A Study of Interrupted Aging in Al-Cu-Mg Alloys

A STUDY OF INTERRUPTED AGING IN AL-CU-MG ALLOYS

by Joseph Ming-Ju Tsai A thesis submitted to the Faculty and Board of Trustees of the Colorado School of Mines in partial fulﬁllment of the requirements for the degree of Doctor of Philosophy (Metallurgy and Materials Engineering).

Golden, Colorado Date:

Signed: Joseph Ming-Ju Tsai

Signed: Michael J. Kaufman Thesis Advisor

Golden, Colorado Date:

Signed: Professor Michael J. Kaufman, Department Head Department of Metallurgical and Materials Engineering

ii ABSTRACT

Recently, a novel interrupted aging treatment, T6I6, was reported to improve the strength and toughness of heat treatable Al alloys. It consisted of the following three stages, namely, 1) artiﬁcially underaging at a typical temperature (150-200 ◦C) and quenching to room temperature, 2) reheating to a lower aging temperature (60-100 ◦C) for times that result in hardening, and 3) reheating to the initial aging temperature again and holding until peak hardness is observed. The improvements were reported to be due to precipitate reﬁnement and secondary precipitation during the low temperature age.

However, the improvements in properties have not been consistently reproduced in the literature, and an explanation for the discrepancies remains unclear. Therefore, the focus of this study was to investigate the interrupted aging processes and the eﬀects of each aging stage, as well as the eﬀect of cold work, on precipitate development in commercial 2024 Al and a high-purity variant of the commercial 2014 alloy, Al-4.2Cu-0.4Mg (wt%). The aging responses, precipitation behavior, and microstructural changes were examined using Vickers microhardness, electrical resistivity, and electron microscopy. Small angle X-ray scattering (SAXS) was also used to characterize the precipitate size and volume fraction during the initial underaging and subsequent stages of interrupted aging.

It was found that, although secondary hardening was observed during the low temperature aging, apparent softening upon re-heating was observed and the peak hardness showed little improvement. In addition, in 2024 alloys which were cold worked after SHT or after underaging, no further strengthening was observed during secondary aging at 65 ◦C. Likewise, subsequent aging to peak hardness showed no improvements. In order to understand these results, an attempt was made to examine the mechanism of hardening during the underaging and secondary hardening treatments using SAXS. Formation of 1 nm diameter solute clusters up to ∼3 vol% was found to contribute primarily to the initial rapid hardening and secondary strengthening during interrupted aging. The subsequent softening was attributed to the reversion of these 1 nm clusters.

Further analyses of the observed cluster hardening involved estimating the modulus and/or surface/chemical strengthening. The shear modulus of the clusters was estimated to be 2.9-3.3 GPa higher than that of 2024, and additional surface energy from cluster shearing was estimated to be 0.16-0.63 J/m2.

iii TABLE OF CONTENTS

ABSTRACT ...... iii

LIST OF FIGURES ...... vii

LIST OF TABLES ...... xii

ACKNOWLEDGEMENTS ...... xiii

CHAPTER 1 INTRODUCTION ...... 1

CHAPTER 2 LITERATURE REVIEW ...... 3

2.1 History of Age Hardening in Aluminum Alloys ...... 3 2.2 Solution Heat Treatment SHT and Thermodynamic Basis of Precipitation ...... 3 2.3 Classical Nucleation Theory ...... 6 2.4 Solute Clustering ...... 8 2.5 Precipitate Strengthening Mechanisms ...... 11 2.5.1 Strengthening through Penetrable Particles ...... 11 2.5.2 Strengthening through Particle Bypassing: Orowan Looping ...... 14 2.6 Artiﬁcial Aging and Precipitate Hardening in Al-Cu and Al-Cu-Mg Alloys ...... 15 2.6.1 Precipitation and Hardening during Aging of Al-Cu Binary Alloys ...... 15 2.6.2 Precipitation and Hardening during Aging of Al-Cu-Mg Alloys ...... 18 2.6.3 Initial Rapid Hardening in Al-Cu-Mg Alloys ...... 22 2.6.4 Proposed Mechanisms for Initial Rapid Hardening in Al-Cu-Mg Alloys . . . . 24 2.6.5 Clusters, GPB zones, S’, S Phases and Precipitation Sequence in Al-Cu-Mg Alloys Revisited ...... 25 2.7 Novel Heat Treatment: Interrupted Aging ...... 27 2.7.1 Secondary Hardening and Precipitation ...... 28 2.7.2 Interrupted Aging ...... 28 2.7.3 Evolution of Clusters and GP zones during Low Temperature Dwell . . . . . 32 2.8 Basic Principles of Small-Angle X-ray Scattering, SAXS ...... 36 2.8.1 Small Angle Scattering ...... 36 2.8.2 Interpretations of SAXS Intensity ...... 37

iv 2.8.3 Volume Fraction Calculation Using SAXS Intensity ...... 38 2.9 Scope of Thesis ...... 39

CHAPTER 3 EXPERIMENTAL PROCEDURE ...... 42

3.1 Sample Procurement and Preparation ...... 42 3.2 Introducing Cold Work ...... 42 3.3 Heat Treatment ...... 42 3.4 Hardness Measurements ...... 45 3.5 Electrical Resistivity Measurements ...... 45 3.6 TEM Analysis ...... 46 3.7 Small-Angle X-ray Scattering, SAXS ...... 47

CHAPTER 4 RESULTS ...... 48

4.1 Single-Stage Aging Behavior of 2024 Alloys ...... 48 4.1.1 Hardness and Resistivity Changes of 2024 Alloys during T6 and T8 Aging . . 48 4.1.2 TEM Microstructures of 2024-T6 Alloys at Initial Plateau and Peak Hardening 50 4.1.3 TEM Microstructures of 2024-T8 Alloys at Initial Plateau and Peak Hardening 53 4.2 Interrupted Aging of 2024 Alloys ...... 57 4.2.1 Interrupted Aging (T6I6) of 2024 Alloys ...... 59 4.2.2 TEM Analyses of 2024 After T6I4 Treatment ...... 62 4.2.3 Interrupted Aging Treatments T8I6 and T6I8 ...... 62 4.2.4 Stretching Behavior of 2024 Alloys in As Solution Heat Treated and Under- aged Condition ...... 67 4.3 Examination of Interrupted Aging in Al-4.2Cu-0.4Mg Alloy ...... 67 4.3.1 Single-Stage Aging of Al-4.2Cu-0.4Mg ...... 67 4.3.2 Interrupted Aging of Al-4.2Cu-0.4Mg ...... 68 4.3.3 TEM Analyses of Al-4.2Cu-0.4Mg ...... 72 4.4 Summary and Comment on Interrupted Aging of Al-Cu-Mg Alloys ...... 72 4.5 Data Interpretation of Small Angle X-Ray Scattering (SAXS) ...... 74 4.5.1 Obtaining Experimental SAXS Intensity and Converting to Absolute Units . 74 4.5.2 Fitting of Small-Angle X-Ray Scattering Data ...... 74 4.6 SAXS Analyses on the Evolution of Solute Clusters during Initial Rapid Hardening and Interrupted Aging ...... 79

v 4.6.1 Evolution of Cluster Formation During Initial Hardening of Naturally Aged 2024 Alloys ...... 83 4.6.2 SAXS Analyses for 2024 Alloys Aged to Plateau Hardness ...... 85 4.6.3 SAXS Analyses of 2024 Alloys After T6I4 and T6I4-Reheated Treatments . . 88 4.6.4 Relationship Between Electrical Resistivity and 1 nm Clusters ...... 89 4.6.5 SAXS Analysis of Al-4.2Cu-0.4Mg Alloys ...... 92 4.7 Summary ...... 93

CHAPTER 5 DISCUSSION ...... 94

5.1 The Effect of Cold Work on Interrupted Aging ...... 94 5.2 The Relationship Between Resistivity and Volume Fractions of Clusters ...... 96 5.3 Reasoning Behind the Assumptions of Cluster Analysis Using SAXS ...... 96 5.4 Mechanism(s) of Cluster Hardening ...... 98 5.5 Predicted Strengthening of the T6I4 Material and the Subsequent Re-aging Curves . 99 5.6 Volume Fractions of Clusters and Interparticle Spacing in 2024 Alloys at Plateau Hardening, After Secondary Aging, and Re-aging...... 101 5.7 Estimating the Shear Modulus Difference and Excess Surface Energy by Shearing . . 102 5.8 Predicted Volume Fractions of 1 nm Clusters for Al-4.2Cu-0.4Mg Alloys Using Re- sistivity and Their Predicted Strengthening ...... 103 5.9 Aging Temperatures and Volume Fractions at Early Aging Plateau ...... 104 5.10 Possible SAXS Interference Effect in Naturally Aged 2024 Alloy ...... 105 5.11 Conclusions on Interrupted Aging of Al-Cu-Mg Alloys ...... 107

CHAPTER 6 SUMMARY ...... 109

6.1 Single Stage and Interrupted Aging of 2024 Alloys ...... 109 6.2 Single Stage T6 and Interrupted Aging T6I6 of Al-4.2Cu-0.4Mg Alloys ...... 110 6.3 Resistivity, Strengthening Mechanisms for Cluster Hardening, and Conclusion . . . . 110

CHAPTER 7 FUTURE WORK ...... 112

REFERENCES CITED ...... 113

vi LIST OF FIGURES

2.1 The ﬁrst age hardening curve published by Alfred Wilm on Duralumin...... 4 2.2 Schematic illustration of solution heat treatment and aging temperatures...... 5 2.3 Schematic illustration of the change of Gibbs free energy (overall driving force) for decomposition/precipitation from SSSS at Tb in blue and Tc in red (Tb < Tc) as illustrated in Figure 2.2...... 6

2.4 Schematic illustration of Gibbs free energies for phase with miscibility gap at Tb and Tc (Tb < Tc) shown overall driving force for decomposition of the SSSS...... 7 2.5 TEM micrograph of the morphology of precipitates formed from spinodal decomposition in a Cu-33.5 wt% Ni-15 wt% Fe alloy after aging at 775 ◦C for 15 minutes. After Newbury et al...... 8

2.6 Plot of ∆G with respect to r from Equation 2.1 and ∆Gd = 0 for homogeneous nucleation. ∆G∗ is the free energy barrier, or activation energy, for nucleation. ∆G = −V (∆GV − ∆Gs) + Aγ...... 9 2.7 Schematic illustration of total change of Helmholtz free energy vs number of clustering solute atoms in a two-phase region. After Nie et al...... 10 2.8 Increase in CRSS vs the precipitate particle radius r. Adapted from Martin...... 12 2.9 Schematic illustration of particle shearing by a moving dislocation. Reproduced after Martin...... 13 2.10 Schematic illustration of Orowan looping (a) a dislocation bowing around impenetrable particles; (b) continuing motion of dislocation with loops of dislocations left around the precipitate particles...... 15 2.11 Al-rich corner of the Al-Cu phase diagram with metastable solvi of GP, θ00 and θ0, and the equilibrium θ in bold...... 16 2.12 Schematic free energy diagram for Al-Cu binary alloys...... 17 2.13 Schematic illustration of (a) GP zone and (b) θ00. After Sato et al...... 18 2.14 High resolution electron micrographs of GP zone (A) and θ00 (C). Intermediate stage between GP zone and θ00 is also observed (B). Image was taken at B={011}. After Karlik et al...... 19 2.15 Schematic illustration of the crystal structures of (a) θ0 and (b) θ. ◦: Al; •: Cu. The representative TEM (BF at B={011}) and OM (by arrows) micrographs are shown in (c) and (d), respectively...... 20 2.16 Vickers hardness changes of various Al-Cu alloys (in wt%) during artiﬁcial aging at (a) 130 ◦C and (b) 190 ◦C...... 21

vii 2.17 Isothermal ternary phase diagram of Al-Cu-Mg at 190 ◦C. The bold line represents the α+S phase boundary at 500 ◦C...... 22 2.18 Bright-ﬁeld image taken near B={001} of 2024 Al-Cu-Mg alloys at peak aging condition: (a) S’ lath and (b) GPB zones (arrows). Image (a) from Shih et al...... 23

2.19 Reconstructed crystal structure of a GPB zone with Al2CuMg as proposed by Wolverton. 24 2.20 SAED pattern of S-phase at B=<001>...... 25 2.21 Aging of 2024 alloys at 170, 190, and 240 ◦C. After Shih et al...... 26 2.22 An example of DSC curves (heating rate = 10K/min) of an Al-1.2Cu-1.2Mg (at%) alloy after naturally aged for several time intervals. From Starink et al...... 27 2.23 Schematic illustration of a T6I6 interrupted aging treatment. After Lumley et al..... 28 2.24 TEM micrographs of a Al-5.9Zn-2.9Mg (wt%) solution heat treated at 465 ◦C and then (a) directly quenched to and aged at 180 ◦C for 24 h (×36,000) and, (b) directly quenched to and aged at 180 ◦C for 1 h and aged at 135 ◦C for 24 h (×99,000). After Embury et al.. 29 2.25 Examples of T6I6 treatments on (a) Al-4Cu (wt%); (b) 6061 (Al-Mg-Si); (c) 7050 (Al-Zn- Mg-Cu) and (d) 8090 (Al-Cu-Mg-Li). 6061 alloys (b) were initially aged at 177 ◦C; 7050 alloys (c) at 130 ◦C, and 8090 alloys at 185 ◦C. Low temperature aging was conducted at 65 ◦C for all alloys. The vertical broken line in each diagram represented secondary hardening after low temperature aging. After Lumley et al...... 31 2.26 Interrupted aging of a 2014 (Al-4Cu-0.5Mg, wt%) alloy: (a) the hardness responses during T6, T6I6, and secondary aging at 65 ◦C; (b) bright-ﬁeld TEM image of a T6, peak-aged 2014 alloy; (c) bright-ﬁeld TEM image of a T6I6, peak-aged 2014 alloy. After Lumley et al...... 32

2.27 Bright-ﬁeld TEM images of a 6061 alloy at B = <001>Al (a) SHT; (b) underaged at 177 ◦C for 20 minutes; (c) after T6I4 (underaging + 65 ◦C for two weeks); (d) T6; (e) T6I6. The dark dotted features in images (b) and (c) were GP zones. Images were taken at the same magniﬁcation. After Buha et al...... 34 2.28 Schematic illustration of an energy beam with a wavelength of λ being scattered by a particle with size 2R...... 36

2.29 Schematic illustration of how the scattered intensity ratio I/I0 varies with the momentum transfer q for various particle sizes...... 37 2.30 Dependence of scattered intensity I on sample thickness t. After Glatter...... 38 2.31 Aging responses of 6111 alloys undergoing (a) T6 at 180 ◦C and T6I6 at 180/65/180 ◦C and (b) T6I6 with underaging times of 20 seconds and 2 minutes. After Goh et al..... 40 3.1 Schematic illustrations of (a) conventional single stage aging of T6 and T8; (b) interrupted aging T6I6 originally proposed by Lumley et al.; (c) interrupted aging variant “T8I6”; (d) interrupted aging variant “T6I8”...... 44 3.2 Schematic illustration of the four-point probe method developed for the electrical resistivity measurements...... 46 4.1 (a) Hardness and (b) resistivity changes of 2024 alloys during single stage aging at 177 ◦C for determining the T6 and T8 conditions...... 49

viii 4.2 Engineering stress-strain curve of as-solution heat treated (SHT) 2024 alloy during the 5% stretch...... 51 4.3 Representative BF TEM micrograph of 2024 alloys after SHT. Image was taken at B = <001>...... 52 4.4 Representative BF TEM micrograph of 2024 alloys underaged at 177 ◦C for 0.5 h. Image was taken at a two-beam condition with g = 200 near B = <011>. The large particles indicated by arrows were T-phase, Al20Cu2Mn3...... 53 4.5 (a) Selected area electron diffraction (SAED) pattern of 2024 alloys aged to peak hardness near B = <001> at two-beam condition, g = 220, and reflections from precipitates were clearly seen; (b) BF image using the conditions set in (a); (c) SAED pattern of 2024 alloys at peak hardness at B = <001>; (d) CDF image using the precipitate reflections circled in (c)...... 54 4.6 (a) BF image of 2024 alloy underaged at 177 ◦C for 0.5 h. Image was set to a two- beam condition with g = 200 near B = <001>. The arrows indicate the T-phase ◦ (Al20Cu2Mn3); (b) High-resolution TEM image of 2024 alloy underaged at 177 C for 0.5 h. Image was taken at B = <011>...... 55 4.7 BF image near B = <112> taken at a two-beam condition with g = 111¯ of as-stretched 2024 alloys...... 56 4.8 BF image near B = <011> taken at a two-beam condition with g = 200 of 2024 alloys stretched to 5% and aged at 177 ◦C for 0.5 h...... 57 4.9 (a) SAED pattern of 2024-T8 alloys aged to peak hardness at two-beam condition. g = 111¯ near B = <112>; (b) BF image using the condition set in (a); (c) WBDF image of (b) taken at g-3g condition; (d) CDF image using the precipitate reflections indicated by the white circle in (a)...... 58 4.10 (a) Hardness and (b) resistivity changes of 2024 alloys during secondary aging at 65 ◦C. 60 4.11 (a) Hardness and (b) resistivity changes of 2024 alloys during interrupted aging T6I6. . 61 4.12 (a) BF image of a 2024 alloy after a T6I4 treatment (g = 200 near B = <001>); (b) BF image of 2024 after a T6I4 treatment and reheated with “fast” heating rate. (g = 200 near B = <011>); (c) BF image of 2024 after a T6I4 treatment and reheated with “slow” heating rate. (g = 200 near B = <011>)...... 63 4.13 (a) Hardness and (b) resistivity changes during 65 ◦C dwell for two weeks for T6I6, T8I6, and T6I8 interrupted aging treatment of 2024 alloys...... 64 4.14 (a) Hardness and (b) resistivity changes of 2024 alloys during interrupted aging T8I6. . 65 4.15 (a) Hardness and (b) resistivity changes of 2024 alloys during interrupted aging T6I8. . 66 4.16 Stress-strain curves of 5% stretching of as solution heat treated (SHT) and underaged for 0.5 h at 177 ◦C (30 UA) 2024 alloys...... 68 4.17 (a) Hardness and (b) resistivity changes of as-cast Al-4.2Cu-0.4Mg alloys during single stage T6 aging...... 69 4.18 Natural aging response of the Al-4.2Cu-0.4Mg alloy illustrated by (a) Vickers hardness and (b) resistivity...... 70

ix 4.19 (a) Hardness and (b) resistivity changes of as-cast Al-4.2Cu-0.4Mg alloys during secondary aging at 65 ◦C...... 71 4.20 (a) Hardness and (b) resistivity changes of as-cast Al-4.2Cu-0.4Mg alloys during T6I6 interrupted aging...... 73 4.21 (a) As-scanned SAXS data from a 2024 alloys aged at 65 ◦C for 22h; (b) the same scan from (a) with colored contrast gradient...... 75 4.22 SAXS intensities for pure aluminum and 2024 alloys aged at RT and 177 ◦C for 18 h. . 76 4.23 Absolute SAXS intensities for 2024 samples naturally aged at RT for times ranging from 0.5 to 24 h...... 77 4.24 SAXS intensity (bold black curve) of a specimen consisting of clusters with radii of 0.5, 1, 2, and 3 nm, and volume fractions of 0.6, 0.1, 0.1, and 0.2, respectively: (a) the measured data is first shown with scattering from clusters with radius of 0.5 nm and volume fraction of 0.6; (b) the following portion is then shown with scattering from clusters with radius of 1 nm and volume fraction of 0.1. Subsequent illustrations are shown for clusters with (c) radius of 2 nm and volume fraction of 0.1, and (d) radius of 3 nm and volume fraction of 0.2...... 78 4.25 SAXS intensity with different volume fractions of cluster distributions. The bold black curve consists of clusters with volume fractions of 0.6 (0.5 nm), 0.1 (1 nm), 0.1 (2 nm), and 0.2 (3 nm), and the blue bashed curve consists of clusters with volume fractions of 0.7 (0.5 nm), 0.1 (1 nm), 0.1 (2 nm), and 0.1 (3 nm)...... 79 4.26 SAXS data fitting (naturally aged 2024 alloys): (a) absolute intensity data; (b) the fitting curve with Id and Porod slope C; (c) 1 nm size clusters with proper relative volume fraction to best fit the raw data was introduced; (d) clusters with sizes from 2 to 30 nm were incrementally fitted to give the best fit to the raw data...... 80 4.27 Examples (naturally aged 2024 alloys) of SAXS data fitting with cluster sizes of (a) 0.8 nm and (b) 1.2 nm compared to fitting with 1 nm (blue dash line) with a fixed relative volume fraction...... 81 4.28 Examples (naturally aged 2024 alloys) of SAXS data fitting with a fixed cluster size of 1 nm with a volume fraction of (a) 2 vol% and (b) 4 vol% compared to fitting with 3 vol% (blue dash line)...... 82

4.29 Relative fractions (vi in Equation 2.19) of 1 to 30 nm size clusters from SAXS analyses of 2024 alloys naturally aged up to 506 h...... 83

4.30 (a) Change of hardness (with respect to SHT, 94 HV ) and (b) 1 nm cluster volume fractions changes with aging time (up to 24 h)...... 84 4.31 Hardness evolution of 2024 alloys aged to hardness plateau at temperatures from RT to 190 ◦C...... 85 4.32 Absolute SAXS intensities for 2024 alloys aged from 25 to 190 ◦C. Pure Al was also included for comparison and showed very weak diﬀuse scattering (Id), consistent with little or no atomic scale density ﬂuctuations in a pure metal...... 86 4.33 SAXS intensities of 2024 alloys aged at RT (25 ◦C) for 506 h, 150 ◦C for 0.5 h, 177 ◦C for 0.5 h, and 177 ◦C for 18 h...... 87

x 4.34 Relative fractions (vi in Equation 2.19) of 1 to 30 nm size clusters from SAXS analyses of 2024 alloys aged to respective hardness plateau at temperatures from RT to 190 ◦C. . 88

4.35 Relative fractions (vi in Equation 2.19) of 1 to 30 nm size clusters from SAXS analyses of 2024 alloys aged at 177 ◦C (for 0.5 h), T6I4, T6I4 reheated to 177 ◦C for 0.5 h (T6I4-reheat), and at 177 ◦C for 18 h...... 89 4.36 SAXS intensities of 2024 alloys aged at 177 ◦C for 0.5 h (underaged), T6I4, and T6I4- reheat to 177 ◦C for 0.5 h (T6I4-reheated)...... 90 4.37 Plateau volume fractions of 1 nm size clusters measured in 2024 alloys plotted against the change in resistivity (with respect to SHT, 5.46 µΩ−cm) at temperatures from 25 to 190 ◦C. The linear relationship (with R2 = 0.948) indicates that the change in resisitivity is directly related to the change in the volume fraction of the 1 nm size clusters. Alloys aged to T6I4 and T6I4 + re-aging to 177 ◦C show a good fit to this relationship, indicating partial reversion of the 1 nm size clusters after re-aging to 177 ◦C. The intercept of the linear regression is set to zero...... 91 4.38 Example of SAXS image plate obtained from the as-cast Al-4.2Cu-0.4Mg alloys (a) with standard holder and (b) with 45◦ tilt. Sample shown in the figure was heat treated to T6I4...... 92 5.1 Stress-strain curve of a 2024 alloy after SHT and underaged at 177 ◦C for 0.5 h. . . . . 95 5.2 Total volume fractions in vol% of clusters 1 to 30 nm) and change of resistivity ∆ρ in µΩ−cm...... 97 5.3 Relationships between change in hardness (relative to that for the SHT condition, namely 94 HV ) and F(r,f) at the hardness plateau at different temperatures: (a) using the chemical strengthening model from Equation 5.4 and (b) using the modulus strengthening model from Equation 5.7. Analyses were conducted for clusters of all sizes (F(r,f)total) and 1 nm only (F(r,f)1 nm). F(r,f) is calculated with r in A˚ and f in volume fraction. . . 100 5.4 Aging temperatures and volume fractions of clusters at early hardness plateau...... 104 5.5 Softening of alloys re-aged at 177 ◦C. The alloys were previously aged at RT and 65 ◦C into their respective plateau...... 105 5.6 Absolute SAXS intensity vs momentum transfer q of 2024 alloys naturally aged for 506 h. Blue curve shows the fitting and the red oval indicates the region where data cannot be properly fitted...... 106 5.7 Schematic illustration of the interference effect with interparticle spacing for 1 nm clusters formed at RT...... 107

xi LIST OF TABLES

2.1 Twelve variants expected between the S phase and the Al-matrix...... 24 2.2 Eﬀect of 5% cold work on tensile properties of a Al-4Cu (wt%) and a 8090 alloy after T8 and T8I6 tempers. After Lumley et al...... 30 2.3 3DAP analyses of 6061 alloys tempered at UA (underaged condition, 20 minutes @ 177 ◦C), T6I4 (UA + 2 weeks @ 177 ◦C), T6 (177 ◦C), T6I6/150, and T6I6/177. After Buha et al...... 35 4.1 Nominal chemical compositions of 2014, 2024 wrought alloys and as-cast Al-4.2Cu-0.4Mg (wt%) ...... 48 4.2 Quantitative SAXS analyses on the volume fractions and cluster size at hardness plateau of naturally aged 2024 alloys and the respective hardness change. The hardness of as solution heat-treated (SHT) alloys was 94 HV ...... 84 4.3 Quantitative SAXS analyses on the volume fractions and cluster sizes for 2024 alloys aged to hardness plateau at temperatures from 25 to 190 ◦C. Units for Q in 1021 cm−4 and f in vol%. C is the Porod slope, which is due to features 30 nm in size...... 88 4.4 Quantitative SAXS analyses of the volume fractions and cluster sizes for 2024 alloys aged at 177 ◦C for 18 hours and tempered to T6I4 and T6I4-reheated to 177 ◦C. Units for Q in 1021 cm−4, and f in vol%...... 89 5.1 Calculated F(r,f) (×10−3) of T6I4 and T6I4-reheat treatments for chemical and modulus strengthening...... 99 5.2 Measured and estimated HV of T6I4 and T6I4-reheat treatments for chemical and modulus strengthening using relationships from Figure 5.3...... 99 5.3 Population densities and interparticle spacings of 1 nm and total clusters for plateau hardened, T6I4, and T6I4-reheat 2024 alloys...... 101 5.4 Calculated surface energies using the chemical strengthening model and modulus diﬀer- ences between matrix and clusters from modulus strengthening model for 2024 aluminum alloys...... 102 5.5 Measured and estimated 1 nm cluster volume fractions and HV of 30-minute underaging, T6I4-2days and T6I4-12days treatments for Al-4.2Cu-0.4Mg alloys using relationships from Figure 4.37 for resistivity and Figure 5.3 for chemical and modulus strengthening. ρ and ∆ρ are in µΩ−cm...... 103

xii ACKNOWLEDGEMENTS

I want to express my sincere gratitude to my advisor, Professor Michael Kaufman, for his genuine support and guidance throughout the study. Special thanks for his patience and willingness to spend time on top of his very busy schedule as the Department Head with me on the microscope and individual meetings. It truly means a lot.

Thanks also to Professors David Olson, John Speer, and Steven Liu from the Metallurgical and Materials Engineering Department, and Professor Mark Eberthart from the Department of Chemistry at Colorado School of Mines for being on my thesis advising committee.

I would also like to thank Dr. Herb Doty at General Motors Powertrain, the NSF/UCRC Center on Advanced Non-Ferrous Structural Alloys (CANFSA), and the Graduate School at Colorado School of Mines for the ﬁnancial support of the study. It truly means a lot, particularly during the ﬁnancial hardship.

Special thanks must be given to Professor Don Williamson for his help with the SAXS analyses, and Professor Kip Findley for his extremely valuable perspective from mechanical metallurgy.

Thanks to Mr. Gary Zito for his assistance on the TEM, and Dr. John Chandler in the Electron Microscopy Laboratory for his assistance on scheduling the instruments. I would also like to thank Dr. Hugh King for his help during the development of the LabView program. Additionally, I’d like to thank Dr. Rishi Raj of the Department of Mechanical Engineering at University of Colorado at Boulder for the use of the “SAXSess” equipment.

Throughout my studies at the Colorado School of Mines, I had the opportunity to work with many outstanding graduate students, some of whom shared with me their research and provided valuable discussions from a variety of perspectives. Here I thank Justin Raines (SSAB Americas R&D), John Gibbs (Northwestern University), Paul Gibbs (Los Alamos National Laboratory), and Drs. Grant Thomas (AK Steel), Joseph Ronevich (Sandia National Laboratory), Mered- ith Heilig (Sundrop Fuels), and Ryan Regier (Stress Engineering Services).

I would like to take this opportunity to thank a few former colleagues at SSAB Americas, who were instrumental in my development as a metallurgist, namely R&D director/manager Rick Bod- nar, Doctors Sunday Abraham, Keith Taylor, Dengqi Bai, and Yufeng Wang for fruitful discussions and constructive suggestions about everything I did at SSAB during my internship.

I want to thank my fellow graduate students in MME, Karem Tello, Matt Kirsch, Scott Cochran, Grant Hudish, and Logan Ward for help and fruitful discussions, which made the rather diﬃcult journey more enjoyable.

xiii I want to thank my elder brother, Dr. Julius Tsai, who has constantly asked this question, or in a similar fashion, “How’s your thesis writing and when are you graduating?” in the past few months when I was rigorously collecting/analyzing data and working on my thesis.

Most importantly, I thank my parents, Jacque and Cristine, for their understanding, counsel- ing, and support.

xiv To my family and friends. CHAPTER 1

INTRODUCTION

Since its accidental discovery in 1906 by Alfred Wilm [1–3], age hardening has been an important process for strengthening aluminum alloys. Typically, aging treatments are performed either at a single temperature to the desired strength, or via two stages consisting of low-high temperatures called double aging. Age hardenable aluminum alloys are of particular interest in areas such as automotive and aerospace industries because of their high strength to weight ratios. Age hardening processes allow these alloys for the applications ranging from engine blocks to spaceframes for space rockets, with alloys that can be designed to meet the application requirements [3]. Re- cently, a novel interrupted aging treatment was reported and shown to further improve the alloy strength and toughness [4–13]. Such improvements would be beneﬁcial to applications because strength/toughness requirements could be met by alloys with lower alloying element contents and could potentially be cost-eﬀective at the industrial level.

The interrupted aging consisted of three stages of aging, namely, 1) artificially underage at a typical temperature (150-200 ◦C) and quench to room temperature, 2) reheat to a lower aging temperature (60-100 ◦C) for times that result in hardening, and 3) reheat to the initial aging temperature until peak hardness is observed. An example of their results using 2014 (Al-4Cu-0.5Mg, wt%) alloy [4] indicates a considerable increase in peak hardness (27%) and fracture toughness (35%) from T6; this increase is believed to be related to a refined distribution of the precipitate phases at peak hardness. The authors suggested that the effectiveness of the heat treatments should be applicable to all heat treatable aluminum alloys, and that such improvements might also be realized if cold work is introduced prior to aging, with some sacrifice in elongation [4].

These improvements were considered to be due to secondary precipitation during the low temperature age. Subsequently, studies [8–10] using 6061 (Al-1.0Mg-0.6Si, wt%) alloys showed that, for alloys undergoing interrupted aging treatments, the secondary precipitation during low temperature aging plays an important role in refining the distribution of the precipitate phases upon peak aging. Atom probe tomography was used to show significant changes in the volume fractions and relative ratios between GP zones and pre-GP zone solute clusters in alloys briefly underaged at high temperature and then held at low temperature for extended times.

By using 3DAP, Lumley and co-workers [7–10] observed two cluster evolutions during the secondary aging: 1) coarsening of the solute clusters formed during the initial high temperature aging; 2) formation of new clusters. The authors deﬁned clusters as undersized (1-2 nm) and irregularly shaped precursors to GP zones, which were larger (2-4 nm) and more regularly shaped (∼spherical). As such, upon reheating, the undersized clusters are expected to dissolve, while those that coarsened to about GP zone sizes remain stable. This results in a higher number density of

1 GP zone-like precipitates. These precipitates may become potential nucleation sites for subsequent microstructural development. The mechanical property improvements at the peak-aged condition of the alloys undergoing interrupted aging treatments, e.g. T6I6, were then attributed to this change in microstructure during the low-temperature aging.

While the improvements in properties reported by Lumley et al. are quite encouraging, they have not been consistently reproduced by subsequent researchers [14–17], who frequently reported contrary results. For example, in 6061 alloys, compared to the T6 treatment, no increase in strength or toughness was observed after the T6I6 treatment [16]. Similar results were also reported in 2024 alloys [17]. However, an explanation for these discrepancies remains unclear.

The focus of this study was to look into the interrupted aging processes and the eﬀects of each aging stage on the precipitate development in commercial 2024 alloys and an Al-4.2Cu-0.4Mg (wt%) alloy which is a high purity version of the 2014 alloy studied by Lumley et al. [4]. The hope was to facilitate a better understanding of the hardening mechanisms at each stage, and thus help explain the discrepancies in the interrupted aging results.

The structure of this thesis is shown as follows:

Chapter 2 starts with a brief summary of the ﬁrst age hardening observed in history, followed by a summary of the theoretical basis for age hardening. The strengthening mechanisms associated with solute clusters and precipitate formation, and examples of age hardening in Al-Cu and Al- Cu-Mg alloys are also brieﬂy summarized. The Chapter ends with a review of the literature on interrupted aging and the scope of this study.

Chapter 3 provides an introduction to the methods and instruments utilized in this study. Procedures for sample preparation and heat treatment are given. Also described are hardness and electrical resistivity techniques used to follow the aging responses of the alloys, and the microstructural analyses that were carried out using TEM and Small Angle X-Ray Scattering (SAXS).

Chapter 4 includes the aging responses of conventional single-stage aging treatments, and the results from interrupted aging studies including additional variants introduced during this study. It also includes microstructural observations and results from SAXS analyses, and occasionally brief discussions are given following the respective results.

Discussion of results are carried out in Chapter 5, including observations not immediately discussed and explained in Chapter 4. Furthermore, the formation of solute clusters and their potential strengthening mechanisms are analyzed and discussed.

Finally, Chapter 6 includes the summary of the discoveries on interrupted aging shown in this study. Recommended further research is included in Chapter 7.

2 CHAPTER 2

LITERATURE REVIEW

This chapter is intended to provide reviews and analyses of published literature on related subjects for the intended research in novel interrupted aging of aluminum alloys. It begins with a review of the classic example of precipitation hardening of Al-Cu binary alloys, followed by reviews of literature regarding aging treatments and precipitation sequences in heat-treatable aluminum alloys. Particular attention is paid to reviewing past and current literature on precipitation hardening in Al-Cu-Mg aluminum alloys. Important literature pertaining to novel interrupted aging is also reviewed. In addition, an overview of precipitation hardening mechanisms is given. Finally, a summary of some present challenges are given at the end of this chapter, as well as the motivation for this study.

2.1 History of Age Hardening in Aluminum Alloys

The precipitation hardening phenomenon in aluminum alloys was first discovered in Berlin in 1906 by German Metallurgist Alfred Wilm [1–3]. While working on replacing the brass used in ammunition, he was attempting to increase the hardness of an aluminum alloy with 3.5-5.5 wt% Cu and less than 1 wt% of Mn and Mg by quenching it from 520 ◦C in hopes of getting similar strengthening effects typically observed in carbon steels. To his frustration, he was unable to obtain the anticipated hardness increases and the alloy became softer the faster he quenched it, which was contrary to the behavior in steels. It was on an opportune Saturday morning when he and his assistant were finishing up a heat treatment, and there was only enough time to make a quick hardness measurement before the lab closed at noon.

The remaining hardness measurements were performed on the following Monday and, to their surprise, the hardness was signiﬁcantly higher than the one recorded two days before. The hardness machine was checked and experiments were repeated, and it was conﬁrmed that the hardness of the alloy increased over a period of time before reaching a constant, as shown in Figure 2.1 [3]. This discovery was published in 1911 and subsequently patented. The alloy he discovered was known as “Duralumin”. Although the alloy was used extensively in early aircraft, the mechanism of hardening was not known until 1919 by Merica et al. [18, 19].

2.2 Solution Heat Treatment SHT and Thermodynamic Basis of Precipitation

Age hardening of aluminum alloys originated from the decomposition of, or precipitation from, a supersaturated solid solution (SSSS) through phase transformations of intermediate phases. The transformations are driven by the reduction of Gibbs free energy. The solution heat treatment

3 Figure 2.1 The ﬁrst age hardening curve published by Alfred Wilm on Duralumin in 1911 [3].

(SHT) involves heating and holding an alloy with composition X0 to a temperature, Ta, above the α + β solvus and just below the eutectic temperature for a period of time to allow solute atoms to dissolve and form a homogeneous phase α, as illustrated by point a in the schematic binary phase diagram with components A and B in Figure 2.2. If the alloy is cooled slowly to point b

(Tb) or point c (Tc), the phases at that point would simply be α + β and the phase fractions can be calculated by the lever rule. However, in an aluminum alloy, i.e. A = Al, if the alloy is quenched rapidly enough so that limited solute diﬀusion takes place, the alloy becomes a solid solution “supersaturated” with solute B, which allows for subsequent decomposition of the SSSS to occur.

The overall driving force for precipitation is the free energy diﬀerence at X0 between the Gibbs energies of α and the equilibrium tie line. The free energies of α and β and their equilibrium tie lines at Tb and Tc are shown in Figure 2.3 If the supersaturated alloy is held, or aged, at Tc, the driving force will be ∆Gc as shown in Figure 2.3. Likewise, the driving force for aging at Tb would be ∆Gb. It is clear that, the lower the aging/holding temperature, the higher the driving force for precipitation/decomposition.

Precipitation typically occurs in three stages: i) nucleation of new phases; ii) growth of the nucleated phases; iii) coarsening of the phases. A higher driving force typically leads to more nucleation and a denser distribution of precipitates, which provides more potent strengthening to

4 Figure 2.2 Schematic illustration of solution heat treatment and aging temperatures. the alloy. Conversely, a lower driving force leads to less nucleation and coarser precipitates, and less strengthening to the alloy.

A special case may also be found in alloys that have solute with the same structure as the matrix, but where a “miscibility gap” exists. The free energy expressions are derived previously [20, 22–

26] and are graphically illustrated in Figure 2.4, with temperature Tb < Tc as shown in Figure 2.2.

If the composition of the alloy X0 is such that it is between the two inflection points on the mixing curve, as indicated by ×’s and arrows, small local concentration fluctuations can lead to an “uphill” diffusion, which can proceed spontaneously without nucleation barriers [20, 22–25]. This is known as “spinodal decomposition”, and coherent precipitates formed through such transformation are uniformly spaced and homogeneously distributed, as shown in Figure 2.5 observed for a Cu-Ni- Fe alloy [27].

Although recent simulation work [28] shows that spinodal decomposition is thermodynamically possible in Al-Ag alloys, it is generally considered that precipitation/decomposition of SSSS in Al alloys is almost exclusively via nucleation and growth.

5 Figure 2.3 Schematic illustration of the change of Gibbs free energy (overall driving force) for decomposition/precipitation from SSSS at Tb in blue and Tc in red (Tb < Tc) as illustrated in Figure 2.2 [20, 21].

2.3 Classical Nucleation Theory

The precipitation of B-rich β phase, shown in Figure 2.2, from A-rich α supersaturated solid solution involves nucleation and growth. The nucleation of a β precipitate with volume V and surface area A can take place at heterogeneous nucleation sites, typically a defect such as a dislocation or an interface such as a grain boundary. It can also take place within the matrix through homogeneous nucleation. For heterogeneous nucleation, if the speciﬁc volume free energy of the precipitate is ∆GV , the strain energy between the precipitate and the matrix per speciﬁc volume is ∆Gs, the free energy change due to destruction of defects is ∆Gd, and the interfacial energy is γ, then the total free energy change to form the precipitate can be expressed in Equation 2.1 [20].

∆G = − V ∆GV + Aγ + V ∆Gs − ∆Gd (2.1)

For spherical precipitates that homogeneously nucleate from the SSSS in an isothermal con- 4 3 2 dition, the volume of the precipitate V = 3 πr , surface area A = 4πr , and free energy of defects ∗ ∆Gd = 0, a critical radius, r , of a nucleus can be derived algebraically from Equation 2.2 or

6 Figure 2.4 Schematic illustration of Gibbs free energies for phase with miscibility gap at Tb and Tc (Tb < Tc) shown overall driving force for decomposition of the SSSS [20, 21]. graphically as in Figure 2.6. The free energy barrier for nucleation, ∆G∗, could also be derived (Equation 2.3).

2γ r∗ = (2.2) (∆GV − ∆Gs)

3 ∗ 16πγ ∆G = 2 (2.3) 3 (∆GV − ∆Gs)

Therefore a nucleus must possess a radius equal to or greater than the critical radius r∗ at the given isothermal condition in order to be energetically stable. A cluster with a radius smaller than r∗ is energetically unstable. It is also known as an “embryo” and will tend to re-dissolve into the solid matrix.

The activation energy for nucleation, ∆G∗, can be reduced by heterogeneous nucleation on defects such as dislocations and grain boundaries. By doing so, the free energy associated with the defects, ∆Gd, can have a negative contribution to ∆G, and the free energy barrier can be further

7 Figure 2.5 TEM micrograph of the morphology of precipitates formed from spinodal decomposition in a Cu-33.5 wt% Ni-15 wt% Fe alloy after aging at 775 ◦C for 15 minutes. After Newbury et al. [27]. reduced. This is the predominant form of nucleation in solids. However, the critical radius r∗ of the nucleus remains identical to that of a homogeneous nucleated nucleus.

2.4 Solute Clustering

Nucleation may also be aided through a special type of intragranular defect called “solute clusters”. These clusters can be deemed as ﬂuctuations of solute distributions in solid solutions, and their existence can be predicted by pair probability function [29].

8 Figure 2.6 Plot of ∆G with respect to r from Equation 2.1 and ∆Gd = 0 for homogeneous nucleation [20]. ∆G∗ is the free energy barrier, or activation energy, for nucleation. ∆G = −V (∆GV − ∆Gs) + Aγ.

Clusters are later defined as solute-rich groupings of atoms whose compositions are different from the random solute distributions in a SSSS. They are also smaller than a critical nucleus and can be quite different in composition, crystallography, and shape to a nucleus [30].

Thermodynamic distinctions between clusters, embryos, and critical nuclei can be made using the standard Helmholtz free energy change, ∆F ◦, and the number of atoms involved in these clusters in a two-phase region, as shown schematically in Figure 2.7 [30].

A region with negative ∆F ◦ can be seen at small n, which is due to the entropy of mixing term (−T ∆S). These entities with small n can be regarded as “clusters”. As n increases, the interfacial energy becomes signiﬁcant and contributes positively to ∆F ◦ because of the unfavorable bonding at the cluster/matrix interface. Similarly, volume strain energy also contributes positively to ∆F ◦,

9 Figure 2.7 Schematic illustration of total change of Helmholtz free energy vs number of clustering solute atoms in a two-phase region. After Nie et al. [30]. leading to an increasing positive total free energy. These clusters are regarded as embryos, which are energetically unstable and prone to dissolve.

If the solute-solute bonds are more favorable than the solute-solvent bonds, the formation of solute-solute pairs (solute clusters) will have a negative contribution to ∆F ◦, as well as the volume free energy of formation. As n continues to increase, the sum of the volume free energy change and the volume strain energy becomes predominant over the interfacial energy, leading to a maximum, ∆F ∗ at n∗, and cause ∆F ◦ to start decreasing. Clusters with n > n∗ become nuclei and are energetically stable.

Cohen [31] suggested that the transition from clusters to critical nuclei was continuous, and the clusters were rapidly and constantly forming and dissolving in times shorter than that required for the formation of critical nuclei.

Formation of clusters may also assist intragranular nucleation. Experimentally, it was ﬁrst suggested by Silcock et al. [32] that clusters could be composed of trace solute elements, which could reduce the mismatch between precipitate and matrix, and reduce the interfacial energy at

10 the interphase boundaries. It was later shown experimentally that, in Al-Cu alloyed with Mg and Ag, Mg-Ag co-clusters promoted the nucleation of the Ω phase [33]. Although the formation of small clusters is energetically favored, they can fail to appear due to reasons such as low diﬀusivity of solute atoms and short survival times [30].

2.5 Precipitate Strengthening Mechanisms

Strengthening in an aged-hardened Al alloy is based on the interactions between moving dislocations and precipitates [1, 3]. The dislocations can interact with the precipitate itself, as well as the possible strain ﬁeld associated with the precipitate.

The combined eﬀects increase the critical resolved shear stress (CRSS, τc), which results in increasing the yield stress of the alloy. The relations between the yield stress σ and τc in polycrystalline materials can be correlated as follows in Equation 2.4 using the Taylor factor M, whose value is inﬂuenced by crystal structure and orientation texture [1, 34]. For isotropic textures, M ≈ 3.1 in fcc and 2.75 for bcc metals.

σ = M · τc (2.4)

The retardation of dislocation motion by precipitate particles can simultaneously involve several interaction mechanisms [1]. The mechanisms ﬁt into two general categories [1] based on the size of the precipitate particles: 1) particle shearing; 2) particle by-passing or Orowan looping. The CRSS increase with particle size is shown in Figure 2.8 [1]. Particle shearing also depends on the modulus of the particle relative to the matrix. In aluminum alloys, however, it was shown by Nicholson et al. [35, 36] that dislocations tend to pass through coherent and semi-coherent particles, but not incoherent particles.

If the precipitates are small and densely populated, the interparticle spacing is small, and the moving dislocations are most likely to break away from the precipitates by shearing. If the particles are large and more spread out, the dislocations are more likely to bypass the particles by bowing, as proposed by Orowan [37]. Obviously, the optimum strengthening from precipitate particles can be obtained at particle radius of rc.

2.5.1 Strengthening through Penetrable Particles

Consider an array of particles is sheared by a moving dislocation, as shown in Figure 2.9 [1], where L is the interparticle spacing, T is the dislocation line tension and φ is the dislocation breaking angle. The “resistance” of the particle to the shearing is F. Therefore for a particle that is sheared at the illustrated condition, F can be expressed by T and φ, as shown in Equation 2.5.

11 Figure 2.8 Increase in CRSS vs the precipitate particle radius r. Adapted from Martin [1].

φ F = 2T cos (2.5) 2

1 2 2 The typical T is 2 µmb [1]. However, it can be 1.6 times greater, i.e. 0.8µmb , in the early stage of aging in Al-Zn-Mg alloys [38] (µm = shear modulus of the matrix material and b = Burgers vector).

There are generally ﬁve properties that aﬀect the ease of particle shearing by the moving dislocation [39]: 1) Coherent strain; 2) Ordered structure; 3) Precipitate Modulus; 4) Surface or chemical; 5) Stacking fault energy.

1. Coherent strain strengthening Coherent strain results from the misﬁtting precipitates that distort the matrix. This causes a stress ﬁeld in the region surrounding the precipitate. This additional stress can interact with the stress from a moving dislocation, and impede its motion. Therefore the force required to

shear the precipitate particle rises and causes τc to increase. The increase in τc is shown in Equation 2.6 [1, 39].

12 Figure 2.9 Schematic illustration of particle shearing by a moving dislocation. Reproduced after Martin [1].

r 1/2 3 1/2 τ = k3/2 µ 3/2 f 1/2 (2.6) c b 2π

where k is a material constant, is the precipitate misﬁt, µ is the shear modulus of the precipitate, f is the volume fraction of the precipitate, r is the radius of the precipitate, and b is the Burgers vector.

2. Order strengthening If the precipitate has an ordered structure, the passing of a dislocation may create an anti- phase boundary (apb), a region where the structural order is disrupted. The additional energy

resulting from creating the apb, γapb, will increase the resistance of the particle being sheared, 0 and therefore will lead to an increase in τc. A typical example of such strengthening is the γ

phase in Ni-based superalloys [39]. The increase in τc can be expressed in Equation 2.7 [1].

γ 6γ fr 1/2 τ = apb apb (2.7) c b π T

In addition, if the dislocations travel in pairs, a more complicated form is shown in Equa- tion 2.8 [1]. " # γ 3π2 γ fr 1/2 τ = apb apb − f (2.8) c 2b 32T

3. Modulus strengthening Because the energy of a dislocation depends linearly on the local modulus, particles which have a modulus diﬀerent from the matrix will interact with the dislocation by locally raising

13 or lowering its energy as it passes through [39]. The relation between the change in τc and the modulus is given by Equation 2.9 [1, 38].

T ∆µ3/2 2 < r >−3/2 τ = 0.9 (< r > f)1/2 2b ln (2.9) c 1/2 b µm f b

where ∆µm is the shear modulus of the matrix and ∆µ is the shear modulus diﬀerence between

the precipitate and matrix, i.e. ∆µ = µ - µm.

4. Surface or chemical strengthening When a precipitate particle is sheared, a step of one Burgers vector is created at the particle- matrix interface for every dislocation that passes. If the surface-to-volume ratio of the particle is relatively high, i.e. when the size of the particle is small and a high population density, this increase in surface area becomes signiﬁcant and can lead to a substantial increase in the surface

energy. This will reﬂect in an increase in τc via the relation shown in Equation 2.10 [1, 39]. GP zones and θ00 precipitates in Al-Cu alloys are typical examples of this type of strengthening. γ is the surface energy of the new surface area created by the sheared precipitates.

3 1/2 γ 3/2 b τ = 2 µ f 1/2 (2.10) c π µb r

5. Strengthening through stacking fault energy If the precipitate has a signiﬁcantly diﬀerent stacking fault energy than that of the matrix, the interaction between an extended dislocation and a precipitate may be dominated by the local variation in the fault width when the moving dislocation enters and is contained within the precipitate [39]. This occurs in close-packed structures, in which it is energetically favorable for dislocations to split into two partials and the stacking fault energy is inversely proportional to the spacing of the partials.

2.5.2 Strengthening through Particle Bypassing: Orowan Looping

As precipitate particles grow, they become increasingly difficult to shear. As a result, the moving dislocation will first “bow” around the particles, as shown in Figure 2.10a. As dislocation motion continues, it may bypass the particles by looping and leave a “loop” of dislocations around them, as shown in Figure 2.10b. This mechanism was first proposed by Orowan [1, 34, 37, 39, 40] and the change in τc with the particle radius r, volume fraction f, and shear modulus of the matrix

µm is shown in Equation 2.11 [1].

3 1/2 µ b τ = · m · f 1/2 (2.11) c 2π r

14 (a) (b)

Figure 2.10 Schematic illustration of (a) a dislocation bowing around impenetrable precipitate particles; (b) continuing motion of dislocation with loops of dislocations left around the particles [1].

2.6 Artiﬁcial Aging and Precipitate Hardening in Al-Cu and Al-Cu-Mg Alloys

Artiﬁcial aging is commonly used to strengthen the 2000, 6000, and 7000 series heat-treatable aluminum alloys with strengthening solutes such as Cu, Zn, Mg, and Li. It results in precipitation of secondary phases from SSSS matrices. A conventional aging treatment includes a single stage heat treatment, e.g. T6 temper, after solution heat treating and rapid quenching.

In the 7000 series Al-Zn-Mg alloys, a two-stage aging treatment is commonly used to enhance the strength over that of single stage. In this case, the low temperature aging promotes precipitation 0 of denser GP zones, and provides more nucleation sites for the strengthening η phase (MgZn2) to precipitate upon heat treating [41]. This section reviews the precipitation in Al-Cu binary alloys and Al-Cu-Mg alloys.

2.6.1 Precipitation and Hardening during Aging of Al-Cu Binary Alloys

The strengthening phenomenon in Al-Cu binary alloys during isothermal aging is a classical example of decomposition of SSSS. During isothermal aging, the decomposition has been shown in

15 many studies [42–50] to occur by the precipitation of transition phases according to of the following sequence:

00 0 αSSSS → α1 + GP zones → α2 + θ → α3 + θ → α4 + θ (2.12)

The metastable solvus for each phase is shown schematically in the Al-rich corner of the Al-Cu binary phase diagram in Figure 2.11 [51], with the driving force indicated as shown in Figure 2.12 [20]. The GP zones in Al-Cu binary alloys were first discovered by Guinier [52] and Preston [53] in Al-4Cu (wt%) binary alloys naturally aged for 6 months. Using different X-ray techniques, they found that small disk-shaped zones formed from aggregation of copper atoms on matrix {100} planes. Other metastable phases, θ00 and θ0, can also precipitate out of the SSSS alloy matrices without first forming GP zones.

Figure 2.11 Al-rich corner of the Al-Cu phase diagram with metastable solvi of GP, θ00 and θ0, and the equilibrium θ in bold [51].

16 Figure 2.12 Schematic free energy diagram for Al-Cu binary alloys [20].

The disc-shaped GP zones in Al-Cu alloys consist of one or more layers of Cu atoms on the Al {100} planes. A single-layer GP zone is shown schematically in Figure 2.13a [54]. The θ00 phase contains two layers of Cu atoms separated by three layers of Al atoms, as schematically shown in Figure 2.13b [54]. Both are observable using high-resolution electron microscopy (HREM) (see Figure 2.14 [49]). An additional intermediate stage (B) also shows the transition between GP zones and θ00. Both GP zones and θ00 are fully coherent with the Al matrix and provide potent strengthening eﬀects.

The θ0 and θ phases, on the other hand, have diﬀerent structures than the Al, as shown in Figure 2.15. The θ0 (Figure 2.15a [20]) has a tetragonal structure with an approximate composition 0 0 of Al2Cu. A TEM micrograph of the θ is shown in Figure 2.15c [55]. Since a of θ is similar to a of Al and c is diﬀerent, θ0 tends to form thin plates on Al {001} planes with coherent interfaces on the (001) and semicoherent interfaces on the orthogonal faces.

The θ phase Al2Cu has a complex body-centered tetragonal structure (Figure 2.15b [20]) and does not have a good matching to the Al matrix. Thus, it usually forms coarse particles and can be observed using SEM, as indicated in Figure 2.15d [57].

The hardening behaviors of Al-Cu alloys with Cu contents from 2 to 4.5 wt% aged at 130 and 190 ◦C are shown in Figure 2.16 [20, 42]. The aging of Al-Cu alloys at 130 ◦C exhibits two-stage hardening, as shown in Figure 2.16a for alloys with Cu content more than 2 wt%. At the early

17 (a) (b)

Figure 2.13 Schematic illustration of (a) GP zone and (b) θ00. After Sato et al. [54]. stage of aging, the formation of GP zones is the cause of hardening. The hardness then comes to a plateau for an extended amount of time before secondary hardening occurs due to the formation of θ00. As aging continues, both θ00 and θ0 continue to form and hardness continues to increase to a peak value. Continued aging results in softening of the alloy, during which θ00 disappears and θ0 becomes the predominant phase. The softening is caused by the loss of coherency from fully coherent θ00 to semi-coherent θ0. For aging at 190 ◦C, the alloys are above the GP zone solvus (see Figure 2.11), making GP zone formation energetically unfavorable. As a result, only single-stage hardening is seen, with θ00 the ﬁrst metastable phase to precipitate out of the SSSS Al matrix.

2.6.2 Precipitation and Hardening during Aging of Al-Cu-Mg Alloys

Alloys containing Cu and Mg are the primary basis for the 2xxx wrought alloys and 3xx cast alloys. The equilibrium phase ﬁelds with Cu and Mg are shown in the section of an isothermal (190 ◦C) ternary phase diagram in Figure 2.17 [55]. Alloys with certain Cu and Mg contents (e.g.

2024) lie in the equilibrium α+S phase ﬁeld where the S phase has a composition of Al2CuMg. Little characterization work has been done for alloys in the α+S+T ﬁelds [51]. Note that the

T-phase (Al6CuMg4) described here is diﬀerent from the T-phase (Al20Cu2Mn3) mentioned later in 2024 alloys.

18 Figure 2.14 High resolution electron micrographs of GP zone (A) and θ00 (C). Intermediate stage between GP zone and θ00 is also observed (B). Image was taken at B={011}. After Karlik et al. [49].

The strengthening sequence and microstructural development are originally considered to be similar to that of the Al-Cu binary alloys. The ﬁrst precipitates, named GPB zones, were ﬁrst proposed by Bagaryatsky to be closely related to S phase and fully coherent with the Al matrix [58, 59]. In fact, the precipitation sequence was believed to be similar to that in Al-Cu binary alloys [58, 59]:

00 0 αSSSS → α1 + GP B zones → α2 + S → α3 + S → α4 + S (2.13)

However, only two distinctive phases, GPB2 and S’, were identified using X-ray diffraction in the early work of Silcock [60]. In her study, the GPB zones showed no X-ray streaks, and no S” phase could be distinguished. The structural differences between S’ and S also could not be distinguished. In addition, Silcock mentioned that a GPB2 zone was observed and was believed to

19 (a) (b)

Figure 2.15 Schematic illustration of the crystal structures of (a) θ0 and (b) θ [20]. ◦: Al; •: Cu. The representative TEM (BF at B={011}) micrograph [55] for θ0 and SEM backscattering electron image [56] for θ phase (bright spots) are shown in (c) and (d), respectively. be closely related to the GPB zone. In fact, at peak aging, only a rod-like precipitate and a lath S phase, were observed, as shown in Figure 2.18 [61].

In some studies, the GPB zones were considered to be the rod-like precipitates mentioned earlier [43, 55]. Their crystal structure, however, is not fully understood. A low-energy conﬁguration of a layered Cu-Al-Mg-Al-Cu (Al2CuMg) model was proposed by Wolverton [62] using ﬁrst principle calculations, and the structure is reconstructed in Figure 2.19, with its a and b equal to that of Al (4.04 A),˚ and its c-axis twice as large (8.08 A).˚ The structure of S phase is generally agreed to be

20 (a)

(b)

Figure 2.16 Vickers hardness changes of various Al-Cu alloys (in wt%) during artiﬁcial aging at (a) 130 ◦C and (b) 190 ◦C [20, 42]. orthorhombic [63]. It has a space group of Cmcm, with lattice parameters are a = 4 A,˚ b = 9.23 A,˚ and c = 7.14 A.˚

The S phase nucleates on the Al {120} habit plane and grows along the <001> [60, 61, 63]. It is also suggested that S phase can coarsen along the <210> direction through a ledge mechanism [61].

21 Figure 2.17 Isothermal ternary phase diagram of Al-Cu-Mg at 190 ◦C. The bold line represents the α+S phase boundary at 500 ◦C [55].

This gives S phase an orientation relationship shown as follows:

< 100 >Al//< 100 >S; < 021¯ >Al//< 010 >S; < 012 >Al//< 001 >S (2.14)

This orientation relationship results in 12 variants with the Al-matrix [63]. These twelve variants are listed in Table 2.1, and the SAED pattern typically observed at B = <001> is shown in Figure 2.20.

2.6.3 Initial Rapid Hardening in Al-Cu-Mg Alloys

The hardening of Al-Cu-Mg based alloys exhibits a similar two-stage aging behavior to that of the Al-Cu binary alloys, as shown in Figure 2.21. Two-stage hardening is clearly observed in alloys aged at 170 and 190 ◦C, whereas alloys aged at 240 ◦C show a single stage hardening. The initial hardening typically accounted for approximately 50% of the total hardness increase in alloys with Cu/Mg close to 1 (atomic ratio) [13, 64]. In Al-Cu alloys, the ﬁrst hardening reaches a plateau after a day of aging at 130 ◦C, whereas the ﬁrst hardening plateau in Al-Cu-Mg alloys occurs rapidly, usually within one minute. This initial rapid hardening was also reported in various studies [12, 13, 51, 57, 61, 63, 65–74]. The Al-Cu-Mg alloys also show potent strengthening under natural aging and a hardening plateau can be observed after one day at RT [57, 67].

22 (a)

(b)

Figure 2.18 Bright-ﬁeld image taken near B={001} of 2024 Al-Cu-Mg alloys at peak aging condition: (a) S’ lath and (b) GPB zones (arrows). Image (a) from Shih et al. [61].

23 Figure 2.19 Reconstructed crystal structure of a GPB zone with Al2CuMg as proposed by Wolverton [62].

Table 2.1 – Twelve variants expected between the S phase and the Al-matrix [63].

Variant Equivalent orientation relationship ¯ 1 [100]Al//[100]S, [021]Al//[010]S, [012]Al//[001]S ¯ ¯ 2 [100]Al//[100]S, [021]Al//[010]S, [012]Al//[001]S ¯¯ ¯ 3 [100]Al//[100]S, [012]Al//[010]S, [021]Al//[001]S ¯ ¯ ¯¯ 4 [100]Al//[100]S, [012]Al//[010]S, [021]Al//[001]S ¯ 5 [001]Al//[100]S, [210]Al//[010]S, [120]Al//[001]S ¯ ¯ 6 [001]Al//[100]S, [210]Al//[010]S, [120]Al//[001]S ¯ ¯ ¯ ¯ 7 [010]Al//[100]S,[201]Al//[010]S,[102]Al//[001]S ¯ ¯ ¯ 8 [010]Al//[100]S,[201]Al//[010]S,[102]Al//[001]S ¯ ¯ ¯¯ 9 [001]Al//[100]S, [120]Al//[010]S,[210]Al//[001]S ¯¯ ¯ 10 [001]Al//[100]S,[120]Al//[010]S, [210]Al//[001]S ¯ 11 [010]Al//[100]S,[102]Al//[010]S, [201]Al//[001]S ¯ ¯ 12 [010]Al//[100]S, [102]Al//[010]S,[201]Al//[001]S

2.6.4 Proposed Mechanisms for Initial Rapid Hardening in Al-Cu-Mg Alloys

It was previously believed that the formation of GPB zones was the cause for the early stage hardening in Al-Cu-Mg alloys [32, 42, 58–61]. However, no GPB zones were reported during the early stage of aging.

Reich et al. [65] and Ratchev et al. [75–77] showed S” formation on pre-existing dislocations during early stage of aging through TEM observations. Nagai et al. [66] also showed that vacancy- Mg complexes could diﬀuse rapidly in Al, thus facilitating the rapid migration of solute atoms to dislocations. Although subsequent research [78–82] showed preferential precipitate formation on dislocations, no direct correlations were made to the strengthening eﬀects due to solute-dislocation segregation.

24 Figure 2.20 SAED patterns of S-phase at B=<001>.

To explain the initial rapid hardening, a solute clustering theory was proposed by Ringer and Polmear, who considered it as an exaggerated or super form of solid solution strengthening [33, 83– 85]. By using a variety of analytical techniques such as DSC, FIM, NMR, and Atom Probe, it was shown that solute clustering was prominent in the early stages of aging [12, 13, 51, 67–71, 73, 74, 86– 93]. The solute clustering only required short range diffusion of solute atoms, which would be more fitting to the short times allowed for solute to diffuse during rapid hardening. It was speculated that dislocation-cluster interactions are the principal cause of the initial rapid hardening.

2.6.5 Clusters, GPB zones, S’, S Phases and Precipitation Sequence in Al-Cu-Mg Alloys Revisited

The determination of the crystal structure of GPB zones has been a challenge [62, 72, 94–101]. In fact, DSC studies [61, 67, 68, 70, 71, 102] suggested that at early stage of SSSS decomposition, only one exothermic peak was present. An example of the DSC studies is shown in Figure 2.22 [67] where an Al-1.2Cu-1.2Mg (at%) alloy was naturally aged at diﬀerent times up to a week.

There are four regions observed in Figure 2.22, with I and III exothermic and II and IV endothermic. In addition, as natural aging time increases, the peak at region I becomes weaker and eventually disappears after natural aging for 10 h or longer.

25 Figure 2.21 Aging of 2024 alloys at 170, 190, and 240 ◦C. After Shih et al. [61].

Region II can be interpreted as the dissolution of precipitates from region I, and region IV the dissolution of those from region III. In short, only two distinct peaks (exothermic peaks) are observed.

The first exothermic peak (Region I) at lower temperature may correspond to the formation of clusters and/or GPB zones. It is suggested that the rod-like GPB zones, as mentioned in Section 2.6.2, are formed through continuous formation of large solute clusters [74]. The appearance of these rod-like GPB zones corresponds to the secondary hardening of the alloy during artificial aging [61, 72, 74], which suggests that GPB zones provide potent strengthening effects to the alloy.

However, the distinction between GPB zones and S” is not as clear as GP zones and θ00 in Al-Cu. Silcock reported two types of GPB zones, namely GPB (no diﬀraction streaks) and GPB2 (with diﬀraction streaks), in the early stage of aging Al-Cu-Mg alloys [60]. Recent evidence on the formation of solute clusters [67, 68, 70, 71, 75, 76] during early stage of aging suggests that solute clusters are the “GPB zones”, whereas the rod-like GPB zones are the GPB2/S”.

The second exothermic peak (region III) occurring at higher temperature suggests the formation/precipitation of S’ and S phase. Earlier observations by Silcock [60] and Gupta [103] suggest that S’ phase is only a slightly distorted version of S phase. The precipitation sequence is revised as follows: [33, 74, 83]

26 Figure 2.22 An example of DSC curves (heating rate = 10K/min) of an Al-1.2Cu-1.2Mg (at%) alloy after naturally aged for several time intervals. From Starink et al. [67].

0 0 αSSSS → α1 + clusters → α2 + clusters + GP B → α3 + GP B + S → α4 + S /S (2.15)

2.7 Novel Heat Treatment: Interrupted Aging

In certain alloys, such as the 7000 series Al-Zn-Mg alloys, a two-stage, low-high temperature “double aging” treatment was commonly used to achieve desirable strength levels [104]. It is believed that the low temperature aging treatment more densely distributed GP zones, which act as precursors for the strengthening precipitate η0. Although partial reversion was typically observed [105–107], it still gives an overall increase in GP zones once the alloy was heated to a higher aging temperature, which altered the η0 distribution and therefore the strengthening.

Recently, Lumley and co-workers [4–11] introduced a novel aging treatment called “interrupted aging”. Unlike double aging in which the alloys are given a low-high temperature aging, it contained a high-low-high multi-stage of temperature variations. The overall thermal history of the interrupted aging is shown schematically in Figure 2.23 [7].

A variety of heat-treatable alloys, both wrought and cast, were subjected to such interrupted aging treatments. The results showed that all the alloys responded positively to the multi-stage interrupted aging process, which was shown to enhance the hardness, mechanical strength, and fracture toughness of heat-treatable aluminum alloys compared to a conventional T6 treatments, with minimal eﬀects on ductility. The improvements were attributed to the secondary precipita-

27 Figure 2.23 Schematic illustration of a T6I6 interrupted aging treatment. After Lumley et al. [7]. tion occurring during the interrupted aging and its eﬀects on microstructure at the peak aging conditions.

2.7.1 Secondary Hardening and Precipitation

It was found that in Al-5.9Zn-2.9Mg (wt%) alloys, continuous microstructural developments occurred after solution heat treatment at 465 ◦C if the alloy was aged briefly at 180 ◦C, and then aged at 135 ◦C for 24 h [105]. Lath-shaped η precipitates were observed if the alloy was aged at 180 ◦C for 24 h (Figure 2.24a). However, if the alloy was aged at 180 ◦C for 1 to 3 h and then aged at 135 ◦C for 24 h, both lath η (arrows in Figure 2.24b) and fine η0 (between the lath η) could be observed. Löffler et al. [106] pointed out that, in supersaturated Al-Zn-Mg alloys artificially aged briefly at elevated temperatures, the hardness continued to change after extended natural aging at lower temperatures. This phenomenon is now termed “secondary precipitation” [108]. In the interrupted aging treatments, the low temperature aging was believed to lead to such secondary precipitation and have significant impacts on the subsequent microstructural development [4, 6–10].

2.7.2 Interrupted Aging

Interrupted aging includes a solution heat treatment and water quenching, followed immediately by high temperature underaging for a period of time less than the peak aging time, and an interruption to a lower temperature (60-100 ◦C) at which aging continues for a long period of time from a few hours to several weeks. Afterward, the aging is resumed at an elevated temperature.

28 (a) (b)

Figure 2.24 TEM micrographs of a Al-5.9Zn-2.9Mg (wt%) solution heat treated at 465 ◦C and then (a) directly quenched to and aged at 180 ◦C for 24 h (×36,000) and, (b) directly quenched to and aged at 180 ◦C for 1 h and aged at 135 ◦C for 24 h (×99,000). After Embury et al. [105].

This process has been designated as T6IX, where “I” refers to the interruption and “X” the resumed aging condition. A “T6I6” designation, for example, indicates that the alloy was artiﬁcially underaged at elevated temperature (Ta), interrupted to lower temperature Tb for an extended period of time, and re-aged at elevated temperature Tc (usually equal to Ta unless speciﬁed otherwise).

If the heat treatment would end after the interruption at lower temperature, it was denoted as “T6I4”. Additional information could be added after the TXIX condition, as well. For example, “T8I6” indicated that the alloy was cold-worked prior to artiﬁcial aging. “T6I6/150” indicated the alloy was re-aged to a diﬀerent temperature, i.e. 150 ◦C, than the initial aging temperature.

In all heat-treatable aluminum alloys tested by Lumley et al., signiﬁcant improvements in 0.2% yield stress, ultimate tensile stress, as well as toughness were observed in alloys undergoing the T6I6 treatment compared to the single-stage T6. Examples of the aging responses of selected alloys are shown in Figure 2.25 [4]. The stages of aging are described below:

Underaging High temperature aging was interrupted after 0.5 h for 7050 alloys (Figure 2.25c) and 10 h for 8090 alloys (Figure 2.25d) depending on the aging response of the alloys.

Interruption, Secondary Aging, or Low Temperature Dwell Following interruption, aging at a lower temperature, denoted as “secondary aging” or “low temperature dwell”, was used to further increase the hardness. This is believed to be caused by secondary precipitation of clusters/GP zones [7–11].

29 Re-aging The subsequent re-aging at the original aging temperature yielded a diﬀerent aging response, resulting in higher peak hardness than the T6 tempered alloys. The change in hardness was attributed to microstructural changes. An example is shown in Figure 2.26 [4], where the 2014 alloy (Al-4Cu-0.5Mg, wt%) showed a hardness increase of approximately 27% (from

∼150 HV to ∼190 HV ) of total peak hardness after T6I6 treatment, as shown in Figure 2.26a. √ It also showed an increase in fracture toughness of approximately 35%, from 26.9 MPa m √ after T6 to 36.2 MPa m after T6I6. Both of these increases in the 2014 alloy were among the highest observed in the alloys tested. The precipitate distributions after peak hardening to T6 and T6I6 are shown in Figures 2.26b and 2.26c, respectively.

Compared to the T6 condition, the T6I6 treatment exhibited ﬁner, less plate-like precipitates, as well as a denser distribution. It was concluded that the interrupted aging treatment did not appear to result in precipitation of new phases, but rather a re-distribution of the expected phases [4–7].

In all examples, alloys were underaged at elevated temperatures, “interrupted” at a low temperature, typically 65 ◦C, for an extended period of time, and then aging was resumed at the elevated temperatures to peak hardness or beyond. Additionally, in all alloys tested, continuous secondary hardening was observed during the entire period of low temperature aging.

Lumley also investigated the eﬀects of pre-aging strain on interrupted aging using Al-4Cu (wt%) and 8090 alloys [4], a variant of interrupted aging termed “T8I6”. It was found that, by introducing 5% cold work, both yield and tensile strength increased, without any change in elongation. The results of T8I6 treatments on a Al-4Cu (wt%) and a 8090 alloy are shown in Table 2.2 [4]. It was concluded that the advantages of interrupted aging were retained if alloys were cold worked prior to aging.

Table 2.2 – Eﬀect of 5% cold work on tensile properties of a Al-4Cu (wt%) and a 8090 alloy after T8 and T8I6 tempers. After Lumley et al. [4].

0.2% Proof Stress Tensile Strength Elongation Alloys Tempers MPa MPa % T8 242 339 7 Al-4Cu T8I6 265 358 7 T8 414 495 5 8090 T8I6 441 518 5

30 (a) (b)

Figure 2.25 Examples of T6I6 treatments on (a) Al-4Cu (wt%); (b) 6061 (Al-Mg-Si); (c) 7050 (Al-Zn-Mg-Cu) and (d) 8090 (Al- Cu-Mg-Li). 6061 alloys (b) were initially aged at 177 ◦C; 7050 alloys (c) at 130 ◦C, and 8090 alloys at 185 ◦C. Low temperature aging was conducted at 65 ◦C for all alloys. The vertical broken line in each diagram represented secondary hardening after low temperature aging. After Lumley et al. [4].

31 (a)

(b) (c)

Figure 2.26 Interrupted aging of a 2014 (Al-4Cu-0.5Mg, wt%) alloy: (a) the hardness responses during T6, T6I6, and secondary aging at 65 ◦C; (b) bright-ﬁeld TEM image of a T6, peak-aged 2014 alloy; (c) bright-ﬁeld TEM image of a T6I6, peak-aged 2014 alloy. After Lumley et al. [4].

2.7.3 Evolution of Clusters and GP zones during Low Temperature Dwell

These improvements in mechanical properties from the T6I6 treatment were believed to be related to the secondary precipitation during low temperature aging [4–7]. Subsequent studies showed that, for alloys undergoing interrupted aging treatments, the secondary precipitation during low temperature aging promoted re-distribution of GP zones [7–11]. It was shown that the number of GP zones increased after the low temperature dwell in 6061 (Al-1.0wt%Mg-0.6wt%Si) alloys that underwent T6I4 treatment.

32 The GP zones in 6061 alloys were believed to be the precursor for subsequent strengthening by the β00 phase. TEM micrographs of the 6061 alloy from SHT to peak aged at T6 and T6I6 are shown in Figure 2.27. It was shown in bright-ﬁeld TEM images that the GP zone (dark dots) distributions had increased from underaged at 177 ◦C for 20 minutes (Figure 2.27b) to the T6I4 condition (177 ◦C for 20 minutes + 65 ◦C for two weeks, Figure 2.27c) [8]. The subsequent T6 and T6I6 peak-aged microstructures are shown in Figures 2.27d and 2.27e, respectively. The precipitates in the T6I6 peak-aged condition were seemingly smaller and more densely distributed than in T6, which was attributed to the re-distribution of GP zones after secondary aging at 65 ◦C for two weeks.

3DAP studies [7, 9, 10] showed that during the initial high temperature underaging, both solute clusters and GP zones were formed. Solute clusters were deﬁned as irregular in shape and contained about 10 to 30 solute atoms, whereas GP zones were deﬁned as more symmetrical, or nearly spherical, and contained more solute atoms. The 3DAP results of total numbers of precipitates and precipitate fractions are shown in Table 2.3 [8–10]. The T6I6/150 and T6I6/177 designations indicate that the alloy was re-aged to 150 and 177 ◦C, respectively.

It was also shown that, after T6I4 treatment, about a 40% increase in total precipitate volume fraction/number density was observed (up from 4.8×1024/m3 in the underaged condition to 6.7×1024/m3 after T6I4) [9]. The authors believed that the secondary aging at 65 ◦C promoted secondary precipitation of clusters and GP zones. Also, the relative ratio of clusters/GP zones had also lowered from 60/40 to 37/63 [9]. The authors therefore concluded that, in addition to the secondary precipitation, coarsening of the existing clusters into GP zones also occurred during the secondary aging at 65 ◦C.

The total number of precipitates (clusters + GP zones + β00) at peak aging had also increased from 0.8 (T6) to 1.1 (T6I6) ×1024/m3, with signiﬁcant change in precipitate ratios, as shown in Table 2.3. The β00 observed after T6I6 contained fewer solute atoms than that after T6, as well as the total numbers of atoms, which indicated a smaller β00 in the peak aged T6I6 alloy. The additional strengthening from T6I6 was then attributed to the change in β00, increase in total numbers of precipitates, and changes in precipitate fraction of the strengthening β00, GP zones, and clusters (60/30/10%, respectively) compared to the conventional T6 (47/18/56%).

An additional re-aging temperature of 150 ◦C was used to compare to the original aging temperature of 177 ◦C [10]. It was shown that, by re-aging at 150 ◦C, although the toughness increased, no added strengthening was observed (also see Table 2.3). 3DAP analyses showed that re-aging at 150 ◦C increased the fraction of clusters, which resulted in the lower fractions of GP zones and β00.

A double aging treatment with low-high (65-177 ◦C) temperature variants was also employed to compare its eﬀect on age hardening of 6061 alloys. It was found that double aging did not have any measurable eﬀects on further strengthening the 6061 alloy compared to a standard T6 treatment [9].

33 (a)

(b) (c)

(d) (e)

Figure 2.27 Bright-ﬁeld TEM images of a 6061 alloy at B = <001>Al (a) SHT; (b) underaged at 177 ◦C for 20 minutes; (c) after T6I4 (underaging + 65 ◦C for two weeks); (d) T6; (e) T6I6. The dark dotted features in images (b) and (c) were GP zones. Images were taken at the same magniﬁcation. After Buha et al. [8, 9].

34 Table 2.3 – 3DAP analyses of 6061 alloys tempered at UA (underaged condition, 20 minutes @ 177 ◦C), T6I4 (UA + 2 weeks @ 177 ◦C), T6 (177 ◦C), T6I6/150, and T6I6/177. After Buha et al. [8–10].

Total numbers 0.2 proof Damage of precipitates Clusters GP zones β00 stress Tolerance √ Temper 1024/m3 Morphology Irregular Spherical Elongated MPa MPa m Fractions (%) 60 40 n/a Solute atoms 10-30 30-100 n/a UA 4.8 n/a n/a Total atoms 40-100 100-300 n/a Size 1-2 nm 2-4 nm n/a Fractions (%) 37 63 n/a Solute atoms 10-30 30-100 n/a T6I4 6.7 n/a n/a Total atoms 35-120 145-430 n/a Size 1-2 nm 2-4 nm n/a Fractions (%) 35 18 47 Solute atoms 10-30 30-100 100-700 T6 0.8 311 30.5 Total atoms 45-75 220-240 1200-2500 Size 1-2 nm 2-4 nm (2-5 nm)2× > 15 nm Fractions (%) 54 23 23 Solute atoms 10-30 30-85 85-170 T6I6/150 1.5 302 40.6 Total atoms 15-52 100-300 300-460 Size 1-2 nm 2-3 nm (2-4 nm)2×11 nm Fractions (%) 10 30 60 Solute atoms 10-30 30-100 120-360 T6I6/177 1.1 335 37 Total atoms 45-50 180-430 850-1300 Size 1-2 nm 2-3 nm (2-4 nm)2×11 nm

35 2.8 Basic Principles of Small-Angle X-ray Scattering, SAXS

Small-Angle X-ray scattering (SAXS) is used to determine the size and volume fraction of precipitates formed in alloys during the early stages of aging, as well as after T6I4 treatments. The basic principles of SAXS, a brief description of the instrumentation, and interpretation of SAXS intensity are shown in the following sections.

2.8.1 Small Angle Scattering

Consider an incident beam, with a given wavelength of λ, scattered by a particle with a radius R. The scattered beam, as a function of “momentum transfer”, q, can then be measured at an angle of 2θ, as shown in Figure 2.28. q is given by Equation 2.16 [109]. The scattering intensity, for a spherical particle with a radius of R, is proportional to a function F(qR), as given by Equation 2.17 [109].

4π sin θ q = (2.16) λ 2 2 2 sin (qR) − qR cos (qR) I = (∆ne) V 3 (2.17) (qR)3 where ∆ne is the electron density diﬀerence between the particle and surrounding medium, and V is the volume of the particle.

Figure 2.28 Schematic illustration of an energy beam with a wavelength of λ being scattered by a particle with size 2R.

36 The relations between intensity ratio, I/I0, where I0 is the maximum scattering intensity equal 2 2 to (∆ne) V , and the transfer momentum q can therefore be plotted with varying R, as shown in Figure 2.29.

Figure 2.29 Schematic illustration of how the scattered intensity ratio I/I0 varies with the momentum transfer q for various particle sizes.

This is useful in determining the size of small features because they scatter the beam to a larger angle of 2θ and allow for easier measurements. The intensity of the measured q at a given 2θ is dependent upon the amount, or volume fraction of the particles with size of 2R. One can show that F(qR) falls from a maximum of unity at q = 0 to a value of 1/2 when qR = 1.815. Thus, when 2R = 3.63/q, the intensity falls to 50% of the maximum and shows the inverse relationship between particle size and scattering angle. By applying Equation 2.16 and if the scattering angle is small so that sin (θ) ≈ θ, the relationship between particle size 2R and scattering angle can be given by Equation 2.18:

3.63 λ 2R = D = (2.18) 2π 2θ

2.8.2 Interpretations of SAXS Intensity

Because of the thickness diﬀerences between each sample scanned, the raw intensity collected must be normalized according to their respective absorption coeﬃcients. The intensity scattered at small angle is proportional to te−µt, where t is the thickness of the sample and µ is the transmission

37 −1 coeﬃcient for the Cu Kα x-rays [109], which is calculated to be 128.25 cm for pure Al and 138.84 cm−1 for 2024 alloys. As shown in Figure 2.30 [109], the thickness for optimum intensity is 1/µ.

Figure 2.30 Dependence of scattered intensity I on sample thickness t. After Glatter [109].

The normalized SAXS intensity, I(q) (in cm−1), is a function of particle size, geometry, and volume fraction. This relation is expressed in Equation 2.19 [109, 110].

X C I(q) = I + I + I = A· v (R )3 F (qR ) + + I (2.19) N L d i i i q3 d where q is the momentum transfer defined in Equation 2.16, IN is the contribution from nanos- tructural features (1-30 nm), modeled by a distribution of spherical particles with relative volume fraction vi and radius Ri.IL is the contribution from larger scale features (>30 nm), such as grain 3 boundary scattering, modeled by the Porod scattering function C/q , and Id is a constant diffuse scattering due to atomic scale electron density fluctuations. A and C are constants.

2.8.3 Volume Fraction Calculation Using SAXS Intensity

Once the function I(q) is determined by ﬁtting the experimental data, the total volume fraction of particles, f, can be calculated from Equation 2.20 [110], where Q is the integrated intensity calculated from Equation 2.21 [110].

38 Q f(1 − f) = (2.20) 2 2 2π Ie (∆ne)

q Z Q = s q [I(q) − I ] dq (2.21) 2 d

−7 In these expressions, qs/2 is a system-dependent constant and is determined to be 2.70×10 −1 (cm ) for the Anton-Paar “SAXSess” system used in this study. Ie is the Thomson scattering per electron and is equal to 7.94×10−26 (cm2) [110]. ∆n is the electron density diﬀerences between the Al-matrix and the clusters, assuming a composition of Al2CuMg based on the ﬁrst-principle calculation proposed by Wolverton [62]. The electron densities for Al and the clusters can be calculated using Equation 2.22 and are 7.83×1023 and 10.13×1023 (cm−3), respectively.

electrons NDN n = = 0 (2.22) e cm3 M where ne is the electron density of the atom/molecule, N = number of electrons/atom or molecule,

N0 = Avogadro number, D = mass density, and M = molecular weight.

2.9 Scope of Thesis

While secondary hardening after a T6I4 treatment was observed in heat treatable alloys such as the 7050 (Al-Zn-Mg-Cu) [11] and Al-Cu-Mg [12, 13, 64], the anticipated improvements after T6I6 treatments over conventional T6 treatments have not been consistently reproduced [14–17].

Goh et al. followed the interrupted aging using a 6111 alloy [14] at 180/65 ◦C and reported that the hardness of the 6111 alloys tested increased from 128 HV after T6 to 132 HV after T6I6. The authors also tested a variety of underaging times from 15 s to 20 minutes at 180 ◦C and found that, regardless of the underaging time, the peak hardness of 132 HV was achieved as long as it was short, although the kinetics, i.e. time required to reach peak hardness after re-aging, was aﬀected. The interrupted aging T6I6 behavior compared to T6, as well as two underaging times (30 s and 2 minutes) are shown in Figures 2.31a and 2.31b, respectively.

As shown in Figure 2.31, hardening occurred in all 6111 alloys tested after T6I4 treatments. Although the authors claimed that measurable improvements in 6111 alloys were found after T6I6, the said improvements were not obvious because a) an improvement of 4 HV (128 to 132 HV ) was not a signiﬁcant improvement, and b) typical variations of Vickers hardness could easily be

4 HV as mentioned by Ref [12, 64] and this study, and the “improvement” could easily be within the variation of measurements. The TEM micrographs shown by the authors also did not exhibit any signiﬁcant change in the morphology and precipitate distributions as previously shown by Buha et al. [8–10]. Therefore, the work was not convincing that the T6I6 aging treatment had, in fact, been beneﬁcial to further strengthening this age-hardenable alloy.

39 (a) (b)

Figure 2.31 Aging responses of 6111 alloys undergoing (a) T6 at 180 ◦C and T6I6 at 180/65/180 ◦C and (b) T6I6 with underaging times of 20 seconds and 2 minutes. After Goh et al. [14].

Conversely, a study of interrupted aging was done using three Al-Cu-Mg alloys with varying Cu/Mg ratios [15]. The authors investigated a variety of underaging times, interrupted times, and re-aging times at 190 ◦C for the three alloys tested. It was shown that hardness increases up to

10 HV were observed after T6I6 compared to conventional T6. However, a question could be raised to the T6 peak aging time reported by the authors, which was consistently, and significantly, lower than reported times. One example was that for an Al- 4.34Cu-1.37Mg (wt%, similar to that of 2024 alloys), the peak aging time was reported to be 6 h at 190 ◦C. In similar Al-Cu-Mg alloys, the peak aging time would be ∼12 h at 190 ◦C [61]. This could be a serious experimental flaw because at an aging time of 6 h, the alloys were essentially underaged and would continue to harden upon re-aging. After careful review of the hardness, the reported increase corresponded well to the hardness of the T6 tested alloys aged to the total time of underaging + re-aging time. As a result, the reported benefits from T6I6 were once again unconvincing.

The controversy was further worsened through recent reports by Risanti et al. [16, 17], who used two underaging temperatures and times, i.e. 65 ◦C and 105 ◦C from 65 h to two weeks, on 6061

40 alloys [16]. Also, different re-aging temperatures, 150 ◦C and 177 ◦C, were used. It was shown that the underaging time played a significant role in controlling the extent of hardening achieved during low temperature aging, while toughness remains relatively unaffected. However, for all conditions, the strength and toughness did not show any substantial improvements, in contrast to reports by Lumley and co-workers [4–11]. Similar examinations were performed on 2024 alloys and, again, no improvements were observed after T6I6 at 180/(RT or 65 ◦C)/180 ◦C [17].

If real, the interrupted aging could be beneficial to alloy designers. However, because of the inconsistencies in reported results, the question regarding the reproducibility of interrupted aging T6I6 arose. Consequently, further examination of this heat treatment would not only be important for academic purposes, but for industrial applications as well. Therefore, the prime emphasis of this study was to examine the effectiveness of interrupted aging through carefully controlled heat treatments and advanced characterization. Particular emphasis was paid to 2024 alloys (Al-Cu-Mg) because of their broad commercial use. Their ability to substantially harden at room temperature, rapidly hardening upon artificial aging, and response to cold work prior to aging [3] make them an attractive candidate for new heat treatment design.

The cause for the rapid hardening, as well as secondary hardening after T6I4 treatments in Al- Cu-Mg alloys had been attributed to the formations of solute clusters [12, 13, 64, 73, 92]. However, the strengthening mechanisms involving clusters have not been clearly shown in the literature. As such, this study would allow for examination of the fundamental and physical nature associated with solute cluster hardening. The signiﬁcance of this study is to provide certain physical properties of these clusters for subsequent simulation and modeling calculations by alloy designs.

Experiments were planned according to the original interrupted aging by Lumley et al. [4–13]. Since Lumley et al. also concluded that the advantages of interrupted aging were retained if alloys were cold worked prior to aging [4], namely the T8I6 treatment, two variants of the interrupted aging with cold work introduced prior to underaging, as well as prior to secondary aging, were introduced. The latter is termed “T6I8”. This allows for examination of the eﬀects of mechanical cold work on interrupted aging. In addition, upon re-aging, two heating rates were used to examine their eﬀects on retaining the secondary hardening after low temperature aging.

41 CHAPTER 3

EXPERIMENTAL PROCEDURE

Heat treatments were conducted using both oil bath for short-time aging (less than 10 minutes) and a conventional air furnace for long-time aging. Vickers microhardness was used for testing the aging responses of the alloys, and electrical resistivity was used to examine the kinetics and microstructural changes at various stages of aging. TEM was employed to examine the microstructural evolution. Finally, Small Angle X-ray scattering (SAXS) was used in combination with all the data obtained previously for cluster analysis.

3.1 Sample Procurement and Preparation

The 2024 aluminum alloys used in the study were 2024-T351 1 inch×1 inch square bars provided by Kaiser Aluminum. The bar alloys were homogenized at 495 ◦C for 1 h and quenched to water before being sectioned by a diamond blade saw into 1.5×3×50 mm strips. The strips were manually ground using SiC papers and polished to 1 µm diamond following standard metallographic procedures. The strips were then solution heat treated at 495-500 ◦C for 0.5 h and then water quenched. The SHT samples were then stored in either liquid nitrogen or containers cooled by dry ice if not immediately heat treated or cold worked. The ﬁnal dimensions of the strips were measured using a precision caliper and recorded for resistivity calculations. As-cast Al-4.2Cu-0.4Mg alloys were also prepared using vacuum arc melter.

3.2 Introducing Cold Work

An MTS Alliance RT/100 Tensile Testing machine with a 500 lbf load cell was used to introduce cold work via pulling in tension under a constant strain rate of 1.67×10−3·s−1 for the T8I6 and T6I8 heat treatments. The strain was monitored using a one-inch extensometer, which was calibrated using a standard 1-inch spacer and a precision caliper prior to each test. Approximately 5% tensile elongation was introduced and the load vs elongation curves were recorded and exported as text ﬁles for subsequent analysis.

3.3 Heat Treatment

The heat treatments were based on previous interrupted aging work, with some added variations and consisted of the following conditions: unstrained (T6I6), 5% pre-strain prior to aging (T8I6), and 5% strain after 0.5 h of aging at 177 ◦C (T6I8). Schematics of conventional single stage and the three interrupted aging variants are shown in Figure 3.1. Hardness measurements

42 were made on the same samples as the electrical resistivity measurements at various stages during these heat treatments.

Conventional single-stage T6 aging includes a SHT and water quench (WQ) to room temperature, followed immediately by artiﬁcial aging at 177 ◦C to peak hardness and beyond. The aging was conducted in an oil bath for short times (>10 minutes) and an air furnace for longer times. The T8 studies included a SHT and water quench to room temperature, followed immediately by pulling in tension to 5% strain, and artiﬁcial aging at 177 ◦C. The thermal history is schematically shown in Figure 3.1a. The hardness and resistivity measurements were made immediately after the samples were removed from oil bath or the furnace. Prior to measurements, acetone was used to quickly remove the oil on the samples after aging in the oil bath.

The interrupted aging T6I6 was based on the treatment originally published by Lumley et al. [4]. After SHT and WQ, samples were immediately underaged for 0.5 h at 177 ◦C, followed by a secondary aging at 65 ◦C for two weeks (T6I4). Once the secondary aging was completed, a re-aging to 177 ◦C with two heating rates, “fast” and “slow”, was performed to peak and beyond. The thermal history is illustrated schematically in Figure 3.1b.

After SHT and WQ, samples for T8I6 treatment were immediately stretched to 5% strain, followed by underaging for 0.5 h at 177 ◦C. Afterwards a secondary aging at 65 ◦C for two weeks was conducted. Similarly, once the secondary aging was completed, re-aging at 177 ◦C using the same two heating rates, “fast” and “slow”, was performed. The thermal history is schematically illustrated in Figure 3.1c.

The interrupted aging T6I8 included a SHT and WQ, followed by immediate underaging for 0.5 at 177 ◦C. Afterwards, a 5% stretch was introduced prior to secondary aging at 65 ◦C for two weeks. Once the secondary aging was completed, a re-aging to 177 ◦C with two heating rates, “fast” and “slow”, was performed. The thermal history is illustrated schematically in Figure 3.1d. Note that only 2024 alloys were subsected to T8I6 and T6I8 treatments, and two heating rates upon re-aging.

The “fast” heating rate was performed by placing the samples to a large aluminum block pre-heated to 177 ◦C in the furnace. Thermocouples were placed on the block and the samples, and readings showed that 177 ◦C was reached typically within 2 minutes. “slow” heating rate was controlled by the furnace to be approximately 1 ◦C/minute. This allows for comparing the effect of heating rate on the interrupted aging treatments and, assuming there is a difference, may be useful in understanding hardness variations in larger sections where such differences might be encountered.

43 (a) (b)

Figure 3.1 Schematic illustrations of (a) conventional single stage aging of T6 and T8; (b) interrupted aging T6I6 originally proposed by Lumley et al. [4]; (c) interrupted aging variant “T8I6”; (d) interrupted aging variant “T6I8”.

44 3.4 Hardness Measurements

A LECO microhardness tester was used for the microhardness measurements. Although not a precise way of evaluating the mechanical properties of the alloys, it provides a good measure of the strengthening sequences occurring in the material. All samples were mounted in bakelite and polished to 1 µm diamond using standard metallographic procedures as mentioned previously in Section 3.1.

The tester is equipped with multiple optical magnifying lenses up to 500X, a Vickers pyramid indenter, and a rotatable caliper to measure the diagonals of the indentations. The hardness values are calculated using Equation 3.1 [111]:

P HV = 1854.4 · (3.1) d2 where P is the set-load in gram force and d is the average diagonal spacing of the indent in µm. In this study, loads of 100g for SHT conditions up to 300g at peak aged conditions were used, with the dimensions of the indentation ranging from 40 to 45 µm. Each reported hardness is an average 2 of at least 15 indentations. The unit is abbreviated in HV (kgf/mm ).

The standard deviations for hardness measurements were less than 4 HV , typically ranging from 1-4 HV , for 2024 alloys and less than 7 HV for as-cast Al-4.2Cu-0.4Mg alloys.

3.5 Electrical Resistivity Measurements

The electrical resistivity measurements were made using a custom-built apparatus based on a four-point probe method, as shown schematically in Figure 3.2 [112]. Because of the low resistivity in metals, the unit is displayed in µΩ−cm. The resistivity changes as precipitation behavior changes [61, 86, 113–122] and should be supplemental to hardness measurements as aging proceeds.

The sample strips described above were subjected to a DC current of 0.5 Amperes. The temperature of the samples during these measurements was typically within 1 ◦C of ambient, roughly 25 ◦C. The cross-sectional dimensions were measured and recorded using a precision digital caliper. The distance between the two voltage leads was controlled by a spacer of 18 mm. The voltage drop between the leads was collected and the resistivity calculated using the equation given in Figure 3.2.

Data were collected using a National Instruments DAQ PCI interface with a chassis-fit sig- nal filter/amplifier in conjunction with a terminal block of eight channel inputs and a built-in cold-junction compensation (CJC) for temperature measurements using Type K (chromel-alumel) thermocouples. This apparatus was interfaced with LabView software and the collected data were stored in a Microsoft Excel spreadsheet format for subsequent processing. Tests using strips of 99.9 wt pct aluminum show that this custom-built resistivity setup has standard deviation

45 Figure 3.2 Schematic illustration of the four-point probe method for the electrical resistivity measurements [112]. around 0.002-0.004 µΩ−cm within a strip, and a standard deviation of 0.018 µΩ−cm between the strips. The temperature variation could also aﬀect the measurement and was calculated to be 0.012 µΩ−cm/ ◦C using data from Ref [123].

3.6 TEM Analysis

TEM-based techniques were used for microstructural and crystallographic analyses. Conven- tional techniques of Bright Field (BF), Centered Dark Field (CDF), and Weak Beam Dark Field (WBDF) imaging were used, along with selected area diffraction in order to observe precipitate reflections and to determine orientation relationships of these precipitates with the Al matrix. High resolution phase contrast imaging, HREM, was also utilized in an attempt to image the fine clusters that will be discussed below.

Samples for TEM analyses were prepared as follows. Sections of interest were mechanically ground to thicknesses ranging from 100-150 µm on 600 grit SiC paper. Three-millimeter diameter discs were then punched using a standard punch. These discs were then electrolytically thinned in a solution with 30 vol% nitric acid (HNO3) and 70 vol% methanol (CH3OH) that was cooled by liquid nitrogen to approximately -25 ◦C using a voltage of approximately 20 Volts. This process is conducted in a E. A. Fischione Instruments Electropolishing Unit Model-120 equipped with a light sensor that is used to detect perforation in the disc.

46 Once complete, the samples are removed from the electrolyte and immediately rinsed by dipping sequentially into three separate beakers with methanol. They were removed from the holder and brieﬂy rinsed in methanol and followed by wicking the solution oﬀ on Kimwipes and drying in air. Unused discs, as well as the TEM specimens that were not immediately examined were stored in a freezer kept at -25 ◦C in order to avoid natural aging at room temperature. The stored specimens were taken out of the freezer and put in a desiccator for approximately 30 minutes to avoid contamination from moisture prior to TEM analysis.

Specimens in the as-solution heat treated and quenched condition were specially prepared. Slices of the SHT alloy were ﬁrst thinned to 150 µm foils, and brieﬂy solution heat treated for an additional 10 minutes, quenched, and then stored immediately under liquid nitrogen prior to disc preparation. This is to avoid the relatively fast natural aging in 2024 alloys [67], and the short solution heat treat time is to avoid excess oxidation [124].

Conventional TEM observations were made using a Phillips CM12 TEM with a Tungsten filament operating at 120kV. Images and diffraction patterns were recorded on film which was then transformed into digital images using an Epson Perfection V750 Pro image scanner. Adobe Photoshop was used for scale bar addition, picture area selections and contrast adjustments if necessary.

HRTEM observations were carried out in a Phillips CM200 TEM with a Lanthanum Hexa- boride (LaB6) ﬁlament operating at 200 kV. The CM200 is also equipped with a CCD camera, which captures the micrographs in digital JPEG format.

3.7 Small-Angle X-ray Scattering, SAXS

Preparations of SAXS samples were similar to the TEM specimen. They were carefully thinned to thickness under 100 µm at 600 grit SiC papers. The ﬁnal thickness was measured by a precision digital caliper and recorded for intensity normalization. Dry ice, with sublimation temperature of 195K, was used to keep the aged samples from natural aging during transportation of samples.

Anton-Paar “SAXSess” SAXS system with collimated Cu Kα X-ray, courtesy of the Depart- ment of Mechanical Engineering at the University of Colorado at Boulder, was used in studying the cluster evolution during aging. The collimated line X-ray has a length of 20 mm, and a height of approximately 0.4 mm. Samples with thickness ranging between 80-100 nm were prepared. The x- ray sensitive image plate was placed in the sample chamber and surfaced an 2θ range from 0 to 40◦. The plate was exposed to the scattered X-rays for 30 minutes, and the intensity was retrieved by a Perkin-Elmer “Cyclone” readout system. Calibration to absolute intensity unit, i.e. cross-sectional area per unit volume, or cm−1, was made using a polyethylene standard.

47 CHAPTER 4

RESULTS

Results of both conventional single stage and interrupted aging behaviors of two Al-Cu-Mg alloys, i.e. a 2024 commercial alloy (Al-4.5Cu-1.5Mg) and an as-cast Al-4.2Cu-0.4Mg (in wt%) alloys, are in this chapter. Their nominal compositions, and that of 2014 alloy, are listed in Table 4.1. The inclusion of the Al-4.2Cu-0.4Mg alloys was to further examine the reproducibility of interrupted aging on alloys with similar strengthening solute concentrations to that of the 2014 commercial alloys used by Lumley et al. [4]. TEM analyses were also performed, if necessary, on selected samples.

Table 4.1 – Nominal chemical compositions of 2014, 2024 wrought alloys and as-cast Al-4.2Cu-0.4Mg (wt%)

Alloy Cu Mg Mn Si Cr Zn Ti Fe Al 2024 3.8-4.9 1.2-1.8 0.3-0.9 < 0.5 < 0.1 < 0.25 < 0.15 < 0.5 bal. Al-Cu-Mg 4.19 0.43 ------bal. 2014 3.9-5.0 0.2-0.8 0.4-1.2 0.5-1.2 < 0.1 < 0.25 < 0.15 - bal.

4.1 Single-Stage Aging Behavior of 2024 Alloys

Aging of the commercial 2024 alloys was performed at 177 ◦C for single-stage T6 and T8 treatments, and the thermal histories are schematically illustrated in Figure 3.1a. As described previously, hardness and resistivity changes were used to establish the baseline for interrupted aging tests on the 2024 alloy, a commonly used high-strength aluminum alloys, not previously tested by Lumley et al..

4.1.1 Hardness and Resistivity Changes of 2024 Alloys during T6 and T8 Aging

Hardness and resistivity changes during single stage T6 and T8 aging treatments are shown in Figures 4.1a and 4.1b, respectively. The T6 behavior is represented by the solid squares and T8 by solid triangles. As mentioned previously, the short-term aging (under 30 minutes) was conducted in an oil bath, whereas the longer term aging was conducted in a conventional box furnace. As mentioned previously, the thermal histories of the single stage T6 and T8 were shown in Figure 3.1a.

From Figure 4.1a, it is apparent that about 50% of the total hardness increase during the T6 treatment was achieved in the ﬁrst few minutes of aging. In fact, a hardness plateau in the un-

48 (a)

(b)

Figure 4.1 (a) Hardness and (b) resistivity changes of 2024 alloys during single stage aging at 177 ◦C for determining the T6 and T8 conditions.

49 stretched alloy was then observed which extended to ∼2 h before secondary hardening was observed. This rapid initial aging behavior in 2024 and its Al-Cu-Mg equivalents has been well documented [51, 61, 65, 66, 69, 72, 73, 75]. The corresponding resistivity changes in Figure 4.1b showed a slight increase from 5.46 to 5.61 µΩ−cm, which was previously attributed to the formation of GPB zones [61]. However, recent discoveries discussed earlier suggest that this increase is due to the formation of solute clusters, which has been proposed to account for the initial rapid hardening. This increased resistivity maintained a plateau for about 1-2 hours before a clear drop occurred, which also corresponded to the secondary hardening.

The secondary hardening during the T6 treatment on 2024 alloys started at 2 h at 177 ◦C, corresponding to the beginning of the drop in resistivity. A peak hardness of 150 HV was reached after 24-28 hours, and the value dropped relatively sharply thereafter. As aging proceeded, the resistivity showed an inverse sigmoidal-shaped curve. The overall resistivity descended to 3.97 µΩ−cm after 330 h of aging at 177 ◦C.

On the other hand, the alloy subjected to the T8 treatment did not respond initially to artiﬁcial aging as rapidly as did the T6 samples. After the 5% stretch, a hardness increase from 94 HV (at

SHT) to 132 HV was observed, with a resistivity increase from 5.46 to 5.63 µΩ−cm. The similarity of the increase in resistivity to that of the T6 sample at the initial stages was purely coincidental.

As aging proceeded, the hardness remained relatively unchanged, or with a minimal increase, during the first hour of aging, and showed a significant hardening afterwards. The resistivity showed a slight drop to 5.53 µΩ−cm (0.1 µΩ−cm drop from the as-stretched value) and leveled off during the first hour of aging. This drop is probably due to solute segregation to dislocations, which removed solute atoms adjacent to dislocations after the cold work introduction. The serrated flow observed before the end of 5% stretch (Figure 4.2) also suggests a strong solute-dislocation interaction.

When compared to the T6 treatment, the alloy subjected to T8 treatment showed a higher ◦ and prolonged peak hardness of 170 HV , which was achieved at 12 to 24 h of aging at 177 C. The descending resistivity associated with the age hardening was similar to that of the T6 samples. However, after 330 hours of aging, the resistivity in the overaged alloy was signiﬁcantly higher, at 4.20 µΩ−cm, than that of the T6 alloys, most likely due to the higher dislocation density introduced by the 5% cold work.

4.1.2 TEM Microstructures of 2024-T6 Alloys at Initial Plateau and Peak Hardening

To understand the diﬀerences in the microstructure at diﬀerent points on the hardness curve, TEM micrographs after SHT, 177 ◦C for 0.5 h, and 177 ◦C for 24 h are shown in Figures 4.3, 4.4 and 4.5, respectively. It was anticipated that at the early stages of aging, the hardening would be attributed to small-scale clusters, whereas at peak hardness, S’ phase would be the predominant strengthening phase.

50 Figure 4.2 Engineering stress-strain curve of as-solution heat treated (SHT) 2024 alloy during the 5% stretch.

In the early stage of aging, despite the increase in hardness and resistivity, there were no observable reflections from the clusters/precipitates. Subsequently, the bright field image at the two-beam condition shown in Figure 4.4 revealed no strain fields from the clusters/precipitates. In fact, the TEM investigation showed that the microstructures of the alloy appeared to be almost identical to that of a typical SHT condition for aluminum alloys.

The microstructures at peak hardness were predominated by S’ phase based on the diﬀraction pattern shown in Figure 4.5a. The coarse S’ phase was clearly observed, as shown in Figure 4.5b. A closer examination of Figure 4.5d revealed small and densely populated precipitates (GPB zones [55] or S” [63, 70]) in the Al-matrix between the coarse S’ phases. This observation was similar to that of a classic example of peak-aged Al-Cu binary alloys, where both θ00 and θ0 are observed.

Further analyses of the underaged 2024 alloys (0.5 h at 177 ◦C) using BF-TEM (Figure 4.6a) and HRTEM (Figure 4.6b) were performed. Although the mottled background was observed in the bright ﬁeld image at two-beam condition (Figure 4.6a) near B = <001> zone with g = 200,

51 Figure 4.3 Representative BF TEM micrograph of 2024 alloys after SHT. Image was taken at B = <001>. subsequent HREM phase contrast (Figure 4.6b) taken at B = <011> revealed no signs of clusters/precipitates.

This observation, or lack of pictorial evidence of clusters/precipitates, was attributed to the fact that these Cu-Mg solute clusters had a less defined size, shape, composition, degree of order, orientation and structure compared to GP zones, and therefore made observations of these clusters difficult [84]. Kovarik et al. [101] also suggested that the volume fraction of these clusters might be so small that the diffraction intensity from them was below the detectable limit, or possibly overshadowed by diffuse diffraction such as that sometimes produced by aluminum oxide surface films that can form during TEM thin foil preparation.

52 Figure 4.4 Representative BF TEM micrograph of 2024 alloys underaged at 177 ◦C for 0.5 h. Image was taken at a two-beam condition with g = 200 near B = <011>. The large particles indicated by arrows were T-phase, Al20Cu2Mn3.

4.1.3 TEM Microstructures of 2024-T8 Alloys at Initial Plateau and Peak Hardening

A TEM micrograph of as-stretched 2024 alloys is shown in Figure 4.7. Densely populated dislocations are apparent in the as-stretched alloys, as shown in Figure 4.7. The increase in dislocation density after tensile stretching should give rise to both increased hardness and resistivity. After 0.5 h of underaging at 177 ◦C, the microstructure appears identical to the as-stretched state, i.e. no observable precipitates were found (Figure 4.8).

Combining with the previous observation of the minimum increase in hardness, the drop in resistivity, and the serrated ﬂow observed during deformation, it is conceivable that the dislocations

53 (a) (b)

Figure 4.5 (a) Selected area electron diffraction (SAED) pattern of 2024 alloys aged to peak hardness near B = <001> at two-beam condition, g = 220, and reflections from precipitates were clearly seen; (b) BF image using the conditions set in (a); (c) SAED pattern of 2024 alloys at peak hardness at B = <001>; (d) CDF image using the precipitate reflections circled in (c).

54 (a)

(b)

Figure 4.6 (a) BF image of 2024 alloy underaged at 177 ◦C for 0.5 h. Image was set to a two- beam condition with g = 200 near B = <001>. The arrows indicate the T-phase ◦ (Al20Cu2Mn3); (b) High-resolution TEM image of 2024 alloy underaged at 177 C for 0.5 h. Image was taken at B = <011>.

55 Figure 4.7 BF image near B = <112> taken at a two-beam condition with g = 111¯ of as-stretched 2024 alloys. had already interacted with solute atoms in the Al-matrix and segregation of these solutes to dislocations would be plausible.

At peak aging, however, the distribution of S’ phase was quite diﬀerent than that of the T6 peak aged alloys. The S’ appeared to be much smaller and more densely distributed throughout the matrix. To better understand the diﬀerences in distributions of precipitates, further microstructural analyses using two-beam condition were performed and the results are shown in Figure 4.9.

In Figure 4.9b and 4.9c, contrasts of both the entangled dislocations (wavy contrast) and precipitates (planar contrast) are apparent. Centered dark field image using the precipitate reflections revealed the contrast from precipitates that can be differentiated from Figure 4.9b and 4.9c, as

56 Figure 4.8 BF image near B = <011> taken at a two-beam condition with g = 200 of 2024 alloys stretched to 5% and aged at 177 ◦C for 0.5 h. shown in Figure 4.9d. This suggests that the S’ phase in 2024-T8 alloys grew on dislocations as expected. In addition, little to no GPB/S” phases were observed at this condition.

4.2 Interrupted Aging of 2024 Alloys

Interrupted aging (T6I6, T8I6, and T6I8) of 2024 alloys was performed according to the schematics shown in Figures 3.1b, 3.1c, and 3.1d, respectively. Additional variants using two heating rates for re-aging to 177 ◦C after secondary aging at 65 ◦C were introduced to further examine their eﬀects on interrupted aging. The two heating rates, one at approximately 1 ◦C/minute (labeled as “slow”) and the second, relatively fast direct heating to 177 ◦C (labeled as “fast”) were employed in the re-aging step.

57 (a) (b)

Figure 4.9 (a) SAED pattern of 2024-T8 alloys aged to peak hardness at two-beam condition. g = 111¯ near B = <112>; (b) BF image using the condition set in (a); (c) WBDF image of (b) taken at g-3g condition; (d) CDF image using the precipitate reﬂections indicated by the white circle in (a).

58 It was shown that the two cold-worked variants, T8I6 and T6I8, exhibited little to no hardness increase during secondary aging at 65 ◦C. Further, the overall aging response from these two variants was similar to that of the single-stage T8. Although secondary hardening was observed after a T6I4 treatment, it reached a plateau within the ﬁrst week of secondary aging. This was a diﬀerent behavior than that reported by Lumley et al., who showed continuous hardening throughout. Ap- parent softening upon re-heating to 177 ◦C was also observed in the unstretched T6I6 treatment. Compared to T6, subsequent aging to T6I6 showed no additional strengthening.

4.2.1 Interrupted Aging (T6I6) of 2024 Alloys

2024 alloys were underaged for 0.5 h at 177 ◦C, interrupted and subjected to secondary aging at 65 ◦C. The hardening behavior is shown in Figure 4.10. The heat treatment condition at the end of the secondary aging was denoted as “T6I4” [4].

During secondary aging at 65 ◦C, both hardness and resistivity increased in the early stage before reaching a plateau. A recently proposed cluster hardening theory suggested that solute clusters were responsible for the rapid increase in hardness [64]. It is conceivable that secondary precipitation of solute clusters occurs during secondary aging at 65 ◦C and accounts for the increase in both hardness and resistivity. However, it is evident that both hardness and resistivity reached a plateau after about ∼180 h of aging at 65 ◦C. This was diﬀerent than the behavior published by Lumley et al. [4–7] for a 2014 alloy, in which the hardness continued to increase during the whole two-week (∼340 h) treatment (Figure 2.26a).

After secondary aging, the alloys were subsequently re-heated to 177 ◦C (the re-aging step) at the two heating rates mentioned earlier. Hardness and resistivity measurements of the T6 and T6I6 alloys and their respective aging behaviors are shown in Figures 4.11a and 4.11b, respectively.

The “slow” heating rate took approximately 1.5 h to go from 65 ◦C to 177 ◦C. As shown in Figure 4.11a, in the first 0.5 h, a higher hardness was observed compared to that when aged at the faster heating rate, which showed a rapid drop in hardness (“softening”). As the temperature reached 177 ◦C, alloys re-heated at both heating rates showed slightly higher hardness and resistivity values than the values achieved before interrupted aging, as shown in Figure 4.11a and 4.11b. Although for both heating rates the resistivity was slightly higher than that of the T6 at the early stage of re-aging, continued aging into the later stages, up to and beyond peak hardness, did not show any significant increase in hardness in contrast to the results reported by Lumley and co- workers [4–11]. In fact, both resistivity and hardness leveled off to that of a standard T6 treatment in spite of the secondary hardening and resistivity increase that had occurred.

59 (a)

(b)

Figure 4.10 (a) Hardness and (b) resistivity changes of 2024 alloys during secondary aging at 65 ◦C.

60 (a)

(b)

Figure 4.11 (a) Hardness and (b) resistivity changes of 2024 alloys during interrupted aging T6I6.

61 4.2.2 TEM Analyses of 2024 After T6I4 Treatment

Based on the results from the T6I4 treatment of 2024 alloys, it was of interest to characterize any differences in microstructure between the alloy after two weeks at 65 ◦C and after reheating and softening at 177 ◦C. However, TEM analyses revealed no resolvable differences in microstructures under the various conditions, namely SHT+WQ (Figure 4.3), 0.5 h underaging at 177 ◦C, after T6I4, and after reheating to 177 ◦C for 0.5 h, (Figures 4.3, 4.12a, 4.12b, and 4.12c, respectively). Specifically, there were no observable signs of clusters/precipitates. Mottled background was observed after all three aging conditions.

Clearly, this result indicates that the size of whatever is causing the increased hardness and the expected strain fields is below the resolution of the instruments used in this study. It has been speculated that this might be related to the size differences of Cu and Mg atoms relative to Al and the possibility that they have a compensating effect and produce little matrix strain upon forming small clusters.

4.2.3 Interrupted Aging Treatments T8I6 and T6I8

The schematic thermal and cold work histories of the T8I6 and T6I8 were shown previously in Figures 3.1c and 3.1d, respectively. Specifically, a 5% RT stretch was introduced at different points in the process in order to detemine the effect of strain on the interrupted aging behavior. In addition, both “fast” and “slow” heating rates after the low temperature aging were also examined as in the T6I6 study. The hardness and resistivity changes during secondary aging at 65 ◦C for two weeks (∼ 340 h) are shown in Figure 4.13 along with the data for the T6I6 treatments. Clearly, there is much less hardening or increase in resistivity in the alloy that was strained prior to the low temperature age.

The hardness and resistivity after completing the T8I6 and T6I8 treatments are shown in Figures 4.14 and Figure 4.15, respectively, along with the more general T8 data. Unlike in the case of T6I6, the hardness and resistivity of T8I6 and T6I8 alloys showed negligible changes during secondary aging, as shown in Figure 4.13, i.e. all hardness and resistivity changes were within the standard variations of measurements (within 4 HV for hardness and 0.02 µΩ−cm for resistivity). It was expected that, due to the lack of improvements after secondary aging at 65 ◦C, both “fast” and “slow” heating rates would have negligible eﬀects on the hardening behavior of the alloys; this was conﬁrmed by comparing with the standard T8 single-stage aging treatments (Figure 4.14a and 4.15a).

Despite the fact that no major diﬀerences were found in either the diﬀerent heating rates, TEM microstructural analyses were performed after the secondary aging at 65 ◦C, as well as after re-aging at 177 ◦C. Similar to the behavior during the T6I6 sequences, no precipitates were observable at any point before re-aging took place. Since hardness increases were observed at any point for alloys

62 (a)

(b) (c)

Figure 4.12 (a) BF image of a 2024 alloy after a T6I4 treatment (g = 200 near B = <001>); (b) BF image of 2024 after a T6I4 treatment and reheated with “fast” heating rate. (g = 200 near B = <011>); (c) BF image of 2024 after a T6I4 treatment and reheated with “slow” heating rate. (g = 200 near B = <011>).

63 (a)

(b)

Figure 4.13 (a) Hardness and (b) resistivity changes during 65 ◦C dwell for two weeks for T6I6, T8I6, and T6I8 interrupted aging treatment of 2024 alloys.

64 (a)

(b)

Figure 4.14 (a) Hardness and (b) resistivity changes of 2024 alloys during interrupted aging T8I6.

65 (a)

(b)

Figure 4.15 (a) Hardness and (b) resistivity changes of 2024 alloys during interrupted aging T6I8.

66 in either the T8I6 or T6I8 interrupted aging treatments, it was evident that cold work was not complementary to the hardness increases during the low temperature dwell as suggested previously by Lumley et al. [4].

4.2.4 Stretching Behavior of 2024 Alloys in As Solution Heat Treated and Underaged Condition

The stress-strain curves for 2024 in two different conditions strained 5% are shown in Fig- ure 4.16. As is apparent, dynamic strain aging (DSA), as characterized by serrated flow, was observed in the as-solution heat treated material towards the end of the stretch. This indicates the presence of mobile solutes in the SHT condition resulting in solute-dislocation interactions. After underaging at 177 ◦C for 0.5 h, the tensile curve showed a higher yield point than that after SHT and no serrated flow over the strain levels tested. This observation agrees well with the work by Rosen et al. [125], who showed that smooth tensile curves were observed once the alloys were aged. This could be due to the formation of solute clusters/precipitates that effectively “trap” the responsible solutes and make them virtually “immobile”.

4.3 Examination of Interrupted Aging in Al-4.2Cu-0.4Mg Alloy

In the previous sections, it was shown that the interrupted aging treatments on 2024 alloys produced no additional strengthening, contrary to the results reported by Lumley et al.. Therefore, the reproducibility of this “novel” aging treatment seems questionable. To further examine the eﬀects of interrupted aging on Al-Cu-Mg alloys, an alloy closer to the 2014 alloy reported initially by Lumley et al. was selected for interrupted aging studies. Speciﬁcally, a high purity (low Fe and Si) version of 2014, namely a Al-4.2Cu-0.4Mg (wt%), was prepared by non-consumable arc melting, and then solution heat treated before aging at 177 ◦C for single stage T6 and 177-65-177 ◦C for interrupted aging T6I6. Similar to the 2024 study, both hardness and resistivity changes were used to determine the aging response of the alloys and whether the interrupted aging response reported by Lumley et al. could be reproduced.

4.3.1 Single-Stage Aging of Al-4.2Cu-0.4Mg

The hardness and resistivity behavior of the as-cast Al-4.2Cu-0.4Mg alloys are shown in Fig- ures 4.17a and 4.17b, respectively. Signiﬁcantly, the hardness did not show the rapid initial hardening seen in 2024, but did increase continuously from ∼68 HV (SHT) to ∼123 HV (at peak hardness hardness) before softening (overaging). This continuous hardening behavior was similar to that shown for the 2014 alloy by Lumley et al. (Figure 2.26a), but had a slightly smaller overall increase in hardness (∼55 HV vs. ∼ 70 HV for 2014).

67 Figure 4.16 Stress-strain curves of 5% stretching of as solution heat treated (SHT) and underaged for 0.5 h at 177 ◦C (30 UA) 2024 alloys.

The resistivity (Figure 4.17b), on the other hand, was relatively constant for the ﬁrst hour of aging (4.14 µΩ−cm) and then dropped steadily to 3.33 µΩ−cm during the next 160 h of aging in an inverse sigmoidal (S-shape) fashion. The initial constant resistivity could suggest that the formation of clusters/precipitates and the loss of solute supersaturation in the Al-matrix had a counter-balancing eﬀect on resistivity.

The Al-4.2Cu-0.4Mg alloy also exhibited potent strengthening after extended natural aging, as shown in Figure 4.18. Based on the resistivity increase, the hardening mechanism is most likely the same as that in 2024 alloys and is speculated to be the formation of clusters.

4.3.2 Interrupted Aging of Al-4.2Cu-0.4Mg

Similar to the T6I6 experiments on Alloy 2024, the SHT Al-4.2Cu-0.4Mg alloys were ﬁrst underaged at 177 ◦C for 0.5 h, water quenched to room temperature and then subjected to secondary aging at 65 ◦C. Cold work and re-heating rate variations were not employed because of their ineﬀectiveness in the 2024 alloy.

68 (a)

(b)

Figure 4.17 (a) Hardness and (b) resistivity changes of as-cast Al-4.2Cu-0.4Mg alloys during single stage T6 aging.

69 (a)

(b)

Figure 4.18 Natural aging response of the Al-4.2Cu-0.4Mg alloy illustrated by (a) Vickers hardness and (b) resistivity.

The Al-4.2Cu-0.4Mg alloys responded to the secondary aging with a signiﬁcant increase in both hardness and resistivity (Figure 4.19). The increase in resistivity suggested a similar hardening mechanism to the 2024 alloy, which was believed to be the secondary precipitation of solute clusters. The similarity between the increasing trends of hardness and resistivity was an additional indication that this secondary precipitation of solute clusters would result in increasing strength. The increase in hardness reached a plateau before the 12-day secondary aging was completed. This is contrary to the behavior of 2014 alloys reported by Lumley et al. [4], that the hardness continued to increase during the whole secondary aging period (see Figure 2.26a).

70 (a)

(b)

Figure 4.19 (a) Hardness and (b) resistivity changes of as-cast Al-4.2Cu-0.4Mg alloys during secondary aging at 65 ◦C.

71 Two samples from the secondary aging treatment, namely, one heat treated for 2 days to a hardness of ∼102 HV and a second for 12 days to the plateau hardness of ∼116 HV were re-heated to 177 ◦C to complete the T6I6 aging treatment. For both secondary aging times, the material softened (Figure 4.17a) and the resistivity dropped (Figure 4.17b), similar to the behavior in the 2024 alloy. If the resistivity increase corresponded to secondary precipitation of clusters, the drop of resistivity would suggest dissolution of small clusters/precipitates, i.e., cluster reversion, and softening of the alloy. This was in contrast to the ﬁnding of Lumley et al. and once again showed the ineﬀectiveness of interrupted aging on an alloy with similar strengthening solute concentrations to the 2014 commercial alloy.

4.3.3 TEM Analyses of Al-4.2Cu-0.4Mg

TEM analyses of the Al-4.2Cu-0.4Mg were not performed due to time constraints. However, studies such as Ref [126] showed that the strengthening precipitates in Al-4Cu-0.4Mg (wt%) were Cu-rich GP zones and/or θ00 during the early stages of aging. It is conceivable that the formation of GP zones/θ00 is the cause of the initial hardening in the Al-4.2Cu-0.4Mg alloy, whereas the secondary hardening results from cluster formation.

4.4 Summary and Comment on Interrupted Aging of Al-Cu-Mg Alloys

Thus far, the novel interrupted aging T6I6 treatment was found to be ineffective in obtaining improved strength in the Al-Cu-Mg alloys examined. This is in agreement to the findings by Risanti et al. [16, 17], but in contrast to those reported in the original interrupted aging literature by Lumley et al. [4–10]. Although a substantial increase in hardness can be achieved during secondary aging at 65 ◦C, softening occurs upon re-heating to 177 ◦C and no significant improvement to the conventional single-stage T6 behavior is observed in the subsequent aging at 177 ◦C. Although results from resistivity suggested possible reversion, it is unclear as to what had specifically dissolved during the re-aging process.

Lumley et al. [4] suggested that the advantages from interrupted aging could be retained if cold work was introduced prior to the aging cycle. However, our ﬁndings indicated that the introduction of cold work prior to the underaging and secondary aging treatments was uncomplementary, resulting in the ineﬀectiveness of both the T8I6 and T6I8 treatments.

Based on these ﬁndings, the hardening during secondary aging at 65 ◦C is speculated to be caused by the formation of additional solute clusters. It was shown, both experimentally and by simulation, that quenched-in vacancies play an important role in the formation of solute clusters [7, 9–13, 30, 33, 51, 66, 68, 73, 84, 93, 101, 106, 122, 127–129]. This could explain the ineﬀectiveness of cold work on interrupted aging. Dislocations are said to act as vacancy sinks [66], and it is conceivable that the quenched-in vacancies are removed during stretching and, as a result, little to

72 (a)

(b)

Figure 4.20 (a) Hardness and (b) resistivity changes of as-cast Al-4.2Cu-0.4Mg alloys during T6I6 interrupted aging.

73 no cluster formation could take place due to the lack of excess vacancies. This would also explain the lack of initial rapid hardening in pre-stretched alloys.

However, characterizing these solute clusters was diﬃcult. Although recent studies [12, 13, 51, 73, 74, 89, 92, 130, 131] showed promising results using 3D atom probe tomography, no obvious conclusions were drawn in terms of the strengthening eﬀects/mechanisms of the solute clusters. To facilitate better understanding, small angle X-ray scattering (SAXS) was utilized to obtain the size and volume fractions of the clusters and their evolution during initial rapid hardening, as well as interrupted aging. SAXS has been used to characterize microstructures of Al-based alloys [132– 137], and was recently applied to in situ study of cluster formations in a Al-2.5Cu-1.5Mg (wt%) alloy during the initial rapid hardening [69]. The results are described in the following sections.

4.5 Data Interpretation of Small Angle X-Ray Scattering (SAXS)

This section includes the experimental observation, interpretations of SAXS data, and cluster size and volume fraction determinations. The basic principle of small angle scattering and the necessary instrumentation was described previously in Section 2.8.

4.5.1 Obtaining Experimental SAXS Intensity and Converting to Absolute Units

An example of the original scan obtained from a SAXS image plate is shown in Figure 4.21a. A good scan consists of a symmetrical exposure at small angles. A color gradient map of the original scan is shown in Figure 4.21b for better visualization of the contrast gradient. It was then averaged over the height of the rectangle from left to right, which corresponded to scattering angles from 0 to about 40◦. Examples of the integrated intensity vs the scattering momentum transfer q are shown in Figure 4.22 using pure aluminum and 2024 alloys aged at RT for 24 h and 177 ◦C for 18 hours. These intensities were then converted to absolute units using sample thicknesses, absorption coeﬃcients, and a calibration factor based on a polyethylene standard. Examples of the absolute intensities (naturally aged 2024 alloys for up to 24 h) are shown in Figure 4.23.

4.5.2 Fitting of Small-Angle X-Ray Scattering Data

Fitting of the SAXS data is crucial in determining the volume fractions and sizes of the clusters. First, one must calculate the electron density of the clusters. This, however, could be potentially difficult because the exact cluster compositions are unknown [12, 13, 73, 92]. Kovarik et al. [101] recently suggested, based on first principle calculations, that for the clusters to have a potent strengthening effect, the Mg/Cu ratios must be in the range of 1 to 2.2. Furthermore, recent research done by Deschamps et al. [69] also suggested that the clusters formed at the early stage of aging consisted of ∼30 at% of Cu. Therefore, the composition of the clusters is assumed to be

Al2CuMg (S-phase composition) with about 25 at% Cu in the clusters. The variations of Mg and

74 (a)

(b)

Figure 4.21 (a) As-scanned SAXS data from a 2024 alloys aged at 65 ◦C for 22h; (b) the same scan from (a) with colored contrast gradient.

Al were not taken into consideration for simpliﬁcation and, because of the similarity of electron density between Al and Mg, interchanging Al and Mg atoms would not signiﬁcantly vary the overall calculated electron density of the clusters.

An example of the SAXS intensity consisting of clusters with different sizes and fractions is shown in Figure 4.24. The scattered intensity comes from a contribution of clusters with radii of 0.5, 1, 2, and 3 nm, and volume fractions of 0.6, 0.1, 0.1, 0.2, respectively. The contribution of clusters with a certain size is clearly specific to a range of q. Fitting the example with different volume fractions of clusters, i.e. 0.7 for 0.5 nm and 0.1 for 3 nm, will result in a significant deviation from the original SAXS scattering intensity, as shown in Figure 4.25.

The fitting procedures are shown in Figure 4.26. First the absolute intensity data are imported into SigmaPlot software (Figure 4.26a), and Id is determined based on the intensities at the highest q’s, as well as a first estimate of the Porod slope C, as shown in Figure 4.26b. Accordingly, relative volume fractions of 1 nm sized clusters are assumed and fitted to the experimental data through trial and error to find the best fit using Equation 2.19 (Figure 4.26c). Finally, all diameters from 2 to 30 nm are incrementally introduced by trial and error to find the best fit for the remainder of the experimental data (Figure 4.26d). Some minor adjustments are made in the value of A to get a final best fit.

75 Figure 4.22 SAXS intensities for pure aluminum and 2024 alloys aged at RT and 177 ◦C for 18 h.

76 Figure 4.23 Absolute SAXS intensities for 2024 samples naturally aged at RT for times ranging from 0.5 to 24 h.

77 (a) (b)

Figure 4.24 SAXS intensity (bold black curve) of a specimen consisting of clusters with radii of 0.5, 1, 2, and 3 nm, and volume fractions of 0.6, 0.1, 0.1, and 0.2, respectively: (a) the measured data is ﬁrst shown with scattering from clusters with radius of 0.5 nm and volume fraction of 0.6; (b) the following portion is then shown with scattering from clusters with radius of 1 nm and volume fraction of 0.1. Subsequent illustrations are shown for clusters with (c) radius of 2 nm and volume fraction of 0.1, and (d) radius of 3 nm and volume fraction of 0.2.

78 Figure 4.25 SAXS intensity with diﬀerent volume fractions of cluster distributions. The bold black curve consists of clusters with volume fractions of 0.6 (0.5 nm), 0.1 (1 nm), 0.1 (2 nm), and 0.2 (3 nm), and the blue bashed curve consists of clusters with volume fractions of 0.7 (0.5 nm), 0.1 (1 nm), 0.1 (2 nm), and 0.1 (3 nm).

More detailed analyses on the cluster size are shown in Figure 4.27. The “best-fit” data suggest the majority of clusters have a size of 1 nm. Variations of ± 0.2 nm showed a noticeable divergence from the “best fit”, in particular, when oversized features (1.2 nm) were used to fit the raw data (Figure 4.27b. This suggests that the trial and error fitting methodology for cluster size determination can be used to obtain cluster sizes with reasonable accuracy and sensitivity to within 0.2 nm in diameter.

Similarly, cluster volume fraction can also be detailed using the same alloy, as shown in Figure 4.28. Following the analysis of cluster size, the volume fraction can be calculated to be 3% of 1 nm clusters for the “best-fit”. Variations of volume fractions show a noticeable divergence from the “best fit”, in both shape and intensity. Based on such divergences, a variation of less than 0.3 vol% can be estimated. It was concluded that the trial and error fitting methodology for size can then be reasonably utilized for quantitative analyses of volume fractions.

4.6 SAXS Analyses on the Evolution of Solute Clusters during Initial Rapid Hard- ening and Interrupted Aging

To better realize the potential applications of the SAXS technique in understanding the cluster formation/evolution, a trial was conducted using alloys in the following conditions: a) aged at room temperature for 24 h, b) aged at 177 ◦C for 18 h, and c) pure aluminum. The raw SAXS intensity

79 (a) (b)

Figure 4.26 SAXS data fitting (naturally aged 2024 alloys): (a) absolute intensity data; (b) the fitting curve with Id and Porod slope C; (c) 1 nm size clusters with proper relative volume fraction to best fit the raw data was introduced; (d) clusters with sizes from 2 to 30 nm were incrementally fitted to give the best fit to the raw data.

80 (a)

(b)

Figure 4.27 Examples (naturally aged 2024 alloys) of SAXS data fitting with cluster sizes of (a) 0.8 nm and (b) 1.2 nm compared to fitting with 1 nm (blue dash line) with a fixed relative volume fraction.

81 (a)

(b)

Figure 4.28 Examples (naturally aged 2024 alloys) of SAXS data fitting with a fixed cluster size of 1 nm with a volume fraction of (a) 2 vol% and (b) 4 vol% compared to fitting with 3 vol% (blue dash line).

82 was shown previously in Figure 4.22, and the differences were clearly seen between the naturally aged alloy and after 177 ◦C +18 h, particularly in the larger q region (smaller D), where the naturally aged alloy showed a significant intensity increase. Since the shapes of the SAXS intensity curves are dependent upon the physical nature of the clusters in the alloys, quantitative SAXS analyses show that the differences must be due to the differences in the cluster distributions under these conditions.

4.6.1 Evolution of Cluster Formation During Initial Hardening of Naturally Aged 2024 Alloys

To further understand the correlations between the measured SAXS intensities, cluster evolution and alloy hardening, naturally aged 2024 alloys were examined. As mentioned previously in Section 3.7, the collection of each SAXS spectra takes 30 minutes. However, because of the rapid hardening in 2024 alloys, in situ measurements at elevated temperatures can not be made. Therefore, room temperature (25 ◦C) aging is selected for the experiment.

The absolute SAXS spectra of naturally aged 2024 alloys in the ﬁrst 24 h are shown in Figure 4.23. By using the analytical approach discussed here, the size and volume fraction of the clusters were calculated. It was shown that during natural aging, the primary change in clusters was the increasing volume fraction of clusters with 1 nm diameter (Figure 4.29), while the volume fractions of clusters with sizes between 2 to 30 nm appeared to be broadly and rather randomly distributed. The results are summarized in Table 4.2.

Figure 4.29 Relative fractions (vi in Equation 2.19) of 1 to 30 nm size clusters from SAXS analyses of 2024 alloys naturally aged up to 506 h.

The average cluster size/diameter can then be calculated using the weighted method, as shown in Equation 4.1. 1/3 X 3 < D > = vi D i (4.1)

83 Table 4.2 – Quantitative SAXS analyses on the volume fractions and cluster size at hardness plateau of naturally aged 2024 alloys and the respective hardness change. The hardness of as solution heat-treated (SHT) alloys was 94 HV .

time (h) HV ∆HV f1 nm % ftotal % f2−30 nm % 0.5 107 13 0.2 1.4 1.2 1 113 19 0.5 1.7 1.2 2 116 22 1.7 2.9 1.2 5.25 120 26 2.5 3.8 1.3 24 132 38 3.1 3.4 0.3 212 133 39 3.1 3.7 0.6 506 133 39 3.1 3.8 0.7

In the early stage of natural aging, the correlation between the increase in volume fraction f1 nm and hardness appeared to be direct. As hardness reached a plateau, so did both f1 nm and resistivity. However, the variations of the volume fractions of larger clusters, f2−30 nm, are unknown. Although sample preparation could be a potential cause, the experiment was not repeated due to time constraint. In addition, the relationship between resistivity and cluster volume fractions is discussed further in Section 4.6.4.

As noted earlier (see Table 4.2), a hardness plateau was reached after naturally aging of 2024 alloys for 24 h and remained relatively unchanged after extended aging. The hardness increase and volume fraction of 1 nm clusters with aging time up to 24 h are shown in Figures 4.30a and 4.30b, respectively.

(a) (b)

Figure 4.30 (a) Change of hardness (with respect to SHT, 94 HV ) and (b) 1 nm cluster volume fractions changes with aging time (up to 24 h).

84 Thus, it appears that there are clear correlations between cluster evolution and hardness behavior, i.e. there is a clear increase in hardness with increasing volume fraction of 1 nm clusters. This implies that the initial strengthening of the alloys is controlled by these 1 nm size clusters.

4.6.2 SAXS Analyses for 2024 Alloys Aged to Plateau Hardness

The hardness behaviors of 2024 alloys aged at room temperature (25 ◦C), 65, 100, 130, 150, 177, and 190 ◦C are shown in Figure 4.31. Rapid hardening of alloys was observed within the ﬁrst minute of aging at temperatures above 100 ◦C. Hardness plateau can be reached within 24 h at RT and 1∼2 h at 65 ◦C. Overall, the hardness values at the diﬀerent plateaus decrease as aging temperature increases.

Figure 4.31 Hardness evolution of 2024 alloys aged to hardness plateau at temperatures from RT to 190 ◦C.

The strengthening eﬀect with changes in volume fraction of clusters in the 2024 alloy can be established by analyzing the change in plateau hardness of samples aged at diﬀerent temperatures from 25 to 190 ◦C. The normalized SAXS intensities for the respective temperatures are shown in Figure 4.32.

85 Figure 4.32 Absolute SAXS intensities for 2024 alloys aged from 25 to 190 ◦C. Pure Al was also included for comparison and showed very weak diﬀuse scattering (Id), consistent with little or no atomic scale density ﬂuctuations in a pure metal.

A comparison of the SAXS results for alloys aged at 25 ◦C (naturally aged for 506h), 150 ◦C (for 0.5 h), and 177 ◦C (for 0.5 h and 18 h) is shown in Figure 4.33. Significant differences are clearly seen, particularly for 1≤q≤10. Subsequent analyses showed that these differences are the result of volume fraction changes of clusters of 1 nm diameter.

Once again, by using the ﬁtting method discussed previously in Section 4.5.2, the correct function of I(q) can be obtained and the Q, f (volume fraction), and cluster sizes can be calculated using Equations 2.20 and 2.21. Fitting was made for cluster sizes from 1 to 30 nm, and the size distributions are compiled in Figure 4.34. It is clear that the fraction of 1 nm clusters shows a decreasing trend as aging temperature increases, whereas the fractions of clusters from 2 to 30 nm size do not show any trends with aging temperatures.

86 Figure 4.33 SAXS intensities of 2024 alloys aged at RT (25 ◦C) for 506 h, 150 ◦C for 0.5 h, 177 ◦C for 0.5 h, and 177 ◦C for 18 h.

The results, including sample thicknesses and the values of e−µt, are listed in Table 4.3. The average size/diameter of clusters was calculated using Equation 4.1, and their respective relative size distributions are shown in Figure 4.34.

The volume fraction of the 1 nm clusters, f1 nm, decreased with increasing aging temperature, and total volume fraction, ftotal, also decreased with increasing aging temperature. The subtle decrease in the Porod slope also suggests that there are small diﬀerences in the bulk material, possibly changes outside the sensitivity range of SAXS measurements. Overall, a clear trend of cluster hardening was observed, i.e. as fractions of clusters increased, the hardness increased.

87 Figure 4.34 Relative fractions (vi in Equation 2.19) of 1 to 30 nm size clusters from SAXS analyses of 2024 alloys aged to respective hardness plateau at temperatures from RT to 190 ◦C.

Table 4.3 – Quantitative SAXS analyses on the volume fractions and cluster sizes for 2024 alloys aged to hardness plateau at temperatures from 25 to 190 ◦C. Units for Q in 1021 cm−4 and f in vol%. C is the Porod slope, which is due to features 30 nm in size.

◦ −µt T( C) t (µm) e CQ1−nm Qtotal f1−nm ftotal f2−30 nm 25 77.2 0.3424 0.034 2.46 2.97 3.1 3.8 0.7 65 82.8 0.3168 0.03 2.21 2.91 2.8 3.7 0.9 100 83.0 0.3159 0.03 1.95 2.47 2.4 3.1 0.7 130 94.0 0.2711 0.04 1.85 2.83 2.3 3.6 1.3 150 71.4 0.3711 0.027 1.03 1.73 1.3 2.2 0.9 177 51.4 0.4899 0.024 0.83 1.28 1.0 1.6 0.6 190 66.2 0.3989 0.02 0.66 1.47 0.8 1.8 1.0

4.6.3 SAXS Analyses of 2024 Alloys After T6I4 and T6I4-Reheated Treatments

In all interrupted aging variants tested, only the unstretched variant showed hardness increases during secondary aging at 65 ◦C. However, the obvious softening, accompanied by a resistivity drop, made the overall T6I6 treatment ineffective. It was speculated that the softening was due to the reversion of the secondary clusters formed at 65 ◦C. In an effort to confirm this speculation, SAXS analyses were conducted on alloys which had undergone the secondary aging at 65 ◦C, i.e. T6I4 treatment, as well as the T6I4 alloys re-aged to 177 ◦C for 0.5 h (T6I4-reheated). The absolute SAXS intensities of the alloy in these two conditions are shown in Figure 4.36, while the relative volume fractions of underaged (aged at 177 ◦C for 0.5 h), T6I4, T6I4-reheat to 177 ◦C for 0.5 h (T6I4-reheat), and near peak aged (at 177 ◦C for 18 h) are shown in Figure 4.35. The integrated intensity Q and the volume fractions f of different size clusters are tabulated in Table 4.4.

The SAXS data indicate that, compared to samples underaged at 177 ◦C (0.5 h), a higher volume faction of 2.0% for 1 nm size clusters, f1 nm, was present in the T6I4 alloys. This is almost a 1% increase compared to the underaged alloy, and showed that the secondary aging indeed promoted secondary precipitation of clusters of 1 nm diameter. Furthermore, an increase from 1.6

88 Figure 4.35 Relative fractions (vi in Equation 2.19) of 1 to 30 nm size clusters from SAXS analyses of 2024 alloys aged at 177 ◦C (for 0.5 h), T6I4, T6I4 reheated to 177 ◦C for 0.5 h (T6I4-reheat), and at 177 ◦C for 18 h.

Table 4.4 – Quantitative SAXS analyses of the volume fractions and cluster sizes for 2024 alloys aged at 177 ◦C for 18 hours and tempered to T6I4 and T6I4-reheated to 177 ◦C. Units for Q in 1021 cm−4, and f in vol%.

Tempering Q1−nm Qtotal f1−nm ftotal f2−30 nm 177 ◦C-30 min 0.83 1.28 1.0 1.6 0.6 177 ◦C-18 h 0.00 0.66 0.0 0.8 0.8 T6I4 1.61 2.60 2.0 3.3 1.3 T6I4-reheat 1.19 2.19 1.4 2.7 1.3 to 3.3% of the total volume fraction was found, which translated into a volume fraction increase in the larger clusters, i.e. size 2 to 30 nm, f2−30 nm, from 0.6 to 1.3%. These data suggest that, the T6I4 treatment, both secondary precipitation of new clusters and cluster coarsening occurred. It also suggested that the additional strengthening after the T6I4 treatment must be the result of this secondary clustering and coarsening.

On the other hand, a drop in overall volume fraction was observed in the T6I4 alloys after ◦ re-aging at 177 C (T6I4-reheat). In particular, the f1 nm dropped from 2.0 to 1.4%. Based on tests done on naturally aged alloys, in which the 1 nm clusters were predominant and believed to contribute to the observed strengthening, it appears that the decrease in f1 nm directly results in the softening observed. It was noted that f2−30 nm remained unchanged, and was therefore not considered to cause softening.

4.6.4 Relationship Between Electrical Resistivity and 1 nm Clusters

The volume fraction of 1 nm clusters and the resistivity of samples aged to plateau hardness at the various temperatures examined are plotted in Figure 4.37. In addition, results from T6I4 and T6I4 re-aged to 177 ◦C samples are also included.

89 Figure 4.36 SAXS intensities of 2024 alloys aged at 177 ◦C for 0.5 h (underaged), T6I4, and T6I4-reheat to 177 ◦C for 0.5 h (T6I4-reheated).

Changes of resistivity with respect to the as SHT condition, i.e. 5.46 µΩ−cm, were used as the ordinate to better reﬂect the “change” in volume fraction of the 1 nm clusters. By assuming that the amount of clusters present after SHT was negligible, a direct, well-ﬁtted (with R2 = 0.948) linear relationship was then established between resistivity and volume fraction of 1 nm clusters by setting the intercept to zero. In other words, the increases in resistivity observed in both the early stages of aging and after secondary aging at 65 ◦C in interrupted aging treatment are directly related to the formation of the 1 nm clusters. With this relationship, the volume fractions of the 1 nm size clusters could be predicted using resistivity changes.

The volume fraction of 1 nm clusters and the respective resistivity of T6I4 and T6I4-reaged samples are also shown in Figure 4.37. The good ﬁt indicates the decrease in resistivity in T6I4- reaged samples must be related to partial reversion of the 1 nm clusters upon re-heating to 177 ◦C.

90 Figure 4.37 Plateau volume fractions of 1 nm size clusters measured in 2024 alloys plotted against the change in resistivity (with respect to SHT, 5.46 µΩ−cm) at temperatures from 25 to 190 ◦C. The linear relationship (with R2 = 0.948) indicates that the change in resisitivity is directly related to the change in the volume fraction of the 1 nm size clusters. Alloys aged to T6I4 and T6I4 + re-aging to 177 ◦C show a good ﬁt to this relationship, indicating partial reversion of the 1 nm size clusters after re-aging to 177 ◦C. The intercept of the linear regression is set to zero.

This reversion must be largely responsible for the softening observed (see Figure 4.11a and Fig- ure 4.20a) during the T6I6 treatment of Al-Cu-Mg alloys, and is likely the reason for the ineﬀec- tiveness of the interrupted aging treatments.

91 4.6.5 SAXS Analysis of Al-4.2Cu-0.4Mg Alloys

Similar SAXS analyses were attempted on the Al-4.2-0.4Mg alloys, in which the ineﬀectiveness of interrupted aging was also observed, with similar resistivity and hardness behaviors to those observed in 2024 alloys. Samples underaged at 177 ◦C, T6I4, and T6I4-reheated were measured by SAXS, but unexpected streaking was observed in all samples. An example is shown in Figure 4.38a.

(a)

(b)

Figure 4.38 Example of SAXS image plate obtained from the as-cast Al-4.2Cu-0.4Mg alloys (a) with standard holder and (b) with 45◦ tilt. Sample shown in the ﬁgure was heat treated to T6I4.

The streaking greatly aﬀected the intensity ranges of interests, and no intensity integration could be completed without including the streaks. After tilting the sample by 45◦, a change in the streaking pattern was observed (Figure 4.38b). However, asymmetrical patterns for small angles were still obvious. The cause of the streaking is speculated to be that there was a strong anisotropy inﬂuence to the scattered X-rays. A possible solution to removing the streaking would be to homogenize, roll and recrystallize the as-cast alloys. However, it was not tested due to time constraints, and the fact that the actual cause of the streaking is unknown.

92 4.7 Summary

In this study three variations of interrupted aging treatments of 2024 aluminum alloys were investigated using microhardness and resistivity measurements and compared with single-stage T6 and T8 treatments, with additional “fast” and “slow” heating rates for the re-aging treatments in an effort to understand the differences in properties. TEM microstructural analyses were performed at various points of interest during the aging processes and combined with SAXS studies. An alloy similar to 2014 with different Mg and Cu concentrations was also investigated for comparison. The results are summarized as follows:

In 2024 alloys, almost 50% of the total hardness increase occurs during an initial rapid hardening period when performing single-stage aging treatments using aging temperatures below ∼190 ◦C. A corresponding increase in resistivity suggests the formation of small precipitates, possibly GPB zones/clusters.

Similar rapid initial hardening is not observed during T8 treatments after the addition of a small amount of cold work, although there is a small drop in resistivity in the early stages of artiﬁcial aging. TEM examination revealed no signs of dislocation recovery, thus the small resistivity drop may indicate that some amount of solute is removed from the matrix, most likely by segregation to dislocations.

Solution heat treated alloys display dynamic strain aging effects (serrated flow) during the tensile tests used to introduce the pre-strain, whereas alloys interrupted after underaging into the plateau mentioned above show no such effects.

The interrupted aging studies did not show any signiﬁcant impact on the ﬁnal T6 or T8 hardness of 2024 alloys due to softening that occurs upon re-aging regardless of the heating rates used.

Increases in hardness and resistivity are observed in unstretched alloys during secondary aging at 65 ◦C. However, only marginal hardness increases are observed in deformed alloys, with little change in resistivity.

SAXS analysis is used to show that the initial rapid hardening appears to be the result of the formation of ﬁne (1 nm) clusters of volume fractions of up to 3%.

Further, the hardness increase during the T6I4 treatment is also due to the formation and coarsening of these ﬁne clusters.

Reversion of the 1 nm clusters formed during the interrupted aging treatment at 65 ◦C appears to be the cause for the observed softening upon re-aging to 177 ◦C.

The observed resistivity increase at the hardness plateaus appears to be due to the increase in the volume fraction of 1 nm clusters.

93 CHAPTER 5

DISCUSSION

The results of this study showed that interrupted aging of two Al-Cu-Mg based alloys did not exhibit any significant beneficial strengthening over that obtained via single-stage aging, in contrast to the significant improvements reported previously by Lumley and co-workers [4–13]. Discussions on the effects of cold work and heating rates, as well as cluster analyses using SAXS and resistivity, are detailed in this chapter.

5.1 The Eﬀect of Cold Work on Interrupted Aging

The formation of clusters was shown both from literature [4–13, 51, 63, 66, 67, 69, 73, 84, 85, 91, 122] and this study (see Chapter 4.6) to be the cause for the initial rapid hardening and secondary strengthening after T6I4 treatment.

Quenched-in vacancies are known to assist the formation of clusters [64, 66, 101]. It is also shown that vacancies can assist solute diﬀusion by forming vacancy-solute pairs [66], which may be attracted to dislocations, thus removing them from the matrix. Introducing cold work increases the dislocation density in the material, which may remove the necessary quenched-in vacancies. Therefore, estimation of the dislocation density may be important in explaining the ineﬀectiveness of cold work on interrupted aging observed in this study. The dislocation density can be calculated using the Taylor equation (Equation 5.1 [34, 40]):

√ ∆τ = αµb ρ (5.1) or √ ∆σ = M · ∆τ = M·αµb ρ (5.2) where ∆τ is the increase in the critically resolved shear stress of the material (MPa), ∆σ is the increase in yield stress (MPa), M is the Taylor factor, α is a dimensionless constant of the material that can range from 0.05 to 1.3 [138], µ is the shear modulus of the material (GPa), b is the Burgers vector (m), and ρ is the dislocation density (m−2).

Equation 5.2 can be re-arranged to calculate dislocation density ρ:

∆σ 2 ρ = (5.3) αµbM

The strength increase through cold work can be estimated using the linear relationships between CRSS, yield stress, and hardness. The yield stresses of a 2024 alloy after SHT and after

94 underaging at 177 ◦C for 0.5 h can be obtained from Figure 5.1 to be ∼135 MPa and ∼245 MPa, respectively. The stress, after pulling to 5% tension, at the two conditions are ∼254 MPa and ∼351 MPa. The stress increase is 119 MPa for SHT and 106 MPa for underaged condition.

Figure 5.1 Stress-strain curve of a 2024 alloy after SHT and underaged at 177 ◦C for 0.5 h.

Several values of α were reported in the literature, namely 0.25 [139], 0.3 [140–142], and 1/3 [143]. Let Taylor factor M = 2.8, α = 0.25, b = 2.86 A,˚ and µ = 27 GPa. The dislocation density can be calculated using Equation 5.3 to be 4.9×1010 cm−2 at the SHT condition and 3.9×1010 cm−2 for the underaged condition. These values agree well with the experimental observation of a rapid increase to 1010 ∼ 1011 [144, 145]. In both conditions, the dislocation density increases by four orders of magnitude from the unstretched condition (∼106 cm−2). The average spacing between √ dislocations can also be calculated to be approximately 50 nm (1/ ρ).

Quenched-in vacancies also have a tendency to aggregate together and form voids/dislocation loops [20, 146–151]. These voids may be annihilated during cold work by dislocation glide [151].

The observed serrated ﬂow (Figure 5.1) during the 5% pre-aging strain suggests solute- dislocation interactions in the vicinity of the moving dislocations, which makes solute segregation

95 to dislocations plausible. It is conceivable that formation of clusters eﬀectively “locks” solutes from the SSSS and results in the lack of mobile solutes to be “dragged” by the mobile dislocations. Therefore, a smooth stress-strain curve is observed in underaged alloys, in which cluster formation is prevalent. In fact, it was shown that for 2024 alloys, serrated ﬂow was only observed under the SHT condition, but not for alloys naturally aged for 3 days under varying strain rates (1.5×10−3 to 1.5×10−5 s−1) [125]. In addition, because of the high dislocation densities in both conditions, heterogeneous nucleation of S’ phase on dislocations is thought to be favorable.

5.2 The Relationship Between Resistivity and Volume Fractions of Clusters

The linear regression of the SAXS and resistivity data in Figure 4.37 provides a correlation between the volume fraction of 1 nm clusters and resistivity change. A similar analysis was carried out to show the total volume fractions of clusters between sizes 1 to 30 nm and resistivity in Figure 5.2. For the latter, the overall regression of the data does not show as good a fit to that of the volume fraction of 1 nm clusters. Therefore, the effects from the larger clusters on resistivity are not as significant as the smaller 1 nm clusters. Although the resistivity in this study is made using a four-point probe apparatus (Section 3.5) and specially prepared samples, the changes in resistivity are usually large enough that they can be measured by common hand-held Eddy current testers, which makes resistivity a quick, useful, and non-destructive tool in monitoring the precipitation behavior.

5.3 Reasoning Behind the Assumptions of Cluster Analysis Using SAXS

A necessary part of calculating volume fraction of clusters using SAXS is the determination of the electron density differences (see Equation 2.20). In other words, the SAXS technique is less sensitive to clusters that are primarily Mg and Al compared with those involving Cu (electron density of copper = 2.46×1024 cm−3) because of the smaller electron density differences between Al (7.83×1023 cm−3) and Mg (5.17×1023 cm−3). Although Mg clusters are present during the heat treatment according to the APT studies by Marceau et al. [12, 13], it is unlikely that they are effective strengtheners based on the study by Kubota et al. [152, 153], in which a much higher concentration of Mg (7 wt% or higher) is required to precipitate the necessary strengthening phases.

Because the electron density of the clusters corresponds to their composition, it is necessary to assume a composition to fit the SAXS data. In this study, the composition of the clusters is assumed to be Al2CuMg, i.e. the same as the S’ and S phases. Further, since nothing (strain fields, diffuse intensity, etc.) was observed in the TEM after the early stage hardening, the clusters must be compositional variations in the fcc structure where the clusters must have compensating sizes.

Interestingly, the average radius of a dimer consisting of a Cu atom (atomic radius = 128 pm) and a Mg atom (atomic radius = 160 pm) is 144 pm, almost identical to that of an Al atom (atomic

96 Figure 5.2 Total volume fractions in vol% of clusters 1 to 30 nm) and change of resistivity ∆ρ in µΩ−cm. radius = 143 pm). Using this argument, the Mg/Cu ratio might be expected to be ∼1 which is the case for the proposed Al2CuMg stoichiometry. This assumption also agrees with a recent ﬁnding by Deschamps et al. [69] using NMR and SAXS, in which Cu concentration is found to be ∼ 30 at% in these ﬁne features found during early stages of aging a Al-2.5Cu-1.5Mg (wt%).

In 2024 alloys, the Cu and Mg contents can be assumed to be 1.7 at% each, which equals to 1.02×1021/cm3 atoms for both Cu and Mg. Mass balance calculations show that free solutes are present after underaging at all temperatures tested from RT to 190 ◦C, thus making the coarsening of clusters/precipitates possible.

97 If all Cu and Mg atoms are in the Al2CuMg clusters, the maximum population density of the 1 nm clusters is calculated to be 12.75×1019/cm3, well above the maximum population density obtained in this study (5.87×1019/cm3, see Section 5.6).

Suppose the clusters are composed mainly from Cu and Mg, the maximum population density of clusters is then calculated to be 6.375×1019/cm3, which is still above the measured population density in this study (5.87×1019/cm3, see Section 5.6). Therefore the measured volume fraction is within the anticipated maximum.

5.4 Mechanism(s) of Cluster Hardening

Various precipitation hardening mechanisms were discussed in Section 2.5 and the appropriate expressions for the diﬀerent hardening mechanisms were given. Based on the experimental results obtained to date however, only chemical strengthening and modulus strengthening seem plausible and these will be discussed further below.

The change in dislocation breaking angle can be diﬃcult to estimate, and to simplify this issue, it is assumed to be unchanged. Because of the identical materials and testing conditions, the strengthening expressions for chemical (Equation 2.10) and modulus (Equation 2.9) strengthening could be re-written and simpliﬁed into Equations 5.4 and 5.7, respectively, as a function of the cluster radius r and volume fraction f. Cc and Cm are the grouped material constants for chemical and modulus strengthening, respectively.

- Chemical Strengthening:

f 1/2 ∆H = C ∝ F (r, f) (5.4) V c r

f 1/2 F (r, f)(for chemical strengthening) ≡ (5.5) r

3 1/2 γ 3/2 M C ≡ 2 µ b (5.6) c π µb A where A is the conversion factor between yield stress (YS, σ) and Vickers hardness (HV), i.e.

∆σ = A · ∆HV , and is calculated to be 1.97. M is the Taylor factor, µ is the shear modulus, b is the Burgers vector (2.86 A),˚ and γ is the surface energy of the new surface area created by the sheared clusters.

- Modulus Strengthening:

2 < r >−3/2 ∆H = C (< r > f)1/2 2b ln ∝ F (r, f) (5.7) V m f 1/2b

98 2 < r >−3/2 F (r, f)(for modulus strengthening) ≡ (< r > f)1/2 2b ln (5.8) f 1/2b

T ∆µ3/2 M Cm ≡ 0.9 (5.9) b µm A again, A is the conversion factor (1.97) between yield stress (YS, σ) and Vickers hardness (HV).

M is the Taylor factor, µ is the shear modulus of the precipitate, µm is the shear modulus of the 2 matrix (27 GPa), b is the Burgers vector (2.86 A),˚ and T is the dislocation line tension (0.8µmb ). Because of the linear relationships between CRSS, yield stress, and Vickers hardness, the change in hardness is also proportional to F(r,f). Using the simpliﬁed equations, relationships between chemical and modulus strengthening and cluster size, volume fraction, and change in hardness can be determined as shown in Figures 5.3a and 5.3b, respectively. Sizes of 1 nm and total volume fractions of clusters 1-30 nm are both used with their respective volume fractions for these plots. Since linear relationships are expected for both mechanisms, the regression analyses were made under the assumption that only the individual strengthening mechanisms are operating in the material.

5.5 Predicted Strengthening of the T6I4 Material and the Subsequent Re-aging Curves

The F(r,f) of T6I4 and T6I4-reheat can be calculated using the results from SAXS analyses. The calculated results of F(r,f) and the respective estimated HV from the individual strengthening mechanisms and clusters involved are listed in Tables 5.1 and 5.2, respectively.

Table 5.1 – Calculated F(r,f) (×10−3) of T6I4 and T6I4-reheat treatments for chemical and modulus strengthening.

chemical modulus chemical modulus Temper F(r,f)total F(r,f)total F(r,f)1 nm F(r,f)1 nm T6I4 3.87 7.60 2.82 4.00 T6I4-reheat 2.87 7.11 2.42 3.21

Table 5.2 – Measured and estimated HV of T6I4 and T6I4-reheat treatments for chemical and modulus strengthening using relationships from Figure 5.3.

chemical modulus chemical modulus Temper HV-measured HVtotal HVtotal HV1 nm HV1 nm T6I4 134 129 129 127 127 T6I4-reheat 126 124 125 124 124

99 (a)

(b)

Figure 5.3 Relationships between change in hardness (relative to that for the SHT condition, namely 94 HV ) and F(r,f) at the hardness plateau at diﬀerent temperatures: (a) using the chemical strengthening model from Equation 5.4 and (b) using the modulus strengthening model from Equation 5.7. Analyses were conducted for clusters of all sizes (F(r,f)total) and 1 nm only (F(r,f)1 nm). F(r,f) is calculated with r in A˚ and f in volume fraction.

100 It is noticeable that the estimated hardening from 1 nm clusters has lower hardness values than the measured T6I4 value for both chemical and modulus strengthening mechanisms, indicating that additional strengthening may be considered. In addition, the secondary strengthening during T6I4 must be contributed from at least chemical and modulus strengthening involving volume fractions of total clusters.

5.6 Volume Fractions of Clusters and Interparticle Spacing in 2024 Alloys at Plateau Hardening, After Secondary Aging, and Re-aging.

The interparticle spacing between the clusters can be calculated if the volume fraction and sizes of the clusters are known. In 1 cm3 of Al, let N = population of clusters in this 1 cm3 Al (will −3 4 3 carry a unit of cm ) and the volume of a cluster = 3 πr . The volume density can therefore be expressed in Equation 5.10.

4 1·f = πr3·N (5.10) 3

Because the unit volume of Al is used for the calculation, N is now equal to the population density and can be determined using Equation 5.11. The interparticle spacing, L, between the clusters is estimated as L = N −1/3. To estimate the average interparticle spacing of clusters from 1 to 30 nm, simply replace r with average , where = /2. Using the results in Table 4.3 and 4.4, the population densities and interparticle spacings can be calculated for 2024 alloys hardened to plateau, T6I4, and T6I4-reheat. The results are shown in Table 5.3.

3f N = (5.11) 4πr3

Table 5.3 – Population densities and interparticle spacings of 1 nm and total clusters for plateau hardened, T6I4, and T6I4-reheat 2024 alloys.

19 −3 19 −3 Temper N1 nm ×10 cm Nt ×10 cm nm L1 nm nm Lt nm 25 ◦C 5.87 13.18 4.07 2.57 19.65 65 ◦C 5.25 10.19 4.41 2.67 21.41 100 ◦C 4.62 8.60 4.41 2.79 22.66 130 ◦C 4.38 5.25 5.45 2.84 26.71 150 ◦C 2.41 3.72 5.16 3.46 29.95 177 ◦C 1.94 2.13 5.61 3.72 36.06 190 ◦C 1.54 1.94 6.06 4.02 37.22 T6I4 3.79 7.67 4.66 2.98 23.53 T6I4-reheat 2.79 3.43 5.75 3.30 30.79

As aging temperature increases, both interparticle spacings L1 nm and Lt decrease. Despite a less than 1 nm change in L1 nm, it is also obvious that there is a signiﬁcant decrease in to-

101 tal interparticle spacing Lt after the T6I4 treatment compared to the underaged counterpart at ◦ 177 C. Likewise, partial reversion of these 1 nm clusters in T6I4-reheat also causes Lt to increase. Since the linear density of these clusters is inversely proportional to L, the number of clusters a dislocation will intercept as it glides through the matrix will also be inversely proportional to L. Therefore, the subtle changes in 1 nm clusters in T6I4 and T6I4-reheat treatments can, in fact, have profound effects on strengthening and softening. In addition, the change in interparticle spacing could also have a profound effect on the dislocation breaking angle φ, which could greatly affect the strengthening through cluster-dislocation interactions.

5.7 Estimating the Shear Modulus Diﬀerence and Excess Surface Energy by Shearing

The diﬀerence between the shear modulus of the 1 nm clusters and Al, calculated from the above analysis, was from 2.94 to 3.34 GPa, as shown in Table 5.4. This calculated range of values agrees well with the estimation of Starink et al. [67, 91], in which a diﬀerence of 3 GPa was given.

The excess surface energy caused by cluster shearing by dislocations can also be estimated using a varying Taylor factor between 2.87 [38] to 3.1 [1]. The results of the calculations made with 1 nm clusters and total clusters are also included in Table 5.4. The values calculated from 1 nm clusters range from 0.16 to 0.17 J/m2. However, the interfacial energy increased to about 0.6 J/m2 using the total clusters from 1 to 30 nm size, and is approximately three times larger than that calculated from 1 nm clusters.

Several literature values of coherent interfacial energies using ab initio calculations were reported in different alloy systems. The coherent interfacial energy between the Al-matrix and the GP zones in Al-Ag alloys at 450 K was calculated to be ∼0.03 J/m2 [154], whereas values of 2 ∼0.2 J/m for the Al3Sc/Al interfacial energy in Al-Sc alloys were reported [155]. In addition, Wolverton et al. [156, 157] calculated the coherent and semi-coherent interfacial energies of a θ0 in Al-Cu alloys to be 0.235 and 0.615 J/m2, respectively. The later value agrees well with the interfacial energy calculated from total clusters using the fitting from Figure 5.3a, and although the explanation is unclear, it can be speculated that the larger clusters may have a different dislocation- particle interaction.

Table 5.4 – Calculated surface energies using the chemical strengthening model and modulus diﬀerences between matrix and clusters from modulus strengthening model for 2024 aluminum alloys.

C ∆µ, GPa γ, J/m2 Chemical, 1 nm 715.9 - 0.16 - 0.17 Chemical, total 4912.8 - 0.57 - 0.63 Modulus, 1 nm 3583.8 3.17 - 3.34 - Modulus, total 3196.9 2.94 - 3.09 -

102 5.8 Predicted Volume Fractions of 1 nm Clusters for Al-4.2Cu-0.4Mg Alloys Using Resistivity and Their Predicted Strengthening

As stated in Chapter 4.6.5, SAXS analyses of the volume fractions of clusters in Al-4.2Cu- 0.4Mg were not successful due to possible anisotropic eﬀects on X-ray scattering. Even so, the volume fractions of 1 nm clusters can be estimated by using the change of electrical resistivity to the 1 nm cluster fraction relationships obtained from Chapter 4.6.4, where f (in %) = 6.686·∆ρ (in µΩ−cm) and the results of the 30-minute-underaging at 177 ◦C, T6I4 for 2 days, and T6I4 for 12 days are listed in Table 5.5

Table 5.5 – Measured and estimated 1 nm cluster volume fractions and HV of 30-minute underaging, T6I4-2days and T6I4-12days treatments for Al-4.2Cu-0.4Mg alloys using relationships from Figure 4.37 for resistivity and Figure 5.3 for chemical and modulus strengthening. ρ and ∆ρ are in µΩ−cm.

chemical modulus Temper HV-measured ρ ∆ρ est. f1 nm % HV HV SHT 68 4.14 - - - - 30 min @ 177 ◦C 81 4.16 0.02 0.13 - - T6I4-2days 104 4.27 0.13 0.87 94 94 T6I4-12days 116 4.31 0.17 1.14 96 96

It is notable that the hardness increase in this early stage only comprises 24% of the total hardness increase, significantly lower than that in 2024 alloys. This agrees with the finding on the magnitude of hardening in the early stage in Al-Cu-Mg alloys, which is found to be increasing as Mg content increases [13]. In addition, there is a minimal increase in resistivity in the first

30 minutes of aging, accompanied by a 13 HV increase. Also, as shown in Chapters 4.6.4 and 5.2, the resistivity is only sensitive to 1 nm clusters and therefore some larger clusters that might not be otherwise detected by resistivity might be the cause for the hardening in the underaged alloy. Further, it was shown that the early precipitates in Al-4Cu-0.3Mg (wt%) alloys to be GP zones and/or θ00 [90] rather than GPB zone formation. It is conceivable that the precipitate formation in the Al-4.2Cu-0.4Mg alloy during underaging to be GP zones and/or θ00. Therefore, additional strengthening by coherent strain between GP zones/θ00 and Al matrix is likely to contribute to the hardness increase.

The large resistivity increases found after the T6I4-2days and T6I4-12days treatments suggest the secondary precipitation of signiﬁcant amounts of 1 nm clusters during aging at 65 ◦C. The estimated volume fraction using change in resistivity suggests a maximum of 0.98 vol% of 1 nm clusters formed in the alloy, about 1 vol% less than that found in the 2024 alloy after T6I4 treatment. This lower volume fraction of 1 nm clusters is expected because of the lower Mg content present.

However, the estimated hardness values using the strengthening mechanisms are consistently lower than the measured hardness values. This also indicates that, similar to 2024 alloys, cluster

103 coarsening may play an important role in the strengthening of alloys undergoing secondary aging at low temperatures. Of course, the strengthening advantages gained during secondary aging are eliminated after re-aging to 177 ◦C due to the obvious softening caused by, in this case, complete reversion of the 1 nm clusters.

5.9 Aging Temperatures and Volume Fractions at Early Aging Plateau

The volume fraction of clusters at early plateau hardness from Table 4.3 can be plotted with respect to aging temperatures, as shown in Figure 5.4.

Figure 5.4 Aging temperatures and volume fractions of clusters at early hardness plateau.

The volume fractions of total clusters seem to have a sharper change when the aging temperature is above 130 ◦C, whereas the volume fraction of 1 nm clusters makes a steadier decline as aging temperature increases. Interestingly, both fractions are projected to drop to zero at aging temperatures around 220 ◦C. The projected result agrees well with the lack of hardness plateau observed in Al-Cu-Mg alloys aged at 240 ◦C [61]. Furthermore, softening can also be observed when alloys that are previously aged at lower temperatures, i.e. RT and 65 ◦C, to the plateau are

104 aged to elevated temperatures, as shown in Figure 5.5. The softening mechanism, as discussed in Chapter 5.5, is the partial if not complete reversion of 1 nm clusters.

Figure 5.5 Softening of alloys re-aged at 177 ◦C. The alloys were previously aged at RT and 65 ◦C into their respective plateau.

5.10 Possible SAXS Interference Eﬀect in Naturally Aged 2024 Alloy

An interesting difference in the shape of the normalized intensity curves between naturally aged (NA) and artificially aged alloys was observed during SAXS analyses. In the NA condition, a portion of the Intensity-q data cannot be fitted as well as those aged at elevated temperature, as shown in Figure 5.6 by the red oval. Consequently, a “peak” like characteristic is observed.

This “peak” like behavior is not observed in alloys aged at higher temperatures such as 190 ◦C. It occurs at q values of 2 to 3. This behavior is speculated to be due to interparticle interference. Speciﬁcally, upon forming the clusters, the solute adjacent to the clusters must be consumed, leaving a solute “depletion” zone, as shown in Figure 5.7. Now if similar analyses on interparticle spacing is applied to solute spacing, it is estimated that in 2024 alloys, the “uniform” spacing between Cu/Mg

105 Figure 5.6 Absolute SAXS intensity vs momentum transfer q of 2024 alloys naturally aged for 506 h. Blue curve shows the fitting and the red oval indicates the region where data cannot be properly fitted. and Mg/Cu is around 8 A˚ (roughly 2 unit cells between each other), which is plausible for short- range diffusion to occur at RT to form clusters. The 1 nm clusters have a volume equivalent to 8 unit cells, and if the chemical composition of Al2CuMg is applied to the clusters, it would contain 16 Cu+Mg solute atoms, which would have to come from the nearby matrix, roughly about a volume of 100 unit cells, or a radius of 1.2 nm. This is consistent with a roughly uniform interparticle spacing of 2.57 nm as shown in Figure 5.7. Simply using Bragg’s law, d = λ/2sinθ = 2π/q, and inserting the rough position of the “peak”, 2 to 3 nm −1, yields a “d-spacing” of 3 to 2 nm, spanning the range of interparticle spacings listed in Table 5.4 and the value shown in Figure 5.7.

It is also considered that the “peak” like characteristic may be attributed to the diﬀraction of the crystalline T-phase, Al20Cu2Mn3 (Bbmm, a = 2.41 nm, b = 1.25 nm, and c = 0.78 nm).

106 However, this possibility is ruled out because a distinctively “sharp” peak would be observed instead of the diﬀuse eﬀect seen in Figure 5.6.

Figure 5.7 Schematic illustration of the interference eﬀect with interparticle spacing for 1 nm clusters formed at RT.

5.11 Conclusions on Interrupted Aging of Al-Cu-Mg Alloys

Introduction of cold work was found to be uncomplimentary to interrupted aging, contrary to the original studies on interrupted aging by Lumley et al. [4, 6, 7]. No increase in hardness is found in either variants, T8I6 and T6I8, during secondary aging at 65 ◦C. This ineﬀectiveness is believed to be caused by the elimination of excess vacancies and, therefore, solute segregation to dislocations becomes the prevalent mechanism because of the excess amount of dislocations present in the alloy.

107 The solute segregation itself is most likely not a strong mechanism because the hardening gained during aging shortly after stretching is minimal.

However it might be inﬂuential to the subsequent microstructural development. Much smaller S’ phases that are also more densely distributed are present around dislocations at peak aging. This indicates that heterogeneous nucleation and growth of the S’ phase on dislocations is prevalent. The hardening observed in stretched alloys during aging is then the result of the S’ phase formation. Segregation to dislocations also eﬀectively removes solutes from the matrix for secondary precipitation of clusters when aging at 65 ◦C and, therefore, leads to the lack of hardness increase in stretched alloys upon secondary aging.

On the other hand, secondary strengthening is observed in part of the interrupted aging, namely T6I4 treatments, on unstretched alloys. In the two Al-Cu-Mg alloys tested in this study, the secondary strengthening happens during the ﬁrst few days of aging at 65 ◦C, and reaches an extended plateau for the remainder of the aging period. Resistivity also follows a similar trend. SAXS analyses show that secondary precipitation and coarsening of clusters during secondary aging is directly related to the hardness increase, with resistivity being sensitive to the formation/reversion of 1 nm clusters. However, retaining the secondary strengthening cannot be achieved and is removed by softening upon re-aging.

The ineffectiveness of the complete T6I6 interrupted aging conducted in this study is shown to be due to the partial or complete reversion of 1 nm clusters. Apart from formation of new clusters, the secondary aging at 65 ◦C indeed promotes coarsening of clusters, but it appears that they do not have any significant effect on subsequent strengthening upon re-aging at 177 ◦C. After secondary aging, the hardness is equal to that of materials aged at room temperature for over 24 h. However, the volume fraction of 1 nm clusters is significantly lower by 1 vol%. Subsequent analyses show that an increase in 1 nm cluster volume fraction and coarsening of clusters also influence the effectiveness of chemical strengthening, which increases the “friction” of the matrix to dislocations to propagate. The coarsened clusters, on the other hand, do not seem to lead to improved strength after re-aging at 177 ◦C to peak hardness.

108 CHAPTER 6

SUMMARY

Several variants of the interrupted aging treatments developed by Lumley et al. [4–7] were investigated in this study using a commercial wrought 2024 alloy as well as an as-cast Al-4.2Cu- 0.4Mg (wt%) alloy that has the Mg and Cu concentrations corresponding to that of the 2014 alloy examined by Lumley [4]. In addition, the re-heating rate to the artificial aging temperature was examined in an effort to determine whether the secondary strengthening observed after a T6I4 temper could be retained during ensuing artificial aging treatment if the heating rate was controlled. SAXS analysis was also used to determine the effects of cluster size and volume fraction associated with the initial rapid hardening that is observed, as well as the secondary strengthening that develops during interrupted aging treatments. The results are summarized in the following sections.

6.1 Single Stage and Interrupted Aging of 2024 Alloys

During single stage T6 aging treatments at typical aging temperatures (e.g. 175 ◦C), about 50% of the total hardness increase is achieved during the ﬁrst minute or so of aging. The increase in resistivity indicates the formation of small solute clusters, which is considered to be the source of hardening. The microstructure at peak aging, achieved after aging at 177 ◦C for ∼24-30 h, contains relatively coarse S’ phase, with ﬁne S”/GPB zones dispersed in the matrix.

On the contrary, no rapid initial hardening is observed during a T8 treatment after the addition of cold work. However, a small resistivity drop is observed with a minimum hardness increase. TEM examinations show no signs of dislocation recovery, thus, the small resistivity drop suggests that some amount of solute is removed from the matrix through segregation to dislocations. The increase in dislocation density and the serrated ﬂow observed during the stretching of solution heat treated alloys also indicate the likelihood of solute segregation to dislocations by short range diﬀusion.

Although underaged alloys show a smooth curve during the tensile tests, the interactions between solute clusters and dislocations by shearing may also result in similar segregation. The microstructure at peak aging is predominantly S’ phase that nucleated heterogeneously and grew on dislocations with denser populations and smaller precipitate sizes than T6. It is also conceivable that the increasing presence of dislocations eliminates the formation of solute clusters, which explains why no hardness increases are observed in stretched alloys, with little change in resistivity during secondary aging at 65 ◦C. In both stretched alloys, the peak hardening is achieved after aging at 177 ◦C for ∼12 h.

109 Secondary hardening is observed in unstretched alloys during secondary aging at 65 ◦C (T6I4 temper). It is shown through SAXS analyses that the secondary hardening results from the formation of 1 nm solute clusters and cluster coarsening during aging at 65 ◦C. However, after completing the full thermal cycle for interrupted aging, no additional strengthening is found after re-aging to 177 ◦C from T6I4 temper due to obvious softening after re-heating, regardless of heating rate. This softening correlates via the SAXS analyses with the partial reversion of solute clusters formed during secondary aging at 65 ◦C. In addition, hardness and resistivity both reach a plateau after about 100 h (4∼5 days)of secondary aging at 65 ◦C, contrary to the continuous hardness increase during the whole 2-week period shown previously by Lumley. Overall, the T6I6 interrupted aging treatment does not provide any signiﬁcant improvements in the hardening of the 2024 alloys.

6.2 Single Stage T6 and Interrupted Aging T6I6 of Al-4.2Cu-0.4Mg Alloys

Al-4.2Cu-0.4Mg alloys were used to examine the eﬀectiveness of interrupted aging on Type 2014 alloys, as originally reported by Lumley et al. [4] as their example of the Al-Cu-Mg based aluminum alloys where higher hardnesses can be achieved using T6I6 treatments. The single stage T6 aging behavior of the Al-4.2Cu-0.4Mg alloys is similar to that of the 2014 alloys used by Lumley. About 24% of total hardening is achieved in the ﬁrst 30 minutes of aging at 177 ◦C with a marginal resistivity increase. The peak hardening at 177 ◦C can be achieved after aging for ∼24-30 h, similar to that of 2024 alloys.

Similar to its 2024 counterpart, the Al-4.2Cu-0.4Mg alloy also shows secondary hardening during aging at 65 ◦C and reaches a plateau after secondary aging at 65 ◦C for ∼100 h. Similar softening is observed after re-heating to 177 ◦C from selected T6I4 conditions, and the overall aging behavior of T6I6-tempered alloys does not exhibit any additional strengthening at peak hardness. Based on the resistivity behavior, the softening is believed to be caused by the partial or complete reversion of secondary clusters. This examination proves the irreproducibility of T6I6 temper on 2014 alloys.

6.3 Resistivity, Strengthening Mechanisms for Cluster Hardening, and Conclusion

The increase in resistivity observed during the initial rapid hardening and secondary aging of unstretched alloys is shown to be closely related to the formation of solute clusters. In particular, the resistivity is only sensitive to the changes of 1 nm clusters, and insensitive to clusters with larger sizes.

The initial rapid hardening and secondary strengthening is shown to be the result of an increase in volume fraction of solute clusters, particularly those that are ∼ 1 nm in size; this has been referred to as cluster hardening. The overall hardening of alloys, however, comes from clusters of all sizes from 1 to 30 nm. SAXS analyses, combined with TEM observations and hardness

110 measurements, show that the strengthening mechanisms for cluster hardening are modulus and chemical/surface strengthening. The normally negligible [1, 39, 91] chemical/surface strengthening cannot be ruled out from the strengthening and softening observed before and after re-heating from the T6I4 temper.

Because these ﬁne clusters are not stable upon reheating and suﬀer dissolution (reversion), it is concluded that the interrupted aging T6I6 and its cold-worked variants, T8I6 and T6I8, do not exhibit the anticipated additional strengthening compared to their single-stage counterparts. The lack of thermal stability of 1 nm clusters that partially or completely revert upon reheating results in the obvious softening.

111 CHAPTER 7

FUTURE WORK

Continued TEM characterization and SAXS analysis on the Al-4.2Cu-0.4Mg alloy, as well as possible alloys with diﬀerent Mg contents, is of interest to further understand the potential eﬀects of Mg content on the size and volume fraction of clusters. Based on the results from Al- 4.2Cu-0.4Mg alloys in this study and Ref [90], it can be postulated that GP zones/θ00 are present at initial underaging at 177 ◦C. If so, it will suggest that there might potentially be a correlation between Mg content and aging temperature for Cu-Mg cluster formation. This may allow for better understanding of the cluster formation in Al-Cu-Mg based alloys and therefore may also allow for new alloy designs.

It may be possible to take advantage of the cluster hardening in alloy designs. An example can be made for 7xxx series alloys (Al-Zn-Mg). It is well known that faster higher peak hardening can be achieved in 7xxx alloys by undergoing a double aging treatment, i.e. low-high stages of temperature [41]. The strengthening phases in 7050 alloys, for example, contain primarily Zn and Mg, with a certain amount of Cu. It may be possible that with additions of excess Mg, there might be some remaining Mg to cluster with Cu after the peak aging is achieved after double aging without any significant Cu precipitation. By doing so, it may be possible for the Cu-modified 7xxx alloys to continue hardening at RT or slightly elevated temperatures such as 65 ◦C after peak aging. However, the thermal instability of the solute clusters would most likely limit the potential utilization of this modified alloy, if RT-hardenable as proposed, to low-temperature applications due to the reversion issue.

In this study, the chemical composition of clusters is assumed to be Al2CuMg. However, various authors [8–13, 73] using 3DAP suggest that the clusters do not have a fixed composition, nor do they exhibit a certain size. This may be because of the criteria used in cluster identification from the atom probe data. Therefore it would be of scientific interest to use the sizes and fractions obtained from SAXS as the cluster identification criteria in 3DAP, and compare the chemical compositions of the clusters to observe any consistency in compositions.

The assumption of fcc clusters can also be possibly re-examined through simulations. It is suggested that the strengthening of solute clusters can be from as small as Cu-Mg dimers [91]. However it is also suggested that significant strengthening by clustering cannot be achieved without an ordered structure [101]. Therefore simulations using strengthening mechanisms identified in this study, modulus and chemical, could be performed on cluster hardening. For example, simulations can be performed on different structures such as the layered structure proposed by Wolverton [62] and the 1D-GPB structures proposed by Kovarik et al. [101] to compare theories and experimental data in hope of facilitating a better understanding.

112 REFERENCES CITED

[1] J. Martin, Precipitation Hardening, 2nd ed. Oxford: Butterworth-Heinemann, 1998.

[2] I. Polmear, “Aluminium Alloys–A Century of Age Hardening,” Materials forum, vol. 28, pp. 1–14, 2004.

[3] I. Polmear, Light Alloys - From Traditional Alloys to Nanocrystals, 4th ed. Oxford: Butterworth-Heinemann, 2006.

[4] R. N. Lumley, I. J. Polmear, and A. J. Morton, “Interrupted aging and secondary precipitation in aluminium alloys,” Materials Science and Technology, vol. 19, no. 11, pp. 1483–1490, Nov. 2003.

[5] R. Lumley, I. Polmear, and A. Morton, “Development of properties during secondary ageing of aluminium alloys,” in Materials Science Forum, vol. 426, no. 1. Trans Tech Publ, 2003, pp. 303–308.

[6] R. N. Lumley, I. J. Polmear, and A. J. Morton, “Development of mechanical properties during secondary aging in aluminium alloys,” Materials Science and Technology, vol. 21, no. 9, pp. 1025–1032, Sep. 2005.

[7] R. Lumley, J. Buha, I. J. Polmear, A. Morton, and A. Crosky, “Secondary Precipitation in Aluminium Alloys & Its Role in Modern Heat Treatment,” Materials Science Forum, vol. 519-521, pp. 283–290, 2006.

[8] J. Buha, R. Lumley, and A. Crosky, “Microstructural development and mechanical properties of interrupted aged Al-Mg-Si-Cu alloy,” Metallurgical and Materials Transactions A, vol. 37, no. 10, pp. 3119–3130, 2006.

[9] J. Buha, R. Lumley, A. Crosky, and K. Hono, “Secondary precipitation in an AlMgSiCu alloy,” Acta Materialia, vol. 55, no. 9, pp. 3015–3024, May 2007.

[10] J. Buha, R. N. Lumley, and A. G. Crosky, “Precipitation and solute distribution in an interrupted-aged AlMgSiCu alloy,” Philosophical Magazine, vol. 88, no. 3, pp. 373–390, Jan. 2008.

[11] J. Buha, R. Lumley, and A. Crosky, “Secondary ageing in an aluminium alloy 7050,” Materials Science and Engineering: A, vol. 492, no. 1-2, pp. 1–10, Sep. 2008.

[12] R. Marceau, G. Sha, R. Lumley, and S. Ringer, “Evolution of solute clustering in AlCuMg alloys during secondary ageing,” Acta Materialia, vol. 58, no. 5, pp. 1795–1805, Mar. 2010.

[13] R. K. W. Marceau, G. Sha, R. Ferragut, A. Dupasquier, and S. P. Ringer, “Solute clustering in AlCuMg alloys during the early stages of elevated temperature ageing,” Acta Materialia, vol. 58, no. 15, pp. 4923–4939, 2010.

[14] S. Goh, J. Nie, L. Bourgeois, B. Muddle, and J. Embury, “Improved Multiple Ageing Treat- ments for AA6111,” Materials Science Forum, vol. 519-521, pp. 1499–1504, Jul. 2006.

113 [15] N. Gao, M. J. Starink, N. Kamp, and I. Sinclair, “Application of uniform design in optimi- sation of three stage ageing of AlCuMg alloys,” Journal of Materials Science, vol. 42, no. 12, pp. 4398–4405, Feb. 2007.

[16] D. Risanti, M. Yin, P. D. del Castillo, and S. van der Zwaag, “A systematic study of the eﬀect of interrupted ageing conditions on the strength and toughness development of AA6061,” Materials Science and Engineering: A, vol. 523, no. 1-2, pp. 99–111, Oct. 2009.

[17] D. Risanti, M. Yin, J. Chen, and S. van der Zwaag, “The Mechanical Properties of AA2024 as Function of the Interrupted Ageing Conditions,” Materials Science Forum, vol. 638-642, pp. 449–454, Apr. 2010.

[18] P. Merica, R. Waltenberg, and H. Scott, “Heat Treatment of Duralumin,” Bull. Am. Inst. Min. Metall. Eng., vol. 150, pp. 913–949, 1919.

[19] P. Merica, R. Waltenberg, and H. Scott, “Heat Treatment of Duralumin,” Sci. Pap. Bur. Stand, vol. 15, pp. 271–316, 1919.

[20] D. Porter and K. Easterling, Phase Tranformations in Metals and Alloys, 2nd ed. Boca Raton: CRC Press, 1992.

[21] R. DeHoﬀ, Thermodynamics in Materials Science, 2nd ed. CRC Press, 2006.

[22] J. W. Cahn and J. E. Hilliard, “Free Energy of a Nonuniform System. I. Interfacial Free Energy,” The Journal of Chemical Physics, vol. 28, no. 2, p. 258, Feb. 1958.

[23] J. W. Cahn, “Free Energy of a Nonuniform System. II. Thermodynamic Basis,” The Journal of Chemical Physics, vol. 30, no. 5, p. 1121, May 1959.

[24] J. W. Cahn and J. E. Hilliard, “Free Energy of a Nonuniform System. III. Nucleation in a Two-Component Incompressible Fluid,” The Journal of Chemical Physics, vol. 31, no. 3, p. 688, Sep. 1959.

[25] K. Rundman and J. Hilliard, “Early stages of spinodal decomposition in an aluminum-zinc alloy,” Acta Metallurgica, vol. 15, no. 6, pp. 1025–1033, Jun. 1967.

[26] E. Ma, “Alloys created between immiscible elements,” Progress in Materials Science, vol. 50, no. 4, pp. 413–509, May 2005.

[27] D. Newbury and D. Williams, “The electron microscope: the materials characterization tool of the millennium,” Acta Materialia, vol. 48, no. 1, pp. 323–346, Jan. 2000.

[28] C. Li, C. Niu, Z. Du, C. Guo, and Y. Jing, “The thermodynamic analysis of GP zones in aged-supersaturated Al-Ag alloys,” CALPHAD: Computer Coupling of Phase Diagrams and Thermochemistry, vol. 34, no. 1, pp. 120–128, 2010.

[29] J. Christian, The theory of transformations in metals and alloys, 1st ed. Oxford: Pergamon Press, 1965.

[30] J. Nie, B. Muddle, H. Aaronson, S. Ringer, and J. Hirth, “On the roles of clusters during intragranular nucleation in the absence of static defects,” Materials Transactions A, vol. 33, no. June, pp. 1649–1658, 2002.

114 [31] J. B. Cohen, “The early stages of solute distribution below a transition temperature,” Met- allurgical Transactions A, vol. 23, no. 10, pp. 2685–2697, Oct. 1992.

[32] J. Silcock, T. Teal, and H. Hardy, “No Title,” Journal of the Institute of Metals, vol. 84, p. 23, 1955.

[33] S. Ringer, K. Hono, I. Polmear, and T. Sakurai, “Nucleation of precipitates in aged AlCuMg (Ag) alloys with high Cu: Mg ratios,” Acta Materialia, vol. 44, no. 5, pp. 1883–1898, 1996.

[34] G. E. Dieter, Mechanical Metallurgy, 3rd ed. New York: McGraw-Hill, 1986.

[35] R. Nicholson, G. Thomas, and J. Nutting, “Electron-microscopic studies of precipitation in aluminum alloys,” Journal of the Institute of Metals, vol. 87, pp. 429–438, 1959.

[36] R. Nicholson, G. Thomas, and J. Nutting, “The interaction of dislocations and precipitates,” Acta Metallurgica, vol. 8, no. 3, pp. 172–176, Mar. 1960.

[37] E. Orowan, “Theory of Yield Without Shear,” Institute of Metals Symposium on Internal Stresses in Metals and Alloys, p. 451, 1948.

[38] A. Melander and P. A. k. Persson, “The strength of a precipitation hardened AlZnMg alloy,” Acta Metallurgica, vol. 26, no. 2, pp. 267–278, Feb. 1978.

[39] J. Williams and A. W. Thompson, “Strengthening of Metals and Alloys,” in Metallurgical Treatises, J. K. Tien and J. F. Elliott, Eds. New York: American Institute of Mining, Metallurgical, and Petroleum Engineers, Inc, 1981, pp. 487–504.

[40] A. H. Cottrell, Dislocations and Plastic Flow in Crystals. Oxford University Press, 1953.

[41] S. V. Emani, J. Benedyk, P. Nash, and D. Chen, “Double aging and thermomechanical heat treatment of AA7075 aluminum alloy extrusions,” Journal of Materials Science, vol. 44, no. 23, pp. 6384–6391, Sep. 2009.

[42] J. Silcock, T. Heal, and H. Hardy, “Structural aging characteristics of binary aluminium- copper alloys,” Journal of the Institute of Metals, vol. 82, p. 239, 1953.

[43] L. F. Mondolfo, Aluminium alloys : structure and properties. London: Butterworth, 1976.

[44] F. LeGoues, R. Wright, Y. Lee, and H. Aaronson, “Inﬂuence of crystallography upon critical nucleus shapes and kinetics of homogeneous fcc-fcc nucleationV. The origin of GP zones in AlAg and AlCu alloys,” Acta Metallurgica, vol. 32, no. 10, pp. 1865–1870, 1984.

[45] E. Matsubara and J. Cohen, “The GP zones in AlCu alloysII,” Acta Metallurgica, vol. 33, no. 11, pp. 1957–1969, 1985.

[46] M. Starink and P. Van Mourik, “Cooling and heating rate dependence of precipitation in an AlCu alloy,” Materials Science and Engineering: A, vol. 156, no. 2, pp. 183–194, 1992.

[47] M. Karlik and B. Jouﬀrey, “High resolution electron microscopy study of Guinier-Preston (GP1) zones in Al—Cu based alloys,” Acta materialia, vol. 45, no. 8, pp. 3251–3263, 1997.

[48] T. Konno and K. Hiraga, “Guinier-preston (GP) zone revisited: atomic level observation by HAADF-TEM technique,” Scripta materialia, vol. 44, pp. 2303–2307, 2001.

115 [49] M. Karlk, A. Bigot, B. Jouﬀrey, P. Auger, and S. Belliot, “HREM, FIM and tomographic atom probe investigation of GuinierPreston zones in an Al1.54at% Cu alloy,” Ultramicroscopy, vol. 98, no. 2-4, pp. 219–230, Jan. 2004. [50] S. Son, M. Takeda, M. Mitome, Y. Bando, and T. Endo, “Precipitation behavior of an AlCu alloy during isothermal aging at low temperatures,” Materials Letters, vol. 59, no. 6, pp. 629–632, Mar. 2005. [51] S. Ringer and K. Hono, “Microstructural Evolution and Age Hardening in Aluminium Alloys:: Atom Probe Field-Ion Microscopy and Transmission Electron Microscopy Studies,” Materials characterization, vol. 44, no. 1-2, pp. 101–131, 2000. [52] A. Guinier, “Structure of Age-Hardened Aluminium-Copper Alloys,” Nature, vol. 142, no. 3595, pp. 569–570, Sep. 1938. [53] G. D. Preston, “Structure of Age-Hardened Aluminium-Copper Alloys,” Nature, vol. 142, no. 3595, pp. 570–570, Sep. 1938. [54] T. Sato and A. Kamio, “High resolution electron microscopy of phase decomposition microstructures in aluminium-based alloys,” Materials Science and Engineering: A, vol. 146, no. 1-2, pp. 161–180, Oct. 1991. [55] J. Hwang, “Characterization and Mechanical Properties of Nanoscale Precipitates in Modiﬁed Al-Si-Cu Alloys Using Transmission Electron Microscopy and Three-Dimensional Atom Probe Tomography,” Ph.D. dissertation, University of North Texas, 2007. [56] J. Tsai and M. Kaufman, “Development of an Electrical Resistivity Apparatus Using Pure Aluminum - Quarterly Report to GM, Feb. 2009,” Colorado School of Mines, Golden, Tech. Rep., 2009. [57] J. Tsai, H. Doty, and M. Kaufman, “In Preparation.” [58] Y. Bagaryatsky, “Structural Changes on Aging Al-Cu-Mg alloys,” Doklady Akademii S.S.S.R, vol. 87, pp. 397–401, 1952. [59] Y. Bagaryatsky, “Y.” Doklady Akademii S.S.S.R, vol. 87, pp. 559–562, 1952. [60] J. Silcock, “The Structural Ageing Characteristics of Al-Cu-Mg Alloys with Copper : Mag- nesium Weight Ratios of 7 : 1 and 2.2 : 1,” Journal of the Institute of Metals, vol. 89, pp. 203–210, 1960. [61] H. Shih, N. Ho, and J. Huang, “Precipitation Behaviors in Al-Cu-Mg and 2024 Aluminum Alloys,” Metallurgical and Materials Transactions A, vol. 27, no. 9, pp. 2479–2494, 1996. [62] C. Wolverton, “Crystal structure and stability of complex precipitate phases in AlCuMg(Si) and AlZnMg alloys,” Acta Materialia, vol. 49, no. 16, pp. 3129–3142, Sep. 2001. [63] S. C. Wang and M. J. Starink, “Precipitates and intermetallic phases in precipitation hardening AlCuMg(Li) based alloys,” International Materials Reviews, vol. 50, no. 4, pp. 193–215, Aug. 2005. [64] R. K. Marceau, “Design in Light Alloys by Understanding Solute Clustering Processes During the Early Stages of Age Hardening in Al-Cu-Mg Alloys,” Ph.D. dissertation, The University of Sydney, 2008.

116 [65] L. Reich, S. Ringer, and K. Hono, “Origin of the initial rapid age hardening in an Al-1.7 at.% Mg-1.1 at.% Cu alloy,” Philosophical magazine letters, vol. 79, no. 9, pp. 639–648, 1999.

[66] Y. Nagai, M. Murayama, Z. Tang, T. Nonaka, K. Hono, and M. Hasegawa, “Role of vacancy- solute complex in the initial rapid age hardening in an Al-Cu-Mg alloy,” Acta Materialia, vol. 49, no. 5, pp. 913–920, 2001.

[67] M. Starink, N. Gao, and J. Yan, “The origins of room temperature hardening of AlCuMg alloys,” Materials Science and Engineering: A, vol. 387-389, pp. 222–226, Dec. 2004.

[68] M. J. Starink, N. Gao, L. Davin, J. Yan, and A. Cerezo, “Room temperature precipitation in quenched AlCuMg alloys: a model for the reaction kinetics and yield strength development,” Philosophical Magazine, vol. 85, no. 13, pp. 1395–1417, May 2005.

[69] A. Deschamps, T. Bastow, F. de Geuser, A. Hill, and C. Hutchinson, “In situ evaluation of the microstructure evolution during rapid hardening of an Al2.5Cu1.5Mg (wt.%) alloy,” Acta Materialia, vol. 59, no. 8, pp. 2918–2927, May 2011.

[70] M. J. Starink, A. Cerezo, J. L. Yan, and N. Gao, “Reply to the comments on Room- temperature precipitation in quenched AlCuMg alloys: a model for the reaction kinetics and yield-strength development,” Philosophical Magazine Letters, vol. 86, no. 4, pp. 243–252, Apr. 2006.

[71] S. Wang, M. Starink, and N. Gao, “Precipitation hardening in AlCuMg alloys revisited,” Scripta Materialia, vol. 54, no. 2, pp. 287–291, Jan. 2006.

[72] L. Kovarik, S. Court, H. Fraser, and M. Mills, “GPB zones and composite GPB/GPBII zones in AlCuMg alloys,” Acta Materialia, vol. 56, no. 17, pp. 4804–4815, Oct. 2008.

[73] R. K. W. Marceau, N. Tsafnat, D. Haley, and S. P. Ringer, “Solute Diﬀusion Characteristics of a Rapid Hardening Al-Cu-Mg Alloy during the Early Stages of Age Hardening,” Metallurgical and Materials Transactions A, 2010.

[74] G. Sha, R. Marceau, X. Gao, B. Muddle, and S. Ringer, “Nanostructure of aluminium alloy 2024: Segregation, clustering and precipitation processes,” Acta Materialia, vol. 59, no. 4, pp. 1659–1670, Feb. 2011.

[75] P. Ratchev, B. Verlinden, P. De Smet, and P. van Houtte, “Precipitation hardening of an Al-4.2 wt%Mg-0.6 wt%Cu alloy,” Acta Materialia, vol. 46, no. 10, pp. 3523–3533, 1998.

[76] P. Ratchev, B. Verlinden, P. De Smet, and P. Van Houtte, “Eﬀect of Cooling Rate and Pre- deformation on the Precipitation Hardening of an Al-4.2wt.%Mg-0.6wt.%Cu Alloy,” Scripta Materialia, vol. 38, no. 8, pp. 1195–1201, Mar. 1998.

[77] P. Ratchev, B. Verlinden, P. De Smet, and P. Van Houtte, “Artiﬁcial ageing of Al-Mg-Cu alloys,” Materials Transactions JIM, vol. 40, no. 1, pp. 34–41, Jan. 1999.

[78] I. Samajdar, P. Ratchev, B. Verlinden, P. Van Houtte, and P. De Smet, “Dislocation cell formation and hot ductility in an AlMgCu alloy,” Materials Science and Engineering: A, vol. 247, no. 1-2, pp. 58–66, Jun. 1998.

[79] Q. Chen, X.-Y. Liu, and S. Biner, “Solute and dislocation junction interactions,” Acta Ma- terialia, vol. 56, no. 13, pp. 2937–2947, Aug. 2008.

117 [80] N. G. Kioussis and N. M. Ghoniem, “Modeling of Dislocation Interaction with Solutes, Nano- Precipitates and Interfaces: A Multiscale Challenge,” Journal of Computational and Theo- retical Nanoscience, vol. 7, no. 8, pp. 1317–1346, Aug. 2010.

[81] Z. Feng, Y. Yang, B. Huang, M. Han, X. Luo, and J. Ru, “Precipitation process along dislocations in AlCuMg alloy during artiﬁcial aging,” Materials Science and Engineering: A, vol. 528, no. 2, pp. 706–714, Dec. 2010.

[82] Z. Feng, Y. Yang, B. Huang, X. Luo, M. Li, M. Han, and M. Fu, “Variant selection and the strengthening eﬀect of S precipitates at dislocations in AlCuMg alloy,” Acta Materialia, vol. 59, no. 6, pp. 2412–2422, Apr. 2011.

[83] S. Ringer, K. Hono, I. Polmear, and T. Sakurai, “Precipitation processes during the early stages of ageing in Al-Cu-Mg alloys,” Applied Surface Science, vol. 94-95, no. null, pp. 253– 260, Mar. 1996.

[84] S. Ringer, I. Polmear, K. Hono, and T. Sakurai, “Cluster hardening in an aged Al-Cu-Mg alloy,” Scripta Materialia, vol. 36, no. 5, pp. 517–521, 1997.

[85] S. Ringer, T. Sakurai, and I. Polmear, “Origins of hardening in aged AlCuMg(Ag) alloys,” Acta materialia, vol. 45, no. 9, pp. 3731–3744, 1997.

[86] A. Charai, T. Walther, C. Alfonso, A. Zahra, and C. Zahra, “Coexistence of Clusters, GPB Zones, S”, S’, and S-phase in an Al-0.9Cu-1.4Mg Alloy,” Acta Materialia, vol. 48, no. 10, pp. 2751–2764, 2000.

[87] P. Gouma, D. Lloyd, and M. Mills, “Precipitation processes in AlMgCu alloys,” Materials Science and Engineering: A, vol. 319, pp. 439–442, 2001.

[88] T. Sato, S. Hirosawa, K. Hirose, and T. Maeguchi, “Roles of microalloying elements on the cluster formation in the initial stage of phase decomposition of Al-based alloys,” Metallurgical and Materials Transactions A, vol. 34, no. 12, pp. 2745–2755, Dec. 2003.

[89] A. Serizawa, S. Hirosawa, and T. Sato, “Three-Dimensional Atom Probe Characterization of Nanoclusters Responsible for Multistep Aging Behavior of an Al-Mg-Si Alloy,” Metallurgical and Materials Transactions A, vol. 39, no. 2, pp. 243–251, Jan. 2008.

[90] S. Ringer, B. Sofyan, K. Prasad, and G. Quan, “Precipitation reactions in Al4.0Cu0.3Mg (wt.%) alloy,” Acta Materialia, vol. 56, no. 9, pp. 2147–2160, May 2008.

[91] M. Starink and S. Wang, “The thermodynamics of and strengthening due to co-clusters: General theory and application to the case of AlCuMg alloys,” Acta Materialia, vol. 57, no. 8, pp. 2376–2389, May 2009.

[92] R. Marceau, L. Stephenson, C. Hutchinson, and S. Ringer, “Quantitative atom probe analysis of nanostructure containing clusters and precipitates with multiple length scales,” Ultrami- croscopy, vol. 111, no. 6, pp. 738–742, May 2011.

[93] T. E. M. Staab, B. Klobes, I. Kohlbach, B. Korﬀ, M. Haaks, E. Dudzik, and K. Maier, “Atomic structure of pre-Guinier-Preston and Guinier-Preston-Bagaryatsky zones in Al-alloys,” Jour- nal of Physics: Conference Series, vol. 265, p. 012018, Jan. 2011.

118 [94] L. Kovarik, P. I. Gouma, C. Kisielowski-Kemmerich, S. Court, and M. J. Mills, “High Res- olution Transmission Electron Microscopy Study of the Early Stages of Aging in Al-Mg-Cu Alloys,” Materials Science Forum, vol. 396-402, pp. 845–850, 2002.

[95] A. Zahra, C. Y. Zahra, K. Raviprasad, and I. J. Polmear, “Eﬀects of minor additions of Mg and Ag on precipitation phenomena in Al4 mass% Cu,” Philosophical Magazine, vol. 84, no. 24, pp. 2521–2541, Aug. 2004.

[96] L. Kovarik, P. Gouma, C. Kisielowski, S. Court, and M. Mills, “A HRTEM study of metastable phase formation in AlMgCu alloys during artiﬁcial aging,” Acta Materialia, vol. 52, no. 9, pp. 2509–2520, May 2004.

[97] A. Zahra, C. Zahra, and M. Dumont, “Eﬀects of Ag or Si on precipitation in the alloy Al-2.5 mass% Cu-1.5 mass% Mg,” Philosophical magazine, vol. 85, no. 31, pp. 3735–3754, 2005.

[98] a. M. Zahra, C. Y. Zahra, and B. Verlinden, “Comments on Room-temperature precipitation in quenched AlCuMg alloys: a model for the reaction kinetics and yield-strength development,” Philosophical Magazine Letters, vol. 86, no. 4, pp. 235–242, Apr. 2006.

[99] L. Kovarik, S. Court, and M. Mills, “STEM HAADF Observations of GPB Zones in Al-Cu-Mg Alloys,” Microscopy and Microanalysis, vol. 13, no. S02, pp. 1064–1065, Aug. 2007.

[100] L. Kovarik and M. Mills, “Structural relationship between one-dimensional crystals of GuinierPrestonBagaryatsky zones in AlCuMg alloys,” Scripta Materialia, vol. 64, no. 11, pp. 999–1002, Jun. 2011.

[101] L. Kovarik and M. Mills, “Ab initio analysis of GuinierPrestonBagaryatsky zone nucleation in AlCuMg alloys,” Acta Materialia, vol. 60, no. 9, pp. 3861–3872, May 2012.

[102] M. Starink, “Analysis of aluminium based alloys by calorimetry: quantitative analysis of reactions and reaction kinetics,” International Materials Reviews, vol. 49, no. 3, pp. 191–226, Jun. 2004.

[103] A. K. Gupta, P. Gaunt, and M. C. Chaturvedi, “The crystallography and morphology of the S’-phase precipitate in an Al(CuMg) alloy,” Philosophical Magazine A, vol. 55, no. 3, pp. 375–387, Mar. 1987.

[104] D. Wang, D. Ni, and Z. Ma, “Eﬀect of pre-strain and two-step aging on microstructure and stress corrosion cracking of 7050 alloy,” Materials Science and Engineering: A, vol. 494, no. 1-2, pp. 360–366, Oct. 2008.

[105] J. Embury and R. Nicholson, “The nucleation of precipitates: The system Al-Zn-Mg,” Acta Metallurgica, vol. 13, no. 4, pp. 403–417, Apr. 1965.

[106] H. L¨oﬄer,I. Kovacs, and J. Lendvai, “Decomposition processes in Al-Zn-Mg alloys,” Journal of Materials Science, vol. 18, no. 8, pp. 2215–2240, 1983.

[107] H. Okuda, M. Tanaka, K. Osamura, and Y. Amemiya, “Partial Reversion of Small GP Zones in an Al-Zn Binary Alloy,” Journal of Applied Crystallography, vol. 30, no. 5, pp. 592–596, 1997.

[108] C. Macchi, A. Somoza, A. Dupasquier, and I. Polmear, “Secondary precipitation in AlZnMg(Ag) alloys,” Acta Materialia, vol. 51, no. 17, pp. 5151–5158, Oct. 2003.

119 [109] O. Glatter and O. Kratky, Small Angle X-Ray Scattering. New York: Academic Press, Inc., 1982. [110] S. Acco, D. Williamson, P. Stolk, F. Saris, M. van den Boogaard, W. Sinke, W. van der Weg, S. Roorda, and P. Zalm, “Hydrogen solubility and network stability in amorphous silicon,” Physical Review B, vol. 53, no. 8, pp. 4415–4427, Feb. 1996. [111] ASM Handbook Volume 8, 2003. [112] O. Rios, “Advanced High-temperature Shape-memory Alloy Development and Thermome- chanical Characterization of Platinum and Palladium Modified NiTi Based SMAs,” Master’s Thesis, University of Florida, 2006. [113] P. Wilkes, “On the anomalous electrical resistance of GP zones,” Acta Metallurgica, vol. 16, no. 6, pp. 863–866, 1968. [114] A. Hillel, “Calculations on the Resistivity of Alloys During Pre-Precipitation,” Acta Metal- lurgica, vol. 18, no. 2, pp. 253–259, 1970. [115] G. Bischof, V. Groger, G. Krexner, and R. M. Nieminen, “Investigation of the composition and structure of GP zones in-Ag by means of positron annihilation,” Journal of Physics: Condensed Matter, vol. 8, pp. 7523–7537, 1996. [116] S. Esmaeili, W. J. Poole, and D. J. Lloyd, “Electrical Resistivity Studies on the Precipitation Behaviour of AA6111,” Materials Science Forum, vol. 331-337, pp. 995–1000, 2000. [117] R. Ferragut, A. Somoza, and I. Torriani, “Pre-precipitation study in the 7012 AlZnMgCu alloy by electrical resistivity,” Materials Science and Engineering A, vol. 334, no. 1, pp. 1–5, 2002. [118] S. Esmaeili, D. Lloyd, and W. Poole, “Modeling of precipitation hardening for the naturally aged Al-Mg-Si-Cu alloy AA6111,” Acta Materialia, vol. 51, no. 12, pp. 3467–3481, Jul. 2003. [119] S. Esmaeili, D. Lloyd, and W. Poole, “Effect of natural aging on the resistivity evolution during artificial aging of the aluminum alloy AA6111,” Materials Letters, vol. 59, no. 5, pp. 575–577, Feb. 2005. [120] S. Esmaeili, D. Vaumousse, M. W. Zandbergen, W. J. Poole, a. Cerezo, and D. J. Lloyd, “A study on the early-stage decomposition in the AlMgSiCu alloy AA6111 by electrical resistivity and three-dimensional atom probe,” Philosophical Magazine, vol. 87, no. 25, pp. 3797–3816, Sep. 2007. [121] H. Seyedrezai, “Early Stages of Ageing in Al-Mg-Si Alloys Early Stages of Ageing in Al-Mg-Si Alloys,” Ph.D. dissertation, McMaster University, 2007. [122] H. Seyedrezai, D. Grebennikov, P. Mascher, and H. S. Zurob, “Study of the early stages of clustering in AlMgSi alloys using the electrical resistivity measurements,” Materials Science and Engineering: A, vol. 525, no. 1-2, pp. 186–191, Nov. 2009. [123] P. Desai, H. James, and C. Ho, “Electrical resistivity of aluminum and manganese,” Journal of Physical Chemistry Reference Data, vol. 13, no. 4, pp. 1131–1172, 1984. [124] A. Beck, M. Heine, E. Caule, and M. Pryor, “The kinetics of the oxidation of Al in oxygen at high temperature,” Corrosion Science, vol. 7, no. 1, pp. 1–22, Jan. 1967.

120 [125] A. Rosen and S. Bodner, “Repeated Discontinuous Yielding of 2024 Aluminum Alloy,” Ma- terials Science and Engineering, vol. 4, no. 2-3, pp. 115–122, 1969.

[126] S. Ringer, K. Prasad, and G. Quan, “Internal co-precipitation in aged Al1.7Cu0.3Mg0.1Ge (at.%) alloy,” Acta Materialia, vol. 56, no. 9, pp. 1933–1941, May 2008.

[127] H. Zurob and H. Seyedrezai, “A model for the growth of solute clusters based on vacancy trapping,” Scripta Materialia, vol. 61, no. 2, pp. 141–144, Jul. 2009.

[128] R. Ferragut, A. Dupasquier, C. Macchi, A. Somoza, R. Lumley, and I. Polmear, “Vacan- cysolute interactions during multiple-step ageing of an AlCuMgAg alloy,” Scripta Materialia, vol. 60, no. 3, pp. 137–140, Feb. 2009.

[129] T. Honma, S. Yanagita, K. Hono, Y. Nagai, and M. Hasegawa, “Coincidence Doppler broaden- ing and 3DAP study of the pre-precipitation stage of an AlLiCuMgAg alloy,” Acta Materialia, vol. 52, no. 7, pp. 1997–2003, Apr. 2004.

[130] J. M. Hyde, A. Cerezo, and T. J. Williams, “Statistical analysis of atom probe data: detecting the early stages of solute clustering and/or co-segregation.” Ultramicroscopy, vol. 109, no. 5, pp. 502–509, 2009.

[131] T. Ogura, S. Hirosawa, A. Cerezo, and T. Sato, “Atom probe tomography of nanoscale microstructures within precipitate free zones in AlZnMg(Ag) alloys,” Acta Materialia, vol. 58, no. 17, pp. 5714–5723, Oct. 2010.

[132] K. Osamura, N. Otsuka, and Y. Murakami, “Resistivity maximum during Guinier-Preston zone formation in an Al-4 wt% Cu alloy,” Philosophical Magazine Part B, vol. 45, no. 6, pp. 583–599, Jun. 1982.

[133] K. Osamura, Y. Murakami, T. Sato, T. Takahashi, T. Abe, and K. Hirano, “Structure of G.P. Zones in an Al-1.7at.%Cu alloy aged for 14 years at room temperature,” Acta Metallurgica, vol. 31, no. 10, pp. 1669–1673, Oct. 1983.

[134] P. Gomiero, F. Livet, Y. Brechet, and F. Louchet, “Microstructure and mechanical properties of a 2091 AlLi alloyI. Microstructure investigated by SAXS and TEM,” Acta Metallurgica et Materialia, vol. 40, no. 4, pp. 847–855, Apr. 1992.

[135] D. Dumont, A. Deschamps, and Y. Brechet, “On the relationship between microstructure, strength and toughness in AA7050 aluminum alloy,” Materials Science and Engineering: A, vol. 356, no. 1-2, pp. 326–336, Sep. 2003.

[136] M. Nicolas and A. Deschamps, “Characterisation and modelling of precipitate evolution in an AlZnMg alloy during non-isothermal heat treatments,” Acta Materialia, vol. 51, no. 20, pp. 6077–6094, Dec. 2003.

[137] M. DUMONT, A. STEUWER, A. DESCHAMPS, M. PEEL, and P. WITHERS, “Microstruc- ture mapping in friction stir welds of 7449 aluminium alloy using SAXS,” Acta Materialia, vol. 54, no. 18, pp. 4793–4801, Oct. 2006.

[138] B. Viguier, “Dislocation densities and strain hardening rate in some intermetallic com- pounds,” Materials Science and Engineering: A, vol. 349, no. 1-2, pp. 132–135, May 2003.

121 [139] S. C. Baik, Y. Estrin, H. S. Kim, and R. J. Hellmig, “Dislocation density-based modeling of deformation behavior of aluminium under equal channel angular pressing,” Materials Science and Engineering: A, vol. 351, no. 1-2, pp. 86–97, Jun. 2003.

[140] M. F. Ashby, “The deformation of plastically non-homogeneous materials,” Philosophical Magazine, vol. 21, no. 170, pp. 399–424, Feb. 1970.

[141] K. Nord-Varhaug, B. r. Forbord, J. Benestad, T. Pettersen, B. Rø nning, a. Bardal, S. Benum, K. Marthinsen, and E. Nes, “Substructure Strengthening in Aluminium Alloys,” Materials Science Forum, vol. 331-337, pp. 1387–1392, 2000.

[142] S. Wang, Z. Zhu, and M. Starink, “Estimation of dislocation densities in cold rolled Al-Mg- Cu-Mn alloys by combination of yield strength data, EBSD and strength models,” Journal of Microscopy, vol. 217, pp. 174–178, Feb. 2005.

[143] F. Barlat, “A simple model for dislocation behavior, strain and strain rate hardening evolution in deforming aluminum alloys,” International Journal of Plasticity, vol. 18, no. 7, pp. 919–939, Jul. 2002.

[144] M. Hordon and B. Averbach, “X-ray measurements of dislocation density in deformed copper and aluminum single crystals,” Acta Metallurgica, vol. 9, no. 3, pp. 237–246, 1961.

[145] M. Hordon and B. Averbach, “Precision density measurements on deformed Copper and Aluminum single crystals,” Acta Metallurgica, vol. 9, no. 3, pp. 247–249, Mar. 1961.

[146] H. Wilsdorf and D. Kuhlmann-Wilsdorf, “Considerations on the Interactions Between Ther- mal Vacancies and Dislocations,” Physical Review Letters, vol. 3, no. 4, pp. 170–172, Aug. 1959.

[147] D. Kuhlmann-Wilsdorf and H. G. F. Wilsdorf, “On The Behavior of Thermal Vacancies in Pure Aluminum,” Journal of Applied Physics, vol. 31, no. 3, p. 516, Mar. 1960.

[148] M. Kiritani, “Formation of Voids and Dislocation Loops in Quenched Aluminum,” Journal of the Physical Society of Japan, vol. 19, no. 5, pp. 618–631, May 1964.

[149] M. Kiritani, Y. Shimomura, and S. Yoshida, “Shape of Voids in Quenched Aluminum,” Journal of the Physical Society of Japan, vol. 19, no. 9, pp. 1624–1631, Sep. 1964.

[150] M. Kiritani and S. Yoshida, “An Evidence of Voids in Quenched Aluminum,” Japanese Jour- nal of Applied Physics, vol. 4, no. 2, pp. 148–149, Feb. 1965.

[151] K. Sezaki and D. Kuhlmann-Wllsdorf, “Evidence for the destruction of voids through glide in aluminum,” Acta Metallurgica, vol. 14, no. 9, pp. 1131–1133, Sep. 1966.

[152] M. Kubota, J. Nie, and B. Muddle, “Characterisation of precipitation hardening response and as-quenched microstructures in Al-Mg (-Ag) alloys,” Materials transactions, vol. 45, no. 12, pp. 3256–3263, 2004.

[153] M. Kubota, “Characterisation of precipitate microstructures of natural and artiﬁcial ageing in Al-Mg (-Ag) alloys,” Materials transactions, vol. 46, no. 2, pp. 241–250, 2005.

[154] M. Asta and J. Hoyt, “Thermodynamic properties of coherent interfaces in f.c.c.-based AgAl alloys: a ﬁrst-principles study,” Acta Materialia, vol. 48, no. 5, pp. 1089–1096, Mar. 2000.

122 [155] M. Asta, S. Foiles, and A. Quong, “First-principles calculations of bulk and interfacial thermodynamic properties for fcc-based Al-Sc alloys,” Physical Review B, vol. 57, no. 18, pp. 11 265–11 275, May 1998.

[156] C. Wolverton and V. Ozoliˇs, “Entropically Favored Ordering: The Metallurgy of Al2Cu Revisited,” Physical Review Letters, vol. 86, no. 24, pp. 5518–5521, Jun. 2001.

[157] V. Vaithyanathan, C. Wolverton, and L. Chen, “Multiscale Modeling of Precipitate Mi- crostructure Evolution,” Physical Review Letters, vol. 88, no. 12, pp. 125 503–, Mar. 2002.

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