Laser Surface Treatment of 7075 Aluminium Aerospace Alloy

Excimer Laser Surface Melting Treatment on 7075-T6

Aluminium Alloy for Improved Corrosion Resistance

A thesis submitted to The University of Manchester for the degree of

Doctor of Philosophy

in the Faculty of Engineering and Physical Sciences

2012

Bader M H M Elkandari

School of Materials

Table of Contents

Table of Contents ...... 2 List of Tables ...... 6 List of Figures ...... 7 List of Abbreviations ...... 16 Abstract ...... 19 Declaration ...... 21 Copyright Statement ...... 22 Acknowledgments ...... 23 Chapter 1 Introduction ...... 24 1.1 Rational of the Research Project ...... 24 1.2 Research Objectives ...... 25 1.3 Thesis Layout ...... 26 Chapter 2 Literature Review ...... 28 2 Introduction to Aluminium Alloys ...... 28 2.1 7xxx Aluminium Alloys ...... 30 2.2 Elements and Impurities in 7xxx Aluminium Alloys ...... 34 2.3 Stages of Heat Treatment of 7xxx Aluminium Alloys ...... 38 2.3.1 Age Hardening ...... 39 2.3.2 Precipitate Free Zones (PFZ) ...... 40 2.3.3 Hydrogen Effects...... 41 2.3.4 Temper Effects ...... 43 2.4 List of Figures ...... 48 Chapter 3 Corrosion ...... 52 3 Introduction to Corrosion ...... 52 3.1 Surface Oxide Film ...... 53 3.2 Electrochemical Polarization ...... 53 3.2.1 Thermodynamics ...... 55 3.2.2 Kinetics...... 57 3.2.3 Passivity and Breakdown of Passivity ...... 58

Table of Contents

3.3 Types of Corrosion ...... 59 3.3.1 Uniform Corrosion ...... 60 3.3.2 Localized Corrosion ...... 60 3.3.3 Pitting Corrosion ...... 60 3.3.4 Crevice Corrosion ...... 63 3.3.5 Filiform Corrosion...... 63 3.3.6 Galvanic Corrosion ...... 64 3.3.7 Cavitation Corrosion ...... 64 3.3.8 Waterline Corrosion ...... 65 3.3.9 Intergranular Corrosion (IGC) ...... 65 3.3.10 Stress Corrosion Cracking (SCC) ...... 66 3.4 Corrosion at High Strength Aluminium Alloys ...... 69 3.5 Corrosion Behaviour of Intermetallic Particles ...... 70 3.6 Aluminium Alloy Surface Protection Methods ...... 73 3.6.1 Non-galvanic Methods ...... 74 3.6.2 Chemical Methods...... 74 3.6.3 Electrolytic Methods ...... 74 3.7 List of Figures ...... 76 Chapter 4 Laser ...... 82 4 Introduction ...... 82 4.1 Basic Principle of Laser Works ...... 82 4.2 Laser Light Characteristics ...... 84 4.3 Laser Surface Melting (LSM) ...... 86 4.3.1 Heating Principle ...... 86 4.3.2 Formation of the Melt Pool ...... 87 4.3.3 Cooling and Solidification ...... 88 4.3.4 Nucleation ...... 89 4.3.5 Growth ...... 90 4.4 Effect of LSM Interaction ...... 92 4.5 Effect of LSM Processing Parameters ...... 93 4.5.1 Energy ...... 94 4.5.2 Wavelength...... 95 4.5.3 Pulse Duration and Interaction Time ...... 96 3

Table of Contents

4.5.4 Number of Pulses per Unit Area ...... 97 4.5.5 Process Environment ...... 98 4.6 Microstructure of LSM Alloys ...... 98 4.7 Corrosion Behaviour of LSM Alloys...... 102 4.8 Excimer Laser for LSM ...... 105 4.9 Safety of Using Laser ...... 107 4.9.1 Laser Classification ...... 107 4.9.2 Risk Assessment ...... 109 4.9.3 Protection from Laser ...... 109 4.10 List of Figures ...... 111 Chapter 5 Experimental Procedures ...... 119 5 Introduction ...... 119 5.1 Substrate Materials ...... 120 5.2 Surface Preparation for LSM Treatment ...... 120 5.3 Anodising ...... 121 5.4 Corrosion Investigation...... 122 5.4.1 Potentiodynamic Polarization Test...... 122 5.4.2 Immersion Test ...... 123 5.4.3 Exfoliation (EXCO) Test ...... 123 5.5 Specimen Preparation for Metallographic Examination ...... 123 5.6 Field Energy Gun Scanning Electron Microscopy (FEG-SEM) ..... 124 5.7 Transmission Electron Microscopy (TEM) ...... 125 5.8 X-ray Diffraction (XRD) ...... 125 Chapter 6 Results ...... 127 6 Microstructure and Corrosion of the As-Received AA 7075-T6 Alloy .. 127 6.1 Microstructural Examination ...... 127 6.2 X-Ray Diffraction ...... 128 6.3 Corrosion Test Results ...... 128 6.3.1 Potentiodynamic Polarization Test...... 128 6.3.2 Immersion Test ...... 129 6.3.3 Exfoliation Test ...... 130 6.4 Conclusions ...... 131 6.5 List of Figures ...... 132

Table of Contents

Chapter 7 ...... 144 7 Laser Surface Melting (LSM) Operation Conditions ...... 144 7.1 List of Figures ...... 149 Chapter 8 Microstructure and Corrosion of the LSM AA 7075-T Alloy 159 8 Microstructural Examination ...... 159 8.1 X-ray Diffraction ...... 162 8.2 Corrosion Test Results after LSM ...... 162 8.2.1 Potentiodynamic Polarization Test...... 162 8.2.2 Immersion Test ...... 163 8.2.3 Exfoliation Test ...... 163 8.3 Conclusions ...... 165 8.4 List of Figures ...... 166 Chapter 9 Microstructure and Corrosion of the LSM AA 7075-T6 Alloy after Anodising ...... 188 9 Introduction ...... 188 9.1 LSM and Anodising ...... 188 9.2 Surface and Cross Section Results ...... 189 9.3 Corrosion Test Results ...... 190 9.3.1 Potentiodynamic Polarization Test...... 190 9.3.2 Exfoliation Test ...... 191 9.4 Conclusions ...... 192 9.5 List of Figures ...... 193 Chapter 10 Discussion ...... 207 10 AA 7075-T6 Alloy ...... 207 10.1 AA 7075-T6 Alloy after LSM ...... 211 10.1.1 Corrosion of LSM Alloy ...... 215 10.2 LSM of AA 7075-T6 Alloy after Anodising ...... 219 Chapter 11 Conclusions and Suggestions for Future Work ...... 223 11.1 Conclusions ...... 223 11.2 Suggestions for Future Work ...... 227 Chapter 12 Appendix 1 ...... 228 References ...... 230 Word count 45,930

List of Tables

Table 2-1: Four-digit designation system for Aluminium Alloys [29]...... 48 Table 2-2: Mechanical properties of AA7xxx alloys with different temper processes [28]...... 49 Table 2-3: Solid solubility and influence of elements in Al properties [28]...... 50 Table 4-1: Examples of LSM types and their wavelengths [164] ...... 113 Table 5-1: The chemical composition of the as-received AA 7075-T6 alloy ... 120

Table 6-1: The corrosion potential (Ecorr), pitting potential (Epit), passive

current density (Ipass) and passive range potential (Epass) from Fig 6-8. Potentials are given with respect to SCE...... 140 Table 7-1: Conditions of laser operation (high pulse energy density) ...... 149 Table 7-2: Conditions of laser operation (low pulse energy density) ...... 149

Table 7-3: Corrosion potential (Ecorr), pitting potential (Epit), passive

current density (Ipass) and passive range potential (Epass) for both as-received and laser treated alloys in Fig. 7-8. Potentials are given with respect to SCE...... 158

Table 8-1: Corrosion potential (Ecorr), pitting potential (Epit), passive

current density (Ipass) and passive range potential (Epass) for both as-received and laser treated AA 7075-T6 alloy determined from the potentiodynamic polarisation curves of Fig. 8-17. Potentials are given with respect to SCE...... 181

Table 9-1: Corrosion potential (Ecorr), pitting potential (Epit), passive

current density (Ipass) and passive range potential (Epass) for the as-received (AR), laser-treated (LT), AR and anodised, and LSM and anodised alloys determined from the potentiodynamic polarisation curves of Fig. 9-4. Potentials are given with respect to SCE...... 202

List of Figures

Figure 2-1: Three main stages of heat treatment (SHT refers to solution heat treatment, Q is quenching, and A* is aging) with precipitation scheme for the 7xxx aluminium alloys ...... 51

Figure 3-1: Schematic diagram showing the thin oxide layer (Al2O3) that protects the metal surface (aluminium) from environmental attack...... 76 Figure 4-1: Energy level diagram for a three electron energy level of a laser and a pulse system [126] ...... 111 Figure 4-2: Energy level diagram for a four electron energy level of a continuous (CW) laser system [126] ...... 111 Figure 4-3: Variation in total free energy of solid-liquid system as the size of the solid changes. The solid is stable nuclei with radius above the critical value r* [132] ...... 112 Figure 4-4: Various growth structures that show planar (a), cellular (b), columnar dendrites (c), and equiaxed dendrites (d) [130] ...... 112 Figure 4-5: Schematic representing the solidification transformation from cell (a), cellular dendrites (b), columnar (c), and to columnar dendrite branches (d) [125] ...... 113 Figure 4-6: Schematic illustration of the laser beam interaction with the surface of a material [138] ...... 114 Figure 4-7: Influence of the pulse duration in LSM of pure metals on the maximum depth of the melting layer [159] ...... 114 Figure 4-8: Cross-section of CW-Nd:YAG LSM on AA 2014-T6 alloy, following etching, that shows different regions of the solidification structure [161] ...... 115 Figure 4-9: Modelling results of excimer LSM of AA 2024 alloy for a single pulse that show the solidification growth velocity as a function of position within the modified layer, including different layer thicknesses [166] ...... 115

List of Figures

Figure 4-10: Scanning electron micrographs of 3 kW CW-Nd:YAG LSM AA 2014-T6 alloy showing the variation in growth rate and dendrite spacing with position in modified layer [161] ...... 116 Figure 4-11: Electron probe micro-analysis X-ray mapping of 3 kW CW- Nd:YAG LSM AA 2014-T451 alloy, show element segregation throughout the modified layer [161] ...... 116 Figure 4-12: Variation in cell spacing with melt depth for AA 2024-T3 alloy using a variety of pulsed treatment methods with scanning electron micrographs of the microstructure at different depths [166] ...... 117 Figure 4-13: High resolution EBSD map showing LSM layer morphology obtained with excimer laser treatment in Zr-free alloy (AA 2024) and Zr-containing alloys (AA 7150) [166] ...... 117 Figure 4-14: Cross-section micrographs through excimer LSM treated layers formed over (a) AA 2096 alloy containing Zr and (b) AA 7075 Zr-free alloys, showing the layer grain structure with behaviour opposite to that found in Fig 4-13 [166] ...... 118 Figure 6-1: Scanning electron micrographs showing the surface of the AA 7075-T6 alloy at different magnification after grinding and polishing up to 1 µm diamond paste. (a) Secondary electrons. (b- d) Backscattered electrons...... 132 Figure 6-2: Scanning electron micrographs for the AA 7075-T6 alloy. The cross-section is normal to the rolling direction (a,c). Particles are shown in two different regions, while (b,d) are given the dimensions of the particles and cavities. (a,b) Secondary electrons. (b,d) Backscattered electrons...... 133 Figure 6-3: SEM-EDX line scan measurements that show the elements present in the constituent particles and precipitates in AA 7075- T6 alloy of the cross-section that is shown also in Fig.(8-2). The scans show Cu and Fe in the large particles, while Mg and Si are present at the cavities where residual particles are present. Zn appears to be distributed in both the alloy matrix and the particles...... 134

List of Figures

Figure 6-4 : SEM-EDX line scan measurement on intermetallic particles on the AA 7075-T6 alloy surface. (a,b) Scanning electron micrographs (backscattered electrons). (c) EDX line-scan analysis...... 135 Figure 6-5: SEM-EDX mapping analysis of AA 7075-T6 alloy. (a) Micrograph of the surface revealing the presence of intermetallic particles. (b) Aluminium. (c) Zinc. (d) Silicon. (e) Magnesium. (f) Manganese. (g) Iron. (h) Copper...... 136 Figure 6-6: SEM-EDX mapping analysis of AA 7075-T6 alloy at the region shown previously in Fig 6-3. (a) Micrograph of the surface (b) Aluminium. (c) Zinc. (d) Manganese. (e) Copper. (f) Iron. (g) Magnesium. (h) Silicon...... 137 Figure 6-7: Results of SEM-EDX spot analysis of different area of the matrix of AA 7075-T6 alloy. (a) and (c) Scanning electron micrographs showing the location of the analysis region. (b) and (d) EDX spectrum...... 138 Figure 6-8: Results of low angle X-ray diffraction measurments for the as- received AA 7075-T6 alloy...... 139 Figure 6-9: Anodic potentiodynamic polarization scan of the as-received and polished AA 7075-T6 alloys, in deaerated 0.1 M NaCl. The scan starts at point 1 and terminates at point 2. Point A shows the

corrosion potential (Ecorr), while point B shows pitting potential

(Epit). Point C shows second breakdown potential on AR- polished alloy...... 140 Figure 6-10: Scanning electron micrographs (BSE) of cross-sections of AA 7075-T6 alloy after immersion in 0.1 M NaCl, open to the air, for 24 h. (a, c) Low magnification images. (b, d) Increased magnification images of (a, c). Images (a, b) reveal pitting corrosion adjacent to intermetallic particles on the alloy surface. Images (c, d) show pits free of second phases...... 141 Figure 6-11: Scanning electron micrographs of the top surface of the AA 7075-T6 alloy before and after immersion in EXCO solution for 180 minutes. (a, b) AR alloy before immersion. (c, d) AR alloy after immersion...... 142 9

List of Figures

Figure 6-12: Scanning electron micrographs (BSE) of cross-sections of the AA 7075-T6 alloy after immersion in EXCO solution for (a) 30, (b) 60, (c) 90 and (d) 120 minutes...... 143 Figure 7-1: Schematic illustration of the excimer laser that shows the focusing of the laser beam onto the surface of the alloy sample that is placed on a motorised holder allowing movements in the X and Y directions...... 150 Figure 7-2: Schematic diagram representing the laser scanning pattern and the incidence of the laser pulse on specimen surface...... 151 Figure 7-3: Scanning electron micrographs (backscattered electron) of cross-sections of the LSM AA 7075-T6 alloy after treatment with high pulse energy density (10 J/cm2), using 10 (a), 25 (b), and 50 (c) pulses...... 152 Figure 7-4: Scanning electron micrographs (secondary electron) of the top surface of LSM AA 7075-T6 alloy, after treatment with low pulse energy density. (a) 10 pulses, (b) 25 pulses, (c) 40 pulses and (d) 50 pulses ...... 153 Figure 7-5: Scanning electron micrographs (BSE) of cross-sections of the LSM AA 7075-T6 alloys treated with low pulse energy density (3.3 J/cm2), using 70 (a), 80 (b) and 90 (c) pulses...... 154 Figure 7-6: Average thickness of the laser melted layer and the largest pore size in LSM AA 7075-T6 alloy vs. the number of pulses (NOP), as determined from scanning electron micrographs of cross-sections for laser-treatments, using high and low energy densities...... 155 Figure 7-7: Potentiodynamic polarisation scans in deaerated 0.1 M NaCl solution of the as-received (AR) AA 7075-T6 alloy and after LSM using a high energy density (10 J/cm2), with 10, 25 and 50 pulses...... 156 Figure 7-8: Potentiodynamic polarisation scans in deaerated 0.1 M NaCl condition of the as-received (AR) AA 7075-T6 alloy and after LSM with high energy density (10 J/cm2), using 10, 25 and 50 pulses, and with low energy density (3.3 J/cm2), using 70, 80 and 90 pulses...... 157 10

List of Figures

Figure 8-1: Scanning electron micrographs (SE) of the top surface of the AA 7075-T6 alloy after LSM using 80 pulses. (a), (b), (c) and (d) show images at increasing magnification. The laser track direction and the overlap region (10 %) are indicated...... 166 Figure 8-2: Scanning electron micrographs (BSE) of the AA 7075-T5 alloy after LSM. (a) Alloy with and without treatment by LSM. (b) Untreated alloy, revealing the presence of intermetallic particles. (c) Alloy after LSM treatment...... 167 Figure 8-3: Scanning electron micrographs (SE) of the top surface of the AA 7075-T6 alloy after LSM using 80 pulses. (a) and (b) show laser tracks, re-heating and overlap regions. (c) Higher magnification in overlap region. (d) and (e) Higher magnification in re-heating area...... 168 Figure 8-4: Results of SEM-EDX spot analysis of the AA 7075-T6 alloy after LSM. (a) EDX spectrum. (b) Scanning electron micrograph showing the location of the analysis region. (c) Elements in weight and atomic percentages...... 169 Figure 8-5: Results of SEM-EDX spot analysis of a different area of the LSM AA 7075-T6 alloy from that shown in Fig 8-4. (a) EDX spectrum. (b) Scanning electron micrograph showing the location of the analysis region. (c) Elements in weight and atomic percentages...... 170 Figure 8-6: Scanning electron micrographs of cross-sections of the AA 7075-T6 alloy after LSM using 80 pulses. (a) and (c) Secondary electrons. (b) and (d) Backscattered electrons...... 171 Figure 8-7: High magnification scanning electron micrographs showing a cross-section of AA 7075-T6 alloy after LSM using 80 pulses...... 172 Figure 8-8: Scanning electron micrographs (BSE) of a cross-section of the AA7075-T6 alloy after LSM, with the presence of a thin melted layer on the alloy surface indicated in (a) and (b). Higher magnification images in (c) and (d) show the thinner melted layer produced above a large second phase particle that was partially dissolved after LSM treatment...... 173

List of Figures

Figure 8-9: SEM-EDX mapping analysis of Fig 8-8 (d) showing the LSM layer and the alloy matrix in (a), and the imaging of Al (b), Zn (c), Cu (d), Fe (e), and Mg (f)...... 174 Figure 8-10: Scanning electron micrographs (BSE) of cross-sections of the AA 7075-T6 alloy after LSM using 80 pulses, showing the laser melted layer in more detail, disclosing porosity in the layer with pores of dimension about 1 µm...... 174 Figure 8-11: Transmission electron micrographs of ultramicrotomed cross- sections of the AA 7075-T6 alloy after LSM using 80 pulses. Micrographs (a) and (b) reveal the LSM layer and alloy matrix with pores indicated by arrows in laser-treated layer. Micrograph (c) shows bands in LSM layer, which are due to the presence of precipitates. The bands are separated by distances of 0.3 to 0.4 μm...... 175 Figure 8-12: Transmission electron micrographs of an ultramicrotomed cross-section of the AA 7075-T6 alloy after LSM using 80 pulses. (b) Higher magnification of (a). The Fig. shows the interface region, as indicated by arrows in red, between the LSM layer and the alloy matrix. The arrows in white and black show pores and scratches in the LSM layer...... 176 Figure 8-13: TEM-EDX spot analysis of the micrograph cross-section of Fig. 8-12 (b); the 5 locations of EDX analysis are indicated in (a)...... 177 Figure 8-14: High angle annular dark field (HAADF) images of an ultramicrotomed cross-section of the AA 7075-T6 alloy after LSM using 80 pulses. (b) Higher magnification of (a). (c) Results of of EDX analysis of (b)...... 178 Figure 8-15: Low angle XRD analysis of the AA 7075-T6 alloy after LSM using 80 pulses...... 179 Figure 8-16: Low angle XRD analysis for both as-received and LSM AA 7075-T6 alloy treated with 70, 80, and 90 pulses...... 180 Figure 8-17: Potentiodynamic polarization scans in deaerated 0.1 M NaCl for AA 7075-T6 alloy and after LSM using 80 pulses...... 181

List of Figures

Figure 8-18: Scanning electron micrographs of the AA 7075-T6 alloy after LSM using 80 pulses and immersion in 0.1 M NaCl, open to air, for 24 h. The Fig. reveals that no evidence of corrosion product in low magnification (a, b) and high magnification (c, d) images. .... 182 Figure 8-19: Scanning electron micrographs (BSE) of the AA 7075-T6 alloy after LSM using 80 pulses and immersion in naturally- aerated EXCO solution for (a) 30, (b) 60, (c) 120, (d) 150 and (e) 180 min. The cross-sections of the laser surface treatment alloys (a-d) reveal no corrosion products after immersion for 150 min. With increasing time of immersion, corrosion causes delamination of the LSM layer, as evident in (e)...... 183 Figure 8-20: Scanning electron micrographs of cross-sections of the AA 7075-T6 alloy in the as-received condition and after laser-surface treatment (LT) using 80 pulses, following immersion in (EXCO) solution for 30 min (a, b) and 60 min (c, d)...... 184 Figure 8-21: Scanning electron micrographs of cross-sections of the AA 7075-T6 alloy in the as-received condition and after laser surface treatment alloys using 80 pulses and immersion in EXCO solution for 120 min (a, b) and 150 min (c, d)...... 185 Figure 8-22: Scanning electron micrographs of cross-sections of the AA 7075-T6 alloy in the as-received condition (a, c) and after laser surface treatment (b, d) using 80 pulses immersion in EXCO solution for 180 min...... 186 Figure 8-23: Scanning electron micrographs of LSM AA 7075-T6 revealing micro-cracks in the re-heating area of the 1st laser track as in (a) and (b), while (c) and (d) show micrographs of cross- sections after exfoliation testing for 180 min...... 187 Figure 9-1: Current density (I) vs time (s) behaviour of the as-received (AR) and LSM AA 7075-T6 alloy during anodising at 12 V

(SCE) in 0.46 M H2SO4 at 295 K for (a) 240 s, (b) 720 s, and (c) 1200 s...... 193 Figure 9-2: Scanning electron micrographs of the top surface of the LSM

AA 7075-T6 alloy after anodising at 12 V in 0.46 M H2SO4 at 295 K for 1200 s; (a) the overlap area in the second-laser track 13

List of Figures

and the reheating area in the first-laser track; (b) high magnification of the reheating area; (c) high magnification of the overlap area. The arrows in (c) show the micro-cracks in the re- heating region...... 194 Figure 9-3: Results of SEM-EDX spot analysis of the LSM AA 7075-T6 alloy following anodising. (a) Scanning electron micrograph showing the location of the analysis region. (b) EDX spectrum. (c) Elements in weight and atomic percentages...... 195 Figure 9-4: Results of SEM-EDX spot analysis of the LSM AA 7075-T6 alloy following anodising. (a) Scanning electron micrograph showing the location of the analysis region. (b) EDX spectrum. (c) Elements in weight and atomic percentages...... 196 Figure 9-5: Results of SEM-EDX spot analysis of the LSM AA 7075-T6 alloy following anodising. (a) Scanning electron micrograph showing the location of the analysis region. (b) EDX spectrum. (c) Elements in weight and atomic percentages...... 197 Figure 9-6: Scanning electron micrographs of cross-sections of the LSM AA 7075-T6 alloy following anodising in sulphuric acid for 1200 s. (b) High magnification of (a)...... 198 Figure 9-7: Transmission electron micrograph of an ultramicrotomed cross-section of LSM AA 7075-T6 alloy anodised at 12 V in 0.46

M H2SO4 at 295 K for 1200 s. (b) High magnification of (a)...... 199 Figure 9-8: Scanning electron micrographs of cross-sections of the as- received AA 7075-T6 alloy following anodising in sulphuric acid for 1200 s. (a) to (d) show images at different locations of the alloy surface...... 200 Figure 9-9: Potentiodynamic polarization tests in deaerated 0.1 M NaCl solution of the as-received (AR) AA 7075-T6 alloy, laser surface

treated (LSM) alloy, and following anodising in 0.46 M H2SO4 for 1200 s. (LSM+Anod., AR+Anod.)...... 201 Figure 9-10: Scanning electron micrographs of top surface of the anodised AA7075-T6 alloy after EXCO corrosion test for an exposure time of 180 min. (b) shows high magnification image of (a). (c)

List of Figures

shows high magnification image of the location marked with yellow arrow...... 203 Figure 9-11: Scanning electron micrographs of cross-sections of the anodised AA7075-T6 alloy after EXCO corrosion test for an exposure time of 180 min. (a), (b) and (c) are images at increasing magnification...... 204 Figure 9-12: Scanning electron micrographs of the top surface of the anodised LSM AA 7075-T6 alloy after EXCO testing for a time of 180 min. (a) overlap and reheating areas. (b, c) middle of laser-treated layer...... 205 Figure 9-13: Scanning electron micrographs of cross-sections of the anodised LSM AA7075-T6 alloy after EXCO corrosion test for an exposure time of 180 min. (a) to (d) are images at increasing magnification...... 206

List of Abbreviations

A* Ageing A Area AA Aluminium alloy AR As received AEL Accessible emission limit ASM American society of materials ASTM American standard for testing materials B Brightness BCC Body centred cubic BSD Berkeley software distribution BSE Back-scattered electron c Velocity of light CR Corrosion rate

CO2 Carbon dioxide CW Continuous wave D Density DC Direct current

Ecorr Corrosion potential E° Standard potential

Ee Electrode potential

Epit Pitting potential

Epass Passive potential ranges EDX Energy dispersive X-ray EXCO Exfoliation corrosion f Focal length F Farady’s constant FCC Face centred cubic FEG-SEM Field energy gun scanning electron microscopy G Thermal gradients

GL Thermal gradient of liquid

GS Thermal gradient of solid ∆G Total Gibbs free energy GP Guinier preston zone HAADF High angle annular dark field

List of Abbreviations

HAZ Heat affected zone HRTEM High resolution transmission electron microscopy I Current density

Icorr Corrosion current density

Ipass Passive current density IGC Intergranular corrosion K Equilibrium constant KrF Krypton fluoride lcoh Coherence length LSM Laser surface melting M Mole MEP Maximum permissible exposure n Index of refraction Nd-YAG Neodymium-doped yttrium aluminium garnet NOP Number of pulses OCP Open circuit potential P Power

Pp Peak power PFZ Precipitate free zone pH Hydrogen ion activity r* Critical growth radius R Gas constant R* Growth rate PRF Pulse repetition frequency (Hz) q Emitted power Q Quenching RRA Retrogression and re-ageing SCC Stress corrosion cracking SCE Saturated calomel electrode SDZ Solute-depleted zone SE Secondary electron SHT Solution heat treatment SKPFM Scanning Kelvin probe force microscopy SSSS Supersaturated solid solution

List of Abbreviations t Pulse duration tcoh Coherence time T Temperature

Tb Boiling temperature

Tf Freezing temperature

Tm Melting temperature ∆T Thermal undercooling

Tin Interaction time TEM Transmission electron microscopy UV Ultraviolet V Vapour XeCl Xenon monochloride xe- Number of electrons XRD X-ray diffraction θ Scattering angle 2θ Diffracted angle

θB Brewster angle ω Beam waist radius σ Stefan-Boltzmann constant λ Wavelength

βa Tafal anodic slop

βc Tafal cathodic slop η Over-voltage polarization Ω Solid angle at.% Atomic percent wt.% Weight percent

Abstract

High strength 7xxx aluminium alloys are used extensively in the aerospace industry because the alloys offer excellent mechanical properties. Unfortunately, the alloys can suffer localised corrosion due to the presence of large intermetallic particles at the alloy surface that are aligned in the rolling direction. Laser surface melting (LSM) techniques offer the potential to reduce and/or to eliminate the intermetallic phases from the surface of the alloy without affecting the alloy matrix.

The present study concerns the application of LSM using an excimer laser to enhance the corrosion resistance of AA 7075-T6 aluminium alloy. The initial stage of the project was aimed at optimising the laser conditions for production of a uniform microstructure, with the increase in the corrosion resistance of the alloy being determined by potentiodynamic polarization measurements in sodium chloride solution. Low and high laser energy densities were used with a different number of pulses per unit area to treat the alloy surface, which were achieved by changing both the laser fluence and the pulse repetition frequency. A laser fluence of 3.3 J/cm2 with 80 pulses was subsequently selected as the optimum condition to treat the surface of the alloy. The composition and microstructure of the alloy before and after LSM treatment, and following corrosion tests, were characterized by scanning electron microscopy (SEM), transmission electron microscopy (TEM) and X-ray diffraction (XRD).

After the laser treatment, the surface and the cross-sections of the alloy showed a significant reduction in the number of large intermetallic particles and a

List of Abbreviations relatively homogenous melted layer was generated that provided significant improvement in the resistance of the alloy against corrosion, as assessed by several corrosion test methods, including exfoliation corrosion (EXCO) tests.

However, delamination of the melted layer was observed after extended testing in the EXCO solution which is possibly related to the formation of bands of fine magnesium and zinc-rich precipitates within the melted layer. Therefore, anodising in sulphuric acid was applied to the LSM alloy, in order to further increase the corrosion resistance and to protect the laser treated layer from delamination by generating a thin oxide film over the LSM layer. The results revealed that the anodic treatment increased the resistance of the alloy to exfoliation attack.

Declaration

No portion of the work referred to in this thesis has been submitted in support of an application for another degree or qualification of this or any other university or any other institution of learning.

Bader M H M Elkandari (2012)

Copyright Statement i. The author of this thesis (including any appendices or schedules to this thesis) owns any copyright in it (the Copyright) and he has given The University of Manchester the right to use such Copyright for any administrative, promotional, educational and/or teaching purposes. ii. Copies of this thesis, either in full or in extracts, may be made only in accordance with the regulation of the John Rylands University Library of Manchester. Details of these regulations may be obtained from the Librarian. This page must form part of any such copies made. iii. The ownership of any patents, designs, trade marks and any and all others intellectual property right except for the Copyright (the “Intellectual Property Rights”) and any reproductions of copyright works, for example graphs and tables (“Reproduction”), which may be described in this thesis, may not be owned by the author and may be owned by third parties. Such intellectual Property Right and Reproduction cannot and must not be made available for use without the prior written permission of the owner(s) of the relevant Intellectual Property Right and/or Reproductions. iv. Further information on the conditions under which disclosure, publication and commercialisation of this thesis, the Copyright and any Intellectual Property and/or Reproductions described in it may take place is available in the University IP Policy in any relevant Thesis restriction declarations deposited in the University Library, The University Library‘s regulations (see http://www.manchester.ac.uk/library/aboutus/regulations) and in The University‘s policy on presentation of Theses.

Acknowledgments

First, I would like to gratefully acknowledge my supervisor, Professor Peter

Skeldon, for his guidance and patience throughout my graduate study at the

University of Manchester. I appreciate his time available and the time he gave me to learn from him. I would be impossible to finish this thesis without his inspiration, encouragement and advice. I would like to thank my second advisor,

Dr. Zhu Liu, for her valuable comments and suggestions, precisely on the laser field. I would like also to thank the staff and technicians of the school of materials, particularly, the technicians of the Corrosion and Protection Centre for their help, guidance and cooperation during the experimental work (namely Chris

Wilkins for SEM, Teruo Hashimoto for TEM, Judith for XRD, Shirley for SEM, and Gary for XRD). I would like to thank my friends who always encouraged me during my study until I finished my thesis (namely Zakria Aburas, Ramadan

Abuaisha, Mohsen Rakhes, Abdeslam Mhich, Ahmed Limnifi, Ahmed Shieki, and Wei Go). I would like to thank my father, mother and wife for their supportive and patience with all my problems. I appreciated their understanding and hospitality. Finally, I would like to take this opportunity to thank the Public

Authority for Applied Education and Training (PAAET) back in Kuwait for their financial support of my PhD study at the University of Manchester in the United

Kingdom.

Chapter 1 Introduction

1.1 Rational of the Research Project

7xxx series aluminium alloys have many applications, notably in aerospace construction, due to the excellent properties of the alloys, including high strength and low density. Unfortunately, these alloys may be highly susceptible to localised corrosion in the presence of aggressive environments containing chlorides. It has been reported that different alloying elements such as magnesium, zinc, copper and iron can have great influences on the corrosion properties of the alloy [1-5].

Aluminium alloy 7075-T6 contains a large number of intermetallic particles, which are either developed purposely to improve the mechanical properties or are present as impurities that can lead to localised corrosion, i.e. pitting and intergranular attack [6-12]. In order to reduce the localised corrosion of the alloy, a number of surface protection methods have been applied [13, 14]. However, none of these protecting methods have significantly reduced the corrosion problems without affecting the other properties of the alloy. For example, methods of surface protection using thermal and heat treatments can affect the composition of the alloy (bulk properties) and reduce its strength due to difficulties in controlling the temperature [15-17]; cladding increases the resistance to stress corrosion cracking, but causes a reduction in fatigue life and microhardness [18, 19]. Further chemical protecting techniques, such as chromic

Introduction acid anodising (CAA), can be highly toxic and potentially damaging to the environment.

Therefore, it is important to improve the corrosion resistance of the AA 7075-T6 alloy by eliminating or reducing the sizes of the large second phases on the surface of the alloy, without affecting its bulk properties. At the same time, a process is required that has no damaging effect on the environment.

1.2 Research Objectives

In this study laser surface melting (LSM) is employed, using an excimer laser, which allows modification of the alloy surface structure without reducing the bulk properties of the alloy [20-26]. The use of the excimer laser associated with a high cooling rate, hence leading to rapid solidification, can significantly reduce the number of intermetallic particles on the surface of the alloy. Such particles have a high cathodic activity compared with the alloy matrix. For example, the presence of large intermetallic particles rich in copper and iron can promote the formation of micro-galvanic couples with the adjacent matrix, leading to preferential dissolution of the alloy matrix [27]. On the other hand, the microstructure of the alloy surface is homogenised under the short pulse duration

(~ 13 ns) of UV radiation of the excimer laser that can dissolve the large second phase particles and redistribute the elements within the solidification layer in very short times (less than 1 μs). Further improvement in the corrosion resistance may be expected if LSM is combined with other protection methods, such as anodising. Thus the objectives of the present project are to optimise the laser processing conditions to generate uniform melted layers, free of large second

Introduction phases; to assess the corrosion resistance of the laser treated surface in comparison with the untreated alloy using several corrosion test methods; and to explore the possibility of further improving the corrosion resistance by anodising of the laser treated surface.

1.3 Thesis Layout

The first part of this thesis, after the introduction (Chapter 2), is a literature review of aluminium and aluminium alloy properties, with a focus on 7xxx aluminium alloys including heat treatments, age hardening, tempering and hydrogen effects. The fundamental understanding of localised corrosion of 7xxx alloys, including the effect of microstructure, potential, temper and environment will be discussed in Chapter 3. This part will further explain the general forms of corrosion attack and corrosion problems of 7xxx alloys, and the methods of corrosion prevention. Chapter 4 will consider LSM techniques and their use with aluminium alloys. Methods of heating, cooling and rapid solidification by LSM will be described and reviewed. Safety aspects of using a laser are briefly explained also. Chapter 5 is devoted to the experimental study, including the characterization of the alloy that is used in this project (AA 7075-T6). The surface preparation of the alloy for the laser treatment, as well as anodising details of the procedure, will be explained. Furthermore, the methods of corrosion evaluation, including potentiodynamic polarization, immersion and exfoliation (EXCO) tests will be described. The methods used in material characterisation before and after LSM, as well as before and after corrosion tests, including field energy gun scanning electron microscopy (FEG-SEM), transmission electron microscopy (TEM) and X-ray diffraction (XRD) will be

Introduction explained. Chapter 6 presents the results of the surface and microstructure examinations of the as-received (AR) AA 7075-T6 alloy before and after corrosion testing. Chapter 7 shows the results of optimisation of the laser operation conditions of the AA 7075-T6 alloy. Chapter 8 presents the results of the LSM AA 7075-T6 alloy in terms of surface and cross-section analyses before and after corrosion evaluation tests. The results are compared with those of the as-received alloy results to show the differences in corrosion behaviour of the laser-treated and untreated alloys. Chapter 9 provides the results for the anodising of the LSM alloy in 0.4 M sulphuric acid (H2SO4) and the evaluation of the corrosion behaviour of the LSM alloy with and without anodising. Chapter

10 will discuss the overall findings for the AR, LSM and the LSM anodised alloys. Finally, the conclusions and suggestions for future work are given in

Chapter 11.

Introduction to Aluminium Alloys

Chapter 2 Literature Review

2. Introduction to Aluminium Alloys

Aluminium is the second most used metal after steel. It is bright, shiny, and light.

It is in group 3 on the periodic table. Pure aluminium has a face centred cubic

(FCC) crystal structure. Aluminium has an excellent electric conductivity, high strength to weight ratio, with a specific weight of 2.7 g/cm³, which is about one- third that of steel and has high corrosion resistance [4]. It has many applications, especially in combination with other elements to form high strength alloys.

Aluminium alloys such as 7xxx alloys and 2xxx alloys have the ability to withstand high temperature and pressure, which has advanced the use of such materials in aerospace applications, where aluminium alloys play an important part in aeroplane structures. For example, AA 7075-T6 is used in aeroplane wing structures and AA 2024-T3 is used in the fuselage [4, 28].

Aluminium alloys are divided into wrought and cast products. Cast alloys are directly cast into their final form by one of various methods, such as sand- casting, die or pressure die casting. Wrought alloys are initially cast as ingots or billets and subsequently hot and/or cold worked mechanically into the final desired form. There are 9 series of wrought alloys. These are designated by a 4 digit number that may be preceded or followed by letters. The first digit indicates the alloy group number. The second digit indicates alloy modifications. The third and fourth digits indicate the aluminium purity. For example, 7xxx alloys are in group number 7 as zinc is the main alloying element with additions of minor

Introduction to Aluminium Alloys percentages of other elements, such as silicon, magnesium, copper...etc. Table 2-

1 provides details of the aluminium alloy classification system [29].

In order to specify the thermal processing in the alloys, a system of thermal temper nomenclature has been designed to distinguish between one alloy temper and another. The letter T and accompanying digits are added as suffixes to the alloy number to distinguish the temper. For example, T6 in aluminium alloy

7075-T6 is used to distinguish solution heat treatment and then artificially ageing, whereas T7 in aluminium alloy 7150-T7 is related to solution heat- treatment and artificially over-ageing. The mechanical properties of 7xxx alloys at different tempers are given in Table 2-2. Further discussion will be given of the stages of heat treatment of AA 7xxx alloys and temper effects later in this

Chapter.

Adding elements such as copper, zinc and magnesium to the aluminium to form a super-saturated solid solution (SSSS), by different heat treatment cycles, can improve the strength properties, but at the same time this will alter other properties, such as corrosion resistance. For example, alloys of the AA 7xxx series are very strongly heat treatable, due to precipitation hardening based on combination of zinc and magnesium, but the corrosion resistance is decreased.

Adding other elements, such as copper with a small amount of chromium, into aluminium-zinc-magnesium alloys can result in the highest strength aluminium alloys available [28]. However, copper decreases the general corrosion resistance. In contrast, the stress corrosion cracking (SCC) resistance increases

Introduction to Aluminium Alloys with a small addition of chromium. Table 2-3 presents more details on the solid solubility of the major elements and their effects on the alloy properties.

Different types of localised corrosion, such as pitting, intergranular (IGC) and exfoliation corrosion, and stress corrosion cracking (SCC), have been found in different areas of wing structures made of 7xxx aluminium alloys [28]. These forms of localised corrosion initiate at local sites and accelerate to cause a failure of the structure. For that reason, the microstructures of 7xxx aluminium alloys have been investigated by many researchers and much information has been reported that will be reviewed in detail in the next section.

2.1 7xxx Aluminium Alloys

7xxx series aluminium alloys, or more precisely Al-Zn-Mg-Cu alloys, have been investigated extensively over several decades. Rosenhain and his colleagues at the National Physical Laboratory in Britain in 1917 obtained a high tensile strength of around 580 MPa for an alloy composition of Al-20%Zn-2.5%Cu-

0.5%Mg-0.5%Mn. In comparison, the tensile strength of Duralumin (Germany) was around 420 MPa. For the next two decades, the alloy seemed to be unsuitable for structural use due to a high susceptibility to intergranular corrosion

(IGC) and stress-corrosion cracking (SCC). However, because of the major importance of Al-Zn-Mg-Cu alloys for aerospace construction, these corrosion problems have attracted continuing research [28].

AA 7075 in the T6 temper condition has an heterogeneous microstructure, consisting of an Al matrix, second phase particles and grain boundary regions.

Introduction to Aluminium Alloys

Mg(ZnCuAl)2 particles, with size about 50-100 nm, have been observed at the grain boundaries using transmission electron microscopy (TEM) [5, 30]. On either side of the grain boundary is a precipitate-free zone (PFZ) or solute- depleted zone (SDZ), about 30-70 nm wide and zinc and copper were usually found to be depleted in this region [10, 31].

Second phase particles in the alloy matrix usually fall into three categories, hardening precipitates, dispersoids, and constituent particles. Hardening precipitates are formed by nucleation and growth from supersaturated solid solution (SSSS); their sizes range from several to tens of nanometres [4]. For example, η and other precipitates forms, such as Mg2Al3, Al2Cu, and Al32Zn49, have been observed in the grain boundary regions [32]. Dispersoids are used to retard recrystallization and control the grain size; they contain aluminium and other elements such as chromium, titanium, manganese, and zirconium. Their sizes are bigger than precipitates, with sizes up to several hundred nanometres.

Typical dispersoids are Al3Ti, Al6Mn, Al20Cu2Mn3 and Al3Zr [4, 32]. Constituent particles are relatively large and can be seen easily by optical microscopy or scanning electron microscopy (SEM). These particles form during solidification and survive subsequent heat treatment. Their sizes range from a few tenths of a micron up to tens of microns, and these particles often break into smaller particles during hot or cold working and align along the working direction. Such particles in AA 7075 alloys consist of Al3Fe, Al2Cu, Al7Cu2Fe, and Mg2Si [4, 32,

33].

Introduction to Aluminium Alloys

The alloys have been studied over a wide range of heat treatment processes. AA

7075 with T73 temper shows a fine dispersion of η precipitates through nucleation from pre-existing GP zones, with improvement in SCC; however, the tensile properties are about 15% below those AA 7075 with the T6 temper [34].

Another duplex ageing treatment, T76, has been applied successfully to 7xxx alloys to increase the resistance to exfoliation corrosion, but again the alloy strength is decreased [28]. Other researchers have successfully achieved a combination of high resistance to SCC with maximum levels of tensile properties by adding 0.25-0.4% silver to Al-Zn-Mg-Cu alloys, enabling a high response to age-hardening to be achieved in single ageing treatment at 160-170˚C [28, 50].

More recently, AA 7050 (Al-6.2%Zn-2.25%Mg-2.3%Cu-0.12%Zr) was developed in the United States in which the level of copper present in the alloy was raised in order to increase the strength with a second stage of T73 treatment.

The presence of large intermetallic compounds can be minimised by increasing the limit of copper content, but with adverse effect on toughness and tensile properties. Modelling indicates that the sum of Cu + Mg wt% should be < 3.6% for an alloy containing around 6% Zn. Another heat treatment, known as retrogression and re-ageing (RRA), has been developed that enables alloys such as AA 7075 to exhibit a high level of tensile properties expected of the T6 condition combined with a resistance to SCC equal to that of the T73 condition

[35, 51]. The process has been given a temper designation T77 and can be applied to highly alloyed compositions such as AA 7055 (Al-8%Zn-2.05%Mg-

2.3%Cu-0.16%Zr) that has been used for structural members of the Boeing 777 wings [35].

Introduction to Aluminium Alloys

As mentioned above, heat treatment processing of 7xxx aluminium alloys can result in improved resistance to SCC but with adverse influence on toughness and strength. At the same time, difficulties in controlling the temperatures and the level of deformation can cause a problem in the material decomposition during age-hardening processes [36].

Compositional changes such as a reduction in the levels of iron and silicon impurities in alloy AA 7075 (0.9%), with a maximum combined amount of

0.22% in the higher purity alloy AA 7475, can improve the alloy toughness. [37]

Furthermore, in AA 7475 the manganese content is reduced, from 0.3% to 0.06%

(maximum), whereas the content of other alloying elements is essentially the same as in AA 7075. The size and number of intermetallic particles, that assist crack propagation, is reduced and similar tensile properties are obtained as with the T735 heat treatment [35, 36]

Another modification has utilized 0.08-0.25% zirconium as a recrystallization inhibitor, either in place of chromium or manganese or in combination with a smaller amount of these elements, in order to reduce quench sensitivity so that slower quench rates can be used without loss of strength on subsequent ageing.

Examples are the alloys 7050, 7150, 7055, and the British alloy 7010 [37, 38, 39]

Modifications to alloying additions, purity, and heat treatment have allowed high values of fracture toughness to be sustained in AA 7150-T651 and AA 7055-

T7751. However, the corrosion behaviour has not shown a significant

Introduction to Aluminium Alloys improvement. One compositional trend has been to raise the Zn:Mg ratio from

2.25 for the early AA 7075 to 5.0 for the latest AA 7085. Another trend has been to reduce the allowable limits on Mn and Cr, as well as the Fe and Si impurities, each to below 0.1%. For example, AA 7085 alloy, with Fe < 0.08% and Si <

0.06%, has a low quench sensitivity, which enables very thick sections (150 mm) to be heat treated, with tensile properties higher than early alloys such as AA

7050. This has been a major factor for selection of AA 7085 for large extruded wing spars and die-forged wing rib components for the European A380 Airbus

[37, 38, 39].

Reference has already been made to the fact that 7xxx alloys tend to show a high strength and that depends on the elements and impurities in alloy composition and the method of heat treatment [38, 39]. However, none of these approaches to alloy development has achieved a significant improvement of the corrosion resistance without affecting the other alloy properties.

2.2 Elements and Impurities in 7xxx Aluminium Alloys

Alloying elements are added to 7xxx alloys mainly for precipitation hardening

[4]. Zinc and magnesium offer the highest potential for strengthening through age-hardening, but also lead to a high susceptibility to stress corrosion cracking

(SCC) [4, 40]. Copper is added to these alloys to improve the age-hardening, as well as to increase the resistance to SCC [2]. The process of heat treatment and age-hardening will be described in detail in the next section of this Chapter.

Introduction to Aluminium Alloys

Zinc, magnesium, and copper are important alloying elements in AA 7xxx alloys.

Each of these elements results in different corrosion properties of Al when present in solid solution. Muller and Galvele studied binary Al-Zn alloys, with

Zn content ranging from 1.07 to 4.85% [7]. They measured the breakdown potential of these alloys with polarization in chloride solution and they found that the shapes of the polarization curves were similar to those found with pure Al.

However, the breakdown potential decreased when Zn was present [7]. Other studies observed similar results [8, 9, 41]. Muller and Galvele also measured the breakdown potential of Al-Mg alloys in chloride solution [7]. Their results showed that the breakdown potential is independent of Mg content and is close to that pure Al. They found only aluminium inside the pit. Therefore, they suggested that Mg would be dissolved preferentially from Al-Mg alloys. The corrosion potential of Al-Cu alloys in chloride environment was found to increase with increase of the Cu content [6] It is reported that addition of Cu to

Al has the opposite effect on the corrosion behaviour of adding Zn to Al [6,7].

Ramgopal and Frankel measured the repassivation potential of binary solid solution Al-Cu alloys and observed that even with an amount as small as 0.022 wt% Cu can raise the repassivation potential by 30 mV relative to that of pure Al

[9]. Adding Cu was found to increase the dissolution kinetics because the potential required to achieve a given current density increased with Cu content.

In addition to change in the electrochemical property, Cu also plays an important role in the microstructure as it can alter the kinetics of precipitation reactions

[11]. Moreover, Muller and Galvele observed Cu enrichment inside the pit in a

Al-Cu binary alloy [6]. This led to the suggestion that Al is preferentially

Introduction to Aluminium Alloys dissolved and Cu stays in the alloy which increases the Cu concentration locally and affects the corrosion mechanism.

From the above studies in chloride environment, Zn lowers the breakdown potential of Al, and Cu increases the breakdown potential of Al; however, Mg has little effect on the breakdown potential of Al. It is reported that Mg is dissolved preferentially from Al-Mg alloys and Al is dissolved preferentially from Al-Zn and Al-Cu alloys [6, 7, 42]. Since Zn and Cu have opposite effects on Al dissolution, it is possible that these two alloying elements have different corrosion products at AA 7xxx alloys. However, the relationship is still not clear.

Second phase particles usually have different electrochemical characteristics and different effects on the corrosion behaviour of the surrounding matrix [12]. They can be cathodic or anodic to the alloy matrix. The presence of these phases can lead to non-uniform attack at specific areas of the alloy surface and are often associated with corrosion problems [4, 45]. The corrosion potential of MgZn2 (η) phase in 1 M NaCl was measured between -1035 and -1100 mV (SCE), which was more negative than the corrosion potential of Al (-700 mV) [43, 44].

Therefore, MgZn2 phase is active with respect to the alloy matrix. Ramgopal and his colleagues studied the η phase as a function of Cu, Zn and Al concentration by using a flash evaporation technique to generate thin film compositional analogs of the intermetallics [9]. For MgZn2 with no Cu content, they observed a sharp current increase at a potential which was close to the reversible potential of

Zn/Zn2+. They suggested that the breakdown potential of the η phase is associated with the dissolution of Zn. By adding Cu to more than 17 at %, the

Introduction to Aluminium Alloys corrosion potential increases as well as the breakdown potential. They did not observe any change in the breakdown potential with adding to 10 at % of Cu [9].

Iron and silicon are reported to be the main impurities in commercial grade aluminium alloys [45, 46]. The solubilities of iron and silicon are high in molten aluminium, but are low in solid aluminium [46]. Therefore, most of iron and silicon will precipitate out during solidification and form large second phase particles. The iron rich particles, such as Al3Fe and Al7Cu2Fe, usually act as cathodes on the surface of the alloy, promoting electrochemical attack. Murray et al. reported that the presence of Cu ions in the solution would cause metallic copper to plate onto the iron-rich particles [47]. If the matrix adjacent to the iron- rich particles contains copper, aluminium will be dissolved and the copper will remain inside pits and plate onto iron-rich particles, which makes the particles more cathodic with respect to aluminium [47].

From the above investigations, second phase particles can be noble and /or active to the adjacent matrix. Cu and Fe-rich particles can be more noble than the matrix and act as cathodes. Mg and Zn- rich particles are more active than the matrix and are usually dissolved preferentially. Adding aluminium to the MgZn2 phase has little effect on corrosion and breakdown potentials, while adding copper to the MgZn2 phase will increase its corrosion and breakdown potential.

Some elements will leave the pit after being dissolved, such as magnesium, but some will stay, such as zinc and silicon. The change in concentration inside the pit will affect the subsequent corrosion. Overall, the presence of different second phase particles in AA 7xxx alloys affects their corrosion behaviour strongly.

Introduction to Aluminium Alloys

2.3 Stages of Heat Treatment of 7xxx Aluminium Alloys

A key purpose of heat treatment is to increase the strength and hardness

(mechanical properties) of the aluminium alloys in order to achieve a final product that can be employed in different environmental applications. For 7xxx series aluminium alloys, the heat treatment process usually comprises three main stages, as follows:

1- Solution heat treatment (SHT) at high temperature within a single

phase region to dissolve the alloying elements in the α-Al matrix.

2- Quenching (Q) in order to form a supersaturated solid solution (SSSS)

of the dissolved elements in the α-Al matrix.

3- Ageing (A*) at temperature below the solvus in order to control the

decomposition of the supersaturated solid solution to form a fine

dispersion of precipitates (GP zones) → (η) → (η’)

The temperature-time diagram of Figure 2-1 shows the three main stages of the heat treatment process, with the precipitation scheme, for the AA 7xxx alloys.

Solution heat treatment involves heating the aluminium alloy above the solvus temperature so that all the constituents are taken into the solid-solution (α-Al).

During quenching, if the alloy were slowly cooled, β-phase would nucleate and grow at heterogeneous sites to form an equilibrium α + β structure. A faster quenching rate reduces the time for diffusion and freezes the alloying elements in the non-equilibrium α-phase structure, forming a supersaturated solid solution

(SSSS). Upon ageing, the SSSS will tend to transform toward an equilibrium

Introduction to Aluminium Alloys structure and form a fine dispersion of Guinier Preston zones (GP zones). In AA

7xxx alloys, the GP zones are reported usually, to have spherical shapes, while η

(MgZn2) is hexagonal [48]. Commonly used heat treatments in AA 7xxx series alloys are T6 and T7 tempers. The T6 temper provides a peak aged condition, which has the maximum strength [4]. It is achieved by SHT with about 465 °C and Q, followed by a single stage heat treatment at temperatures in the range

120-135 °C for 24 h [49]. The tensile strength can reach around 550-610 MPa, depending on the alloy composition [4].

Unfortunately, AA 7xxx alloys in the T6 temper have a relatively low resistance to SCC. To improve the SCC resistance, a duplex heat treatment or over-aging

T7 temper was developed. It involves two steps of artificial aging treatment. The first step is at 107-120 ˚C for a few hours to nucleate GP zones [49]. The second step is at 160-170 ˚C and produces a refined dispersion of particles (η’) from the pre-existing GP zones [49]. The refined dispersion of particles reduces the strength by about 15% from that in T6 temper, but was reported to increase the resistance to SCC [4, 52].

2.3.1 Age Hardening

In age hardening, the complete decomposition of the SSSS usually occurs by a complex process, which may involve several stages, initially involving formation of Guinier-Preston (GP) zones, a solute rich groups of atoms. They retain the structure of the alloy and are coherent with the matrix. Diffusion associated with their formation involves the movement of atoms over short distances, assisted by vacant lattice sites that are present due to quenching [28]. Depending on the alloy

Introduction to Aluminium Alloys composition, the rate of nucleation of the GP zones may be influenced by the presence of the excess vacant lattice sites. On the other hand, the intermediate precipitates are much larger in size than the GP zones and are only partly coherent with the lattice planes of the matrix. When a critical dispersion of the

GP zones or intermediate precipitates, or a combination of both, occurs, maximum hardening in commercial alloys can be achieved. For increased stability of the alloys, the GP zones solvus lies between the quench-bath temperature and the ageing temperature. This gives an advantage of nucleation of intermediate precipitates from pre-existing GP zones (duplex ageing treatment) that can improve the SCC resistance in high strength aluminium alloys [28].

2.3.2 Precipitate Free Zones (PFZ)

Precipitate free zones (PFZ) can be formed in alloys during ageing. A narrow region (~50 nm) on either side of the grain boundary is depleted of solute due to diffusion of solute atoms to the boundary where relatively large precipitates are formed [10, 50]. The PFZ arises because of a depletion of vacancies to levels below that needed to assist with nucleation of precipitates at the particular ageing temperature. Higher solution heat treatment temperature and faster quenching rate to a lower aging temperature, usually below the solvus, increase the excess vacancy content and narrow the PFZs. At the same time, the concentration of the solute becomes higher, that means that smaller nuclei are stable. This will reduce the critical vacancy concentration required for nucleation to occur. Therefore, the vacancy depleted part of PFZ cannot be formed in alloys aged at temperatures below the GP zones solvus, as the GP zones can form homogeneously with no need of vacancies. Viana et al. observed the microstructure of AA 7075-T6 by

Introduction to Aluminium Alloys

TEM and found very fine precipitates distributed homogeneously inside the grains and coarser continuous precipitates distributed at the grain boundaries

[51]. Puiggali [52] and Ramgopal [10] found precipitates in AA 7010-T6 and

AA7150-T6 to be of size up to 50 nm in the matrix and from 50 to 120 nm at the grain boundaries. The width of the PFZ is about 30 nm [10, 52]. The precipitates at the grain boundaries for the T7 temper are coarser and more spaced than with the T6 temper, while inside the grains, the precipitates are rather coarse and uniformly distributed [51]. It is suggested that the precipitates in the T7 temper are essentially η. The elongated η precipitates in AA 7150-T7 were reported to be of sizes up to 70 nm and 250 nm in the matrix and the grain boundaries respectively [52]. The width of PFZ at the grain boundaries was measured to be between 40 and 50 nm [52]. Overall, the PFZ may be influenced by temper reactions; as has been discussed, the T6 temper maximises the tensile strength of

AA 7xxx alloys, while the T7 temper improves the SCC resistance. Fine precipitates inside the grains of alloys in the T6 temper are responsible for a high mechanical property, while coarser and more spaced grain boundary precipitates in the T7 temper alloys seem to enhance the SCC resistance.

2.3.3 Hydrogen Effects

Hydrogen embrittlement is one of the main factors that can cause SCC.

Generally, hydrogen acts to weaken interatomic bounds in the plane strain region at the crack tip [53]. In aluminium alloys, localised corrosion is generally accompanied by hydrogen evolution. The pH is reduced inside pits, owing to cation hydrolysis as well as drops in the ohmic potential. Therefore, the potential inside pits or with other forms of localised corrosion is below the reversible

Introduction to Aluminium Alloys potential for hydrogen evolution. Furthermore, the hydrogen evolution reaction is

a cathodic reaction ( 2H 2e H 2 ) that consumes a portion of the electrons generated by the localised corrosion. This portion of electrons was reported to be about 20% and even higher in some cases [54, 55]. As a result, hydrogen evolution inside pits occurs at a high rate during localised dissolution of the alloy. Moreover, the presence of hydrogen atoms either at the crack tip or after transport into the alloy ahead of the crack tip can result in loss of ductility of the alloy and brittle cracking [44, 53]. Several authors have demonstrated hydrogen embrittlement in aluminium alloys [44, 45, 56]. The occurrence of intergranular fracture and a small reduction area in a tensile test usually indicates the loss of ductility [45]. Albrecht et al. performed a series of tensile tests on equiaxed AA 7075 alloy and found a significant embrittling effect, as manifested by a decreased reduction area, when precharging with hydrogen [56]. They interpreted the embrittlement as a result of the presence of cathodic-charged internal hydrogen and further straining enhanced the transportation of additional hydrogen by the movement of dislocations. Therefore, a large amount of hydrogen can be transported by dislocations deeply into the interior of alloys.

This provides a more rapid transport mode than the lattice transport mode of hydrogen in aluminium alloys, which can explain the existence of hydrogen embrittlement.

In aluminium alloys, the grain boundaries are the preferred locations of hydrogen accumulation [56, 57]. It was reported that Mg segregation promotes hydrogen entry into the grain boundaries in aluminium [58]. Pickens and Langan studied

Al-Zn-Mg alloys and observed the presence of segregated Mg at the grain

Introduction to Aluminium Alloys boundaries [59]. Pickens et al. proposed that the segregated Mg can promote hydrogen to concentrate at the grain boundaries in sufficient quantities to cause embrittlement [60].

From the above discussion, hydrogen can be transported by dislocations and accumulate at the grain boundaries, causing embrittlement in aluminium alloys that enhances localised corrosion attack. However, mechanisms of hydrogen effects in SCC and IGC in AA 7xxx alloys are still being investigated.

2.3.4 Temper Effects

As mentioned earlier the decomposition of the SSSS of the alloy can be achieved with different tempers during hard ageing. The most common tempers of the AA

7xxx alloys are T6 and T7. T6 provides the highest strength for AA 7xxx alloys, but results in susceptibility to localised corrosion [4]. T7 is an overaging temper that provides improved SCC resistance, but reduced strength [52, 59, 61-66].

Najjar used slow strain rate experiments in 3% NaCl solution with AA 7075 alloys in T6 and T7 tempers [67]. The results showed that AA 7075-T6 was highly sensitive to SCC, while overaged AA 7075-T7 had a lower IGC and improved SCC resistance. He observed the possibility that overaging induces a reduction of the following factors.

A- Localised anodic dissolution at the grain boundary regions

B- Hydrogen diffusion into the grain boundaries.

C- Local stress concentrations through the homogeneity of deformation at

the grain boundaries.

Introduction to Aluminium Alloys

Therefore, overaged grain boundary regions are less sensitive to SCC and the fracture has to propagate by discontinuous, cleavage-like cracks [67].

Coarsening of the grain boundary precipitates is another factor that has been reported to be responsible for the improvement in the resistance to SCC [52, 66,

68]. Rajan and his colleagues studied the grain boundary precipitates of AA 7075 in T6, T73, and RRA (retrogression re-aging) tempers and performed SCC testing [66]. The results showed that the heat treatments that produced more SCC resistance have larger grain boundary precipitate sizes than the T6 condition.

Tsai reported similar observations in AA 7475 alloys between T6 and RRA tempers [63, 69]. It seemed that the RRA temper can produce larger and more spaced grain boundary precipitates than the T6 temper [63, 66, 69]. The wide spaced grain boundary precipitates can delay the localised anodic dissolution and can act as trap sites for hydrogen, where hydrogen bubbles can form. Therefore, the hydrogen concentration at grain boundaries can be reduced to below the critical value for inhibiting of hydrogen embrittlement and improved SCC resistance. As mentioned earlier, dislocations can increase hydrogen transportation to grain boundary precipitates during quenching following solution treatment and that reduces the SCC resistance of the alloys. Talianker and Cina observed a high density of dislocations adjacent to grain boundaries and in the bulk of the grains of several 7xxx alloys in the T6 temper [70]. However, those dislocations disappeared after an RRA temper [70]. They suggested that the short time and high temperature of retrogression could annihilate dislocations without causing overaging. The disappearance of dislocations results in retarded transport of hydrogen into the alloy and improved SCC resistance. They believed that a

Introduction to Aluminium Alloys major change in dislocation density, rather than a minor change in the precipitate structure, could increase the SCC resistance of 7075 aluminium alloys with an

RRA temper.

Potential measurements are a useful tool to evaluate the effects of solution heat treatment and the subsequent aging process. Anodic polarization curves for AA

7xxx-T6 in chloride solution typically exhibit two breakdown potentials, whereas the T7 exhibits only one [10, 42, 71, 72]. Maitra observed only IGC between the two breakdown potentials and both pitting and IGC above the second breakdown potential in AA 7075-T6 [71]. This finding is the same as that of Galvele and

Micheli in Al-Cu alloys. The alloy in the T7 temper was immune to IGC and pitting only occurred above the breakdown potential [73]. Therefore, Maitra suggested that the disappearance of the second breakdown potential in T7 alloy resulted in an improvement of the IGC resistance. However, Ramgopal and his colleagues observed different results from Maitra. They found only pitting occurred in AA 7075-T6 between the two breakdown potentials [42]. They further investigated AA 7150 in T6 and T7 tempers. Similar to AA 7075 alloy, two breakdown potentials were shown in the T6 temper and only one breakdown potential in the T7 temper [10]. This time they performed a short time (1 h) and a long time (24 h) potentiostatic tests separately on their samples. In short time tests, the results were the same as those in the AA 7075: only pitting in T6 samples between the two breakdown potentials and both pitting and IGC appeared above the more noble breakdown potential while T7 samples were immune. However, in long time tests (24 h), both pitting and IGC were shown on all their specimens. As a result, they assumed that IGC of AA 7150 is a time-

Introduction to Aluminium Alloys dependent phenomenon [10]. Meng and his colleagues made other observations of the corrosion process in relation to the breakdown potentials in AA 7xxx-T6

[72]. Between the two breakdown potentials, dark uniform corroded regions were observed on the sample surface. These regions were examined by SEM and no large or deep pits were seen on the surface. Above the second breakdown potential, selective grain attack and IGC were evident. They concluded that the first breakdown potential corresponds to transient dissolution, while stable dissolution occurs above the second breakdown potential. Issacs ground

AA7075-T6 to 600 grit and exposed the sample at the OCP (open circuit potential) for 1 h before polarization. He used image subtraction and found streaking attack along the scratched lines on the surface. The streaking attack was not observed in the T7 temper. He suggested that streaking attack was caused by mechanical grinding and resulted in two breakdown potentials [74].

However, he did not explain why AA 7xxx alloys in the T7 temper are immune to streaking attack, since the sample preparation by grinding was the same as for the AA 7xxx-T6 temper. Ramgopal examined the grain boundary precipitates of

AA 7150 alloy in T6 and T7 tempers. He found that the copper concentration in the grain boundary precipitates was higher in the T7 temper than that in the T6 temper [10, 42]. He presumed that the presence of a high concentration of copper ions in solution due to dissolution of T7 grain boundary precipitates may be responsible for preventing dissolution of the PFZs near the grain boundary region and making the T7 temper less susceptible to IGC. As a result, he suggested that the breakdown potential in the T7 temper is due to pitting in the matrix.

However, a high copper environment in the PFZs is needed to dissolve the grain boundary precipitates first. If no dissolution of grain boundary precipitates

Introduction to Aluminium Alloys occurs, the PFZs are still more susceptible than the matrix. Since the T7 temper has a higher IGC resistance, it could be suggested that the breakdown potential is due to dissolution of precipitates in this type of temper.

From the above discussions, the corrosion resistance for an over-aging (T7) temper is higher than with a peak-aging (T6) temper for AA 7xxx alloys. It is an evident that the AA 7xxx-T6 alloy exhibits two breakdown potentials after examined in 0.1 M NaCl deaerated solution using potentiodynamic polarization test as well discuss later in details in corrosion test results of the as-received alloy section on chapter eight of this report. The first breakdown potential is associated with pitting in the matrix, pitting in the solute enriched region of the grain boundary, or streaking attack created by mechanical grinding. Some authors proposed that coarsening of precipitates and more space between precipitates at the grain boundary in AA 7xxx alloy with T7 temper can improve the corrosion resistance, although others proposed that this may occur because of annihilation of dislocations and retarding of hydrogen transportation into the alloys. It seems that heat treatment cycles can decrease the corrosion susceptibility to SCC in

7xxx aluminium alloys.

Introduction to Aluminium Alloys

2.4 List of Figures

Table 2-1: Four-digit designation system for Aluminium Alloys [29].

Alloy Type Four digit designation Wrought alloys Casting alloys 99% Aluminum 1XXX 1XX.X Copper 2XXX 2XX.X Manganese 3XXX 3XX.X Silicon 4XXX 4XX.X Magnesium 5XXX 5XX.X Magnesium and silicon 6XXX Zinc 7XXX 7XX.X Lithium 8XXX Others 9XXX Note: Designations are based on aluminium content or main alloying elements [29].

Introduction to Aluminium Alloys

Table 2-2: Mechanical properties of AA7xxx alloys with different temper processes [28].

Thickness Tensile St. Yield St. Elongation Condition Temper (mm) (MPa) (MPa) (%) Definition

O 0.38-50.8 276 145 9-10 Fully Annealed SHT at 465˚C, T6 0.20-6.32 510-538 434-476 5-8 quenched then artificially aged SHT, quenched, stress relieved by stretching at T651 6.35-101.6 462-538 372-462 3-9 0.5–3%, permanent then naturally aged SHT, quenched, T76 3.18-6.32 503 427 8 two steps of artificially aged SHT, quenched, stabilised, 3% T73 1.02-6.32 462 386 8 cold worked, two steps of artificially aged SHT, quenched, cold worked, stabilised, stress relieved by T7351 6.35-101.6 421-476 331-393 6-7 stretching 0.5- 3% permanent set and two steps artificially aged

Introduction to Aluminium Alloys

Table 2-3: Solid solubility and influence of elements in Al properties [28].

Solid Element Symbol Solubility Effect on Aluminium properties in Al Increases strengthening, permits precipitation hardening, reduces Copper Cu Low corrosion resistance, ductility and weldability Increases strength through solid solution Magnesium Mg High strengthening, improves work hardening ability. Cause intergranular corrosion Increases strength, improves work Manganese Mn Low hardening ability Increases strength and ductility, in combination with Mg produce Silicon Si Low precipitates hardening, cast-ability and ductility. Increases strength, permits precipitation Zinc Zn High hardening, cause stress corrosion Chromium Cr Low Increases stress corrosion resistance Improves the strength at elevated Nickel Ni Low temperature Added as a grain-refining element Titanium Ti Low particularly in filler metals Zirconium Zr High Inhibiting re-crystallization

Introduction to Aluminium Alloys

SSSS (GP zones) → (η) → (η’)

Figure 2-1: Three main stages of heat treatment (SHT refers to solution heat treatment, Q is quenching, and A* is aging) with precipitation scheme for the 7xxx aluminium alloys

Chapter 3 Corrosion

3. Introduction to Corrosion

All common metals used in engineering are polycrystalline, being composed of many crystallites of varying size and orientation. These crystallites meet together to form grains and grain boundaries in the solid material. Corrosion attack according to the ASM-Materials Engineering Dictionary [78] is a chemical or electrochemical reaction between a material, usually a metal and its environment, that produces dissolution on the metal grains and their boundaries which leads to a deterioration of the material properties. Aluminium and its alloys are usually protected by a thin film of oxide, aluminium oxide (Al2O3) that acts as a barrier to protect the surface in various environments. In acidic and/or alkaline environments, this thin film may undergo dissolution, with consequent corrosion of the substrate. In general corrosion, oxidation and reduction reactions take place on the metal surface. The oxidation reaction results in the generation of the electrons at the anodic areas of the corroding metal, while the reduction reaction consumes these electrons at the cathodic areas of the same metal. The section on electrochemical polarization in this chapter is concerned with the reactions in detail. Different corrosion forms can be initiated on aluminium alloys of the 7xxx series. The type of corrosion depends to a large extent on the alloy composition and the chemical nature of the environment, as will be reviewed in this Chapter.

Corrosion

3.1 Surface Oxide Film

Aluminium and its alloys are relatively stable in most natural environments due to the rapid formation of the natural oxide film of alumina on their surfaces that acts as barrier to environmental attack. Figure 3-1 depicts the thin oxide film on the aluminium surface [75]. The oxide film is stable in aqueous solutions of pH in the range 4.5-8.5, while it is soluble in strong acid or alkaline environments

[77, 78]. The film thickness is typically several nanometres (5-15 nm). The rate of film growth increases at higher temperatures and higher relative humidities.

Much thicker surface oxide films can be produced by various chemical and electrochemical treatments. For example, the natural film can be thickened 500 times by immersion of the aluminium in certain hot acid or alkaline solutions. In highly aggressive environments containing chloride ions, the oxide layer may be dissolved, once the chloride ions have replaced oxygen anions, and the film loses its protective ability [79, 80]. In order to understand the chemical reactions between the environment and the surface of the alloy, electrochemical polarization is reviewed in the next section.

3.2 Electrochemical Polarization

Electrochemical polarization measurements employ an anode, cathode, and an electrical conductive solution between them. A DC current is applied in the electrolyte solution making electrons transfer from the anode site to the cathode site. The hydrated ions flow in the opposite direction in the electrolyte, from the anode through the electrolyte to the cathode. In this case, the anode will become

Corrosion positive due to loss of electrons and the cathode will become negative due to gain of electrons. The oxidation reaction on the anode area can be represented by the following reaction;

M M x xe (1) where M represents the metal or alloy, M x is the metal ion, and xe is the number of electrons involved in the oxidation process. On the cathodic area, a reduction reaction takes place, and which may be represented by the following reaction;

M x xe M (2)

The anodic and cathodic reactions must maintain charge neutrality [81].

M M x xe (3)

The potential difference between the anodic and cathodic reactions on the metal surface is the driving force for the corrosion reaction and therefore; current flows through the corrosion cell. For example, corrosion of aluminium in acid solution involves the transfer of electrons from the anodic site into the cathodic site. The oxidation reaction is Al Al 3 3e and the reduction reaction in acid solution is the reduction of hydrogen ions. This reaction can be written as follows,

2H 2e H 2 . The current on the other hand is a measure of the number of electrons flowing per second, and the greater the current the more electrons transfer between the anode and the cathode. Similarly, more hydrogen ions react to form uncharged hydrogen atoms. Therefore, the potential and current are of the primary interest in the study of corrosion. Potential provides a value of the thermodynamic driving force of the corrosion, and the current provides a measure of the electrochemical reaction rates that take place during the corrosion.

Corrosion

3.2.1 Thermodynamics

For corrosion to be spontaneous, the change in the Gibbs free energy ∆G must be negative. The free energy for the individual oxidation and reduction reactions is related to a reversible electromotive force or cell potential (Ec) through the equation

G xeFEc where xe is the number of electrons, and F is faraday’s constant. Therefore, E is directly related to the driving force (change in the Gibbs free energy) for the reaction. For example, for a chemical reaction

aA+bB = cC+dD where A and B are reactants and C and D are products, the equilibrium constant for the reaction is given by;

and;

G G RT lnK

Then;

 xeFEc xeFE RT lnK

By rewriting this equation;

Corrosion

This equation is called the Nernst equation, that relates the cell potential Ec to the standard potential E˚ and activity/concentration of the electro-active ions [82].

Furthermore, the Nernst equation is a powerful tool in constructing Pourbaix

Diagrams (Potential-pH) that predict the regions of thermodynamic stabilities of metal species at different combinations based on measurement of the potential and pH values, such that the corrosion behaviour of a metal in an aqueous solution can be predicted. However, Pourbaix Diagrams provide no guidance on the rate of the reactions within a region where corrosion is possible [81]. In general, the metal can exist in one of the three states in an aqueous solution.

These correspond to regions of corrosivity, immunity, and passivity in the relevant Pourbaix Diagrams as shown in Figure 3-2. In corrosive regions, corrosion occurs by general dissolution of the metal, while in the immune region, the potential of the metal is so far depressed that the oxidation reaction is not thermodynamically possible. In the passive region, the potential of the metal is elevated and the metal becomes covered with a protective film, isolating the metal from its environment. The metal is possibly resistant to corrosion in areas where the film is stable, because this thin layer, usually based on oxide, leads to a reduction in the rate of anodic dissolution on the metal surface.

Corrosion

3.2.2 Kinetics

Kinetics is the study of the rate at which reactions will occur. The potentials of the oxidation and reduction reactions on a metal surface are polarized from their equilibrium values. For example, when corrosion occurs and current flows between anodic and cathodic areas, the potential of the anode increases to more positive values (noble), and the potential of the cathode decreases to more negative values (active). Therefore, both anodic and cathodic potentials will approach a common value called the corrosion potential (Ecorr). A better way of visualising the relationship between potential and current is by means of an

Evans Diagram, where potential is plotted on the vertical Y axis and log current density is plotted on the horizontal X axis (Fig.3-4) [83].

In equilibrium, the corrosion current density (icorr) at the corrosion potential

(Ecorr) is directly proportional to the corrosion rate (CR). This current can be converted to an actual corrosion rate (CR) using Faraday’s law [84].

iM CR ZFD where i is the current density, M is the atomic weight of the metal, F is Faraday’s constant (96,500 C/g equivalent), and D is the density of the metal. By returning to the Evans Diagram, the slopes of the anodic and cathodic lines are referred to as Tafel slopes and have units of mV or V per decade of current. In general, the corrosion rate (CR) is determined by extrapolating the anodic and cathodic Tafel regions to the corrosion potential (Ecorr). At the corrosion potential, the extrapolated anodic and cathodic Tafel regions of the respective curves intersect.

Corrosion

At this intersection, the rate of the cathodic reaction equals the rate of the anodic reaction and the value of the current is the corrosion current density (icorr).

This is the Butler-Volmer equation, where I is the applied current density, icorr is the corrosion current density, is the over-voltage polarization (E – Ecorr),

and a and c are the anodic and cathodic Tafel slopes that are represented in

Figure 3-3 [85].

3.2.3 Passivity and Breakdown of Passivity

Since both driving force and kinetic considerations are crucial in determining the extent of corrosion on a metal surface, the phenomena of passivity and its breakdown are crucial factors in corrosion initiation, such as by pitting for certain metals and alloys [86]. Because of the presence of the protective film that naturally forms on aluminium and aluminium alloys in the passive region during an electrochemical procedure, the breakdown of passive region can vary during a polarization scan. The most common way to observe the breakdown potential

(Ep) in a wide range of metals and alloys is by potentiodynamic polarization. In this process, Ep is the potential at which the current dramatically increases after a period of passivity, marking the point above which pits initiate and grow [53].

An effective film is one that resists the breaching (breakdown) of the passivity.

Corrosion

Breakdown of the passivity leads to a localized corrosion that can lead to corrosion failures by pitting, crevice, and intergranular corrosion, and stress corrosion cracking. Figure 3-4 shows Ep, Ecorr, and the passive region in a potentiodynamic polarization curve. All the breakdown mechanisms involve a damaging species. For example, one of the major species causing breakdown of passivity and lowering of Ep is the concentration of chloride ions. Competing with passive film breakdown is the passive film repassivation. Thus, an effective alloy for resisting localized corrosion would be one that its surface not only forms a passive film but also is capable of repassivating at a rate sufficiently high. Once breakdown has occurred, the alloy surface is exposed to the corrosive environment [83]. However, Ep is often a function of experimental parameters, such as potential scan rate, and may consequently not be a real measurement for assessing the exact pitting potential [87, 88]. Therefore, it is important to undertake further corrosion tests to clarify the result of the investigations.

3.3 Types of Corrosion

Corrosion can take different forms depending on the alloy composition and its environments. In this part of the report, different types of corrosion will be reviewed and the circumstances in which these areas are indicated. In general for

7xxx alloys, localized forms of corrosion, such as pitting, intergranular (IGC), and stress corrosion cracking (SCC), are the most types of corrosion of the alloy surfaces.

Corrosion

3.3.1 Uniform Corrosion

This type of attack means that the metal lost to form corrosion products is about the same over the entire surface of the metal. It is the simplest form of corrosion, occurring in the atmosphere, in solution, and in soil under normal service conditions. Passive materials such as aluminium alloys are generally subjected to localized corrosion. Uniform corrosion commonly occurs on the metal surface and starts in the oxide layer before continuing to attack the underlying material.

Only in very acidic solution or very alkaline solution (4 > pH > 8), is aluminium homogenously corroded [89].

3.3.2 Localized Corrosion

Localized corrosion can be defined as the selective removal of metal by corrosion at small areas on the metal surface in contact with a corrosive environment, usually liquid [90]. It takes place when small local sites are attacked at a much higher rate than the rest of the original surface. There are many different types of localized corrosion, including pitting, crevice, filiform, galvanic, cavitation, waterline, intergranular, and stress corrosion cracking, which will be described next.

3.3.3 Pitting Corrosion

This type of attack is a very severe form of localized corrosion that leads to very small holes in the metal surface. The lack of oxygen around these holes is the driving power for pitting corrosion. This area becomes anodic while the area with excess of oxygen becomes cathodic, leading to very localized galvanic

Corrosion corrosion. The attack penetrates the mass of the metal, with limited diffusion of ions further promoting the localized lack of oxygen [91]. The initiation of pitting often starts with the presence of local defects at the metal surface, such as flaws in the oxide or segregates of alloy elements. In the presence of an aggressive environment containing chlorides, Cl‾ anions are believed to disrupt the oxide layer at pre-existing defects resulting in micro-cracks [92]. Recent studies show that the pits initiate and propagate in aluminium when the following chemical reaction takes place;

1- In pit initiation, aluminium is dissolved, according to the following reaction;

Al 3 cations react with Cl‾ anions to form AlCl4-. Hydrolysis of this species results in acidification of the bottom of the pit to pH < 3 (due to formation of H+ ions). The highly aggressive environment results in auto-propagation of the pit.

2- cations concentrated at the bottom of the pit, diffuse out the pit where an alkaline environment is generated by the cathode process of the hydrogen gas evolution, or oxygen reduction reaction;

As a result, aluminium hydroxide (Al(OH)3) is formed and precipitates at the pit borders blocking the pit hole for further ionic exchange processes [92]. In polarization, pitting develops at potentials more positive than the pitting or breakdown potential (Ep). The potential to which the aluminium is polarized by specific cathode reaction determines the corrosion current density and corrosion

Corrosion rate [89]. Aluminium alloys of the 7xxx series are normally clad to protect against pitting [89, 93, 94]. However, it has been noted that the strength of the alloys will be reduce by about 15% compared with un-clad alloys [4, 95]. It is clear that pitting initiation reduces the passivity of the alloys, which lowers the corrosion resistance and enables other corrosion types such as intergranular corrosion (IGC), exfoliation corrosion (EXCO), and stress corrosion cracking

(SCC) to propagate along the grains and their boundaries resulting in catastrophic failure. Several authors have classified the stages of pitting as breakdown of the passive film (breakdown potential), metastable pit formation, and pit growth [87, 96]. As mentioned above, breakdown of the passive film in pure aluminium generally occurs with weak points in the material surface that often occur at flaws in the oxide layer, such as locations of scratches and impurities, in association with the presence of an aggressive environment. On the other hand, aluminium alloys usually contain second phase particles and solute- rich regions that may disrupt the protective oxide layer and increase local corrosion attack on the surface, especially in the presence of a chloride- containing environment [4, 97-99]. However, some initiated pits do not propagate; they can be repassivated, hence causing metastable pitting. Once a pit is stable, propagation is autocatalytic as the process forms metal cations that attract the solution anions (Cl-) into the pit, the cations undergo hydrolysis which lowers the pH value that adds further aggressiveness to the pit solution [100,

101]. The highly aggressive solution and the low pH values favour more dissolution of the metal that causes pit growth and pit propagation [86].

Corrosion

3.3.4 Crevice Corrosion

Another localized type of corrosion attack develops between metal-to-metal or non metal-to-metal components contacting areas due to the formation of a crevice with an oxygen differential cell in a corrosive environment [28, 89].

Recent researches indicate that both the potential drop and the change in composition of the crevice electrolyte are caused by de-oxygenation of the crevice and a separation of electro-active areas, with a net anodic reaction occurring within the crevice and a net cathodic reaction occurring on the metal surface. Examples of crevices are gaps and contact areas between parts, under gaskets or seals, and inside cracks and seams [101].

3.3.5 Filiform Corrosion

Filiform corrosion is another case of localized corrosion that may occur on an aluminium surface under an organic coating [89]. The source of initiation is usually a defect or scratch in the coating. Filiform corrosion is associated with an oxygen concentration cell in which the anodic active area is the head of the filament and the cathode is the area surrounding it, including the tail. When aluminium ions hydrolyse, hydrogen ions are formed and reduced at the tip forming hydrogen gas. The pH value of the tip becomes more acidic. The cathodic reaction (oxygen reduction on the intermetallics that are present at the surface) takes place behind the anode. Aluminium ions and hydroxide ions form insoluble aluminium hydroxide in the tail. The active anodic reaction at the tip results in propagation of the filament and causes detachment of the coating from the substrate [102].

Corrosion

3.3.6 Galvanic Corrosion

Another type of localized corrosion attack may be initiated when electrical contact occurs between two different metals that are immersed in an electrolyte.

Such galvanic corrosion is an electrochemical process that forms between metals with electrode potentials that differ by at least 50 mV and are indirect contact with each other [89]. The electrolyte (environment) provides a means for ion migration, with metallic ions moving from the anode to the cathode. This leads to the anodic metal corroding more quickly than it otherwise would be. For example, aluminium and its alloys under most environment conditions are the anodes in galvanic cells with most other metals, protecting them by corroding sacrificially. Contact of aluminium with more cathodic metals results in an increase of the breakdown potential of the aluminium and this must be avoided in any environment in which aluminium itself is subjected to pitting corrosion [89].

3.3.7 Cavitation Corrosion

This type of corrosion usually combines electrochemical action with mechanical damage. For example, a turbulent flow of liquid below the vapour pressure can initiate voids (gas bubbles) that interact with component surfaces to dislodge the protective film during service and even alter the state of work hardening of the metal at the surface. The severity of this type of corrosion depends on the aggressiveness of the environment. Weight loss is the standard result of such attack [28].

Corrosion

3.3.8 Waterline Corrosion

Waterline corrosion can affect semi-submerged structures such as those in sea water. With aluminium alloys, waterline corrosion may arise due to a difference in the chloride level between the sea water at the air/water parting line. This effect is weak in water that is in motion because the meniscus is constantly being renewed and not as concentrated with chloride [28].

3.3.9 Intergranular Corrosion (IGC)

Intergranular corrosion (IGC) is recorded as localized attack along the grain boundaries, or immediately adjacent to the grain boundaries, while the bulk of the grains remain largely unaffected. This type of corrosion is usually associated with segregation effects (impurities enriched at the grain boundaries) or specific phases precipitated on the grain boundaries [89]. Such precipitation can produce zones with anodic behaviour that reduce the corrosion resistance in their immediate vicinity [103, 104]. The attack usually progresses a narrow path along the grain boundaries. Aluminium alloys are susceptible to intergranular attack on account of either phases anodic to aluminium being presented along the grain boundaries or due to depleted zones of copper adjacent to the boundaries in copper-containing alloys. Dix et al. studied the corrosion behaviour of Al-4%Cu in NaCl-H2O2 environment [105, 106]. They observed Al2Cu precipitates along grain boundaries and Cu-depleted zones were observed adjacent to the grain boundaries. They found that the Cu-depleted zone has a higher corrosion potential, by about 44 mV, than that of the grain matrix. They concluded that

Corrosion

IGC in Al-Cu alloys is caused by the difference in the corrosion potential between these regions. In AA 7xxx alloys, such as AA 7075 and AA 7150, both zinc and copper were found to be depleted in the PFZ [10, 68]. As was discussed in the previous chapter, zinc and copper have opposite effects on the dissolution of aluminium. Depletion of Cu makes the PFZ more susceptible than the matrix, while depletion of Zn can be beneficial to the PFZ. Therefore, it is not easy to determine the role of the PFZ in the process of IGC in AA 7xxx alloys compared with binary alloys. In AA 7xxx alloys, the grain boundary precipitates of η are reported to be more active than the surrounding matrix [42-44]. Therefore, η would be preferentially dissolved, leaving the alloy matrix with more corrosion products. It has been observed that the IGC morphology in AA 7xxx alloys in

NaCl solution is irregular and complicated and that the matrix near the grain boundary is also dissolved. This is probably due to the acidic and crevice environment that Buchheit et al. described in the IGC mechanism in Al-Cu-Li alloys [107]. Overall, the composition change at the grain boundary due to formation of either precipitates or a PFZ changes the electrochemical properties.

The difference in potential between grain boundary precipitates, the PFZ and the grain matrix makes aluminium alloys such as 7xxx susceptible to IGC attack.

3.3.10 Stress Corrosion Cracking (SCC)

SSC is a phenomenon which results in brittle failure in alloys normally considered ductile, when the alloys are subject to a tensile stress in a particular corrosive environment. It involves time-dependent interaction between the alloy microstructure, mechanical deformation, and environmental conditions. Only aluminium alloys with a high level of Cu, Mg, Si, Zn, or Li in solution may be

Corrosion susceptible to SCC. In practice, Al-Cu-Mg (2xxx series) and Al-Zn-Mg-Cu

(7xxx series) may suffer SCC [108]. When a crack occurs, it is characteristically intergranular and involves the presence of an active anodic constituent in the region of the grain boundaries [89]. The recrystallization area in the alloys is the most susceptible to SCC, and it is for this reason that composition, work hardening, and elevation in heat treatment temperatures are normally adjusted for wrought alloys to prevent recrystallization. However, the resistance of Al-alloys to SCC depends upon the direction of stressing with respect to the elongated grain structure. Maximum attack occurs if stressing is normal to the grain direction, i.e. in the short transverse direction of the components, because the crack path along grain boundaries is clearly defined [28].

Solution heat treatment and quenching introduces residual stresses into the alloy component. Quenching places surface regions of a component into compression, with the centre in tension. If the central regions are exposed for example by machining away sections in a component that has not been stress relieved, then the internal residual tensile stresses will be added to any tensile stress imposed in service, increasing the probability of SCC. Therefore, alloy products are usually given a mechanical stretch of 1 to 3% after quenching in order to reduce the level of residual stresses and increase the level of resistance onto the SCC.

In precipitation-hardening aluminium alloys aged at elevated temperatures, the resistance to SCC varies with the ageing condition. Understanding the mechanisms of SCC in aluminium alloys has involved consideration of the following microstructural features.

Corrosion

Precipitate-free zones (PFZ) at the grain boundaries. In a corrosive

environment, PFZs are considered to show as anodic behaviour with respect

to the grain centres. Moreover, strain is concentrated in these zones because

they are relatively soft.

Nature of the matrix precipitate, since cracks occur in an alloy when GP

zones (initial formation zones of precipitates as they form from the solid

solution) are present. In this situation, deformation tends to be concentrated

in discrete slip bands. Stress generated where these bands impinge upon grain

boundaries can contribute to intercrystalline cracking under SCC conditions.

Dispersion of precipitate particles in grain boundaries. In some aged

aluminium alloys, SCC occurs more rapidly when particles in grain

boundaries are closely spaced.

Solute concentrations in the region of the grain boundaries. During ageing of

aluminium alloys, local electrochemical potentials due to the solute elements

are varied due to differences in solute levels. For example, it has been

observed that a higher magnesium content in grain boundary regions results

in the adjacent oxide layer having an increased MgO content, which is a less

effective barrier against environmental influences [28].

Hydrogen embrittlement may occur due to rapid diffusion of hydrogen atoms

along the grain boundaries, as has been explained in Chapter 2 [109].

The chemical nature of atom species at the crack surface may result in

lowering of the cohesive strength of the interatomic bonds in the crack-tip

region resulting in an advancing crack.

Corrosion

Recent studies have shown that stress corrosion cracking (SCC) at grain boundaries initiates in brittle regions [35, 92, 108, 109]. Furthermore, the presence of hydrogen diffusion along the grain boundaries plays a vital part in

SCC. However, the overall process is complex and it seems probable that one or more of the above microstructural factors are involved. The relative importance of each of these factors may depend upon the particular combination of alloy and environment.

3.4 Corrosion at High Strength Aluminium Alloys

The presence of a highly stable oxide film on pure aluminium (reaction 1) increases the corrosion resistance in neutral conditions (4.5< pH < 8.5).

+ - 2Al + 3H2O → Al2O3 + 6H + 6e (1)

The oxide film is an insulator and therefore, decreases the cathodic reactivity by blocking electron conduction [97]. However, the oxide can be decreased in thickness and lost due to the presence of aggressive species, such as chloride ions, or by changes in the pH. For example, in acidic and/or alkaline atmosphere, the oxide is removed by reactions 2 and 3.

+ 3+ Al2O3 + 6H → 2Al + 3H2O (2)

- - Al2O3 +3H2O + 2OH → 2 Al(OH)4 (3)

Once the passive oxide breaks down, aluminium dissolution occurs anodically at all pH values as can be seen in reactions 4 and 5.

Al → Al3+ + 3e- (4)

- - - Al + 4OH → Al (OH) 4 + 3e (5)

Corrosion

The metal ions can be released into the solution and undergo hydrolysis as shown in reaction 6. This increases the acidity and further aluminium dissolution due to reaction 4.

3+ 2+ + Al + H2O → AlOH + H (6)

The pH of the environment plays a significant role in determining both the anodic and cathodic reactions. In neutral and alkaline conditions, the cathodic reaction is oxygen reduction as given by reaction 7. In acidic environment, the cathodic reaction is hydrogen evolution, as shown by reaction 8.

- - O2 + 2H2O + 4e → 4OH (7)

+ - 2H + 2e → H2 (8)

In general, corrosion reactions depend on the environment in which the alloy is employed, which in turn affects the stability of the surface oxide. Also, alloying of aluminium affects the surface oxide, as the presence of the intermetallic particles can decrease the protectiveness of the surface oxide film and increase the electron conductivity between the anodic and cathodic sites [110]. All of the electrons produced by the anodic reaction are consumed in the cathodic reaction which means that the anodic and cathodic rates must be the same.

3.5 Corrosion Behaviour of Intermetallic Particles

As mentioned earlier, pure aluminium has a high corrosion resistance due to having a highly stable oxide film. However, the corrosion resistance is decreased for aluminium alloys due to the presence of constituent particles. Generally, intermetallic particles are classified as anodic or cathodic relative to the alloy matrix. Particles that have corrosion potentials (Ecorr) more negative than that of alloy matrix are considered anodic (active), while those which are less negative

Corrosion than that alloy matrix are considered cathodic (noble). More details are giving on the physical properties of the intermetallic particles at 7xxx alloys in Appendix 1.

As mentioned in the electrochemical polarization section, the OCP (Ecorr) is functioning of the anodic and cathodic reaction kinetics on each of the individual components of the alloy and their respective proportions in the bulk of the alloy.

In polarization, these particles are polarized towards the bulk alloy OCP and that leads to formation of micro-galvanic cells, which normally results in the protection of the more noble component at the expense of corroding the other [1,

96]. Consequently, active (anodic to the matrix) particles undergo anodic self- dissolution, while noble (cathodic to the matrix) particles corrode the surrounding matrix. Figure 3-6 demonstrates the concepts of this mechanism

[111].

The preferentially dissolved particles which are deemed anodic to the matrix do not reveal a breakdown potential [1]. That indicates that they are unable to maintain a passive film, either because the films are soluble in the aggressive solution or the particles are polarized to values at which they are rapidly dissolved. These particles are always at risk of attack. The more noble particles exhibit different forms of corrosion mechanisms to the alloy matrix. Due to galvanic coupling [112, 113] or elevation of the pH [110, 114], these particles lower the breakdown potential of the surrounding matrix and lead to localised attack. These particles always increase the attack of the surrounding matrix. High strength aluminium aerospace alloys, such as 7xxx series are susceptible to corrosion attack. The major alloying additions, zinc, magnesium, and copper, form a number of precipitates of varying composition. The precipitates of interest

Corrosion to the corrosion behaviour contain copper as these are preferential sites for the cathodic reaction and cause pit initiation. Another precipitate of interest is the

MgZn2 phase, which can continuously precipitate along the grain boundaries; this phase is highly reactive and can dissolve rapidly resulting in IGC attack [10].

Much attention has been focused on potential differences between the matrix and intermetallics [1, 72, 87, 88, 96, 115-117]. Recently, Andreatta et al. used a micro-capillary cell and scanning Kelvin probe force microscopy (SKPFM) measurements at random positions on AA 7075-T6 and solution heat-treated AA

7075. They found Al7Cu2Fe, (Al,Cu)6(Fe,Cu), and Mg2Si intermetallic particles distributed on both alloy surfaces. They measured the potentials for each phase and found that Al7Cu2Fe intermetallics induce a higher cathodic breakdown potential than the (Al,Cu)6(Fe,Cu) intermetallics that corresponded to the higher

Volta potential differences for the former particles (-475 mV) compared with the later particles (-282 mV versus SCE) [118]. They concluded that the breakdown potential of areas containing intermetallics is related to the Volta potential difference between the intermetallics and the surrounding matrix [118]. On the other hand, Birbilis investigated the electrochemical properties on a phase-by- phase basis by fabricating test alloys that could provide a population of the intermetallics of interest large enough to be studied independently by the potentiodynamic polarization microcell. He studied the intermetallic particles in

AA 7075 in 0.1 M NaCl and found differences in potential between them and that the corrosion rates for each phase can be extracted from polarization curves by inspection of the current at the OCP of the specific alloy in the same solution

[1].

Corrosion

It is clear from the above discussion that the presence of various second phase particles in 7xxx alloys affects their corrosion behaviour strongly and leads to localised (micro-galvanic coupling with the matrix) corrosion attack which at the end reduce the alloy resistance to corrosion and causes more dissolution of the surrounding matrix.

The reduction or elimination of all of these particles from the alloy surface without affecting other alloy properties would significantly increase the corrosion resistance of the aluminium alloys. That improvement may be expected to result from the use of laser surface melting (LSM) techniques and that will discuss later in the next Chapter.

3.6 Aluminium Alloy Surface Protection Methods

There are several surface protection methods that can be applied to aluminium alloy surfaces in order to increase their corrosion resistance for a given application. The basic idea for surface treatment is to provide the aluminium alloy surface with a stable layer or coating that protects the alloy from environmental attack. The processes for surface protection of aluminium alloys, including non-galvanic, chemical, and electrolytic methods, will be reviewed in this section [13, 14].

Corrosion

3.6.1 Non-galvanic Methods

Surface enhancement of aluminium that does not include chemical or electrolytic action is considered to be a non-galvanic method. It includes mechanical finishes, such as grinding, polishing, buffing, rolling or blasting to remove any scratches or dirt from the surface, and the application of organic finishes, such as primer/top-coats, plastic coatings, protective coatings or powder coatings. These processes are mostly used in architecture and light components to increase the corrosion resistance of the alloy and give it a suitable selected colour [13, 14].

3.6.2 Chemical Methods

Surface treatment of aluminium by a chemical reaction, with no external source of electric power used, is considered to be a chemical method. Etching, chemical polishing, chemical conversion coating and electroless processes are all examples of chemical methods. For applications such as aerospace, chemical conversion coating is extensively used to protect alloys against corrosion [13, 14].

3.6.3 Electrolytic Methods

Surface treatment of aluminium including electrochemical action and external power supply is considered to be an electrolytic method; examples include elctropolishing, electroplating and anodising. Anodising is the more commonly used electrolytic method for surface treatment of aluminium and aluminium alloys. This is due to a number of properties offered by this technique, such as increase corrosion resistance, provision of hard and wear resistant surfaces, and electrical and thermal insulation, and the surface can be used as a base for

Corrosion organic finishing and plating. This technique has an advantage that, with contain anodising processes, it does not require use of toxic chemical or heavy metals

[14, 119].

In most protection methods, the presence of different constituent particles and precipitates, which are the source of most corrosion problems, cannot be removed and/or decreased at the surface of the alloys. In the presence of an aggressive environment containing chloride ions, these particles may be harmful, for example leading to IGC and SCC. It is often difficult to detect the corrosive attack and repair it before it causes massive failures. Therefore, it is important to eliminate or reduce the number of intermetallic particles from the alloy surface without affecting the alloy bulk properties in order to increase the corrosion resistance which can be done using LSM as can be shown later in this work. The principles of the lasers technique and the theory of rapid solidification of the alloy surface will be discussed and reviewed in the next chapter.

Corrosion

3.7 List of Figures

Figure 3-1: Schematic diagram showing the thin oxide layer (Al2O3) that protects the metal surface (aluminium) from environmental attack.

Corrosion

acid corrosion alkaline corrosion

passivity

immunity

o Figure 3-2: Pourbaix diagram for the Al-H2O system at 25 C that shows regions of immunity, passivity, and corrosion [120].

Corrosion

Figure 3-3: Corrosion polarisation diagram showing Tafel extrapolated slopes for the both anodic and cathodic current density. The intersecting point indicates the both corrosion potential (Ecorr) and the corrosion current density (Icorr) [85].

Corrosion

Figure 3-4: Schematic diagram of the polarization behaviour of a metal showing the transformation from active to passive states and the breakdown of passivity at point F [121, 122].

Corrosion

Figure 3-5: Intergranular corrosion (IGC) of a failed aircraft component made of 7075-T6 aluminium alloy (width of image = 500 µm) [104].

Corrosion

Figure 3-6: Schematic diagram showing the corrosion behaviour due to intermetallic particles that are anodic and cathodic with respect to the alloy matrix [111].

Chapter 4 Laser

4. Introduction

The surface treatment of metals and alloys using a laser is a subject of considerable interest in present because it offers the chance to save strategic materials or to improve components by optimizing surface and bulk properties

[123, 124]. This Chapter will present the basic principle of the laser, laser beam characteristics, LSM including heating principles, the formation of a melt pool, cooling and solidification of the melted surface layer, nucleation and growth. The microstructure as well as the corrosion behaviour of LSM alloys will be discussed and reviewed. Finally, the safety aspects of the use of lasers, including laser classifications and laser protection forms, will be explained.

4.1 Basic Principle of Laser Works

The name Laser is an acronym for Light Amplification by Stimulated Emission of Radiation [123, 124]. Light of the laser is generally accepted to be electromagnetic radiation ranging from 1 nm to 1000 µm in wavelength. It is reported that the wavelength of the visible laser light ranges from approximately

400 to 700 nm. From 700 nm to 10 µm is considered the near-infrared (NIR), and beyond this limit is the far-infrared (FIR). In contrast, 200 to 400 nm is considered ultraviolet (UV); while below 200 nm is deep ultraviolet (DUV) [124,

123].

Laser

General lasers employ a resonator that typically consist of two mirrors are placed parallel to each other in order to form an optical oscillator in which a light can move back and forth between the mirrors in an active medium that consists of high purity material state(gas, liquid or solid) which is capable of amplifying the laser light. An external source of energy (electric current) pumps in the gain medium, laser radiation is generated then in the medium which consists of atoms or molecules that have specific electron energy levels. The movement of electrons between the energy levels causes a spontaneous emission of radiation.

A minimum of three electron energy levels are required for a laser to generate

[123, 126]. These energy levels are ground level, a metastable level and a high level, as indicated in the Figure 4-1. When an electron is excited from the ground state to the high level with energy input, laser then can produce. The number of electrons in the high energy level must be increased to more than the number of electrons in the ground level, and this is referred to as population inversion. To invert the electron population requires that more than half of the electrons to be raised to the high energy level. Following the population inversion, electrons then lose energy and consequently reduce the energy levels (electron decay), and therefore, the photons are emitted by the movement of the electron from the metastable level down to the ground level. To amplify the amount of photons produced, the reduction of electrons is stimulated by the interaction of additional photons excited atoms in the gain medium. The radiation emission of laser light can be generated then in the form of a continuous wave (CW) or pulses. For a continuous laser, the energy required to invert the electron population must be

Laser reduced so that the inversion can be easily maintained, and that can be achieved by increasing the number of electron energy levels to a minimum of four with two metastable levels, as illustrated in Figure 4-2, while in a pulsed laser, the energy input must be high enough to maintain the population inversion and that can be achieved in three electron energy levels, as this can provide the same peak energy to invert the electron population, but the overall average input energy is lower than that in four electron energy levels [123, 126]. It is mentioned that once continuous radiation has been produced, it can be changed to provide radiation of pulses by using Q-switching that enables storage of enough radiation and controls the release of this radiation in a form of pulses [123]. The length of the laser pulse can vary from tenths of seconds to a few femtoseconds (10-15 s).

The average power may vary between milliwatt and kilowatt levels, with peak power attaining the order of gigawatts [127].

4.2 Laser Light Characteristics

Laser light has four main characteristics that differentiate it from the other form of lights. It is monochromatic, has low divergence, coherent, and high brightness

[123, 124, 126]. Monochromatic laser light is employed in applications such as measurement, alignment, and holography. Low divergence is the property that enables a laser beam to retain high brightness over a long distance, and is the basis of alignment systems. Coherent radiation comprises waves travelling with the same wavelength, amplitude and wave-front. It is a measure of the degree to which light waves are in phase in both time and space. The light of the laser radiation has a high coherence. Coherent laser light is up to 100000 times higher in intensity than incoherent light of the same equivalent power, since the

Laser divergence or dispersion of energy is very low as the beam propagates from the

laser. The coherence length ( lcoh ) depends on the wavelength (λ), and the bandwidth (∆λ),

lcoh

The coherence time ( tcoh ) depends on the ( ), and the velocity of light (c),

l t coh coh c

Brightness (B) is a measure of intensity of light at a particular location. It is the emitted power (q) per unit area (A) to the solid angle (Ω),

q . A

The Intensity (I) obtained by focussing a beam of light, and is directly proportional to the brightness (B).

I = π w2 / f2

Where f is the optical focal length, and w is the beam waist radius. The beam intensity can be related to the temperature through the Stefan-Boltzmann law,

. 4

Where σ is Stefan-Boltzmann constant (5.67 x 10-8 J/m2s K-4) and T is the absolute temperature of the emitting surface. By focusing, the intensity of a beam can be increased as well as the temperature of the emitting structure that causes more phonon-electron energy exchanges. These are more likely to interact with structure rather than oscillate and reradiate. Therefore, the absorptivity on the structure surface will increase and reflectivity will decrease.

Laser

Thermal mechanisms of material processing take advantage of the high brightness (high power density) of a laser beam. Athermal (photonic) mechanisms are based on the short wavelength (high energy) of the beam, and the short duration of the pulses that can be produced. The beam characteristics influence the beam propagation and focusability, and therefore have an important effect on the suitability of the beam for material processing. The polarization of a laser beam also affects the amount of light observed in many material processing applications. Light sources interact differently with a material surface as the angle of incidence, θB, increases (The angle at which complete absorption occurs and is called the Brewster angle). It is related to the index of refraction, n, by

-1 θB = tan (n). Beam polarization affects the amount of energy observed by the material, and therefore enhances the efficiency and quality of laser processing

[126,127].

4.3 Laser Surface Melting (LSM)

LSM improves the corrosion resistance of the alloy surface microstructure through localised melting and rapid solidification. To understand the change from an heterogeneous microstructure into an homogeneous by LSM, requires understanding the laser heating source, formation of melt pool, cooling and solidification, as well as nucleation and growth.

4.3.1 Heating Principle

The energy input to the substrate is low in comparison with a conventional means of surface melting. Heating rates associated with laser surface melting are

Laser therefore orders of magnitude higher than conventional methods. It is reported that the heating rates for typical continuous laser surface melting are on the order of 105 K s-1 for an energy input of 0.1 J m-1 associated with large-scale processing, while for a nanosecond duration excimer laser can reach 109 K s-1 with energy of 10 J. [124] Therefore, melting occurs rapidly with an increase in the absorptivity of the material to the focusing laser beam.

4.3.2 Formation of the Melt Pool

The melting of the surface results in formation of a melt pool. It is reported that temperature gradients of an order of 104-106 K m-1 develop between the centre and the cooler solid/melt interface at the bottom of the pool [124]. In most materials, the coefficient of surface tension increases with a decrease in temperature. Therefore, surface tension gradients induce fluid flow from the centre of the melt towards the edges (called Marangoni flow) [124]. Marangoni flow is the dominant convection mechanism in the melt pool since there are no other forces induced on moving charged particles in the presence of magnetic and electric field. The Marangoni number (Ma) is the mass transfer along an interface between two fluids due to surface tension gradients. Therefore, it is the ratio of the convection rate and the conduction rate of molten material. Typical

Marangoni numbers for metals during melting lie in the range 103-106 [124]. For metal solutes in molten metals, it is reported that the diffusion coefficients of the liquid state lie in the range of 10-6-10-7 m2 s-1. [124] Therefore, the diffusion length for a very short melt times is considerably smaller than the melt depth, such that homogenisation occurs mainly by convection. Material flow by

Laser convection due to surface tension gradients is preferential in LSM as the dispersion rate of the solute is faster.

4.3.3 Cooling and Solidification

Cooling rates during LSM are significantly high. It is reported that cooling rates

~ 103 K s-1 for a mutikilowatt surface melting, rising to 1011 K s-1 for very low energy input surface skin processing [128]. In rapid solidification of metals and alloys, thermal and constitutional undercooling temperatures are important parameters to be considered (Temperature gradients). Thermal undercooling

(∆T) is the degree of liquid cooling below the equilibrium freezing temperature

(Tf) before solidification takes place (∆T = Tf – T), and occurs during rapid solidification in pure metals due to high cooling rates. The composition of the alloy changes over the liquidus range in the solidification process. As the composition changes ahead of the solidification interface, the actual temperature

(T) may not match the temperature required by the alloy liquidus, and this leads to constitutional undercooling. Therefore, the constitutional undercooling temperature has an influence on solidification structure of alloys, which may cause by thermal gradients being less steep than melting point gradients [129,

130]. On the other hand, the driving force for solidification is the thermal gradients (G) within both solid (GS) and liquid (GL) at the interface that also influence the solidification structure. GL is more critical than GS since it controls stability and the morphology at the interface. Therefore, any movement of the material within melt pool formation will affect GL and solidification microstructure. The rate of solid/liquid interface with respect to the liquidus is the interface growth rate (R*). It is reported that G influences the solidification

Laser morphology microstructure, and R* influences the scale of solidification and solute distribution. Therefore, G and R* are important, as they influence the solidification morphology, and the scale of solidification substructure respectively [130, 131]. Nucleation and growth scenarios will take place during solidification and that explain in the next sections.

4.3.4 Nucleation

In LSM application, once the melt in pool area of laser treated surface is mixed in the liquid state, the nucleation of solid is required (amorphous or crystalline).

The formation of a solid particle from a liquid can be controlled by the change in total free energy of the applied system (∆G). It is reported that the difference in the volume free energy of the particle and surface free energy of the solid/liquid created interface. Figure 4-3 shows the relationship between surface, volume, and total free energy [132]. The Fig. shows that the volume free energy is negative and supporting the nucleation process (growth of cluster), whereas the surface free energy is positive and is therefore an activation barrier for the formation of a stable nucleus. An increase in thermal undercooling (∆T) may then increase the energy barrier for nucleation and reduces the critical size for stable nuclei (r*) (dissolving of cluster) [132]. Two types of nucleation can form, namely homogeneous and heterogeneous. Homogeneous nucleation occurs when atoms cluster together to form a stable nucleus, without the aid of foreign bodies such as impurities, inclusions, or interaction with the substrates. A large amount of undercooling is required in homogeneous nucleation to supply the energy needed for the formation of the new surfaces, which described above. In heterogeneous nucleation, less undercooling is required as this type of nucleation

Laser uses the surfaces of foreign bodies (e.g. particles suspended within the melt) to significantly reduce the surface energy requirement [130, 131]. Heterogeneous nucleation is the standard situation for most cases of solidification [131].

Heterogeneous nucleation occurs as the liquid fully wets the solid substrate it is in contact with, so there is no nucleation energy barrier and nucleation occurs spontaneously. The full wetting of the surface by the liquid also means that the solid that grows is influenced by the substrate [133].

4.3.5 Growth

Once nucleus formation has been stabilised, growth of solid starts at the interface with addition of atoms from the liquid. The growth structure morphology that occurs with more adding of atoms is controlled by the shape and stability of the liquid/solid interface. The stable interface shape for growth depends on a number of conditions, including heat flow rate, thermal gradients, mass flow rate, composition gradients and growth rate [125, 130, 131]. Typical growth morphologies are planar, cellular, and dendritic; Figure 4-4 shows different growth morphologies that can be formed during solidification [125]. The direction of growth in the planar mode is perpendicular to the flat solidification front and parallel to the thermal gradient (G). In planar growth, the role of crystallography is limited but can affect the growth rate (R*), as some crystallographic directions can grow faster than others. The aligned grains in the faster growth direction will develop at a faster rate and can dominate the microstructure [125, 130]. Cellular growth occurs where the direction of grain growth is controlled by heat flow conditions with very small influence from crystallography. Cellular growth is not the same as planar growth since it occurs

Laser at faster cooling rates. Therefore, the solute in the alloy cannot redistribute fast enough, leading to rejection of solute ahead of the interface. This rejection in solute results in a change in composition ahead of the interface and constitutional undercooling occurs. Constitutional undercooling means that the liquid ahead of solidification exists below its freezing temperature and solid protrusion start to develop on the planar interface. The protrusion can then grow into the liquid with more stability, following the maximum thermal gradient. The growth of a single protrusion rejects solute laterally, that lowers the solidification temperature and produces cavities that help the formation of more protuberances. These protuberances break down the planar interface and form a cellular structure with more uniformly spaced cells that grow parallel to the maximum thermal gradient.

An increase in the cooling rate decreases the cell size dimension. Therefore, growth like cellular can be formed to each other [125, 130, 131]. The transition from cellular to dendritic solidification is not completely recognised, but it is thought to be due to constitutional undercooling between the cells that causes instability perpendicular to the growth direction. This unstable growth of the cell sides causes to ramify the cells into dendrites, as shown in Figure 4-5 [130]. The increase in the interface instability forming dendrites from cells is related to an increase in the cooling rate. Fast cooling rates can influence the crystal structure and growth direction. Therefore, a dendritic growth morphology occurs in specific crystallographic directions and has been identified in two forms, columnar and equiaxed. Columnar dendrites are not affected by heat flow rate and can grow along the same crystallographic direction parallel to each other

(Figs. 4-4 (c) and 4-5 (d)), while equiaxed dendrites grow in different orientation depending on crystallography and neighbouring dendrites are not necessarily

Laser parallel to each other. Therefore, each separate grain can form equiaxed dendrites

(Fig. 4-4 (d)) [125, 130, 134]. As the dendritic structure is influenced by crystallography, the increase of the cooling rate can influence the dendrite spacing, resulting in a fine structure. In the other way, there is a limit to the refinement of the dendritic structure with the increasing cooling rate. This is due to the decrease in atom diffusion time with very high cooling rates [129, 135].

For stable growth dendrite structure, the liquid/solid interface must be able to form stable protuberances along the interface. For stable and growing protuberances, it must be able to accumulate atoms by diffusion that means that the distance between interface protuberances must be less than the atomic diffusion distance during solidification. Increasing the solidification cooling rates reduces the diffusion distance and interface protuberances distance and this requires very high surface energy. Once the high surface energy for stable protuberances can no longer be reached, planar growth again becomes stable and that means there is a limit of stability for such planar growth [130, 131].It is reported that above this limit, planar growth produces a segregation free microstructure by trapping the solute in solid solution. Therefore, a uniform composition across the melted layer can be produced with no signal of second phases and extended element solubility limits.

4.4 Effect of LSM Interaction

Some lasers operate continuously, and are termed continuous wave (CW) lasers, while other lasers operate discontinuously and are known as pulse lasers. For

CW lasers, their interaction time is controlled by the translation speed of either the laser or the sample. For pulse lasers, the pulse duration is the interaction time.

Laser

Different lasers produce different wavelengths. Table 4-1 shows examples of

LSM types and their wavelengths [164]. The laser gain medium is responsible for producing these wavelengths which then can be used as a selection criterion for particular applications of LSM [136, 137,164].

In LSM, radiation is absorbed into a thin surface region of metal or alloy and thermal energy is then conducted away from the surface to the material bulk

[138]. Figure 4-6 shows a schematic illustration of the laser beam interaction with the surface of the material [138]. The red arrows in the Fig. represented the thermal diffusion on the surface. The high temperatures that are produced by the laser are enough to melt the surface and a fusion interface can develop that propagates from the substrate. Further increasing in the heating rates can result in vaporisation of the material that causes problems in the LSM as it creates an intense plasma which can deflect the laser beam and lower the energy density at the surface. As vaporised material escapes from the molten material, it can cause poorer and un homogeneous surface with an increase in the surface tension sideways movement in the melted area [139-141]. Therefore, rapid cooling occurs once the radiation is removed from the surface that freezes the material to form non-equilibrium structure as there is not sufficient time for diffusion or precipitation [142-144]. The high cooling rates occur as the substrate material acts as a heat sink and it is reported that the cooling rates in excimer LSM can

9 11 3 5 reach 10 -10 K/s in comparison with CW-CO2 laser that reach 10 -10 K/s

[145, 146].

Laser

4.5 Effect of LSM Processing Parameters

As have been mentioned in the previous section, different lasers produce different wavelengths that depend on the laser gain medium, which are then used as laser selection criteria for an application such as LSM. The selection process of LSM is usually based on its inherent characteristics, including pulse duration and wavelength. In some lasers, the pulse duration and the wavelength can vary due to change in the gain medium excitation that alters the resultant pulse duration [147]. The changes in these characteristics, as well as other processing variables such as energy, a number of pulses per unit area, scan rate, atmosphere, spot size, and the beam overlap area, can affect the surface of the laser treated layer [148]. In excimer LSM, the laser energy density and the number of pulses per unit area have great influences on the laser melted layer than that of the other laser processing variables.

4.5.1 Energy

The objective of LSM processing on metals and alloys is to modify their surface structures and that depends strongly on the laser energy density or laser fluence

(J/cm2) as well as the number of pulses of the laser used. Laser processing with a very low fluence may not induce material modification as there is not enough energy to cause melting. However, there might be sufficient energy available for removing any contamination from the surface, such as strongly bound dirt or paint [141, 149]. In laser processing with a high fluence (high energy density), the degree of surface heating causes more material melting that results in an increase in the thickness of the melting layer [150-153]. An increase in the

Laser fluence means an increase in the energy input to the substrate to a level higher than required for melting alone. At this higher level, evaporation occurs for the melted regions that affect the thickness of the modified layer and produces poor melting. Therefore, there will be a maximum layer thickness that can be reached by increasing the laser fluence [148]. To induce material modification through surface melting requires an increase in the laser fluence up to the melting threshold value [147, 148, 150]. Under this value, there is a step change in the material state as well as enhancement of radiation absorptivity [147]. The change in the radiation absorptivity of the melted regions means that in the LSM process, the energy absorbed is not constant over the pulse duration. If the changes in absorption during LSM could be controlled by varying the fluence over the pulse duration, there would be an increase in process efficiency and decrease in material vaporisation [147].

4.5.2 Wavelength

For CW-LSM, the absorption of the laser radiation into a surface is related to the radiation wavelength, while in pulsed-LSM there is no clear relationship between wavelength and absorption due to the influence of the pulse duration, but generally an increase in pulse duration will increase absorptivity [135, 152]. The surface to be treated can also affect the radiation absorptivity in different ways including roughness, oxidation, temperature and state [150, 152, 154]. For example, there is higher radiation absorptivity when the substrate surface is rough or in the presence of surface oxide [155]. In term of highly reactive materials, such as aluminium, the application of an absorbent coating such as

Laser graphite may be used to enhance the radiation absorptivity on their surfaces

[135].

4.5.3 Pulse Duration and Interaction Time

It is mentioned in the section 4-5 that the interaction time for continuous wave

(CW) lasers is controlled by the translation speed of either the laser or the sample whereas in pulse lasers the interaction time is determined by the pulse duration.

Generally, the interaction time of CW-lasers is longer than that of most pulsed lasers due to limitations in the scanning speed systems that have to be employed.

It is reported that longer interaction times for CW-lasers tend to form thicker modified layers with slower solidification rates and more segregation than with layers produced by pulsed lasers [142-157]. Increased segregation in a thicker solidification structure will cause a decrease in corrosion resistance, so thinner layers produced by pulsed lasers with more rapid cooling rates would be beneficial in terms of corrosion resistance [158-163]. The layer thickness produced in LSM processes is influenced by both the interaction time and the fluence [145-150]. This influence can be explained using the surface temperature gradients and vaporisation threshold [150-152]. Initially, the laser surface melting temperature is higher than the melting temperature but less than vaporisation threshold and therefore, the melt depth can be increased by an increase in energy density and melting until the temperature at the surface exceeds the vaporisation temperature (boiling temperature) and the maximum possible melt depth is reached. Since any further increases in energy density will cause vaporisation and loss of material, the pulse duration becomes important as an increase in pulse duration will further increase the melt depth at constant

Laser power [159]. Figure 4-7 shows the influence of pulse duration in the LSM of pure metals with respect to the maximum depth of the melting layer [159]. On the other hand the interaction time of the laser beam is influenced by the pulse repetition frequency (PRF), and if the PRF is decreased, the interaction time will be increased at constant power. Therefore, an optimum processing interaction time can be controlled under PRF. Note that the fluence value required for the surface temperature to reach the vaporisation temperature may be dependent on the pulse duration too [158].

4.5.4 Number of Pulses per Unit Area

The number of pulses per unit area used in LSM processes has been observed to affect the melting structure in a number of ways. It is reported that an increase in the number of pulses per unit area causes an increase in melt depth [180, 201-

204]. If each pulse interacts with the surface material to create an increase in melt depth, several pulses can increase the region of melted depth and subsequently the thickness of the melted layer. However, as an increase in the number of pulses causes movement to be induced in the liquid material

(explained in section 4.4.2). This movement in the liquid state pushes the material away from the centre of the laser pulse to the edges and with increasing the repeated number of pulses per unit area results in an increase in the peak layer thickness and the surface roughness [159, 160, 180]. Also, to apply a large number of pulses to a single location would lead to an extremely slow process time when treating large areas of material. This will then impose a limit on the number of pulses and hence a level of homogenisation that can be achieved

[161]. Further studies of samples treated with a large number of pulses per unit

Laser area showed that they can suffer from severe porosity [162, 180]. Generally, metals need to be treated with a large number of pulses per unit area practically influenced by laser energy density, pulse duration, interaction time, and scan rate. If these parameters are set correctly with respect to material composition, then the quality of the modified-melted layer can be improved.

4.5.5 Process Environment

LSM processes can be applied in different environments; each has its own advantages [123, 139, 142, 177-179]. For example, excimer LSM can be used in air and/or gas atmospheres. In both environments, the structure of the substrate surfaces can be modified. In order to compare between these modifications, it is important to know the nature of the reactions at their surfaces. In a nitrogen environment, the nature of the chemical reaction is to increase the amount of nitride precipitates within the modified layer, which is beneficial in increasing wear resistance and hardening, while in an open environment, the modified layer is beneficial in terms of corrosion resistance as the main principle in using excimer laser in this project is to eliminate the large number of intermetallic particles and precipitates from the alloy surface and produce a uniform melted layer that assist in improving the corrosion resistance without affecting the alloy bulk properties.

4.6 Microstructure of LSM Alloys

At very high solidification rates (higher than the local diffusion rate), equilibrium at the liquid/solid interface cannot be established and the solute atoms are frozen

Laser into the solid at the same composition as they arrive at the interface (solute trapping). However, at low solidification rates (less than the diffusion rate), local equilibrium can be established and interface crystals can achieve the same composition as the melt [124, 163]. Therefore, a high cooling rate, as in LSM results in a fine solidified microstructure, which may contain non-equilibrium phases, new precipitates and extended solid solubility. Rapid solidification microstructures have been investigated after using different LSM techniques

[129, 135, 153, 164, 165]. Studies, by TEM, of microstructures produced by

CW-CO2 LSM of Al-Cu alloy confirmed that the microstructure morphology was a function of the dendrites growth rate in rapid solidification [153]. Furthermore, for LSM of Al-Cu alloy, a map has been created that correlates the microstructure to both the growth rate and the material composition [165]. A variation in solidification structure due to composition and growth rate changes has been shown to occur through the thickness of modified layers. For example, a planar growth region was observed at the base of modified layers in CW-

Nd:YAG LSM of AA 2014-T6, whereas the rest of the layer had a segregated cellular structure, as can be shown in Figure 4-8 [161]. The base of the modified layer is often referred to as the fusion boundary. Solidification always starts and accelerates from the fusion boundary, that means that the depth of the modified layer has an effect on the growth rate and the fastest solidification front velocity that can be achieved within it. Ryan et al. have modelled the variation of solidification in excimer LSM of AA 2024 for a single radiation pulse [166].

They observed that there is an initial transient region on re-solidification following each pulse, and the growth front is rapidly accelerating. They found that the growth front travels at speed of 5 m s-1 after a distance of ~500 nm,

Laser whereupon it appears to obtain stability. They concluded there will be a narrow initial region of growth at start of re-solidification of each pulse where the front can become morphological unstable, resulting in the formation of solute segregation or banded structures. Figure 4-9 shows the results from this modelling [166]. CW-Nd:YAG LSM of AA 2014-T6 alloy showed that the variation in the growth rate changed the dendrite arm spacing from 5 µm at fusion boundary into 2 µm near the top surface as illustrated in Figure 4-10

[161]. However, more segregation of elements was observed at the dendrite boundaries, as can be clearly seen in the BSD image in Figure 4-11 [161]. It is clear that the growth rate has been varied throughout the modified layer depth and this would indicate that changing the modified layer thickness would also change the growth rate along the solidification structure. Variations in modified layer thickness have been attributed to the solidification structure scale [161]. An increase in LSM layer thickness causes a decrease in the growth rate which in turn increases the cell spacing and therefore, dendrites can be growth, and that can be seen clearly in Figure 4-12, where different LSM processes have been used to treat AA 2024-T3 [161]. A strong influence of the substrate material on solidification was shown in studies of excimer LSM on AA 2024 and AA 7150

[166, 167]. LSM of these two alloys under identical processing conditions produced different solidification structures. Planar growth was observed in AA

2024, and cellular growth observed in AA 7150 as shown in Figure 4-13. The authors explained the differences in the solidification structures as being due to differences in the alloying compositions. In AA 7150, cellular growth grew from an equiaxed grain region that nucleated from either seed grains that were coherent with the base layer or Al3Zr particles that solidified ahead of the

100

Laser solidification front. Therefore, the presence of zirconium in AA 7150 was considered to be the main cause for cellular growth initiation and the dramatically changed solidification structure compared with AA 2024 [166]. In contrast, the authors investigated excimer LSM on both Zr-free AA 7075 and AA

2096 containing Zr/Li. The solidification layers on both alloys showed opposite behaviour relative to AA 7150 containing Zr and AA 2024. A fine columnar grain structure formed over AA 2096, which contains Zr, and a thin planar on Zr- free AA 7075. Figure 4-16 shows the cross-sections of both alloys after LSM treatments [166].

Several workers have carried out investigations on Al-Cu binary alloys after

LSM [168-170]. The solidification of the melted layer has shown to be epitaxial, starting with a very thin layer of planar front growth [168, 169, 170]. For alloys containing less than 5 wt% of Cu, the growth after the planar front region has been reported to consist of dendrites of α-Al with either CuAl2 precipitates [168,

169] or CuAl2-Al eutectic as interdendritic phase. Muntiz showed the existence of solute banding within the resolidified layer, which results from the differences in the microstructure of the precipitated interdendritic phase due to local changing in the solidification rate [173]. Milewski et al. used Nd:YAG LSM on

AA 2024 and observed similar solidification microstructure results to those reported by Muntiz [176]. Noordhuis an De Hosson studied the nucleation of precipitates in this alloy following 1.3 kW CO2 LSM and noted that the nucleation could be induced after surface melting by shot peening for as little as

10 s [177]. In terms of Al-Zn alloys, Lasek et al. treated Al-Zn alloys with a CO2 laser operating at 200-1400 W. In Al 30 wt.% Zn, small spherical GP zones were

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Laser seen to form in the partially melted grains at the heat affected zone (HAZ) region

[178]. Epitaxial planar front solidification was seen to occur at the maximum melt depth, and transformation from planar front solidification to dendritic growth was preceded at the region of planar front growth also contained solute enriched droplets that solidified and precipitated β-Zn spheroids on the final cooling [178]. For Al with 1 wt.% Zn, Lasek et al. found that there was no cellular structure within the resolidified structure. In alloys containing 1, 6 and

10 wt.% Zn, a banded structure was observed [179]. Lasek mentioned that cracking occured along the grain boundaries into the melted region and the presence of high residual stresses concentrated in the HAZ and at grain boundaries enabled the heterogeneous nucleation of equilibrium β-Zn particles during room temperature ageing. Roosz et al. [181] investigated the surface melting of Al 6 wt.% Zn and 2 wt.% Mg alloy. The microstructure of the modified layer, in this case, consisted of planar front region at the maximum melt depth that extended to the melt by as much as one quarter of the total melt depth

(~40/150 µm). The rest of the melt zone then solidified with a coarse cellular dendritic structure [181].

Overall, the solidification microstructure of LSM on aluminium alloys can be varied with different forms and shapes depending on the material composition and the laser surface melting used. Planar, cellular and dendritic are the most common solidification structure types that are formed in alloy surfaces. Planar is the most appropriate structure for more corrosion resistance due to production of uniform and homogenous melting layer above the bulk alloy. The variation of the growth rate with rapid solidification can affect the melted layer thickness as well

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Laser as the structure morphology. As mentioned above, decreasing growth rate will increase the layer thickness and cell spacing that favours the formation of more dendrites.

4.7 Corrosion Behaviour of LSM Alloys

In this part, corrosion investigations will be reviewed for different LSM of aluminium alloys. Some workers noted an improvement of alloy surface properties regarding corrosion attack. Bonora et al. were the first to study the effect of LSM on the corrosion behaviour of pure Al using a Q-switched ruby laser [154].Comparison between laser treated Al and untreated Al, showed that the passive current density was lower in LSM Al than untreated Al under both potentiostatic and potentiodynamic control, while the breakdown potential was unchanged after LSM. They concluded that the lower corrosion rates were due to a chemically inert, amorphous aluminium surface layer produced by LSM. This effect was also produced for pure Al after LSM using a Q-switched Nd:YAG laser, with corrosion tests carried out in boric-borate solution containing 1 g/l

NaCl [171]. Moore et al. studied the effect of LSM on the electrochemical behaviour of AA 3003. They melted the alloy surface with a single pulse CO2 laser in vacuum. In anodic polarization tests, they observed that the rate of anodic dissolution was less for the melted samples. However, the pitting potential in deaerated 0.1 M NaCl was not improved and the free corrosion potential was increased by 300 mV, having the unexpected effect of decreasing the difference between the free corrosion potential and the pitting potential [182].

Moore et al. also used CO2 LSM to study the effect of laser surface modification on the pitting behaviour of AA 2024. No improvement in the pitting resistance after LSM was found due to many cracks and pores that were present in the

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Laser melted layer that increased localised corrosion [182]. Another problem with cracks and localised corrosion is that once the base of the melted layer is breached, i.e. the transient region zone, there is an accelerated corrosion perpendicular to the surface along this region causing layer delamination [162,

165]. It is reported that even with no cracks, layer delamination has been observed once uniform corrosion has occurred over the surface and exposed the transient region (interface) that contains different forms of solidification microstructure [165]. McMahon [184] and Watkins et al. [185] reported that the corrosion resistance was increased in LSM AA 2014, as the potential required to cause appreciable pitting (Ep) was up to 100 mV higher after LSM. Ferreira et al.[186] showed that for LSM of AA 2024-T351, the form of corrosion in 3 wt.%

NaCl solution was changed compared with the untreated alloy. For the as received alloy, both pitting and intergranular corrosion were observed, while for the LSM alloy only pitting corrosion was present. In AA 7075, there was a decrease in the corrosion potential and an increase in the passive range during polarization in deaerated 0.1 M NaCl solution. The increase of the passive range may be that the LSM has low cathodic reactivity. Also, galvanic corrosion is less of a risk if the LSM layer acts as a sacrificial layer for zinc-containing aluminium alloys, where zinc can go into the solid solution providing more anodic dissolution which led to the LSM layer being dissolved and the exposed substrate being cathodically protected [157]. Further to these investigations, LSM can remove surface precipitates by dissolving them into the melted solid solution, which can reduce corrosion behaviour by protecting the alloy matrix. For example, increase of Cu in the matrix will decrease the anodic activity [24].

However, in wide range of LSM studies, complete removal of precipitates and

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Laser generation of a homogenous surface layer on base alloy has not been achieved

[145, 162, 178, 179]. Therefore, LSM reduces or eliminates the number of precipitates in the surface region and modifies the structure [161, 162, 179].

Also, it has been mentioned that the corrosion in LSM layer is associated with segregated regions within the melted layer [24, 166, 178]. Certain solidification structures, such as cellular and dendritic, are more likely to produce segregation that arises from low cooling rates. Therefore, reducing dendrite or cell spacing by high cooling rates can lead to improvement in the corrosion resistance [20-24,

180-188].

In general, corrosion behaviour may vary with the material composition used and the laser type, as well as laser processing parameters, particularly the cooling rate induced by LSM that affects the solidification microstructure and any element segregation within the modified layer [130]. Ideally, to improve the resistance of

LSM to corrosion, the melting layer has to be smooth, crack-and pore-free, and uniform, with complete precipitate removal and containing no sensitive regions from processing. The sensitive regions are always found in the overlap area between the two tracks of pulse movements. This is due to the re-heating of the laser area inducing further microstructure changes and causing micro-segregation within very small regions [190].

4.8 Excimer Laser for LSM

The name Excimer laser derives from the excited dimer molecules which are the lasing species. This laser is slightly different in that the gain is so strong that it does not need an oscillator. In KrF excimer laser, an electric discharge is

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Laser generated in a gas mixture of krypton (Kr) and fluorine (F). An excited dimer

KrF is formed which undergoes the stimulated emission process. It generates ultraviolet photons in brief pulse for each discharge of the condenser bank into the gas mixture. The generated pulses are usually very short, around 10 to 30 ns, but are very powerful. The energy range of the light pulses vary between milliJoules to Joules, average power up to several hundred watts, pulse repetition frequency rate between 1 to 100 Hz. Operational limits are determined by the rate of the high speed switching and the resonator length. The maximum gas flow velocity determines the maximum pulse frequency obtainable. The short wavelength ultraviolet (UV) light has high absorption, high spatial resolution, and high photon energy. Therefore, it can be focused to a small spot size and positioned with accuracy to process a wide range of materials including metals and alloys. The short pulse width and high peak power of an excimer laser can reduce the heat affected zone (HAZ) in most materials to a minimum. A further advantage of the excimer laser is the generation of a very high cooling rate, around 109-1011 K/s that can produce rapid solidification structure in the melted area without affecting the bulk properties of the material being used. Studies of the interaction of the KrF excimer laser with aluminium alloy surfaces, such as

Al-Cu-Mg alloys [24, 26, 201] indicated a melted layer with a uniform depth of about 20 µm. The Laser beam was homogenized producing a square spot area.

The laser-treated surfaces exhibited a higher corrosion resistance than as received material. Yue et al. [20, 21, 192] used an excimer laser to treat AA 7075 alloy.

The melted layer was about 8 µm thick. They also used the excimer laser with another alloy, AA 6013. A laser melted layer of about 5 µm thickness was formed at the surface [23]. The excimer laser was found to be effective in

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Laser improving the pitting corrosion resistance and reducing intergranular corrosion cracking onto the aluminium alloy surfaces. Ryan and Prangnell [166] indicated that the pulse laser surface melting treatment by an excimer laser produces a chemically more uniform layer on the top of AA 2024 friction stir weld material and thus can give more protection of the alloy surface from environmental corrosion attack. Overall, optimization of the operating parameters of the excimer laser, such as energy density, number of pulses per unit area, pulse repetition frequency and speed rate, will result in a highly refined surface layer of the alloy without affecting the bulk properties. This is the aim of using KrF excimer laser in the present work for treatment of the AA 7075-T6 alloy, which is explained in detail in subsequent sections of the thesis.

4.9 Safety of Using Laser

4.9.1 Laser Classification

Lasers are categorized into four classes from low to high risk. The higher the category, the more powerful and dangerous the laser is. Basically, all lasers used for material processing fall into highest categories. The Accessible Emission

Limit (AEL) is the maximum emission level permitted within a particular class, and is used to define the laser class. The Maximum Permissible Exposure (MPE) is the maximum level of the laser radiation to which a human can be exposed without adverse biological effects to the eye or skin. A pulsed laser is defined as one of that delivers energy in single or multiple pulses which are less than or equal to 0.25 s in duration.

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Laser

4.9.1.1 Class 1 Laser

Class 1 laser systems are considered incapable of producing damaging levels of radiation during operation since the output of the laser is very low or because of installed safeguards. The emitted radiation is enclosed and is not accessible. The lasers can be installed anywhere and no eye protection is required for workers.

Class 1M lasers present no risk for skin damage, but could result in damage to the eye if viewed with optical instruments such as binoculars.

4.9.1.2 Class 2 Laser

Class 2 laser systems emit visible radiation (370-750 nm) of low power (below 1 mW). They cannot damage a person accidently, but could damage the retina if viewed directly for more than 0.25 s. For continuous wave (CW) lasers, the power must not exceed 1 mW. It is the only class that applies exclusively to visible lasers operating in CW and pulsed modes.

4.9.1.3 Class 3 Laser

Class 3 lasers are subdivided into class 3R and class 3B. Class 3R lasers may emit visible or invisible radiation or both. If the beam is focused it presents a hazard. Outside the visible range, these lasers cannot exceed five times the class

1 AEL. In the visible range, they may not exceed five times class 2 AEL (5 mW).

Class 3B lasers mark the border between medium and high power devices. The beam of the visible class 3B laser may be safe to view on diffuse surface for less than 10 seconds at a distance greater than 130 mm. However, eye protection

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Laser should be worn. The power of visible class 3B laser must not exceed 0.5 W.

Exposure to 0.5 W for periods longer than 0.25 s is considered dangerous.

4.9.1.4 Class 4 Laser

Class 4 lasers operate above 0.5 W and produce radiant energy greater than 0.125

J within an exposure time of less than 0.25 s. Therefore, eye protection must be worn. Fire and skin hazards are also associated with this class type of laser.

4.9.2 Risk Assessment

Risk assessment is a management tool that allows the user to be familiar with the basis of both the specification of safety equipment and the content of safety procedures. After checking with the safety manager, the user can ensure that most important problems are recognized. The precautions taken to minimize risk from these potential problems are noted. Once the necessary precautions have been taken and noted, the risk assessment is kept for future reference.

4.9.3 Protection from Laser

4.9.3.1 Screening

All laser products require a protective housing that prevents access to radiation in excess of the limits of class 1. When class 3B or class 4 lasers are in use, screens, curtains, roller blinds, and window blocks are designed as passive guards to enclose these areas. Window blocks are acrylic filters that enable the processing zone to be viewed while stopping laser radiation from coming out. Shutters and

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Laser emission indicators (visible or audible) are required in all lasers above class 1 level. Operating controls must be located such that the user is not exposed to radiation. Safety and remote interlocks must be installed to prevent access to radiation above class 1 level. Also, systems such as 3B and 4 classes must have a key control that cannot be removed during operation.

4.9.3.2 Protective Eyewear

The eye can detect and focus light in the range of wavelength 350-750 nm that corresponding to the spectrum of colours in the rainbow. The eye also has an ocular focus region that includes wavelengths from the end of the red to 1400 nm. CW lasers can cause damage by thermal processes that overheat the tissue.

Pulsed lasers can cause damage from acoustic and vibrational shocks that rupture blood vessels in the retina. Absorption eyewear is the most commonly used type of eye protection. The laser beam is blocked by absorption of the relevant frequency. Glass or plastic absorption filters, with incorporated organic dye for the plastic filters, can easily be made specific to a particular wavelength or power, which reduces colour distortion and protects the eye from any damage.

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Laser

4.10 List of Figures

Figure 4-1: Energy level diagram for a three electron energy level of a laser and a pulse system [126]

Figure 4-2: Energy level diagram for a four electron energy level of a continuous (CW) laser system [126]

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Laser

Figure 4-3: Variation in total free energy of solid-liquid system as the size of the solid changes. The solid is stable nuclei with radius above the critical value r* [132]

Figure 4-4: Various growth structures that show planar (a), cellular (b), columnar dendrites (c), and equiaxed dendrites (d) [125]

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Laser

Figure 4-5: Schematic representing the solidification transformation from cell (a), cellular dendrites (b), columnar (c), and to columnar dendrite branches (d) [130]

Table 4-1: Examples of LSM types and their wavelengths [164]

Laser Wavelength Pulse length Operation

Excimer (KrF) 248 nm 13-25 ns Pulsed Excimer (XeCl) 308 nm 25 ns Pulsed Ruby 0.694 μm 0.2-5 ms Pulsed Nd:YAG 1.06 μm 200 ns Repetitively (Q-switched) CW-Nd:YAG 1.06 μm - Continuous

CO2 10.6 μm 50-200 ns Repetitively (Q-switched)

CW-CO2 10.6 μm - Continuous

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Laser

Figure 4-6: Schematic illustration of the laser beam interaction with the surface of a material [138]

Figure 4-7: Influence of the pulse duration in LSM of pure metals on the maximum depth of the melting layer [159]

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Laser

Figure 4-8: Cross-section of CW-Nd:YAG LSM on AA 2014-T6 alloy, following etching, that shows different regions of the solidification structure [161]

Figure 4-9: Modelling results of excimer LSM of AA 2024 alloy for a single pulse that show the solidification growth velocity as a function of position within the modified layer, including different layer thicknesses [166]

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Laser

Figure 4-10: Scanning electron micrographs of 3 kW CW-Nd:YAG LSM AA 2014-T6 alloy showing the variation in growth rate and dendrite spacing with position in modified layer [161]

Figure 4-11: Electron probe micro-analysis X-ray mapping of 3 kW CW- Nd:YAG LSM AA 2014-T451 alloy, show element segregation throughout the modified layer [161]

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Laser

Figure 4-12: Variation in cell spacing with melt depth for AA 2024-T3 alloy using a variety of pulsed treatment methods with scanning electron micrographs of the microstructure at different depths [166]

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Laser

Figure 4-13: High resolution EBSD map showing LSM layer morphology obtained with excimer laser treatment in Zr-free alloy (AA 2024) and Zr- containing alloys (AA 7150) [166]

Figure 4-14: Cross-section micrographs through excimer LSM treated layers formed over (a) AA 2096 alloy containing Zr and (b) AA 7075 Zr-free alloys, showing the layer grain structure with behaviour opposite to that found in Fig 4-13 [166].

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Chapter 5

Experimental Procedures

5. Introduction

This project aims to increase the resistance of AA 7075-T6 aluminium alloy to corrosion, without affecting its bulk properties, by laser surface melting (LSM), using an excimer laser. In order to optimise the corrosion improvement, the excimer laser was used with different numbers of pulses per unit area and varying energy density of the laser pulses.

After LSM, anodising was used with the aim of adding more protection to the laser-treated area. The corrosion susceptibility of the alloy in the as-received, laser-surface melted, and laser surface melting followed by anodising conditions was assessed by electrochemical measurements (potentiodynamic polarization), immersion testing in sodium chloride solution and exfoliation testing in EXCO solution in order to evaluate the pitting and intergranular corrosion resistances of the as-received alloy in comparison with the alloy treated by LSM with and without anodising. This Chapter describes the procedures of surface preparation, corrosion test evaluations and the characterisation techniques.

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Experimental Procedures

5.1 Substrate Materials

The material used in the experiments was aluminium alloy AA 7075-T6 sheet of

4 mm thickness supplied by Dolgarrog Aluminium Limited, North Wales, UK.

Table 5-1 shows the chemical composition in wt% of the alloy that was provided by the supplier.

Table 5-1: The chemical composition of the as-received AA 7075-T6 alloy

Element Zn Mg Cu Cr Fe Mn Si Ti

Wt% 5.83 2.45 1.67 0.21 0.24 0.08 0.12 0.03

5.2 Surface Preparation for LSM Treatment

The following preparation of the specimens was carried out before applying

LSM.

Specimens of dimension 20x20 mm, were cut from the as-received AA 7075-

T6 alloy sheet to provide L-ST faces for testing.

Specimens were then wet ground on the top surface using emery papers in the sequence 400, 600, 800 and 1200 grit. A new paper was used for grinding each specimen surface.

Specimens were then cleaned in an ultrasonic bath, using deionised water followed by ethanol for about 20 minutes in each case.

Specimens were then dried in air for about 24 h before applying the LSM treatment.

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Experimental Procedures

Laser surface melting (LSM) was carried out in air using a Lumonics IPEX-848 excimer laser, operated at 80 W with a wavelength of 248 nm. The conditions of

LSM are given in Chapter 7, which shows the results of studies carried out to optimize the laser parameters. After LSM, the specimens were cleaned with deionised water and dried in air for 24 h. An aluminium rod of 1 mm in diameter was attached for electrical connection for electrochemical tests. Lacquer (45- stopping-off) was used to mask regions not treated by LSM.

5.3 Anodising

The alloy after LSM, with an area of about 80 mm2, was anodised in order to further protect the laser surface melting layer from corrosion. In the anodising procedure, stirred 0.46 M sulphuric acid (H2SO4) solution at a temperature of 297

K was used, with anodising carried out potentiostatically at 12 V (SCE) for a total time of 240, 720 and 1200 s. A two-electrode cell was employed with the specimen as the anode and an aluminium sheet as the cathode. All the electrodes were electrically connected to a potentiostat (Solartron SI 1287 Electrochemical

Interface). During anodising, the current-density (I) vs. time (t) was monitored.

The selected anodising condition ensures the generation of a well-developed porous anodic oxide [198]. After anodising, sealing the anodic films was accomplished by immersion of the alloy specimens in boiling deionised water for

30 min. In addition, the as-received alloy was ground to 1200 SiC, then anodised and sealed under similar conditions to those used for the LSM specimens, to provide a reference condition for comparison purposes.

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Experimental Procedures

5.4 Corrosion Investigation

Three different types of corrosion tests have been employed. Each was repeated for a minimum of three times. Potentiodynamic polarization and immersion testing in sodium chloride solutions, and exfoliation testing in EXCO solution, were used, with comparison made between as-received and LSM alloy specimens.

5.4.1 Potentiodynamic Polarization Test

In potentiodynamic polarization measurements, each specimen of AA 7075-T6 alloy, in the as-received, laser-treated and laser-treated followed by anodising conditions, was masked with (45-stopping-off) lacquer, leaving about 20 mm2 unmasked. A second layer of masking was applied after 2 h. The two mask layers were left to dry in air for (24 h) at room temperature before commencement of the electrochemical measurements. Deaerated 0.1 M NaCl was used as the test solution. Deaeration was achieved by passing nitrogen in a closed cell containing the electrolyte for 2400 s before immersion of the specimens. Each specimen was then immersed in the solution for 40 minutes to stabilize the open circuit potential (OCP), with continued purging of the cell with nitrogen carried out throughout the test. The working electrode (specimen), counter electrode

(platinum), and reference electrode (saturated calomel electrode (SCE)) were connected to a potentiostat (1280 Solatron), which allowed the change of the potential of the specimen in a controlled manner and measurement of the current flow, with a current resolution of 0.1 μA, as a function of potential. Anodic polarization was applied from about 100 mV below the open circuit potential,

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Experimental Procedures with a potential scan rate of 0.1667 mV/s, according to the ASTM standard [78,

195].

5.4.2 Immersion Test

In immersion tests, each of the laser-treated and untreated substrates was immersed for 24 h in 0.1 M NaCl solution, open to air, at room temperature. This type of corrosion test is used to evaluate the susceptibility of the alloy to pitting and intergranular corrosion.

5.4.3 Exfoliation (EXCO) Test

The as-received and laser-treated AA 7075-T6 alloy was immersed in the EXCO test solution for various times, up to several hours. The test solution was prepared according to the ASTM G-34-01 standard test method for exfoliation corrosion susceptibility in 2XXX and 7XXX series aluminium alloys [196]. The test was carried out at room temperature and the solution contains 4.0 M NaCl,

0.5 M KNO3 and 0.1 M HNO3 [196].

5.5 Specimen Preparation for Metallographic Examination

In order to prepare the specimen after LSM for surface and microstructure examinations, the following steps were undertaken.

Specimens were cut in cross-section, using a Struers Accutom-5 disc cutting

machine, with a high speed coolant fluid to prevent temperature rises,

generated due to friction that could affect the LSM layer.

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Experimental Procedures

Cut specimens were mounted in conducting resin, containing graphite.

Mounted specimen surfaces were ground with SiC paper in the sequence 400,

600, 800, 1200, 2500 and 4000 grit.

Ground specimens were polished with 6, 3, and 1 µm diamond paste.

Specimens were washed with deionised water after each step of polishing,

with changing of the clothes for each paste.

After polishing, specimens were cleaned in an ultra-sonic bath containing

ethanol for 15 minutes.

Finally, specimens were cleaned with deionised water and dried in a stream of

air for several hours prior to investigation.

5.6 Field Energy Gun Scanning Electron Microscopy (FEG-

SEM)

Scanning electron microscopy (SEM) is capable of producing high-resolution images of the specimen surface. Secondary electron (SE) images were taken, in which most of the electrons collected by the detector originate from the near- surface region of the sample. Backscattered electron (BSE) images convey information on the atomic number from a greater depth in the sample. An EDX spectrometer was used to analyse elemental compositions and the distribution of the elements, the latter using the mapping analysis option. In this project, a

Philips XL30 FEG-SEM microscope, operated at 20 kV, was employed for examining the surface and cross-sections of specimens.

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Experimental Procedures

5.7 Transmission Electron Microscopy (TEM)

In transmission electron microscopy, a beam of electrons is transmitted through a specimen, then an image is formed, either on a florescent screen or layer of photographic film, or using a sensor such as a CCD camera [193]. The main use of the transmission electron microscope is to examine in sub-microscopic detail the structure and composition of specimen. The specimen for TEM investigation was prepared using a Leica Ultracut ultramicrotome. Sections of the specimen

(approx. 20 nm thick) were cut using a diamond knife and collected on nickel grids for examination. In the present work, JEOL 2000 FX II and FEI F30 Tecnai microscopes were employed. The microscopes were operated at 120 and 300 kV respectively. The Tecnai microscope was equipped with high angle annular dark field (HAADF) imaging and EDX analysis systems.

5.8 X-ray Diffraction (XRD)

X-ray diffraction is an analytical technique used to obtain information about crystallographic structure, chemical composition, and physical properties of materials. XRD is based on observing the scattered intensity of the X-ray beam of a selected wavelength (λ) hitting a sample as a function of the incident and scattered angles (θ) [194]. The intensity of the diffracted beam as a function of

2θ, the angle between the incident ray and the scattering planes, is recorded and peaks are observed at the Bragg’s law condition (λ = 2dsinθ). The results are compared with the powder diffraction index that contains the d-spacing between planes in the atomic lattice and the relative intensity values for a large number of elements and compounds in order to determine the phase structure. An X' Pert

125

Experimental Procedures

MPD Diffractometer was used in the present work, with low angle X-ray scattering (~2º), which allows to identify compounds in thin layer on the surface of the alloy.

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Chapter 6 Results

6. Microstructure and Corrosion of the As- Received AA 7075-T6 Alloy

6.1 Microstructural Examination

The FEG-SEM images of Fig. 6-1 show that relatively large constituent particles

(examples are circled in yellow) are present at the surface of the as-received AA

7075-T6 alloy specimens after mechanical grinding and polishing of the sheet surface up to 1 µm diamond paste. The particles are aligned in the rolling direction. Fig. 6-2 shows constituent particles on a polished alloy cross-section of the sheet normal to the direction of rolling. The particles range in size from about

1.5-12 µm. Examination at high magnification shows also the presence of cavities on the alloy surface (see arrow in Fig. 6-2 (a)), where particles have been either removed or dissolved during polishing of the specimens. Some cavities appear to contain remnants of second phase particles, as indicated by the arrow in

Fig 6-2 (c). The EDX line scans of Fig. 6-3 and 6-4, indicate the presence of aluminium, copper and iron in the large particles, while the cavities reveal the presence of magnesium and silicon. The arrows in Fig. 6-3 (c) show a region of relatively high copper, and the presence of iron, suggesting the presence of copper-iron-rich precipitates at these area locations, while the green arrows in the

Fig. 6-4 (c) show a region rich in iron and relatively low in copper. The EDX mapping analysis, shown in Figs. 6-5 and 6-6, confirms the presence of copper and iron in the large constituent particles, while magnesium and silicon are present at the locations of cavities. An iron-rich particle of low copper content is also evident. The EDX spectrum analysis, as shown in Fig. 6-7, indicates that

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Microstructure and Corrosion of the As-received AA 7075-T6 Alloy zinc and magnesium are distributed relatively uniformly in the alloy matrix. It is reported in the literature that the large constituent particles consist of Al7Cu2Fe,

Al2Cu (θ-phase) and Al3Fe (β-phase), while smaller particles containing magnesium and silicon consist of Mg2Si (β-phase) [4, 12, 32, 33,]. Magnesium and zinc precipitates along the grain boundaries as MgZn2 (η-phase) [3, 5, 7, 32,

33].

6.2 X-Ray Diffraction

Fig. 6-8 shows the results of the XRD measurements on the as-received AA

7075-T6 alloy. The low angle X-ray diffractometer analysis on the alloy surface indicates the presence of two types of second phases (Al7Cu2Fe and Mg2Si), evident from low intensity peaks. No peaks are detected from Al2Cu and Al3Fe, suggesting that such phases are minor constituents of the alloy. Al(Cu,Fe) second phase was also observed in the EDX line scan analysis of the large second phase particles, using FEG-SEM, confirming that most of the large intermetallic particles of the AA 7075-T6 alloy contain copper and iron. The presence of magnesium and silicon was also evident by EDX analysis in cavities on the alloy surface, where remnants of Mg2Si particles inside.

6.3 Corrosion Test Results

6.3.1 Potentiodynamic Polarization Test

Fig. 6-9 shows the anodic potentiodynamic polarization scans of the AA 7075-T6 alloy, in the as-received condition and in the polished condition, in 0.1 M NaCl

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Microstructure and Corrosion of the As-received AA 7075-T6 Alloy deaerated solution. The potential scan starts from point 1 and progresses in the positive direction until termination at point 2. At point A, the open circuit potential (OCP), the sum of the anodic and cathodic currents on the specimen surface of the working electrode is equal to zero. The corrosion potentials (Ecorr) of the as-received and as-polished alloy were recorded as -0.920 and -0.930 V vs

SCE, respectively. As the potential increases, the current increases and a low current density of about 1.0x10-7 A/cm2 for the as-received alloy and 3.3x10-7

A/cm2 for the polished alloy is recorded in the passive region that terminates at point B. The passive film then breaks down and the current density rapidly increases. This increase in the current density is due to pitting corrosion. The pitting potential (Epit) was recorded as about -0.710 V (SCE) for the as-received alloy and -0.695 V (SCE) for the as-polished alloy. Point C shows a second breakdown potential of the polished alloy. With further increase in potential, pits propagate on the alloy surface [27, 197].

6.3.2 Immersion Test

Fig. 6-10 shows scanning electron micrographs of the as-received AA 7075-T6 alloy surface after immersion in 0.1 M NaCl solution, open to the air, for an exposure time of 24 h. Figs. 6-10 (a) and (b) show pits adjacent to intermetallic particles on the alloy surface, while Figs. 6-10 (c) and (d) show pits free of second phase, possibly due to particles being removed by corrosion during immersion in the NaCl solution. Corrosion around particles may also lead to their undermining, with the particles then falling from the alloy surface. In the examples shown in Fig. 6-10, the size and depth of the pitting adjacent to intermetallic particles were slightly greater than those at sites where intermetallic

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Microstructure and Corrosion of the As-received AA 7075-T6 Alloy particles were absent (10 and 5 μm in Fig. 6-10 (b) and 5 and 4 μm in Fig. 6-10

(d)). Cavities may also result from dissolution of Mg2Si phase and possibly also from loss of particles during preparation of the cross-section.

6.3.3 Exfoliation Test

Fig. 6-11 shows the top surface of the as-received AA 7075-T6 alloy after exfoliation corrosion testing for 180 min. Fig. 6-11 (a) and (b) show the surface of the alloy before testing, while Fig. 6-11 (c) and (d) show the surface of the alloy after testing. The arrows in the Fig. 6-11 (c, d) show that the corrosion occurred due to the presence of second phase particles. Fig. 6-12 shows a cross- section of the as-received AA 7075-T6 alloy after exfoliation corrosion testing for different times (30, 60, 90 and 120 minutes). Fig. 6-12 (a) shows the presence of corrosion products after the shortest time of immersion (30 min). The corrosion extends laterally under the alloy surface. The depth of the attack increases with increasing the time of immersion, and results in relatively thick, cracked layers of corrosion product, with the presence of directional attack along the rolling direction as in Fig. 6-12 (c) and (d). The depth of the attack was measured to be 15 μm after 30 min, 75 μm after 60 min, 100 µm after 90 min and

140 μm after 120 min of immersion The occurrence of intergranular corrosion

(IGC) on the AA 7075-T6 alloy is generally accepted to be due to precipitates of

MgZn2-η phases on the grain boundaries [3, 5, 7, 10]. The anodic behaviour of these precipitates allows preferential dissolution along the grain boundary regions.

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Microstructure and Corrosion of the As-received AA 7075-T6 Alloy

6.4 Conclusions

The AA 7075-T6 alloy after surface and cross-section examinations revealed different sizes of intermetallic particles that are aligned in the rolling direction.

Large particles were copper-and iron-rich, and small particles were magnesium- and silicon-rich. XRD disclosed the presence of Mg2Si and Al7Cu2Fe phases. In potentiodynamic polarization tests in 0.1 M NaCl (deaerated condition), the corrosion potential (Ecorr) and the pitting potential (Epit) of the alloy were measured to be -0.930 V and -0.695 V (vs SCE) respectively, and the average passive range of potential was 0.235 V. After immersion testing in 0.1 M NaCl, the alloy exhibited pitting adjacent to second phase particles, associated with dissolution of the surrounding matrix. Exfoliation (EXCO) tests show that corrosion occurred within 30 min of immersion, and extended into the matrix with increasing time of immersion. The aggressiveness of the test solution with presence of second phase particles on the alloy surface allows intergranular corrosion (IGC) to propagate along selected grain boundaries.

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6.5 List of Figures

Figure 6-1: Scanning electron micrographs showing the surface of the AA 7075-T6 alloy at different magnification after grinding and polishing up to 1 µm diamond paste. (a) Secondary electrons. (b-d) Backscattered electrons.

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Figure 6-2: Scanning electron micrographs for the AA 7075-T6 alloy. The cross-section is normal to the rolling direction (a,c). Particles are shown in two different regions, while (b,d) are given the dimensions of the particles and cavities. (a,b) Secondary electrons. (b,d) Backscattered electrons.

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Microstructure and Corrosion of the As-received AA 7075-T6 Alloy

Figure 6-3: SEM-EDX line scan measurements that show the elements present in the constituent particles and precipitates in AA 7075-T6 alloy of the cross-section that is shown also in Fig.(8-2). The scans show Cu and Fe in the large particles, while Mg and Si are present at the cavities where residual particles are present. Zn appears to be distributed in both the alloy matrix and the particles.

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Microstructure and Corrosion of the As-received AA 7075-T6 Alloy

Figure 6-4 : SEM-EDX line scan measurement on intermetallic particles on the AA 7075-T6 alloy surface. (a,b) Scanning electron micrographs (backscattered electrons). (c) EDX line-scan analysis.

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Microstructure and Corrosion of the As-received AA 7075-T6 Alloy

Figure 6-5: SEM-EDX mapping analysis of AA 7075-T6 alloy. (a) Micrograph of the surface revealing the presence of intermetallic particles. (b) Aluminium. (c) Zinc. (d) Silicon. (e) Magnesium. (f) Manganese. (g) Iron. (h) Copper.

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Microstructure and Corrosion of the As-received AA 7075-T6 Alloy

Figure 6-6: SEM-EDX mapping analysis of AA 7075-T6 alloy at the region shown previously in Fig 6-3. (a) Micrograph of the surface (b) Aluminium. (c) Zinc. (d) Manganese. (e) Copper. (f) Iron. (g) Magnesium. (h) Silicon.

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Microstructure and Corrosion of the As-received AA 7075-T6 Alloy

Figure 6-7: Results of SEM-EDX spot analysis of different area of the matrix of AA 7075-T6 alloy. (a) and (c) Scanning electron micrographs showing the location of the analysis region. (b) and (d) EDX spectrum.

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Figure 6-8: Results of low angle X-ray diffraction measurments for the as- received AA 7075-T6 alloy.

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Microstructure and Corrosion of the As-received AA 7075-T6 Alloy

Figure 6-9: Anodic potentiodynamic polarization scan of the as-received and polished AA 7075-T6 alloys, in deaerated 0.1 M NaCl. The scan starts at point 1 and terminates at point 2. Point A shows the corrosion potential

(Ecorr), while point B shows the pitting potential (Epit). Point C shows a second breakdown potential on AR-polished alloy.

Table 6-1: The corrosion potential (Ecorr), pitting potential (Epit), passive current density (Ipass) and passive range potential (Epass) from Fig 6-8. Potentials are given with respect to SCE. 2 Sample Ecorr (V) Epit (V) Ipass (A/cm ) Epass (V)

AR -0.920 -0.710 1.0x10-7 0.210

AR-polished -0.930 -0.695 3.3x10-7 0.235

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Microstructure and Corrosion of the As-received AA 7075-T6 Alloy

Figure 6-10: Scanning electron micrographs (BSE) of cross-sections of AA 7075-T6 alloy after immersion in 0.1 M NaCl, open to the air, for 24 h. (a, c) Low magnification images. (b, d) Increased magnification images of (a, c). Images (a, b) reveal pitting corrosion adjacent to intermetallic particles on the alloy surface. Images (c, d) show pits free of second phases.

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Microstructure and Corrosion of the As-received AA 7075-T6 Alloy

Figure 6-11: Scanning electron micrographs of the top surface of the AA 7075-T6 alloy before and after immersion in EXCO solution for 180 minutes. (a, b) AR alloy before immersion. (c, d) AR alloy after immersion.

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Figure 6-12: Scanning electron micrographs (BSE) of cross-sections of the AA 7075-T6 alloy after immersion in EXCO solution for (a) 30, (b) 60, (c) 90 and (d) 120 minutes.

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Chapter 7

7. Laser Surface Melting (LSM) Operation Conditions

Laser surface melting was carried out in air using a Lumonics IPEX-848 excimer laser, operated at 80 W with a wavelength of 248 nm. Krypton fluoride (KrF) was the laser active medium. The pulse duration was fixed at 13 ns. Different numbers of laser pulses were used on the alloy surface, with an incident focused laser beam of either high or low energy density. The laser spot size was measured to be approximately 1.0 × 3.0 mm. Figures 7-1 and 7-2 show schematic illustrations of excimer LSM and the incident of laser pulses at the surface of the alloy specimen.

In the high pulse energy density condition, the fluence was 10 J/cm2. 10, 25 or 50 pulses were applied with pulse repetition frequencies (PRF) of 4, 10 and 20 Hz respectively. In order to treat a relatively large area of the alloy surface sufficient for corrosion evaluation, two laser tracks were used with an overlap ratio of 10% between them to generate a total treated area of 100 mm2. Table 7-1 summarises the laser operation conditions using the high pulse energy density.

In the low pulse energy density condition, the fluence was equal to 3.3 J/cm2. 70,

80 or 90 pulses per unit area were employed with pulse repetition frequencies of

25, 28 and 31 Hz respectively. Two tracks, with an overlap ratio of 10%, were

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LSM Operation Conditions used to generate a treated area of 100 mm2. Table 7-2 summarises the laser operation condition using the low pulse energy density.

The peak power of the laser pulse (Pp), is giving by;

E P [W] (1) p t where E is the pulse energy in Joules and t is the pulse duration in seconds.

The power density or irradiance which refers to the power of the pulse applied per unit area is given by;

P Irradiance [W / cm2] (2) A

The fluence refers to the total amount of pulse energy applied over the area of the laser spot beam.

E fluence [J / cm2] (3) A

During LSM, the interaction time of the laser pulse, Tin, can be calculated by the following equation;

Tin = length of laser pulse / scan rate velocity (4) where the length of laser pulse in mm and the scan rate velocity in mm / s. The number of pulses (NOP) produced is determined by the product of the laser pulse repetition frequency (PRF) and interaction time;

NOP = PRF x Tin (5)

From the above equations, there is a direct relationship between the number of pulses and the interaction time. For example, an increase in the number of pulses increases the interaction time, and oppositely, a decrease in the number of pulses

145

LSM Operation Conditions decreases the interaction time. Scanning electron micrographs of cross-sections of the LSM AA 7075-T6 alloy treated with a high pulse energy density (10

J/cm2) shown in Fig. 7-3, disclosed a large number of pores and cracks in the melted layers, which were more numerous with increase in the number of pulses.

The thickness of the melted layer increased with increasing of the number of pulses from about 5 µm to about 15 µm between 10 and 50 pulses. The high fluence of the laser pulses appeared to cause vaporisation of the alloy, and therefore, porosity in the melted layer. Treating the alloy surface with excimer

LSM using a low pulse energy density needed a relatively high number of pulses

(i.e 70, 80 or 90) to achieve diffusion and mixing in the melt pool area of the alloy surface. In contrast, treatments using a low fluence and a low number of pulses (≤ 50 pulses) were insufficient to melt the large intermetallic particles at the surface of the alloy; some of these particles (highlighted by yellow circles) still remained in the melted layer, as evident in Fig. 7-4. Scanning electron micrographs of cross-sections of the alloy treated with a low pulse fluence (3.3

J/cm2), using 70, 80 and 90 pulses, showed a relatively low number of pores and a relatively uniform melted layer was produced at the surface of the alloy, especially when using 80 pulses, as illustrated in Fig. 7-5, although some porosity still remained. Fig. 7-6 demonstrated the influence of the number of pulses of the excimer laser using both high and low energy densities on the average thickness of the laser-melted layer and the size of the largest pores in the melted layers, as measured from the scanning electron micrographs of the cross- sections. It is shown from the graph that the alloy treated with 80 pulses at low energy density has the lowest ratio of the pore size to the average thickness.

After LSM, each of the laser-treated areas of the alloy produced using high and

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LSM Operation Conditions low laser pulse energy densities was examined using potentiodynamic polarisation tests in deaerated 0.1 M NaCl solution. The tests reveal the corrosion potential, the pitting potential, the passive potential range and the passive current density, and therefore, are a useful tool for corrosion evaluation. The AA 7075-

T6 alloy treated with a high pulse fluence, using 10, 25 and 50 pulses, showed a reduction in the corrosion potential by about -181, -230 and -220 mV in comparison with the corrosion potential of the as-received alloy, as shown in Fig.

7-7. This reduction in the corrosion potential could be attributed to a decrease in the cathodic reactivity of the alloy surface after LSM by dissolution of the large cathodic particles. Further, there was a reduction by one order of magnitude in the passive current density for 10 and 25 pulses in comparison with the as- received alloy, from 1.0x10-7 to 4.6x10-8and 7.3x10-8 A/cm2 respectively, and less than one order of magnitude for 50 pulses from 1.0x10-7 to 1.6x10-7 A/cm2.

However, the pitting potential is not affected greatly, being about -670, -680 and

-695 mV (SCE) for 10, 25 and 50 pulses and -710 mV (SCE) for the as-received alloy. Potentiodynamic polarisation tests, as presented in Fig 7-8, revealed a further reduction in the passive current density of the LSM alloy treated with low fluence, using 70, 80 and 90 pulses (3.4x10-9, 5.7x10-9 and 5.1x10-9 A/cm2) respectively in comparison with the as-received alloy as well as the LSM alloy treated with high fluence. Table 7-4 presents the measured values of the corrosion potentials (Ecorr), pitting potentials (Epit), passive current densities (Ipass) and passive potential ranges (Epass) from polarisation tests for specimens treated with both high and low fluence; the potentials are expressed with respect to

(SCE). The alloy treated using 80 pulses and low fluence (3.3 J/cm2) shows the

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LSM Operation Conditions greatest reduction in the passive current density in comparison with the as- received alloy and the alloy treated with high fluence.

From the previous conditions, it is important to select the appropriate parameter values for LSM treatment in order to produce enough melting to dissolve constituent particles at the substrate surface, since the aim of this project is to use

LSM to eliminate those large particles, which have been reported to increase the corrosion in environments containing chloride ions [4, 12, 32, 33, 45-47].

From the above observations of the cross-sections and the results of potentiodynamic polarisation tests of the excimer laser-treated on AA 7075-T6 alloys, it is more appropriate to use the low pulse energy density, rather than the high pulse energy density, to treat the surface of the AA 7075-T6 alloy for further improvement in corrosion resistance. The condition of low fluence with

80 pulses was selected for further evaluation of the corrosion properties of the laser-treated alloy. The results are presented in Chapter 8.

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7.1 List of Figures

Table 7-1: Conditions of laser operation (high pulse energy density)

NOP Fluence J/cm² 10 10 25 10 50 10

NOP: Number of pulses per unit area

Table 7-2: Conditions of laser operation (low pulse energy density)

NOP Fluence J/cm² 70 3.3 80 3.3 90 3.3

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LSM Operation Conditions

Figure 7-1: Schematic illustration of the excimer laser that shows the focusing of the laser beam onto the surface of the alloy sample that is placed on a motorised holder allowing movements in the X and Y directions.

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LSM Operation Conditions

Figure 7-2: Schematic diagram representing the laser scanning pattern and the incidence of the laser pulse on specimen surface.

151

LSM Operation Conditions

Figure 7-3: Scanning electron micrographs (backscattered electron) of cross-sections of the LSM AA 7075-T6 alloy after treatment with high pulse energy density (10 J/cm2), using 10 (a), 25 (b), and 50 (c) pulses.

152

LSM Operation Conditions

Figure 7-4: Scanning electron micrographs (secondary electron) of the top surface of LSM AA 7075-T6 alloy, after treatment with low pulse energy density. (a) 10 pulses, (b) 25 pulses, (c) 40 pulses and (d) 50 pulses

153

LSM Operation Conditions

Figure 7-5: Scanning electron micrographs (BSE) of cross-sections of the LSM AA 7075-T6 alloys treated with low pulse energy density (3.3 J/cm2), using 70 (a), 80 (b) and 90 (c) pulses.

154

LSM Operation Conditions

Figure 7-6: Average thickness of the laser melted layer and the largest pore size in LSM AA 7075-T6 alloy vs. the number of pulses (NOP), as determined from scanning electron micrographs of cross-sections for laser- treatments, using high and low energy densities.

155

LSM Operation Conditions

Figure 7-7: Potentiodynamic polarisation scans in deaerated 0.1 M NaCl solution of the as-received (AR) AA 7075-T6 alloy and after LSM using a high energy density (10 J/cm2), with 10, 25 and 50 pulses.

156

LSM Operation Conditions

Figure 7-8: Potentiodynamic polarisation scans in deaerated 0.1 M NaCl condition of the as-received (AR) AA 7075-T6 alloy and after LSM with high energy density (10 J/cm2), using 10, 25 and 50 pulses, and with low energy density (3.3 J/cm2), using 70, 80 and 90 pulses.

157

LSM Operation Conditions

Table 7-3: Corrosion potential (Ecorr), pitting potential (Epit), passive current density (Ipass) and passive range potential (Epass) for both as-received and laser treated alloys in Fig. 7-8. Potentials are given with respect to SCE.

2 Sample Ecorr (V) Epit (V) Ipass (A/cm ) Epass (V)

AR -0.920 -0.710 1.0x10-7 0.210

10 P -1.111 -0.670 4.6x10-8 0.441

25 P -1.160 -0.680 7.3x10-8 0.480

50 P -1.150 -0.695 1.6x10-7 0.454

70 P -0.920 -0.700 3.4x10-9 0.216

80 P -1.030 -0.680 5.7x10-9 0.329

90 P -0.940 -0.660 5.1x10-9 0.280

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Chapter 8 Microstructure and Corrosion of the LSM AA 7075-T6 Alloy

8. Microstructural Examination

It is evident from scanning electron micrographs, shown in Fig. 8-1, that the top surface of the AA 7075-T6 alloy after excimer LSM treatment, using 80 pulses with an energy density of 3.3 J/cm2, has a wavy morphology. The micrographs display the two tracks of the laser and the region of overlap, which is 300 µm wide. No large second phase particles are evident on the melted area, as can be seen by comparison with an untreated area of the alloy shown in Fig. 8-2, (b) and

(c). Fig. 8-3 shows another area of the alloy after LSM treatment. The Fig. discloses the overlap and the re-heating regions. The overlap region forms part of second laser track, while the re-heating region lies within the first laser track, as indicated in (a) and (b). The width of the re-heating region is determined by the heat-transfer properties of the alloy; modelling studies are required to evaluate the temperature profile in the reheated zone. Fig. 8-3 (c) shows a higher magnification image in the overlap region. No cracking is evident in the overlap region. Fig. 8-3 (d) and (e) show higher magnification images of the re-heating area. Micro-cracks are evident in this area, which are marked with red arrows in

Fig. 8-3 (e). EDX spectra of Figs. 8-4 and 8-5 reveal that aluminium, zinc, copper, and oxygen are present in the LSM layer. The EDX spot focused on very small area of LSM alloy and because of low content of magnesium, this element may not show in the spectrum analysis of the selected regions. The EDX analysis shows a high of oxygen (12.4 and 10.8 wt.%), with copper and zinc contents appearing to be similar to those in the bulk alloy (1.6 and 1.7 wt.% for Cu and

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Microstructure and Corrosion of the LSM AA 7075-T6 Alloy

5.6 and 5.8 wt.% for Zn respectively). The high oxygen content suggests the presence of an oxide film on the laser-melted surface. The percentages of the elements that resulted from the EDX analysis are recorded in the Tables that accompany both Figs. Micrographs at relatively high magnification of cross- sections of the LSM area, shown in Fig. 8-6, disclose a relatively rough surface with a melted layer of about 3 μm thickness. The rapid solidification of the molten materials produced by excimer laser pulses causes the melted layer to be frozen with a wavy surface. As shown in Chapter 7, an increase of the number of laser pulses per unit area, resulted in a rougher melted layer, which is attributed to the turbulent flow of molten fluid from the centre of the melt pool to the adjacent areas [180]. Fig. 8-7 shows details of the roughness; the peak-to-trough distance is approximately 1.6 µm.

Scanning electron micrographs of cross-sections of the melted layer and the alloy matrix are shown in Fig. 8-8. They reveal a thin melted layer on top of the alloy, and a partially-melted second phase particle. The thickness of the melted layer is about 2.5 μm in the particular region of the specimen, compared with about 3.0

µm in Fig. 8-6. The arrows in the scanning electron micrograph of Fig. 8-8 (d), show thin solute-rich bands in the melted layer, particularly above the partially- melted particle. Fig. 8-9 shows a cross-section and EDX mapping analysis of a region of the melted layer at a location of a partially-dissolved second phase particle. The EDX maps indicate that the particle consists of Al-Cu-Fe phase.

The backscattered scanning electron micrograph of Fig. 8-10 (a) shows local enrichment of copper in the melted layer, above the partially-melted particle.

Cross-sections of another region of the melted layer are shown in Fig. 8-10. The

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Microstructure and Corrosion of the LSM AA 7075-T6 Alloy melted layer above both the matrix and partially-melted second phase particle is revealed in the backscattered electron images with a thickness between 2 to 3

µm. It appears that the pulses of the excimer LSM melt the intermetallic particles and redistribute the alloying elements within the solid solution formed on cooling of the melted layer. Porosity on the other hand is also evident in the melting layer, with pores of diameter up to ~ 1 µm. Porosity may arise due to release of hydrogen and evaporation of magnesium from the bulk alloy, or from absorption of gas in the atmosphere during LSM treatment [180,166-167].

Fig. 8-11 shows transmission electron micrographs of an ultramicrotomed cross- section of the LSM layer and the alloy matrix. It appears that the melted layer in the sectioned regions contains a number of pores, indicated by white arrows. The thickness of the melted layer is between 2.8 to 3.0 μm, in agreement with findings from SEM. Fig. 8-11 (c) shows thin bands within melted layer, which are considered later; these bands contain fine precipitates. The bands are separated by distances of about 0.3 to 0.4 μm. Fig. 8-12 shows another area of the ultramicrotomed cross-section of the LSM layer and alloy matrix. Fig. 8-12

(a) reveals pores, indicated by white arrows, in the LSM layer. The average thickness of the LSM layer is about 3.0 μm. Cavities and scratches are shown in the laser treated layer, which are due to damage produced by sectioning, as indicated with black arrows. Fig. 8-12 (b) shows a higher magnification image of the cross-section micrograph of Fig. 8-12 (a). It reveals the region between the

LSM layer and alloy matrix (interface), as indicated with red arrows. The vertical striations in images are chatter lines which are formed during ultramicrotomy due to the hardness of the melted layer. For more detailed examination, EDX

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Microstructure and Corrosion of the LSM AA 7075-T6 Alloy spot analyses were taken at five locations in the cross-section, as indicated in Fig.

8-13 (a). EDX analysis at the locations of Figs. 8-13 (b) and (c) shows enrichment of zinc and magnesium at the interface region in comparison with the matrix (Fig. 8.13 (f)). Fig. 8-13 (d) shows enrichment of zinc and magnesium also in the precipitates within the LSM region The EDX analysis of Fig. 8-13 (e) shows zinc-magnesium-rich and chromium-rich precipitates in the matrix. Fig. 8-

14 (a) shows a high magnification image of an ultramicrotomed cross-section of the laser-melted layer in the high angle annular dark field mode. It is evident that distances of approximately 0.3 to 0.4 µm separated the precipitate bands, as also evident in Fig. 8-11 (c). Fig. 8-14 (c) shows the results of EDX analysis of a band of Fig. 8-14 (b). The band reveals enrichment of zinc and magnesium.

8.1 X-ray Diffraction

Fig. 8-15 shows the results of low angle X-ray diffraction of the AA 7075-T6 alloy after LSM. It reveals peaks due to the Al of the matrix and the laser-melted layer. No significantly high peaks from large second phases were observed. In comparison with the XRD of the LSM alloy, the as-received alloy (Fig. 8-16) shows relatively large peaks due to second phase particles.

8.2 Corrosion Test Results after LSM

8.2.1 Potentiodynamic Polarization Test

Fig. 8-17 shows potentiodynamic polarization curves for the laser-treated and the as-received AA 7075-T6 alloy in deaerated 0.1 M NaCl solution. The curves

162

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy reveal that there is a reduction by two orders in magnitude in the passive current density for the LSM specimen in comparison with the as-received alloy, from

1.0x10-7 to 5.7x10-9 A/cm2. At the same time, there is a significant increase in the passive range of the LSM alloy. Table 8-1 shows the measured values of the corrosion potential (Ecorr), pitting potential (Epit), passive current density (Ipass) and passive potential range (Epass) for the both laser-treated and the as-received alloy. Laser melting of the alloy surface results in a reduction of the corrosion potential by about -100 mV. However, the pitting potential is not significantly affected, being about -685 mV (SCE) for the laser-treated alloy and -710 mV

(SCE) for the as-received alloy.

8.2.2 Immersion Test

Fig. 8-18 shows scanning electron micrographs of the LSM AA 7075-T6 alloy after immersion in 0.1 M NaCl, open to the air, for 24 h. The cross-sections, examined at a range of magnifications, reveal no evidence of corrosion products after LSM. The thickness of the laser-melted layer was similar to that before the immersion test. The findings indicate an improvement in corrosion resistance of the alloy after laser surface melting treatment in comparison with as-received alloy. The latter showed pitting corrosion, (see Fig. 6-10), due to the presence of second phase particles on the alloy surface.

8.2.3 Exfoliation Test

Fig. 8-19 shows scanning electron micrographs of cross-sections of the LSM AA

7075-T6 alloy tested in naturally-aerated EXCO solution for times of 30, 60,

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Microstructure and Corrosion of the LSM AA 7075-T6 Alloy

120, 150 and 180 min. The cross-sections of the laser treated (LT) alloys, as shown in Fig. 8-19 (a-d), reveal no corrosion product after immersion for up to

150 min. However, with increasing time of immersion, up to 180 min, the alloy shows corrosion underneath the laser-treated layer (Fig. 8-19 (e)).

Figs. 8-20 and 8-21 show higher magnification scanning electron micrographs of cross-sections of the as-received and laser-treated alloy after immersion testing.

It is evident that no significant corrosion attack of the laser-treated alloy occurs within exposure times of 30, 60, 120 and 150 min, in comparison with the alloy without laser surface treatment. On the other hand, Fig. 8-22 shows scanning electron micrographs of cross-sections of the laser-treated and untreated alloys after immersion for 180 min. The high magnification micrograph of Fig. 8-22 (c) reveals severe IGC of the as-received alloy. However, corrosion was limited in the laser-treated alloy and located within the melted layer, which caused the layer to delaminate from the alloy matrix, as shown in Fig. 8-22 (d). The corroding solution may find a way to attack the base of the melted layer due to the presence of defects and/or microcracks that allow access of the solution to the substrate region. Cracks were observed in the re-heated region of the LSM layer, as shown in the scanning electron micrograph of Fig. 8-23. However, LSM reduces the severe exfoliation corrosion attack that occurs with the as-received alloy in a similar test condition.

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Microstructure and Corrosion of the LSM AA 7075-T6 Alloy

8.3 Conclusions

AA 7075-T6 alloy after excimer LSM shows a relatively thin melted layer with a wave-like surface. The thickness of the LSM layer produced under the selected

LSM conditions was 2.5 to 3.5 μm. Rapid solidification of the melted metal caused dissolution of constituent particles and generated a relatively uniform melted layer that was free from large second phases. After LSM, the alloy exhibits a reduction in the passive current density by two orders in magnitude in comparison with the untreated alloy when examined under similar conditions of electrochemical polarization in sodium chloride solution. The LSM alloy shows no significant attack after immersion in 0.1 M sodium chloride solution, open to air, for 24 h. In comparison, the as-received alloy shows pitting corrosion in the same environment. The results from exfoliation corrosion tests show a large reduction in the corrosion damage of alloy following LSM. However, the presence of cracks in the re-heated region caused by the influence of the second laser track on the first laser track enables electrolyte to penetrate to the matrix.

The presence of a zinc and magnesium rich bands in the melted layer, and between the LSM layer and the bulk alloy, facilitates corrosion along these bands which causes delamination of the LSM layer from the alloy.

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8.4 List of Figures

166

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy

Figure 8-2: Scanning electron micrographs (BSE) of the AA 7075-T5 alloy after LSM. (a) Alloy with and without treatment by LSM. (b) Untreated alloy, revealing the presence of intermetallic particles. (c) Alloy after LSM treatment.

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Microstructure and Corrosion of the LSM AA 7075-T6 Alloy

Figure 8-3: Scanning electron micrographs (SE) of the top surface of the AA 7075-T6 alloy after LSM using 80 pulses. (a) and (b) show laser tracks, re- heating and overlap regions. (c) Higher magnification in overlap region. (d) and (e) Higher magnification in re-heating area.

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Microstructure and Corrosion of the LSM AA 7075-T6 Alloy

Figure 8-4: Results of SEM-EDX spot analysis of the AA 7075-T6 alloy after LSM. (a) EDX spectrum. (b) Scanning electron micrograph showing the location of the analysis region. (c) Elements in weight and atomic percentages.

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Microstructure and Corrosion of the LSM AA 7075-T6 Alloy

Figure 8-5: Results of SEM-EDX spot analysis of a different area of the LSM AA 7075-T6 alloy from that shown in Fig 8-4. (a) EDX spectrum. (b) Scanning electron micrograph showing the location of the analysis region. (c) Elements in weight and atomic percentages.

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Microstructure and Corrosion of the LSM AA 7075-T6 Alloy

Figure 8-6: Scanning electron micrographs of cross-sections of the AA 7075- T6 alloy after LSM using 80 pulses. (a) and (c) Secondary electrons. (b) and (d) Backscattered electrons.

171

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy

Figure 8-7: High magnification scanning electron micrographs showing a cross-section of AA 7075-T6 alloy after LSM using 80 pulses.

172

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy

Figure 8-8: Scanning electron micrographs (BSE) of a cross-section of the AA7075-T6 alloy after LSM, with the presence of a thin melted layer on the alloy surface indicated in (a) and (b). Higher magnification images in (c) and (d) show the thinner melted layer produced above a large second phase particle that was partially dissolved after LSM treatment.

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Microstructure and Corrosion of the LSM AA 7075-T6 Alloy

Figure 8-9: SEM-EDX mapping analysis of Fig 8-8 (d) showing the LSM layer and the alloy matrix in (a), and the imaging of Al (b), Zn (c), Cu (d), Fe (e), and Mg (f).

Figure 8-10: Scanning electron micrographs (BSE) of cross-sections of the AA 7075-T6 alloy after LSM using 80 pulses, showing the laser melted layer in more detail, disclosing porosity in the layer with pores of dimension about 1 µm.

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Microstructure and Corrosion of the LSM AA 7075-T6 Alloy

Figure 8-11: Transmission electron micrographs of ultramicrotomed cross- sections of the AA 7075-T6 alloy after LSM using 80 pulses. Micrographs (a) and (b) reveal the LSM layer and alloy matrix with pores indicated by arrows in laser-treated layer. Micrograph (c) shows bands in LSM layer, which are due to the presence of precipitates. The bands are separated by distances of 0.3 to 0.4 μm.

175

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy

Figure 8-12: Transmission electron micrographs of an ultramicrotomed cross-section of the AA 7075-T6 alloy after LSM using 80 pulses. (b) Higher magnification of (a). The Fig. shows the interface region, as indicated by arrows in red, between the LSM layer and the alloy matrix. The arrows in white and black show pores and scratches in the LSM layer.

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Microstructure and Corrosion of the LSM AA 7075-T6 Alloy

Figure 8-13: TEM-EDX spot analysis of the micrograph cross-section of Fig. 8-12 (b); the 5 locations of EDX analysis are indicated in (a).

177

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy

Figure 8-14: High angle annular dark field (HAADF) images of an ultramicrotomed cross-section of the AA 7075-T6 alloy after LSM using 80 pulses. (b) Higher magnification of (a). (c) Results of of EDX analysis of (b).

178

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy

Figure 8-15: Low angle XRD analysis of the AA 7075-T6 alloy after LSM using 80 pulses.

179

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy

Figure 8-16: Low angle XRD analysis for both as-received and LSM AA 7075-T6 alloy treated with 70, 80, and 90 pulses.

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Microstructure and Corrosion of the LSM AA 7075-T6 Alloy

Figure 8-17: Potentiodynamic polarization scans in deaerated 0.1 M NaCl for AA 7075-T6 alloy and after LSM using 80 pulses.

Table 8-1: Corrosion potential (Ecorr), pitting potential (Epit), passive current density (Ipass) and passive range potential (Epass) for both as-received and laser treated AA 7075-T6 alloy determined from the potentiodynamic polarisation curves of Fig. 8-17. Potentials are given with respect to SCE. 2 Sample Ecorr (V) Epit (V) Ipass (A/cm ) Epass (V)

AR -0.920 -0.710 1.0x10-7 0.210

LT -1.030 -0.685 5.7x10-9 0.349

181

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy

182

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy

Figure 8-19: Scanning electron micrographs (BSE) of the AA 7075-T6 alloy after LSM using 80 pulses and immersion in naturally-aerated EXCO solution for (a) 30, (b) 60, (c) 120, (d) 150 and (e) 180 min. The cross-sections of the laser surface treatment alloys (a-d) reveal no corrosion products after immersion for 150 min. With increasing time of immersion, corrosion causes delamination of the LSM layer, as evident in (e).

183

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy

Figure 8-20: Scanning electron micrographs of cross-sections of the AA 7075-T6 alloy in the as-received condition and after laser-surface treatment (LT) using 80 pulses, following immersion in (EXCO) solution for 30 min (a, b) and 60 min (c, d).

184

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy

Figure 8-21: Scanning electron micrographs of cross-sections of the AA 7075-T6 alloy in the as-received condition and after laser surface treatment alloys using 80 pulses and immersion in EXCO solution for 120 min (a, b) and 150 min (c, d).

185

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy

Figure 8-22: Scanning electron micrographs of cross-sections of the AA 7075-T6 alloy in the as-received condition (a, c) and after laser surface treatment (b, d) using 80 pulses immersion in EXCO solution for 180 min.

186

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy

Figure 8-23: Scanning electron micrographs of LSM AA 7075-T6 revealing micro-cracks in the re-heating area of the 1st laser track as in (a) and (b), while (c) and (d) show micrographs of cross-sections after exfoliation testing for 180 min.

187

Chapter 9 Microstructure and Corrosion of the LSM AA 7075-T6 Alloy after Anodising

9. Introduction

Anodising, as mentioned in Chapter 3, is an electrochemical procedure that is used to increase the thickness of the protective oxide layer on the surface of aluminium and aluminium alloys in order to improve their corrosion resistances.

In this section of the thesis, results are presented for anodising of laser surface melted AA 7075-T6 alloy. The anodising was carried out at 12 V in stirred 0.46

M sulphuric acid (H2SO4) electrolyte at a temperature of 295 K for times of 240,

720, and 1200 s. The aim was to protect the laser surface melted layer from delamination during exfoliation testing. Specimens are examined by SEM and

TEM, with the corrosion resistance of the laser-treated and anodised alloy evaluated in sodium chloride solution using potentiodynamic polarization and exfoliation immersion tests.

9.1 LSM and Anodising

The current density / time behaviour of the LSM specimen, as well as the as- received alloy, was recorded during anodising at 12 V (SCE) in 0.46 M H2SO4, as shown in Fig. 9-1. The specimens, with an area of approximately 80 mm2, show an initial surge in the current density followed by a decrease to a minimum value and an increase to a steady current density. The initial surge of the current density is related to the growth of a barrier film. The steady current density corresponds to the growth of a porous film [26]. A significant reduction of the

188

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy after Anodising steady current density, by about 50 %, was evident after LSM in comparison with the as-received AA 7075-T6 alloy following anodising for times of 240, 720 and 1200 s. The values of the current density for the as-received and laser-treated alloys are approximately 1.7x10-3 and 7.7x10-4 A/cm2 respectively. Following anodising the film was sealed by immersion in deionised boiling water for 30 min.

9.2 Surface and Cross Section Results

Figure 9-2 shows scanning electron micrographs of the top surface of LSM AA

7075-T6 alloy after anodising in 0.46 M H2SO4 for 1200 s. Fig. 9-2 (a) shows the overlap area in the second laser-track and the re-heating area in the first laser- track. Figs. 9-2 (b) and (c) show higher magnification images of the re-heating area and the overlap area, with the presence of micro-cracks evident in the re- heated region of the LSM layer, as indicated by arrows in Fig. 9-2 (b). Fig. 9-3

(a) shows a scanning electron micrograph and the results of EDX analysis of the re-heating area after anodising. Oxygen, aluminium, magnesium and zinc are indicated by EDX analysis, as shown in Fig. 9-3 (b), with oxygen at about 27 wt% , magnesium at about 1.6 wt% and zinc at about 4.5 wt% (Fig. 9-3 (c)). Fig.

9-4 (a) shows a scanning electron micrograph with an EDX spectrum taken on the overlap region. Fig. 9-4 (b) shows the EDX analysis, with oxygen at about

26.1 wt%, magnesium at about 1.4 wt% and zinc at about 4.5 wt% (Fig. 9-4 (c)) .

Fig. 9-5 (a) shows a scanning electron micrograph with an EDX spectrum taken on the middle of the laser-track. Fig. 9-5 (b) shows the EDX analysis, with oxygen at about 39.8 wt%, magnesium at about 1.6 wt% and zinc at about 3.9 wt% (Fig. 9-5 (c)). From the EDX measurements, the anodising of the LSM AA

189

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy after Anodising

7075-T6 alloy in 0.46 M H2SO4 for 1200 s has increased the level of oxygen, and slightly reduced the content of magnesium and zinc in the LSM layer in comparison with LSM alloy without anodising (see Figs. 8-4 and 8-5). Fig. 9-6

(a) shows a scanning electron micrograph of a cross-section of the LSM AA

7075-T6 alloy after anodising in sulphuric acid for 1200 s. Fig 9-6 (b) shows a high magnification image of Fig. 9-6 (a). The cross-sections reveal a thin anodic film of about 313 to 332 nm thickness above the laser-melted layer. Fig. 9-7 shows a transmission electron micrograph of an ultramicrotomed cross-section of the anodised LSM AA 7075-T6 alloy. The micrographs in Fig. 9-7 (a) and (b) disclose a bubble-like texture of the anodic film, which has a relatively uniform thickness of about 380 nm over the laser-treated layer. For the purpose of comparison, Fig. 9-8 shows high resolution scanning electron micrographs of the as-received alloy cross-sections after anodising in H2SO4 for 1200 s. The micrographs in Figs. 9-8 (a), (b) and (c) reveal a thin anodic film over the as- received alloy with a thickness of about 200 nm. Fig 9-8 (d) shows anodic film over a large intermetallic particle with a thickness of about 2 µm. The yellow arrows in the Figure disclosed cracks in the anodic film formation.

9.3 Corrosion Test Results

9.3.1 Potentiodynamic Polarization Test

A potentiodynamic polarization scan of the LSM alloy after anodising is shown in Fig. 9-9, revealing a reduction in the passive current density (7.4x10-10 A/cm2) and increase in the corrosion potential (-0.780 V) (SCE) in comparison with the as-received alloy (1.0x10-7A/cm2, and -0.930 V) (SCE), AR and anodised alloy

190

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy after Anodising

(1.2x10-7 A/cm2, and -0.790 V) (SCE) and LSM alloy (5.7x10-9 A/cm2, and -

0.103 V) (SCE).The pitting potential was not significantly affected, being around

-0.670 to -0.700 V (SCE) The reduced passive current density could be attributed to the anodic film formation which increases the thickness of the oxide layer in comparison with those on the as-received alloy before and following anodising and after LSM. Table 9-1 shows the measurement values of the Ecorr, Epit, and

Ipass determined from Fig. 9-9.

9.3.2 Exfoliation Test

Figures 9-10 (a) and (b) showed scanning electron micrograph of the top-surface of the anodised as-received alloy after immersion in naturally aerated EXCO solution for 180 min, with the presence of area of corrosion and cracks evident.

The high magnification scanning electron micrograph, of Fig. 9-10 (c) reveals cracks of width about 1 to 1.65 µm. Fig. 9-11 presents scanning electron micrographs of cross-sections of the anodised alloy, disclosing the presence of exfoliation corrosion in the as-received alloy matrix. The depth of the exfoliation attack is about 65 µm. Fig. 9-12 shows scanning electron micrographs of the top- surface of the anodised LSM alloy after immersion in naturally aerated EXCO solution for 180 min. Fig. 9-12 (a) discloses the re-heating area and the overlap area. The red arrows in the first-track of the laser-treated layer show corrosion products. Figs. 9-12 (b) and (c) show scanning electron micrographs in the middle of the anodised and laser treated layer after immersion testing, where corrosion products are marked with red rings. Fig. 9-13 shows scanning electron micrographs of cross-sections of the LSM AA 7075-T6 alloy following anodising, after exfoliation testing for 180 min. Fig. 9-13 (b) is higher

191

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy after Anodising magnification of Fig. 9-13 (a) and Fig. 9-13 (d) is a higher magnification of Fig.

9-13 (c). The thickness of the laser-melted layer is about 2.8 to 3.1 µm. Pores with less than 0.4 µm diameter are shown with arrows in Fig. 9-13 (c) and (d).

No evidence of delamination of the laser-melted layer from the alloy surface occurred after corrosion testing for 180 min.

9.4 Conclusions

Anodising of LSM AA 7075-T6 alloy in sulphuric acid results in improvement in the corrosion performance of the alloy. Potentiodynamic polarization testing revealed a decrease by one order of magnitude of the passive current density and increase in the corrosion potential in comparison with the LSM alloy without anodising. LSM resulted in an increase the thickness of the anodic film from about 200 nm to 320-380 nm. The anodic film formed on the LSM alloy acted as an effective barrier, and significantly limited the exfoliation attack. Therefore, there was no evidence of exfoliation or the occurrence of delamination of the anodised-LSM alloy after immersion testing for 180 min, in contrast to the behaviour of the LSM alloy without anodising which has been described in

Chapter 8.

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Microstructure and Corrosion of the LSM AA 7075-T6 Alloy after Anodising

9.5 List of Figures

Figure 9-1: Current density (I) vs time (s) behaviour of the as-received (AR) and LSM AA 7075-T6 alloy during anodising at 12 V (SCE) in 0.46 M

H2SO4 at 295 K for (a) 240 s, (b) 720 s, and (c) 1200 s.

193

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy after Anodising

Figure 9-2: Scanning electron micrographs of the top surface of the LSM

AA 7075-T6 alloy after anodising at 12 V in 0.46 M H2SO4 at 295 K for 1200 s; (a) the overlap area in the second-laser track and the reheating area in the first-laser track; (b) high magnification of the reheating area; (c) high magnification of the overlap area. The arrows in (c) show the micro-cracks in the re-heating region.

194

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy after Anodising

Figure 9-3: Results of SEM-EDX spot analysis of the LSM AA 7075-T6 alloy following anodising. (a) Scanning electron micrograph showing the location of the analysis region. (b) EDX spectrum. (c) Elements in weight and atomic percentages.

195

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy after Anodising

Figure 9-4: Results of SEM-EDX spot analysis of the LSM AA 7075-T6 alloy following anodising. (a) Scanning electron micrograph showing the location of the analysis region. (b) EDX spectrum. (c) Elements in weight and atomic percentages.

196

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy after Anodising

Figure 9-5: Results of SEM-EDX spot analysis of the LSM AA 7075-T6 alloy following anodising. (a) Scanning electron micrograph showing the location of the analysis region. (b) EDX spectrum. (c) Elements in weight and atomic percentages.

197

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy after Anodising

Figure 9-6: Scanning electron micrographs of cross-sections of the LSM AA 7075-T6 alloy following anodising in sulphuric acid for 1200 s. (b) High magnification of (a).

198

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy after Anodising

Figure 9-7: Transmission electron micrograph of an ultramicrotomed cross- section of LSM AA 7075-T6 alloy anodised at 12 V in 0.46 M H2SO4 at 295 K for 1200 s. (b) High magnification of (a).

199

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy after Anodising

Figure 9-8: Scanning electron micrographs of cross-sections of the as- received AA 7075-T6 alloy following anodising in sulphuric acid for 1200 s. (a) to (d) show images at different locations of the alloy surface.

200

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy after Anodising

Figure 9-9: Potentiodynamic polarization tests in deaerated 0.1 M NaCl solution of the as-received (AR) AA 7075-T6 alloy, laser surface treated

(LSM) alloy, and following anodising in 0.46 M H2SO4 for 1200 s. (LSM+Anod., AR+Anod.).

201

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy after Anodising

Table 9-1: Corrosion potential (Ecorr), pitting potential (Epit), passive current density (Ipass) and passive range potential (Epass) for the as-received (AR), laser-treated (LT), AR and anodised, and LSM and anodised alloys determined from the potentiodynamic polarisation curves of Fig. 9-4. Potentials are given with respect to SCE.

202

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy after Anodising

Figure 9-10: Scanning electron micrographs of top surface of the anodised AA7075-T6 alloy after EXCO corrosion test for an exposure time of 180 min. (b) shows high magnification image of (a). (c) shows high magnification image of the location marked with yellow arrow.

203

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy after Anodising

Figure 9-11: Scanning electron micrographs of cross-sections of the anodised AA7075-T6 alloy after EXCO corrosion test for an exposure time of 180 min. (a), (b) and (c) are images at increasing magnification.

204

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy after Anodising

Figure 9-12: Scanning electron micrographs of the top surface of the anodised LSM AA 7075-T6 alloy after EXCO testing for a time of 180 min. (a) overlap and reheating areas. (b, c) middle of laser-treated layer.

205

Microstructure and Corrosion of the LSM AA 7075-T6 Alloy after Anodising

Figure 9-13: Scanning electron micrographs of cross-sections of the anodised LSM AA7075-T6 alloy after EXCO corrosion test for an exposure time of 180 min. (a) to (d) are images at increasing magnification.

206

Chapter 10 Discussion

10. AA 7075-T6 Alloy

The microstructure of the as-received AA 7075-T6 alloy surface, after being polished up to 1 μm diamond paste, revealed intermetallic particles of different size and composition that are often aligned along the rolling direction of the alloy sheet (Fig. 6-1). The dimensions of the large intermetallic particles typically varied between 3 to 11 µm. Some of these intermetallic particles appeared to be broken into small particles during rolling. The other particles were clustered with measured dimensions of the groups of closely-spaced particles being between 20 and 25 µm (Fig. 6-2). The presence of the particles increases the susceptibility of the alloy surface to corrosion, as such particles are reported to be the main cause of corrosion due to their different electrochemical properties with respect to the alloy matrix [32-33, 44-47]. SEM and EDX line scans and mapping revealed copper and iron in the particles. Other particles of size between 1.5 to 5 µm revealed magnesium and silicon (Figs. 6-3, 6-4, 6-5 and 6-6). SEM-EDX analysis detected zinc and magnesium that are distributed uniformly in the matrix

(Fig. 6-7). It is reported that the alloy contains MgZn2 particles and their sizes are very small in comparison with other second phases (hundreds of nanometres) [4,

32]. In addition to the alloy matrix, low angle XRD on the as-received alloy (Fig.

6-8) indicated the presence of two second phases (Al7Cu2Fe and Mg2Si). The overall findings are consistent with previous reports that the second phases of

AA 7075 alloy consist of Al3Fe, Al7Cu2Fe, and Mg2Si with the presence of

MgZn2 [4, 32, 33, 44-47].

207

Discussion

Intermetallic particles on the surface of the alloy usually have different electrochemical characteristics and different effects on the corrosion behaviour of the surrounding matrix [12]. They can exhibit either cathodic or anodic behaviour with respect to the alloy matrix. The presence of these phases can lead to non-uniform attack at specific areas of the alloy surface and are a key reason for the initiation of pitting corrosion (Fig. 6-10) [4].

The potentiodynamic polarisation scan of the polished AA 7075-T6 alloy in deaerated 0.1 M NaCl solution (Fig. 6-9) revealed two breakdown potentials. The first break-down potential may be attributed to pitting. The second break-down potential may be related to pitting propagation and dissolution of a reactive surface layer [27, 197]. The process of mechanical polishing of an alloy surface can be responsible for production of a thin surface layer with an increased electrochemical activity. Other investigators have made similar observations of the potentiodynamic polarisation in deaerated sodium chloride solution of AA

7075 alloy after polishing to those of the present work [27, 197]. They concluded that the formation of a thin zinc-rich layer due to the polishing operation results in a higher anodic electrochemical reactivity than that of the underlying alloy.

According to the observations of the potentiodynamic polarisation found in the present study and of other authors, it seems that the first breakdown potential is responsible for the pitting in the matrix, while the second may be associated with the dissolution of the electrochemical active layer. In the present work, the pitting potential was not changed significantly after LSM, which eliminates the active surface layer, and a second breakdown potential was not observed.

208

Discussion

Scanning electron micrographs shown in Fig. (6-11) disclosed the as-received alloy surface after an exfoliation corrosion test. Corrosion initiates near sites of second phase particles. These intermetallic particles act as cathodes, with the pH rise due to the cathodic reaction resulting in corrosion of the adjacent matrix

(Fig. 6-12). However, other particles are preferentially dissolved after corrosion testing, leaving cavities in the alloy surface. The former phases mainly contain copper and/or iron, while the latter phases mainly contain magnesium, zinc and/or silicon.

Table B of Appendix 1 provides details of the intermetallic particles that are formed in AA 7075-T6 alloy and their corrosion behaviour. The corrosion potential (Ecorr) vs (SCE) and the average current density at Ecorr of the intermetallic particles with respect to the matrix of the alloy are recorded in the

Table. The data in the Table were determined elsewhere using a microcapillary electrochemical cell that facilitated investigation of intermetallic particles in the micrometer size range in 0.1 M NaCl solution [1]. The anodic phases, such as

MgZn2 (η-phase) and Mg2Si (β-phase), have a low Ecorr (-1029 and -1538 mV

(SCE)) in comparison with the cathodic phases, Al3Fe (β-phase) and Al2Cu (θ- phase) (-539 and -665 mV (SCE)) as well as the alloy matrix (-799 mV (SCE)).

Iron and silicon are reported to be basic impurities in commercial grade aluminium alloys [45]. The solubilities of iron and silicon are high in molten aluminium, but are low in solid aluminium [46]. Therefore, most of the iron and silicon will precipitate out during solidification and form large second phase

209

Discussion

particles. The iron-rich particles, such as Al3Fe and Al7Cu2Fe, usually act as cathodes on the surface of the alloy, promoting electrochemical attack. Murray et al. reported that the presence of copper ions in the solution would cause metallic copper to plate onto the iron-rich particles [47]. If the matrix adjacent to the iron- rich particles contains copper, aluminium can be dissolved and the copper will remain inside pits, which makes the particles more cathodic with respect to the aluminium matrix [47]. On the other hand, constituent particles containing magnesium, such as Mg2Si and MgZn2 phases, are more active than the matrix of the alloy and are usually dissolved preferentially. Some elements of an active particle will leave the pit after being dissolved, such as magnesium, but some will stay, such as silicon. The change in concentration of the elements inside the pit will affect the further corrosion mechanism.

Overall, the presence of different second phase particles in AA 7xxx alloys affects their corrosion behaviour strongly. Therefore, reduction or elimination of such intermetallic particles from the surface of the alloy, without affecting the properties of the alloy matrix, by using LSM can lead to an improved corrosion resistance of the alloy by reducing the micro-galvanic coupling between the aluminium matrix and the second phase particles. The next section will discuss the AA 7075-T6 alloy surface after laser treatment using an excimer laser.

210

Discussion

10.1 AA 7075-T6 Alloy after LSM

In the present excimer LSM, adequate solute dispersion, with an increase in corrosion resistance, was achieved by using a laser pulse fluence of 3.3 J/cm2, as was assessed from the results of SEM examination of LSM cross-sections and of the potentiodynamic polarisation tests (Figs. 7-5, 8). An increase in the thickness of the melted-layer by increasing the fluence of the excimer laser led to an increased number of defects, such as a pores and cracks (Fig. 7-3), which can reduce the corrosion resistance. Ryan et al. showed that a large degree of solute dispersion occurs by diffusion and mixing in excimer LSM of AA 7075 and AA

2024 alloys, which provided an improvement of the corrosion resistance [166]. It is evident that the excimer laser pulse heats and melts the surface of the alloy and due to rapid solidification associated with a high cooling rate, a wave-like melted layer can be generated (Figs. 8-1, 2). However, the production of such a melted layer can be affected by a number of different processes, particularly vaporisation and surface tension gradients. A high fluence received from the laser pulse can produce a high temperature on the surface of the alloy. It is reported that vaporisation requires high temperatures and surface tension gradients depend on the temperature differences in the melt pool [143, 144, 167].

Once the temperature of the alloy surface is at the vaporisation temperature, large thermal stresses can be induced, which may be sufficient to produce defects such as pores caused by evaporation of elements, i.e. zinc and magnesium, that have a boiling temperature lower than the temperatures that result from the laser pulse, and subsequent cracking can reduce the improvement of the corrosion resistance.

Details of the physical properties of the elements in the AA 7075-T6 alloy,

211

Discussion including their melting and boiling temperatures, are provided in Table A of

Appendix 1. Ideally, the laser fluence is required to be controlled to produce a melted layer with few defects. This was achieved in the present work by using a low pulse energy density instead of a high pulse energy density. Therefore, an excimer laser fluence of 3.3 J/cm2 and 80 pulses per unit area was selected as the preferred condition for further improvement to corrosion resistance after evaluation of several different laser treatment conditions, as explained in Chapter

Following LSM treatment of the AA 7075-T6 alloy, the melted layer exhibited bands of precipitates in the solidified structure (Figs. 8-11, 12, 13, 14). Other investigators have made similar observations after treating AA 7075, AA 2024 and AA 2025 alloys using an excimer LSM [166, 167, 180]. The maximum number of bands that can be present within a modified melted layer and still provide corrosion resistance improvement is unknown, but the corrosion resistance is more likely to be improved if the bands can be minimised. EDX analysis (Figs. 8-13, 14) disclosed zinc and magnesium enrichment in these bands which may be expected to dissolve if the corrosive electrolyte reaches these regions within the melted layer. Dissolution along the bands may then cause delamination of the laser treated layer. Other investigators found bands with enrichment of copper and magnesium after treating AA 2024 alloy using excimer LSM [166, 167]. They suggested that these bands within the melted layer could be generated as the result of an initial transient caused by the acceleration of the advancing growth front, combined with the solute enriched trails, that extends the interface stability criteria to greater growth velocity [166,

212

Discussion

167]. Other works found fine discontinuous bands, containing the main elements of precipitates of the alloy after excimer LSM of AA 2024 and AA 2050 alloys

[26, 180]. They concluded that the short duration of each laser pulse prevented complete diffusion of the alloying elements of the intermetallic precipitates within a melted layer. It might be possible that these bands result from the short time spent in the liquid state after each laser pulse, which does not give sufficient time for the main alloying elements to be mixed. Therefore, precipitates as they arrived in the melt pool area leave behind band of discontinuous precipitates of bands within laser melted layer.

The high resolution TEM micrographs of Figs. 8-11 and 8-14 revealed a distance of around 0.3 to 0.4 µm separating two bands. It might be suggested that increasing the thickness of the melted layer by increasing the laser fluence may increase the number of bands [166, 180]. However, the presence of defects, such as large pores, may increase the localised attack and reduce the corrosion resistance in environments that contain chloride ions. In contrast, generating a thin melted layer by reducing the laser fluence and increasing the number of pulses could minimise the number of bands in the melted layer and that may enhance the improvement of the corrosion resistance. Therefore, the laser fluence was reduced to 3.3 J/cm2 and the laser pulses were increased to 80 pulses to treat the surface of the alloy in this study. Other investigators found that increasing the number of laser pulses from 10 to 50 can reduce the precipitation bands in LSM

AA 2024 and AA 2050 alloys and subsequently improve the corrosion resistance and reduce the localised attack in the melted layer [202]. It has also been reported that in LSM the solidification front moves from the base of the melted

213

Discussion layer to the surface of the melt pool, following the thermal gradient. There is a region of slow growth of the solidification front at the base of the melted layer, which has been reported as an interface [166]. The authors suggested that excimer LSM treatment provides a mixing of the aluminium matrix and the second phase particles in solid solution of the melted layer. The solid solution should then solidify quickly to prevent interface instability, trapping the solute in solid-solution. However, due to the short melt time (nano-seconds) after each repeated number of pulses, it might be possible that precipitation can occur

[166]. Other investigators showed that due to the short duration of each laser pulse (~13 ns), it is difficult to achieve complete dissolution of the intermetallic precipitates in the melted layer and, therefore, precipitation and segregation may develop in the melt pool area [180, 201]. From the above discussion, it is possible that the high cooling rate of the excimer laser, with the presence of different alloying elements that have different physical properties in the solid solution of the melted layer, influences the solidification microstructure by generating thin bands of precipitates, which are enriched with the alloying elements.

The XRD for both the as-received and the LSM alloy are compared in Fig. 8-16.

It revealed that Al7Cu2Fe and Mg2Si were present in the as-received alloy, and that these phases are reduced after LSM treatment, with increasing reduction being achieved by increasing the number of pulses. Other investigators made the same observation of a reduction in the signal intensities of the large second phases after treating AA 2050-T8 alloy using excimer LSM [180]. It is reported

214

Discussion also that increasing the number of the laser pulses enhanced the dissolution of the intermetallic phases, as revealed by XRD [201, 202].

10.1.1 Corrosion of LSM Alloy

The results of potentiodynamic polarisation tests in deaerated 0.1 M NaCl solution (Fig. 8-17), as summarised in Table (8-1), showed a reduction in the passive current density by two orders of magnitude for the LSM alloy in comparison with the as-received alloy. There was also a significant increase in the passive range of the LSM alloy due to the formation of the 3-4 µm thick melted layer. The increased range can be attributed to a reduction in the cathodic reactivity of the alloy surface after laser surface treatment due to reduction in the presence of copper-and iron-rich intermetallic phases. Other work reported similar observations after treating AA 2024-T531 friction stir welds by excimer

LSM [203]. The findings of an extended passive range and the reduction in the passive current density are also consistent with the results found by Viejo et al. on AA 2025-T8 alloy and Liu et al. on AA 2024-T351 alloy after excimer LSM

[201, 202] In contrast, the breakdown potential (pitting potential) was not significantly affected by the LSM process. Other authors made similar observations after excimer of LSM AA 7449-T7951 alloy [204]. They suggested that the incorporation of zinc in solid solution in the LSM layer creates a relatively active layer, which is anodic with respect to the base alloy and detremines the electrochemical behaviour [204]. The reduction in the cathodic reactivity and the increase in the anodic activity of the alloy surface may be responsible for the change in the electrochemical behaviour of the present LSM alloy, which revealed only one breakdown potential. Yue et al. studied the effect

215

Discussion of excimer LSM on the corrosion resistance of AA 7075-T651 alloy using potentiodynamic polarisation in 3.5% NaCl solution open to air. They found a reduction by six-times in the passive current density following LSM, and a passive region was obtained, in contrast to the behaviour of the untreated alloy.

The untreated alloy suffered from widespread pitting corrosion [192]. They investigated the surface of laser-treated and untreated alloys and concluded that the improvement in the corrosion resistance was due to presence of an oxide layer on top of the laser-melted zone, which is a chemically stable phase (α-

Al2O3) that serves as an effective barrier to protect the matrix against corrosion

[192]. The present work also showed (Figs. 8-4, 5, 13, 14) an increased oxygen content after treating the alloy with the excimer laser, as evident by EDX analysis of the alloy surface.

The results of immersion tests of the LSM AA 7075-T6 alloy in 0.1 M NaCl solution, open to the air, for 24 and 48 h, revealed no evidence of pitting corrosion, with the laser melted layer being similar to that before the immersion test (Fig. 8-18). The findings are consistent with an improvement in the corrosion resistance of the LSM alloy. The improvement can be related to dissolution/removal of the large (noble) second phase particles that are copper- rich and/or iron-rich elements from the surface of the alloy, which may act as preferential sites for oxygen reduction or hydrogen evolution in chloride environments, and the absence of the other active constituent particles that are rich with magnesium element which can be dissolved uniformly. In comparison, the untreated alloy disclosed pitting corrosion due to the presence of these second phase particles on the alloy surface (see Fig. 6-10). Other investigators found an

216

Discussion improvement in the pitting corrosion resistance of LSM AA 2024-T351 friction stir welds after immersion in the same solution [203]. They concluded that LSM using an excimer laser decreases and homogenises the anodic and the cathodic reactivity in both the weld region and the parent material by dissolution of the intermetallic particles and redistribution of the alloying elements in the laser melted layer [203].

Treating a sufficiently large area of the alloy surface, using an excimer laser, required the use of two adjacent laser tracks. High magnification scanning electron micrographs of the top surface of the melted area (Fig. 8-3 (e)) disclosed defects, such as micro-cracks on the first-laser-track direction that occurred due to tempering or re-heating from the second-laser-track movement. The SEM micrographs of cross-sections of the laser melted layer (Figs. 8-6, 7-10) show that the cracks do not penetrate deeply, suggesting that these cracks may only occur in a thin surface layer of melted alloy or may only penetrate the surface of an oxide layer. It appears that during exfoliation testing of the LSM AA 7075-T6 alloy, partial delamination of the melted layer occurred after 180 min (Figs. 8-22 and 8-23). The test solution is able to propagate along the melted layer interfaces without attacking the alloy matrix significantly. In contrast, observations of the as-received alloy revealed deeply penetrating intergranular corrosion (Figs. 8-21,

22, 23). The zinc and magnesium precipitates in the melted layer and at the interface may dissolve preferentially and subsequently reduce the corrosion resistance, causing localised attack in the presence of an aggressive environment.

It is reported that the different microstructure in the interface region, together with presence of bands of precipitates in the melted layer, makes LSM alloys

217

Discussion more susceptible to corrosion and can lead to delamination of the melted layer

[162, 165, 180, 201, 202, 203, 204]. The delamination of the melted layer can not be avoided once the top layer is breached by the corrosion solution, and the solution has access to the reactive precipitation bands.

Optimising the energy density and selecting a suitable number of pulses of the excimer LSM is a one of the key factors in increasing the corrosion resistance of the AA 7075-T6 alloy, and that starts with careful control of the laser pulse energy density to cause only melting of the alloy and mixing of the alloying elements in solid solution and hence to produce a uniform melted layer.

From the above discussion of LSM, the improvement in corrosion resistance was due to formation of a thin homogenous melted layer by dissolution of the intermetallic particles into the α-Al solution. The solute is then maintained in the surface layer through planar solidification due to the high cooling rate that is a characteristic of the excimer laser. However, due to very short melt times, the melted layer exhibited bands of precipitates that contain zinc and magnesium.

Generating more than one laser-track can influence the treated area due to over- heating that may cause micro-cracks at the previous laser-track that may reduce the resistance of the laser treated alloy to corrosion. Therefore, protecting of the

LSM alloy is required to improve the resistance of the laser melted layer to delamination after exfoliation testing, Anodising of the LSM alloy in 0.46 M sulphuric acid to further protect the alloy will be discussed in the next section.

218

Discussion

10.2 LSM of AA 7075-T6 Alloy after Anodising

Anodising of aluminium and aluminium alloys followed by the application of an organic coating, such as paint, is commonly used in the aerospace industry to protect the substrates from corrosion. The oxidation behaviour depends on the composition of the substrate and the nature of the electrolyte. The as-received

AA 7075-T6 alloy showed large second phases rich with copper and/or iron elements, aligned on the rolling direction, of the alloy surface. In comparison, the

LSM AA 7075-T6 alloy showed zinc, magnesium and copper in solid solution or as fine precipitates in the melted layer, with the presence of oxygen that is attributed to alumina formation on the alloy surface (Figs. 8-4, 5, 13, 14). During anodising of LSM alloys at 12 V (SCE) in 0.46 M H2SO4 for 240, 720 and 1200 sec, the steady current density was reduced by about 50 % in comparison with the as-received alloy, respectively (Fig. 9-1), suggesting that the LSM alloy can grow a porous film more efficiently than the as-received alloy. The presence of copper-rich and iron-rich intermetallic particles on the as-received alloy decreases the efficiency of the anodising due to favouring secondary electrochemical reactions during the anodising process, in particular oxygen generation. As a result, the current density is increased. It is reported that during anodising of Al-Cu and Al-Cu-Fe second phases in acidic solution, the presence of copper leads to extension oxygen evolution [199, 200]. In contrast, as the large intermetallic particles are dissolved by laser surface melting, the process of electron discharge is less favoured and, therefore, less oxygen evolution takes place. This results in a reduction of anodic current density. Sealing the alloys after anodising with deionised boiling water is important. It is reported that the

219

Discussion

hydration of alumina in boiling water produces aluminium hydroxide (Al(OH)3) that blocks the pores of the anodic film, hence preventing aggressive ions from the corrosive solution accessing the underlying alloy [200, 205].

The EDX analysis of the LSM AA 7075-T6 alloy after anodising for 1200 s,

(Figs. 9-3, 4 and 5) revealed that the level of oxygen is increased on the anodised alloy, in comparison with the level of oxygen in the LSM alloy without anodising

(Figs. 8-4 and 5).The high magnification of SEM cross-sections revealed a thin film generated over the LSM alloy (Fig. 9-6). TEM of ultramicrotomed cross- sections (Fig. 9-7) disclosed a bubble-like texture of the anodic film. The film was about 330 to 380 nm thick, according to the results of SEM and TEM, and had developed uniformly above the laser treated layer. Elements of the melted- layer and the electrolyte are incorporated into the anodic film, which is composed mainly of amorphous alumina. The oxygen is generated by oxidation of O2- ions of the alumina and is present in bubbles in the anodic film [200], as well as being released to the electrolyte. The as-received alloy after anodising showed an anodic film about 200 nm thick developed over the alloy matrix and a

2 µm thick layer developed over the large second phase particles, with the presence of defects, such as micro-cracks, at both locations (Fig. 9-8). It is expected that the presence of intermetallic particles enriched with copper and/or iron increases the level of oxygen evolution and, therefore, results in less efficient formation of the anodic film compared with the anodic film generated after LSM.

220

Discussion

The potentiodynamic polarisation measurements in sodium chloride, (presented in Fig. 9-9 and in Table 9-1), revealed that the LSM AA 7075-T6 alloy after anodising showed a reduction in the passive current density by more than two orders of magnitude in comparison with the as-received alloy before and after anodising, and by one order of magnitude in comparison with the LSM alloy.

Thus, the anodising provided an improvement in the resistance of the laser- treated alloy to corrosion initiation. Further corrosion investigation, using EXCO test, showed that the anodised as-received alloy after exfoliation testing for 180 min (Figs. 9-10 and 9-11) suffered from exfoliation corrosion. The depth of the exfoliation attack was about 65 µm. In contrast, scanning electron micrographs of both top surface and cross-section of the anodised LSM alloy after exfoliation testing (Figs. 9-12 and 9-13) revealed that the alloy was not attacked significantly. Corrosion products were found only locally in a few locations and no evidence of delamination was found, unlike the behaviour of the LSM alloy without anodising following a similar EXCO corrosion test (Fig. 8-23). It is evident that anodising can reduce further the corrosion susceptibility of the laser- treated alloy due to the generation of a uniform oxide film above the alloy, which acts as a protective layer that prevents the corrosive solution from attacking the melted layer. Other investigators made the same observations of the corrosion resistance of the LSM AA 2050 and LSM AA 2024 after anodising in sulphuric acid [201, 202]. They concluded that the improvement was attributed to the formation of thick oxide layer that protected the alloys from corrosion attack in environments containing chloride ions.

221

Discussion

From the above discussion, LSM of the AA 7075-T6 alloy followed by anodising in H2SO4 produced a stable film above the melted layer, which acted as an effective barrier that protects the substrate surface from further exfoliation corrosion. The dissolution of the large intermetallic particles combined with refinement of the alloy surface after LSM is responsible for the improvement of the anodising performance in comparison with as-received alloy.

222

Conclusions and Suggestions for Future Work

Chapter 11 Conclusions and Suggestions for Future Work

11.1 Conclusions

The AA 7075-T6 alloy after surface and cross-section examinations revealed different sizes of intermetallic particles that are aligned in the rolling direction.

Large particles were copper-and iron-rich, and small particles were magnesium- and silicon-rich. XRD disclosed the presence of Mg2Si and Al7Cu2Fe phases.

After immersion testing in 0.1 M NaCl, the alloy exhibited pitting adjacent to second phase particles, associated with dissolution of the surrounding matrix.

Exfoliation (EXCO) tests show that corrosion occurred within 30 min of immersion, and extended into the matrix with increasing time of immersion. The aggressiveness of the test solution with the presence of second phase particles on the alloy surface allows intergranular corrosion (IGC) to propagate along selected grain boundaries.

In terms of optimization of the LSM, the results from SEM examinations of cross-sections and potentiodynamic polarisation tests of the excimer laser-treated

AA 7075-T6 alloy, showed that a low pulse fluence (3.3 J/cm2) reduced the porosity and generated a more homogenous melted layer. A reduction was observed in the passive current density in comparison with the alloy treated using a high pulse fluence (10 J/cm2). Therefore, the condition of low fluence (3.3

J/cm2) with 80 pulses was selected for further evaluation of the corrosion properties of the laser-treated alloy.

223

Conclusions and Suggestions for Future Work

In LSM, using an excimer laser, of the AA 7075-T6 alloy, a thin homogenous melted layer, between 3 to 4 µm thick, was produced on the alloy surface.

Increase of the number of laser pulses per unit area and a reduction of the laser fluence level (from 10 to 3.3 J/cm2) caused dissolution of the large intermetallic particles and precipitates on the alloy surface. The high cooling rate (nanosecond) associated with the rapid solidification of the excimer laser caused a wave-like top surface of the melted alloy. Generating more than one laser-track can influence the treated area, with re-heating causing defects, such as micro- cracks, in the previous laser-track. The cracks were restricted to a thin surface layer of the laser-treated material, possibly being limited to the oxide film on the alloy surface. Additionally, cross-sections of the melted layer showed bands of discontinuous solutes that were rich in magnesium and zinc. The width of the individual bands was about 10 to 15 nm.

In general, LSM AA 7075-T6 alloy displayed improved passivity in comparison with the as-received alloy, with a significant reduction, by two orders of magnitude, in the passive current density as assessed by potentiodynamic polarization testing. SEM and TEM micrographs of the surface and cross- sections showed a large reduction in the number of large intermetallic particles for the alloy surface treated by LSM in comparison with the as-received alloy.

XRD measurements on the LSM alloy supported the observed reduction of the presence of large intermetallic phases on the laser-treated alloy surface.

Corrosion evaluation after immersion testing in sodium chloride open to air showed an improvement in the resistance to pitting initiation on the LSM alloy

224

Conclusions and Suggestions for Future Work compared with the as-received alloy at the same time of exposure to the solution.

Further, evaluation of the exfoliation corrosion resistance by immersion of both the LSM and as-received alloy in EXCO solution (4.0 M NaCl, 0.5 M KNO3 and

0.1 M HNO3) for 3 h showed a much greater resistance of the LSM alloy to intergranular corrosion (IGC) than the as-received alloy, which revealed severe

IGC. The LSM layer revealed delamination of the melted layer after EXCO test, which is possibly related to the generation of magnesium- and zinc-rich bands within melted layer that are readily attacked by the solution.

Combining the LSM of the alloy with anodising in sulphuric acid electrolyte gave a further improvement in the exfoliation corrosion resistance, due to formation of a protective oxide film, about 380 nm thick, which can heal defects present after the LSM treatment. Consequently, delamination of the melted layer in the EXCO test was prevented. In comparison, the thickness of the anodic film on the as-received alloy was about 200 nm. The presence of intermetallic precipitates enriched with copper and/or iron resulted in a less protective film.

Therefore, the anodised as-received alloy after the EXCO test showed exfoliation corrosion in the alloy matrix. However, the depth of attacked on the as-received alloy matrix after anodising was generally reduced (65 µm) in comparison with the depth of attacked in as-received alloy without anodising (140 µm).

Previous work on 7xxx series using excimer LSM including investigations of the microstructures of welds by Ryan and Prangnell. The work focused 2000 series aluminium alloy, with more limited study of alloys of the 7000 series. Unlike the present work, the corrosion resistance of the alloys was not given attention.

225

Conclusions and Suggestions for Future Work

Notably, the study revealed the formation of copper-rich precipitate bands in AA

2024 aluminium alloy. Only two works have examined the corrosion resistance of 7000 series alloys treated by an excimer laser. Davenport carried out both short- and long-term corrosion tests on friction stir welds (FSW), which disclosed delamination of the melted layer due to corrosion. This was suggested to due to the presence of precipitate banding, similar to that reported by Ryan and Prangnell, although no microstructural evidence for the banding was presented. The corrosion resistance of LSM AA 7075 aluminium alloy was investigated by Yue et al, but the work was limited to electrochemical impedance spectroscopy and potentiodynamic polarisation tests in sodium chloride solution.

The corrosion previous studies have generally indicated an improvement in the corrosion resistance following LSM treatments.

The present study has added to the knowledge on LSM of 7000 series aluminium alloys, presenting evidence of the influence of different laser processing conditions on the morphology of the melted layer and showing a benefit of using a lower fluence than used in earlier studies. Furthermore, microstructural examination has shown that MgZn2-rich bands are formed in the LSM layer, which have not been reported before. The corrosion testing regime has also been extended to relatively long-term exfoliation corrosion tests, revealing the benefits from LSM of the alloy. Finally, it has been demonstrated that a further improvement in the resistance of the alloy to delamination can be achieved post- treatment of the laser-treated surface using anodizing to generate a protective oxide film.

226

Conclusions and Suggestions for Future Work

11.2 Suggestions for Future Work

The investigations of LSM of the AA 7075-T6 alloy showed an improvement in the corrosion resistance of the alloy due to the generation of a thin homogenous melted layer that was free from large intermetallic particles. However, the presence of defects in the melted layer, particularly in the re-heating region of the first laser-track, suggests scope for further improvement the corrosion resistance.

i- Investigate the relationship between the multiple laser track movements,

number of pulses per unit area and crack initiation due to over-heating.

ii- Modelling the rate of heat diffusion with respect to the number of laser

pulses to identify the location of heat affected zones of the laser treated

area.

iii- Investigate the effect of different anodising conditions on the LSM layer

with respect to the alloy matrix to identify the best solution for anodising

and the optimum thicknesses of the resulting oxide.

iv- Investigate the relationship between the number of the lateral bands in the

melting layer and their corrosion behaviour with respect to the alloy

matrix before and after anodising.

v- In the present wok, precipitation bands have been identified by TEM in

the melted layers above the matrix and solute trails have been identified

by SEM above the copper-rich second phase particles. Further studies can

investigate by TEM the distribution of precipitates above the latter

particles in order to determine the distribution of fine precipitates

throughout the melted layer.

227

Chapter 12 Appendix 1

A- Physical properties of the elements in 7xxx alloys.

Element Wt. Melting Boiling Vapour Heat.of Crystal % Point Point Pressure. Vap. Structure. (K) (K) (Pa/K) (kJ/mol) Zn 5.83 693 1180 610 123.6 Hexagonal

Mg 2.45 923 1363 701 128 Hexagonal

Cu 1.56 1358 2835 1509 300 FCC

Fe 0.24 1811 3134 1728 340 BCC

Cr 0.21 2180 2944 1656 339.5 BCC

Si 0.12 1687 3538 1908 359 DC

Mn 0.08 1519 2334 1228 221 BCC

Ti 0.03 1914 3560 1982 425 Hexagonal

Al Bal. 933.5 2792 1482 294 FCC

228

Appendix 1

B- Corrosion behaviour of the intermetallic particles on AA 7075-T6 alloy after immersion test in 0.1 M NaCl

Intermetallic Phase Ecorr Icorr Behaviour Comments Particles (mV) (A/cm2) with respect (SCE) to matrix

-4 Al7Cu2Fe S-phase -551 -3.1 x 10 Cathodic Constituent Particles

-4 Al2Cu θ -665 -4.7 x 10 Cathodic Constituent

-5 Al3Fe β -539 -9.9 x 10 Cathodic Constituent

-4 Mg2Si β -1538 1.9 x 10 Anodic Precipitates and Constituent

-3 MgZn2 η -1029 1.0 x 10 Anodic Precipitates along grain boundaries

Matrix α -799 -8.1 x 10-5 Anodic ------

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