Corrosion Inhibition Mechanisms of Aluminum Alloy 2024-T3 by Selected Non-Chromate Inhibitors

CORROSION INHIBITION MECHANISMS OF ALUMINUM ALLOY 2024-T3 BY SELECTED NON-CHROMATE INHIBITORS

DISSERTATION

Presented in Partial Fulfillment of the Requirements for the Degree Doctor of Philosophy in the Graduate School of The Ohio State University

Omar A. Lopez-Garrity, M.S.

Graduate Program in Materials Science and Engineering

The Ohio State University

2013

Dissertation Committee:

Professor Gerald S. Frankel, Advisor

Professor Rudolph G. Buchheit

Professor John Morral

Omar Lopez-Garrity

2013

ABSTRACT

The pursuit to find a chromate-alternative has led to the development of several chromate-free aerospace primers and coating systems that offer good protection.

However, fundamental understanding of the functionality of the chromate-free pigments that are embedded within these coating systems is lacking. The objective of this study was to understand the fundamental mechanism of corrosion inhibition of aluminum alloy

2- 2- 3+ 2024-T3 by molybdate (MoO4 ), silicate (SiO3 ), and praseodymium (Pr ) with the goal of developing the kind of understanding that was accomplished for chromate.

Furthermore, since most inhibiting conversion coatings and pigments act by releasing soluble species into the local environment, it was of interest to understand the mechanism of inhibition in aqueous 0.1 M NaCl solution.

The mechanism of inhibition of AA2024-T3 by the select non-chromate inhibitors was investigated using various electrochemical, microscopic and spectroscopic techniques. Naturally aerated polarization curves showed that molybdate provided mixed inhibition in near-neutral pH and at a threshold concentration of 0.1 M. The largest effect was a 250 mV increase in the breakdown potential associated with pitting and a 350 mV decrease in the open-circuit potential (OCP). In addition, electrochemical impedance indicated that the corrosion inhibition mechanism is oxygen-dependent owing to the protection afforded by Mo(VI) species. It was proposed that the corrosion inhibition of

AA2024-T3 by molybdate may occur following a two-step process whereby molybdate is ii rapidly reduced to MoO·(OH)2 over the intermetallic particles and is subsequently oxidized to intermediate molybdenum oxides (e.g. Mo4O11) in the presence of oxygen which is reduced. This in turn may lead to a local acidification, promoting the condensation and polymerization of molybdate species in solution to form polymolybdate

6- 4- species (Mo7O24 and Mo8O26 ). Furthermore, S-phase particle dissolution is decreased, suppressing surface copper enrichment and significantly lowering oxygen reduction kinetics.

Electrochemical polarization curves show that silicate provides strong anodic inhibition in high alkaline conditions at a threshold concentration of 0.01 M. The largest effect was a 1 V increase in the breakdown potential at a silicate concentration of 0.025

M. The corrosion inhibition mechanism involves the formation of aluminosilicate compounds by the reaction of silicate anions in solution and the aluminate ions that form during oxide dissolution. The Na+ ions adsorb to the negatively charged surface and coordinate with the hydroxyl group of the aluminosilicate anions, thereby forming a protective thin-film over the Al matrix. At near-neutral pH, silicate partially blocks attack of the intermetallic particles by a precipitation mechanism that results in the formation of silica- and silicate- based derivatives. In acidic solution, it was shown that activation of the aluminum surface promotes the formation of a mixed silica/aluminosilicate film over the surface that is porous in nature, providing very little corrosion protection. Furthermore, strong synergistic behavior was observed when molybdate was combined with silicate resulting in much lower threshold concentrations for corrosion inhibition.

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Naturally aerated polarization curves show that praseodymium provides cathodic inhibition in near-neutral pH solution. The largest effect was an order of magnitude decrease in the oxygen reduction kinetics at an optimum praseodymium concentration of

0.0002 M. Corrosion inhibition of AA2024-T3 by praseodymium involves the formation of an insoluble film composed of over the intermetallic particles in response to an increase in the pH at the metal/electrolyte interface. As the cathodes on the surface become blocked by a thick-oxide, new sites are activated resulting in the formation of a thick film across the whole surface. Furthermore, electrochemical impedance in decarbonated solution showed that the absence of CO2 decreased the total resistance relative to the same aerated solution. The hydroxycarbonate protective layer formed in aerated solution is essential for the protection of the alloy indicating that CO2 plays an important role in the inhibition mechanism by praseodymium. In low and high solution pH, praseodymium renders poor inhibition. In the latter case, it was shown that praseodymium still forms a thick oxide film over the intermetallic particles as a result of oxygen reduction, but is unable to stop uniform dissolution of the Al-oxide film.

Dedicated to my wife Meng

And to my parents

ACKNOWLEDGEMENTS

I would like to express my sincerest gratitude to my advisor, Dr. Gerald Frankel, for his advice, guidance and encouragement throughout the course of my PhD career. I especially would like to thank Dr. Frankel for allowing me to have the opportunity to be a part of the Fontana Corrosion Center. I also acknowledge his patience and support during the completion of this document. I would also like to thank Dr. Rudy Buchheit for his help and constructive criticism during the course of my research. Furthermore, I would also like to thank Dr. John Morral for agreeing to serve on my oral examination committee. I sincerely appreciate his time and patience during the completion of this document.

I would also like to acknowledge a number of people who helped me during the course of my PhD work, including Dr. Lisa Hommel, Dr. Saikat Adhikari, Dr. Belinda

Hurley, Dr. Anusha Chilukuri, Dr. Brendy Rincon-Troconis, Meng Tong Lopez-Garrity,

Jimwook Seong, and Pitichon Klomjit. I would also like to extend my thanks to Mrs.

Christine Putnam and Mr. Mark Cooper for all of the administrative and logistic support.

I also acknowledge the help of Ali Lopez-Garrity with sample preparation and lab assistance. I would also like to thank current and past FCC members for their support and constant assistance, including Dr. Bastion Maier, Dr. Federico Gambina, Dr.

Marianno Kappes, Dr. Jesus Vega, Dr. Liu Cao, Dr. Huang Lin, Shanshan Wang, and the rest of the group. vi

Most of all, I would like to thank my whole family for their love and support, especially my parents, Kevin Garrity and Aida Lopez-Garrity, and my grandmother,

Luisa Quintero, for their unconditional love and encouragement throughout my academic career. I would also like to thank my mother in-law, Yinghua Zhao, for her support during the completion of this thesis. And last but not least, I would like to thank my wife, Meng, for all of her love, support, and encouragement during the course of this work.

Finally, I would like to acknowledge the Strategic Environmental Research and

Development Program for funding this work under project number WP-1620.

vii

VITA

July 30th, 1986 ...... Born – Valencia, Venezuela

2008...... B.S. Mechanical Engineering, The Ohio

State University

2010...... M.S. Materials Science and Engineering,

The Ohio State University

2008 to present ...... Graduate Research Associate, Department

of Materials Science and Engineering, The

Ohio State University

FIELDS OF STUDY

Major Field: Materials Science and Engineering

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TABLE OF CONTENTS

ABSTRACT ...... ii ACKNOWLEDGEMENTS ...... vi VITA ...... viii TABLE OF CONTENTS ...... ix LIST OF TABLES ...... xiv LIST OF FIGURES ...... xv CHAPTER 1: INTRODUCTION ...... 1 CHAPTER 2: LITERATURE REVIEW ...... 5 2.1 Corrosion of AA2024-T3 ...... 6

2.2 Corrosion Inhibition by Chromate-Protection Schemes ...... 9

2.2.1 Chromate Background and Toxicity ...... 10

2.2.2 Chromate Inhibition ...... 11

2.2.3 Chromate Conversion Coatings and Pigments ...... 12

2.3 Molybdates ...... 15

2.3.1 Background and Toxicity ...... 15

2.3.2 Speciation in Aqueous Solution ...... 16

2.3.3. Corrosion Inhibition of Ferrous Alloys ...... 17

2.3.4 Corrosion Inhibition of Non-Ferrous Alloys ...... 18

2.3.5 Molybdate Conversion Coatings, Pigments, and Synergisms ...... 19

2.4 Silicates ...... 20

2.4.1 Background and Toxicity ...... 21

2.4.2 Aqueous Chemistry ...... 22 ix

2.4.3 Corrosion Inhibition in Ferrous and Non-Ferrous Alloys ...... 22

2.4.4. Organically Modified Silicate Coatings and Synergisms ...... 23

2.5 Rare-Earth-Metal Compounds ...... 25

2.5.1 Background and Toxicity ...... 25

2.5.2 Inhibition by Rare-Earth Salts and Synergisms ...... 26

2.5.3 Mechanism of Inhibition ...... 27

2.5.4 Rare-Earth Conversion Coatings on Aluminum Alloys ...... 29

2.6 Unresolved Issues ...... 30

CHAPTER 3: CORROSION INHIBITION OF AA2024-T3 BY SODIUM MOLYBDATE ...... 43 3.1 Introduction ...... 43

3.2 Experimental ...... 46

3.2.1 Materials and sample preparation: ...... 46

3.2.2 Chronoamperometry and potentiodynamic polarization curves: ...... 47

3.2.3 Electrochemical Impedance Spectroscopy: ...... 47

3.2.4 Free-corrosion experiments coupled with secondary electron microscopy:.... 48

3.2.5 X-ray photoelectron spectroscopy: ...... 48

3.2.6 Raman spectroscopy: ...... 48

3.2.7 Atomic force microscopy: ...... 49

3.3 Results and Discussion ...... 49

3.3.1 Aqueous solution chemistry: ...... 49

3.3.2 Polarization in aerated sodium chloride solution: ...... 50

3.3.3 Free corrosion morphology of AA2024-T3 after exposure: ...... 51

3.3.4 Chronoamperometry ...... 53

3.3.5 Raman Spectroscopy: ...... 58

3.3.6 Evaluation of inhibition in deaerated electrolyte ...... 59

3.3.7 Electrochemical Impedance Spectroscopy ...... 61

3.3.8 Mechanism of Inhibition ...... 63

3.3.9 In Situ Atomic Force Microscopy ...... 69

3.4 Conclusions ...... 73

CHAPTER 4: CORROSION INHIBITION OF AA2024-T3 BY SODIUM SILICATE ...... 106 4.1 Introduction ...... 106

4.2 Experimental ...... 108

4.2.1 Materials and sample preparation: ...... 108

4.2.2 Potentiodynamic polarization curves and electrochemical impedance spectroscopy: ...... 109

4.2.3 Free-corrosion experiments coupled with scanning electron microscopy:.... 109

4.2.5 X-ray photoelectron spectroscopy: ...... 110

4.2.5 Atomic force microscopy ...... 110

4.3 Results and Discussion ...... 111

4.3.1 Aqueous solution chemistry: ...... 111

4.3.2 Polarization in aerated sodium chloride solution: ...... 112

4.3.3 Electrochemical Impedance Spectroscopy: ...... 114

4.3.4 Morphology of AA2024-T3 after free corrosion exposure: ...... 116

4.3.5 X-ray photoelectron spectroscopy: ...... 117

4.3.6 Effect of Solution pH: ...... 119

4.3.7 Mechanism of Inhibition:...... 122

4.3.8 Synergistic inhibition of silicate and molybdate ion species: ...... 127

4.3.9 In Situ Atomic Force Microscopy:...... 131

4.4 Conclusions ...... 134

CHAPTER 5: CORROSION INHIBITON OF AA2024-T3 BY PRASEODYMIUM CHLORIDE ...... 170 5.1 Introduction ...... 170

5.2 Experimental ...... 172

5.2.1 Materials and sample preparation: ...... 172

5.2.2 Potentiodynamic polarization curves and electrochemical impedance spectroscopy: ...... 173

5.2.3 Free-corrosion experiments coupled with secondary electron microscopy:.. 174

5.2.4 Atomic force microscopy: ...... 174

5.2.5 X-ray photoelectron spectroscopy: ...... 175

5.3 Results and Discussion ...... 176

5.3.1 Aqueous solution chemistry: ...... 176

5.3.2 Polarization in aerated sodium chloride solution: ...... 177

5.3.3 Morphology of AA2024-T3 after free-corrosion exposure: ...... 178

5.3.4 X-ray photoelectron spectroscopy: ...... 181

5.3.5 Long term free-corrosion exposure: ...... 183

5.3.6 Evaluation of inhibition in decarbonated and deaerated electrolyte: ...... 184

5.3.7 Effect of Solution pH: ...... 189

5.3.8 Mechanism of Inhibition:...... 190

3.4.9 In Situ Atomic Force Microscopy:...... 192

5.4 Conclusions ...... 197

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CHAPTER 6: CONCLUSIONS AND RECOMMENDATIONS FOR ...... 226 FUTURE WORK ...... 226 6.1 Conclusions ...... 226

6.2 Implications of Research Findings ...... 232

6.3 Recommendations for Future Work ...... 234

BIBLIOGRAPHY ...... 238

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LIST OF TABLES

Table 4.1. Chemical quantification analysis obtained from high-resolution XPS spectra after 1 day exposure in 0.1 M NaCl solution with and without 25 m Na2SiO3...... 146

Table 4.2. Chemical quantification analysis obtained from high-resolution XPS spectra after 3 day exposure in pH adjusted 0.1 M NaCl solution with 25 m Na2SiO3...... 153

Table 4.3. Chemical quantification analysis obtained from high-resolution XPS spectra after 10 days of exposure in 0.1 M NaCl solution with 1 m Na2SiO3 + 1 mM Na2MoO4.

...... 158

Table 5.1. Chemical quantification analysis obtained from high-resolution XPS spectra after 2 day exposure in different environments...... 210

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LIST OF FIGURES

Figure 2.1. Schematic representation of silicate anion deposition of aluminum oxide surface...... 33

Figure 3.1. Schematic representation of chronoamperometry cell configuration (WE – working electrode CE – counter electrode RE – reference electrode). Solution was continuously bubbled with air to reduce convection effects after injection...... 75

Figure 3. 2. Chemical equilibrium diagram for (a) 10 µM, (b) 1 mM, and (b) 100 mM

2- TM MoO4 in 0.1 M NaCl solution. Specie diagram generated using Medusa software.

The dashed black line signifies maximum solubility...... 76

Figure 3.3. Naturally aerated polarization curves for AA2024-T3 in 0.1 M NaCl solution at varying Na2MoO4 concentrations...... 79

Figure 3.4. Open-circuit potential Eocp and pitting potential Epit with incremental amounts of Na2MoO3...... 80

Figure 3.5. Open-circuit potential macrographs after 1 day exposure to (a) NaCl-only solution and (b-h) 0.1 M NaCl + 125 mM Na2MoO4 solution at varying time...... 81

Figure 3.6. Optical micrographs after 1 d exposure to (a) NaCl-only solution and (b) 0.1

M NaCl + 125 mM Na2MoO4 solution at open-circuit potential...... 82

Figure 3.7. SEM-EDS analysis showing maps for Al, Cu, Mo, and O after 30 min exposure in 0.1 M NaCl solution with 125 mM Na2MoO4...... 83

Figure 3.8. SEM-EDS analysis of S-phase particle after 2 day exposure in 0.1 M NaCl +

125 mM Na2MoO4. (a) Secondary electron micrograph. Red line is reference for (b)

EDS line profiles...... 84

Figure 3. 9. SEM-EDS analysis of Fe-containing particle after 2 day exposure in 0.1 M

NaCl + 125 mM Na2MoO4. (a) Secondary electron micrograph. Red line is reference for (b) EDS line profiles...... 85

Figure 3.10. Mo 3d spectra after 1 day exposure in naturally aerated solution with 0.1 M

NaCl + 125 mM Na2MoO4...... 86

Figure 3.11. Cathodic chronoamperometry of AA2024-T3 in air-bubbled 0.1 M NaCl. 87

Figure 3.12. Electron micrograph of surface after 2 h exposure to 0.1 M NaCl + 125 mM

Na2MoO4 solution at (a) -900 mV SCE- and (b) -1100 mV SCE fixed potential. (c) EDS spectrum of sample surface after -1100 mV SCE fixed potential...... 88

Figure 3.13. Anodic chronoamperometry of AA2024-T3 in air-bubbled 0.1 M NaCl. .. 89

Figure 3.14. (a-b) Optical micrographs and (c) SEM-EDS analysis of substrate after

Na2MoO4 injection at -535 mV SCE. There is evidence of oxide formation inside corroded areas as suggested by Mo and O maps...... 90

Figure 3.15. Mo 3d spectra of sample surface immediately after injection of concentrated

Na2MoO4 to cell solution...... 91

Figure 3.16. Raman spectra at different microstructural features of a sample after 2 days of exposure in naturally aerated 0.1 M NaCl solution with 0.125 M Na2MoO4...... 92

xvi

Figure 3.17. Cyclic polarization curves in deaerated 0.1 M NaCl solution with incremental amounts of Na2MoO4...... 93

Figure 3.18. Cathodic polarization curves of AA2024-T3 in aerated and deaerated 0.1 M

NaCl solution with 125 mM molybdate...... 94

Figure 3.19. Sample macrograph and SEM-EDS analysis after 1 day exposure in deaerated 0.1 M NaCl + 125 mM Na2MoO4 solution. Substrate was partially exposed indicated by dashed line...... 95

Figure 3.20. Mo 3d spectra of sample surface after 1 day exposure in deaerated 0.1 M

NaCl + 125 mM Na2MoO4 solution...... 96

Figure 3.21. Bode magnitude and phase angle plot of 2024-T3 coupons immersed in 0.1

M NaCl with and without 125 mM Na2MoO4 in (a) aerated and (b) deaerated solution. 97

-5 Figure 3.22. In situ AFM scratching in 0.1 M NaCl with 10 M Na2MoO4.

Scan size=70µm. Top left images is Volta potential map before exposure with z range

=500 mV. The other images are topographic maps with z range=200 nm...... 98

-3 Figure 3.23. In situ AFM scratching in 0.1 M NaCl with 10 M Na2MoO4.

Scan size=65µm. Top left images is Volta potential map before exposure with z range

=500 mV. The other images are topographic maps with z range=200 nm...... 99

Figure 3.24. In situ AFM scratching in 0.1 M NaCl with 0.1 M Na2MoO4.

Scan size=35µm. Top left images is Volta potential map before exposure with z range

=500 mV. The other images are topographic maps with z range=300 nm...... 100

Figure 3.25. Ex situ SEM-EDS analysis after AFM scratching in 0.1 M NaCl with 0.1 M

Na2MoO4...... 101

xvii

Figure 4.1. Effect of silicate concentration on solution pH...... 137

Figure 4.2. Chemical equilibrium diagram for (a) 1 mM and (b) 100 mM Si(OH)4 in 0.1

M NaCl solution. Speciation diagram generated using MedusaTM software. The dashed black line signifies maximum solubility and ‘cr’ in parenthesis denotes crystalline species...... 138

Figure 4.3. Naturally aerated polarization curves for AA2024-T3 in 0.1 M NaCl solution at varying Na2SiO3 concentrations...... 139

Figure 4.4. Corrosion potential Ecorr and pitting potential Epit with incremental amounts of Na2SiO3...... 140

Figure 4.5. Bode magnitude and phase angle plots of 2023-T3 coupons immersed in (a)

0.1 M NaCl (pH 6 – natural pH) (b) 0.1 M NaCl (pH 12.8 – pH adjusted) and (c) 0.1 M

NaCl + 25 mM Na2SiO3 (pH 12.8 – natural pH)...... 141

Figure 4.6. Impedance data summary of (a) total resistance Rt and (b) film capacitance Cf with time...... 143

Figure 4.7. Images of samples after 1 day exposure in 0.1 M NaCl solution with and without 25 mM Na2SiO3 (a) open-circuit potential macrographs and (b) optical micrographs...... 144

Figure 4.8. (a) Secondary electron micrographs before and after 3 day exposure in 0.1 M

NaCl solution with 25 mM Na2SiO3, and EDS spectrum of S-phase particle after exposure (b) EDS mapping of selected area after exposure...... 145

xviii

Figure 4.9. Figure 4.9. Spectra showing the (a) Si 2p (b) Al 2p (c) Na 1s and (d) Mg 1s peaks after 1 day exposure in A – 0.1 M NaCl and B – 0.1 M NaCl + 25 mM Na2SiO3 solution...... 147

Figure 4.10. Sputter depth profile after 1 day immersion in 0.1 M NaCl + 25 mM

Na2SiO3...... 148

Figure 4.11. Naturally aerated polarization curves in 0.1 M NaCl at varying Na2SiO3 concentrations in pH adjusted (a) 7.3 and (b) 4 solutions...... 149

Figure 4.12. (a) Open-circuit potential macrograph of coupon exposed to 0.1 M NaCl +

25 mM Na2SiO3 solution at pH 7.3 (b) Secondary electron image of surface (c)

Secondary electron image of partially covered large intermetallic particle with accompanying EDS spectrum of particulates. Evidence of small amount of Si detected over particulate covered area...... 150

Figure 4.13. (a) Open-circuit potential macrograph of coupon exposed to 0.1 M NaCl +

25 mM Na2SiO3 solution at pH 4 (b) Secondary electron image of site A (refer to macrograph) (c) Backscatter electron image of site A with accompanying EDS spectrum of labeled S-phase particle (d) Secondary electron image of site B with accompanying

EDS spectrum of film...... 151

Figure 4.14. Spectra showing the (a) Si 2p and (b) Al 2p peaks after 3 day exposure in pH 7.30 and pH 4 solution containing 0.1 M NaCl + 25 mM Na2SiO3...... 154

Figure 4.15. Stability diagram of a ternary alloy system consisting of Al, Cu, and Mg in silicate solution (generated using OLI AnalyzerTM)...... 155

xix

Figure 4.16. Naturally aerated polarization curves in 0.1 M NaCl at varying Na2MoO4

(compound A) and Na2SiO3 (compound B) concentrations mixed together in solution. 156

Figure 4.17. (a) Open-circuit potential macrograph of coupon exposed to 0.1 M NaCl

-3 -3 solution with 10 M Na2SiO3 + 10 M Na2SiO3 for 10 days (b) Secondary electron image of surface with elementals maps of scanned area (c) Secondary electron image with EDS spectrum of large intermetallic particle covered with Si particulates...... 157

Figure 4.18. Spectra showing the (a) Si 2p (b) Mo 3d (c) Al 2p and (d) O 1s peaks after

10 days of exposure in 0.1 M NaCl solution with 1 mM Na2SiO3 + 1 mM Na2MoO4. . 159

Figure 4.19. OCP measurements in 0.1 M NaCl solution with and without 10-3 M

-3 Na2SiO3 + 10 M Na2MoO4...... 160

-5 Figure 4.20. In situ AFM scratching in 0.1 M NaCl + 10 M Na2SiO3.

Scan size = 70 µm. Top left image is Volta Potential map before exposure with z range =

300 mV. The other images are topographic maps with z range = 300 nm. The first is the map prior to exposure and the rest are in situ images during rastering at indicated times and setpoint voltages...... 161

Figure 4.21. SKPFM and SEM analysis of the selected area before in situ AFM scratching in 0.1 M NaCl + 0.025 M Na2SiO3...... 162

Figure 4.22. In situ AFM scratching in 0.1 M NaCl with 0.025 M Na2SiO3.

Scan size=30µm. Z range=300nm...... 163

Figure 4.23. In situ AFM scratching in 0.1 M NaCl + 0.1 M Na2SiO3.

Scan size = 75 µm. Top left image is Volta Potential map before exposure with z range =

500 mV. The other images are topographic maps with z range = 300 nm. The first is the

xx map prior to exposure and the rest are in situ images during rastering at indicated times and setpoint voltages...... 164

Figure 4.24. Ex situ SEM-EDS analysis after AFM scratching in 0.1 M NaCl with 0.1 M

Na2SiO3...... 165

Figure 5.1. Chemical equilibrium diagram for 1 mM Pr3+ in 0.1 M NaCl solution in the

(a) absence and (b) presence of atmospheric CO2. Specie diagram generated using

MedusaTM software. The dashed black line signifies maximum solubility...... 199

TM Figure 5.2. Effect of PrCl3 concentration on solution pH. Generated with Medusa . 200

Figure 5.3. Naturally aerated polarization curves of AA2024-T3 in 0.1 M NaCl solution at varying PrCl3 concentrations...... 201

Figure 5.4. (a) Open-circuit potential macrographs and (b) optical micrographs of samples after 1 day exposure in 0.1 M NaCl solution with and without 0.2 mM PrCl3. 202

Figure 5.5. Secondary electron micrographs of scanned area containing S-phase particles before and after 2 days exposure in 0.1 M NaCl solution with 0.2 mM PrCl3...... 203

Figure 5.6. (a) Volta-potential and (b) topography images of S-phase intermetallic particles before and after 5 hours of exposure to 0.2 mM PrCl3 in 0.1 M NaCl containing solution. Scan size = 30 x 30 µm2. Z range = 550 nm (topography); 500 mV (Volta- potential map)...... 204

Figure 5.7. (a) Topography and (b) Volta-potential images of Fe-containing intermetallic particles before and after 5 hours exposure to 0.2 mM PrCl3 in 0.1 M NaCl containing

xxi solution. Scan size = 65 x 65 µm2. Z range = 500 nm (topography); 500 mV (Volta- potential map)...... 205

Figure 5.8. Spectra showing the (a) Pr 3d5/2 and (b) Pr 4p3/2 peaks after 2 days of exposure in A – 0.1 M NaCl and B – 0.1 M NaCl + 0.2 mM PrCl3 solution...... 206

Figure 5.9. Spectra showing the (a) O 1s and (b) C 1s peaks after 2 days of exposure in

0.1 M NaCl + 0.2 mM PrCl3 solution...... 207

Figure 5.10. (a) Macrograph of sample after 30 days immersion in 0.1 M NaCl solution with PrCl3 (b) Secondary electron images of surface after exposure and (c) XRD spectrum of surface film...... 208

Figure 5.11. (a) Open-circuit potential macrographs and (b) optical micrographs of partially exposed samples after 2 day immersion in 0.1 M NaCl + 0.2 mM PrCl3 under decarbonated and deaerated solutions. Dashed lines signify exposed area...... 209

Figure 5.12. Spectra showing the Pr 3d5/2 and O 1s peaks after 2 days of exposure in (a) decarbonated and (b) deaerated 0.1 M NaCl solution with 0.2 mM PrCl3...... 211

Figure 5.13. Spectra showing the Pr 4p3/2 peaks after 2 days of exposure in A – decarbonated and B – deaerated 0.1 M NaCl solution with 0.2 mM PrCl3...... 212

Figure 5.14. (a) Bode magnitude and (b) phase angle plots of 2024-T3 coupons immersed in 0.1 M NaCl with and without PrCl3 in naturally aerated and decarbonated solution...... 213

Figure 5.15. Bode magnitude and phase angle plot of 2024-T3 coupons immersed in deaerated 0.1 M NaCl with and without 0.2 mM PrCl3...... 214

xxii

Figure 5.16. Naturally aerated polarization curves in 0.1 M NaCl with 0.2 mM PrCl3 in

(a) pH 3 and (b) pH 10 solutions...... 215

Figure 5.17. SEM-EDS analysis after 3 days of exposure in 0.1 M NaCl with 0.2 mM

PrCl3 in pH 3 solution...... 216

Figure 5.18. SEM-EDS analysis after 3 days of exposure in 0.1 M NaCl with 0.2 mM

PrCl3 in pH 10 solution...... 217

-5 Figure 5.19. In situ AFM scratching in 0.1 M NaCl with 10 M PrCl3.

Scan size = 20 µm. Top left image is Volta Potential map before exposure with z range =

500 mV. The other images are topographic maps with z range = 200 nm. The first is the map prior to exposure and the rest are in situ images during rastering at indicated times and setpoint voltages...... 218

Figure 5.20. Ex situ SEM-EDS analysis after AFM scratching in 0.1 M NaCl with 10-5 M

PrCl3...... 219

Figure 5.21. SKPFM and SEM analysis of the selected area before in situ AFM scratching in 0.1 M NaCl + 0.2 mM PrCl3...... 220

Figure 5.22. In situ AFM scratching in 0.1 M NaCl with 0.2 mM PrCl3.

Scan size=35µm. Z range=500 nm...... 221

Figure 5.23. Ex situ SEM-EDS analysis after AFM scratching in 0.1 M NaCl with 0.2 mM PrCl3...... 222

xxiii

CHAPTER 1: INTRODUCTION

Corrosion has large societal implications with respect to cost and safety in today’s modern world. Its effects are felt throughout many engineering fields including the aerospace industry where the cumulative action of aggressive environments and mechanical stresses can lead to localized corrosion of aluminum alloys 1-4. Therefore, it is imperative to develop protection schemes that improve the service life of aircraft and mitigate failures associated with corrosion.

Multi-layered coating protection schemes involving conversion coatings and pigmented coatings have been primarily used to mitigate corrosion of Al alloys employed in aerospace aircraft. These coating systems release soluble inhibiting species into the local environment and provide active corrosion inhibition to the underlying metal substrate by suppressing either the anodic or cathodic reaction, or both. Traditionally, chromate-based pigments and conversion coatings have been used to prevent the corrosion of high-strength aluminum alloys 5, 6. They effectively reduce the rate of oxygen reduction over cathodes and moderately hinder the anodic dissolution kinetics 7, 8.

Furthermore, chromate conversion coatings have the distinct ability to release soluble

Cr6+ into the local environment at a defect providing a “self-healing” characteristic to various systems 9-11. However, the carcinogenic effect of soluble hexavalent chromate has forced governments to impart regulations that restrict its use in a wide-range of applications. Therefore, efforts are underway to find environmentally-friendly 1 alternatives that offer the same level of protection and reliability as chromate protection schemes 6.

The pursuit to find a chromate-alternative has led to the development of several chromate-free aerospace primers and coating systems that offer good protection. For instance, a three-step production of a conversion coating on aircraft aluminum alloys has met the US military standard for salt-fog tests 12. However, fundamental understanding of the functionality of the chromate-free pigments that are embedded within these coating systems is lacking. Therefore, the objective of this dissertation was to investigate how selected chromate-free inhibitors impart corrosion inhibition on aerospace aluminum alloy 2024-T3, with the intent of developing the kind of understanding that was accomplished with chromate. Advances in the scientific understanding of these issues will help facilitate improvements to many non-chromate technologies.

The inhibitors selected for this study are soluble reagent grade forms of

2- 2- 3+ molybdate (MoO4 ), silicate (SiO3 ), and praseodymium (Pr ) compounds. These inhibitors release during field exposures by leaching from current chromate-free conversion coatings and pigments, thereby providing good protection to the underlying substrate. However, an understanding of the fundamental inhibition mechanism provided by these non-chromate inhibitors is lacking and thus warrants further studies.

Furthermore, experience has shown that the same approaches that were used to study chromate can be extended to investigations concerning the inhibition mechanism of non- chromate inhibitors 13-15. Therefore, various experimental techniques that have been used to identify the corrosion inhibition by chromate were employed in this work, including

2 standard electrochemical measurements, chronoamperometry, microscopic and spectroscopic techniques such as x-ray photoelectron spectroscopy and Raman spectroscopy, and atomic force microscopy.

This dissertation includes six chapters. Chapter 2 is a review that discusses the mechanism of inhibition of select chromate-free inhibitors as described in the literature and pertaining to aluminum alloy 2024-T3. In addition, important aspects relating to the alloy metallurgy and its susceptibility to localized corrosion, chromate inhibition, and the background and toxicity of the select non-chromate inhibitors is discussed in the review.

Chapters 3 through 5 present the investigations on the corrosion inhibition mechanism of the selected non-chromate inhibitors. Mechanisms are proposed for under various exposure conditions and are substantiated by results pertaining to aqueous chemistry, performance, and surface analytical techniques. They are written as stand-alone papers, and will be submitted individually for publication. Finally, Chapter 6 summarizes the technical findings of this dissertation and makes conclusions regarding the corrosion inhibition mechanism of the aforementioned inhibitors. Future work has been suggested to advance the understanding of non-chromate inhibitors and to help in the development of feasible alternatives that offer the same level of protection and reliably as chromate protection schemes.

REFERENCES

1. J. R. Galvele and Demichel.Sm, Corros Sci 10, 795-& (1970). 2. I. L. Muller and J. R. Galvele, Corros Sci 17, 179-& (1977). 3. G. S. Chen, M. Gao and R. P. Wei, Corrosion 52, 8-15 (1996). 4. C. M. Liao and R. P. Wei, Electrochim. Acta 45, 881-888 (1999). 5. L. Xia, E. Akiyama, G. Frankel and R. McCreery, Journal of the Electrochemical Society 147, 2556-2562 (2000). 6. M. W. Kendig and R. G. Buchheit, Corrosion 59, 379-400 (2003). 7. G. O. Ilevbare and J. R. Scully, Corrosion 57, 480-480 (2001). 8. G. O. Ilevbare, J. R. Scully, J. Yuan and R. G. Kelly, Corrosion 56, 227-242 (2000). 9. W. J. Clark and R. L. McCreery, Journal of the Electrochemical Society 149, B379-B386 (2002). 10. J. Zhao, L. Xia, A. Sehgal, D. Lu, R. L. McCreery and G. S. Frankel, Surf Coat Tech 140, 51-57 (2001). 11. J. Zhao, G. Frankel and R. L. McCreery, Journal of the Electrochemical Society 145, 2258-2264 (1998). 12. R. N. Miller, Vol. 5221371, US, 1993. 13. K. D. Ralston, T. L. Young and R. G. Buchheit, Journal of the Electrochemical Society 156, C135-C146 (2009). 14. M. Iannuzzi and G. S. Frankel, Corrosion Science 49, 2371-2391 (2007). 15. M. Iannuzzi and G. S. Frankel, Corrosion 63, 672-688 (2007).

CHAPTER 2: LITERATURE REVIEW

Aluminum alloys are used extensively in aerospace applications due to their high strength-to-weight ratio 1. Their improved mechanical properties are attributed to the addition of alloying elements that precipitate into fine particles during age-hardening after a solid-solution heat treatment. However, during the solidification process, coarse constituent particles, or intermetallic compounds, also segregate throughout the alloy with a resulting heterogeneous microstructure that is susceptible to localized corrosion 2.

This is accomplished by discussing the basic metallurgy and microstructure of the alloy along with its susceptibility to localized corrosion. In addition, the mechanism by which chromate-based protection schemes inhibit corrosion is described, and a detailed review of the inhibition provided by several selected non-chromate inhibitors is discussed.

2.1 Corrosion of AA2024-T3

The susceptibility of AA2024-T3 (nominal composition 3.9-4.9% Cu, 1.2-1.8%

Mg, 0.3-0.9% Mn, 0.5% Fe, 0.5% Si, balance Al) 2 to localized corrosion is generally attributed to the heterogeneity of the surface microstructure. The ‘T3’ designation indicates that the alloy was solution-annealed, quenched, and naturally aged to a stable condition. The microstructure consists of a solid solution matrix that contains a variety of intermetallic and second phase particles which can be classified into three different categories 2:

1.) Hardening Precipitates. These strengthening particles are typically less than 0.01

µm in size and form as a result of precipitation during the natural aging process.

They initially form at grain boundaries before precipitating throughout the grains,

and typically consist of a mixture of S-phase (Al2MgCu) and theta phase (Al2Cu).

A precipitate-free zone (PFZ) can form adjacent to the grain boundaries which,

along with the presence of second phase particles, often make the grain boundary

region of AA2024-T3 susceptible to intergranular corrosion 2, 3.

2.) Constituent particles or inclusions: These particles are greater than 1 µm in size

and are present in commercial alloys due to residual elements like Fe and Si 4.

They form as a result of eutectic decomposition during ingot solidification and do

not re-dissolve during subsequent sheet processing. In addition, because

aluminum is highly electronegative and trivalent, most of these constituent

particles are in the form of intermetallic compounds. Although these particles do

not enhance the mechanical properties of the alloy, they play an important role in

5-10 the corrosion process . Al2Cu2Fe, Mg2Si, FeMnAl6 and Al2Cu(Mn,Fe)3 are

typical examples of constituent particles found in AA2024-T3 2, 4, 11. In the

following discussion constituent particles will be referred to as Fe-containing

particles. It should be noted that larger S-phase particles ranging in size from 1-

10 µm have been observed using electron microscopy techniques, and are

considered to belong to this class of intermetallic particles 12. Although these

larger particles do not contribute to the mechanical strength, they do play a major

role in the corrosion susceptibility of AA2024-T3.

3.) Dispersoid particles: These are 0.01-0.1 µm in size and exhibit higher melting

points than the aluminum alloy like the constituent particles. As a result, they

precipitate out of the liquid phase and are unaffected by subsequent heat treatment

2. Furthermore, due to their small size, homogeneous dispersion, and comparative

electrochemical inertness, dispersoids do not have a significant effect on the

corrosion processes of high-strength aluminum alloys 11. Typical examples of

2, 4, 11 dispersoids found in AA2024-T3 are Al3Ti, Al6Mn and Al20Cu2Mn3 .

The susceptibility of AA2024-T3 to localized corrosion is generally attributed to the presence of local galvanic cells that form between the precipitated Cu-rich particles and the Al matrix 8, 11, 13-15. However, the formation of galvanic cells on the alloy surface is not sufficient to explain the corrosion behavior of aluminum alloys. In addition to the heterogeneous microstructure, the composition of the passive film present over the intermetallic and S-phase particles varies in comparison to the passive film formed over 7 the Al matrix 16. The difference is associated with impurities and/or defects within the oxide that alter the electrochemical behavior of the particle. Moreover, these changes in behavior are reflected by differences in the open-circuit potential and pitting potential, which are themselves reflections of the catalytic reactions that occur over the particles.

Thus, it is the formation of galvanic cells on the alloy surface and the rate of oxygen reduction over the intermetallic particles that describe the mechanism for overall damage accumulation 11, 13-15, 17.

Electrochemical characterization of several artificial intermetallic particles revealed that the larger Fe-containing particles are intrinsically nobler than the matrix 11,

18. Polarization curves performed by Birbilis and Buchheit showed that these particles have a higher open-circuit potential (also concomitant with a higher pitting potential) with respect to the Al matrix and other intermetallic particles 11. Similarly, LeBlanc and

Frankel showed that Fe-containing particles have a nobler Volta-potential than the Al matrix and interacted strongly with their surroundings 19. They act as local cathodes on the surface supporting the oxygen reduction reaction. In acidic (Eq. 1) and neutral or alkaline solution (Eq. 2), the oxygen reduction reaction is described by the following equations, respectively:

O Eq. 1

Eq. 2

Thus, the matrix adjacent to the particle corrodes owing to the polarization provided by the nobler Fe-containing intermetallic particles and/or localized increase in pH associated with hydroxyl ion formation 19-23.

The predominant type of intermetallic particle found in the structure of AA2024-

8 T3 is the S-phase precipitate with Al2CuMg composition . Electrochemical characterization revealed that the S-phase particles have corrosion potentials negative to that of the matrix 11, 14, 15. However, an altered surface film that is different in composition and microstructure to the bulk material is formed over these particles during polishing 15. Therefore, anodic dissolution of the S-phase is delayed until the altered surface film is removed during exposure to aggressive solution. Once activated, S phase particles undergo preferential dissolution of Mg and Al 8. This leads to de-alloying and the formation of nobler Cu-rich remnants that promote the anodic dissolution of the adjacent matrix 20. Furthermore, the de-alloyed Cu remnants release metallic Cu into the solution that can oxidize and subsequently reduce on the nearby surface of the alloy 8.

The resulting Cu-enrichment of the surface increases the effective area for the oxygen reduction reaction and increases the susceptibility to localized corrosion 8, 24. Therefore, an effective way to increase the corrosion resistance of AA2024-T3 is to suppress the dissolution of S-phase particles 25.

2.2 Corrosion Inhibition by Chromate-Protection Schemes

An organic coating is typically applied to a metal substrate to form a protective barrier against corrosion. However, an organic coating alone is not sufficient to protect the underlying metal substrate due to the presence features within the coating (i.e. micro- pores, low cross-link density, high pigment volume concentration, etc.) that provide pathways for corrosive agents like water, oxygen, and chloride ions to diffuse to the

9 metal/coating interface 26. Therefore, it is often necessary to incorporate inorganic or organic inhibitors that leach from the coating in the presence of water and protect the underlying substrate. In this regard, chromate has been the most effective inhibitor for ferrous and non-ferrous metals in a wide-range of applications.

2.2.1 Chromate Background and Toxicity

Chromate has been used to effectively protect various metals since the early

1900s 26. It has been introduced in coating systems using various strategies including pigmentation (typically consisting of strontium chromate or zinc chromate, i.e. SrCrO4,

H2CrO4Zn), chromate conversion coatings, and chromic acid anodization of the alloy substrate. The latter technique involves the electrochemical growth of an oxide bi-layer consisting of a thick porous layer on top of a thinner non-porous film 27. In the final stages of anodization, chromates are introduced by sealing the porous layer with chromic acid (H2CrO4). Despite the superior corrosion protection afforded by this method, chromate conversion coatings (CCC) are preferentially used because of their ease of application and lower cost. A CCC is a chemically grown oxide layer on the metal surface that provides an active barrier against corrosion. Additionally, chromates are commonly used as pigments in organic coatings in conjunction with a CCC 27.

Despite the strong protection afforded by chromate protection-schemes, the use of chromates has been limited due to their carcinogenic effect 26, 28, 29. It appears that the same properties that make chromate a strong corrosion inhibitor also make it environmentally unsafe. According to toxicology studies, the most common ailment is

10 lung cancer 30, with the primary form of exposure being ingestion and inhalation.

Although chromate per se is not responsible for the DNA damage that leads to cancer, the molecular debris associated with the reduction of chromate species induces critical changes in the DNA 30, 31.

The carcinogenic effects associated with chromate have led to the development of chromate-free protection schemes 26, 29. However, in order to formulate new technologies that are suitable to replace chromate, it is necessary to understand the mechanism by which chromates provide corrosion protection. The following section will review some of the most important aspects of chromate inhibition.

2.2.2 Chromate Inhibition

Inhibitors can be classified as anodic or cathodic depending on which reaction they more strongly reduce. In this regard, chromate is considered to be a powerful cathodic inhibitor, even at high chloride-to-chromate concentration ratios 32. For

-5 2- instance, the addition of 10 M Cr2O7 to oxygen-bubbled 1 M NaCl solution decreased the oxygen reduction kinetics by an order of magnitude. According to Clark et al. 33, the inhibition mechanism involves reduction of Cr6+ to Cr3+ and the subsequent formation of an adherent hydroxide layer on the metal surface of near-monolayer thickness. The layer provides a barrier for electron transfer and effectively hinders further Cr6+ reduction and oxygen reduction kinetics. On high-strength Al alloys, Cr6+ rapidly adsorbs at the Cu- rich intermetallic particles and subsequently reduces to Cr3+ forming an irreversible barrier 28, 33, 34. The adsorption of Cr6+ at the cathodes prevents the adsorption of oxygen molecules, thereby hindering the catalytic step for oxygen reduction. 11

Anodic inhibition of Al alloys by soluble chromate has been thoroughly studied 18,

32, 35. The extent of anodic inhibition seems to be dependent on a number of environmental factors such as the chloride-to-chromate concentration ratio, pH, degree of aeration, and on the alloy chemistry and temper. An increase in the pitting potential, which can be considered to be a measure of stable pit growth susceptibility, was observed at chloride-to-chromate ratios below 10:1 to 1:1 36. Although the exact mechanism by which chromate imparts anodic inhibition is still not understood, it is suggested that

2- chromate anions (CrO4 ) in solution migrate to a pit formed in the passive film where it is adsorbed and forms a mixed Al3+/Cr6+ oxide 37. Modeling has also suggested that the adsorption of chromate anions neutralizes the charge on the surface and thereby prevents chloride ion adsorption 38. The anodic inhibition effect by chromate is further supported by its ability to decrease the metastable pit nucleation and pit growth rates 28, 39.

2.2.3 Chromate Conversion Coatings and Pigments

The advent of CCCs can be traced to the early part of the twentieth century 26.

Today, these coatings have shown exceptional versatility in a wide range of applications and effectively inhibit corrosion by reducing the rate of the oxygen reaction 26, 28.

Furthermore, CCCs have the distinct ability to ‘self-heal’ by releasing soluble Cr6+ into the local aqueous environment at the presence of a chemical or mechanical defect 40-42.

This self-healing effect by CCCs was clearly demonstrated by Zhao et al. using the artificial scratch cell technique 40. In these experiments, Raman spectroscopy was used to monitor the migration of chromate species from a CCC to an initially untreated Al alloy sample. After several hours, Cr6+ leached from the CCC to the fresh alloy surface. The 12 initially untreated alloy became more corrosion resistant with time which was made evident by a two-order increase in the polarization resistance and a 60 mV increase in the pitting potential. It was determined that the self-healing mechanism involves liberation of Cr6+ from the CCC and migration of the Cr6+ to an incipient defect. This is followed by reduction of the leached Cr6+ to insoluble Cr3+ hydroxide, or interaction of Cr6+ with previously formed corrosion product to stifle further corrosion 40, 43.

Chromate conversion coatings are formed within minutes by spray or full immersion, and can be as thick as 3 µm 44. Modern coating baths contain activators and accelerators that help to expedite the formation process and increase film thickness 44, 45.

The first stage of film formation, after removal of the surface oxide, involves the simultaneous reactions of Al oxidation and reduction of Cr6+ to Cr3+ at the metal surface

37. In addition, the local pH at the metal-electrolyte interface increases due to the hydrogen evolution reaction on the activated Al surface. This leads to coating formation based on a sol-gel mechanism, involving hydrolysis of Cr3+ and the formation of a chromium-hydroxide polymer “backbone” composed of edge- and corner-sharing Cr3+ octahedral units 28, 33, 46, 47. As this process continues, labile Cr6+ ions reversibly attach to the insoluble chromium backbone to form Cr3+-O- Cr6+ bonds, which are readily detected by Raman spectroscopy 6, 46, 47. Furthermore, the adsorption of Cr6+ from solution to the

Cr3+ oxide/hydroxide can be described by the following reaction 43:

– ( ) ( ) – – ( ) Eq. 3

According to Eq. 3, the adsorption and release of Cr6+ is pH dependent, with the latter occurring in alkaline conditions 43. Thus, the local increase in pH associated with oxygen

13 reduction near a defect site will trigger Cr6+ release from the coating, allowing the desorbed Cr6+ to migrate to the actively corroding site and stop corrosion following a similar mechanism as described above.

Chromate is also implemented as a pigment in paints and primers 48. Hydration of pigmented coatings will activate inhibitor release and provide corrosion protection to the underlying substrate 49. A usable pigment must exhibit appropriate solubility and inhibition efficacy; if too soluble, the pigment will promote coating degradation via osmotic blistering 48. Thus, the sparingly soluble chromates of the alkaline earth metals

(i.e. compounds with Ca, Sr, Ba, and Zn cations) are typically used as pigments in organic coatings. They provide a source for active Cr6+ that can leach out from the primer upon contact with water and reduce over defects to form a mixed Al/Cr3+ hydroxide. However, it is important to note that the mechanism of Cr6+ release in chromate pigmented primers is significantly different from a CCC. The former will leach

Cr6+ into solution until a finite solubility is reached, while the latter is a reversible pH- dependent reaction.

Non-chromate conversion coatings and inhibiting pigments have been developed with some success by trial-and-error approaches. However, the fundamental understanding of their functionality is still lacking, thereby inhibiting further advances and creating risk associated with their use 26, 28, 48, 50. In the following sections, a detailed review of the inhibition provided by several selected non-chromate inhibitors is discussed.

2.3 Molybdates

The carcinogenic effect of soluble hexavalent chromate has led to the desire to find environmentally-friendly inhibitors that offer the same level of protection and reliability as chromate protection schemes 28. In this regard, molybdate and molybdate- based compounds have been used in the past to replace chromate and have shown promising behavior as corrosion inhibitors in various environments 51. Furthermore, recent studies have shown that molybdate provides a synergistic effect in combination with other compounds, effectively enhancing corrosion protection for both ferrous and non-ferrous alloys 51-53.

2.3.1 Background and Toxicity

Molybdenum is predominantly used in metallurgical applications, such as stainless steel and cast iron alloys 54. Its addition to steel alloys increases their strength and thermal resistance, and improves the corrosion resistance of the alloy. Furthermore, molybdenum-based compounds have been established as practical corrosion inhibitors for both ferrous and non-ferrous metals since they were introduced in 1939 for the corrosion control of automobile cooling systems 55-63.

Molybdenum-based compounds are considered to have extremely low or negligible toxicity 51. The low degree of toxicity can be attributed to the fact that molybdenum is a necessary trace element in the body and an essential micronutrient to plant life and animal health 64, 65. It has been reported that absorption of molybdenum from the diet is the most important source of molybdenum; the concentration in ambient air and natural waters is usually negligible 66. According to the World Health 15

Organization, the minimum daily requirement for molybdenum is 2 µg Mo/kg-body weight 54. However, excessive concentration of molybdenum in the diet can lead to a form of toxicity in animals that is similar in symptoms to copper deficiency.

Surprisingly, the effect of molybdenum on human beings has not been studied thoroughly. Despite the lack of information that exists on human toxicity, neither a carcinogenic nor a teratogenic effect has been demonstrated in animal experiments 54.

Therefore, the use of molybdenum compounds is considered to be less dangerous relative to chromate, since they are not suspected carcinogens.

2.3.2 Speciation in Aqueous Solution

The aqueous chemistry of molybdate has been studied in detail. It has been found that the molybdate speciation is highly dependent on the concentration and pH of the solution 67-69. In alkaline and neutral conditions, the dominant species is the monomeric

2- 2- [MoO4 ] oxy-anion. At lower concentrations, the protonation of MoO4 promotes the

- formation of HMoO4 and H2MoO4 species. With increasing acidity, the anion becomes protonated and polymerization-condensation occurs at higher Mo concentrations.

67 6- 4- According to Krishnan et al. , hepta- [Mo7O24 ] and octa- [Mo8O26 ] molybdate anions form in the pH range 3-5 and below pH 2, respectively. These are the only polymeric species that form in acidic solution. At pH 0.9, MoO3 precipitates.

Measurements performed using X-ray absorption spectroscopy confirmed no

2- complexation between MoO4 and chloride in near-neutral and alkaline solutions with chloride concentration up to 5.5 M 70. Under these conditions, molybdenum speciation is

- 2- dominated by the tetrahedral molybdate species (i.e. HMoO4 or MoO4 ). However, in 16 concentrated hydrochloric acid, molybdenum speciation is dominated by distorted

6-2m-n octahedral, oxo-chloro complexes [MoOmCln ].

2.3.3. Corrosion Inhibition of Ferrous Alloys

Molybdate is classified as an oxygen-dependent anodic inhibitor for both ferrous and non-ferrous metals 56-60. It is structurally similar to chromate; however, it is only

2- effective in moderately alkaline conditions where the MoO4 oxy-anion is stable but practically redox-inactive 48. As a consequence, molybdate does not promote the spontaneous passivation of a metal surface and thus may require the presence of an oxidant in the environment, such as dissolved oxygen 56. In the context of Fe corrosion, the inhibition mechanism involves the formation of a hydrated mixed oxide film which provides a barrier for anodic dissolution. It is generally agreed that a non-protective ferrous (Fe2+) - molybdate complex forms during the early stages of corrosion and is subsequently oxidized in the presence of oxygen 57. The resulting ferric (Fe3+) - molybdate complex is insoluble and enhances the stability of the hydrated Fe2O3 films

2- that develop over actively corroding sites. Furthermore, it has been proposed that MoO4 species adsorb on the outermost part of the hydrated oxide layer (by ion exchange mechanism), thereby imparting a negative charge on the surface 71. As a consequence, a barrier effect impedes both the ingress of Cl- to the underlying substrate and the transport of Fe2+ away from the surface. In addition, the adsorbed species can leach from the surface in response to a break in the film and retard pit growth by precipitating either

58 FeMoO4 or condensed Mo species inside the pit .

2.3.4 Corrosion Inhibition of Non-Ferrous Alloys

The impact of molybdate on the corrosion resistance of aluminum has been explored extensively in the literature 51, 55, 61-63, 72-78. Molybdate improves the pitting resistance of aluminum and its alloys by increasing the pitting potential to more noble values with respect to the open-circuit potential 73. However, the extent of inhibition is highly dependent on the ratio of molybdate-to-chloride ion concentration 77, 78. Current- time measurements of an aluminum electrode that was polarized slightly above the pitting potential showed that the current fluctuations associated with pitting decayed with incremental additions of molybdate 76. Furthermore, the addition of 15,700 ppm

Na2MoO4 to 1000 ppm chloride solution at pH 7 increased the pitting potential of pure aluminum by 1300 mV 62. However, the large ratio of molybdate-to-chloride ion concentration may not be suitable for organic coating applications. According to Meer-

Lerk et al. 79, excessive solubility of inhibitor pigments necessary for corrosion inhibition can promote coating degradation via osmotic blistering. Therefore, recent studies have focused in studying the synergistic behavior of molybdate with other inorganic compounds 52, 53.

The mechanism by which molybdate imparts inhibition is not well understood and different mechanistic proposals have been put forward over the years 62, 63, 76, 78.

According to Moshier et al. 62, molybdate forms an adsorbed layer on the surface that is concentration dependent. At molybdate concentrations less than 100 ppm, a thick film forms consisting mostly of MoO2. In contrast, at higher concentrations, the ratio of

2- Mo(IV) to Mo(VI) decreases and the molybdate anion MoO4 is preferentially adsorbed

18 on the surface, thereby impeding the ingress of chloride ions. An alternative mechanism proposed by Breslin et al. involves an oxidation-reduction process whereby aluminum is oxidized and molybdate is reduced 63. It was suggested that there is competitive adsorption between molybdate and chloride ions at flawed regions in the passive film. At sufficiently high molybdate concentrations, reduction of Mo(VI) to Mo(IV) is favored resulting in the formation of a protective MoO2 film.

The effect of molybdate on the complex corrosion processes of AA2024-T3 has been investigated by some authors. Jakab et al. studied the effect of molybdate in solution on the Cu replating, anodic behavior, and oxygen reduction reaction (ORR) on

AA2024-T3 in NaCl solution 77, 78. It was found that molybdate suppresses the localized corrosion of AA2024-T3 by decreasing oxygen reduction kinetics and significantly increasing the pitting potential. However, large critical concentrations were necessary to achieve inhibition 77.

2.3.5 Molybdate Conversion Coatings, Pigments, and Synergisms

Recent developments involving molybdate-based conversion coatings have shown promise as possible CCC replacements 50. For instance, a molybdate conversion coating developed by Rodrigues et al. ennobled the corrosion potential of AA2024 and protected the underlying substrate via anodic inhibition 80, 81. It was determined that the coating consisted of two layers; a thin top layer of Mo(VI) and Mo(V), and an inner layer of reduced Mo(IV) 81. Furthermore, it is hypothesized that the Mo(VI) present on the surface can migrate and reduce over an active site thereby providing a self-healing effect similar to chromate. However, no evidence exists to support this claim. 19

It should be noted that molybdate is considered to be inferior to chromate in suppressing the cathodic reaction on aluminum, which is particularly important for copper-bearing alloys and the ability to maintain good adhesion of paint coatings 82.

Therefore, numerous cathodic inhibitors are used in conjunction with molybdate to provide a synergistic effect in protecting both steel and aluminum alloys 51. In this regard, cerium- and zinc- molybdate have been considered to be promising substitutes for chromate pigments 50, 83. In addition, many examples are reported of organic inhibitors that are synergistic with molybdate 84-89, including a molybdate-triazole combination used to protect steel and aluminum in cooling water 85.

The effect of synergism has also been extended to the development of composite conversion coatings 90. Mansfield and colleagues developed a treatment that involves chemical passivation in hot cerium salt solution followed by subsequent anodic

91 polarization in Na2MoO4 to protect aluminum alloys . The modified surface layer produced exceptional corrosion resistance, lowering the anodic passive current density and decreasing oxygen reduction kinetics by three-orders and one-order of magnitude, respectively. Panels subjected to 0.5 M NaCl for 30 days or in a salt-spray chamber

(ASTM B-117) did not show any signs of degradation.

2.4 Silicates

Alkali silicate solutions (i.e. water-glass) are widely used chemicals for a variety of applications. In particular, they are used in the pre-treatment stages of aluminum alloys as cleaners and corrosion inhibitors 92. Furthermore, natural alkali earth silicates

(e.g. CaSiO3, MgSiO3) are commonly used as commercial pigments in coating systems. 20

It is important to note that silicate is redox inactive and its mechanism is highly dependent on the conditions of the environment.

2.4.1 Background and Toxicity

Soluble alkaline silicates have been used in the past to effectively inhibit the corrosion of various metals in water distribution systems 93. Generally, silicates impart anodic inhibition by increasing the resistance to localized corrosion through the formation a hydrated metal- silicate oxide on the surface 93, 94.

Although the biological role of silicon (silica, silicate) has been studied extensively in literature, there is controversial evidence surrounding its effect as a carcinogen 95. However, most studies suggest that when silica enters the body, some type of toxicity will invariably result via the deposition of silica in various parts of the body. Inhalation has been shown to be the most common way by which silica exerts a serious effect with many studies involving workers who are occupationally exposed to crystalline silica and suffer from silicosis. Despite the controversies that exist in the data, the International Agency for Research on Cancer (IARC) has classified crystalline silica as a Group 1 carcinogen (i.e. that there sufficient evidence for carcinogenicity in experimental animals and humans), with the lung as the sole tumor site 96. Inhalation or exposures to other forms of silica based compounds have shown limited effect in comparison to crystalline silica. Interestingly, a review of the toxicity of inhaled amorphous silica indicated that its effects on the respiratory tract may be reversible upon termination of the exposure 97. In addition, depending on the concentration, the silica-to-

21 alkali ratio, sensitivity of the exposed tissue, and the length of exposure, soluble silicates can induce effects ranging from irritation to degradation of cell tissue 98.

2.4.2 Aqueous Chemistry

The nature of silicate anions species in aqueous solution has been extensively studied in the past. A majority of these investigations involved potentiometric titrations, infrared, Raman and 29Si NMR spectroscopy 99-105. It has been shown that the anionic complexity of silicate in solution is related to the concentration and the silicon-to-cation

29 ratio (SiO2/Na2O). After combining the results obtained by Si NMR and infrared spectroscopy, Bass et al. determined that larger quantities of polymeric silicates were

101 present in solution with higher SiO2:Na2O ratio and higher SiO2 concentrations . In contrast, monomer, dimer, and cyclic trimer silicate species were dominant in diluted alkaline solutions (SiO2:Na2O < 1.65, 0.4 M). Recently, Halasz et al. assigned the infrared and Raman bands to silicate species existing in sodium metasilicate solutions and claimed that silicate monomers predominated in 0.2-3 M aqueous metasilicate solutions

100.

2.4.3 Corrosion Inhibition in Ferrous and Non-Ferrous Alloys

The effects of silicates and silica deposits on corrosion have been reported extensively in literature. Adsorption models have been proposed for various metals including iron 106-110, copper 111, zinc 112, 113, and aluminum 114-119. In the case of steel, both Fe2+ and Fe3+ participate in the formation of a protective layer by reacting with silicate 120. Yang et al. investigated the adsorption of silicate on synthetic magnetite

(Fe3O4) and proposed an adsorption mechanism in accordance with a ligand exchange reaction 106. It was proposed that monodentate and bidentate complexes exist at low surface loading levels with subsequent polymerization occurring on the surface at lower pH.

A few studies have been conducted on the interaction of silicates with aluminum oxides 114-117, 119. Firment et al. proposed a deposition mechanism for monomeric silica coatings on alumina particles in aqueous solution 119. A random growth mechanism was proposed involving adsorption of silica units approaching the surface where they first contact the surface. Complete coverage of the surface is achieved as the silica loading increases. Gaggiano et al. studied the interaction of soluble sodium silicates on porous anodic alumina 115-117, 121. It was proposed that aluminosilicate anions in solution react with the aluminate ions formed during oxide dissolution at high pH. The sodium cations act as a coagulating agent between the negatively charged aluminum oxide surface and aluminosilicate anions in solution. A schematic model of this mechanism is depicted in

Figure 2.1.

2.4.4. Organically Modified Silicate Coatings and Synergisms

Silica-based sol-gel derived coatings are being considered as potential candidates for chromate replacement 122-125. These coatings hydrolyze and condense upon deposition to form a continuous three-dimensional oxide matrix. The resulting coating provides good protection by forming an inert hydrophobic barrier. However, the internal stresses of the system limit film thicknesses to less than 1 µm; otherwise cracking will 23 occur 26. Thus, in order to promote defect free films, metal oxides are being combined with organic segments to lower the intrinsic stresses and reduce shrinkage of the coating.

This has led to the advent of organically modified silicate (Ormosil) coatings for the protection of aluminum based alloys 26, 122, 126-128.

The sol-gel process of Ormosil coatings involves the chemical bonding of an organic component to a silica matrix 122. These coatings provide increased protection due to their barrier properties, excellent adhesion, chemical inertness, and ease of application.

For instance, Metroke et al. determined that the adhesive bond strength of an epoxy primer coated onto an Ormasil coating exceeded 2000 psi according to ASTM D4541 125.

Using the same test, the average bond strength calculated in the presence of a chromate conversion coating (Alodine 1200) was approximately 1450 psi. In another example, the use of polyorganosiloxane-modified natural products (potato starch) improved the impedance response of AA2024-T3 by two orders of magnitude and resulted in panels that were corrosion-free after 288 hours of 5% salt fog exposure 129. Furthermore, the flexibility of the sol-gel process allows for the incorporation of corrosion inhibitors that enhance the protective capabilities of the organically-modified coatings 125.

Other studies involving the corrosion protection of aluminum alloys have focused in studying the synergy exhibited by various inhibitor combinations. Taylor and

Chambers developed high-throughput methods to assess binary pairing of 12 inorganic chemistries 130. While some systems exhibited antagonistic behavior, others demonstrated synergies that were comparable to or better than chromate. These chemistries included pairings of silicate with rare-earth metal cations and vanadates. In

24 addition, studies involving corrosion of carbon steel in binary mixed solutions of silicate and molybdate yielded improved corrosion resistance 131. The results indicate that molybdate and silicate are incorporated in the surface layer, but the extent of molybdate uptake was limited by the preferential reaction of silicate with the surface.

2.5 Rare-Earth-Metal Compounds

A different approach to finding a suitable replacement for chromate involves using cathodic or mixed inhibitors 83. In this respect, rare-earth (RE) metal compounds have been investigated and have shown to be promising candidates as chromate replacements. In this section, a review of the information available on this topic is presented.

2.5.1 Background and Toxicity

Rare-earth (RE) metal compounds, principally cerium, were used initially for the corrosion protection of metals at high temperature 83. The use of these compounds was stimulated in response to the regulations imposed on the use of chromates. Since chromate is used extensively in the protection of aerospace aluminum alloys, most of the research has involved studying the effect of RE salts on the corrosion inhibition of these alloys.

The use of RE metal compounds has substantially increased in a number of fields.

As a result, larger concentrations of these compounds are entering the environment and are accumulating in the human body 132. However, the effects of RE compounds are still unclear. According to Haley 133, these compounds have low toxicity, and are not harmful

25 if ingested or inhaled depending on the stability of the complexes. For instance, chelated

RE compounds can be excreted rapidly via urine, while unchelated ionic RE easily form colloids in the blood that can target the liver and spleen 134. In addition, long-term exposure of RE salts seems to cause pneumoconiosis (i.e. lung disease) in humans134.

Despite the extensive use of RE metals in a wide-range of applications, many toxicologists agree that further tests are required to ensure their use is safe for practical purposes 132, 134.

2.5.2 Inhibition by Rare-Earth Salts and Synergisms

The early work by Hinton et al. 135-139 demonstrated that the corrosion rate of aluminum alloy 7075-T6 (Al-Zn) decreased by an order of magnitude when 100 ppm of

RE compounds were added to 0.1 M NaCl. It was determined that the oxygen reduction kinetics were reduced and the corrosion potential shifted to lower potentials 136. Cerium and praseodymium solutions provided better inhibition, exhibiting larger shifts in the corrosion potential and lower pit densities in comparison to other RE salts. Furthermore, the evolution of the corrosion rate in the presence of 1000 ppm RE chlorides has been

136 compared with other salts such as FeCl2, CoCl2 and NiCl2 . According to the results, cerium exhibited the best degree of inhibition. Experiments performed on other aluminum alloys such as AA5083 (Al-Mg), showed similar results with 500 ppm of

140 CeCl3 providing the best inhibition against corrosion . In the case of AA2014 (Al-Cu), the addition of 1000 ppm CeCl3 in 3.5% NaCl increased the polarization resistance and decreased the corrosion rate by an order of magnitude 141. Furthermore, studies have shown that binary mixtures of RE salts display a synergistic effect 142. For instance, a 26 binary chloride solution doped with 250 ppm CeCl3 + 250 ppm LaCl3 had a larger effect on the oxygen reduction kinetics and polarization resistance compared to their individual additions.

Recently, studies have focused in developing innocuous hybrid inhibitors of organic-inorganic composites that exhibit synergistic behavior. For instance, Forsyth et al. 143, 144 investigated the feasibility of coupling RE metals with multifunctional organic components such as diphenyl- (dpp-) or dibutyl- (dbp-) phosphate. Polarization curves conducted in aerated solution indicated that Ce(dpp)3 is an effective cathodic inhibitor, decreasing the oxygen reduction kinetics by an order of magnitude and decreasing the corrosion potential by more than 150 mV 145. It was proposed that the onset of corrosion increases the local pH high enough to dissociate the composite inhibitor, allowing Ce3+ to deposit over the intermetallic particles of AA2024-T3 and dpp- to form an insoluble complex over the aluminum matrix.

In summary, RE salts and composites show an efficiency that is similar to chromate. Their ability to effectively hinder cathodic kinetics for a large variety of aluminum alloys makes them promising candidates for the replacement of the highly effective, but toxic chromate.

2.5.3 Mechanism of Inhibition

A simple method proposed by Aldykewicz et al. confirmed the cathodic nature of cerium salts on AA2024 146. An aluminum-copper galvanic couple was used to mimic the electrochemical behavior of the intermetallic particles present within the Al matrix.

Their results showed that the corrosion inhibition is related to the formation of a cerium- 27 rich film over the copper surface. In addition, free-corrosion experiments conducted in cerium chloride solution showed that cerium was concentrated only at the copper containing intermetallic particles 135, 146. Thus, a mechanism of corrosion inhibition has been proposed which involves the formation of an insoluble film over the intermetallic particles in response to an increase in the local pH at the metal/electrolyte interface. The solubility of the RE salt decreases in alkaline environments resulting in the formation of an insoluble oxide/hydroxide that adheres strongly to the surface. This film provides a barrier to the diffusion of oxygen and the transport of electrons to the oxygen reduction reaction. Hence, the overall corrosion rate is decreased. A similar proposal was made by

Yasakau et al. concerning the deposition of cerium and lanthanum over S-phase in

AA2024 147. It was determined that the precipitation of a hydroxide over the S-phase particles inhibits the anodic and cathodic processes, and decreases the redeposition of copper on the surface.

Despite the proposed mechanism of inhibition by RE metal compounds, a better explanation for the inhibition by cerium has been described involving the oxidation of

Ce3+ to Ce4+ under sufficiently alkaline conditions 148. Due to the complexities of the oxygen reduction reaction, a two-electron reduction pathway may proceed in which oxygen is reduced to produce hydrogen peroxide at the vicinity of the intermetallic particles. Oxidation of Ce3+ will proceed in the presence of hydrogen peroxide resulting in the precipitation of insoluble CeO2. In the case where oxygen reduction occurs by a four-electron reduction pathway, oxidation of cerium is likely to occur away from the

28 intermetallic particles, leading to inefficient film deposition. Evidence of Ce3+ oxidation was also supported by X-ray photoelectron spectroscopy studies on AA7075-T6 136.

2.5.4 Rare-Earth Conversion Coatings on Aluminum Alloys

The strong barrier properties afforded by RE oxides/hydroxides have led to the development of alternate conversion coatings on aluminum alloys 50, 83. First attempts by

Hinton et al. involved prolonged exposure of AA7075 by full immersion in 1000 ppm

149 CeCl3 solution . According to the authors, a minimum corrosion rate was observed after 90 hours of immersion. Although it was determined that the resulting coating provided strong corrosion protection and improved adherence of deposited epoxy layers, such a prolonged treatment of the surface is impractical for commercial use 150.

Therefore, in order expedite the film forming process, a treatment method called

‘cerating’ was developed and implemented for both ferrous and non-ferrous alloys 151. In this process, the alloy is immersed in cerium chloride solution containing numerous oxidizing agents and organic additives. Further improvements led to the formation of

‘cerate coatings’ that were successfully applied to AA2024. These coatings are obtained by full immersion for 10 minutes in cerium chloride solution with 0.3% (v/v) hydrogen- peroxide, at 316 K and pH 1.9. In addition, a prior cleaning and deoxidizing treatment is applied, and the coating is sealed in a post-treatment stage using silicate solution 152. It was determined that the cerate coating lowers the corrosion potential by almost 200 mV, and protects the underlying metal for 336 hours in salt-spray exposure.

Other approaches based on electrolytic activation have been proposed for the development of new coatings. For instance, Hinton and co-workers deposited coatings 29 on AA7075 using galvanostatic measurements with cathodic currents up to 0.2 mA·cm-2

153 in CeCl3 solution . Although this technique reduced the time required for deposition, the degree of protection was also reduced by the presence of small holes in the coating due to hydrogen evolution. A better performance was obtained by Mansfeld et al. 91 for coatings developed by immersion in boiling Ce(NO3)3 and CeCl3, followed by anodic polarization in 0.1 M Na2MoO4 solution. The resulting film was determined to be very stable and effectively hindered the oxygen reduction kinetics.

In summary, comparable protection levels to chromate have been achieved by either immersion or electrolytic activation methods using RE metal compounds.

However, these methods have limitations that restrict their use in practice. Therefore, further research is necessary to develop feasible treatments for chromate replacement.

2.6 Unresolved Issues

The pursuit to find a chromate-alternative has led to the development of several chromate-free aerospace primers that offer good protection. However, fundamental understanding of the functionality of the chromate-free pigments that are embedded within these primers is lacking, thereby hindering further advances and creating risk associated with their use. Therefore, it is critical to investigate the inhibition provided by current non-chromate inhibitors to develop the kind of understanding that was accomplished with chromate. Furthermore, considering that the number of non-chromate corrosion inhibiting components is very large, it is necessary to develop reliable methods that can test various aspects of non-chromate protections systems including inhibitor

30 release and transport within primer layers, inhibition mechanism of underlying substrate, and the effect of multi-component inhibitor pigments.

One approach to study the inhibition mechanism of non-chromate inhibitors has focused in utilizing the techniques and understanding that was obtained by the earlier work on chromate. This was particularly effective for studies involving the inhibition mechanism of AA2024-T3 by vanadates 154-156. For instance, utilizing the methods that were developed for chromate, it is now understood how vanadate inhibition is greatly influenced by vanadate speciation. In another example, it has been determined that the particular phase of praseodymium oxide added to an epoxy-polyamide primer matrix had a significant effect on the activation and transport of the inhibiting species 157. These examples provide proof of the usefulness of the lessons learned from the work on chromate and how fundamental studies on important aspects of non-chromate inhibition can lead to improvements in practical applications.

Despite the improved corrosion protection afforded by molybdate, silicate, and praseodymium in non-chromate protections schemes, significant questions remain regarding their functionality on the corrosion inhibition of high-strength aluminum alloys. For instance, it is not completely understood whether molybdate imparts anodic inhibition by an adsorption or reduction mechanism. In addition, the effect of corrosion inhibition on aluminum alloy 2024-T3 by silicate and praseodymium has not been studied in detail. Therefore, a similar approach to the one described above can be employed to help advance the understanding of leading non-chromate conversion coatings and pigments, including the aforementioned inhibitors. In all three cases, the exact form of

31 the species that is responsible for inhibition is not known, and the environment influence

(oxygen, CO2, pH, etc.) on their inhibition performance has not been thoroughly ascertained.

Furthermore, recent studies involving the corrosion protection of aluminum alloys have focused in studying the synergy exhibited by various inorganic and organic inhibitor combinations 130. However, detailed mechanistic studies regarding the inhibition of mixed inhibitor components have not been actively pursued. Advances in the scientific understanding of these issues will undoubtedly help facilitate improvements to many non-chromate technologies.

FIGURES

Figure 2.1. Schematic representation of silicate anion deposition of aluminum oxide surface.

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CHAPTER 3: CORROSION INHIBITION OF AA2024-T3 BY SODIUM MOLYBDATE

3.1 Introduction

High strength aluminum alloy 2024-T3 (AA2024-T3) is used widely in aircraft structure and design 1-3. The major alloying element is Cu and the nominal composition of AA2024-T3 is 3.9-4.9% Cu, 1.2-1.8% Mg, 0.3-0.9% Mn, 0.5% Fe, 0.5% Si, and balance Al 4. The presence of Cu and Fe render the alloy susceptible to localized corrosion in the form of pitting and intergranular corrosion 5-7. In general, the susceptibility to localized corrosion increases with Cu content, an effect that is attributed to the presence of local galvanic cells that form between the precipitated Cu-rich particles and the Al matrix 8-12. The most predominant type of intermetallic particle found in the structure of AA2024-T3 corresponds to the S-phase precipitates with Al2CuMg composition 10. Previous studies have shown that S-phase particles undergo preferential dissolution of Mg and Al when exposed to aggressive chloride solutions, leaving behind

Cu-rich remnants that are nobler than the surrounding matrix. The local galvanic interactions established between the de-alloyed S-phase particles and Al matrix can then lead to a trench-like attack or pitting on the matrix 13. Furthermore, non-faradaic liberation of metallic Cu occurs as the dealloyed particle remnant coarsens, which can then oxidize to form ions that can be replated on the surrounding matrix 10, 14. Therefore, suppressing Mg dissolution from S-phase particles will prevent the formation of local Cu 43 cathodes capable of supporting oxygen reduction, thereby effectively increasing the resistance of AA2024-T3 to localized corrosion 15.

Traditionally, chromate-based pigments and conversion coatings have been used to prevent corrosion of high-strength aluminum alloys 16, 17. They effectively reduce the rate of oxygen reduction over cathodes and moderately hinder anodic dissolution kinetics

18, 19. Furthermore, chromate conversion coatings (CCC) have the distinct ability to release soluble Cr(VI) into the local environment at a defect providing a “self-healing” characteristic on various metals 20-22. However, the carcinogenic effect of soluble hexavalent chromate has led to the desire to find environmentally-friendly inhibitors that offer the same level of protection and reliability as chromate protection schemes 17. In this regard, molybdate and molybdate-based compounds have in the past been used to replace chromate and have shown promising behavior as corrosion inhibitors in various environments 23. Furthermore, recent studies have shown that molybdate provides a synergistic inhibition effect in combination with other compounds, effectively enhancing corrosion protection for both ferrous and non-ferrous alloys 23-25.

The impact of molybdate on the corrosion resistance of aluminum has been explored extensively in the literature 23, 26-36. Molybdate improves the pitting resistance of aluminum and its alloys by increasing the pitting potential to more noble values with respect to the open-circuit potential 30. However, the extent of inhibition is highly dependent on the ratio of molybdate to chloride ion concentration 35, 36. Current-time measurements in which an aluminum electrode was polarized slightly above the pitting potential in 0.1M NaCl solution showed that the current fluctuations associated with

44 pitting decayed with incremental amounts of molybdate 34. Furthermore, the addition of

15,700 ppm Na2MoO4 to 1000 ppm chloride solution at pH 7 increased the pitting potential of pure aluminum by 1300 mV 29. However, the large ratio of molybdate-to- chloride ion concentration may not be suitable for organic coating applications where high concentrations of inhibitor release necessary for corrosion inhibition may lead to coating degradation via osmotic blistering 37. Therefore, recent studies have focused in studying the synergistic behavior of molybdate with other inorganic compounds 24, 25.

AA2024-T3 in NaCl solutions 35, 36. It was found that molybdate suppresses the localized corrosion of AA2024-T3 by decreasing oxygen reduction kinetics and significantly increasing the pitting potential. However, large critical concentrations were required to achieve inhibition 35.

The mechanism by which molybdate imparts inhibition is not well understood and different mechanistic proposals have been put forward over the years 29, 32, 34, 36.

According to Moshier et al. molybdate forms an adsorbed layer on the surface that is concentration dependent 29. At molybdate concentrations less than 100 ppm, a thick film forms consisting mostly of MoO2. In contrast, at higher concentration, the ratio of

2- Mo(IV) to Mo(VI) concentration decreases and the molybdate anion MoO4 is preferentially adsorbed on the surface thereby impeding the ingress of chloride ions. An alternative mechanism proposed by Breslin et al. involves an oxidation-reduction process

45 whereby aluminum is oxidized and molybdate is reduced 32. It was suggested that there is competitive adsorption between molybdate and chloride ions at flawed regions in the passive film. At sufficiently high molybdate concentrations, reduction of Mo(VI) to

Mo(IV) is favored resulting in the formation of a protective MoO2 film.

Despite the improved corrosion protection afforded by non-chromate protection schemes that incorporate molybdate as a corrosion inhibitor for high-strength aluminum alloys, significant questions remain regarding the functionality of molybdate. The purpose of this study was to understand the mechanism of corrosion inhibition provided by molybdate for AA2024-T3, with the goal of developing the level of understanding that was accomplished with chromate. Furthermore, since most inhibiting conversion coatings and pigments act by releasing soluble species into local aqueous environments, it was of interest to understand the mechanism of inhibition provided by molybdate dissolved in aqueous NaCl solution. A mechanism detailing the functionality of molybdate on the localized corrosion processes of AA2024-T3 is proposed.

3.2 Experimental

3.2.1 Materials and sample preparation: Reagent-grade sodium molybdate and sodium chloride were used for all experiments. Solutions were prepared using 18.2 MΩ-cm deionized water.

Samples of solution heat treated, naturally aged AA2024-T3 (nominal composition 3.9-4.9% Cu, 1.2-1.8% Mg, 0.3-0.9% Mn, 0.5% Fe, 0.5% Si) 4 were mechanically abraded with SiC paper to 1200 grit in a nonaqueous slurry (Blue Lube from Struers) to minimize the onset of corrosion. Samples analyzed microscopically 46 were polished to 1 μm diamond paste. All samples were cleaned with ethyl-alcohol in an ultrasonic bath, air dried, and stored overnight in a desiccator. Disk samples approximately 1 cm in diameter and 0.3 cm thick were cut from an AA2024-T3 sheet for atomic force microscopy (AFM) experiments. For electrochemical and free-corrosion experiments, sample dimensions of 2.5 x 2.5 x 0.3 cm were employed.

3.2.2 Chronoamperometry and potentiodynamic polarization curves: Aerated chronoamperometry experiments were performed to study the effects of molybdate injection at constant applied potentials. Solutions were continuously bubbled with air to enhance convection after injection. A platinum-mesh counter electrode and a saturated reference electrode (SCE) were used for both chronoamperometry and electrochemical tests. Figure 3.1 shows the chronoamperometry cell set up.

Cathodic and anodic polarization experiments were carried out in either aerated or deaerated solutions. Aerated experiments were performed in quiescent naturally aerated solution. Deaeration was achieved by 2 h Ar degassing of both the solution reservoir and the electrochemical cell. After 2 h, solution was forced into the cell by pressurizing the reservoir. Cathodic and anodic polarization curves were acquired separately under each condition and were repeated at least in triplicate. The potential sweep was conducted starting at the open-circuit potential (OCP), using a scan rate of 10 mV/min for all experiments.

3.2.3 Electrochemical Impedance Spectroscopy: Impedance measurements were performed with and without molybdate additions to 0.1 M NaCl solution. Experiments were conducted in aerated and deaerated electrolyte following the same procedure as

47 described above. A 10 mV sinusoidal voltage was applied at OCP with the frequency ranging from 105 to 10-2 Hz. The impedance response was fitted using an equivalent circuit to extract total resistances and film capacitances.

3.2.4 Free-corrosion experiments coupled with secondary electron microscopy:

AA2024-T3 coupons were exposed in aerated NaCl solution with and without molybdate.

Samples immersed in solution were rested against the beaker wall, with the unpolished surface facing the wall and covered with electroplating tape (3M). At the end of the exposure, all samples were rinsed using 18.2 MΩ-cm deionized water, air-dried, and stored overnight in a desiccator. They were visually inspected and analyzed using either a Quanta 200 or Sirion scanning electron microscope (SEM) integrated with energy dispersive X-ray spectroscopy (EDS).

3.2.5 X-ray photoelectron spectroscopy: XPS measurements were performed to investigate the presence of molybdate at the surface of AA2024-T3 and its possible oxidation state. Measurements were conducted on samples that were immersed under free-corrosion or chronoamperometric conditions. XPS spectra were acquired using a

Kratos AXIS Ultra spectrometer with a monochromated Al x-ray source operated at 130

W and calibrated to the adventitious carbon 1s peak at 284.6 eV. Chemical state assessemnt was achieved by curve-fitting the spectra using CasaXPSTM software. Fitting of the Mo 3d spectra was performed by constraining the orbital-split components with an energy separation of 3.2 eV and an intensity ratio of 1.5 29.

3.2.6 Raman spectroscopy: Raman spectra were acquired with a Renishaw inVia Raman microprobe system using a 514.5 nm laser. A video charged coupled device (CCD)

48 camera was used to focus the laser (spot size ~50 µm) over different microstructural features on the surface. All spectral results were repeated at least in duplicate.

Furthermore, characterization of the surface was carried out before each measurement using SEM/EDS.

3.2.7 Atomic force microscopy: In situ atomic force microscopy (AFM) scratching experiments were conducted with a Multimode AFM coupled with a Nanoscope V controller (Bruker Corporation). Naturally aerated electrolyte was pumped through a glass cell with a volume of 0.1 mL at a rate of 10 mL/h. Scratching was performed in contact mode with commercially available Si cantilever tips at a constant scan rate of 2.5

Hz. The force applied by the tip on the sample surface is proportional to the deflection of the cantilever which is kept constant and controlled by the set-point voltage. The set- point voltage was adjusted from 0.2 to 10 V which correlates to tip-sample pressures ranging from 100 to 525 nN, respectively. Therefore, increasing the set-point voltage is equivalent to increasing the scratching force on the sample. Tips were replaced after each experiment due to wear from hydrogen gas evolution and other corrosion products that form during scratching. Prior to scratching, scanning Kelvin probe force microscopy

(SKPFM) was conducted in air allowing for the simultaneous measurement of the surface topography and Volta-potential distribution. Furthermore, characterization of the scratched surface was carried out before and after each experiment using SEM/EDS.

3.3 Results and Discussion

3.3.1 Aqueous solution chemistry: The Mo species expected in dilute sodium chloride solution were determined using chemical equilibrium diagrams computed by MedusaTM. 49

Embedded within the software is a database that contains chemical equilibrium constants at 25°C. Figure 3.2 shows the generated species diagram for 10 µM, 1 mM, and 100 mM

2- MoO4 in 0.1 M NaCl solution. The natural pH of solution is concentration dependent,

2- increasing with incremental addition of MoO4 . The calculated natural pH at the concentrations studied is between 7.0 and 8.4. The increase in pH is associated with the

- formation of HMoO4 ions in solution. However, under all conditions, the dominant

2- species is the molybdate anion (i.e. MoO4 ) at pH values greater than 6.

In acidic solution, a number of differences are observed. At the lowest

2- - concentration studied, protonated forms of the MoO4 ion (i.e. HMoO4 , H2MoO4) become the preferred species at pH less than 4.5. However, at higher concentrations the dominant species in acidic solution are protonated forms of the heptamolybdate ion

6- Mo7O24 , which forms at pH values lower than 6 for the highest concentration studied.

The observations obtained from the species calculations are consistent with previous reports in literature 38-40. Based on these studies, it was concluded that the molybdate species in solution undergoes condensation polymerization at pH values lower than 6. The first species to form is the heptamolybdate species; at higher acidification the

-4 octamer Mo8O26 forms; above 1.5 protons bound per molybdenum, more aggregated species occur.

3.3.2 Polarization in aerated sodium chloride solution: The corrosion inhibition effect of molybdate on AA2024-T3 was determined using potentiodynamic polarization curves.

Figure 3.3 shows naturally aerated polarization curves in 0.1 M NaCl solution with varying Na2MoO4 concentration. The natural pH measured in these solutions was

50 between 7.1 and 8.3. The results show that molybdate provides a mixed inhibition effect.

It causes a decrease in the OCP and an increase in the breakdown potential, Ebr, associated with pitting at concentrations larger than 100 mM Na2MoO4. The largest effect was observed at 125 mM where Eocp decreased by 350 mV and the breakdown potential, Ebr, increased by 250 mV. At concentrations lower than 100 mM Na2MoO4,

Eocp did not change and Ebr slightly increased. Figure 3.4 summarizes the effect of molybdate concentration on Ebr and Eocp. A larger separation between Ebr and Eocp at higher molybdate concentrations is indicative of improved corrosion resistance by lowering the tendency for pit initiation 41. Based on these results, it can be inferred that the threshold concentration for corrosion inhibition of AA2024-T3 in 0.1 M NaCl solution is between 50 and 100 mM Na2MoO4.

3.3.3 Free corrosion morphology of AA2024-T3 after exposure: Since the maximum effect of corrosion inhibition by molybdate was observed at 125 mM concentration, it was of interest to study the free corrosion morphology of AA2024-T3 after exposure in

0.1 M NaCl + 125 mM Na2MoO4 solution. Figure 3.5a shows a macrograph of a sample exposed to 0.1 M NaCl solution after 1 day as a control. In the absence of inhibitor, optical microscopy reveals evidence of attack associated with both the intermetallic particles and matrix (Figure 3.6a). After 1 day exposure in 0.1 M NaCl + 125 mM

Na2MoO4, the extent of attack is lowered (Figure 3.6b). Furthermore, there is evidence of a thick oxide on the intermetallic particles.

Figures 3.3b-h shows macrographs after OCP exposure in naturally aerated solution with molybdate at varying times. After 30 min of exposure, the surface turns

51 light-brown in color. SEM-EDS analysis shows evidence of a film present over the intermetallic particles (Figure 3.7). The large concentration of Mo and O detected by

EDS over the Cu intermetallic particles suggests the formation of a Mo-based oxide. In addition, no localized attack was observed on the surface including within or at the periphery of S-phase and Fe-containing particles.

Figure 3.8 shows SEM-EDS analysis of an S-phase particle after a 2 day exposure under the same solution conditions. As for the one day exposure, the secondary electron micrograph (Figure 3.8a) shows the presence of an oxide film over the particle that is rich in Mo and O as detected by EDS (Figure 3.8b). Furthermore, Mg and Al were detected at the particle, which suggests that molybdate suppresses the dealloying of S-phase particles. Interestingly, a larger concentration of Mo and O was detected at the periphery of the particle. Similar results were obtained after analysis of a Fe-containing intermetallic particle under the same solution conditions as shown in Figure 3.9.

X-ray photoelectron spectroscopy was used to investigate the oxidation state of

Mo on the surface after 1 day exposure in naturally aerated solution containing molybdate. Figure 3.10 shows the resulting Mo 3d spectrum, which reveals a single Mo

3d5/2 peak at a binding energy of 232.6 eV. The higher binding energy peak at 235.8 eV

42 is associated with the corresponding Mo 3d3/2 spin-orbit-split component . These measured binding energies are consistent with Mo(VI) species 29, 42. It should be noted that several compounds could be identified at the aforementioned Mo 3d binding energies, including MoO3, MoOCl4, and Al2(MoO4)3. Although it is not possible to negate the existence of MoO3, the latter compounds were rejected as possible candidates

52 because chlorine was not detected on the surface, and the Al 2p spectrum did not reveal a peak that was consistent with Al2(MoO4)3.

3.3.4 Chronoamperometry: To investigate the mechanism of corrosion inhibition by molybdate, the current evolution at different fixed potentials was monitored before and after molybdate injections. Concentrated Na2MoO4 solution was injected to 0.1 M NaCl solution to make 125 mM molybdate after mixing. Figure 3.11 shows cathodic chronoamperometric measurements at two different potentials. Injection of aerated 0.1

M NaCl at -900 mV SCE had no effect on the cathodic current as expected. In contrast, injection of concentrated molybdate solution resulted in a cathodic current peak followed by a decrease to a net anodic current. At -900 mV SCE, the anodic current reached a maximum at +6.5 µA/cm2 at about 3 min after injection, before decreasing to a steady state value of +1.3 µA/cm2. The shift from cathodic to anodic current suggests cathodic inhibition at -900 mV SCE which is supported by the polarization curves shown in Figure

3.2. Similar trends were observed immediately after injection at -1100 mV SCE.

However, the current reached a steady state cathodic current after decreasing to a net anodic current. After 60 min at -1100 mV SCE, the current was 70% of the value prior to injection.

Figure 3.12a shows a secondary electron micrograph after 2 hour exposure to solution containing molybdate at -900 mV SCE. The micrograph shows a thick oxide present over the cathode supporting intermetallic particles. Attack in the form of micro pits can be observed throughout the matrix with preferentially segregation along the polishing lines. Figure 3.12b shows a secondary electron micrograph after 2 hour

53 exposure at -1100 mV SCE. Unlike in the previous condition, a thick film is present across the whole surface that is rich in Mo as made evident by the supporting EDS spectrum (Figure 3.10c).

The reversible potential for the reduction of molybdate depends on pH and total molybdate concentration. According to the chemical diagrams shown in Figure 3.2, the

2- dominant species in as-prepared solution is the molybdate anion MoO4 . Therefore, the

2- reversible potential for the MoO4 /MoO2 reduction reaction was calculated at varying concentrations based on the results obtained from the chemical speciation diagrams. The reversible potential varies from -570 to -640 mV SCE at molybdate concentrations between 10-3 M and 0.1 M. Consequently, lower applied potentials should increase the driving force for this reaction to occur. The presence of a current peak during injection of molybdate at fixed cathodic potentials suggests that inhibition results from a surface film that forms by reduction. Furthermore, evidence obtained by the electron micrographs in Figure 3.12 indicates that reduction occurs preferentially over the intermetallic particles.

It should be noted that, although the above experimental conditions are different than what is observed at OCP, the same phenomena may still exist at OCP. The global

OCP of AA2024-T3 in 0.1 M NaCl solution is approximately -550 mV SCE and does not change with the addition of molybdate at low concentrations. In contrast, at concentrations of 0.1 M molybdate or greater, there is a significant decrease in the global

OCP, which could allow for the formation of MoO2. The decrease in OCP at these larger concentrations could be attributed to an increase in pH observed with incremental

54 additions of molybdate, which was found to vary from 7.1 to 8.3 in the solutions studied.

2- Alternatively, there could be a competitive adsorption mechanism, where MoO4 in solution is adsorbed on the intermetallic particles and competes with O2 molecules.

Hence, a larger concentration of molybdate may be necessary to displace the adsorbed O2 molecules resulting in a decrease in oxygen reduction kinetics and a decrease in the OCP.

Experiments were also performed at applied anodic potentials to investigate the effect of molybdate injection during the onset of localized corrosion. Figure 3.13 shows that injection of aerated 0.1 M NaCl at -535 mV SCE had no effect on the anodic current.

Injection of concentrated molybdate solution after 60 min produced a small current peak.

After the initial transient, the anodic current decreased within seconds to significantly lower anodic currents. The current continued to decrease, resulting in more than 99% reduction in current after 150 min relative to the current before injection. Similar trends were observed after injection at a fixed potential of -525 mV SCE. However, the current before injection and the observed current peak after injection were higher in magnitude compared to the previous condition. Similarly, more than 99% reduction in current relative to the current before injection was observed.

Figure 3.14a-b shows the optical micrographs after Na2MoO4 injection at -535 mV SCE. It should be noted that prior to injection the current was allowed to stabilize at the applied potential in solution without inhibitor. There is evidence of oxide formation over the smaller intermetallic particles (Figure 3.14a). In contrast, no oxide is observed over the Fe containing intermetallic particles except at the peripheries where attack may have occurred prior to injection during anodic dissolution of the adjacent matrix (Figure

3.15b). The secondary electron micrograph shown in Figure 3.14c reveals the presence of an oxide film over S-phase particles, which is rich in Mo and O as detected by EDS.

In addition, attack in the form of pitting and intergranular corrosion (IGC) is observed throughout the surface. Interestingly, Mo and O were detected inside both corrosion features by EDS.

It was of interest to compare the oxidation state of Mo at the surface after injection of concentrated Na2MoO4 solution at applied potentials to what was observed at

OCP (920 mV SCE) as was shown in Figure 3.10. Figure 3.15 shows the Mo 3d spectrum 30 seconds after injection at -525 mV SCE and -900 mV SCE. Under both conditions, a large peak consistent with Mo(VI) species was observed at a binding energy of 232.7 eV, which constitutes roughly 85% of the Mo on the surface. In addition, a smaller peak at a binding energy of 230.4 eV was present at both potentials. The smaller peak is associated with reduced Mo(VI) species, which is consistent with MoO·(OH)2. It should be noted that several compounds could be assigned at this binding energy, including MoCl4. However, this compound was rejected as a possible candidate because chlorine was not detected on the surface. In addition, subsequent SEM-EDS analysis (not shown) revealed the presence of a thick oxide over the intermetallic particles that was rich in Mo and O.

The area of the current peaks observed in Figure 3.11 was used to calculate the amount of Mo(VI) reduced on the surface via Faraday’s law. For this calculation, it was assumed that reduction occurred only over the Cu-rich intermetallic particles, which is consistent with the free corrosion morphology observed after exposure (Figure 3.6). In

56 addition, inspection of the surface 30 seconds after injection at -900 mV SCE using optical microscopy showed the preferential formation of a thick film over the intermetallic particles. According to Buchheit et al. 10, approximately 4.2% of the total

AA2024 surface is covered by Cu-rich intermetallic particles. The average spike area for

2 the potential range of -0.8 to -1.1 V SCE was found to be 11.4 mC/cm . Assuming that

Mo(VI) reduces to Mo(IV) (2 equivalent/mol), the calculated mean spike area is equivalent to 60 nmol of Mo(IV)/cm2. Using an average bond length of ~0.2 nm for Mo–

O bonds 43, a close-packed Mo(IV) monolayer is predicted to contain 1.3 nmol of

Mo(IV)/cm2. Thus, within the uncertainties caused by the alloy microstructure and film geometry, the present results indicate the rapid formation of 46 monolayers of Mo(IV) over the Cu-rich intermetallic particles. The formation of several dozen monolayers over the intermetallic particles suggests that the reduced film is not self-inhibiting. In contrast, injection of chromate resulted in the rapid formation of a single monolayer over the alloy surface which was attributed to the reduction of CrVI to CrIII 20. Assuming that reduction of Mo(VI) occurs over the whole exposed area and not just the intermetallic particles, the coverage of Mo(IV) is calculated to be 2 monolayers, which can be considered a single monolayer given the error in the calculation. Furthermore, it is possible that the inhibition observed by molybdate could be attributed to the presence of adsorbed Mo(VI) species on the whole surface, which was made evident by XPS (Figure 3.10).

The implications of these results are significant in understanding the corrosion inhibition behavior of molybdate on AA2024-T3. The presence of MoO·(OH)2 on the surface supports the chronoamperometric data which, in combination with SEM/EDS

57 analysis, provided evidence of molybdate reduction over the intermetallic particles.

However, the presence of Mo(VI) on the alloy surface after both OCP and chronoamperometric experiments suggests that either the reduced species oxidizes to

Mo(VI), or that there is subsequent adsorption of molybdate from solution over the surface. These effects are explained in more detail below.

3.3.5 Raman Spectroscopy: Raman spectra were acquired to elucidate the molybdenum oxide species present on the alloy surface. Figure 3.16 shows three offset spectra acquired at different microstructural features (Al matrix, S-phase, Fe-intermetallic) of a sample exposed at OCP (-920 mV SCE) for 2 days in naturally aerated 0.1 M NaCl solution with 0.125 M Na2MoO4. Characterization over the Al matrix revealed a peak at

930 cm-1 with a broad shoulder extending to about 600 cm-1. Interestingly, the peak at

-1 6- 930 cm is consistent with the major Raman band of polymolybdate species (Mo7O24 )

44, 45. The broad shoulder encompasses the major Raman bands for lower Mo oxidation

-1 -1 46 state species, including MoO2 (740 cm ) and Mo4O11 (907 cm ) . Note that the existence of these lower oxidation state species over the alloy matrix could be attributed to the presence of submicroscopic intermetallic particles, such as θ-phase (Al2Cu), that promote Mo(VI) reduction. The broad band in the range from 350 to 250 cm-1 is also consistent with these species 44, 46. Similar results are obtained when characterizing S- phase particles on the surface. All the major bands of the spectra concur with the

6- -1 -1 existence of Mo7O24 (930 cm , 350 cm ) and lower oxidation state species like MoO2

(740 cm-1, 350 cm-1). The band present at 870 cm-1 may arise from additional vibration modes associated with the Mo-O-Mo bonds of the polymolybdate species 44. It should be

-1 45, 46 noted that the major Raman band for MoO3 (996 cm ) was not detected in the above spectra.

Characterization of the Fe-containing intermetallic particles (Fe-IMC) revealed the existence of yet another species on the surface. The corresponding spectrum in

Figure 3.16 shows major Raman bands at 960, 920 and 370 cm-1 which are associated

4- 44 -1 with Mo8O26 . In addition, a peak observed at 790 cm can be attributed to one of the

46 major Raman bands consistent with Mo4O11 . Furthermore, MoO3 was also not detected by Raman spectroscopy at the Fe-intermetallic particles.

The results obtained from Raman spectroscopy are in agreement with previous reports in the literature. According to Hu et al. 44, the structure of surface molybdenum oxides present on Al2O3 under ambient conditions consists of polymeric species

6- 4- (Mo7O24 and Mo8O26 ) at high Mo loadings. In addition, Spevack et al. found evidence of multivalent (Mo(IV), Mo(V), and Mo(VI)) species within molybdenum oxide films supported on alumina and graphite 45. Thus, it can be concluded that a final step in the inhibition mechanism may involve adsorption and subsequent polymerization of molybdate species from solution.

3.3.6 Evaluation of inhibition in deaerated electrolyte: Experiments were performed in deaerated electrolyte to analyze the effect of oxygen on the inhibition mechanism and performance. Figure 3.17 shows cyclic anodic polarization curves in deaerated 0.1 M

NaCl solution. As found previously, AA2024-T3 exhibits two breakdown potentials in the absence of inhibitor 7. The breakdown at -605 mV SCE is attributed to the transient dissolution of S phase particles, while the nobler potential at -550 mV SCE is associated

59 with pitting and intergranular corrosion. In the presence of molybdate, only one breakdown potential is observed. At 1 mM molybdate concentration, Ebr was measured to be -550 mV SCE. With increasing molybdate concentration, Ebr increased with the largest effect observed at concentrations greater than 100 mM Na2MoO4. However, an increase in the breakdown potential was also associated with a 10x increase in the passive current density, which can be attributed to an increase in pH at the higher concentrations.

Furthermore, the addition of molybdate had no effect on the repassivation potential, Erp.

It should be noted that Erp was selected at the point where an inflection in the current density was observed during the reverse scan.

Figure 3.18 compares the cathodic polarization curves obtained in aerated and deaerated electrolyte. Interestingly, the addition of molybdate in deaerated solution had little effect on the hydrogen evolution kinetics. In addition, the cathodic polarization curve is essentially identical for aerated and deaerated conditions in the presence of molybdate. The results suggest that molybdate strongly hinders the oxygen reduction reaction, but does not affect hydrogen evolution kinetics.

Figure 3.19 shows a macrograph and the secondary electron micrographs of a substrate after 1 day exposure at open circuit in deaerated 0.1 M NaCl + 125 mM

Na2MoO4 solution. The surface turned visibly brown-yellow after exposure, which was different than the brown surface that was observed under aerated conditions. The backscattered electron micrograph shows the presence of a film that covers both matrix and intermetallic particles as made evident by the brighter areas in the images.

Furthermore, the electron micrographs show evidence of increased attack in the form of

60 pitting on the Al matrix adjacent to the oxide-covered intermetallic particles. It should be noted that in aerated electrolyte an oxide formed only over the intermetallic particles.

The difference in behavior could be attributed to increased Cu replating on the alloy surface as a result of increased attack in deaerated solution. The attack itself may be related to the inability of molybdate to inhibit hydrogen evolution kinetics as observed in

Figure 3.18. The subsequent increase in surface Cu enrichment may allow for subsequent molybdate reduction over the matrix. Therefore, it is important to note that the increased film coverage in deaerated electrolyte does not signify improved corrosion inhibition, but rather manifests from increased corrosion kinetics.

Figure 3.20 shows the Mo 3d XPS spectrum after 1 day exposure at OCP (-920 mV SCE) in deaerated 0.1 M NaCl + 125 mM Na2MoO4 solution. Two Mo 3d5/2 peaks with binding energies at 230.0 eV and 232.0 eV are observed in the spectrum,

+4 +6 corresponding to Mo and Mo species. The corresponding Mo 3d3/2 binding energies were found to be at 233.3 eV and 235.1 eV. These species are consistent with

2- MoO·(OH)2 and MoO4 and make up approximately 60% and 40% of the Mo concentration on the surface, respectively. The large quantity of Mo(IV) on the surface corroborates the cathodic chronoamperometry data, which suggests reduction of

2- molybdate species from solution. In addition, the presence of MoO4 could be attributed to the adsorption of Mo(VI) species in solution over the bare Al matrix as described below.

3.3.7 Electrochemical Impedance Spectroscopy: Impedance spectroscopy was conducted to determine the inhibition performance under aerated and deaerated electrolytic

61 conditions. The Bode magnitude and phase angle plots are shown in Figure 3.21a. After

1 day exposure in aerated solution, the total impedance was over an order of magnitude higher in 0.1 M NaCl + 125 mM Na2MoO4 solution than in the pure NaCl solution. The total resistance in solution with and without inhibitor was measured to be 5 kΩ·cm2 and

210 kΩ·cm2, respectively. Furthermore, a capacitive reactance over a larger domain in the phase angle plot indicates improved corrosion inhibition.

To extract the total resistance and film capacitance, the spectra were fitted using a simplified Randles circuit that included a constant phase element (CPE) to represent the film capacitance. The true capacitance was derived from the CPE parameters using the following equation 47:

Eq. 3.1 where Y and α represent the CPE magnitude and exponent, respectively, Rp is the extracted total resistance, and Cdl is the calculated film capacitance. The film capacitances were measured to be 82 µF·cm-2 for the case of molybdate in solution, and

579 µF·cm-2 for the negative control.

In deaerated solution, a number of differences are observed. Figure 3.21b shows the Bode magnitude and phase angle plots after 1 day exposure in deaerated solution. A

3x lower total impedance compared to the negative control is observed in 0.1 M NaCl +

125 mM Na2MoO4 solution. The total resistance in the absence and presence of inhibitor was measured to be 34 kΩ·cm2 and 105 kΩ·cm2, respectively. In addition, the respective film capacitance values were measured to be 31 µF·cm-2 and 18 µF·cm-2.

Interestingly, the above results suggest that the inhibition mechanism is oxygen- dependent. In the absence of oxygen, the addition of molybdate actually decreases the corrosion resistance of AA2024-T3. On the other hand, in the presence of oxygen where the Mo(VI) species has been shown to be predominant on the surface, there is significant improvement in the corrosion protection of the alloy. The total resistance increased by

40x and the film capacitance decreased by 7x compared to the negative control. Hence, it is clear that molybdate imparts corrosion inhibition only in aerated conditions, where oxygen reduction dominates the cathodic kinetics. In addition, there is further evidence to support that molybdate in deaerated solution is a poor inhibitor of hydrogen evolution kinetics, where MoO·(OH)2 was observed to be the predominant species on the surface.

3.3.8 Mechanism of Inhibition: According to the electrochemical data, molybdate provides strong mixed inhibition at sufficiently large concentrations. The corrosion inhibition mechanism involves a two-step process whereby aluminum is oxidized (Eq.

3.2) and molybdate is initially reduced (Eq. 3.3) over the intermetallic particles to form

MoO·(OH)2:

Eq. 3.2

( ) Eq. 3.3

The presence of MoO·(OH)2 species is supported by the XPS data acquired in deaerated solution. It should be noted that a lack of thermodynamic data for this species does not allow for the calculation of the reversible potential for Eq. 3.3. However, a conservative approximation can be obtained by substituting the reduced species with its oxide derivative (i.e. MoO2). In that case, the reversible potential is calculated to be

63 approximately -640 mV SCE at 0.1 M molybdate concentration. Consequently, at lower potentials the driving force for this reaction increases. This is supported by the chronoamperometric data, which showed a current peak after molybdate injection at fixed cathodic potentials. A precursor to reduction may very well involve competitive adsorption between the molybdate anions and O2 molecules in solution. It should be noted that this step is similar to the mechanism proposed by Breslin et al. 32.

A similar effect may still exist under open circuit conditions. At sufficiently large molybdate concentrations, the global OCP of AA2024-T3 significantly decreases to potentials below the reversible potential of Eq. 3.3, allowing for the formation of Mo(IV) species. The decrease in OCP could be attributed to the increase in solution pH with incremental molybdate additions. Alternatively, increasing the concentration of molybdate may decrease the available area for O2 adsorption thereby lowering oxygen reduction kinetics.

Note that reduction of Mo(VI) to Mo(IV) was considered only over the Cu-rich intermetallic particles. SEM-EDS analysis and Raman spectroscopy of the sample surface following free-corrosion and chronoamperometric experiments did not suggest the formation of a reduced film over the Al matrix. It is likely that reduction of Mo(VI) is concentrated only over the intermetallic particles since they behave as local cathodes on the surface 8-12, and thus favor reduction of Mo(VI) species. On the other hand, the poor catalytic nature of the oxide-covered Al matrix does not support Mo(VI) reduction.

The results discussed in the previous sections indicate that oxygen plays an important role in the inhibition mechanism. In the presence of oxygen, XPS showed that

64 the dominant species is associated with Mo(VI), while in deaerated solution Mo is predominantly present as MoO·(OH)2 (Mo(IV)). Furthermore, the electrochemical impedance experiments showed a decrease in the total resistance in deaerated solution.

The lack of inhibition observed in oxygen-free environments could be explained by the inhibitive nature of molybdate toward the oxygen reduction reaction. In the absence of oxygen, the low surface potential might decrease the extent of adsorption of molybdate species from solution. In addition, the polarization curves shown in Figure 3.18 indicate that molybdate does not suppress hydrogen evolution kinetics in both aerated and deaerated electrolyte.

The predominance of Mo(VI) species after exposure to aerated electrolyte could be attributed to Mo(VI) adsorption over the alloy surface. The possibility of Mo(VI) adsorption over the Mo-covered intermetallic particles was considered following the same phenomena as observed with chromate. The first stage of CCC formation involves the simultaneous reactions of Al oxidation and reduction of CrVI to CrIII at the metal surface 48. As this process continues, labile CrVI ions reversibly attach to the insoluble

CrIII oxide 6, 49, 50. The adsorption of chromate anions to the chromium oxide that forms over the alloy surface is supported by the isoelectric point of CrIII oxide, which is reported to be ~ 5.5 51. Thus, the adsorption of chromate ions is made favorable by the low pH of chromate baths used during CCC formation. In the case of molybdate, the presence of Mo(VI) over the Cu-rich intermetallic particles was made evident by Raman spectroscopy (Figure 3.16). This may suggest adsorption of Mo(VI) species over the oxide covered intermetallic particles. However, it is important to note that adsorption of

65 molybdate anions at the oxide is not favorable at the bulk pH (8.3) given that the isoelectric point of Mo(IV) oxide is ~2 52, 53. The reported low value of isoelectric point for Mo(IV) oxide is consistent with the general guidelines set by Parks regarding the isoelectric points of metal dioxides 54. Thus, it is unlikely that molybdate anions adsorb on the reduced Mo(IV) oxide present on the intermetallic particles.

In the case of adsorption over the Al matrix, Mo(VI) adsorption was considered following a mechanism proposed by Moshier et al. 29. It was determined that at higher molybdate concentrations, the surface of a pure Al sample is covered by a thin layer of

2- MoO4 species at near-neutral pH. Adsorption of molybdate anions over the aluminum surface is favored given the isoelectric point (pH of zero charge) of oxide-covered aluminum (9.5) 51. In this work, it is proposed that a similar mechanism occurs over the alloy matrix, which is not covered with reduced Mo(IV) like the intermetallic particles as described above. Because the isoelectric point of Al oxide on the matrix is higher than the bulk pH, adsorption of molybdate species in solution onto the Al matrix is promoted.

Furthermore, Raman spectroscopy analysis over the matrix suggests that the adsorbed species condense and polymerize to form polymolybdate species.

The presence of Mo(VI) over the Cu-rich intermetallic particles could be explained by a second step in the inhibition mechanism involving subsequent oxidation of the reduced species. This might lead to the formation of intermediate molybdenum oxides, such as Mo4O11 (Eq. 3.4), and is supported by oxygen reduction (Eq. 3.5):

Eq. 3.4

Eq. 3.5

The generation of protons by Eq. 3.4 could result in local acidification at the film/electrolyte interface (assuming spatial separation of Eq. 3.4 and 3.5). This in turn may promote the subsequent condensation and polymerization of molybdate species from

6- 4- solution to form polymolybdates (Mo7O24 and Mo8O26 ) on the alloy surface, which is supported by Raman spectroscopy (Figure 3.16). The oxidation of the reduced species in

Eq. 3.6 to a higher oxide such as MoO3 was not considered as a likely step in the inhibition mechanism since MoO3 was not detected by Raman spectroscopy after exposure to naturally aerated solution.

Although there is no direct evidence to support the subsequent oxidation of reduced species on the intermetallic particles, this step is supported by various findings in the literature. For instance, the long-term exposure to air of a pure Al surface covered

29 with MoO2 resulted in the formation of Mo(VI) species . Furthermore, Detltombe et al.

55 determined that lower oxides/hydroxides such as MoO2 are easily oxidized in the presence of aerated water to intermediate molybdenum oxides as described in Eq. 3.4

(also known as ‘molybdenum blue’; oxides with molybdenum-to-oxygen ratios of approximately 0.375) 55. In addition, it has been ascertained that these species exist in solution as hydrated compounds 55, 56. Note that a lack of thermodynamic data also does not allow for the calculation of the reversible potential for Eq. 3.4. However, substituting the intermediate molybdenum oxide with hydrated MoO3 species yields a reversible potential of approximately -1800 mV SCE at pH 8, which is well below the global OCP of AA2024-T3 at 0.1 M molybdate concentration, indicating that oxidation of Mo(IV) is possible.

Hence the inhibition mechanism over the intermetallic particles can be summarized by three-steps: the first step involves the simultaneous oxidation of Al and

2- the reduction of MoO4 species to Mo(IV) oxide species. This step is followed by subsequent oxidation of the reduced Mo(IV) species to intermediate molybdenum oxides, such as Mo4O11. These two steps are thermodynamically favored considering that the

OCP (~900 mV SCE) at 0.125 M molybdate concentration lies below and above the

2- reversible potentials for MoO4 reduction and Mo(IV) oxidation, respectively. The last step of the inhibition mechanism involves polymerization of molybdate species from

6- solution in response to a local acidification, resulting in the formation of Mo7O24 and

-4 Mo8O26 . Note that the inhibition mechanism over the Al matrix does not involve the phenomena described above. Instead, the Al matrix is protected following an adsorption mechanism of Mo(VI) species in solution given the higher isoelectric point of oxide- covered aluminum as discussed above.

Another important aspect of the inhibition mechanism involves the suppressed dissolution of S-phase particles. According to SEM-EDS analysis, a molybdenum-based oxide forms over both S-phase and Fe-containing intermetallic particles. S-phase particles initially undergo selective dissolution at the early stages of corrosion before leaving behind Cu-rich remnants that increase oxygen reduction kinetics 10, 14. The formation of an oxide over these particles has shown to effectively reduce anodic dissolution kinetics of AA2024-T3. This was observed electrochemically by an increase in the breakdown potential associated with pitting and an increase in the impedance response in the presence of oxygen. The decreased anodic dissolution kinetics could be

68 attributed to a blockage effect against chloride ion adsorption. Furthermore, a decrease in

S-phase particle dissolution will in effect suppress surface Cu enrichment. This was also captured electrochemically by a decrease in OCP to more electronegative potentials associated with a decrease in oxygen reduction kinetics.

Another possible description of the inhibition mechanism involves direct

2- reduction of MoO4 to both a mixed oxidation state, for example a combination of

MoO·(OH)2 and Mo4O11 species, which have nominal oxidation states of +4 and +5.5, respectively. This can be described by the following equation:

( ) Eq. 3.6

Substituting MoO2 for the reduced species MoO·(OH)2, the reversible potential for the above reaction was calculated to be approximately -660 mV SCE at 0.1 M molybdate concentration. However, this reaction was negated as a prominent step in the inhibition mechanism since it does not account for the role of oxygen. As described above, it is likely that oxygen reduces on the surface to support the subsequent oxidation of reduced

Mo species. The oxidation of the reduced species in Eq. 3.6 to a higher oxide such as

MoO3 was not considered as a likely step in the inhibition mechanism since MoO3 was not detected by Raman spectroscopy.

3.3.9 In Situ Atomic Force Microscopy: In situ AFM scratching involves rastering a region in contact-mode with a Si cantilever at sufficiently high forces to destabilize passivity 11. This technique provides real time information of the early stages of localized corrosion with sub-microscopic resolution. When analyzing the evolution and kinetics of attack, three things have to be considered: the force that initiates attack, the

69 time for attack to nucleate, and the corrosion morphology. Higher forces and longer nucleation times imply greater inhibition performance. This technique was used to further investigate the complex interaction between alloy microstructure and molybdate ions in 0.1 M NaCl solution.

Figure 3.22 shows images of the Volta potential distribution of the scratched area prior to scratching, and the chronological sequence of attack in aerated 0.1 M NaCl

-5 solution with 10 M Na2MoO4. The Volta potential map obtained with SKPFM reveals the presence of a large Fe-containing intermetallic particle and smaller S-phase particles that were confirmed using SEM-EDS. Attack was observed during 15 minutes of OCP stabilization prior to scratching. It should be noted that no attack was observed during

OCP stabilization in inhibitor-free NaCl solution. The form of attack observed was trenching at the periphery of the larger Fe-containing particle, and complete removal of the S-phase particle after 20 minutes of scratching at a constant force (100 nN). Similar results were obtained by Frankel et al.11, 13. After 3 hours of scratching and increasing the scratching force to 150 nN, complete removal of the larger intermetallic particle and pitting in the matrix was observed. It is not clear whether the particles were rapidly dissolved or removed from the surface as a result of undercutting and dislodging from continuous scratching. Furthermore, there is evidence to suggest that attack initiated at the periphery of S phase particles as indicated by an arrow in the 0.5 hour topography image in Figure 3.22.

These results suggest that the dissolution kinetics of AA2024-T3 is exacerbated by the addition of 10-5 M molybdate. Unlike in the case of inhibitor-free solution, attack

70 is observed during OCP stabilization prior to scratching. This effect may be attributed to the increased oxidizing nature of the solution with the addition of molybdate. The

2- reversible potential for the MoO4 /MoO2 reduction reaction at this concentration is approximately -610 mV SCE. The global OCP measured was around -550 mV SCE.

However, local changes in pH at the intermetallic particles may result in a potential shifts, allowing for increased Al oxidation and molybdate reduction. Furthermore, higher molybdate concentrations may be necessary to compete with O2 adsorption over local cathodes before forming a protective oxide film, which is not observed at lower concentrations.

Figure 3.23 shows the obtained SKPFM images prior to scratching and the

-3 evolution of attack in aerated 0.1 M NaCl + 10 M Na2MoO4. The Volta potential map shows the presence of two Fe-containing intermetallic particles and two smaller S-phase particles (characterized beforehand using SEM-EDS). It should be noted that the topography image captured by SKPFM does not resolve the S-phase particles that are present on the surface. Although these particles are harder than the Al matrix, the loss of

57 material by corrosion may balance the slower rate of removal during polishing . Hence, they are initially coplanar with the matrix and are not observable in the topography image prior to exposure. Therefore, the high contrast obtained from the Volta potential map allows for better detection of S-phase particles on the surface. Attack initiated at the periphery of an S-phase particle at the beginning of the experiment, as shown by an arrow in Figure 3.23. After 30 minutes of scratching at a force of 100 nN, the S-phase particle was completely removed from the surface. Slightly increasing the scratching force to 125

71 nN resulted in removal of the second S-phase particle. Increasing the scratching force to the limits of the tip resulted in continuous attack of both S-phase particles. No Mo was detected inside the pits after the experiment using SEM-EDS (not shown). In addition, no attack was observed at the periphery or within the Fe-containing particles throughout the experiment.

The corrosion resistance was drastically improved after increasing the molybdate concentration to 0.1 M (Figure 3.24). Shortly after pumping solution through the fluid cell, a film formed over the intermetallic particles. This was observed with the help of an attached video camera that allows for the observation of an area larger than the scanned area. In addition, the higher topography contrast observed over the intermetallic particles in the 0 hour topography image shown in Figure 3.24 suggests film formation. It is interesting to note that a film forms over both Fe-containing and S-phase particles.

Furthermore, the absence of attack even after increasing the scratching force to the limits of the tip suggests the film is very protective, which considering the previous results, is most likely a mixed oxide consisting of MoO2 and Mo(VI) species. Cross-sectional height measurements with respect to the matrix performed with the AFM software revealed that the oxide thickness over the Fe-containing intermetallic particle does not change with time, which was found to be approximately 10 µm thick. On the other hand, the thickness of the oxide over the S-phase particle increased with time, from about 7 µm at the start of the experiment to about 9 µm when scratching was stopped. Furthermore, the detection of Mo and Mg over the S-phase particle using SEM-EDS after scratching suggests that S-phase particle dissolution is suppressed by oxide formation (Figure 3.25).

However, Cl was also detected over the intermetallic particles and micro pits are observed on the Al matrix from the electron micrograph.

3.4 Conclusions

The mechanism of inhibition of AA2024-T3 by molybdate was investigated using various electrochemical, microscopic, and spectroscopic techniques. The following conclusions can be obtained:

1. Electrochemical polarization curves show that molybdate provides mixed

inhibition at a threshold concentration of 0.1 M. The largest effect observed was

a 250 mV increase in the breakdown potential associated with pitting and a 350

mV decrease in OCP.

2. Corrosion inhibition of AA2024-T3 by molybdate may occur following a two-

step process whereby molybdate is rapidly reduced to form MoO·(OH)2 over the

intermetallic particles, and is subsequently oxidized to intermediate molybdenum

oxides in the presence of oxygen. This in turn may lead to a local acidification,

promoting the condensation and polymerization of molybdate species in solution

6- 4- to form polymolybdate species (Mo7O24 and Mo8O26 ). Furthermore, S-phase

particle dissolution is decreased, suppressing surface copper enrichment and

significantly lowering oxygen reduction kinetics.

3. Electrochemical impedance showed a decrease in the total impedance with the

addition of molybdate in deaerated electrolyte where MoO·(OH)2 was found to be

the predominant species and hydrogen evolution dominate the cathodic kinetics.

In aerated solution, the total impedance increased by 40x in the presence of 73

molybdate. This indicates that corrosion inhibition of AA2024-T3 by molybdate

is oxygen-dependent owing to the protection afforded by Mo(VI) species.

4. The corrosion morphology and kinetics of AA2024-T3 in 0.1 M NaCl solution

with molybdate was investigated using in situ AFM scratching. The addition of

low molybdate concentrations resulted in accelerated attack of S phase particles.

Attack always nucleated at the periphery of the intermetallic particles. Increasing

the molybdate concentration to 0.1 M resulted in drastically improved corrosion

inhibition. Scratching at the maximum tip forces resulted in minor attack of the

matrix and the formation of a hard, insoluble film over both S phase and Fe-

containing intermetallic particles, most likely associated with a mixed

molybdenum oxide consisting of MoO·(OH)2, Mo4O11, and Mo(VI) species.

FIGURES

Figure 3. 2. Chemical equilibrium diagram for (a) 10 µM, (b) 1 mM, and (b) 100 mM 2- TM MoO4 in 0.1 M NaCl solution. Specie diagram generated using Medusa software. The dashed black line signifies maximum solubility.

Continued

Figure 3.2: Continued

Continued

Figure 3.2: Continued

Figure 3.3. Naturally aerated polarization curves for AA2024-T3 in 0.1 M NaCl solution at varying Na2MoO4 concentrations.

Ebr Eocp -200

-400

-600 Potential (mV vs. SCE) vs. (mV Potential -800

-1000 0 20 40 60 80 100 120 140 160

Concentration (mM)

Figure 3.4. Open-circuit potential Eocp and pitting potential Epit with incremental amounts of Na2MoO3.

a b

c d

e f

Figure 3.5. Open-circuit potential macrographs after 1 day exposure to (a) NaCl-only solution and (b-h) 0.1 M NaCl + 125 mM Na2MoO4 solution at varying time. 81

a b

Figure 3.6. Optical micrographs after 1 d exposure to (a) NaCl-only solution and (b) 0.1 M NaCl + 125 mM Na2MoO4 solution at open-circuit potential.

Figure 3.7. SEM-EDS analysis showing maps for Al, Cu, Mo, and O after 30 min exposure in 0.1 M NaCl solution with 125 mM Na2MoO4.

100 O Na 80 Al Mg Al Mo 60 Cl Cu

40 Cu 20 Mo Mg O 0 0 5 10 15 20 25 Position (um)

Figure 3.8. SEM-EDS analysis of S-phase particle after 2 day exposure in 0.1 M NaCl + 125 mM Na2MoO4. (a) Secondary electron micrograph. Red line is reference for (b) EDS line profiles.

100 O Na Al 80 Mo Mn Al Fe 60 Cu

Mo 20 Cu Fe

0 0 5 10 15 20 25 Position (um)

Figure 3. 9. SEM-EDS analysis of Fe-containing particle after 2 day exposure in 0.1 M NaCl + 125 mM Na2MoO4. (a) Secondary electron micrograph. Red line is reference for (b) EDS line profiles.

Figure 3.10. Mo 3d spectra after 1 day exposure in naturally aerated solution with 0.1 M NaCl + 125 mM Na2MoO4.

2 10-5 NaCl injection E = -900 mV SCE app 0 -900 mV (NaCl only) -2 10 -5

-4 10 -5

) 2 -1100 mV

-6 10 -5 i (A/cm i -8 10 -5

-0.0001

-0.00012 MoO -2 injection 4 -0.00014 0 10 20 30 40 50 60

Time [min]

Figure 3.11. Cathodic chronoamperometry of AA2024-T3 in air-bubbled 0.1 M NaCl.

a b

Figure 3.12. Electron micrograph of surface after 2 h exposure to 0.1 M NaCl + 125 mM Na2MoO4 solution at (a) -900 mV SCE- and (b) -1100 mV SCE fixed potential. (c) EDS spectrum of sample surface after -1100 mV SCE fixed potential.

-2 0.0007 MoO injection 4 E = -525 mV SCE app 0.0006

0.0005 MoO -2 injection 4

) -535 mV 2 0.0004 -535 mV

0.0003 NaCl injection i (A/cm i

0.0002

0.0001 > 99% Reduction 0 0 50 100 150

Time [min]

Figure 3.13. Anodic chronoamperometry of AA2024-T3 in air-bubbled 0.1 M NaCl.

a b

Figure 3.14. (a-b) Optical micrographs and (c) SEM-EDS analysis of substrate after Na2MoO4 injection at -535 mV SCE. There is evidence of oxide formation inside corroded areas as suggested by Mo and O maps.

Figure 3.15. Mo 3d spectra of sample surface immediately after injection of concentrated Na2MoO4 to cell solution.

Mo4+ Mo6+

280 930 907 Matrix 740

930 350 870 S-phase 740 920

370 Intensity(Arbitrary Units) Fe-IMC 960 790

200 400 600 800 1000 1200

Raman Shift (cm-1)

Figure 3.16. Raman spectra at different microstructural features of a sample after 2 days of exposure in naturally aerated 0.1 M NaCl solution with 0.125 M Na2MoO4.

-200 125 mM

100 mM -400 10 mM

1 mM - 550 mV SCE E -600 rp - 605 mV SCE

No inhibitor

Potential (mV vs. SCE) vs. (mV Potential -800

-1000 10-10 10-8 10-6 0.0001 0.01

Current Density (A/cm2)

Figure 3.17. Cyclic polarization curves in deaerated 0.1 M NaCl solution with incremental amounts of Na2MoO4.

-400

No Inhibitor Aerated -600

No Inhibitor -800 Deaerated 125 mM MoO 2- 4 Deaerated -1000

125 mM MoO 2- -1200 4 Aerated

-1400 Potential (mV vs. SCE) vs. (mV Potential

-1600

-1800 10-9 10-8 10-7 10-6 10-5 0.0001 0.001 0.01

Current Density (A/cm2)

Figure 3.18. Cathodic polarization curves of AA2024-T3 in aerated and deaerated 0.1 M NaCl solution with 125 mM molybdate.

Figure 3.19. Sample macrograph and SEM-EDS analysis after 1 day exposure in deaerated 0.1 M NaCl + 125 mM Na2MoO4 solution. Substrate was partially exposed indicated by dashed line.

Figure 3.20. Mo 3d spectra of sample surface after 1 day exposure in deaerated 0.1 M NaCl + 125 mM Na2MoO4 solution.

0.1 M NaCl only 106 0.1 M NaCl + 125 mM Na2MoO4 -80

105 -60

104 (degrees) Theta

) 2 -40

1000

-20 IZI (Ohm.cm IZI 100

0 10

1 20 0.01 0.1 1 10 100 1000 104 105

Frequency (Hz)

0 mM - oxygen - Z phase a

125 mM - oxygen - Z phase

0.1 M NaCl only 106 -100 0.1 M NaCl + 125 mM Na2MoO4

105 -80

104 -60

1000 -40

100 -20

10 0

1 20 0.01 0.1 1 10 100 1000 104 105

Frequency (Hz)

0 mM - no oxygen - Z pha b 125 mM - No oxygen - Z phase Figure 3.21. Bode magnitude and phase angle plot of 2024-T3 coupons immersed in 0.1 M NaCl with and without 125 mM Na2MoO4 in (a) aerated and (b) deaerated solution.

-5 Figure 3.22. In situ AFM scratching in 0.1 M NaCl with 10 M Na2MoO4. Scan size=70µm. Top left images is Volta potential map before exposure with z range =500 mV. The other images are topographic maps with z range=200 nm.

-3 Figure 3.23. In situ AFM scratching in 0.1 M NaCl with 10 M Na2MoO4. Scan size=65µm. Top left images is Volta potential map before exposure with z range =500 mV. The other images are topographic maps with z range=200 nm.

Figure 3.24. In situ AFM scratching in 0.1 M NaCl with 0.1 M Na2MoO4. Scan size=35µm. Top left images is Volta potential map before exposure with z range =500 mV. The other images are topographic maps with z range=300 nm.

100

Figure 3.25. Ex situ SEM-EDS analysis after AFM scratching in 0.1 M NaCl with 0.1 M Na2MoO4.

101

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105

CHAPTER 4: CORROSION INHIBITION OF AA2024-T3 BY SODIUM SILICATE

4.1 Introduction

Chromate-based pigments and conversion coatings have been used in the past to prevent corrosion of aluminum-skinned aircraft. They effectively reduce the rate of oxygen reduction over cathodes and moderately hinder anodic dissolution kinetics 1, 2.

Furthermore, chromate conversion coatings (CCC) have the distinct ability to release soluble Cr6+ into the local environment at the presence of a defect providing “self- healing” on various metals 3-5. However, the carcinogenic nature of soluble hexavalent chromate has led to strict regulations that mandate the development of environmentally- friendly alternatives for surface pretreatments and coating pigments 6. In this regard, soluble silicates and silicate-based protection schemes have been studied as corrosion inhibitors for aluminum alloys.

Sodium silicate solutions, commonly known as ‘water glass’, are widely used chemicals for a variety of applications. In particular, they are used in pre-treatments of aluminum alloys as cleaners and corrosion inhibitors 7. Silicates provide corrosion protection to various metals by forming a film of adsorbed species on the surface 8, 9.

The effect of silicates and silica deposits has been reported extensively in the literature. Adsorption models have been proposed for various metals including iron 10-14, copper 15, zinc 16, 17, and aluminum 18-23. In the case of steel, both Fe2+ and Fe3+ 106 participate in the formation of a protective layer by reacting with silicate 24. Yang et al. investigated the adsorption of silicate on synthetic magnetite (Fe3O4) and proposed an adsorption mechanism in accordance with a ligand exchange reaction 10. It was suggested that monodentate and bidentate complexes exist at low surface loading levels with subsequent polymerization occurring on the surface with pH lowering.

A few studies have been conducted on the interaction of silicates with aluminum oxides 18-21, 23. Firment et al. proposed a deposition mechanism for monomeric silica coatings on alumina particles in aqueous solution 23. A random growth mechanism was suggested involving adsorption of silica units where they first contact the surface.

Complete coverage of the surface is achieved as the silica loading increases. Gaggiano et al. studied the interaction of soluble sodium silicates on porous anodic alumina 19-21, 25.

They proposed that aluminosilicate anions in solution react with the aluminate ions formed during oxide dissolution at high pH. The sodium cations act as a coagulating agent between the negatively charged aluminum oxide surface and aluminosilicate anions in solution.

Previous studies involving inhibitor combinations have shown a synergistic effect when combining silicate with other inorganic inhibitors 26-28. Taylor and Chambers developed high-throughput methods to assess binary pairing of 12 inorganic chemistries

26. While some systems exhibited antagonistic behavior, others demonstrated synergies that were comparable or better than the equivalent concentration of Cr6+. These chemistries included pairings of rare earth cations and vanadates with silicate. In addition, studies involving carbon steel corrosion in molybdate and silicate mixed

107 solutions yielded synergistic behavior 28. Interestingly, it was proposed that the silicate anions react preferentially with the surface limiting the incorporation of molybdate species.

Despite the improved corrosion protection afforded by silicate, little is understood about its effect on aluminum alloys. The purpose of this study is to understand the mechanism of corrosion inhibition provided by silicate on aluminum alloy (AA) 2024-

T3. Furthermore, since most inhibiting conversion coatings and pigments act by releasing soluble species into the local aqueous environment, it was of interest to understand the mechanism of inhibition provided by silicate dissolved in aqueous NaCl solution. A mechanism detailing the functionality of silicate at varying pH is proposed.

In addition, the effect of silicate and molybdate mixtures in 0.1 M NaCl solution on the corrosion inhibition of AA2024-T3 was investigated.

4.2 Experimental

4.2.1 Materials and sample preparation: Reagent-grade sodium silicate and sodium chloride were used for all experiments. Solutions were prepared using 18.2 MΩ·cm deionized water. Solution pH was adjusted with diluted solutions of concentrated H2SO4 and NaOH.

Samples of solution heat treated, naturally aged AA2024-T3 (nominal composition 3.9-4.9% Cu, 1.2-1.8% Mg, 0.3-0.9% Mn, 0.5% Fe, 0.5% Si) 29 were mechanically abraded with SiC paper to 1200 grit in a nonaqueous slurry (Blue Lube from Struers) to minimize the onset of corrosion. Samples analyzed microscopically were polished to 1 μm diamond paste. All samples were cleaned with ethyl-alcohol in an 108 ultrasonic bath, air dried, and stored overnight in a desiccator. For electrochemical and free-corrosion experiments, sample dimensions of 2.5 x 2.5 x 0.3 cm were employed.

4.2.2 Potentiodynamic polarization curves and electrochemical impedance spectroscopy:

A platinum-mesh counter electrode and a saturated reference electrode (SCE) were used for all electrochemical experiments. Cathodic and anodic polarization experiments were performed in quiescent, naturally aerated solutions. Cathodic and anodic polarization curves were acquired separately under each condition and were repeated at least in triplicate. The curves shown below are typical curves for each condition. The potential sweep was conducted starting at the open-circuit potential (OCP), using a scan rate of 10 mV/min.

Impedance measurements were performed with and without silicate additions to naturally aerated 0.1 M NaCl solution. A 10 mV sinusoidal voltage was applied at OCP with the frequency ranging from 105 to 10-2 Hz. The impedance response was fitted using an equivalent circuit to extract total resistances and film capacitances. Further details are described below. In addition, all electrochemical experiments were performed with a GamryTM Ref 600 potentiostat using an exposure area of approximately 1 cm2.

4.2.3 Free-corrosion experiments coupled with scanning electron microscopy: AA2024-

T3 coupons were exposed in aerated NaCl solution with and without silicate. Samples immersed in solution were rested against the beaker wall, with the unpolished surface facing the wall and covered with electroplating tape (3M Corp.). At the end of the exposure, all samples were rinsed using 18.2 MΩ·cm deionized water, air-dried, and stored overnight in a desiccator. They were visually inspected and analyzed using either

109 a Quanta 200 or Sirion scanning electron microscope (SEM) integrated with energy dispersive X-ray spectroscopy (EDS).

4.2.5 X-ray photoelectron spectroscopy: XPS measurements were performed to investigate the presence of silicate at the surface of AA2024-T3. Measurements were conducted on samples that were immersed under free-corrosion conditions. XPS spectra were acquired using a Kratos AXIS Ultra spectrometer with a monochromated Al x-ray source operated at 130 W and calibrated to the adventitious carbon 1s peak at 284.6 eV.

Chemical state assessment was achieved by curve-fitting the spectra using CasaXPSTM software. All reported binding energies are rounded to the nearest tenth of an eV and are ascertained a precision value of +/- 0.2 eV.

4.2.5 Atomic force microscopy: In situ atomic force microscopy (AFM) scratching experiments were conducted with a Multimode AFM coupled with a Nanoscope V controller (Bruker Corporation). Naturally aerated electrolyte was pumped into a glass cell with a volume of 0.1 mL at a rate of 10 mL/h. Scratching was performed in contact mode with commercially available Si cantilever tips at a constant scan rate of 2.5 Hz.

The force applied by the tip on the sample surface is proportional to the deflection of the cantilever which is kept constant and controlled by the set-point voltage. The set-point voltage was adjusted from 0.2 to 10 V which correlates to tip-sample pressures ranging from 100 to 525 nN, respectively. Therefore, increasing the set-point voltage is equivalent to increasing the scratching force on the sample. Tips were replaced after each experiment due to wear during scratching. Prior to each experiment, scanning

Kelvin probe force microscopy (SKPFM) was conducted in air allowing for the

110 simultaneous measurement of the surface topography and Volta-potential distribution.

Furthermore, characterization of the scratched surface was carried out before and after each experiment using SEM/EDS.

4.3 Results and Discussion

4.3.1 Aqueous solution chemistry: Many efforts have been reported in the literature to describe the chemistry of soluble silicates 20, 30-35. It has been shown that the anionic complexity of silicate in solution is related to concentration and the silicon-to-cation ratio

(i.e. SiO2/Na2O). In this study, SiO2/Na2O was equal to 1. The solution pH was measured after progressive addition of Na2SiO3 to 0.1 M NaCl solution. As illustrated in Figure 4.1, increasing the silicate concentration increases the solution pH. The increase in pH can be explained by the formation of silicic acid and hydroxyl ions as described in the following chemical equation 35:

( ) Eq. 4.1

29Si-NMR has been recognized as the best method to study the anionic speciation in silicate solutions 20, 30, 31, 33. However, performing these measurements is difficult.

Attempts to conduct 29Si-NMR in this study failed due to intense background resonance.

Nevertheless, Gaggiano et al. used solid-state 29NMR to study anionic species in silicate

20 solution at 0.1 M concentration . At SiO2/Na2O =1 the main species in solution were

found to be silicate monomers, which can be represented as ( ) according to

Swaddle 32.

The effect of pH on the speciation formed in dilute sodium chloride solution was investigated using chemical equilibrium diagrams computed by MedusaTM. Embedded 111 within the software is a database that contains chemical equilibrium constants at 25°C.

Figure 4.2 shows the generated speciation diagram for 1 mM and 100 mM silicic acid in

0.1 M NaCl solution. According to Figure 4.2a, the dominant species in solution

containing 1 mM silicic acid is ( ) at pH greater than 10.3. At lower pH, the dominant species in solution is silica, or SiO2. Similar results are observed in solution containing 100 mM silicic acid (Figure 4.2b). However, the threshold pH for silica formation shifts to around 12.2. Above this point, the predominant species in solution are

monomeric ( ) and ( ) anions. As discussed below, the mechanism of inhibition by silicate is highly dependent on the monomeric silicate anions present in solution.

4.3.2 Polarization in aerated sodium chloride solution: Figure 4.3 shows naturally aerated polarization curves in 0.1 M NaCl solution with varying Na2SiO3 concentration.

The natural pH measured in the silicate containing solutions was between 10.7 and 13.6.

In addition, two control experiments were performed in as-prepared (pH 6) and alkaline

(pH 11) solution. Note that these curves are typical results taken from the replicated experiments. In inhibitor-free solution at a high pH, the corrosion potential, Ecorr, decreases because of a large increase in the passive current density associated with the amphoteric nature of the aluminum oxide film. Furthermore, the curves show that the pitting potential, Epit, remains unaffected after increasing the pH of the inhibitor-free solution. The polarization curves show that the higher apparent Epit at the higher pH is caused by the intersection of the higher passive current with the unchanged current associated with pit growth. In contrast, the addition of silicate increases Epit. The largest

112 effect was observed at concentrations greater than 25 mM where Epit increased by 1 V.

These observations indicate that silicate imparts strong anodic inhibition. However, the passive current density increased by an order of magnitude at the highest concentrations tested. The observed increase in passive current density can be attributed to an increase in pH as the silicate concentration in solution is increased, as was found for the inhibitor- free solution. A decrease in the corrosion potential, Ecorr, was also observed as a result of an increase in the passive current density with higher silicate additions. The largest effect was observed at the highest concentration tested (0.1 M), corresponding to a 100 mV decrease in Ecorr. Figure 4.4 summarizes the effect of silicate concentration on Ecorr and

Epit. Epit is taken as the point at which the current increases sharply from the passive current density. No effect was observed at 1 mM, so 10 mM can be inferred to be the threshold concentration for corrosion inhibition.

It is interesting to note that silicate also increased the cathodic limiting current densities associated with ORR. The limiting current density can be expressed by the following equation:

Eq. 4.2

where and is the boundary layer thickness and diffusion coefficient of oxygen in through the boundary layer in solution, is the number of electrons transferred to reduce

1 mole of oxygen, is the Faraday constant, and is the bulk solution concentration of oxygen. According to Scully et al. 36, in the presence of a film on the surface (oxide, oxyhydroxide, or hydroxide), another term can be added to Eq. 4.2 to account for the diffusion of oxygen through the film : 113

Eq. 4.3

where and are the average film thickness and diffusion coefficient of oxygen through the film. According to Eq. 4.3, the limiting current density is directly proportional to the diffusivity of oxygen through the film on the surface. As discussed below, exposure of a sample to silicate solution results in the dissolution of the aluminum-oxide and formation of a silicate-based thin film on the surface. Thus, the increase in the ORR kinetics observed in Figure 4.3 can be attributed to an increase in the diffusivity of oxygen through the altered surface film that forms after exposure to silicate solution (neglecting differences in film thickness). Furthermore, hydrogen evolution kinetics are shifted to lower potentials, an effect attributed to a decrease in the reversible potential of hydrogen evolution as the pH increases with the addition of silicate.

4.3.3 Electrochemical Impedance Spectroscopy: Figure 4.5 shows the time evolution of the impedance response in 0.1 M NaCl solution with and without 25 mM Na2SiO3.

Experiments were conducted in inhibitor-free solution at pH 6 (natural pH) and pH 12.8

(pH-adjusted). The natural pH of 0.1 M NaCl solution with 25 mM Na2SiO3 was measured to be around 12.8. In as-prepared 0.1 M NaCl solution (pH 6), the impedance decreases with time for the first 8 days and then changes little at longer times (Figure

4.5a). At higher pH (Figure 4.5b), the impedance is very low after the first day of exposure, but then increases with time. The increase in impedance could be attributed to the formation of corrosion product over the alloy surface which may slightly increase the corrosion resistance of the alloy. With the addition of silicate, a number of differences

114 are observed. From the beginning, the impedance is higher in the presence of silicate and it increases with time. The low frequency impedance increases by an order of magnitude after 8 days of exposure. In addition, the capacitive reactance increases over a larger domain in the phase angle plot indicating improved corrosion inhibition.

In order to extract the total resistance and film capacitance, the spectra were fitted using a simplified Randles circuit that included a constant phase element (CPE) to represent the film capacitance. The true capacitance was derived from the CPE parameters using the following equation 37:

Eq. 4.4 where Y and α represent the CPE magnitude and exponent, respectively, Rt is the extracted total resistance, and Cf is the calculated film capacitance. The measured total resistance and calculated film capacitances are summarized in Figure 4.6. It should be noted that these experiments were not replicated. In the absence of inhibitor (pH 6 solution), the total resistance remained relatively constant at around 6 kΩ·cm2, while the film capacitance increased from 130 µF·cm-2 to 310 µF·cm-2 after 8 days of exposure.

With the addition of silicate, the total resistivity increased almost two orders of magnitude from 74 kΩ·cm2 to 5.7 MΩ·cm2. The film capacitance in solution containing silicate slightly decreased from 5.6 µF·cm-2 to 4.1 µF·cm-2.

A difference in the total of three orders of magnitude compared to the negative control after 8 days of exposure indicates a strong inhibition by silicate. Furthermore, the results show a time-dependent evolution of silicate inhibition might be associated with an

115 increase in oxide film thickness on the surface as indicated by a slight decrease in the film capacitance with time.

4.3.4 Morphology of AA2024-T3 after free corrosion exposure: Since the largest corrosion inhibition effect by silicate was observed at 25 mM concentration, it was of interest to study the morphology of AA2024-T3 after exposure in 0.1 M NaCl + 25 mM

Na2SiO3 solution under free corrosion conditions. Figure 4.7a shows macrographs after 1 day exposure in naturally aerated solution with and without sodium silicate. In the absence of inhibitor, there are clear signs of corrosion on the sample surface. Optical microscopy reveals evidence of attack over the intermetallic particles and matrix (Figure

4.7b). With the addition of silicate, the surface is free of corrosion at the magnification shown. No attack is observed over the larger intermetallic particles or matrix. However, there are signs of deposition over the smaller intermetallic particles, which were made evident with subsequent SEM-EDS analysis.

Figure 4.8 shows SEM-EDS analysis of a sample surface after 3 days of exposure in 0.1 M NaCl solution with 25 mM Na2SiO3. Comparison of the before and after secondary electron images shown in Figure 4.8a reveals partial S-phase particle dissolution. Interestingly, higher concentrations of Si and O were detected over the partially dissolved S-phase particles using EDS. The detection of Mg at several particles

(Figure 4.8b) further indicates only partial dissolution of the S-phase. Si and O were also detected within pits on the surface and at the larger Fe-containing intermetallic particles.

The susceptibility of S-phase particle dissolution has been reported extensively in literature 38-41. The S-phase is initially active with respect to the matrix, undergoing

116 preferential dissolution of Mg and Al. Although there is evidence to suggest that S-phase particle dissolution is suppressed, it is likely that dissolution of several particles occurs in response to the extreme alkaline conditions that form with the addition of silicate.

Furthermore, the aggressiveness of the high pH solution should create conditions that result in passive film dissolution. However, the absence of attack over the Al matrix suggests that silicate enhances passivity. In addition, no attack was observed on the matrix adjacent to the Fe-containing intermetallic particles, which is typically observed when 2024-T3 corrodes in naturally aerated, chloride solutions 42-46.

4.3.5 X-ray photoelectron spectroscopy: XPS was used to characterize the surface after 3 days of exposure in 0.1 M NaCl solution with and without 25 mM Na2SiO3. Table 4.1 shows the results for the chemical quantification analysis of the surface after exposure in both solutions. The analysis area was approximately 0.21 mm2. Cl was not detectable on the surface in both cases. Unlike in the presence of silicate, no amounts of Si or Na were detected on the surface exposed to inhibitor-free solution. In addition, Al and Cu are present at much lower concentrations compared to the negative control, which suggests that there is a thicker film on the surface. Furthermore, no Mg was present on the surface after exposure to inhibitor-free solution. The presence of Mg on the surface of the sample exposed to silicate solution may originate from the exposed S-phase particles that remain protected on the surface. The absence of Mg on the surface of the sample exposed to the inhibitor-free solution might result from dissolution or from the strong build-up of corrosion product over the intermetallic particles.

117

Figure 4.9a shows the Si 2p spectra after exposure in 0.1 M NaCl solution with and without 25 mM Na2SiO3. In the presence of silicate, a peak associated with aluminosilicate species is present at a binding energy of 102.2 eV 47-52. According to the literature, this binding energy is consistent with albite (Na(AlSi3O8)) and natrolite

+ 50 (Na2(Al2Si3O10)·2H2O), ionic networks containing Na with tetrahedral aluminum .

This is further supported by the presence of an O 1s peak at a binding energy of 531.7 eV

(not shown), which is also consistent with these ionic network species 50. No Si 2p peak is observed in the case of inhibitor-free solution. Figure 4.9b shows the Al 2p spectra after exposure. In inhibitor-free solution, two peaks corresponding to Al3+ are observed

53, 54 at binding energies of 74.1 and 75.2 eV . The former peak is consistent with Al2O3,

54 and the latter can be ascribed to Al(OH)3 . In the presence of silicate, one peak corresponding to Al3+ is observed at a binding energy of 74.1 eV, and another peak at

71.0 eV is observed corresponding to Al metal. According to the literature, the Al 2p binding energies associated with aluminosilicate compounds range from 74.0 eV to 74.8 eV 47, 48, 50, 51. In particular, the binding energies of the ionic network species are reported to be around 74.14 eV and 74.05 eV (all measured energies in this report are rounded to the nearest tenth of an eV). Thus, it is likely that the Al 2p peak observed after silicate exposure is a convolution of several signals corresponding to Al in Al2O3 and aluminosilicate films on the surface. Furthermore, the detection of Al metal after exposure to silicate suggests that the resulting film is thinner than the negative control.

The observed difference in film thickness may arise from the enhanced passivity observed in the presence of silicate.

118

As mentioned above, no Na or Mg was detected over the sample exposed to inhibitor-free solution. Figure 4.9c and Figure 4.9d show the Na 1s and Mg 1s spectra, respectively. In the presence of silicate, each spectrum exhibited a single peak at binding energies of 1071.8 eV and 1303.6 eV, respectively. Classification of the latter peak is consistent with MgO found in magnesium-silicate compounds 25, 55. Attempts to classify the former peak proved to be difficult since the corresponding binding energy overlapped with many ionic compounds 25.

Figure 4.10 shows the sputter depth profiles of the surface film after 3 days exposure to 25 mM Na2SiO3 in 0.1 M NaCl solution. It should be noted that the Si profile was obtained by monitoring the Si 2s transition region during sputtering since it exhibited less plasma interference compared to the Si 2p photoelectron band. The initial dip in the C concentration after the first etch is attributed to carbon contamination from exposure to the atmosphere. According to the profile, Si, Al, and O are present throughout the film. Na was also detected throughout the film, albeit at much smaller concentrations. Furthermore, analysis of the Si 2s spectrum during etching revealed a single peak consistent with aluminosilicate.

4.3.6 Effect of Solution pH: It was of interest to investigate the effect of solution pH on the inhibition mechanism by silicate. According to the chemical diagrams depicted in

Figure 4.2, the dominant species in solution is SiO2 except in highly alkaline solutions.

Furthermore, the presence of monomeric anionic silicate decreases steadily as the pH is lowered.

119

Figure 4.11 shows naturally aerated polarization curves in pH adjusted 7.3 (near- neutral) and 4 (acidic) solutions. In near-neutral solution, silicate has no apparent effect except at low potentials. Again, these are typical polarization curves for each solution.

The limiting current density for ORR is slightly increased and the potential for hydrogen evolution is shifted to more negative potentials by 300 mV in the presence of silicate. In acidic solution, there is a larger increase in ORR kinetics observed at the highest- concentration tested. No effect is observed on the hydrogen evolution kinetics.

Free corrosion experiments were also performed under these conditions to further assess the inhibition behavior by silicate at low pH. Figure 4.12a shows a macrograph of a sample immersed in pH 7.3 solution containing 0.1 M NaCl + 25 mM Na2SiO3. No signs of attack were observed by visual inspection after exposure for 3 days in solution.

Figure 4.12b shows an electron micrograph of the sample surface. There is evidence of minor attack on the Al matrix and at the periphery of some of the intermetallic particles.

Furthermore, the electron micrographs show preferential deposition of sub-microscopic

Si particulates over the intermetallic particles (Figure 4.12c).

Figure 4.13a shows a macrograph image of a sample immersed in a pH 4 solution of 25 mM Na2SiO3 in 0.1 M NaCl. After 3 days of exposure, the sample surface is covered with a heterogeneous film that is gel-like in appearance and partially covered with a white film. Subsequent SEM analysis reveals corrosion product and preferential film formation over the intermetallic particles (Figure 4.13b). Interestingly, there is a large concentration of Si detected over an S-phase particle. SEM-EDS analysis was also conducted over an area on the surface that was covered with a white film, revealing that

120 the film is thick and porous in nature (Figure 4.13c). Characterization of the film using

EDS shows that it is composed of Si and O.

Surface characterization was also performed using XPS. Table 4.2 shows the results for the chemical quantification analysis after 3 days of exposure in near-neutral pH and acidic solution. It should be noted that Cl was not detectable on the sample surface after exposure to either solution, and a significantly lower concentration of Na was detected in comparison to alkaline conditions (Table 4.1). In addition, there is a larger concentration of oxidized Mg and Cu on the surface of the sample exposed to the acidic solution, which is consistent with increased corrosion activity.

Figure 4.14a shows the Si 2p spectra after 3 days immersion in near-neutral and acidic solution. In the case of near-neutral pH solution, a peak associated with silicate species is present at a binding energy of 102.8 eV 47-52. According to the literature 50, the binding energy of talc (Mg3(Si4O10)(OH)2) and kaolinite (Al2Si2O5(OH)4) are 102.93 eV and 102.78 eV, respectively. Thus, it is probable that the Si 2p peak observed after exposure to silicate in near-neutral solution is a convolution of signals corresponding to

Si atoms in these silicate species. This is further supported by the presence of an O 1s peak at a binding energy of 532.1 eV (not shown), which is also consistent with talc and kaolinite species 50. The corresponding Al 2p spectrum shown in Figure 4.14b reveals three peaks associated with Al3+ at binding energies of 73.7 eV, 74.7 eV, and 75.7 eV.

The peaks corresponding to the lowest and highest binding energy are attributed to aluminum- oxide and hydroxide, respectively. However, the intermediate peak at 74.7 eV is consistent with kaolinite species 47, 48, 50, 51.

121

In acidic solution, a number of differences are observed. The Si 2p spectrum corresponding to acidic solution in Figure 4.14a reveals two peaks at a binding energy of

102.5 eV and 103.3 eV. The peak associated with the lowest binding energy is consistent

47-52 with aluminosilicate species , specifically pyrophyllite (Al2(Si4O10)(OH)2), while the

48, 51 peak at the highest energy corresponds to SiO2 . This is further supported by the presence of an O 1s peak at a binding energy of 531.9 eV (not shown), which is also consistent with pyrophyllite species 50. The corresponding Al 2p spectrum shown in figure 4.14b exhibits two narrow peaks at a binding energy of 74.5 eV and 75.7 eV, and a broad peak at a binding energy of 78.5 eV. The peak associated with 74.5 is consistent with pyrophyllite species 47, 48, 50, 51, and the peak at 75.7 eV is associated with aluminum hydroxide 53, 54. The broad peak observed at 78.5 eV is associated with the Cu 3p transition level.

According to the chemical diagrams depicted in Figure 4.2, the dominant species in lower pH solution is SiO2. However, according to the XPS results shown in Figure

4.14, the majority of Si on the surface is present in the form of aluminosilicate species. It is possible that aluminosilicate species deposit over the intermetallic particles and local cathodes at higher pH. The details of this effect are described below.

4.3.7 Mechanism of Inhibition: In what follows, inhibition mechanisms are proposed for alkaline, neutral, and acidic environments. In alkaline conditions, the aluminum oxide surface is chemically unstable and dissolves in solution. The polarization curves and impedance spectra in alkaline solution suggest that silicate forms a film over the alloy surface, thereby enhancing passivity. The formation of a silicate film likely occurs via an

122 adsorption mechanism as supported by the XPS data which showed evidence of silicate species on the surface. Furthermore, the adsorbed film is thin in nature as no morphology change could be observed over the Al matrix after exposure (Figure 4.7).

However, attack was observed over the S-phase particles as made evident by the SEM images shown in Figure 4.8.

The reduced attack observed over the Al matrix can be explained by the formation of an aluminosilicate thin film over the surface. In alkaline solution, the Al oxide film is

32, 56, 57 chemically unstable and dissolves to form aluminate ions ( ) :

( ) Eq. 4.5

The monomeric silicate anions in solution react with the ( ) species to form aluminosilicate anions. This step is described by the following chemical reaction proposed by Swaddle 32:

( ) ( ) ( ) ( ) ( ) Eq. 4.6

It has been reported that the isoelectric points of oxide-covered aluminum and copper are around 9.5 58, 59. Therefore, in alkaline silicate solutions both the oxide surface and the aluminosilicate anions are negatively charged. As suggested by Iler, the deposition of silicate anions on negatively charged oxide surfaces requires the presence of a coagulating agent, usually a small concentration of polyvalent metal ions 35. Gaggiano et al. 19, 20 suggested that it is possible that Na+ ions detected by XPS within the film adsorb on the negatively charged oxide surface and behave as a coagulating agent. The Na+ ions on the surface coordinate with the hydroxyl group of the aluminosilicate anions acting as a bridge between the anions and the surface. Furthermore, Na might behave this way

123 throughout the film as indicated by the sputter depth profiles shown in Figure 4.10. The resulting silicate layer grows with time and provides a protective film over the aluminum alloy as supported by the electrochemical polarization curves and impedance spectroscopy. This mechanism has been proposed recently by Gaggiano et al. 19, 20.

The susceptibility of S-phase particle dissolution has been reported extensively in literature 40, 44, 60-64. The S-phase is active to the matrix, undergoing preferential dissolution of Mg and Al 40. Although there is evidence to suggest that S-phase particle dissolution is suppressed, attack was still observed over several particles. It is possible that attack occurred in response to the extreme alkaline conditions that form in the presence of silicate. However, according to Figure 4.8, there is evidence of Si present over the attacked S-phase particles. Silicate might react with the dissolved Mg and Al ions in solution as the S-phase particles undergo preferential dissolution. This is supported by the stability diagram shown in Figure 4.15 of a ternary alloy system consisting of Al, Cu, and Mg in silicate solution. The plot was generated with OLI

AnalyzerTM, computer modeling software, which utilizes a thermodynamic framework to simulate multi-phase aqueous systems. Accordingly, there is favorable formation of magnesium-silicate and alkali-aluminosilicate species in the high pH regime. This is supported by the detection of oxidized Mg and aluminosilicate species by XPS in alkaline conditions (Figure 4.9). It should be noted that no Mg was detected after exposure to inhibitor-free solution. Furthermore, the suppressed dissolution of S-phase particles is reflected by an increase in the pitting potentials shown in Figure 4.3.

124

In more acidic media, silicate provides poor inhibition as shown in the polarization curves shown in Figure 4.11. Nevertheless, a mechanism is proposed to account for the effect of silicate at near-neutral and acidic solution.

At near-neutral pH, the aluminum surface is passivated except at the intermetallic particles where the oxide is thinner. These Cu-rich intermetallic particles support oxygen reduction leading to a local increase in pH and an increase in the available monomeric silicate anions adjacent to the intermetallic particles. These reactions are described by the following equations:

Eq. 4.7

( ) ( ) Eq. 4.8

The local availability of silicate anions leads to the deposition of silicate- and silica- based derivatives over the intermetallic particles as depicted in the electron images shown in Figure 4.12. With regards to S-phase particles, it is likely that a precursor for deposition is selective dissolution of Mg and Al to generate metallic ions in solution:

Eq. 4.9

Eq. 4.10

The metallic ions then chemically react with the monomeric silicate anions that form over the intermetallic particles in response to an increase in local pH. These processes can be simplified and illustrated by the following chemical equations in the case where the form

of silicate species present in solution is of the type ( ) :

( ) ( ) ( ) Eq. 4.11

( ) ( ) ( ) Eq. 4.12

125

The resulting metallic-silicate species deposit over the intermetallic particles and perhaps block further attack. However, according to the electron images shown in Figure 4.12, the intermetallic particles are only partially covered with deposits, implying that catalytic activity could still occur over the uncovered areas. This is supported by the polarization curves shown in Figure 4.11 whereby virtually no effect is observed with the addition of silicate at lower pH solution. It is possible that not enough silicate anions form over the intermetallic particle to deposit over the whole area.

Exposure to silicate solution at low pH results in the formation of a heterogeneous film. Characterization of the film by XPS revealed that it is composed of both silica and aluminosilicate species. According to Iler 35, deposition of colloidal silica particles on a surface from a dilute solution may occur when conditions of pH and salt concentration are close to those causing coagulation and precipitation. Deposition occurs when the colloidal particles collide and combine with the solid surface, resulting in the formation of a porous film that is often white and opaque when dried. This was in fact observed after exposure to silicate solution at low pH as shown in Figure 4.13. A porous film rich in Si and O was identified with subsequent SEM-EDS analysis, and was observed to partially cover the surface after exposure.

Deposition of aluminosilicate anions occurs by a process entirely different from the one above. In acidic environments, the entire surface of the aluminum alloy is activated including the intermetallic particles and matrix. During this process, the interfacial electrolyte pH increases as described by the following equation 57:

Eq. 4.13

126

Under these conditions, the alloy is subjected to uniform corrosion where the oxygen reduction reaction is supported by the entire surface. It is possible that this local increase in pH increases the monomeric silicate anion concentration near the surface of the alloy as described by Eq. 4.8. Afterwards, it is proposed that the monomeric silicate anions chemically react with the activated aluminum species forming aluminosilicate anions.

This process can be described by the following equation, which may be an over- simplification of the actual process:

( ) ( ) Eq. 4.14

This process is similar to the reaction described by Eq. 4.11. However, the difference in mechanism lies in the fact that in the more acidic regime the whole aluminum surface becomes activated. This leads to the formation of a mixed silica/aluminosilicate film over the surface that is porous in nature, providing virtually no corrosion protection as indicated by the polarization curves shown in Figure 4.11b.

4.3.8 Synergistic inhibition of silicate and molybdate ion species: The inhibition mechanism by molybdate alone was discussed in the previous chapter. It was shown that molybdate provides mixed inhibition at a threshold concentration of 0.1 M. The largest effect observed was a 250 mV increase in the pitting potential and a 350 mV decrease in

OCP. It was proposed that corrosion inhibition of AA2024-T3 was oxygen-dependent involving two steps whereby molybdate is rapidly reduced to form MoO2 over the intermetallic particles and is subsequently oxidized to MoO3 in the presence of oxygen.

Furthermore, S-phase particle dissolution is decreased, suppressing surface copper enrichment and significantly lowering oxygen reduction kinetics.

127

Figure 4.16 shows naturally aerated polarization curves at various mixtures of

Na2MoO4 (compound A) and Na2SiO3 (compound B) in 0.1 M NaCl solution.

Synergistic behavior was observed at all inhibitor concentrations tested. At a 1:1 mixing ratio with 0.002 M total inhibitor concentration, Epit was observed to increase by 100 mV.

Increasing the total inhibitor concentration to 10 mM increased Epit by 200 mV and decreased Ecorr by 180 mV. In addition, an arbitrary mixing ratio of 5:1 was selected to assess the contribution of each inhibitor to the synergism observed when mixing the inhibitors together. Adding more silicate in solution than molybdate resulted in a 150 mV increase in Epit and a 150 mV decrease in Ecorr. On the other hand, adding more molybdate to solution slightly decreased Ecorr and increased Epit by 100 mV. This suggests that silicate has a greater effect on the corrosion inhibition of 2024-T3 when mixed together in solution with molybdate.

It has been reported that excessive solubility of inhibitor pigments added to an organic coating can promote osmotic blistering, one of the main degradation mechanisms of organic coatings 65. Therefore, it is necessary to incorporate inhibitors into coating systems that require low concentrations for corrosion inhibition. In this regard, it is not suitable to use silicate or molybdate as stand-alone inhibitors in organic coatings since they require large concentrations to impart inhibition. However, according to the polarization curves shown in Figure 4.16, significant inhibition is observed at low total inhibitor concentrations when silicate and molybdate are mixed together in solution. For instance, a 100 mV increase in Epit was observed in a 1:1 mixture containing 0.002 M total inhibitor concentration. In the case where silicate or molybdate were used

128 independently in solution, no inhibition was observed at inhibitor concentrations below

0.01 M and 0.1 M, respectively. Therefore, it can be concluded that the threshold concentration for corrosion inhibition is significantly lowered when silicate and molybdate are mixed together in solution.

To further assess the synergistic behavior between silicate and molybdate, it was of interest to study the free corrosion morphology of a sample exposed to 0.1 M NaCl

-3 -3 solution with 10 M Na2SiO3 + 10 M Na2MoO4. The pH of this solution was measured to be approximately 11. No signs of attack were observed by visual inspection after 10 days of exposure (Figure 4.17a). However, subsequent SEM-EDS analysis revealed deposition over several intermetallic particles with trace amounts of Si, Mo, and O

(Figure 4.17b). Furthermore, Mg was detected at S-phase particles, which suggests that dealloying of these particles is suppressed. SEM-EDS analysis of a large Fe-containing intermetallic particle revealed the presence of sub-microscopic deposits that segregated along the polishing lines (Figure 4.16c). No Mo was detected at the particle using EDS.

It should be noted that the limitations of EDS may prevent detection of a small amount

Mo over the particle.

Characterization of the sample surface after 10 days of exposure was performed using XPS. Figure 4.17a shows the Si 2p spectrum, which reveals a single peak consistent with aluminosilicate species at a binding energy of 102.6 eV 47-52. This is supported by the Al 2p spectrum shown in Figure 4.17b, which reveals two peaks at binding energies of 74.7 and 73.8 eV. The former peak, pertaining to 92% of the Al present on the surface, is associated with aluminosilciate species 47, 48, 50, 51. The latter

129 peak is consistent with Al2O3. In addition, the O 1s spectrum in Figure 4.17c exhibits two peaks at a binding energy of 531.0 and 531.8 eV, which are consistent with Al2O3 and aluminosilicate species, respectively 50. Table 4.3 summarizes the results for the chemical quantification analysis. Interestingly, no Mo or Mg was detected on the surface, and a significant amount of Na was detected by XPS. However, attempts to classify the Na 1s peak at 1071.5 eV (not shown) proved to be difficult since the corresponding binding energy overlapped with many ionic compounds 25.

The results indicate that silicate anions react preferentially with the surface, thereby limiting the interaction between the surface and the molybdate anions in solution.

However, it is possible that molybdate reduces over the intermetallic particles during the initial stages of exposure. According to the OCP measurements shown in Figure 4.19, the global OCP of the working electrode increases steadily from an initial value of -1.2 V

2- SCE. This is below the reversible potential for MoO4 /MoO2 reduction, which was calculated to be -570 mV SCE. Furthermore, SEM-EDS analysis confirmed the presence of Mo and O over both S-phase and Fe-containing intermetallic particles (Figure 4.17).

Therefore, it is likely that molybdate anions reduce over the intermetallic particles prior

2- to the deposition of a thin aluminosilicate film over the surface. Reduction of MoO4 over the intermetallic particles was also supported by the results from the previous chapter.

The synergistic effect between molybdate and silicate anions in solution can be explained by the individual interaction of the respective species over the alloy surface.

Chemical diagrams generated using MedusaTM and OLI AnalyzerTM (not shown) did not

130 reveal complexation between the molybdate and monomeric silicate anions in solution.

Thus, it is proposed that molybdate preferentially reduces over the intermetallic particles during the early stages of corrosion. During this process, the Al matrix will undergo dissolution in response to a high pH, leading to the formation of a thin aluminosilicate film over the surface with the help of Na cations as described above. Furthermore, the results suggest that adsorption also occurs over the Mo-covered intermetallic particles since no Mo was detected using XPS. Note that the limited depth resolution of XPS may not allow for the detection of reduced Mo species that forms prior to the adsorption of a thin film over the intermetallic particles. Thus, the complete coverage of the alloy matrix, and the reinforcement of a mixed Mo/Si-based film over the intermetallic particles, enhances passivity. The improved resistance to localized corrosion is also reflected by an increase in the pitting potential as shown in the polarization curves of

Figure 4.16.

4.3.9 In Situ Atomic Force Microscopy: In situ AFM scratching involves rastering a region in contact-mode with a Si cantilever at sufficiently high forces to destabilize passivity 39. This technique provides real time information of the early stages of localized corrosion with sub-microscopic resolution. When analyzing the evolution and kinetics of attack, three things have to be considered: the force that initiates attack, the time for attack to nucleate, and the corrosion morphology. Higher forces and longer nucleation times imply greater inhibition performance. This technique was used to further investigate the complex interaction between alloy microstructure and silicate ions in 0.1 M NaCl solution.

131

Figure 4.20 shows the Volta potential and topography map prior to the experiment, and the chronological sequence of attack during scratching in 0.1 M NaCl +

-5 10 M Na2SiO3. The Volta potential map of the scratched region reveals the presence of two Fe-containing intermetallic particles and several S-phase particles which were identified using SEM-EDS. It should be noted that the topography image captured by

SKPFM does not resolve the S-phase particles that are present on the surface. Although these particles are harder than the Al matrix, the loss of material by corrosion may

45 balance the slower rate of removal during polishing . Hence, they are initially coplanar with the matrix and are not observable in the topography image prior to exposure.

Therefore, the high contrast obtained from the Volta potential map allows for better detection of S-phase particles on the surface. Attack was observed during 15 minutes of

OCP stabilization prior to scratching. This effect may be attributed to an increase in pH with the addition of silicate to solution. The 0 h topography image shows that the S- phase particles suffer complete removal from the surface prior to scratching. After 1 h of scratching at relatively low forces, further attack is observed in the form of trenching at the periphery of the larger Fe-containing intermetallic particles. With increasing force, attack becomes more severe as the intermetallic particles are slowly removed from the surface. It is not clear whether the particles were rapidly dissolved or removed from the surface as a result of undercutting and dislodging from continuous scratching. At the 3 h mark, corrosion product is deposited at the periphery of a pit following the formation of a hydrogen bubble (indicated by an arrow in Figure 4.20). Scratching at the maximum tip

132 force resulted in complete coverage of the pit by corrosion product. Furthermore, no attack was observed on the Al matrix during the scratching experiment.

It should be noted that no attack was observed during OCP stabilization in inhibitor-free NaCl solution (not shown). The form of attack observed during scratching in inhibitor-free solution was trenching at the periphery of the larger Fe-containing particle, and complete removal of the S-phase particle after 20 minutes of scratching at a low constant force (100 nN). Similar results were obtained by previous researchers39, 42.

Hence, the above results suggest that the resistance to localized corrosion of AA2024-T3 is exacerbated by the addition of 10-5 M silicate.

Increasing the silicate concentration to 0.025 M significantly decreased the dissolution rate of AA2024-T3. This was the optimum concentration for corrosion inhibition observed by the electrochemical polarization curves shown in Figure 4.3.

Figure 4.21 shows the SKPFM images and a secondary electron image of the scratched area prior to the experiment, which contains an S-phase particle and Fe-intermetallic particles. It should be noted that only the latter particles are observable in the SKPFM images. Figure 4.22 shows topography images captured at different times during the scratching experiment. In contrast to the previous condition, no localized attack was observed during OCP stabilization and throughout the scratching experiment. Note that the dark shade in the topography images is not attributed to attack, but rather to an imaging artifact associated with drifting of the AFM scanner. Increasing contrast is also observed over the intermetallic particles with increasing scratching force suggesting that the particles protrude higher from the surface with respect to the matrix. It is not clear

133 whether this effect is attributed to silicate deposition over the intermetallic particles, or to the uniform dissolution of the Al matrix in response to a high pH with the addition of silicate, or both. Ex situ SEM-EDS analysis (not shown) confirmed the absence of localized attack over the surface and the presence of Si over the S-phase particles.

Increasing the silicate concentration to 0.1 M yielded similar results. Figure 4.23 shows the Volta potential and topography map prior to the experiment, and the topography image captured at the start and at the end of scratching in 0.1 M NaCl solution with 0.1 M Na2SiO3. Similarly, the increased contrast over the intermetallic particles with increasing scratching force could be attributed to silicate deposition, or to the uniform dissolution of the Al matrix in response to a high pH. Figure 4.24 shows the subsequent SEM-EDS analysis of the scratched area after the experiment. The secondary electron image shows evidence of partial localized attack on the Al matrix within and outside the scratched region (labeled in Figure 4.23 by arrows). The detection of Mg over the S-phase particles suggests that S-phase dissolution was suppressed.

Furthermore, the presence of Si may indicate deposition of talc (Mg3(Si4O10)(OH)2) over the particles following a similar mechanism as described above.

4.4 Conclusions

The mechanism of corrosion inhibition of AA2024-T3 by silicate was investigated using electrochemical, microscopic, and spectroscopic techniques. The following conclusions can be obtained:

1. Electrochemical polarization curves show that silicate provides strong anodic

inhibition at a threshold concentration of 0.01 M. The largest effect observed was 134

a 1 V increase in the breakdown potential associated with pitting at 0.025 M

silicate concentration.

2. Corrosion inhibition of AA2024-T3 by silicate in alkaline solution conforms to

the mechanism suggested by Gaggiano et al. for pure aluminum 19, 20. It was

proposed that aluminosilicate is formed by the reaction of silicate anions in

solution and the aluminate ions that form during oxide dissolution. The Na+ ions

adsorb to the negatively charged surface and coordinate with the hydroxyl group

of the aluminosilicate anions, thereby forming a protective thin-film over the Al

matrix. Furthermore, S-phase particle dissolution is suppressed owing to the

formation of magnesium-silicate and alkali-aluminosilicate species that deposit

over the particles after the onset of corrosion.

3. At near-neutral pH, silicate partially blocks attack of the intermetallic particles by

a precipitation mechanism that results in the formation of silica- and silicate-

based derivatives. In acidic solution, it was shown that activation of the

aluminum surface promotes the formation of a mixed silica/aluminosilicate film

over the surface that is porous in nature, providing very little corrosion protection.

4. Strong synergy was observed when mixing silicate and molybdate together in

solution. The threshold concentration for corrosion inhibition was lowered by

almost an order of a magnitude in comparison to silicate alone. It is proposed that

molybdate reduces over the intermetallic particles prior to the preferential

deposition of a thin aluminosilicate film over the surface.

135

5. The corrosion morphology and kinetics of AA2024-T3 in 0.1 M NaCl solution

with silicate was investigated using in situ AFM scratching. The addition of low

silicate concentrations resulted in accelerated attack of S phase particles. Attack

also nucleated at the periphery of the larger Fe-containing intermetallic particles.

Increasing the silicate concentration to 0.025 and 0.1 M resulted in drastically

improved corrosion inhibition. In the former case, no localized attack was

observed during scratching. At 0.1 M concentration, partial dissolution of the

matrix was observed within and outside the scratched area. Furthermore, ex situ

SEM-EDS analysis revealed the presence of Si over the S-phase particles, which

could be attributed to the formation of magnesium-silicate species.

136

FIGURES AND TABLES

R2 = 0.99

12 SolutionpH

10 0 20 40 60 80 100

[SiO -2] mM 3

Figure 4.1. Effect of silicate concentration on solution pH.

137

Figure 4.2. Chemical equilibrium diagram for (a) 1 mM and (b) 100 mM Si(OH)4 in 0.1 M NaCl solution. Speciation diagram generated using MedusaTM software. The dashed black line signifies maximum solubility and ‘cr’ in parenthesis denotes crystalline species.

138

800 25 mM 100 mM

400 E 50 mM pit

0 10 mM

-400 E corr 1 mM

-800

No inhibitor (pH 6) Potential (mV vs. SCE) vs. (mV Potential No inhibitor (pH 11) -1200

-1600 10-10 10-9 10-8 10-7 10-6 10-5 0.0001 0.001

Current Density (A/cm2)

Figure 4.3. Naturally aerated polarization curves for AA2024-T3 in 0.1 M NaCl solution at varying Na2SiO3 concentrations.

139

1000 Epit 800 Ecorr

600

400

200

-200

-400 Potential (mV vs. SCE) vs. (mV Potential -600

-800

-1000 0 20 40 60 80 100

Concentration (mM)

Figure 4.4. Corrosion potential Ecorr and pitting potential Epit with incremental amounts of Na2SiO3.

140

105 -100

-80 104 1 day

) Increasing time 2 -60 8 day Increasing time 1000 1 day

-40 IZI (ohm.cm IZI

Theta(degrees) 8 day 100 -20

10 0 4 5 0.01 0.1 1 10 100 1000 104 105 0.01 0.1 1 10 100 1000 10 10 Frequency (Hz) Frequency (Hz)

105 -80

4 -60

) 2

-40 1000

8 day Increasing time IZI (ohm.cm IZI Theta(degrees) Increasing time -20 100 8 day

1 day 0 1 day

10 4 5 0.01 0.1 1 10 100 1000 104 105 0.01 0.1 1 10 100 1000 10 10 Frequency (Hz) Frequency (Hz)

Figure 4.5. Bode magnitude and phase angle plots of 2023-T3 coupons immersed in (a) 0.1 M NaCl (pH 6 – natural pH) (b) 0.1 M NaCl (pH 12.8 – pH adjusted) and (c) 0.1 M NaCl + 25 mM Na2SiO3 (pH 12.8 – natural pH).

141

Figure 4.5: Continued

107 -100 Increasing time

6 8 day 10 -80 8 day

105 ) 2 -60 5 hour Increasing time 104

-40 5 hour IZI (ohm.cm IZI

1000 Theta(degrees)

-20 100

10 0 4 5 0.01 0.1 1 10 100 1000 104 105 0.01 0.1 1 10 100 1000 10 10 Frequency (Hz) Frequency (Hz)

142

7 10 0.1 M NaCl + 25 mM Na SiO 2 3

106

) 2

105

(Ohm.cm

p R

104 0.1 M NaCl-only

1000 0 1 2 3 4 5 6 7 8

Time (day)

0.001

0.1 M NaCl-only

0.0001

)

(F/cm

dl C 10-5 0.1 M NaCl + 25 mM Na SiO 2 3

10-6 0 1 2 3 4 5 6 7 8

Time (day)

Figure 4.6. Impedance data summary of (a) total resistance Rt and (b) film capacitance Cf with time.

143

Figure 4.7. Images of samples after 1 day exposure in 0.1 M NaCl solution with and without 25 mM Na2SiO3 (a) open-circuit potential macrographs and (b) optical micrographs.

144

Figure 4.8. (a) Secondary electron micrographs before and after 3 day exposure in 0.1 M NaCl solution with 25 mM Na2SiO3, and EDS spectrum of S-phase particle after exposure (b) EDS mapping of selected area after exposure.

145

Table 4.1. Chemical quantification analysis obtained from high-resolution XPS spectra after 1 day exposure in 0.1 M NaCl solution with and without 25 m Na2SiO3.

Total Atomic % 0.1 M Mg Cu O Si Al Na Cl C NaCl- Inhibitor- - 0.49 63.34 - 25.18 - - 11.00 free + 25 mM 0.59 0.16 48.02 10.84 13.36 4.18 - 22.34 Na2SiO3

146

a b

c d

Figure 4.9. Figure 4.9. Spectra showing the (a) Si 2p (b) Al 2p (c) Na 1s and (d) Mg 1s peaks after 1 day exposure in A – 0.1 M NaCl and B – 0.1 M NaCl + 25 mM Na2SiO3 solution.

147

60 O

30 Al(0) Atomic% Si Al(3+) 20

10 C

Na 0 0 50 100 150 200

Etch time (sec)

Figure 4.10. Sputter depth profile after 1 day immersion in 0.1 M NaCl + 25 mM Na2SiO3.

148

-400

-600

-800 No inhibitor 1 mM 10 mM -1000 25 mM

-1200

Potential (mV vs. SCE) vs. (mV Potential pH 7.3

-1400

-1600 10-10 10-9 10-8 10-7 10-6 10-5 0.0001 0.001

Current Density (A/cm2)

-400

-600 No inhibitor 1 mM -800 10 mM pH 4 25 mM

-1000

-1200 Potential (mV vs. SCE) vs. (mV Potential

-1400

-1600 10-9 10-8 10-7 10-6 10-5 0.0001 0.001 0.01 0.1

Current Density (A/cm2)

Figure 4.11. Naturally aerated polarization curves in 0.1 M NaCl at varying Na2SiO3 concentrations in pH adjusted (a) 7.3 and (b) 4 solutions.

149

a b

Figure 4.12. (a) Open-circuit potential macrograph of coupon exposed to 0.1 M NaCl + 25 mM Na2SiO3 solution at pH 7.3 (b) Secondary electron image of surface (c) Secondary electron image of partially covered large intermetallic particle with accompanying EDS spectrum of particulates. Evidence of small amount of Si detected over particulate covered area.

150

Figure 4.13. (a) Open-circuit potential macrograph of coupon exposed to 0.1 M NaCl + 25 mM Na2SiO3 solution at pH 4 (b) Secondary electron image of site A (refer to macrograph) (c) Backscatter electron image of site A with accompanying EDS spectrum of labeled S-phase particle (d) Secondary electron image of site B with accompanying EDS spectrum of film.

Continued 151

Figure 4.12: Continued

152

Table 4.2. Chemical quantification analysis obtained from high-resolution XPS spectra after 3 day exposure in pH adjusted 0.1 M NaCl solution with 25 m Na2SiO3.

Total Atomic % pH Mg Cu O Si Al Na Cl C adjusted

7.3 0.20 0.13 48.42 8.46 22.02 0.16 - 20.61

4.0 0.46 1.84 53.57 8.04 17.33 0.26 - 18.52

153

Figure 4.14. Spectra showing the (a) Si 2p and (b) Al 2p peaks after 3 day exposure in pH 7.30 and pH 4 solution containing 0.1 M NaCl + 25 mM Na2SiO3.

154

Figure 4.15. Stability diagram of a ternary alloy system consisting of Al, Cu, and Mg in silicate solution (generated using OLI AnalyzerTM).

155

-200

E pit -400

-600 E No inhibitor corr 0.1 mM A + 0.1 mM B 1 mM A + 1 mM B -800 5 mM A + 1 mM B 1 mM A + 5 mM B

Potential (mV vs. SCE) vs. (mV Potential 5 mM A + 5 mM B

-1000

-1200 10-10 10-8 10-6 0.0001 0.01

Current Density (A/cm2)

Figure 4.16. Naturally aerated polarization curves in 0.1 M NaCl at varying Na2MoO4 (compound A) and Na2SiO3 (compound B) concentrations mixed together in solution.

156

Figure 4.17. (a) Open-circuit potential macrograph of coupon exposed to 0.1 M NaCl -3 -3 solution with 10 M Na2SiO3 + 10 M Na2SiO3 for 10 days (b) Secondary electron image of surface with elementals maps of scanned area (c) Secondary electron image with EDS spectrum of large intermetallic particle covered with Si particulates. 157

Table 4.3. Chemical quantification analysis obtained from high-resolution XPS spectra after 10 days of exposure in 0.1 M NaCl solution with 1 m Na2SiO3 + 1 mM Na2MoO4.

Total Atomic % Mg Cu O Si Mo Al Na Cl C - 0.28 38.29 13.57 - 7.96 4.56 - 35.33

158

a b

c d

Figure 4.18. Spectra showing the (a) Si 2p (b) Mo 3d (c) Al 2p and (d) O 1s peaks after 10 days of exposure in 0.1 M NaCl solution with 1 mM Na2SiO3 + 1 mM Na2MoO4.

159

-400

0.1 M NaCl-only

-600

1:1 (SiO -2 : MoO -2) 3 4 -800 0.002 M total con.

-1000 Potential (mV vs. SCE) vs. (mV Potential -1200

-1400 0 10 20 30 40 50 60

Time (min)

Figure 4.19. OCP measurements in 0.1 M NaCl solution with and without 10-3 M -3 Na2SiO3 + 10 M Na2MoO4.

160

-5 Figure 4.20. In situ AFM scratching in 0.1 M NaCl + 10 M Na2SiO3. Scan size = 70 µm. Top left image is Volta Potential map before exposure with z range = 300 mV. The other images are topographic maps with z range = 300 nm. The first is the map prior to exposure and the rest are in situ images during rastering at indicated times and setpoint voltages.

161

Figure 4.21. SKPFM and SEM analysis of the selected area before in situ AFM scratching in 0.1 M NaCl + 0.025 M Na2SiO3.

162

Figure 4.22. In situ AFM scratching in 0.1 M NaCl with 0.025 M Na2SiO3. Scan size=30µm. Z range=300nm.

163

Figure 4.23. In situ AFM scratching in 0.1 M NaCl + 0.1 M Na2SiO3. Scan size = 75 µm. Top left image is Volta Potential map before exposure with z range = 500 mV. The other images are topographic maps with z range = 300 nm. The first is the map prior to exposure and the rest are in situ images during rastering at indicated times and setpoint voltages.

164

Figure 4.24. Ex situ SEM-EDS analysis after AFM scratching in 0.1 M NaCl with 0.1 M Na2SiO3.

165

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169

CHAPTER 5: CORROSION INHIBITON OF AA2024-T3 BY PRASEODYMIUM CHLORIDE

5.1 Introduction

Aluminum alloys are used extensively in aerospace applications due to their excellent mechanical properties 1. However, the heterogeneous microstructure of aerospace aluminum alloys renders them susceptible to localized corrosion 2-4.

Traditionally, chromate-based pigments and conversion coatings have been used to prevent the corrosion of high-strength aluminum alloys 5, 6. However, the carcinogenic effect of soluble hexavalent chromate has forced governments to impart regulations that restrict their use in a wide-range of applications. Therefore, efforts are underway to find environmentally-friendly alternatives that offer the same level of protection and reliability as chromate protection schemes 6. In this regard, rare-earth (RE) metal compounds have been investigated and have shown to be promising candidates as chromate replacements.

The early work by Hinton et al. 7-9 demonstrated that the corrosion rate of aluminum alloy 7075-T6 (Al-Zn) decreased by an order of magnitude when 100 ppm of

RE salts were added to 0.1 M NaCl. The oxygen reduction kinetics were reduced and the corrosion potential shifted to lower potentials 7. Cerium and praseodymium solutions provided better inhibition than other RE metals, exhibiting larger shifts in the corrosion potential and lower pit densities. 170

The mechanism of corrosion inhibition by RE metal compounds involves the formation of an insoluble film over the intermetallic particles in response to an increase in the local pH at the metal/electrolyte interface 10-14. The solubility of RE metal compounds decreases in alkaline environments resulting in the formation of an insoluble oxide/hydroxide that adheres strongly to the surface. This film provides a barrier to the diffusion of oxygen and the transport of electrons to the oxygen reduction reaction.

Hence, the overall corrosion rate is decreased. A similar proposal was made by Yasakau et al. concerning the deposition of cerium and lanthanum over S-phase in AA2024 15. It was determined that the precipitation of a hydroxide over the S-phase particles inhibits the anodic and cathodic processes, and decreases the redeposition of copper on the surface.

Despite the proposed mechanism of inhibition by RE metal compounds, another explanation for the inhibition by cerium has been described by Aldekiewicz et al. 16, 17.

They proposed that, under sufficiently alkaline conditions, oxygen will oxidize Ce3+ to

Ce4+ 16. Due to the complexities of the oxygen reduction reaction, a two-electron reduction pathway may proceed in which oxygen is reduced to produce hydrogen peroxide at the vicinity of the intermetallic particles. Oxidation of Ce3+ will occur in the presence of hydrogen peroxide resulting in the precipitation of insoluble CeO2. In the case where oxygen reduction occurs by a four-electron reduction pathway, oxidation of cerium is likely to occur away from the intermetallic particles, leading to inefficient film deposition.

171

Rare-earth metal compounds have been introduced as replacements for chromates using various strategies. Hughes et al. 18 developed a cerium-rich conversion coating for

-1 AA2024-T3 by immersing the alloy into solution containing 10 g L CeCl3 and 1 %

H2O2. The thickness of the coating varied across the surface with heavy precipitation over the intermetallic particles. This method is different from another approach that is based on oxidation of the alloy surface. For instance, Mansfeld et al. 19 developed a two- step coating process that involves immersion into boiling Ce(NO3)3 and CeCl3 solution, followed by anodic polarization in 0.1 M Na2MoO4. The resulting film was determined to be very stable and effectively hindered the oxygen reduction kinetics.

Despite the improved corrosion protection afforded by RE metal compounds, the inherent effect of praseodymium on aluminum alloys has not been studied as deeply as the other RE elements. Furthermore, considering that praseodymium oxides/hydroxides are currently employed as pigments in some currently-available commercial non- chromate coatings, it is of interest to study the functionality of praseodymium on the localized corrosion processes of AA2024-T3. In addition, most conversion coatings and pigments act by releasing soluble species into the local aqueous environment. Thus, the purpose of this study is to understand the mechanism of corrosion inhibition of AA2024-

T3 by PrCl3 dissolved in 0.1 M NaCl solution. The functionality of praseodymium in various environmental conditions (oxygen-free, CO2, pH, etc.) is also addressed.

5.2 Experimental

5.2.1 Materials and sample preparation: Reagent-grade praseodymium chloride (PrCl3) and sodium chloride (NaCl) were used for all experiments. Solutions were prepared 172 using 18.2 MΩ·cm deionized water. Solution pH was adjusted with diluted solutions of concentrated H2SO4 and NaOH.

Samples of solution heat treated, naturally aged AA2024-T3 (nominal composition 3.9-4.9% Cu, 1.2-1.8% Mg, 0.3-0.9% Mn, 0.5% Fe, 0.5% Si) 20 were mechanically abraded with SiC paper to 1200 grit in a nonaqueous slurry (Blue Lube from Struers) to minimize the onset of corrosion. Samples analyzed microscopically were polished to 1 μm diamond paste. All samples were cleaned with ethyl alcohol in an ultrasonic bath, air dried, and stored overnight in a desiccator. For all experiments, sample dimensions of 2.5 x 2.5 x 0.3 cm were employed.

5.2.2 Potentiodynamic polarization curves and electrochemical impedance spectroscopy:

A platinum-mesh counter electrode and a saturated reference electrode (SCE) were used for all electrochemical experiments. Cathodic and anodic polarization experiments were performed in naturally aerated solutions and were acquired separately under each condition. Experiments were repeated at least in triplicate and the curves shown below are typical curves for each condition. The potential sweep was conducted starting at the open-circuit potential (OCP) using a scan rate of 10 mV/min.

Impedance measurements were performed with and without PrCl3 additions to 0.1

M NaCl solution. Experiments were conducted in aerated, deaerated, and decarbonated electrolyte. Deaeration was achieved by 2 h Ar degassing of both the solution reservoir and the electrochemical cell. After 2 h, solution was forced into the cell by pressurizing the reservoir. Decarbonation of solutions was performed with Air Ultra ZeroTM (air free of CO2 and containing less than 0.1 ppm of hydrocarbons) degassing following a similar

173 procedure. A 10 mV sinusoidal voltage was applied at OCP with the frequency ranging from 105 to 10-2 Hz. The impedance response was fitted using a simplified Randles circuit to extract total resistance and film capacitance. Further details are described below. In addition, all electrochemical experiments were performed with a GamryTM Ref

600 potentiostat using an exposure area of approximately 1 cm2.

5.2.3 Free-corrosion experiments coupled with secondary electron microscopy:

AA2024-T3 coupons were exposed in aerated, deaerated, and decarbonated NaCl solution with and without PrCl3. Samples immersed in solution were rested against the beaker wall, with the unpolished surface facing the wall and covered with electroplating tape (3M). At the end of the exposure, all samples were rinsed using 18.2 MΩ·cm deionized water, air-dried, and stored overnight in a desiccator. They were visually inspected and analyzed using either a Quanta 200 or Sirion scanning electron microscope

(SEM) integrated with energy dispersive X-ray spectroscopy (EDS). In addition, a long term OCP exposed sample was analyzed by X-ray diffraction using a Scintag XDS 2000 diffractometer with Cu Kα radiation (λ = 1.506 A) and at a scan rate of 0.5 deg./min.

5.2.4 Atomic force microscopy: In situ AFM scratching and scanning Kelvin probe force microscopy (SKPFM) experiments were conducted with a Multimode AFM coupled with a Nanoscope V controller (Bruker Corporation). SKPFM allows for the simultaneous measurement of the surface topography and Volta-potential distribution on a line-by-line basis using metal-coated silicon tips that are electrically conducting. The cantilevers were obtained from Bruker Corporation, and were coated with a PtIr thin-film on both sides. It has been determined that the Volta-potential measured by this technique is a

174 good measure of practical nobility 21; a strong correlation was observed between the measured potential and the corrosion potential measured in solution. Details of this technique can be found elsewhere 21. Furthermore, characterization of the surface was carried out before and after each experiment using SEM/EDS.

In situ AFM scratching experiments were performed in naturally aerated electrolyte that was pumped into a glass cell with a volume of 0.1 mL at a rate of 10 mL/h. Scratching was conducted in contact mode with commercially available Si cantilever tips at a constant scan rate of 2.5 Hz. The force applied by the tip on the sample surface is proportional to the deflection of the cantilever, which is kept constant and controlled by the set-point voltage. The set-point voltage was adjusted from 0.2 to

10 V which correlates to tip-sample pressures ranging from 100 to 525 nN, respectively.

Therefore, increasing the set-point voltage is equivalent to increasing the scratching force on the sample. Tips were replaced after each experiment due to wear during scratching.

Prior to each experiment, SKPFM was conducted in air and the surface was characterized using SEM-EDS. Characterization of the surface after each experiment was also conducted using SEM-EDS.

5.2.5 X-ray photoelectron spectroscopy: XPS measurements were performed to investigate the presence of praseodymium at the surface of AA2024-T3. Measurements were conducted on samples that were immersed under free-corrosion conditions. XPS spectra were acquired using a Kratos AXIS Ultra spectrometer with a monochromated Al x-ray source operated at 130 W and calibrated to the adventitious carbon 1s peak at 284.6 eV. Chemical state assessment was achieved by curve-fitting the spectra using

175

CasaXPSTM software. All reported binding energies are rounded to the nearest tenth of an eV and are ascertained a precision value of +/- 0.2 eV.

5.3 Results and Discussion

5.3.1 Aqueous solution chemistry: Previous studies by Treu et al. showed activation of

Pr2O3 and Pr6O11 pigments embedded within an organic coating after salt spray exposure

22. Dissolution and re-precipitation of these oxides resulted in the formation of praseodymium- hydroxide and/or hydroxycarbonate species on the coating surface and in areas where the underlying substrate was exposed. Therefore, it was of interest to assess the praseodymium species likely to form in dilute sodium chloride solution in the absence and presence of atmospheric CO2. This was accomplished using chemical equilibrium diagrams computed by MedusaTM. Embedded within the software is a database that contains chemical equilibrium constants at 25°C.

-3 Figure 5.1 shows the generated speciation diagram for 10 M PrCl3 in 0.1 M

NaCl solution in the absence and presence of atmospheric CO2 (partial pressure equal to

10-3.5 atm). The natural pH was measured to be around pH 5.7 and pH 5.5, respectively.

3+ In the absence of atmospheric CO2, the solubility of Pr decreases significantly at pH greater than 7.8 owing to the formation of insoluble ( ) species. The overall solubility remains low until pH 10, where further hydrolysis results in the formation of

3+ ( ) anions. In the presence of CO2, the critical pH for Pr solubility is shifted to

6.1. Above this pH, insoluble ( ) is formed. The solubility begins to increase

gradually at pH 8.9 owing to the formation of ( ) . Thus, it can be concluded that

176 decreased solubility of praseodymium species when carbonate species are present in solution could enhance the deposition of precipitates on the surface.

The effect of Pr3+ concentration on the solution pH was also calculated by

MedusaTM. As illustrated in Figure 5.2, the concentration is inversely proportional to the pH; the addition of 1 mM Pr3+ decreased the pH from about pH 5.6 to pH 4.1. A similar value was measured experimentally at the same concentration. The decrease in pH can be explained by the hydrolysis of praseodymium ions in solution as described by the following chemical equation:

( ) Eq. 5.1

Hence, increasing the concentration of Pr3+ leads to further hydrolysis and a decrease in the pH. As discussed below, the effect of concentration on the solution pH has a significant implication on the inhibition performance by PrCl3.

5.3.2 Polarization in aerated sodium chloride solution: The corrosion inhibition effect of praseodymium on AA2024-T3 was determined using potentiodynamic polarization curves. Figure 5.3 shows naturally aerated polarization curves in 0.1 M NaCl solution with varying PrCl3 concentrations. The results show that praseodymium imparts cathodic inhibition by decreasing the limiting current density associated with the oxygen reduction reaction. The largest effect was observed at an optimum concentration of 0.2 mM where the oxygen reduction kinetics decreased by 20x. However, this decrease is not enough to be accompanied by a decrease in the open-circuit potential (OCP), which is pinned at the breakdown potential. At the highest concentration tested (4.0 mM) the decrease in oxygen reduction kinetics is accompanied by a 30 mV downward shift in the OCP. It is

177 interesting to note that the observed decrease in OCP at this concentration is not associated with cathodic inhibition, but is attributed to an increase in the anodic kinetics.

The decrease of inhibition performance can be associated with an increase in pH at higher concentrations of PrCl3 in solution (Figure 5.1). Thus, it can be inferred that the optimum concentration for corrosion inhibition of AA2024-T3 in 0.1 M NaCl solution is 0.2 mM

PrCl3.

5.3.3 Morphology of AA2024-T3 after free-corrosion exposure: Since the largest corrosion inhibition effect by praseodymium was observed at 0.2 mM concentration, it was of interest to study the morphology of AA2024-T3 after exposure in 0.1 M NaCl +

0.2 mM PrCl3 solution under free-corrosion conditions. Figure 5.3a shows macrographs after 1 day of exposure in naturally aerated solution with and without PrCl3. In the absence of inhibitor, there are clear signs of corrosion on the sample surface. Optical microscopy reveals evidence of attack over the intermetallic particles and matrix (Figure

5.4b). With the addition of praseodymium, white deposits on the surface are observed in the low magnification image. At higher magnifications, the extent of attack is significantly lowered, Figure 5.4b.

Figure 5.5 shows SEM images of an area containing S-phase particles before and after 2 days of exposure in 0.1 M NaCl solution with 0.2 mM PrCl3. Comparison of the before and after secondary electron images reveals the precipitation of a thick-oxide over the S-phase particles that is rich in Pr and O as detected by EDS (not shown).

Furthermore, Mg detected at the particles suggests that praseodymium partially suppresses the dissolution of S-phase particles. SEM-EDS analysis of a region

178 containing Fe-intermetallic particles (not shown) did not reveal signs of deposition or corrosion activity.

Figure 5.5a shows Volta-potential distribution maps of a scanned area containing a cluster of S-phase particles (identified by SEM/EDS) before and after 5 hours of exposure to PrCl3 solution. The before image shows a well-defined frontier among the intermetallic particles within the cluster. However, a high potential contrast encompassing an area wider than the cluster region is observed after exposure.

Comparison of the before and after topography images in Figure 5.5b shows increased contrast over the S-phase intermetallic particles associated with deposits on top of the particles. Subsequent SEM-EDS analysis (not shown) revealed that the deposits are rich in Pr and O, which is in agreement with the results obtained after 2 days of exposure in

PrCl3 (Figure 5.5). It should be noted that the S-phase particles in as-polished samples are not distinguished in the topography maps. Although these particles are harder than the

Al matrix, the loss of material by corrosion during polishing balances the slower rate of

12 removal by polishing . Hence, they are initially almost coplanar with the matrix and are not observable in the topography image prior to exposure. Furthermore, according to

Schmutz and Frankel 23, the high potential contrast associated with the S-phase prior to exposure is attributed to an altered surface film that is different in composition and microstructure to the bulk material and forms over these particles during polishing.

The broadening of the potential signal over the cluster region after exposure may indicate partial dissolution of the S-phase particles. According to the literature, S-phase particles undergo preferential dissolution of Mg and Al during exposure to aggressive

179 solution 24. The de-alloying of the S-phase results in the formation of Cu-rich remnants that are nobler than the surrounding matrix, thereby transforming these particles from local anodes to local cathodes. Furthermore, the de-alloyed Cu remnants release metallic

Cu into the solution that can oxidize and subsequently reduce on the surrounding matrix

24. This phenomenon is made evident by the formation of rings around the intermetallic particles that manifest from Cu redeposition (observed in Figure 5.14b) 15. In addition,

SKPFM studies conducted on AA2024-T3 after exposure to CeCl3 solution showed a broadening of the Volta potential peak over the S-phase particles 15. It was proposed that the broadening of the potential peak is associated with replated Cu on the surrounding matrix as the S-phase undergo selective dissolution prior to cerium deposition 15.

Following a similar reasoning, it can be concluded that the high potential contrast encompassing an area wider than the cluster region in Figure 5.5a is attributed to replated

Cu on the surrounding matrix as the S-phase undergo selective dissolution. However, there is no indication in the topographic map of a film deposited over the alloy matrix.

Perhaps due to the detection limits of EDS, no praseodymium was detected over the larger Fe-intermetallic particles after SEM-EDS analysis. However, precipitation over these particles was made evident by SKPFM. Figure 5.6 shows the topography and

Volta-potential distribution maps of Fe-containing intermetallic particles before and after exposure to PrCl3 solution. The before and after topography images suggest non-uniform deposition of precipitates over the particles after exposure to solution. This was supported by the respective Volta-potential map shown in Figure 5.6a which shows variable contrast over the Fe-particles sites. Furthermore, there is evidence of deposition

180 over the Al matrix and at the smaller intermetallic particles associated with S-phase, which is in agreement with the above results.

The electrochemical properties of Fe-containing intermetallic particles have been discussed extensively in literature. These particles are intrinsically nobler than the matrix and act as local cathodes on the surface 11, 13, 25. They support oxygen reduction and promote dissolution of the adjacent matrix 10-14. However, previous studies have revealed that some of the Fe-intermetallic particles exhibit lateral gradients in Volta-potential 21.

These gradients can be attributed to compositional heterogeneities within these particles.

Hence, electrochemical variances over the particles during exposure may result in pH gradients and the non-uniform deposition of corrosion product in the presence PrCl3.

Furthermore, the results suggest that the relatively low amount of deposit observed on these intermetallic particles is enough to decrease the current supplied to the anodic reaction. Hence, the rate of Al dissolution is reduced. Other authors have observed similar effects involving cerium ions in solution. Aldykewicz et al. attributed an absence of Ce detected by EDS to the presence of a more protective oxide over the Fe-containing intermetallic particles that is effective in blocking the cathodic reaction 17.

5.3.4 X-ray photoelectron spectroscopy: Praseodymium 3d- and 4p- photoemission spectra were acquired by XPS to investigate the oxidation state of praseodymium on the surface after 2 days of exposure in naturally aerated solution. The 4p spectra were acquired to help deconvolute the Pr 3d spectra which overlap the energy region of the Cu

2p lines. Furthermore, referencing of the praseodymium peaks with the literature was accomplished with respect to the Pr 3d lines, which are shown in Figure 5.8a. In the

181

3+ presence of 0.2 mM PrCl3, a main peak in the 3d5/2 spectrum associated with Pr is observed at a binding energy of 933.7 eV 26. In addition, a characteristic satellite peak consistent with Pr3+ is also observed at a binding energy of 929.8 eV. According to the literature, rare-earth metal oxides exhibit satellite features at lower binding energy

27, 28 because of solid-state hybridization effects . In the absence of PrCl3, two peaks corresponding to metallic and oxidized copper are observed in the same spectral region, at a binding energy of 933.0 eV and 935.7 eV, respectively. It should be noted that copper was not detected in the presence of praseodymium. This is likely because both the Cu-enriched intermetallic particles and replated Cu are covered by praseodymium- rich deposits. Figure 5.8b shows the Pr 4p3/2 spectra for the same samples. In the presence of praseodymium, two peaks are observed at a binding energy of 216.6 eV and

219.6 eV. The former peak is consistent with shake-down satellites observed by rare- earth metal compounds 26. Quantification of the peak area ratio of Pr- 4p to 3d after correcting for the sensitivity factors was calculated to be 1.07 (within 1 at% accuracy limit of error), which provides support for in the assignment of the peaks in the Pr 3d spectra. No peak is observed in the Pr 4p energy range for the sample exposed to the inhibitor-free solution.

Figure 5.9 shows the O 1s and C 1s photoemission spectra after exposure to

PrCl3-containing solution. The O 1s spectrum (Figure 5.9a) reveals a large peak at a binding energy of 531.6 eV, which encompasses hydroxide and/or carbonate species, but not metal oxides 29. The Al 2p spectrum (not shown) reveals a strong Al3+ peak at 74.1

30 eV which can be attributed to Al(OH)3 . Thus, a large portion of the O 1s peak can

182 assigned to the aluminum hydroxide that is present on the surface 29. Interestingly, the O

1s peak for lanthanum-carbonate species was found to be at a binding energy of 531.6 eV

31. Attempts to find references of praseodymium carbonate/hydroxycarbonate species in the literature were not successful. However, the corresponding C 1s spectrum shown in

Figure 5.9b indicated the presence of carbonate species at a binding energy of 289.3 eV

29 , which can encompass either praseodymium-carbonate ( ( ) ) and/or praseodymium-hydroxycarbonate ( ). In addition, two peaks pertaining to adventitious carbon (from atmospheric contamination) are present at binding energies of

284.6 and 285.6 eV. It should be noted that all samples exposed to air contain carbonate contamination from the atmosphere 31. However, an excess of carbonate species was detected on the sample exposed to PrCl3 compared to the negative control (discussed in more detail below).

5.3.5 Long term free-corrosion exposure: Experiments were also conducted for longer immersion times. Figure 5.10a shows a macrograph of a sample after 30 days of exposure in 0.1 M NaCl + 0.2 mM PrCl3 under free-corrosion conditions. Visual inspection of the surface shows clear signs of film formation, which was confirmed by subsequent SEM-EDS analysis. The SEM images shown in Figure 5.10b reveal mud- cracking and the presence of nodules 5 µm in diameter. Characterization of the film by

XRD (Figure 5.10c) indicated the presence of species on the surface, which supports the excess carbonate species found with XPS. These results suggest that the alloy becomes completely covered with time as old cathodic sites become blocked by precipitates and new sites become activated. It should be noted that the complete

183 coverage of the alloy surface has also been observed by other RE metal inhibitors for high-strength aluminum alloys 32.

5.3.6 Evaluation of inhibition in decarbonated and deaerated electrolyte: Experiments were performed in deaerated and decarbonated electrolyte to assess the effects of oxygen and CO2 on the inhibition mechanism by praseodymium. Figure 5.11a shows macrographs after 2 days of exposure in decarbonated and deaerated 0.1 M NaCl solution with 0.2 mM PrCl3. In the absence of CO2, there is evidence of a film on the surface.

Optical microscopy reveals evidence of film deposition and areas of mud-cracking over the aluminum matrix near several intermetallic particles. This is consistent with the island-like growth typically exhibited by other rare-earth metal oxides 7, 32. In deaerated solution, the sample surface turns light brown and there is evidence of deposition, albeit at smaller amounts, over the intermetallic particles according to the corresponding optical micrograph in Figure 5.11b.

Surface characterization was also performed using XPS. Table 5.1 shows the results for the chemical quantification analysis after 2 days of exposure in aerated, deaerated, and decarbonated solution. The concentrations shown in parenthesis correspond to carbonated species. It should be noted that a relatively high concentration of carbonate was present on the sample exposed to decarbonated solution. This could be attributed to traces of carbonate during exposure or carbonate contamination (from the atmosphere) when transferring the sample to the XPS chamber 31. Furthermore, a significant amount of praseodymium was found to be present on the surface, which can be attributed to the deposition observed over the intermetallic particles (Figure 5.10b).

184

In naturally aerated PrCl3 solution (with CO2), no concentration of Mg or Cu was detected on the sample surface. The presence of excess carbonate species in comparison to the control can be attributed to praseodymium carbonate or hydroxycarbonate as discussed above. In addition, the presence of a thick film on the surface is supported by the relatively low concentration of Al3+ (determined from the Al 2p spectra – not shown) detected on the surface. A thick film over the intermetallic particles or surface can cover the underlying Al-oxide/hydroxide, which would be reflected by a weaker Al3+ signal. A similar effect was observed in the case of decarbonated solution, where a weaker Al3+ signal was also observed in comparison to the negative control.

In deaerated solution, a significant amount of oxidized Mg and Cu (spectra not shown) was detected on the sample surface. This can be attributed to increased corrosion activity, which is consistent with the low impedance magnitude (discussed below) and the lower concentration of deposits observed over the intermetallic particles (Figure

5.10b). In addition, a relatively small concentration of praseodymium was detected on the surface by XPS. These observations suggest that the deposition of praseodymium species is hindered in oxygen-free solution by a decrease in the rate of the cathodic reaction.

Figure 5.12a shows the Pr 3d5/2 and O 1s photoelecton spectra after 2 days of immersion in decarbonated 0.1 M NaCl solution with 0.2 mM PrCl3. Two peaks

3+ associated with Pr are present in the Pr 3d5/2 spectrum at a binding energy of 933.2 and

929.2 eV 26. The accompanying O 1s spectrum reveals a large peak at a binding energy of 531.5 eV, which envelopes hydroxide and/or carbonate species but not metal oxides 29.

185

However, considering that the solution was decarbonated during exposure, it is likely that the peak can be ascribed to the combined signals of aluminum-hydroxide ( ( ) ) and praseodymium-hydroxide ( ( ) ) species. This is also supported by the C 1s spectra

(not shown).

In deaerated solution (with CO2), a number of differences are observed. The Pr

3+ 3d5/2 spectrum shown in Figure 5.12b reveals two peaks associated with Pr at binding energies of 933.8 and 929.6 eV 26. However, two additional peaks are observed at 936.0 and 931.8 eV, which can be ascribed to two different species of oxidized copper; Cu2+ and Cu+, respectively 29. The accompanying O 1s spectrum exhibits a peak at a binding energy of 531.6 eV. As before, this peak envelopes hydroxide and/or carbonate species but not metal oxides 29. While the Al 2p spectrum (not shown) confirmed the presence of

( ) on the surface, it is also possible that praseodymium carbonate species exists on the surface like in naturally aerated solution (with CO2), albeit at smaller concentrations.

The presence of a small amount of praseodymium along with a relatively large concentration of carbonate species on the surface is supported by the chemical quantification analysis shown in Table 5.1.

Figure 5.13 shows the Pr 4p3/2 spectra under similar conditions. In both decarbonated and deaerated solutions, two peaks are present at binding energies of 219.6 and 216.3 eV. As described above, the latter peak is associated with shake-down satellites which are commonly present in RE metal compounds 26. Furthermore, the corrected peak area ratio of Pr- 4p to 3d was found to be close to 1 for both solution conditions, which further corroborates the aforementioned conclusions.

186

Impedance spectroscopy was conducted to determine the inhibition performance under aerated (with atmospheric CO2) and decarbonated electrolytic conditions. The

Bode magnitude and phase angle plots are shown in Figure 5.14. To extract the total resistance and film capacitance, the spectra were fitted using a simplified Randles circuit that included a constant phase element (CPE) to represent the film capacitance. Note that data points associated with a second time constant at lower frequencies are not considered in these fits. The true capacitance was derived from the CPE parameters using the following equation 33:

Eq. 5.2

In the above equation, Y and α represent the CPE magnitude and exponent, respectively,

Rt is the extracted total resistance, and Cf is the calculated film capacitance. After 2 days of exposure in aerated electrolyte, the low frequency impedance was an order of magnitude higher in 0.1 M NaCl + 0.2 mM PrCl3 solution than in inhibitor-free solution.

The total resistance was measured to be 110 kΩ·cm2 and 4 kΩ·cm2 in solution with and without PrCl3, respectively. In addition, the phase angle is larger over a larger frequency domain in the presence of praseodymium indicating improved corrosion inhibition. The

-2 film capacitances were measured to be 38 µF·cm in the case of PrCl3 solution, and 120

µF·cm-2 for the negative control. It should be noted that the duplication of these results was impeded by contamination related issues with the first batch of experiments.

Interestingly, in the absence of CO2, the inhibition performance is diminished. The low frequency response decreases by a factor of 3 in comparison to aerated solution containing PrCl3. In the absence of CO2, the total resistance and film capacitance was 187

2 -2 measured to be 25 kΩ·cm and 60 µF·cm , respectively. These results suggest that CO2 plays an important role in the inhibition mechanism of AA2024 by PrCl3.

Interestingly, the inhibition performance of praseodymium is significantly diminished in oxygen-free solution. Figure 5.15 shows the Bode magnitude and phase angle plots after 2 days of exposure in deaerated 0.1 M NaCl solution with and without

PrCl3. In the presence of praseodymium, the low frequency impedance is decreased by a factor of 3 in comparison to the negative control. The measured total resistance was 98.0

2 2 kΩ·cm and 11.3 kΩ·cm in the absence and presence of PrCl3, respectively. In addition, the film capacitance was measured to be 19.1 µF·cm-2 in inhibitor-free solution and 15.2

µF·cm-2 in the presence of praseodymium. Thus, it is evident that the inhibition mechanism of AA2024 by PrCl3 is not only CO2-dependent but also requires the presence of an oxidant in the environment, such as dissolved oxygen. Moreover, in the absence of oxygen, the addition of praseodymium actually decreases the corrosion resistance of the alloy, which could be attributed to a decrease in pH with the addition of PrCl3. On the other hand, in the presence of oxygen (and CO2) there is a significant improvement in the corrosion protection of the alloy reflected by a lowering of the oxygen reduction kinetics

(Figure 5.3). The total resistance increased by 26x and the film capacitance decreased by a factor of 3 with respect to the negative control. It is likely that in the presence of oxygen, the interfacial metal/electrolyte pH is increased high enough to favor film deposition. However, in oxygen-free environments, film deposition is not favorable and a decrease in the total resistance can be attributed to a decrease in pH. Hence, it is clear that praseodymium requires an environment that is both aerated and carbonated to impart

188 corrosion inhibition, where oxygen reduction dominates the cathodic kinetics and the carbonate species in solution promotes the deposition of a praseodymium- hydroxycarbonate film.

5.3.7 Effect of Solution pH: It has been reported that the pH at the metal/electrolyte interface can vary across the alloy during localized corrosion 34. Thus, it was of interest to investigate the effect of solution pH on the inhibition mechanism by praseodymium.

Figure 5.16 shows naturally aerated polarization curves in pH adjusted 3 (acidic) and 10 (alkaline) solutions. In acidic solution, praseodymium has no apparent inhibition effect compared to the negative control. . However, in alkaline solution, there is a slight decrease in the oxygen reduction kinetics compared to the negative control.

Free corrosion experiments were also performed under these conditions to further assess the inhibition behavior of Pr3+ in acidic and alkaline solution. SEM-EDS analysis after 3 days of exposure in acidic solution containing 0.1 M NaCl + 0.2 mM PrCl3 revealed the presence of a film over the surface that is rich in sulfur and with trace amounts of praseodymium (Figure 5.17). SEM analysis also revealed higher concentration of sulfur and oxygen over the Cu-rich intermetallic particles suggesting a preferential deposition of a sulfur-based film over these particles. The only source of sulfur is from H2SO4, which was added to adjust the pH of the solution. Therefore, it is likely that the film formed over the surface at low pH is not owed to Pr3+ inhibition but to the presence of this incidental sulfur concentration. It is possible that adsorption of sulfur-containing species is favored at low pH given the isoelectric point of oxide- covered aluminum (i.e. pH ~ 9.5) 35. The adsorbed species can then react with metal

189 cations as the surface is attack. Furthermore, according to the polarization curve shown in Figure 5.16b, it is clear that the resulting sulfur-based film does not provide any protection to the alloy substrate.

In alkaline solution, SEM analysis revealed the deposition of corrosion product across the surface rich in praseodymium and pitting in the Al matrix. Furthermore, thick- oxide precipitates rich in praseodymium were present over both S-phase and Fe- containing intermetallic particles (Figure 5.18). However, virtually no corrosion inhibition was observed electrochemically to correlate with the oxide precipitation

(Figure 5.18b). This could be attributed to Cu-enrichment of the surface as the Al passive film is chemically dissolved at high pH, thereby increasing the effective area for oxygen reduction to occur. Furthermore, the enrichment of Cu on the surface leads to enhanced localized attack which is consistent with the pitting observed on the Al matrix.

5.3.8 Mechanism of Inhibition: At near-neutral pH, the aluminum surface is passivated except at the Cu-containing intermetallic particles where the oxide is thinner and more defective due to the presence of different elements in the intermetallic phase 36. These

Cu-rich intermetallic particles support oxygen reduction leading to a local increase in pH associated with hydroxyl ion formation as described by the following reaction:

Eq. 5.3

The polarization curves and free-corrosion experiments conducted in naturally aerated solution suggest that praseodymium imparts cathodic inhibition by forming a film over the Cu-rich intermetallic particles. Thus, as proposed by Hinton et al. 7, 32, the corrosion inhibition mechanism involves the formation of an insoluble film over the intermetallic

190 particles in response to an increase in the local pH at the metal/electrolyte interface. This film provides a barrier to the diffusion of oxygen and the transport of electrons to the oxygen reduction reaction. Furthermore, the decreased corrosion resistance in deaerated electrolyte suggests the inhibition mechanism is oxygen-dependent. This could be attributed to a decrease in the cathodic kinetics in the absence of oxygen, resulting in lower increases in pH at the metal/electrolyte interface and a decrease in the deposition of praseodymium species over the intermetallic particles.

The inhibition mechanism requires the presence of both oxygen and CO2 in the

3+ environment. In the presence of CO2, the critical pH for Pr solubility is shifted to lower values resulting in the formation of ( ) (Figure 5.1b). According to XPS and

XRD, is present over the alloy surface after exposure to aerated (with atmospheric CO2) environments. Thus, it can be concluded that ( ) forms in solution over the intermetallic particles as the local pH increases at the metal/electrolyte interface due to oxygen reduction. This carbonate species react with the hydroxyl ions that generate during oxygen reduction, leading to the formation and deposition of

over the intermetallic particles as described by the following equation:

( ) Eq. 5.4

Furthermore, as made evident by XRD, the surface of the alloy becomes completely covered by the praseodymium-carbonate film as old cathodes become blocked and new sites become activated. It should be noted that the presence of praseodymium- hydroxycarbonate species was previously observed in the scribed areas of coated panels

191 after exposure to salt spray 22. However, the exact form of the hydroxycarbonate species was not confirmed in their study.

In the absence of CO2, characterization of the surface by XPS suggests that praseodymium is present in the form of ( ) . This is consistent with the chemical equilibrium diagram depicted in Figure 5.1a where the dominant species in the alkaline regime is ( ) . However, according to the impedance spectroscopy experiments, the inhibition performance is diminished in the absence of CO2. This could be attributed to the deposition of a denser film in aerated environments. According to the literature, the density of praseodymium carbonate species was found to be 4.686 g/cm3; the density of praseodymium hydroxide is 3.095 g/cm3 37. Hence, it is possible that the formation of

on the surface provides a more effective barrier against oxygen diffusion.

The result is lower oxygen reduction kinetics and a decrease in the overall corrosion rate.

3.4.9 In Situ Atomic Force Microscopy: In situ AFM scratching involves rastering a region in contact-mode with a Si cantilever at sufficiently high forces to destabilize passivity 23. This technique provides real time information of the early stages of localized corrosion with sub-microscopic resolution. When analyzing the evolution and kinetics of attack, three things have to be considered: the force that initiates attack, the time for attack to nucleate, and the corrosion morphology. Higher forces and longer nucleation times imply greater inhibition performance. This technique was used to further investigate the complex interaction between alloy microstructure and praseodymium ions in 0.1 M NaCl solution.

192

Figure 5.19 shows images of the Volta potential and topography distribution of the scratched area prior to scratching, and the chronological sequence of attack in aerated

-5 0.1 M NaCl with 10 M PrCl3. The Volta potential map obtained with SKPFM reveals the presence of a large Fe-containing intermetallic particle and smaller S-phase particles that were confirmed using SEM-EDS. The OCP was allowed to stabilize for 15 min after

PrCl3 solution was introduced into the glass cell before scratching was commenced.

There is evidence of deposition prior to scratching over the intermetallic particles. This is made evident by comparing the topography images in Figure 5.19 for the sample prior to exposure and at 0 h, which is just after immersion. In the latter image, the height of the S-phase and Fe-containing intermetallic particles increased by an average of 40 nm and 100 nm with respect to the matrix, respectively. This suggests that praseodymium begins to deposit over the intermetallic particles early during the corrosion process. In addition, no attack was observed during the first 30 min of scratching at 100 nN.

Increasing the scratching force 125 nN after 30 min resulted in attack of an S-phase particle and on the Al matrix adjacent to the larger Fe-containing intermetallic particle

(attack labeled by arrows in the topography image). No further attack was observed until the scratching force was increased to 150 nN after 3 h of scratching. Increasing the scratching force to much higher values (400 nN) resulted in further attack of the S-phase particle. Note that the dark shade on the right hand side of the intermetallic particles is not attributed to attack, but rather to an imaging artifact associated with drifting of the

AFM scanner.

193

In inhibitor-free solution (not shown), the form of attack observed was trenching at the periphery of the larger Fe-containing particle and complete removal of the S-phase particle after 20 minutes of scratching at a low constant force (100 nN). Similar results were obtained by previous researchers 10, 23. The above results suggest that the resistance

-5 to localized corrosion of AA2024-T3 is increased by the addition of 10 M PrCl3 because no attack was observed until the scratching force was increased to 125 nN. In addition, only one of the S-phase particles in the scratched region was attacked during the experiment. Comparing the 3 h and 8 h topography maps also indicated partial removal of the deposited film over a large and small intermetallic particle (labeled in the images by dashed lines). This is supported by an average height decrease of 30 and 40 nm with respect to matrix for the smaller and larger intermetallic particle, respectively. However, the absence of attack after film removal can be attributed to the reformation of a film over the intermetallic particles because of the continued exposure to inhibitor-containing solution.

Figure 5.20 shows the subsequent SEM-EDS analysis of the scratched area after the experiment. The secondary electron image clearly shows evidence of attack at the periphery of the intermetallic particles and the complete removal of the attacked S-phase particle. In addition, EDS mapping shows that Mg is not present over the intact S-phase particles. Furthermore, no praseodymium was detected over the surface after the scratching experiment which could be related to the detection limits of EDS.

The inhibition performance by praseodymium was improved for a PrCl3 concentration of 0.2 mM. This was the optimum concentration for corrosion inhibition

194 observed by the electrochemical polarization curves (Figure 5.3). Figure 5.21 shows the

SKPFM images and a secondary electron image of the scratched area prior to the experiment which contains both Fe-intermetallic particles and S-phase particles. It should be noted that only the Fe-intermetallic particles are observable in the SKPFM images. Figure 5.22 shows topography images captured at different times during the scratching experiment. There is evidence of strong deposition prior to scratching over the intermetallic particles. This is made evident by comparing the topography image in

Figure 5.21 with the 0 h topography map in Figure 5.22. In the latter image, the height of the S-phase and Fe-containing intermetallic particles increased by an average of 360 and

250 nm with respect to the matrix, respectively, which can be attributed to deposition.

For this experiment, the OCP was allowed to stabilize for 30 min after PrCl3 solution was introduced into the glass cell. In addition, no attack was observed until the scratching force was increased to 200 nN after 4 h of scratching. Attack nucleated at the periphery of the Fe-containing intermetallic particles (attack labeled by arrows in the topography image) and no significant corrosion was observed after increasing the scratching force.

However, increasing the scratching force to the limit of the AFM tip (525 nN) resulted in removal of the deposit present over a smaller intermetallic particle (labeled by an arrow in Figure 5.22). Due to the small size of this intermetallic particle, attempts to identify it by EDS were not successful. Furthermore, it is evident that the particle was removed from the surface during scratching.

Figure 5.23 shows the subsequent SEM-EDS analysis of the scratched area. The secondary electron image clearly makes evident the presence of a thick film over both the

195

S-phase and Fe-containing intermetallic particles. Furthermore, the EDS analysis shows that the film is rich in Pr. The detection of Mg over the S-phase particles suggests that S- phase dissolution was suppressed.

It should be noted that attack always nucleated at the periphery of the Fe- containing intermetallic particles. As described above, electrochemical variances over these particles during exposure may result in pH gradients and the non-uniform deposition of corrosion product in the presence PrCl3. As a result, the film over the Fe- intermetallic particles may not provide a barrier strong enough to stop the anodic dissolution of the adjacent matrix.

The AFM scratching experiments provided strong evidence for the inhibition efficacy of PrCl3 solution. In the absence of inhibitor, corrosion always nucleated at the periphery of the intermetallic particles. In addition, S-phase particles suffered complete removal after 20 minutes of scratching at a low constant force (100 nN). The addition of

-5 10 M and 0.2 mM PrCl3 dramatically changed the kinetics of attack. Deposition prior to scratching suggested that praseodymium begins to deposit over the intermetallic particles

-5 early during the corrosion process. Scratching in 10 M PrCl3 resulted in S-phase particle attack at low tip pressures. In contrast, scratching at the maximum AFM tip force did not result in S-phase particle dissolution with the addition of 0.2 mM PrCl3.

However, attack always nucleated at the periphery of the Fe-containing intermetallic particles at relatively low tip forces, but did not spread with increasing scratching force.

196

5.4 Conclusions

The mechanism of corrosion inhibition of AA2024-T3 by praseodymium was investigated using electrochemical, microscopic, and spectroscopic techniques. The following conclusions can be obtained:

1. Electrochemical polarization curves show that praseodymium provides

cathodic inhibition at an optimum concentration of 0.2 mM. The largest effect

observed was an order of magnitude decrease in the oxygen reduction

kinetics. In addition, at higher PrCl3 concentrations the corrosion inhibition

performance is diminished in response to a decrease in pH.

2. Corrosion inhibition of AA2024-T3 by praseodymium involves the formation

of an insoluble film over the intermetallic particles in response to an increase

in the pH at the metal/electrolyte interface. This film provides a barrier to the

diffusion of oxygen and the transport of electrons to the oxygen reduction

reaction. Furthermore, a thick-film forms across the alloy surface as old

cathodes become blocked and new sites become activated.

3. Surface characterization by XPS and XRD revealed that the film is composed

of in aerated (with atmospheric CO2) solution. In the absence of

CO2, it was suggested that the film is composed of ( ) .

4. Electrochemical impedance showed that the inhibition mechanism by PrCl3

requires the presence of both oxygen and CO2 in the environment. In

deaerated electrolyte, the total impedance decreased with the addition of

PrCl3, which could be attributed to a decrease in pH with the addition of

197

PrCl3. In addition, a lower rate of cathodic kinetics in the absence of oxygen

could lead to lower increases in pH at the metal/electrolyte interface, thereby

establishing conditions where the deposition of praseodymium species over

the intermetallic particles is not favorable. In the absence of CO2, the total

resistance decreased and the film capacitance increased.

5. The corrosion morphology and kinetics of AA2024-T3 in 0.1 M NaCl solution

with praseodymium was investigated using in situ AFM scratching. PrCl3

-5 reduced the dissolution rate of AA2024-T3. The addition of 10 M PrCl3

resulted in the formation of a film over the intermetallic particles, which was

easily removed at lower scratching forces. Increasing the concentration of

PrCl3 to 0.2 mM resulted in the formation of a thicker film over the

intermetallic particles. S-phase dealloying was suppressed only at 0.2 mM

PrCl3, which was confirmed by subsequent SEM-EDS analysis. Furthermore,

attack always nucleated at the periphery of the Fe-containing intermetallic

particles.

198

FIGURES AND TABLES

Figure 5.1. Chemical equilibrium diagram for 1 mM Pr3+ in 0.1 M NaCl solution in the (a) absence and (b) presence of atmospheric CO2. Specie diagram generated using MedusaTM software. The dashed black line signifies maximum solubility.

199

 [ C l ] T O T = 1 0 0 . 0 0 m M L o g P C O =  3 . 5 0 + + 2 [ H ] T O T = 1 . 0 0 E  1 2 M [ N a ] T O T = 1 0 0 . 0 0 m M

6.0

5.5

5.0 pH

4.5

4.0 0.0 0.2 0.4 0.6 0.8 1.0 [ P r 3 + ] M T O T

TM Figure 5.2. Effect of PrCl3 concentration on solution pH. Generated with Medusa .

200

-400

-600 No inhibitor 0.03 mM 0.2 mM -800 0.3 mM 4.0 mM

Potential (mV vs. SCE) vs. (mV Potential -1000

-1200 10-9 10-8 10-7 10-6 10-5 0.0001 0.001 0.01

Current Density (A/cm2)

Figure 5.3. Naturally aerated polarization curves of AA2024-T3 in 0.1 M NaCl solution at varying PrCl3 concentrations.

201

Figure 5.4. (a) Open-circuit potential macrographs and (b) optical micrographs of samples after 1 day exposure in 0.1 M NaCl solution with and without 0.2 mM PrCl3.

202

Figure 5.5. Secondary electron micrographs of scanned area containing S-phase particles before and after 2 days exposure in 0.1 M NaCl solution with 0.2 mM PrCl3.

203

204

Figure 5.7. (a) Topography and (b) Volta-potential images of Fe-containing intermetallic particles before and after 5 hours exposure to 0.2 mM PrCl3 in 0.1 M NaCl containing solution. Scan size = 65 x 65 µm2. Z range = 500 nm (topography); 500 mV (Volta- potential map).

205

Figure 5.8. Spectra showing the (a) Pr 3d5/2 and (b) Pr 4p3/2 peaks after 2 days of exposure in A – 0.1 M NaCl and B – 0.1 M NaCl + 0.2 mM PrCl3 solution.

206

b Figure 5.9. Spectra showing the (a) O 1s and (b) C 1s peaks after 2 days of exposure in 0.1 M NaCl + 0.2 mM PrCl3 solution. 207

Figure 5.10. (a) Macrograph of sample after 30 days immersion in 0.1 M NaCl solution with PrCl3 (b) Secondary electron images of surface after exposure and (c) XRD spectrum of surface film. 208

b Figure 5.11. (a) Open-circuit potential macrographs and (b) optical micrographs of partially exposed samples after 2 day immersion in 0.1 M NaCl + 0.2 mM PrCl3 under decarbonated and deaerated solutions. Dashed lines signify exposed area.

209

Table 5.1. Chemical quantification analysis obtained from high-resolution XPS spectra after 2 day exposure in different environments.

Total Atomic % Solution Mg Cu O Al Cl Pr C* 0.1 M NaCl - 22.40 0.17 0.15 52.51 23.72 1.04 - aerated (1.60) PrCl3 added - 41.79 - - 39.33 16.95 0.83 1.10 aerated (3.20) PrCl3 added - 25.88 0.10 0.30 50.98 21.83 0.58 0.33 deaerated (2.74) PrCl3 added - 37.43 0.05 - 41.55 18.44 1.00 1.53 decarbonated (2.53)

* Concentrations in parenthesis correspond to carbonate phases.

210

Figure 5.12. Spectra showing the Pr 3d5/2 and O 1s peaks after 2 days of exposure in (a) decarbonated and (b) deaerated 0.1 M NaCl solution with 0.2 mM PrCl3.

211

Figure 5.13. Spectra showing the Pr 4p3/2 peaks after 2 days of exposure in A – decarbonated and B – deaerated 0.1 M NaCl solution with 0.2 mM PrCl3.

212

106

5 PrCl added (aerated)

10 3 ) 2 4 PrCl added 10 3 (decarbonated)

1000

IZI (Ohm.cm IZI 0.1 M NaCl (aerated)

100

10 0.01 0.1 1 10 100 1000 104 105

Frequency (Hz)

-80 PrCl added (aerated) 3

-60 PrCl added 3 (decarbonated)

-40 Theta(degrees)

-20 0.1 M NaCl (aerated)

0.01 0.1 1 10 100 1000 104 105

Frequency (Hz)

b Figure 5.14. (a) Bode magnitude and (b) phase angle plots of 2024-T3 coupons immersed in 0.1 M NaCl with and without PrCl3 in naturally aerated and decarbonated solution.

213

106

105 0.1 M NaCl

) 2 104

1000 PrCl added

IZI (Ohm.cm IZI 3

100

10 0.01 0.1 1 10 100 1000 104 105

Frequency (Hz)

0.1 M NaCl 0.1 M NaCl -80 0.1 M NaCl + 0.2 mM PrCl3

-60

-40

Theta(degrees) PrCl added 3 -20

0.01 0.1 1 10 100 1000 104 105

Frequency (Hz)

b Figure 5.15. Bode magnitude and phase angle plot of 2024-T3 coupons immersed in deaerated 0.1 M NaCl with and without 0.2 mM PrCl3.

0.1 M NaCl 0.1 M NaCl + 0.2 mM PrCl3 214

-400

-600

-800 pH 3

-1000

Potential (mV vs. SCE) vs. (mV Potential 0.1 M NaCl -1200 0.2 mM Pr(3+) added

-1400 10-9 10-8 10-7 10-6 10-5 0.0001 0.001 0.01

Current Density (A/cm2)

-200 No inhibitor -400 0.2 mM Pr(3+) added

-600

-800

-1000 Potential (mV vs. SCE) vs. (mV Potential

-1200 pH 10

-1400 10-9 10-8 10-7 10-6 10-5 0.0001 0.001 0.01

Current Density (A/cm2)

Figure 5.16. Naturally aerated polarization curves in 0.1 M NaCl with 0.2 mM PrCl3 in (a) pH 3 and (b) pH 10 solutions. 215

Figure 5.17. SEM-EDS analysis after 3 days of exposure in 0.1 M NaCl with 0.2 mM PrCl3 in pH 3 solution.

216

Figure 5.18. SEM-EDS analysis after 3 days of exposure in 0.1 M NaCl with 0.2 mM PrCl3 in pH 10 solution.

217

-5 Figure 5.19. In situ AFM scratching in 0.1 M NaCl with 10 M PrCl3. Scan size = 20 µm. Top left image is Volta Potential map before exposure with z range = 500 mV. The other images are topographic maps with z range = 200 nm. The first is the map prior to exposure and the rest are in situ images during rastering at indicated times and setpoint voltages.

218

Figure 5.20. Ex situ SEM-EDS analysis after AFM scratching in 0.1 M NaCl with 10-5 M PrCl3.

219

Figure 5.21. SKPFM and SEM analysis of the selected area before in situ AFM scratching in 0.1 M NaCl + 0.2 mM PrCl3.

220

Figure 5.22. In situ AFM scratching in 0.1 M NaCl with 0.2 mM PrCl3. Scan size=35µm. Z range=500 nm.

221

Figure 5.23. Ex situ SEM-EDS analysis after AFM scratching in 0.1 M NaCl with 0.2 mM PrCl3.

222

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CHAPTER 6: CONCLUSIONS AND RECOMMENDATIONS FOR FUTURE WORK

6.1 Conclusions

The pursuit to find a chromate-alternative has led to the development of several chromate-free aerospace primers that offer good protection. However, fundamental understanding of the functionality of the chromate-free pigments that are embedded within these primers is lacking. Therefore, the objective of this dissertation was to investigate how selected chromate-free inhibitors impart corrosion inhibition on aluminum alloy 2024-T3, with the intent of developing the kind of understanding that was accomplished with chromate. The inhibitors selected for this study are found in

2- current commercial non-chromate coatings and include molybdate (MoO4 ), silicate

2- 3+ (SiO2 ), and praseodymium (Pr ). Furthermore, a mechanism detailing the functionality of the aforementioned inhibitors at varying environmental conditions is proposed. The most relevant findings pertaining to the corrosion inhibition by molybdate are:

1. Electrochemical polarization curves showed that molybdate provides mixed

inhibition at a threshold concentration of 0.1 M. The largest effect observed

was a 250 mV increase in the breakdown potential associated with pitting

and a 350 mV decrease in OCP.

2. SEM/EDS analysis of a sample exposed to molybdate-containing NaCl

solution showed suppression of S phase particle attack. Furthermore, a film

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was observed over both S phase and Fe-containing intermetallic particles

that was rich in Mo and O.

3. Cathodic amperometry in aerated 0.1 M NaCl showed a large peak in

cathodic current after injection of concentrated molybdate solution. The

current peak was followed by a decrease to net anodic current before

reaching a steady state cathodic current.

4. Injection of concentrated molybdate solution during anodic amperometry

resulted in a rapid decrease in anodic current. A 99% reduction in current

compared to the pre-injection state was observed after injection indicating

that the inhibitor strongly impedes the growth of active pits.

5. X-ray photoelectron spectroscopy (XPS) after exposure to aerated solution

showed that Mo is present on the surface as Mo(VI). However, in deaerated

conditions Mo is present predominantly in a reduced state pertaining to

MoO·(OH)2.

6. Raman spectroscopy was used to elucidate the species present on the surface

after exposure to aerated solution. The results confirmed the existence of

6- 4- Mo(VI) species pertaining to polymolybdates (Mo7O24 and Mo8O26 ). In

addition, intermediate molybdenum oxides were detected on the surface,

primarily over the S-phase particles and Al matrix.

7. Electrochemical impedance showed a decrease in the total impedance with

the addition of molybdate in deaerated electrolyte where hydrogen evolution

dominate the cathodic kinetics. In aerated solution, the total resistance

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increased by 40x in the presence of molybdate. This indicates that corrosion

inhibition of AA2024-T3 by molybdate is oxygen-dependent owing to the

protection afforded by Mo(VI) species.

8. Corrosion inhibition of AA2024-T3 by molybdate may occur following a

two-step process whereby molybdate is rapidly reduced to form MoO·(OH)2

over the intermetallic particles, and is subsequently oxidized to intermediate

molybdenum oxides (e.g. Mo4O11) in the presence of oxygen. This in turn

may lead to a local acidification, promoting the condensation and

polymerization of molybdate species in solution to form polymolybdate

6- 4- species (Mo7O24 and Mo8O26 ). Furthermore, S-phase particle dissolution

is decreased, suppressing surface copper enrichment and significantly

lowering oxygen reduction kinetics.

9. Inhibition of the Al matrix does not involve the same phenomena as

described above. Instead, the Al matrix is protected following an adsorption

mechanism of Mo(VI) species in solution given the higher isoelectric point

of oxide-covered aluminum. Raman spectroscopy analysis over the matrix

also suggests that the adsorbed species condense and polymerize to form

polymolybdate species.

Silicates are also commonly employed as pigments in current organic primers. In this study, the inhibition mechanism by sodium silicate was studied. The relevant findings pertaining to the corrosion inhibition by silicate are:

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10. Electrochemical polarization curves showed that silicate provides strong

anodic inhibition at a threshold concentration of 0.01 M. The largest effect

observed was a 1 V increase in the pitting potential at 0.025 M silicate

concentration.

11. SEM-EDS analysis revealed partial dissolution of the S-phase particles.

However, high concentrations of Si and O were detected over these particles

using EDS. Furthermore, Si and O were also detected within pits on the

surface and at the larger Fe-containing intermetallic particles.

12. XPS sputter depth profiles of the surface indicated that a film forms on the

surface rich in Si, Al, and O. Na was also detected throughout the film,

albeit at much smaller concentrations. Furthermore, analysis of the Si 2s

spectrum during etching revealed a single peak consistent with

aluminosilicate compounds.

13. Corrosion inhibition of AA2024-T3 by silicate in alkaline solution conforms

to the mechanism suggested by Gaggiano et al. for pure aluminum 1, 2. It

was proposed that aluminosilicate is formed by the reaction of silicate

anions in solution and the aluminate ions that form during oxide dissolution.

The Na+ ions adsorb to the negatively charged surface and coordinate with

the hydroxyl group of the aluminosilicate anions, thereby forming a

protective thin-film over the Al matrix. Furthermore, S-phase particle

dissolution is suppressed owing to the formation of magnesium-silicate and

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alkali-aluminosilicate species that deposit over the particles after the onset

of corrosion.

14. Electrochemical impedance revealed a time-dependent evolution of silicate

inhibition that might be associated with film growth on the surface. This is

supported by an increase in the total resistance, along with a decrease in the

film capacitance, with time.

15. At near-neutral pH, silicate partially blocks attack of the intermetallic

particles by a precipitation mechanism that results in the formation of silica-

and silicate- based derivatives.

16. In acidic solution, it was shown that activation of the aluminum surface

promotes the formation of a mixed silica/aluminosilicate film over the

surface that is porous in nature, providing very little corrosion protection.

17. Strong synergy was observed when mixing silicate and molybdate together

in solution. The threshold concentration for corrosion inhibition was

lowered by almost an order of magnitude in comparison to silicate alone. It

is proposed that molybdate reduces over the intermetallic particles prior to

the preferential deposition of a thin aluminosilicate film over the surface.

Praseodymium hydroxide is currently employed as a pigment in organic coatings.

In this study, the inhibition mechanism by praseodymium chloride was studied. The relevant findings pertaining to the corrosion inhibition by praseodymium are:

18. Electrochemical polarization curves show that praseodymium provides

cathodic inhibition at an optimum concentration of 0.0002 M. The largest

230

effect observed was an order of magnitude decrease in the oxygen reduction

kinetics. In addition, at higher PrCl3 concentrations the corrosion inhibition

performance is diminished in response to a decrease in pH.

19. SEM/EDS analysis of a sample exposed to praseodymium-containing NaCl

solution showed suppression of S phase particle attack. Furthermore, a

thick-oxide that was rich in Pr and O was observed over several

intermetallic particles.

20. Surface characterization by XPS and XRD revealed that the film is likely

composed of after exposure in aerated (with atmospheric CO2)

solution. The same phase was found in deaerated conditions. However, in

the absence of CO2, it was suggested that ( ) was the form of

praseodymium deposited on the surface.

21. Corrosion inhibition of AA2024-T3 by praseodymium involves the

formation of an insoluble film over the intermetallic particles in response to

an increase in the pH at the metal/electrolyte interface. This film provides a

barrier to the diffusion of oxygen and the transport of electrons to the

oxygen reduction reaction. Furthermore, a thick-film deposits across the

alloy surface as old cathodes become blocked and new sites become

activated.

22. Electrochemical impedance showed a decrease in the total impedance with

the addition of PrCl3 in deaerated electrolyte, which suggests that the

inhibition mechanism is oxygen-dependent. This could be attributed to a

231

decrease in pH with the addition of PrCl3. In addition, a lower rate of

cathodic kinetics in the absence of oxygen could lead to lower increases in

pH at the metal/electrolyte interface, thereby establishing conditions where

the deposition of praseodymium species over the intermetallic particles is

not favorable.

23. In the absence of CO2, the total resistance decreased and the film

capacitance increased in comparison to aerated (with CO2) conditions,

indicating that CO2 plays an important role in the inhibition mechanism by

PrCl3.

6.2 Implications of Research Findings

The evidence presented in this work has shed further light on the inhibition

2- 2- 3+ mechanism of molybdate (MoO4 ), silicate (SiO3 ), and praseodymium (Pr ) in bulk aqueous solution. Moreover, the functionality of these non-chromate inhibitors in various environmental conditions (oxygen-free, CO2, pH, etc.) was also addressed. In the case of molybdate, it was inferred that the inhibition mechanism is highly influenced by the presence or absence of oxygen. Electrochemical impedance showed a decrease in the total impedance with the addition of molybdate in deaerated electrolyte. In contrast, the total resistance increased by 40x with respect to the control in aerated solution. This indicates that the corrosion inhibition of AA2024-T3 by molybdate is oxygen-dependent, and the use of molybdate in practice should be considered only in aerated environments.

Silicate is a strong anodic inhibitor in alkaline conditions. At an optimum concentration of 25 mM the pitting potential increased by 1 V. The results showed that 232 the inhibition mechanism involves the dissolution of the aluminum oxide film at a high pH, promoting the formation of a thin-film on the surface composed of an ionic network of Na+ ions and aluminosilicate species. At lower pH, a thick, porous film forms over the surface rendering poor corrosion inhibition. Note that the inhibition mechanism in alkaline solution is governed by adsorption of aluminosilicate species on the surface, while in the acidic regime the formation of a film is dominated by the precipitation of corrosion product. Furthermore, the discrepancy in the inhibition mechanism observed at high and low pH can be explained in terms of the silicate speciation in solution. At low pH, the dominant species is silica, or SiO2. However, the effectiveness of silicate as a corrosion inhibitor for AA2024-T3 depends on the concentration of available monomeric silicate anions in solution. These anions form complex compounds that adsorb on the metal surface and significantly decrease the rate of metal dissolution. Thus, the use of silica-based pigments should be avoided, unless conditions are applied in practice where the pH is maintained high enough to promote silicate anion formation.

Praseodymium imparted cathodic inhibition at near-neutral pH, and at a similar critical concentration as chromate 3. However, in acidic and alkaline solution, praseodymium renders poor corrosion inhibition. At high pH, it was shown that praseodymium forms a thick film over the intermetallic particles as a result of a local increase in pH at the metal/electrolyte interface, but does not stop the uniform dissolution of the Al-oxide film. Optimum inhibition was observed in near-neutral pH, where it is proposed that a film composed of praseodymium-hydroxycarbonate species forms over the whole surface in the presence of CO2. In addition, electrochemical impedance

233 showed a decrease in the total resistance with the addition of PrCl3 in deaerated electrolyte. A similar effect was observed in decarbonated solution. Thus, it can be concluded that in practice it is important to incorporate praseodymium in environments where the pH is near-neutral, and where both oxygen and CO2 are present.

Although the inhibitors studied in this investigation provided good corrosion protection, it is important to consider the effect of these inhibitors beyond the realm of bulk solution inhibition. Moreover, caution must be exercised when extending the use of these inhibitors to coating applications. For instance, corrosion inhibition by molybdate was observed at a threshold concentration of 0.1 M. Under these circumstances, the large ratio of molybdate-to-chloride ion concentration is not suitable for organic coating applications. The high concentration of inhibitor release from pigments can lead to coating degradation via osmotic blistering 4. However, these negative effects can be overcome by developing inhibitor combinations. For example, a strong synergistic effect was observed when mixing molybdate and silicate together in solution, which also resulted in a significant decrease in the critical concentration for corrosion inhibition.

Thus, these results show that is possible to help facilitate the development of multi- component protection schemes that provide the same level of protection and reliability as chromate-based systems.

6.3 Recommendations for Future Work

This dissertation has provided new insight on the inhibition mechanisms of selected chromate-free inhibitors in hopes of helping to facilitate improvements in non-chromate protection schemes. However, during the course of this study, interesting observations

234 were encountered that require further scrutiny. In this regard, the following recommendations should be considered to promote further understanding of non- chromate technologies:

1. One important characteristic of chromate conversion coatings is their ability

to self-heal in the presence of a defect. Therefore, it would be interesting to

test the self-healing ability of the inhibitors investigated in this study. This

can be accomplished using the artificial scratch cell technique.

2. Strong synergy was observed when mixing silicate and molybdate together

in solution. The threshold concentration for corrosion inhibition was

lowered by almost an order of magnitude in comparison to silicate alone.

The low concentration observed is ideal for organic coating applications. It

has been proposed that high concentrations of dissolved species necessary

for inhibition can lead to coating degradation via osmotic blistering 4.

Future investigations should focus on the formation of suitable pigments

that can supply the optimum ratio of silicate-to-molybdate species to the

damaged area. In addition, overall corrosion protection of the pigmented

coatings should be assessed using electrochemical impedance spectroscopy

in combination with standard salt fog tests.

3. Cerium- and zinc- molybdate have been considered to be promising

substitutes for chromate pigments5, 6. Likewise, the use of praseodymium

can be considered in combination with molybdate and other inorganic

inhibitors that do not effectively suppress cathodic kinetics. Furthermore,

235

the data presented in this study showed that praseodymium imparts cathodic

inhibition at low concentrations, a characteristic that is ideal for organic

coating applications.

4. It is not surprising that the heterogeneous microstructure of the alloy has a

large influence on the corrosion inhibition mechanism of AA2024-T3. For

instance, there is a distinct effect of praseodymium deposition between the

S-phase and Fe-containing intermetallic particles. Very little deposition was

observed over the latter particles in comparison to the S-phase. In another

example, free- corrosion exposure experiments coupled with SEM showed

preferential deposition of silica/silicate species over the S-phase particles.

Therefore, future work should involve a phase by phase electrochemical

analysis of the constituent particles present in AA2024-T3 and exposed to

inhibitor-containing solution. This can be accomplished using the

electrochemical microcell technique which was a useful approach for studies

involving the corrosion inhibition of AA2024-T3 by vanadates 7.

236

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