. The equivalent circuit model used to fit EIS data for CCC. R s is solution resistance, R c is CCC coating resistance, and C c is CCC coating resistance. 195

212 10 7 Low Frequency Impedance (ohm cm 2 ) CCC polished CCC pretreated bare polished Cu Content (wt%) Figure 5.7. Coating resistance of CCC on polished and acid pretreated AA7xxx-T6. Polarization resistance of bare polished AA7xxx-T6 was measured as control experiments. All EIS measurements were conducted at OCP after 48 h immersion in aerated 0.5 M NaCl. 196

213 90 µm Figure 5.8. Optical micrograph of acid pretreated surface of AA7075-T6. 197

214 5000 counts/s Al 2p Intensity as-polished pretreated Binding Energy (ev) (a) Al 2p 5000 counts/s Zn 2p 1/2 Zn 2p 3/2 Intensity as-polished pretreated Binding Energy (ev) (b) Zn 2p Figure 5.9. XPS spectra measured from as polished and acid pretreated AA7075-T6 samples. (Continued) 198

215 Figure 5.9: (Continued) 200 counts/s Mg 2p Intensity as-polished pretreated Binding Energy (ev) 40 (c) Mg 2p 5000 counts/s Cu 2p 1/2 Cu 2p 3/2 Intensity pretreated as-polished Binding Energy (ev) (d) Cu 2p (Continued) 199

216 Figure 5.9: (Continued) counts/s O 1s Intensity as-polished pretreated Binding Energy (ev) (e) O 1s Figure 5.9. XPS spectra measured from as polished and acid pretreated AA7075-T6 samples. 200

217 90 µm (a) 90 µm (b) Figure Optical micrograph of CCC on (a) as-polished and (b) acid pretreated AA7075 samples. 201