School of Materials Science and Engineering
A thesis submitted in complete fulfillment of the requirements for the award of the degree of
Doctor of Philosophy in
Materials Science and Engineering
“Accumulative Roll Bonding of Multilayered Aluminium Alloys”
Oday Al-Buhamad August 2009 Acknowledgments ______
First of all, I would like to convey my sincere gratitude to my advisor Professor
Michael Ferry for his continuous guidance, support and patience throughout the past five years, needless to mention his bright scientific ideas for overcoming many hurdles with the progress of this project; I am so honored to be your student. I am also obliged for Professor Paul Munroe my co-advisor for his supportive discussions during this project. Dr. Zakaria Quadir, your teaching, invaluable advice and extended discussions propelled me in the right direction and largely excelled my efforts for finalizing this project, I am so indebted and appreciative to you mate.
I would like also to thank;
Dr. Tania Vodenitcharova for the useful conversation about mechanical properties data interpretations, Dr. Martin Xu for his help in the casting at early stage of the project, Dr. Bulent Gun for his help in the statistical analysis, Nanang Burhan for using his cast Scandium samples, Dr. Nora Mateescu for her support in the Dual beam
FIB, A/Prof. Lori Bassman & Dr. Philip Boughton for their support in tensile testing set-up & discussion, Mathew Pincott for his help in formatting the thesis. Also, Lana
Strizhevsky, Cathy Lau, Anil-Prakash, Jane Gao, Flora Lau, Philip Chatfield and
Dr. George Yang for their support in various occasions.
Finally, my beloved wife Heba and daughters Alia & Rawya for their patience and moral support throughout this project and of course the blessing & supplication of my parents.
ii
Dedicated to my family
(My Darling & the 2-Butterflies)
& for
Cherishing the everlasting
Memories
of splendid Sydney
iii Abstract ______
Multilayered aluminium alloy composites were produced by accumulative roll bonding
(ARB) to very high strain to generate sheet materials consisting of either 32 or 64 alternating layers of Al and Al-0.3w.%Sc alloy. Based on the starting heat treatment condition of the Al(Sc) alloy and the roll bonding temperature, several different
Al/Al(Sc) combinations were produced: (i) SSSS-ARB (Al(Sc) in the supersaturated condition; Tdef = 200 C; 32 layers); (ii) Aged-ARB (Al(Sc) in the artificially aged condition; Tdef = 200 C; 32 layers), and (iii) SSSS-ARB-HT (Al(Sc) in the SSSS condition; Tdef = 350 C; 64 layers). Regardless of the roll bonding conditions, Al(Sc) in the form of a dispersion of ultrafine Al3Sc particles strongly impedes structural changes during thermo-mechanical processing whereas Al readily undergoes extensive dynamic and static restoration.
The major aim of the thesis is to understand the effect of initial microstructure and processing conditions on microstructural development in these multilayered Al/Al(Sc) composites. The microstructures were investigated mainly by backscatter electron (BSE) and ion channeling contrast (ICC) imaging in the DualBeam Platform and transmission electron microscopy (TEM) whereas the crystallographic nature of the microstructures were investigated by electron backscatter diffraction (EBSD) and the various diffraction techniques available in the TEM. The mechanical properties of the materials were investigated by hardness and tensile testing.
iv The deformation microstructure and texture of these two alloy combinations were strongly influenced by both the initial heat treatment condition of the Al(Sc) alloy whereby large-scale shear bands are generated during rolling when a dispersion of fine
Al3Sc particles is present in the Al(Sc) layers. The deformation mechanism of both SSSS-ARB and Aged-ARB was strongly controlled by the relative hardening behaviour of adjacent layers. In Aged-ARB, a higher magnitude of in-plane shear stress, exceeding the flow stress of Al(Sc), was operative at the interfaces between layers; this was shown to cause the shear banding in this material.
All materials were annealed for up to 6h at 350 °C. This extended annealing generated alternating layers of coarse grains (Al layers) and a recovered substructure (Al(Sc) layers) with the substantial waviness of the layers in both Aged-ARB and SSSS-ARB- HT being inherited from the as-deformed material. While the Al(Sc) layers remain unrecrystallized in all materials due to particle pinning effects, the Al layers underwent continuous and discontinuous recrystallization after low and high temperature roll bonding, respectively. Shear banding in Aged-ARB also resulted in a reduction in intensity of the rolling texture components and had a randomizing effect on the recrystallization texture of the Al layers.
The Al/A(Sc) multilayered composites were found to conform to the classic inverse strength/ductility relationship and no significant improvement in ductility (for a given strength) was evident. The barriers to achieving an excellent combination of ductility and strength (i.e. toughness) in these materials were identified to be delamination of the layers, which can be largely reduced (or eliminated) by careful control of starting materials (heat treatment condition and thickness) as well as the processing parameters during ARB.
v Table of Content
______
Chapter 1 INTRODUCTION 1
Chapter 2 LITERATURE REVIEW – PART (I) Severe Plastic Deformation Processing (SPD) 3
2.1 Introduction 3 2.2 The Nature and Characteristics of SPD 5 2.2.1 Industrial applicability of SPD 6 2.3 Accumulative Roll Bonding (ARB) 7 2.3.1 Advantages of ARB 8 2.3.2 ARB of aluminium alloys 9 2.4 The Microstructure of Cold Rolled Metals 9 2.4.1 Cells and subgrains 10 2.4.2 Microbands 11 2.4.3 Shear bands 11 2.4.4 Lamellar bands 12 2.4.5 Typical microstructures of ARB-processed alloys 13 2.5 Texture Development during Rolling and Annealing 15 2.5.1 Introduction 15 2.5.2 Texture measurement 15 2.5.2.1 Electron backscatter diffraction (EBSD) 16 2.5.3 Texture representation in sheet metals 17 2.5.3.1 Pole figures 17 2.5.3.2 Inverse pole figures 18 2.5.3.3 Orientation distribution function (ODF) 18 2.5.4 Texture development in cold-rolled aluminium alloys 20 2.5.4.1 Deformation (rolling) textures 20 2.5.4.2 Recrystallization textures in rolled aluminium alloys 21 2.5.5 Texture of ARB-processed aluminium alloys 22 2.6 Mechanical Behaviour of ARB-Processed Metals 23
Chapter 3 LITERATURE REVIEW–PART (II) Static Restoration Processes Pertinent to SPD 26
3.1 Introduction 26 3.2 Static Restoration Processes 27 3.2.1 Recovery 27 vi 3.2.2 Static recrystallization 28 3.2.2.1 Nucleation of recrystallization 28 3.2.2.2 Growth of nuclei 29 3.3 Classification of Static Restoration Processes 29 3.3.1 Continuous and discontinuous recrystallization 30 3.3.1.1 Continuous recrystallization in highly strained alloys 31 3.3.1.2 Transition from continuous to discontinuous recrystallization 31 3.4 The Effect of Second-Phase Particles on Static Restoration 32 3.4.1 Principles of Zener pinning 32 3.4.2 Grain growth kinetics in the presence of a particle dispersion 33 3.5 Physical Metallurgy of Al-Sc Alloy 34 3.5.1 Attributes of scandium (Sc) in aluminium 35 3.5.2 Formation of Al3Sc during solidification 35 3.5.3 Precipitation of Al3Sc from supersaturated solid solution 36 3.5.4 Effect of Al3Sc on deformation and annealing 37 3.5.4.1 Deformation microstructure 38 3.5.4.2 Recrystallization behaviour 38 3.5.5 Influence of Al3Sc dispersoids on mechanical properties 39 3.6 Summary and Scope of Thesis 40
Chapter 4 EXPERIMENTAL PROCEDURE 43
4.1 Introduction 43 4.2 Materials Processing 43 4.2.1 Preliminary processing of candidate materials 43 4.2.2 Accumulative roll bonding (ARB) 44 4.2.3 Post-deformation annealing 46 4.3 Mechanical Testing 46 4.3.1 Hardness testing 46 4.3.2 Tensile testing 47 4.4 Microstructural Analysis 47 4.4.1 Focused ion beam imaging 47 4.4.2 Electron backscatter diffraction (EBSD) 48 4.4.3 Transmission electron microscopy 49 4.4.3.1 Site-specific TEM sample preparation 49 4.4.3.2 TEM analysis 50
Chapter 5 RESULTS-I Deformation and Annealing Behaviour of Al/Al(Sc) Lamellar Composite 51 (32 Alternating Layers: SSSS-ARB)
5.1 Introduction 51 5.2 Deformation Microstructures 52 5.2.1 Hardness of individual layers during ARB 52 5.2.2 Microstructural development within the Al and Al(Sc) layers 53 5.3 Deformation Textures 56 5.3.1 Texture of the Al layers after ARB 56 5.3.2 Texture of the Al(Sc) layers after ARB 57
vii 5.4 Microstructural Evolution During Annealing 58 5.4.1 Development of microstructure and texture of Al and Al(Sc) layers after annealing for 3 min at 350°C 58 5.4.2 Rapid coarsening of certain grains within the Al layers 62 5.4.3 Microstructure and texture of Al- and Al(Sc) layers after 63 annealing for 6 h at 350 °C 5.4.4 Al/Al(Sc) interface behaviour during annealing at 350°C 64
Chapter 6 RESULTS-II Deformation and Annealing Behaviour of Al/Al(Sc) Lamellar Composite 65 (32 Alternating Layers: Aged-ARB)
6.1 Introduction 65 6.2 Starting Microstructure of Al(Sc) Alloy 66 6.3 Deformation Microstructures 67 6.3.1 Hardness of individual layers during ARB 67 6.3.2 Microstructural development within the Al and Al(Sc) layers 67 6.4 Deformation Textures 69 6.4.1 General texture development 69 6.4.2 Effect of shear banding on local texture development 70 6.5 Microstructural Evolution During Annealing 71 6.5.1 Microstructure and texture of Al and Al(Sc) layers after annealing for 3 min at 350 °C 71 6.5.2 Microstructure and texture of Al and Al(Sc) layers after annealing for 6 h at 350 °C 71 6.6 Summary 73
Chapter 7 RESULTS-III Microstructure and Mechanical Properties of Al/Al(Sc) Lamellar Composite 74 (64 Alternating Layers: SSSS-ARB-HT)
7.1 Introduction 74 7.2 Microstructural Development During Deformation and Annealing 75 7.2.1 General observations 75 7.2.2 Deformation textures within individual deformed layers 76 7.2.3 Texture development within individual layers during annealing 78 7.2.3.1 Microstructure after annealing at 250 C78 7.2.3.2 Microstructure after annealing at 300 C80 7.2.3.3 Microstructure after annealing at 350 C81 7.3 Mechanical Properties 83 7.4 Summary 84
Chapter 8 DISCUSSION 86
8.1 Introduction 86 8.2 Deformation and Stability of Multilayered Structures 89 viii 8.2.1 Microstructural development in the Al and Al(Sc) layers during ARB 89 8.2.2 Effect of scandium on differential hardening of the Al and Al(Sc) layers in SSSS- and Aged-ARB 93 8.2.3 Origin of large-scale shear banding in Aged-ARB during ARB 95 8.2.4 Texture development of the Al and Al(Sc) layers during ARB 99 8.2.4.1 Absence of large-scale shear banding (SSSS-ARB) 99 8.2.4.2 Effect of large-scale shear banding (Aged-ARB) 100 8.3 Annealing of As-Deformed Al/(Al(Sc) Composites 103 8.3.1 Annealing behaviour following ARB at low homologous temperature 103 8.3.1.1 Continuous and discontinuous recrystallization of the Al layers 104 8.3.1.2 Elimination of Brass texture component in Al layers during annealing 106 8.3.1.3 Influence of shear banding on recrystallization textures in Aged-ARB 109 8.3.2 Annealing behaviour following ARB at high homologous temperature 111 8.3.2.1 Microstructural development in Al layers during ARB in the presence of concomitant recovery and recrystallization 111 8.3.2.2 Development of microstructure and texture during annealing 112
8.3.3 Effect of Al3Sc dispersoids on structural stability of the Al(Sc) layers 112 8.3.3.1 Kinetics of coarsening of Al3Sc dispersoids 113 8.3.3.2 Conditions for retardation of recrystallization in the Al(Sc) layers 115 8.3.3.3 Particle-controlled subgrain coarsening in the Al(Sc) layers 117 8.4 Strength and Ductility of Al/Al(Sc) Composites 119 8.5 Control of Annealing Texture and Formability 121
Chapter 9 CONCLUDING SUMMARY 123
Chapter 10 REFRENCES 130
ix List of Publications
The following publications were generated throughout the course of the thesis:
Refereed Papers (DEST-C1) 1. O. Al-Buhamad, M.Z. Quadir, and M. Ferry: “Mechanical Properties of 6-Cycles ARB Processed Al/Al-0.3%Sc Composite”. Materials Science Forum Vols. 618-619 (2009) pp 551-554. 2. M.Z. Quadir, O. Al-Buhamad, KD. Lau, R. Quarfoth, L. Bassman, P.R. Munroe and M. Ferry, “The effect of initial microstructure and processing temperature on microstructure and texture in multilayered Al/Al(Sc) ARB sheets”, International Journal of Materials Research, Vol 100, pp. 1705-1714, 2009. 3. M.Z. Quadir, M. Ferry, O. Al-Buhamad, and P.R. Munroe: “Shear banding and randomization of the recrystallization texture in Al layers in a multilayered Al-based ARB sheet”, Acta Materialia, Vol 57, pp. 29-40, 2009. 4. M.Z. Quadir, O. Al-Buhamad, L. Bassman and M. Ferry: “Development of a recovered/recrystallized multilayered Microstructure in Al Alloys by Accumulative Roll Bonding”, Acta Materialia. Vol 55, pp. 5438-5448, 2007. 5. M. Z. Quadir, O. Al-Buhamad and M Ferry: “Hard and Soft Layered Composite by Accumulative Roll Bonding (ARB)”, Materials Science Forum. Vol. 558-559, pp. 307- 312, 2007. 6. M.Z. Quadir, L. Bassman, O. Al-Buhamad and M. Ferry: “Formation of recrystallized and recovered alternative layers in Al alloys” Materials Science Forum. Vol. 561-565, pp. 131-134, 2007.
Refereed Papers (DEST-E1) 7. O. Al-Buhamad, M.Z. Quadir and M. Ferry: “Some Aspects of Grain Coarsening in an ECAP Deformed and Heat Treated Al-Hf Alloy”, Proceedings of ICOTOM-15 (International Conference of Texture of Materials), CSA Publication, Pittsburgh, USA, 2008. 8. M. Z. Quadir, O. Al-Buhamad, L. Bassman and M. Ferry: “Texture Randomization by Intense Shearing in Layered Al And Al(0.3%Sc) ARB Processed Sheet”. Proceedings of ICOTOM-15 (International Conference of Texture of Materials), CSA Publication, Pittsburgh, USA, 2008.
x Chapter 1 ______INTRODUCTION
Accumulative roll bonding (ARB) is a severe plastic deformation (SPD) process first developed in 1998. It involves the simultaneous rolling of two stacked sheets of material of similar thickness. In the first rolling pass, the thickness of both sheets is reduced to the initial sheet thickness. This procedure is repeated several times with the outcome being the production of roll-bonded sheet whereby very large strains are accumulated thereby resulting in considerable microstructural refinement.
Almost all SPD processes (equal channel angular pressing, high pressure torsion etc.) readily produce ultra-fine-grain (UFG) microstructures in a diverse range of metallic materials. Nevertheless, the amount of material produced in a given processing stage is usually small. A key element in ARB is the ability to continuously produce material in a manner similar to conventional rolling. ARB has other useful advantages as there is no major requirement for forming facilities with large load capacity and expensive dies, the productivity rate can be high, and large volumes of
1 material can be produced. Therefore, ARB is highly appropriate for manufacturing
UFG alloys in the form of sheets, bars and plates; these shapes are the most widely used in many commercial applications.
A major advantage of ARB is the capacity to produce multilayered sheet composites containing alternating layers of two or more dissimilar metals. This enables the creation of sheet products where the overall properties are based on the combined properties of the individual alloys and may therefore provide scope for extending into new regions of property space. A major objective of this thesis is the development of multilayered Al alloy composites by ARB using two alloys that respond differently to thermomechanical processing. Here, commercial purity Al and an Al-0.3% Sc alloy will be roll-bonded in alternating sequence and subsequently annealed. These alloys were selected due to the significant knowledge generated over recent years concerning their behaviour during and after SPD. For example, Al(Sc) generates a dispersion of ultra-fine Al3Sc particles that strongly impede structural changes during thermomechanical processing whereas Al readily undergoes extensive dynamic and static restoration.
This thesis describes the development of multilayered Al/Al(Sc) composites by ARB and annealing where the novelty lies in materials that have never been combined in a manner; this may exploit the useful properties of the individual components. It involves a systematic study, using advanced characterization techniques, of the effect of initial microstructure and processing conditions on microstructural development by exploiting the strong influence of ultrafine particles on inhibiting recrystallization.
It is expected that these multilayered alloys will generate a microstructure and texture specific to the individual layers, thereby generating composite structures that, potentially, have a unique combination of properties.
2 Chapter 2 ______LITERATURE REVIEW – PART I Severe Plastic Deformation Processing
2.1 Introduction
With the advent of nanocrystalline (NC) fabrication processes through the study of
Gleiter (1981) on synthesized nanostructured metals using inert gas condensation, research on nanostructured materials is now a major focus in materials science
(Meyers et al., 2006; Li et al., 2008). The development of the concept of ultrafine grains can be traced to the work of both Embury and Fisher (1966) and Armstrong et al. (1966). The term ultra-fine grain size (UFG) is regularly used to describe grain sizes in the range ~250-1000 nm, which constitutes the upper limit of the wider regime of nanocrystalline (NC)/nanostructured (NS) materials that comprise the grain size scale ranging from 1 to 250 nm; the upper range also encompasses so- called submicron grained (SMG) materials (Koch, 2002; Meyers et al., 2006).
The development of NC/NS and SMG materials has greatly attracted the attention of the materials research community. Here, two complementary approaches have been
3 developed in the attempt to synthesize/fabricate these types of material (Zhu et al.,
2004; Meyers et al., 2006; Viswanathan et al., 2006; Azushima et al., 2008):
(i) Bottom-Up Approach – This involves assembling NS materials from individual atoms or from nano-scale “building blocks” such as nano-particles, ions or molecules. Techniques in this category include inert gas condensation (Gleiter,
1989), electro-deposition (Erb, 1995) and chemical and physical deposition
(Suryanarayana, 1999), sintering (Chen and Wang, 2000), nanoconsolidation
(Vassen and Stover, 2001), and other non-equilibrium methods. The disadvantages here are the generated flaws such as porosity and trapped gases that naturally exist in these consolidated NS materials (Koch, 2002; Lowe and Zhu, 2003; Zhu, 2006).
(ii) Top-Down Approach – This involves processing initially coarse-grained materials to generate substantial grain refinement, that is, by breaking down the microstructure of the starting bulk materials. The most successful processing methods involve severe plastic deformation (SPD) which is a mechanical working technique. SPD methods that utilize plastic deformation have a number of advantages over bottom-up approach processes, by accommodating less contamination and the ability to fabricate bulk components with very low porosity.
However, the prime disadvantage of SPD is the generation of high residual internal stresses that may lead to some degree of instability in the microstructure and result in variable properties (Shaw, 2000; Valiev et al., 2000a,b; Viswanathan et al., 2006).
SPD can be defined as those metal forming processes that generate very large plastic strains into a bulk metal, and requires processing below a temperature where recrystallization does not occur readily (Valiev et al., 1993; Valiev et al., 2000a;
Tsuji et al., 2003; Horita et al., 2006). Such SPD-produced materials have grain sizes in the range 100-1000 nm, and have a deformation substructure consisting of cells or
4 subgrains typically smaller than 100 nm. Hence, they are also often termed NS materials (Jiang et al., 2000).
A major objective of SPD is to fabricate high strength, lightweight components. The typical true plastic strains achievable in conventional metal forming processes (e.g. rolling, forging and extrusion) is usually less than ~ 2, but these processes can also be carried out to generate much larger plastic strains. Unfortunately, such processes produce extremely thin samples in at least one of the dimensions and, thus, are inadequate as structural components.
Therefore, to impose an extremely large strain without drastically changing the shape of the bulk metal, many redundant SPD processes have been developed including equal channel angular pressing (ECAP) (Segal, 2002; Valiev et al., 2006; Valiev and
Langdon, 2006), accumulative roll-bonding (ARB) (Saito et al., 1998a; Tsuji et al.,
1999), high pressure torsion (HPT) (Valiev et al., 1991; Valiev, 1997), repetitive corrugation and straightening (RCS) (Huang et al., 2001a), cyclic extrusion compression (CEC) (Korbel et al., 1983), torsion extrusion (Mizunuma, 2006), severe torsion straining (STS) (Nakamura et al., 2004), cyclic closed-die forging
(CCDF) (Lowe and Valiev, 2000), super short multi-pass rolling (SSMR) (Manabu et al., 2008), and friction stir processing (FSP) (Mishra et al., 2000; Su et al., 2003).
Some of the more common SPD processes are shown schematically in Fig. 2.1.
2.2 The Nature and Characteristics of SPD
The fundamental principle for the foregoing SPD methods is to enforce an enormously high strain into the material such that structural refinement occurs by shear and fracture of phases along with recrystallization processes (most likely dynamic recrystallization). Here, the ultimate microstructural outcome is determined by a balance between the rates of work hardening and dynamic recovery. During
5
Figure 2.1. Schematic illustrations showing the principles of three types of SPD: HPT (high pressure torsion), ECAE/P (equal channel angular extrusion/pressing) and CEC (cyclic extrusion and compression), after Tsuji (2006).
SPD, the main parameters controlling the resultant microstructure includes temperature and strain rate of deformation, imposed pressure, lubrication, intersecting angle (c.f. ECAP) and degree of strain (Valiev et al., 1994; Ferrasse et al., 1997; Kawazoe et al., 1997; Popov et al., 1997; Valiev, 1997; Saito et al., 1998a;
Tsuji et al., 1999; Shaw, 2000). For many metals, these processing routes may generate very fine grain sizes down to 20 nm and dislocation densities as high as
5×10-14 m–2 (Valiev, 1997).
In developing these SPD processing techniques to fabricate UFG and NC/NS materials, certain requirements must also be fulfilled including (Valiev et al., 2000a):
(i) generation of UFG structures with a large fraction of high angle grain boundaries
(HAGBs); (ii) the formation of a uniform nanostructures (NS) throughout the work- piece, and (iii) minimization of cracking or other damage due to the large plastic strains. These requirements can be met through the application of unique mechanical working schemes in relevance to individual SPD processes.
2.2.1 Industrial applicability of SPD
NS/NC materials produced by SPD have the capacity to be used in large-scale industrial applications primarily due to the generation of a combination of superior properties and manufacturability (performance). Furthermore, SPD is attractive as the resultant materials can be tailored to the desired end use (i.e. good machinability, forgability and formability) at potentially low processing cost (Zhu et al., 2004;
Lowe, 2006). Therefore, NS/NC components may be used in specialized structural applications in various industrial sectors such as medical, aerospace, automobile, aircraft, defence, sports, manufacturing and biomedical etc. A notable successful example shown in Fig. 2.2a is the SPD-fabricated Ti-alloy screws by which is used in the biomedical industries (Zhernakov and Yakupov, 1997). Another example is the
‘Piston’ type complex-shaped component (Valiev et al., 1993) made of NS AA1420
6
Figure 2.2. Examples of commercial applications of SPD-produced components: (a) plate implants made of NS Ti and (b) ‘‘Piston’’ component fabricated from NS AA1420 Al, after Azushima et al. (2008).
Al for the automobile industry, Fig. 2.2b. Other effective applications and industrial implementation of SPD-processed nano-materials can be found in the literature
(Lowe and Valiev, 2000; Valiev et al., 2000a,b; Tsuji et al., 2003; Zhu et al., 2004;
Piers Newbery et al., 2006; Azushima et al., 2008).
2.3 Accumulative Roll Bonding (ARB)
Accumulative roll bonding (ARB) was developed initially by Saito and co-workers
(Saito et al., 1998a). The process involves the simultaneous rolling of two stacked sheets carrying the same thickness. After one ARB pass, the thickness of both sheets is reduced to the initial sheet thickness. This procedure is iterated several times whereby the alignment of the sheets relative to each other can be changed between consecutive ARB cycles. The outcome of repeated rolling is the production of sheets bonded together and the considerable microstructural refinement that usually occurs
(Saito et al., 1998b; Tsuji et al., 1999; Tsuji et al., 2004a).
To achieve high-quality bonding of the sheets, individual sheet surface treatment of degreasing and wire-brushing must be carried out ahead of each stacking stage. In some cases, the stacked sheets are heated to below the recrystallization temperature of the material and roll-bonded immediately to capture excellent bonding and to reduce the rolling force (Tsuji et al., 2003). Overall, ARB is characterized by the combined effect of deformation and sheet bonding and the process can be repeated many times to generate very large plastic strains.
ARB has been used extensively by many researchers and successfully implemented to produces NS/UFG in a range of metallic alloys including Al alloys (Tsuji et al.,
2000; Tsuji et al., 2002a; Xing et al., 2002), Cu-Ag alloy (Ohsaki et al., 2007), Zr alloy (Jiang et al., 2007), IF steel (Kamikawa et al., 2003; Tsuji et al., 2002b), LC
7 steel (Tsuji et al., 2002c), Ni alloy (Koizumi et al., 2005), and multilayer combinations such as Zr/Cu (Tsuji et al., 2005), Al/Cu (Eizadjou et al., 2008) and
Cu/Co/P (Sakai et al., 2001).
A major factor in ARB is ensuring a 30-40% crucial rolling reduction per pass
(Tylecote, 1968), with reduction values often ~50% during the initial rolling cycles.
A potential problem associated with ARB is edge cracking which is largely depending on the nature of processed materials e.g. strain-hardened AA5XXX Al
(Fig. 2.3). These edge cracks have the tendency to spread into the centre of ARB sheet thereby terminating the process (Tsuji et al., 2003). Tsuji (2006) discusses several methods for avoiding edge cracking during ARB.
2.3.1 Advantages of ARB
Almost all the major SPD processes (ECAP, HPT, CEC and RSC) readily produce
UFG microstructures in a diverse range of metallic materials. Nevertheless, the materials dimensions remain small and these techniques are better categorized as batch rather than continuous processes. In this context, ARB has numerous advantages over other SPD processes such as: (i) no requirement for forming facilities with large load capacity and expensive dies; (ii) the productivity rate can be high, and (iii) large volumes of material can be produced. Therefore, ARB is highly appropriate for manufacturing UFG sheets, bars and plates; these are the most widely used shapes in commercial applications. A key element in ARB is the ability to continuously produce materials in a manner similar to conventional rolling.
8
.
Figure 2.3. Typical appearances of ARB-processed sheets: (a) AA1100 Al ARB processed by five cycles at room temperature. (b) AA5083 Al ARB processed by two cycles at room temperature showing extensive tearing, after Tsuji (2006).
2.3.2 ARB of aluminium alloys
Throughout the last decade, ARB of a wide range of Al alloys has been heavily exercised with most alloy types processed successfully to various degrees of strain.
Table 2.1 provides a brief summary of some ARB-processed Al alloys.
Table 2.1. ARB-processed aluminium alloys classified per series designation with some relevant references.
Alloy Designation Reference
AA1XXX; including commercial and high (Tsuji et al., 2000, 2002a; Terada et al., purity aluminium 2007) AA2XXX (Tsuji et al., 2004b) (Xing et al., 2002; Chowdhury et al., AA3XXX 2006b; Pirgazi et al., 2008b) AA4XXX No published attempts AA5XXX (Saito et al., 1999; Song et al., 2006) AA6XXX (Park et al., 2001; Lee et al., 2000, 2002) AA7XXX (Tsuji et al., 2003) (Kim et al., 2002; 2005a,b; Karlik et al., AA8XXX 2004) Al-Li (Chowdhury et al., 2006a) Al-Sc (Min et al., 2005) (Eizadjou et al., 2008; Chen et al., 2006; Bi-Metallic; Al/Cu, Al/Mg, Al/Ni Chang et al., 2009; Min et al., 2006)
2.4 The Microstructure of Cold Rolled Metals
It is well-documented that the two main modes of deformation in metals are slip and twining (Humphreys and Hatherly, 2004). The operation of either deformation mode is largely dependent on the stacking fault energy (SFE) of the material. Twinning occurs readily in low SFE cubic metals e.g. Cu alloys, since the dislocations readily dissociate whereas slip is the dominant mode in higher SFE cubic metals. The
9 principal slip system for face centred cubic (FCC) metals is {111}<110> which consists of the most densely packed planes and directions, respectively.
This functional slip system principally relies on the grain orientation and deformation process, commencing with dislocation glide on the mostly favorably oriented slip system/s. Consequently, this generally leads to the creation of slip bands and slip lines on a polished surface. The former denotes a cluster of slip lines shown earlier by Brown (1952), which are typically viewed through the optical microscope. In contrast, they are hardly observable or traceable with TEM and SEM examinations.
Generally, cells, subgrains and microbands significantly compose up to ~50% of the as-rolled microstructure, and shear bands becoming more common beyond ~50% rolling reduction (Dillamore et al, 1979; Yeung and Duggan, 1986; Lee and Duggan,
1993). The microstructural evolution in high SFE FCC metals such as Al alloys are summarized in the following sections.
2.4.1 Cells and subgrains
Statistically trapped dislocations usually rearrange into an equiaxed cell structure.
The cell interiors have a low dislocation density and are enclosed by tangled dislocation walls of fixed width. Normally, between adjacent cells the orientation variance is low (<2 ). A representative cellular substructure examined under TEM is shown in Fig. 2.4 for cold rolled aluminium (Hansen, 2001). At higher strains or higher deformation temperatures, the cell structure may better be described as subgrains whereby cell walls are sharpened and the cell interiors have a very low dislocation density. During room temperature deformation the associated strain level here is relevant to the melting point of the material and the activated slip systems for an individual grain are dissimilar at every location because of the neighbouring grains effect, i.e. creating Geometrically Necessary Boundaries (GNBs) depicted by grouping of “cell blocks”.
10
Figure 2.4. (a) Bright field TEM micrograph and (b) microstructural sketch in a grain of a 10% cold rolled specimen of high purity Al (99.996%) in longitudinal plane view. One set of extended non-crystallographic dislocation boundaries is marked A, B, C, etc., and their misorientations are shown. The extended boundaries are either single, planar Dense Dislocation Walls (DDWs) or Microbands (MBs), which are plate-like regions formed by two closely spaced
DDWs. These boundaries form cell blocks marked CB1, CB2, etc., which are subdivided by cell boundaries marked a, b, c, etc. The rolling direction is marked RD and the dashed lines are traces of {111} planes (Hansen, 2001).
2.4.2 Microbands
When a particular length of certain GNBs possessing discrete dislocation structures begin to divide, this generates elongated cells surrounded by dense dislocation walls
(DDW) leading to microstructural features termed microbands. Based on the identification used by the RISØ group, first generation microbands are lenticular in character and initiate at ~10% rolling reduction. Further straining results in the microbands resembling lamellar structures divided by DDWs raising the angles of alignment with the rolling plane to ~20 -35 . These microbands are distinctive elements and have been reported in a range of metals rolled to medium-high strain levels. Within a single microband, the orientation range is low, with misorientations with adjacent microbands also low (<2 ). Microbands are argued to be crystallographically oriented along their favored slip planes i.e. bands form parallel to {111} slip planes in copper (Hatherly and Malin, 1979), and {110} slip planes in
LC-Steel (Aghan and Nutting, 1980). When additional deformation is applied, other slip systems are operative thereby generating a roughly uniform slip zone called second generation microbands. These were observed by Aghan and Nutting (1980) and Hughes and Hansen (1993) as shear offsets and were termed S-bands.
2.4.3 Shear bands
Shear bands are microstructural features normally appearing after ~50% rolling reduction that generate highly localized flow during deformation (Duggan et al.,
1978; Dillamore et al, 1979; Yeung and Duggan, 1986; Lee and Duggan, 1993). The investigation by Duggan et al. (1978) on 70/30 brass revealed the large magnitude of shear strain carried by shear bands. In general, shear bands are sheet-like structures created parallel to TD and are projected at 25-40 to the rolling plane. Due to the nature of localized shear, the subgrains within shear bands are more elongated than the matrix (Brown, 1972; Grewen et al., 1977). In the study by Dillamore and co- workers (Dillamore et al, 1979), shear bands are displayed under the optical
11 microscope as darkly etched markings passing through the entire sample thickness; these are categorized as macro-shear bands (Grewen et al., 1977; Stuwe, 1978;
Dillamore et al, 1979). In contrast, micro-shear bands are those features confined largely within single grains, as observed by Yeung and Duggan (1986; 1987). The required conditions of shear band formation with its associated instability source is still a controversial issue, although it is understood that shear bands form whenever homogeneous deformation cannot occur.
2.4.4 Lamellar bands
As noted previously, the dislocation structures generated during cold deformation can be classified as GNBs and IDBs. However, at large strains, this evolution leads to structures composed of dislocation boundaries having a wider range of misorientations with spacings in the submicron range. To further explain this, the cell-block (DDWs) and microband boundaries at low and medium strains have boundary planes that follow certain planes that may or may not be a slip plane. With increasing strain there is an increasing tendency for the dislocation boundaries to reorient from a typical cell block structure into a lamellar structure, as shown schematically in Fig. 2.5. Here, the typical cell block structure of GNBs at low strains includes MBs and single DDWs that surround blocks of equiaxed cells.
However, the typical structure generated at larger strains consists of lamellar boundaries (LBs) that are formed by the sandwiching of thin layers of cells and subgrains oriented along the flow direction (Hughes and Hansen, 1997; Malin and
Hatherly 1979).
Hence, LBs are defined as nearly planar boundaries enclosing blocks of IDBs (Fig.
2.6) and, thus, are classified as GNBs (Hansen and Hughes 1995; Liu and Hansen
1995; Hughes and Hansen 1997). This substantial refinement of the microstructure is followed by an increase in the average misorientation of adjacent cell-blocks
12
Figure 2.5. Schematic diagram of the expected deformation microstructure and grain subdivision in high SFE metals: (a) Small to medium strain deformation showing elongated microbands (MBs) and dense dislocation walls (DDWs) surrounding groups of cells in cell blocks; (b) large strain deformation showing lamellar boundaries (LBs) parallel to the deformation direction, sandwiching narrow slabs of cells or equiaxed subgrains (Hughes and Hansen, 1997).
Figure 2.6. Bright field TEM micrographs showing the lamellar structure and the introduction of dislocations within the lamellae after 15% (a) and 50% (b) cold rolling of samples annealed for 30 min at 150 C (Huang et al., 2008a).
(Hansen and Jensen, 1999). Numerous detailed studies have shown that such deformation-induced boundaries display a large angular spread and can reach average values of 5-10 at intermediate strains, and that the angular distributions at large strain contains a significant fraction of deformation-induced high angle grain boundaries (HAGBs) (Gil Sevillano et al., 1981; Oscarsson et al. 1994; Jensen 1995;
Hughes and Hansen 1997; Liu et al. 1998).
2.4.5 Typical microstructures of ARB-processed alloys
The deformed microstructure of metals and alloys at large strains has been studied extensively over many years (Langford and Cohen, 1975; Sevillano et al., 1981;
Hecker and Stout, 1984). It is known that metallic sheet processed by several ARB cycles contains predominantly UFG microstructures. However, during the initial few cycles the microstructure displays a typical deformation microstructure composing of cells and subgrains, i.e. corresponding to 50-75% cold rolling (Tsuji et al., 2000) with most of the observed boundaries being LAGBs. Thereafter, the number and density of the deformation-induced HAGBs increases with each rolling cycle (i.e. strain) and, eventually, the ARB sheet is comprised largely of elongated bands subdivided by HAGBs.
It is worth noting that newly-formed HAGBs appear mainly in the vicinity of initial grain boundaries and their fraction increases with increasing strain, that is, greater than 80% HAGBs have been observed after five ARB cycles in IF-steel and high purity aluminium sheet. The microstructures observed on both TD and ND sections of ARB-processed materials, after several rolling cycles, show relatively similar structures elongated parallel to RD (Fig. 2.7). They resemble the so-called LBs observed in heavily rolled FCC metals (Hansen and Jensen, 1999). Such observations has been further validated by Kikuchi band analysis in TEM and demonstrated that
13
Figure 2.7. (a) Bright field TEM micrograph showing a lamellar structural morphology and dislocation configurations in a NS Al processed by ARB (six- cycles) to a true equivalent strain of 4.8. (b) Frequency histogram showing the distribution of boundary misorientation angles (Huang et al., 2008b). each elongated region has sub-micron dimensions surrounded mainly by HAGBs (Ito et al., 2000; Tsuji et al., 2000).
Recent work by various research groups on a range of metals (Lee et al., 2002a;
Kamikawa et al., 2004; Kamikawa et al., 2007a) have shown that the redundant shear strain introduced at subsurface regions of ARB sheet due to the significant friction between the rolls and the materials effectively promotes the formation of the UFG microstructure. These UFG microstructures are also uniform throughout the sheet thickness. In a recent study (Li et al., 2006), it was shown that various ARB- processed materials are homogeneously filled with elongated UFG structures in all materials and independent of crystal structure and SFE; this confirmed that any microstructural differences due to the redundant shear strain associated with un- lubricated roll-bonding is homogenized by recurring rolling cycles.
The Riso group have made a substantial contribution to the understanding of microstructural evolution during ARB on the basis of grain subdivision (Hansen et al., 2001). The schematic diagram in Fig. 2.5 shows the classic grain subdivision microstructures. Here, deformation-induced boundaries are classified into categories of IDBs and GNBs (Hansen and Jensen, 1999; Hansen et al., 2001).
In summary, ARB rapidly generates a lamellar microstructure, whereby, with increasing strain, LBs eventually dominate the microstructure and these are reduced in thickness and saturate at a thickness depending on the competing processes of dislocation formation and annihilation by recovery. On a final note, the development of large misorientation angles across dislocation boundaries is more rapid during
ARB compared with conventional rolling. Consequently, for an equivalent level of strain the HAGB fraction is higher in ARB-processed materials (Liu and Hansen,
1999; Huang et al., 2001b; Park et al., 2001; Huang et al., 2003).
14 2.5 Texture Development during Rolling and Annealing
2.5.1 Introduction
The individual crystallites in a material may be randomly oriented or aligned in a particular preferred crystallographic orientation. The sum of the crystallographic orientations of the crystallites within a polycrystalline aggregate is known as the texture of the material (Humphreys and Hatherly, 2004). A texture may develop by processes such as electrodeposition, casting, rolling, forging, extrusion, wire drawing and annealing. The particular type and strength of the texture is governed both by the material (crystal structure, phase distribution, purity etc.) as well as the processing route. Since many aspects of the thesis are concerned with texture development during rolling and annealing, this section provides information required to understand and interpret textures. Further details of textures in materials may be found in Hatherly and Hutchinson (1979), Bunge (1982), Randle and Engler (2000) and Humphreys and Hatherly (2004).
2.5.2 Texture measurement
The texture of a material can be measured by a range of techniques (Randle and
Engler 2000; Humphreys and Hatherly 2004). Macro-methods include X-ray and neutron diffraction and micro-methods include electron diffraction in either the scanning (SEM) or transmission electron microscope (TEM) using the techniques of electron backscatter diffraction (EBSD) and TEM microdiffraction. The commonly used X-ray technique usually involves back reflection to generate a pole figure with the mathematical combination of several pole figures required for more detailed texture information. A limitation of the technique is the small volume of material examined due to the limited depth of penetration (< 0.1 mm) of the X-ray beam into the sample. Furthermore, the technique is of very limited use for micro-texture studies since the texture generated is an average over a large number of crystallites.
15 2.5.2.1 Electron backscatter diffraction (EBSD)
An EBSD system interfaced to a scanning electron microscope (SEM) is a powerful method for studying the crystallographic nature of crystalline materials. This technique involves the automated computer analysis of EBSD patterns to generate crystallographic data on a point-by-point basis over a selected area of a sample at a rate of up to 300 points per second. A typical EBSD setup is shown in Fig. 2.8, which comprises a sensitive camera and an image processing system for pattern averaging and background subtraction. The EBSD acquisition software controls the data acquisition, solves the diffraction patterns and stores the data. Further software is required to analyse, manipulate and display the data.
In the EBSD technique, the electron beam of the SEM is focused at a point on a highly tilted (~60-70o) sample which generates a Kikuchi or backscattered electron diffraction pattern at that point (Fig. 2.8). The EBSD pattern supplies all the crystallographic information of that particular region of the sample. The physics of
Kikuchi electron diffraction is beyond the scope of this thesis but details are given elsewhere (Randle and Engler, 2000).
To generate a map of microstructure based on the information extracted from each diffraction pattern, points are usually arranged in a regular grid and either the electron beam of the SEM is programmed to step to each point in turn or the beam is held stationary and the specimen stage programmed to traverse beneath it. At each step, the coordinates of the point and crystallographic information are recorded and stored using the acquisition software. From these data, a map is constructed to reveal various crystallographic features of the microstructure. For example, the data may be used to show the distribution of texture components of an analysed area as well as
16 calculate the distribution of grain misorientation, average grain size and distribution and host of other important parameters (Humphreys, 2001).
2.5.3 Texture representation in sheet metals
The orientation of a grain in a sheet material can be defined as uvwhkl ])[( where
hkl)( is the Miller indices of a particular plane of the grain, which is parallel to the plane containing the rolling (RD) and transverse direction (TD) while uvw][ is the
Miller indices of a particular direction in the grain, which is parallel to RD. Every crystal within a polycrystalline material is oriented in some way with respect to some type of sample coordinate system; see e.g. Fig. 2.9 for rolled sheet. There are many ways to represent groups of crystal orientations with respect to the external axes of a sample.
The textures discussed in this thesis are produced by rolling and annealing and may be described with pole figures or inverse pole figures. For sheet materials produced by thermo-mechanical processing (TMP), textures are best described using the orientation distribution function (ODF) since the major texture components in sheet metals are usually distributed as tubes of orientation or fibres in three dimensional spaces.
2.5.3.1 Pole figures A pole figure is a stereographic projection showing the distribution of poles of a particular set of crystallographic planes in the assembly of crystallites (grains) that constitutes the specimen (Humphreys and Hatherly, 2004). To describe the pole figure, one or more reference directions need to be assigned such as the drawing direction in wire, or ND, TD and RD in sheet produced by rolling. Fig. 2.10 shows this situation for sheet materials showing the distribution of <100> poles in a grain with respect to the sheet axes. The pole figure of a polycrystalline material has a
17
Figure 2.8. Typical EBSD configuration in the SEM, showing a typical Kikuchi EBSD pattern impinging on a phosphorous screen from a point on the sample surface (Humphreys and Hatherly, 2004).
Figure 2.9. Schematic representation of sample axes associated with rolling. distribution of intensity which may be represented either by individual points on the pole figure (for microdiffraction techniques such as EBSD and TEM) or contour lines with respect to a sample containing grains with a random distribution of orientations. Contoured pole figures can be generated both by global texture analysis
(X-ray or neutron diffraction) and discrete texture measurements (EBSD).
2.5.3.2 Inverse pole figures
The inverse pole figure is a useful mode of representing the texture generated in processes that require the specification of only a single axis. This is suitable for axisymmetric deformation processes such as wire drawing or extrusion. The frequency by which a particular crystallographic direction coincides with the specimen axis is plotted in a single triangle of a stereographic projection. A representative example of a single orientation is shown in Fig. 2.11.
2.5.3.3 Orientation distribution function (ODF)
The orientation distribution function (ODF) allows an orientation to be represented in three dimensions, termed Euler space. It is a powerful method of representing texture because a particular texture component {hkl}
18
Figure 2.10. Schematic representation of pole generation in a typical 100 pole figure.
Figure 2.11. Schematic representation of an inverse pole figure. The orientation (hkl)[uvw] of a grain in rolled sheet can be described in terms of three Euler angles using either Roe or Bunge notation (Bunge, 1982). Using the latter, (hkl) and [uvw] (crystallographic axes) are represented in a standard projection
(Fig. 2.12) and the specimen orientation is specified by ND and RD. The angles and 2 completely specify ND with RD lying in the plane normal to ND. The angle
1 therefore completely specifies RD. The orientation of the grain can be displayed in 3-D with the Euler angles as axes. For cubic materials, the orientation is completely represented within a 909090 volume of Euler space. The advantage of the ODF over pole figures is that a given texture component is completely described by the set of Euler angles: 1 and 2 . It is useful to note that mathematical relationships exist for converting between Miller Indices and Euler angles (Bunge, 1982).
The ODF data of rolled FCC metals are normally represented as a series of slices taken through Euler space at 2 = 0, 5, 10...90 . This is a result of the deformation texture usually consisting of a tube of orientations, called the -fibre, which runs through Euler space from {110}<112> (Brass orientation) through {123}<634> (S orientation) to {112}<111> (Copper orientation) (Humphreys and Hatherly 2004).
Another important orientation tube in FCC metals is the FCC-fibre which runs from
{110}<112> to {110}<001> (Goss orientation). Figure 2.13 is a schematic representation of one of the branches of both the fcc - and -fibres in Euler space.
There are other orientation tubes of lesser importance for metals based on cubic crystal systems, Fig. 2.14, and these are described in detail by Bunge (1982) and Humphreys and Hatherly (2004).
2.5.4 Texture development in cold-rolled aluminium alloys
The orientation changes that take place in a crystalline metal during rolling are due to
19
Figure 2.12. Definition of Euler angles in cubic materials with respect to the standard <001> projection, after Humphreys and Hatherly (2004).
Figure 2.13. 90 × 90 × 90° volume of Euler space showing the fcc- and -fibres that develop in rolled FCC materials, after Hirsch and Lücke (1988a).
Figure 2.14. Plot of two important FCC orientation fibres in Euler angle space: - fibre = <110> axis parallel to ND; -fibre = <110> axis tilted 60 from ND towards RD, adapted from Hirsch and Lücke (1988b) by Humphreys and Hatherly (2004). the fact that deformation occurs on certain slip or twinning systems (Humphreys and
Hatherly, 2004). The strength of the texture and the balance between the various texture components after rolling depends on various factors such as the texture of the as-cast material, degree of deformation, temperature of deformation, crystal structure, stacking fault energy, as-cast grain size and shape and the presence of a second-phase. In rolled metal sheets, textures are classified into deformation
(rolling) textures and recrystallization textures. The grains therefore exhibit certain preferred orientations resulting from their rotations during the fabrication process
(Akef and Driver, 1993). FCC metals have been studied extensively with aluminium alloys being a major focus due to the large expansion in their usage as a structural material.
2.5.4.1 Deformation (rolling) textures
It was reported in section 2.5.3.3 that the principal texture components {112}<111>,
{110}<112>, {123}<634> are those commonly used for describing FCC deformation textures (Grewen and Huber, 1978) as they have attained their relevant names and symbols given in Table 2.2. This table demonstrates the first subspace of Euler angle of main texture components in FCC rolled metals (Humphreys and Hatherly, 2004).
Figure 2.13 showed that the texture in cold rolled high SFE, FCC metals such as aluminium is comprised mainly of the classic -fibre.
Table 2.2. Texture components in rolled FCC metals i.e. first sub-space, after Humphreys and Hatherly (2004).
Component, {hkl}
20 During rolling, the homogeneities along the - and -fibres decrease as strain is increased. This initially occurs to the -fibre and, at ~95% reduction; the entire - fibre essentially vanishes whereas the !-fibre remains prominent. Eventually, by further rolling the !-fibre starts to deteriorate as it is replaced with pronounced peaks
(gradually develops).
Here, the S-component strengthens and becomes the major texture component of heavily rolled aluminium (Humphreys and Hatherly, 2004). It can be summarized that the rolling texture of aluminium at low strains is described by the reasonably homogeneous tube of orientations, whereas at large strain, these tubes decompose into peaks equivalent to the texture components given in Table 2.2. The texture of cold rolled Al-Cu alloy is shown in Fig. 2.15.
It is pertinent to note that some age-hardened aluminium alloys encounter a texture transition, particularly in the Al-Li alloy system, whereby a strong {110}<112> rolling texture develops at high strains (Bowen, 1990; Lucke and Engler, 1990). The controlling mechanism here is chiefly related to the onset of planar slip and shear banding, and is promoted by the presence of deformable second-phase particles in the microstructure (Humphreys and Hatherly, 2004).
2.5.4.2 Recrystallization textures in rolled aluminium alloys
The texture that develops when a deformed metal is annealed has been the subject of extensive study (Humphreys and Hatherly, 2004). Similar to deformation, the strength and type of texture components that develop during recrystallization are affected by numerous processing and materials variables. The texture is usually related directly to the nucleation event during recrystallization, the orientation dependence of the rate of nucleation in inhomogeneities of various type and
21
Figure 2.15. Typical FCC deformation texture of 90% cold rolled Al-4.5% Cu (Sebald and Gottstein 2002).
orientation environment, and to the nature, energy and mobility of the boundaries between grains of various orientations.
Discontinuous recrystallization of low-solute Al alloys following a cold rolling often generates a strong cube texture. Nevertheless, the presence of Mg or other elements in solution decreases recovery and generates shear bands during cold rolling (Morii et al., 1985) which collectively decreases the domination of the cube texture orientation due to: (i) suppression of cube texture due to SBs cutting through cube- oriented grains (Ridha and Hutchinson, 1982), (ii) SBs promoting nucleation of grains with texture components such as {011}<100>, {011}<122> and {013}<231>
(Hirsch, 1990; Engler, 2001). These processes tend to generate a more randomized texture in the recrystallized sheet.
2.5.5 Texture of ARB-processed aluminium alloys
In general, the texture development in ARB-processed metals has not been studied extensively compared with the more extensive studies of general microstructural evolution and associated mechanical properties. The available studies of texture evolution in Al by ARB are given as follows: high purity Al (Heason and Prangnell,
2002; Lee et al., 2002b; Kamikawa et al., 2006; Pirgazi and Akbarzadeh, 2009),
AA8011 (Kim et al., 2002, 2005a), Al-Li (Chowdhury et al., 2006a), AA3XXX
(Chowdhury et al., 2006b; Pirgazi et al., 2008b). Texture investigations have also been carried out on Mg/Al ARB composites (Chang et al., 2009).
The early work on commercial purity Al has shown, by analysing {111} pole figures, that the texture developed after ARB is asymmetric and very weak (Saito et al.,
1998a). Furthermore, for the same alloy, Heason and Prangnell (2002) revealed that most of the shear texture is rotated Cu and S textures when incorporated into the centre during ARB process. Another study by Kim et al. (2005a) observed the
22 presence of ideal Cu and Dillamore {4,4,11}<11,11,8> orientations at the centre of the strip after ARB, demonstrating that the texture of ARB-processed sheet has more complex features than that of conventionally rolled sheets. Figure 2.16 shows the texture of ARB-processed AA8011 Al sheet. The strain mode and texture are different in the ARB process with a through-thickness location, especially in case of a high reduction rate per pass and high friction conditions between the roll and rolled sheets ( Kim et al., 2002; Lee et al., 2002b; Kim et al., 2005a).
The basic concept associated with texture evolution during ARB is that the shear texture at the subsurface continuously changes to a standard rolling texture at the sheet centre as a result of additional rolling cycles and this pattern progresses until the final cycle. These textural changes are a result of the continual crystal rotations induced by the plane-strain deformation at the sheet centre. Therefore, most areas throughout the sheet thickness eventually consist of the classic rolling texture, except for those regions very close to the sheet surface (Saito et al., 1998a; Huang et al.,
2003; Kamikawa et al., 2007a).
2.6 Mechanical Behaviour of ARB-Processed Metals
During SPD, the microstructure is considerably refined to generate ultra fine grains and various types of microstructural heterogeneities (see e.g. section 2.4). Figure
2.17 shows the microhardness as a function of accumulated strain for a range of SPD techniques carried out on aluminium (Pi et al., 2006). It can be seen that considerable hardening occurs with hardness increasing in a similar manner regardless of the SPD route. The UFG structures generated by SPD generally also result in a considerable increase in yield strength of a given material that is several times greater than their coarse-grained counterparts (Tsuji et al., 2002a; Zhu et al.,
2003; Huang et al., 2006; Song et al., 2006). However, SPD also drastically reduces the uniform elongation or ductility of the material to values less than 3% (Zhu et al.,
23
Figure 2.16. 111 pole figures and 2 = 45-, 65-, and 90-degree sections of Euler space of the eight-cycle ARB-processed AA8011 Al sheet: (a) subsurface (s = 0.9), (b) quarter thickness (s = 0.5), and (c) centre (s = 0.1). In the ODFs, contour levels are drawn at, 1, 3, 5….25x random (Kim et al., 2005a).
Figure 2.17. Microhardness as a function of accumulated strain for a range of SPD routes on aluminium (Pi et al., 2006). MAC/F denotes multi-axial compression/forging process and CR denotes cold rolling. 2003; Huang et al., 2008c). Figure 2.18 shows nominal stress-strain curves for ARB- produced AA1100 Al showing the significant influence of the number of rolling cycles on flow stress and ductility (Kwan et al., 2008).
The low uniform elongation during tensile straining of ARB-processed materials is often correlated to an increase in strain rate sensitivity (Wei et al., 2004). Such limited uniform elongation is a critical factor for the practical application of UFG/NS materials and therefore needs to be controlled. It is argued that if the rate of strain- hardening can be enhanced during tensile straining, uniform elongation can also be improved. One possible way to increase the rate of strain hardening is via the introduction of fine dispersoids within the UFG microstructure such that a large number of dislocations can be accumulated. In a recent study on an UFG Cu-Cr alloy produced by ARB (Takata et al., 2009), it was shown that strain hardening is considerably enhanced by the addition of finely dispersed second phase particles.
The effect of dislocation density of ARB-processed nanostructured (NS) metals on mechanical behavior has been explored in several recent studies (Huang et al., 2006;
Kamikawa et al., 2007; Huang et al., 2008a,b). It was shown that the presence of a low density of mobile dislocations can affect the hardening response. Here, the relationship between structural parameters and mechanical properties revealed that the dislocation structure associated with LAGBs and the internal dislocations substantially influenced the mechanical properties of the deformed ARB samples.
This is emphasized by the finding that dislocations in LAGBs and within grains contribute to strengthening in addition to the expected Hall-Petch contribution
(Hansen, 2004; Nes et al., 2006).
A somewhat surprising phenomena has recently been reported in annealed UFG/NS metals produced by ARB (Tsuji et al., 2002a; Bowen et al., 2004; Yu et al., 2005;
Huang et al., 2006; Kamikawa et al., 2008). Here, an increase in yield strength
24
Figure 2.18. Nominal stress–strain curves of AA1100 Al after various cycles of ARB: (a) as-rolled condition, (b) after annealing at 250 °C (Kwan et al., 2008). (hardening) occurred after low-temperature annealing in ARB-processed high-purity
Al and interstitial-free steel. These materials then demonstrated a reverse scenario of yield strength softening if the post-ARB annealing treatments was followed by a further small deformation, e.g. 15% cold rolling (Huang et al., 2006; Huang et al.,
2008d; Kamikawa et al., 2008), Such behaviour has been attributed to the fine-scale structure produced by ARB and to the structural alteration of dislocations during subsequent annealing. In very recent work (Huang, 2009) on NS metals (i.e.
AA1100) produced by six-cycle ARB, it was confirmed that the mechanical behaviour was profoundly reliant on the dislocation substructure. Here, a new approach for mechanical properties optimization was highlighted in terms of tailoring the dislocation structure by post-processing treatments (annealing and deformation), which in turn hampers the detrimental plastic instability effect resulting from the lack of mobile dislocations.
In concurrence with this hot topic of mechanical property optimization of nanostructured metals, another extensive study by Kamikawa et al. (2009) explored the structural modification approach by post-ARB annealing that activates a
“dislocation source hardening” mechanism. They developed a two-step annealing route following six-cycle ARB of high-purity Al and achieved a flow stress of 60-70
MPa and a uniform elongation of 20–30% (Fig. 2.19). These workers also discussed their observed Hall–Petch slope deviation and the occurrence of localized shear deformation along the shear bands causing a yield drop and/or premature failure
Indeed, several studies on ARB-processed materials have addressed the conventional
(well-established) strengthening mechanisms and their compliance with the Hall-
Petch relationship (Eizadjou et al., 2008; Pirgazi et al., 2008a; Eizadjou et al., 2009;
Hosseini and Manesh, 2009). Figure 2.20 shows typical results that supports this type of relation in ARB-processed commercial purity aluminium.
25
Figure 2.19. Nominal stress–strain curves of ARB-processed-and-annealed samples of high purity Al. The dashed curve denotes the starting material (Kamikawa et al., 2009).
Figure 2.20. Hall–Petch plot showing flow stress and hardness of ARB-processed commercial purity aluminium (Eizadjou et al., 2009).
Chapter 3 ______LITERATURE REVIEW–PART II Static Restoration Processes Pertinent to Severe Plastic Deformation
3.1 Introduction
Most cold rolled sheet products are usually annealed which results in an increase in ductility with a concomitant decrease in strength (Humphreys and Hatherly, 2004).
The principal material factors that influence static annealing include composition, initial grain size and texture, second phase precipitates and solute distribution. Some important processing variables include deformation mode, strain, strain rate and deformation temperature. The driving pressure for softening by both recovery and recrystallization (either dynamic or static) is the stored energy of deformation.
Recovery reduces the stored energy of the material and, hence, the driving force for recrystallization, by mechanisms such as the annihilation of dislocations and their re- arrangement into low angle grain boundaries. For aluminium alloys, some of the more notable parameters affecting static annealing include initial microstructure and
26 texture, chemical composition, degree and mode of deformation, annealing temperature and heating rate (Humphreys and Hatherly 2004).
The high stored energy of the deformation microstructure renders it thermodynamically unstable and heating to a high temperature (annealing), allows thermally-activated processes such as dislocation climb and cross-slip to occur. This provides the mechanism for the annihilation and rearrangement of these defects to generate a lower energy state. During annealing, the deformation microstructure may experience up to three classic stages of structural alteration: recovery, recrystallization and grain growth (Humphreys and Hatherly, 2004).
3.2 Static Restoration Processes
3.2.1 Recovery
Structural changes taking place during recovery of deformed metals have been a long lasting topic of interest. Most of this research has been focused on metals deformed to low and medium strains (Humphreys and Hatherly, 2004). However, there is renewed interest on the recovery behaviour of highly strained metals (Iwahashi et al.,
1997; Saito et al., 1999; Liu et al., 2002; Tsuji et al., 2002c). A major reason is that the thermal stability decreases as the microstructure is refined and another is the decrease in ductility after large strains whereby low temperature annealing may be required to provide a suitable balance between strength and ductility.
Recovery generally involves only partial restoration of properties prior to recrystallization because the dislocation structure cannot be removed, but reaches a metastable state during annealing. During recovery, dislocations are annihilated and rearranged to reduce the stored energy of the matrix, Fig. 3.1. The original grain boundaries are not affected although significant subgrain growth will occur. The driving force for recovery is the reduction in stored energy due to a decrease in both
27
Figure 3.1. Various stages of recovery in a plastically deformed metal (Humphreys and Hatherly, 2004). vacancies and dislocations. Both recovery and recrystallization are competing processes as they have the same driving forces, as shown by the calorimetric study of both processes in Fig. 3.2. The rate of recovery is affected by annealing temperature with high temperatures resulting in a higher rate of recovery. The rate of recovery is also reduced by solute and fine particles by retarding the movement of dislocations and grain boundaries. If sufficient recovery is allowed to occur, there will be a lower driving force for recrystallization which may result in a slower rate of recrystallization (Beck 1954). Throughout recovery, the strength is slightly decreased but the ductility may be increased significantly (Humphreys and Hatherly, 2004).
Several physical properties, altered during plastic deformation, are also subsequently modified by recovery such as electrical resistivity and density.
3.2.2 Static recrystallization
Recrystallization may occur during the recovery stage but involves an incubation period prior to the nucleation and growth of so-called ‘strain-free’ grains. The nucleation process is characterized by high angle grain boundary (HAGB) formation to generate relatively strain-free grains that subsequently grow and consume the recovering substructure. When compared with recovery, recrystallization results in a more significant increase in ductility and a concomitant decrease in strength.
3.2.2.1 Nucleation of recrystallization
Nucleation corresponds to the formation of a grain of critical size containing a very low dislocation density within the surrounding deformation microstructure. Such a grain has a low internal energy and is separated from the deformation substructure by a high angle grain boundary. The process of nucleation originates from the cells or subgrains generated by deformation, and due to the nuclei being highly misoriented with respect to their neighbouring cells or subgrains. In order for a nucleus to grow
28
Figure 3.2. The recovery of 99.999% Al deformed to a true strain of 6.91 at 350 C, as measured by differential scanning calorimetry (DSC) (Schmidt and Haessner, 1990). into the deformation microstructure, it must be greater than a critical radius, Rc , which is related to the grain boundary energy and stored energy E:
4E R (3.1) c
It is important to note that the distribution of nuclei during recrystallization is not random with nucleation occurring at prior grain boundaries, transition bands, deformation bands, shear bands and in the vicinity of coarse inclusions/particles
(Humphreys and Hatherly, 2004). Nucleation also commences faster in high stored energy texture components of the deformation microstructure.
3.2.2.2 Growth of nuclei
If nuclei are formed at sites such as grain boundaries, shear bands, deformation bands etc., they may continue to grow to consume the deformation microstructure thereby further lowering the stored energy of the system. The growth may occur while new nuclei are forming with the entire process continuing until the growing grains impinge to completely consume the deformation microstructure. Static recrystallization commences with a nucleation stage with further growth of these nuclei consuming the remaining deformation substructure to eventually produce a fully recrystallized microstructure. An understanding of the nucleation and growth stages of recrystallization is extremely important as they affect the kinetics of the process, the final grain size and distribution and produce a recrystallization texture which has a large influence on sheet formability (Humphreys and Hatherly, 2004).
3.3 C lassification of Static R estoration Processes
The foregoing discussion was concerned with the classic restoration processes in a deformed metal: recovery and recrystallization. Table 3.1 shows a classification
29 scheme whereby restoration processes are defined as either continuous or discontinuous, as shown schematically in Fig. 3.3. The former consists of recovery, recrystallization and normal grain growth where the microstructure coarsens gradually with no major nucleation event. The latter consists of discontinuous
(abnormal) subgrain growth, recrystallization and abnormal grain growth and involves nucleation and preferential growth of certain grains throughout the microstructure.
Table 3.1. Summary of static annealing phenomena, after Humphreys (1999).
Restoration Recovery Recrystallization Grain Growth Process Type
Continuous Normal grain Continuous Subgrain growth recrystallization growth
Discontinuous Primary Abnormal grain Discontinuous subgrain growth recrystallization growth
In Fig. 3.4, the transition from a fine to coarse cellular structure (region A), is a continuous process i.e. either subgrain growth or normal grain growth, which is boundary misorientation dependent. In contrast, the discontinuous transformation
(region B) can be described as either primary recrystallization (discontinuous recrystallization) or abnormal grain growth. The processes shown in Table 3.1 occur by the migration of either HAGBs or LAGBs, the differences between them being characterized by the spatial distribution and properties of these boundaries
(Humphreys, 1999).
3.3.1 Continuous and discontinuous recrystallization
Oscarsson et al. (1992) and Jazaeri and Humphreys (2004b) called the homogeneous evolution of a substructure during annealing as continuous recrystallization. An example of continuous recrystallization is shown in Fig. 3.5 for heavily rolled Al-Mg
30
(a) (b)
(c) (d)
(e) (f) Figure 3.3. Schematic diagram of the main annealing processes: (a) deformed state, (b) recovered, (c) partially recrystallized, (d) fully recrystallized, (e) grain growth and (f) abnormal grain growth, after Humphreys and Hatherly (2004).
Figure 3.4. Schematic illustration of continuous and discontinuous annealing phenomena (Humphreys, 1999).
.
Figure 3.5. Series of EBSD maps showing continuous recrystallization in fine-grained Al-0.1Mg alloy deformed to a true strain of 2.6. (a) As-deformed, and annealed for 1h at (b) 200 C, (c) 225 C, (d) 250 C and (e) 300 C (Jazaeri and Humphreys, 2004b).
alloy showing a predominantly HAGB microstructure developing at relatively low annealing temperatures and gradual grain coarsening by minor boundary movements.
There is no significant change in the HAGB fraction, with the deformed grains simply becoming more equiaxed and slightly larger. Generally, continuous recrystallization is mainly affected by process variables like the degree of plastic deformation, strain rate, pre-ageing, annealing time and temperature, processing route and the amount of prior deformation or initial grain size (Humphreys and
Hatherly, 2004).
3.3.1.1 Continuous recrystallization in highly strained alloys
The significant effect of strain and starting microstructure on continuous recrystallization was highlighted by Humphreys et al. (1999), Gholinia et al., (2001) and Prangnell et al., (2004) on Al-Mg and other Al-base alloys. Jazaeri and
Humphreys (2004a,b) studied a range of commercial grade Al alloys and showed there was a distinct transition from discontinuous to continuous recrystallization that was both strain and microstructure dependent. For discontinuous recrystallization, there was a relatively sharp increase in the HAGB fraction as the deformation substructure is consumed by recrystallizing grains. However, continuous recrystallization results in little change in HAGB fraction. Figure 3.6 shows the influence of both strain and annealing temperature on the evolution of HAGB fraction for a range of commercial Al alloys. It is a requirement that such highly strained microstructures contain greater than 60-70% HAGBs for ensuring microstructural stability (Humphreys, 1997; Humphreys et al., 1999).
3.3.1.2 Transition from continuous to discontinuous recrystallization
Humphreys (1997a,b) showed that fine-grain microstructures (evolving by continuous recrystallization) are intrinsically unstable at high annealing temperatures. Several experimental studies have been carried out on the annealing response of heavily
31
Figure 3.6. The changes in the boundary character (HAGB%) during annealing of the medium grain size Al alloys after deformation at true strains up to 3.9: (a) Al-0.1Mg, (b) AA1200 and (c) AA8006 (Jazaeri and Humphreys, 2004b).
. deformed Al alloys that do not contain a significant amount of fine particles i.e. Al-
Mg (Furukawa et al., 1996) and (Morris and Munoz-Morris, 2002), AA3004 and AA
1100 (Horita et al., 2001a,b) , and AA 1050 (Cao et al., 2003) and (Yu et al., 2004).
Despite the large fraction of HAGBs in these alloys, they were prone to normal continuous recrystallization with discontinuous recrystallization commencing at temperatures below ~350 C to eventually produce a coarse (>10μm) grain size.
Ferry et al. (2005) investigated the static restoration of ECAP-deformed Al-0.2Sc alloy containing fine, coherent Al3Sc dispersoids and demonstrated that annealing at temperatures up to 500 C resulted in continuous recrystallization, whereby the initial microstructure evolves gradually with no marked change in the grain size distribution, texture and HAGB fraction, Fig. 3.7. However, extended annealing at
500 C or shorter times at higher temperatures resulted in discontinuous recrystallization (Fig. 3.8). For particle-containing Al-Sc alloys, the annealing temperature dominates the criterion for a transition from continuous to discontinuous recrystallization (Hamilton and Ferry, 2004; Ferry et al., 2005). The influence of fine particles on grain boundary pinning is outlined in section 3.4.
3.4 The Effect of Second-Phase Particles on Static Restoration
3.4.1 Principles of Zener pinning
Zener proposed (Smith, 1948) that a drag force exerted on a grain boundary by particles would counter the driving force for growth from its curvature, that is:
dR PPM (3.2) dt z
The consequence of this proposal was, provided its descriptive parameters remained the same and regardless of the boundary mobility (M), a particle dispersion could completely counter normal grain growth when the driving force for such growth (P)
32
Figure 3.7. Bright field TEM micrographs showing the continuous evolution of microstructure in ECAP-deformed Al-0.2Sc alloy during annealing at 350 °C, after Ferry et al. (2005).
Figure 3.8. ICC micrograph of Al-0.2%Sc alloy pre-aged for 3h at 350 C then annealed for 2h at 500 C showing the onset of discontinuous recrystallization (arrows), after Ferry et al. (2005).
does not exceed the pinning force by the particles (Pz). At this point, the growth rate is zero, i.e. dR/dt = 0, and the grains have reached a critical radius, Rc for a given particle dispersion.
A full derivation of the Zener equation (Manohar et al., 1998) produces, in its general form, the critical grain size (Rc) expressed as a function of average particle radius (r),
Constants (K and m) and particle volume fraction (f):
r R K (3.3) c z f m
The main contention between the many studies carried out over the years is not the general form of the equation, but the nature of the constants Kz and m. The derivation by Manohar et al. (1998) showed that the constants K and m have likely values of 4/3 and 1 respectively, although possibly 2/3 and 1. Figure 3.9 is a summary of the research carried out on Al- and Fe-based systems containing a variety of fine particles of various volume fractions. From this figure, it appears that the pinning nature of the particles is independent of system and dispersoids, but there is a linearity change for f > 0.05. While the reason for this is not fully established, it may be attributed to the non-random correlation of boundaries with particles becoming significant.
3.4.2 Grain growth kinetics in the presence of a particle dispersion
Through the assumption that grain boundary velocity is inversely proportional to its radius of curvature, Hillert (1962) incorporated the principle of Oswald ripening via the distribution of second phase particles to obtain the following equation describing the growth of grains of radius R:
33
\
Figure 3.9. Plot of critical grain radius as a function of volume fraction of fine particles for a wide range of alloys, after Manohar et al. (1998). dR F 11 V cM G W (3.4) dt H cr RR X where c is a small constant, M the boundary mobility, the boundary surface energy and Rcr is a critical radius that varies according to:
dR cM cr (3.5) dt Rcr
Hillert extended this normal grain growth theory to include the principles of particle pinning, and has been subsequently documented in more recent works such as that by
Humphreys and Hatherly (2004). This expresses the growth rate for grains of radius
R in a particle-dispersed system where the pinning pressure (Pz = z) is that offered by the particle dispersion to grain boundary movement:
dR F 11 V cM G cz W (3.6) dt H cr RR X
Eq. (3.6) can be rewritten as follows:
dR F 3 f V M G b b W (3.7) dt H R 2r X
Hence, particles have a significant retarding influence on the kinetics of grain growth
(and recrystallization) with growth offset by both particle size and their volume fraction.
3.5 Physical Metallurgy of Al-Sc Alloys
This thesis is concerned with the production of multilayered lamellar composites where the alternating layers are commercial purity Al and an Al-0.3 wt.% Sc alloy.
The influence of scandium in aluminium therefore needs to be addressed in detail
34 and this section provides an overview of the physical metallurgy of Al-Sc alloys. A more complete review of this alloy system is given by Royset and Ryum (2005).
3.5.1 Attributes of scandium (Sc) in aluminium
The influence of rare earth elements and transition metals in various alloy systems on microstructure and properties has been studied extensively in the last few decades. It has been shown that the mechanical properties, electrical conductivity and corrosion resistance etc. are often improved in these alloys. Scandium is a rare-earth element and one of the most potent alloying elements in aluminium and magnesium. It is a
- 3), with a melting point of 1541 °C. When added to
Al alloys, scandium can significantly increase strength and reduce grain size as the dispersoids that form inhibit recrystallization (Royset and Ryum, 2005). The main effect of scandium in conventional alloys are improved strength, improved resistance to recrystallization and reduced grain size, and improved resistance to hot cracking
(Toropova et al., 1998; Fuller et al., 1999; Kendig and Miracle, 2002; Marquis et al.,
2003; Royset and Riddle, 2004; Royset and Ryum, 2005).
Figures 3.10 and 3.11 show the complete and Al-rich section, respectively, of the Al-
Sc phase diagram. It can be seen that four intermetallic phases are possible: Al3Sc,
Al2Sc, AlSc and AlSc2. It is also evident from Fig. 3.11 that scandium shows partial solid solubility in aluminium with a maximum solubility at the eutectic temperature of ~ 0.35 wt.% Sc (Fujikawa et al. 1979). The reported eutectic temperature is ~ 660
°C and the following reaction occurs, i.e.: L 3Sc.
3.5.2 Formation of Al3Sc during solidification
The crystal structure of the Al3Sc phase is ordered FCC, although, in crystallographic terminology, it is a primitive cubic lattice with four atoms (one Sc and three Al atoms) associated with each lattice point. The particular atomic arrangement is L12,
35
Figure 3.10. The A1-Sc binary phase diagram, after Murray (1998).
Figure 3.11. Enlargement of the Al-rich region of the A1-Sc phase diagram illustrating that the melting point of A1 is reduced by scandium (Murray, 1998)). as found in several other intermetallic compounds such as Ni3Al, Ni3Fe, and Cu3Au.
There are at least four ways that the Al3Sc phase can form in dilute binary Al-Sc alloys:
1. Upon solidification of hypereutectic alloys (Sc > ~ 0.6 wt.%), Al3Sc is the first
phase to form.
2. The solidification of hypo- and hyper-eutectic alloys ends with the formation of
eutectic Al + Al3Sc.
3. Al3Sc can precipitate discontinuously from a supersaturated solid solution.
4. Al3Sc can precipitate continuously (nucleation and growth) from a supersaturated
solid solution.
Following solidification, primary Al3Sc is reported to have a highly faceted morphology in a slowly cooled Al-8.7 wt.% Sc alloy indicating that the facet planes are {100} planes (Blake and Hopkins, 1985). Brodova et al. (1992a,b) investigated the microstructure of a rapidly cooled Al-2 wt.% Sc alloy and found that the dominant morphology of primary Al3Sc was dependent both on the cooling rate and the initial melt temperature. The nucleation mechanism and growth morphology of primary Al3Sc in an Al-0.7 wt.% Sc alloy solidified at a cooling rate of 100 K/s were investigated by Hyde et al. (2001). It was found that the primary Al3Sc particles grow with dendritic morphology, with the primary dendrite arms in the crystallographic <100> directions.
3.5.3 Precipitation of Al3Sc from supersaturated solid solution
Precipitation of Al3Sc from supersaturated solid solution is characterized by nucleation, growth and coarsening stages. It is argued that the reaction occurs largely by homogeneous nucleation and diffusion-controlled growth and that Al3Sc is the only phase to form, that is, the equilibrium phase nucleates directly from the
36 supersaturated solid solution. Even though the nucleation of Al3Sc is frequently reported to occur homogeneously throughout the Al matrix, there are some reports of heterogeneous Al3Sc nucleation on dislocations (Tan et al., 1992; Jones and
Humphreys, 2003) and grain boundaries (Nakayama et al. 1997; Novotny and Ardell,
2001). The shape of continuously precipitated Al3Sc particles is most frequently reported to be spherical. In a systematic TEM study of Al3Sc particle morphology during precipitation of binary Al-Sc alloys, Marquis and Seidman (2001) demonstrated that the precipitates can take a wide range of shapes as they grow depending on the Sc content in the alloy, transformation temperature and transformation time. During the growth of a coherent Al3Sc particle (Fig. 3.12), it eventually reaches a critical size where it will be energetically more favourable to introduce an interface dislocation at the Al/Al3Sc interface than to increase the matrix strain (Fig. 3.13), with a critical radius of ~ 20 nm. Drits et al. (1984) showed that the coarsening rate of the Al3Sc particles increases when they lose coherency.
3.5.4 Effect of Al3Sc on deformation and annealing
Most of the effects of scandium in wrought Al alloys are linked to the formation of the Al3Sc phase. In a typical processing route, Al3Sc particles can form under three different conditions:
1. During solidification after casting or welding, Al3Sc particles can form in the melt
-Al, thus leading to grain refinement.
2. High temperature processing of the alloy in the range 400-600 °C, for instance
homogenization, hot-rolling or extrusion, can generate a dense distribution of
Al3Sc particles of typically 20-100 nm in diameter. The particle distributions
formed under such conditions are reported to lead to good recrystallization
resistance and enhanced superplasticity.
37
Figure 3.12. Bright field TEM micrographs of Al3Sc particles for different aging conditions: (a) 400 C, 52 min, (b) 430 C, 85 min, (c) 460 C, 1.8 min (Iwamura and Miura, 2004).
Figure 3.13. Weak beam TEM micrograph of the dislocation network surrounding an Al3Sc dispersoid (Iwamura and Miura, 2004).
3. Heat treatment in the range 250-350 °C can lead to significant precipitation
hardening of an alloy supersaturated in scandium. The size of strengthening Al3Sc
precipitates is typically in the range 2-6 nm in diameter.
3.5.4.1 Deformation microstructure
In a recent study on ECAP-deformed Al-Sc alloys (Apps et al., 2003, 2005) it was found that Al3Sc in the form of fine, non-shearable, dispersoids tends to homogenise slip, retard the formation of a cellular substructure and inhibit the formation of micro-shear bands during deformation. Collectively, these dispersoids reduce the rate of HAGB generation at low-to-medium strains and, hence, retard the formation of an
UFG structure to higher plastic strains. However, if the Al3Sc particles are coarse, incoherent and widely-spaced, there is little impedance to dislocation rearrangement and a well-defined cellular substructure is generated after straining (Ferry et al.,
2005).
3.5.4.2 Recrystallization behaviour
An important effect of a dense distribution of ultrafine Al3Sc dispersoids is their ability to stabilize the grain/subgrain structure of an alloy through Zener pinning, thereby improving the mechanical properties of the alloy. There are two noticeable dispersoid-related effects that have been reported for Sc-containing alloys. The first is the stabilizing effect of the deformation microstructure and a non-recrystallized structure is preserved in an alloy therefore adding considerably to its strength. A second major effect is the preservation of a small grain size during hot deformation, which is necessary for superplastic processing of an alloy.
Many researchers have reported that Al-Sc alloys are stable against recrystallization to near the melting point of the alloy due to the very high density of coherent, nanosized Al3Sc dispersoids generated in the microstructure either before or during
38 annealing (Nakayama et al., 1997; Toropova et al., 1998; Jones and Humphreys,
2003; Lee et al., 2002c; Seidman et al., 2002; Ferry et al., 2005; Royset and Ryum,
2005; Smola et al., 2007). Royset et al. (2004) reported, for AA6082 and AA7108 alloys, that Sc and Zr in combination leads to particularly improved recrystallization resistance. Figure 3.14 shows the fraction recrystallized in cold rolled sheets of Al-
0.4%Sc and Al-0.4 %Sc-0.15%Zr as a function of annealing temperature showing considerable retardation of recrystallization (Royset and Ryum, 2005). Several studies on more complex Al-Sc-Hf-Zr and Al-Fe-Si-(Sc)-(Hf)-(Zr) alloys have shown that a high density of fine and homogeneously distributed dispersoids results in considerable coarsening resistance of the deformation substructure at temperatures up to 550-600 C (Forbord et al., 2004; Hallem et al., 2006; Hallem et al., 2007).
3.5.5 Influence of Al3Sc dispersoids on mechanical properties
The addition of small amounts of scandium to aluminium (up to Ң0.5%) has a beneficial effect on mechanical properties, weldability and corrosion resistance.
Many researchers have investigated the contribution of Al3Sc on mechanical properties of a range of Al alloys (Torma et al., 1989; Parker et al., 1995; Komura et al., 2000; Fuller et al., 2002; Marquis et al., 2003; Royset and Ryum, 2005). In a recent study (Seidman et al., 2002), the room-temperature flow stress of Al was shown to be increased significantly (from 20 to 140-200 MPa) with the generation of
Al3Sc dispersoids. It was shown that peak strengthening corresponded to a precipitate radius of 1.5 nm (Fig. 3.15). Based on classical dispersion-strengthening theories, strength was argued to be controlled by the precipitate shearing mechanism for small particle sizes and the Orowan dislocation bypass mechanism for larger sizes.
39
Figure 3.14. Fraction recrystallised in cold rolled sheets of Al-0.4Sc and Al-0.4S-0.15Zr as a function of annealing temperature (Royset and Ryum, 2005)
Figure 3.15. Increase in yield stress as a function of precipitate radius for an
Al-0.3% Sc alloy aged at different temperatures. Experimental points are obtained from uni-axial yield stress measurements and microhardness measurements, and the theoretical lines are calculated for the Orowan stress (
), the cutting stress due to the APB energy (Δ ), and the cutting stress due to Or 1 the lattice and modulus mismatch (Δ +Δ ) (Seidman et al., 2002). 2 3
3.6 Summary and Scope of Thesis
In the investigations on severe plastic deformation (SPD) of alloys, a method that is becoming increasingly important is accumulative roll bonding (ARB). This process involves the simultaneous rolling of two stacked sheets of comparable thickness. In the first rolling pass, the thickness of both sheets is reduced to the initial sheet thickness, and the procedure is repeated several times to produce final sheet material of similar thickness to the starting material. The repeated rolling procedure generates very large plastic strains resulting in substantial microstructural refinement.
Therefore, ARB has the ability to produce an ultrafine grained microstructure without any substantial change in sheet dimensions. Such SMG/UFG grain size alloys may also exhibit a combination of useful properties including high strength and hardness, and excellent superplastic formability. ARB has a notable advantage over other SPD methods as the former can easily be scaled up for direct industrial applications in the continuous production of UFG sheet of many types of alloy. An additional advantage of ARB is the ability to produce multilayered sheet composite materials containing alternating layers of two or more dissimilar metals. This type of
ARB route enables the creation of sheet materials where the overall properties are based on the combined properties of the individual alloys and may therefore provide scope for extending into new regions of property space.
The principal objective of the thesis is to produce multilayered aluminium alloy composites by ARB comprising of alternating layers of commercial purity Al and an
Al-0.3% Sc alloy. These alloys were selected due to the considerable knowledge generated over recent years on SPD of Al and Al(Sc) alloys. However, these alloys have never been combined in a way that may exploit the useful properties of the individual alloys. For example, commercial purity Al alloys readily generate an
UFG microstructure by ARB although its thermal stability is low and
40 recrystallization occurs readily during subsequent annealing at relatively low temperatures. This effect is due to the very high stored energy associated with large surface area of grain boundaries within the given volume of material. Al(Sc) alloys can also be deformed to large strains by ARB (and ECAP) but suitable processing strategies can produce a dispersion of nanosized Al3Sc particles throughout the deformation microstructure. High thermal stability of the deformation microstructure can therefore be achieved as these particles are also thermally stable and, therefore, inhibit grain boundary migration associated with recrystallization and grain growth by the process of particle (Zener) pinning.
This thesis describes a systematic study of the effect of initial microstructure and processing conditions on the development of multilayered Al/Al(Sc) composites by exploiting the strong influence of ultrafine particles on inhibiting recrystallization. It is expected that these multilayered materials will be produced containing a microstructure and texture specific to the individual layers, thereby generating composite structures that, potentially, have a unique combination of properties. The multilayered composite materials will be produced by ARB in conjunction with post- deformation annealing treatments. Chapter 4 outlines the experimental procedures concerning the primary production of the alloys, ARB and annealing experiments and methods of examining the microstructures and textures using ion channelling contrast (ICC) imaging, electron back scattered diffraction (EBSD) and transmission electron microscopy (TEM). The results are divided into three chapters, depending on both the initial microstructure of the Al(Sc) alloy and processing variables.
Chapter 5 describes the production and properties of Al/Al(Sc) composites after multi-pass ARB and subsequent annealing where the Al(Sc) alloy is heat treated to generate a supersaturated solid solution prior to roll bonding. Chapter 6 explores the influence of ultrafine Al3Sc particles in the Al(Sc) sheet prior to ARB on
41 microstructural development during both roll bonding and subsequent annealing.
The materials in chapters 5 and 6 will be roll bonded at 200 °C to a large strain (five rolling cycles to produce 32 alternating layers).
Chapter 7 explores the roll bonding behaviour and microstructural development of
Al/Al(Sc) where the alloys are in the same condition as that produced in chapter 5, but where rolling is carried out at a higher temperature (350 °C) and to a larger strain
(six rolling cycles to produce 64 alternating layers). Chapter 8 provides a detailed discussion of the general evolution of the microstructure and texture of all materials produced after ARB and during subsequent annealing with particular emphasis on the fundamental nature of structural evolution during each processing stage, the propensity for shear band formation during ARB, and the general influence of particle pinning on the generation of these unique multilayered microstructures.
Finally, chapter 9 provides a concluding summary of the major findings in this thesis.
42 Chapter 4 ______Experimental Procedure
4.1 Introduction
This chapter describes the experimental procedures for fabricating and analysing the various roll bonded materials produced in this thesis. It includes an outline of the initial processing and heat treatments of the nominated materials, the accumulative roll bonding (ARB) process and associated post-ARB annealing treatments, and sample preparation methodologies and procedures for carrying out both mechanical testing and microstructural analysis.
4.2 Materials Processing
4.2.1 Preliminary processing of candidate materials
The starting alloys used for generating multi-layered composite materials by ARB include commercial purity (99.8 wt.%) aluminium, hereafter termed Al, and an Al-
0.3wt.%Sc alloy, hereafter termed Al(Sc). The various processing routes used in this
43 thesis to produce the composite materials are given in Fig. 4.1. The initial processing of Al was carried out by cold rolling an as-cast slab of the material from a thickness of 13 mm to 1 mm in 10% incremental reductions. The as-rolled material was annealed for 45min at 375 °C to full recrystallization using a preheated air circulating muffle furnace followed by air cooling to room temperature. This procedure generated an equiaxed microstructure of grain size ~50 ± 30 m.
For producing Al(Sc), high purity Al (99.99%) was melted in conjunction with a Al-
33%Sc master alloy to generate 10 mm thick as-cast slabs of nominal composition
Al-0.3%Sc. The as-cast material was homogenized for 48h at 640 °C in an air circulating furnace. The heat treated alloy was cold rolled in multiple passes to a thickness of 1 mm then further solution treated for 48h at 640 °C followed by cold water quenching. This procedure generated a coarse-grained microstructure across the thickness of the Al(Sc) sheet. It also generated Al(Sc) in the supersaturated solid solution (SSSS) condition. Samples of the same alloy were artificially aged for 4h at
300 C to produce a microstructure containing a uniform dispersion of nanosized
Al3Sc dispersoids: this additional processing stage generated sheet material in the artificially aged condition. The commercial purity Al and both SSSS- and Aged-
Al(Sc) were stacked and processed by ARB to generate three different composite materials consisting of alternating layers of each alloy type (Fig. 4.1).
4.2.2 Accumulative roll bonding (ARB)
Figure 4.1 shows a flow diagram outlining the various processing stages for generating the three different multilayered composite materials. For ARB, the Al and Al(Sc) alloys were cut into 1×50×100 mm sheets then brushed and cleaned prior to stacking and subsequent roll bonding. The brushing stage removes both oxide and dirt and consisted of grinding with SiC paper followed by brushing using a steel and
44 Figure 4.1. Flow diagram showing the processing routes for producing the three types of accumulative roll bonded (ARB) samples used in this thesis. stainless-steel brush. All samples were cleaned in acetone prior to further processing.
Figure 4.2 is a schematic diagram showing the stages involved in ARB. Here, a sheet of both Al and the two types of Al(Sc) alloy were stacked and clamped by three steel-wires located on pre-drilled holes. Roll bonding was carried out at a temperature of either 200 C or 350 C (Fig. 4.1); this produced three types of composite material hereafter termed SSSS-ARB, SSSS-ARB-HT and Aged-ARB.
The lower temperature rolling was used for exploring the scientific fundamentals of deformation and annealing whereas the higher rolling temperature was used for exploring the feasibility of scaling up the process by lowering the rolling loads and improving the bonding between layers.
For ARB, the clamped sheets were held for ~5 min at the predetermined temperature in a preheated muffle furnace then rolled at the same temperature without lubrication to 50% reduction in a single pass using a two-high laboratory rolling mill. The as- rolled materials were cut into two similar sized pieces, cleaned and brushed and further roll bonded. The overall roll bonding conditions are summarised in Table
4.1.
Table 4.1. Summary of materials and ARB processing parameters. Average Product Condition Total true Rolling Number of thickness code of Al(Sc) Rolling strain of temperature alternating of a name prior to cycles processed (°C) layers given used in ARB material layer thesis SSSS 200 5 32 ~4.1 ~15.6 μm SSSS-ARB
SSSS + artificial 200 5 32 ~4.1 ~15.6 μm Aged-ARB aged
SSSS-ARB- SSSS 350 6 64 ~4.9 ~7.8 μm HT
45 Figure 4.2. Schematic representation of the ARB processing loop showing: (i) surface cleaning, (ii) stacking and (iii) roll bonding at a predetermined temperature. The number of rolling cycles shown in Table 4.1 includes an additional final rolling cycle to 50% reduction. The total true effective strain of each processed material was calculated using the following relation (Tsuji, 2006):
2 ln1 rn (4.1) 3 where r is the reduction in thickness per rolling cycle and n is the number of cycles.
4.2.3 Post-deformation annealing
The post-deformation annealing treatments were carried out in a muffle furnace exhibiting temperature control of ± 2° using an external thermocouple located at various locations in the furnace. For SSSS- and Aged-ARB, annealing was carried out for both 3 min and 6 h at 350 °C. This temperature was selected based on the detailed knowledge of the restoration processes that occur in both Al and Al(Sc)
(Ferry et al., 2005). For SSSS-ARB-HT, annealing was carried out for 3 min at 250,
300 and 350 °C.
4.3 Mechanical Testing
4.3.1 Hardness testing
Samples were sliced from the RD-ND section of the as-rolled material then ground through successively finer grades of SiC paper using water as lubricant. This was followed by polishing using 3 µm to 1 µm diamond paste in an oil-based lubricant.
Vickers microhardness measurements was conducted utilizing Leco M-400-H1 hardness testing machine equipped with visual display sensor (VDS) monitor and a rotary encoder attached to the indenter. The hardness of individual layers of the materials was determined after every ARB cycle in the RD-ND section using a 10 mg load since the examined section was too thin for higher loads. A total of 5-10
46 indentations per test area were taken to ensure consistency of data. The size of each indentation was accurately measured using SEM.
4.3.2 Tensile testing
ASTM-E8 standardised tensile samples were prepared from the SSSS-ARB samples in both the as-deformed and annealed states. Figure 4.3a shows a typical sample produced for tensile testing with Fig. 4.3b showing detailed sample dimensions. The samples were strained to failure using an INSTRON-1185 Universal Testing System consisting of a mechanical frame including a heavy motor for driving the vertical crosshead, an electronic controller and a desktop computer interfaced with the
INSTRON system controller. Tensile testing was carried out at room temperature using a constant crosshead speed of 0.5 mm/min and force monitored using a 10kN load cell. The displacement during testing was measured by a laser extensometer, whereby the samples were strapped with a special laser reflective tape at both ends of their gauge lengths.
4.4 Microstructural Analysis
4.4.1 Focused ion beam imaging
An FEI Nova NanoLab 200 focused ion beam (FIB) (hereafter called the DualBeam platform) was used in this thesis. The DualBeam platform combines a high resolution
FIB and a FEGSEM. The FIB uses a fine beam (~7nm) of highly energetic gallium ions (~30kV) that scan over the sample surface. At high beam currents the Ga+ beam rapidly sputters away the specimen surface which allows subsurface cross-sections to be viewed. However, at lower beam currents the secondary electrons or secondary ions emitted from the sample surface can be used for high resolution ion channelling contrast (ICC) imaging (see e.g. Fig. 4.4).
47 Figure 4.3. (a) Optical macrograph showing a representative ARB-processed sample (SSSS-ARB-HT) for tensile testing. (b) Schematic diagram of the tensile sample dimensions. Figure 4.4. Representative ion channelling contrast (ICC) micrographs (RD-ND section) showing the multi-layered annealed microstructure of SSSS-ARB with the dotted circle highlighting the Al/Al(Sc) interface located for FIB milling to produce site-specific TEM foils. The DualBeam platform was used in this thesis to examine ND-RD sections of the as-rolled and annealed materials. An area of microstructure suitable for ICC imaging was located on the sample surface and higher beam current (>1000pA) was selected.
A “mill-box” was then drawn onto the area of interest for milling. The beam scans over this area until milling is achieved. A smaller aperture is re-inserted (lower beam current <1000pA) for examination of milled structure. This loop of changing beam current was repeated until a good ICC image was obtained. ICC imaging was also used to locate suitable areas of microstructure for subsequent detailed analysis by electron backscatter diffraction (EBSD) interfaced to the FEGSEM.
4.4.2 Electron backscatter diffraction (EBSD)
A major part of the work carried out in this thesis was an investigation of the crystallographic nature of the deformed state of the ARB materials and the evolution of microstructure and texture during subsequent annealing. For this purpose, the powerful EBSD technique was carried out on all materials. Each sample of both the as-deformed and annealed states was prepared by conventional metallographic methods. Whereby, they were mounted in special (tailor-made) holder then ground on successively finer grades of SiC paper followed by mechanical polishing using 3
m and 1 m diamond paste. Electro-polishing was then carried out in a 20% Nital solution at around -30°C using a voltage of ~15V. Each sample was immersed for
30-40s in the solution followed by rinsing with ethanol and drying with compressed air.
ND-RD sections of the as-deformed and annealed samples were examined in both a
JEOL-JSM-7001F FEGSEM and the DualBeam platform in electron channelling contrast (ECC) and ICC imaging modes. Using these imaging modes, suitable areas of microstructure were selected for analysis by EBSD. Here, a TSL data acquisition system was utilized for determining crystallographic nature of the Al and Al-0.3%Sc
48 layers. The basic operational concept of the EBSD system consists of capturing an
EBSD pattern, creating an EBSD material file, indexing the EBSD pattern, adjusting the Hough parameters and finally collecting the orientation data (Randle and Engler,
2000). Throughout this thesis, a low angle grain boundary (LAGB) was defined as
o the misorientation between adjoining grains of %m < 15 , whereas a high angle grain
o boundary (HAGB) as %m > 15 .
4.4.3 Transmission electron microscopy
4.4.3.1 Site-specific TEM sample preparation
The DualBeam platform was also used to generate thin slices of material in specific areas of microstructure suitable for TEM. Initially, samples were electropolished by the method described for EBSD sample preparation. This enabled the microstructure to be investigated by ICC imaging for locating suitable sites for TEM. The principal regions of interest in this thesis were the interfacial regions of adjacent Al and Al(Sc) layers in partly annealed samples. Figure 4.5 and 4.6 shows some of the steps used for generating site-specific TEM samples from a typical roll-bonded sample.
The following provides an outline of the procedures used to generate TEM foils. A region of interest was located by ICC imaging. An appropriate TEM routine was then selected for automated sample preparation. The next stage involved the deposition of a Pt layer with fiducial markers (X) milled into the Pt pads (Fig. 4.5a): these were required for sample location and alignment during ion beam milling. A sloping trench on either side of the region of interest was milled with a high Ga+ beam current (~5nA), then both sides of the foil were polished further using a reduced beam current (e.g. 3-0.3nA) to reduce the thickness to ~1m. Next, the sample was tilted to 45 and cuts were made to partially free the TEM foil from the bulk sample
(Fig. 4.6). Here, a low beam current was used (e.g. 300 and 100pA) to further thin
49 Figure 4.5. (a) Scanning electron micrograph of a typical specimen prepared for transmission electron microscopy (TEM) by site-specific focused ion beam (FIB) milling showing the progression of the process. (b) TEM foil prepared by FIB ready for pick-up in “lift-out” status after manual thinning using a low beam current (< 50pA). Figure 4.6. 45 inverted-view image of a typical ARB TEM sample prepared by FIB showing the progression of the process in (a) to (c) indicated by the arrow with cuts made to partially free the thin foil. Gradually reduced magnification images in (d) and (e) displaying the hole depth at the sample surface generated by milling. the foil and finally generate a near electron transparent region on the foil. Further manual thinning of the foil was achieved using a lower beam current (< 50pA) to create a “lift-out” TEM cross section of thickness of ~ 100nm (Fig. 4.5b). Finally, the TEM foil was cut free at the holding edges.
The TEM foil lift-out process was carried out ex-situ using a very fine glass probe.
Here, the foil is attracted by electrostatic forces onto the probe and then transferred to the TEM grid in a similar manner. Figure 4.7a shows a top view on the initial stage of selecting the region of interest at the as-polished surface of the sample surface with Fig. 4.7b showing the TEM foil ready to place on the TEM grid.
4.4.3.2 TEM analysis
The TEM foils produced by FIB were examined using a Phillips CM200 TEM operating at 200kV. Samples were imaged in either bright field or dark field modes with scanning TEM (STEM) used for investigating the interfacial regions between layers. Selected area electron diffraction (SAED) patterns were taken as a first approximation to changes in average boundary misorientation and for determining the coherency of Al3Sc particles within the Al matrix. The local orientation changes within the fine features of the deformation substructure were investigated by convergent beam electron diffraction (CBED).
50 Figure 4.7. A top view FIB micrograph of the initial stage of selecting the ROI (dotted circle) for producing a TEM foil on the sample surface with arrows indicating the ARB sample width in (a) versus the final stage of freed TEM foil in a “lift-out” status in (b). Chapter 5 ______RESULTS I Deformation and Annealing Behaviour of Al/Al(Sc) Lamellar Composite (32 Alternating Layers: SSSS-ARB)
5.1 Introduction
A number of detailed studies of severe plastic deformation have been carried out individually on either Al or Al(Sc) alloys (Saito et al., 1999; Tsuji et al., 2002a;
Huang et al., 2003, Min et al., 2005; Kim et al., 2005a,b; Kamikawa et al., 2007b;
Huang et al., 2008a,b,c,d; Kamikawa et al., 2009), but there is no information on the deformation and annealing behaviour of these materials in the form of a multilayered structure. This chapter describes the deformation and annealing behaviour of a multilayered aluminium alloy composite consisting of 32 alternating layers of Al and
Al(Sc). This experimental program for generating this material is outlined in chapter
4. The focus of this chapter is to investigate the rolling and annealing behaviour of the material whereby the Al(Sc) alloy was heat treated to generate a supersaturated solid solution followed by warm rolling. Here, the initial Al(Sc) sheet was heat treated for 48h at 640 C and cold water quenched: this generated a starting hardness
51 of 29 VHN that was comparable with the hardness of the commercial purity Al sheet
(21 VHN), Fig. 5.1. A major aim of chapters 5-8 is to investigate the effect of both the initial hardness and microstructure of adjacent layers on deformation and annealing behaviour with this chapter focusing on materials with comparable work hardening behaviour.
Accumulative roll bonding (ARB) of the Al and Al(Sc) sheet materials was carried out at 200 C. This temperature was chosen initially since it is known that the onset of static precipitation of fine Al3Sc particles from the supersaturated solid solution takes ~2h at this temperature (Toropova et al., 1998). This is not reached during
ARB processing since the total time at 200 C for each rolling cycle is 5 min which equates to 25 min for the entire processing route (five rolling cycles). In chapter 6, however, the starting Al(Sc) alloy sheet was further aged for 4h at 300 C prior to
ARB to generate a starting hardness of 78 VHN and a distribution of nanosized, coherent Al3Sc particles with rolling carried out at the same temperature. Hence, chapter 6 investigates the effect of markedly different starting microstructures on deformation behaviour during ARB and annealing behaviour following processing.
5.2 Deformation Microstructures
5.2.1 Hardness of individual layers during ARB
The effect of scandium in solid solution on the relative hardening behaviour of the Al and Al(Sc) layers during ARB is shown in Fig. 5.1. The hardness measurements were conducted after each rolling pass, with the hardness impressions measured by scanning electron microscopy for improved accuracy. It can be seen that the hardness differential between Al(Sc) and Al increases from 8 to 24 VHN after deformation to a true strain of 4.1 (five rolling passes). Hence, the scandium present in solid solution imparts additional hardening of the Al(Sc) layers relative to the higher
52 Figure 5.1. Vickers hardness (HVN10 mg) as a function of true strain for the Al and Al(Sc) layers measured after each ARB cycle. purity Al layers and this is expected to generate different deformation microstructures.
5.2.2 Microstructural development within the Al and Al(Sc) layers
The deformation microstructure was investigated by both ion channelling contrast
(ICC) imaging and electron backscatter diffraction (EBSD) in the DualBeam platform and transmission electron microscopy (TEM) at particular sites in the roll- bonded material using FIB to generated site-specific TEM foils. Much of the analysis by these techniques was carried out on the ND-RD cross section of the material. A large shear strain was generated in the surface and subsurface regions of the ARB sheet during rolling without lubrication to large individual rolling reductions (~50% per pass). Hence, all investigations were limited to mid-thickness regions of the ARB material sheets where these effects are minimal.
The through-thickness microstructure of the fabricated material (hereafter termed
SSSS-ARB) is shown in Fig 5.2; the Al and Al(Sc) layers are clearly visible as darker and brighter alternating parallel layers. It can be seen that the Al layers are slightly thinner than Al(Sc) layers thereby indicating that a higher strain is accommodated in the former due to their lower hardness for a given rolling pass (Fig.
5.1). The average true strain in the Al and Al(Sc) layers was estimated to be 4.2 and
3.9, respectively. Figure 5.2 also shows small bright islands located at the interface of adjacent layers (arrowed) and these were identified to be oxide debris that probably formed during the intermediate heating stage prior to each rolling pass.
These oxides are sparsely distributed and irregular in shape and, thus, are not expected to have an adverse effect on the quality of the roll bonded interface (Quadir et al., 2007).
53 Figure 5.2. Overview ion channeling contrast (ICC) micrograph of SSSS-ARB (ND-RD section) showing several alternating Al and Al(Sc) layers. The small boxes are described in later figures. Figure 5.3 is an ICC micrograph showing nine alternating Al and Al(Sc) layers. The
Al layers are visible as fine grey scale contrast regions aligned along RD with the
Al(Sc) layers being more diffuse. The difference in layer thickness in Al and Al(Sc) is obvious in this micrograph. As shown in Fig 5.2, the oxide debris generally exist as discontinuous particles at the bonded interfaces (arrowed). However, Fig. 5.3 shows that matrix material flows around the larger oxide particles. It is also evident that the layers show some degree of waviness, which is an indication of the onset of shear banding during ARB. It is pertinent to note that the extent and type of shear banding in SSSS-ARB after five rolling passes is considerably different that produced by ARB using artificially-aged Al(Sc) sheet as the starting material (see e.g. Fig. 6.2). The reason for the marked difference in shear banding behaviour during ARB and its implications in fabrication of the materials is discussed in detail in chapter 8. In general, SSSS-ARB generated micro-shear bands largely restricted within the individual layers, and these microstructural heterogeneities rarely operate across layer boundaries, and are never found to operate across the entire sample thickness.
Figure 5.4 shows a higher resolution ICC micrograph showing an Al layer (middle) bounded by two Al(Sc) layers. The reasonably uniformly distributed strip-like structures within the Al layers are aligned parallel to RD and, thus, resemble the typical lamellar band (LB) structures characteristic of many types of heavily rolled sheet metals (Jazaeri and Humphreys, 2004a,b; Prangnell et al., 2004; Kamikawa et al., 2006). In Fig. 5.4, the substructure within the Al(Sc) layers is not easily resolved although the shear bands (SBs) can be seen to be aligned at ~30 to RD and confined inside the Al(Sc) layers. Since shear banding is restricted to the Al(Sc) layers and rarely observed in the Al layers, they probably initiate in the former. The lamellar substructure within the Al layers is shown more clearly in the higher resolution ICC micrograph in Fig. 5.5. LBs are highly elongated parallel to RD and have an average
54 Figure 5.3. ICC micrograph of SSSS-ARB (ND-RD section) showing nine alternating Al and Al(Sc) layers. Shear bands and oxide debris are labelled.
Figure 5.4. High resolution ICC micrograph (ND-RD section) showing lamellar bands (LBs) in the Al-layer (middle) and shear bands in the adjacent Al(Sc)- layers.
Figure 5.5. High resolution ICC micrograph (RD-ND section) of a typical Al- layer in SSSS-ARB illustrating the elongated nature of the LBs. thickness of ~0.4 m. A more detailed analysis of the LB structure in the Al layers was carried out using TEM.
Figure 5.6a shows a bright field TEM micrograph of a representative region within an Al layer revealing LBs aligned parallel to RD of average thickness of ~0.4 μm although some thicker bands are also present. This micrograph also reveals that the
Al layers have undergone some recovery via structural relaxation, curving and bulging of LB boundaries. The orientation of individual LBs and the misorientation between these features were determined by convergence beam electron diffraction
(CBED). Figure 5.6b shows four typical indexed Kikuchi patterns with Fig. 5.6c showing a more complete analysis of both the LB boundaries in Fig. 5.6a and the misorientation between bands; they are predominantly high angle grain boundaries
(HAGBs), that is, they have misorientations greater than 15°. Section 5.3 provides a more detailed analysis of the orientation distribution within the deformation substructure using EBSD.
The substructure within the Al(Sc) layers was also investigated by bright field TEM, as shown in a representative example in Fig. 5.7a. This micrograph reveals both the arrangement of lamellar bands and the high internal dislocation density of the bands.
The deformation features within the Al(Sc) layers were finer than those in Al layers, with an average thickness of ~0.15 μm. Typical examples of indexed Kikuchi patterns from individual LBs are shown in Fig. 5.7b with Fig. 5.7c showing a reconstructed image of Fig. 5.7a revealing a preponderance of low angle grain boundaries (LAGBs), that is, they have misorientations less than 15°. It is clear that scandium in solid solution prior to ARB has a significant influence on the resultant deformation substructure with LBs in the Al(Sc) layers being more highly dislocated and finer in scale compared with the Al layers. This behaviour is likely to generate a different hardening response in the layers, as revealed in Fig. 5.1. A more thorough
55 Figure 5.6. (a) Bright field TEM micrograph (ND-RD section) of a typical Al- layer in SSSS-ARB showing internal dislocations and lamellar band boundaries. (b) Four typical indexed diffraction patterns from LBs in (a). (c) Boundary misorientation map of (a) where grey and black lines represent boundaries with misorientations <15 and >15 , respectively (dashed lines denoted boundaries for which the misorientations were not measured). Figure 5.7. (a) Bright field TEM micrograph (ND-RD section) of a typical Al(Sc)-layer in SSSS-ARB showing a high internal dislocation density within cells. (b) Four typical indexed diffraction patterns from LBs in (a). (c) Boundary misorientation map of (a) where grey and black lines represent boundaries with misorientations <15 and >15 , respectively (dashed lines denoted boundaries for which the misorientations were not measured). analysis of the difference in microstructure between Al and Al(Sc) layers is given in section 8.2.1.
5.3 Deformation Textures
5.3.1 Texture of the Al layers after ARB
In order to measure the general texture of Al the layers, EBSD area scans were carried out, as demonstrated by the boxes shown in Fig. 5.2. Figure 5.8 shows the set of EBSD micrographs of these regions (4 m layer thickness). For texture analysis, data from the first two surface layers were discarded due to surface- damage, which induced uneven crystallographic randomization by intense shearing.
The acquired data from twelve interior layers in Fig. 5.8 were combined into a single data file and plotted as a 111 pole figure (Fig. 5.9) and 2 sections of Euler space
(Fig. 5.10). Both methods of texture representation reveal the copper-type rolling texture typical of all high stacking fault energy (SFE), face centred cubic (FCC) metals cold rolled to high strains (Humphreys and Hatherly 2004). The texture consists of a well-defined -fibre extending from Brass (B: {011}[112]) to S3
({123}<634>) to Copper (Cu: {112}<110>). Other notable orientations are also present including Cube (C: {001}<100>), Rotated cube (RC: {001}<110>) and Goss
(G: {110}<100>).
The foregoing figures show the global texture of the Al layers. The orientations and boundary misorientations of the deformation features within specific Al layers are given in Figs 5.11 and 5.12. The EBSD micrographs and accompanying 111 pole figures in these figures reveal a clear dominance of {110}<112> (Brass) but with the presence of other rolling components including {4 4 11}<11 11 8> (Taylor),
{123}<634> (S3) and {241}<112> (S1). The white areas in the EBSD micrographs denote regions of all other orientations. The microstructure of these two different Al
56 Figure 5.8. Combined EBSD micrograph of twelve interior Al-layers in SSSS- ARB, as indicated by the boxed regions in Figure 5.2 (some scans are indexed for demonstration purposes). . Figure 5.9. EBSD-generated 111 pole figure of the twelve combined scans of the Al-layers shown in Fig. 5.7.
Figure 5.10. 2 sections of Euler space (Bunge notation) showing the orientation distribution of the Al layers in SSSS-ARB. Orientations of interest are labelled as B (Brass {011}<112>), S3 ({123}<634>), Cu (Copper {112}<110>), C (Cube{001}<100>), RC (Rotated Cube {001}<110>) and G (Goss {110}<100>). Figure 5.11. (a) EBSD micrograph of Al-layer in SSSS-ARB (ND-RD section) showing the LBs bounded by high angle grain boundaries (HAGBs) (black lines > 15o) and the distribution of dominant rolling texture components (15o spread about an ideal texture component). (b) 111 pole figure of area in (a) showing a typical Cu-type rolling texture. (c) Relative misorientation profile along the vertical line in (a) (parallel to ND). Figure 5.12. (a) EBSD micrograph of another Al-layer in SSSS-ARB (ND-RD section) showing the LBs bounded by HAGBs (black lines > 15o) and the distribution of dominant rolling texture components (15o spread about an ideal texture component). (b) 111 pole figure of area in (a) showing a typical Cu-type rolling texture. (c) Relative misorientation profile along the vertical line in (a) (parallel to ND). layers is dominated by RD-aligned LBs that are tens of microns in length and separated by HAGBs (black lines). It can also be seen that the LBs are often bowed at boundaries and triple junctions, thereby indicating that these layers have undergone both static and dynamic recovery during heating to, and rolling at 200 C.
A plot of the change in misorientation along the vertical line in Figs 5.11a and 5.12a shows that the average spacing of the HAGBs is ~0.7 μm in both layers.
5.3.2 Texture of the Al(Sc) layers after ARB
Figure 5.13 is an ICC micrograph of the material after ARB showing an Al layer between two Al(Sc) layers. As noted previously, the Al layer is composed of strip- like LBs aligned along RD. In the Al(Sc) layers, this type of substructure is so fine that it appears diffuse and beyond the resolution of FEGSEM, but there are easily recognizable coarse substructures within the layers, that is, deformation bands and shear bands. The former are visible in the upper Al(Sc) layer as dark and light regions. Again, the shear bands are confined to the Al(Sc) layers.
EBSD data of the lower Al(Sc) layer of Fig. 5.13 is given in Fig. 5.14: this Al(Sc) layer shows both a sharp and gradual transition between (132)[64¯3] and (231)[34¯6] texture components and shows relatively less frequent presence of HAGBs (dark lines) when compared to the Al layers described previously. The S3 variants in Fig.
5.14a are shown in the corresponding 111 pole figure in Fig. 5.14b. The relative misorientation along the vertical line in Fig. 5.14a is plotted in Fig. 5.14c, showing only six HAGBs over a distance of 18 μm. For the top Al(Sc) layer in Fig. 5.13, the grayscale contrast seen in the ICC micrograph is revealed in the EBSD micrograph in
Fig. 5.15. There is clear evidence of grain subdivision where four regions (A-D) are labelled. The boundaries of the A/B and C/D regions are high angle whereas the boundary between B/C is low angle. The distribution of orientations in Fig. 5.15a is shown in 111 pole figure in Fig. 5.15c showing that regions B and C have a spread in
57 Figure 5.13. ICC micrograph of an ND-RD section showing three alternating layers of Al and Al(Sc). Dashed boxes are regions described in next two figures. Figure 5.14. (a) EBSD micrograph of the lower Al(Sc)-layer in Fig 5.13 showing HAGBs and LAGBs as black and white lines, respectively. (b) 111 pole figure of area in (a) showing two complementary S3 texture components. (c) Relative misorientation profile along the vertical line in (a) (parallel to ND). Figure 5.15. (a) EBSD micrograph of the upper Al(Sc)-layer in Fig. 5.13 showing four regions (A-D) separated by HAGBs. (b) 111 pole figure of area in (a) showing brass and copper texture components. (c) Relative and accumulative misorientation profiles along the vertical line in (a) showing the cross-over of the A-D regions (taken parallel to ND). orientations centred at the brass texture components whereas regions A and D have a spread in orientations centred at the copper components. The cumulative misorientation profile, along the line shown in Fig. 5.15a, illustrates the considerable orientation changes across this part of the microstructure and the presence of
HAGBs. These are typical deformation banding structure without having a narrow transition band (~5m) between them.
5.4 Microstructural Evolution During Annealing
5.4.1 Development of microstructure and texture of Al and Al(Sc) layers after annealing for 3 min at 350°C
The as-deformed material was annealed for up to 6 h at 350 C (section 4.2.3).
Figures 5.16a and b shows an ICC and corresponding EBSD micrograph of the initial stages of annealing (3 min at 350 ºC) exhibiting a central Al layer bounded by two
Al(Sc) layers. In the Al layer, there is significant spheroidization of the deformation substructure and clear delineation between the Al and Al(Sc) layers can be seen. In the latter, very little microstructural change is evident from the as-deformed state
(see e.g. Figs 5.3, 5.4, 5.13) and the rapidly growing grains originating within the Al layer do not propagate into the Al(Sc) layers. The 111 pole figures of the orientation distribution in various regions in Fig. 5.16 are given in Fig. 5.17. Figure 5.17a gives the overall distribution of orientations showing the Cu-type rolling texture generated during ARB.
Despite the substantial grain coarsening in the Al layer, Fig. 5.17b reveals that no new major orientations are evident which is unexpected if conventional discontinuous recrystallization was operative in this layer. The texture of the Al(Sc) layers (Fig. 5.17c) indicates that very little microstructural change has occurred during this short annealing procedure and it can be seen in the figure that two major
58 Figure 5.16. (a) ICC micrograph and (b) corresponding EBSD micrograph (ND- RD section) of a typical area of microstructure in SSSS-ARB showing a central Al layer bounded by two Al(Sc) layers after annealing for 3min at 350 °C. The figures reveal considerable spheroidization and subgrain coarsening in the central Al-layer and a retained deformation substructure in the adjacent Al-Sc layers. In (b), the black and white lines represent boundaries of < 15o and > 15o misorientation, respectively. Figure 5.17. EBSD-generated 111 pole figures of the area shown in Fig. 5.16: (a) entire area; (b) central Al-layer and (c) the adjacent Al(Sc) layers. texture components, located at the Brass (011) [2 11] and Copper (121) [111] orientations, are present.
Figure 5.18 shows another region of microstructure in the lightly annealed material with Fig. 5.19 showing the texture in the various layers. Again, there is clear evidence of considerable spheroidization of the substructure within the Al layer with preferential coarsening of a few grains. Furthermore, the well-defined boundary between Al and Al(Sc) layers is evident and the upper and lower Al(Sc) layers show shear bands and a more undulated structure within these bands. Again, no major microstructural changes associated with annealing are evident. The 111 pole figures of Fig. 5.19 shows many similarities to the pole figures in Fig. 5.17, particular within the Al layer. In the Al(Sc) layers, however, there are three major orientations indicating that the microstructural division during ARB is different to that of the Al layers with considerable deformation banding and shear banding occurring thereby breaking the layers up into specific texture components.
Figure 5.20 is a high resolution ICC and corresponding EBSD micrograph of the boxed area in Fig. 5.18 for further analysis. There is significant spheroidization at many locations along the LB boundaries, and grains are separated within a horizontal band by LAGBs. However, EBSD reveals larger relative rotations within the layer compared with the as-rolled material. Therefore, a larger angular range is needed to encompass an entire band of material, and this leads to overlapping of orientations along the !-fibre. To illustrate the spread from the mean orientation within each band, the orientations of the data points from each band (numbered 1-8 in Fig. 5.20b) are plotted in separate 111 pole figures. The central orientation of each spread has been identified, and the nearest standard orientations are marked in the pole figures with connecting lines between poles. Band 4 is an exception as the central point lies
59 Figure 5.18. (a) ICC/EBSD micrograph combination (ND-RD section) of another area of microstructure in SSSS-ARB showing the central Al layer undergoing substantial grain coarsening and the adjacent as-deformed Al(Sc) layers (3min at 350 °C). In (b), the black and white lines represent boundaries of < 15o and > 15o misorientation, respectively. The boxed area will be analyzed in detail in Fig. 5.20. Figure 5.19. EBSD-generated 111 pole figures of the area shown in Fig. 5.18: (a) entire area; (b) central Al-layer and (c) the adjacent Al(Sc) layers. Figure 5.20. (a) and (b) ICC/EBSD micrograph combination of the boxed regions in Figure 5.18. In (b), white and black lines indicate LAGBs and HAGBs, respectively. Eight horizontal bands are also marked in the EBSD micrograph in (b) each centered at one or two orientations with an orientation spread shown in the corresponding 111 pole figures (right). between {124}<211> (S1) and {110}<112> (Brass) and, hence, both are shown in the pole figure.
Bands 1-5 contain mainly LAGBs, and some of these boundaries are curved towards the neighbouring bands. These grains are primarily of Cu ({112}<111>) and S3
({321}<436>) orientations. Bands 6 to 8 consist mainly of the Brass ({110}<112>) or S3 ({321}<436>) orientations and do not contain internal LAGBs. Again, despite the slight structural changes after 3 min at 350 °C, there is no major change to the rolling texture.
5.4.2 Rapid coarsening of certain grains within the Al layers
It can be seen in Figs 5.16, 5.18 and 5.20 that some regions of microstructure within the Al layers are in a more advanced stage of grain coarsening after 3 min at 350°C.
This is illustrated by the four representative ICC images in Fig. 5.21 where they all show coarser (~10 m) grains surrounded by fine-grained regions; the latter are ~ 8-
10× smaller in size than the rapidly-growing grains. The rapidly growing grains are located either at Al/Al(Sc) layer boundaries or slightly away from boundary
(inwards). In order to analyse the individual orientation of these coarse grains, circles were milled on specific grains using FIB in the DualBeam platform and EBSD carried out in manual mode in the usual manner.
Figure 5.22 shows an example of the FIB-milled coarse grain after tilting to 70° for
EBSD analysis. It is pertinent to note that a large sample area was examined since these coarse grains after 3 min at 350 °C were infrequent in the microstructure.
Figure 5.23 shows the orientations of 64 coarse grains in a ~0.6 mm2 area of microstructure; the rolling texture is clearly retained in the Al layers but the 111 pole figure also shows the absence of the Brass component and a larger spread in
60 Figure 5.21. Series of ICC micrographs showing substantial grain coarsening in Al-layers after annealing for 3 min at 350 °C. The circles represent focused ion beam (FIB) markings used to generate orientation data by EBSD of these rapidly growing grains. Figure 5.22. Representative SEM secondary electron micrograph of a rapidly growing grain in the Al-layer after stage tilting to 70° in readiness for EBSD.
Figure 5.23. (a) 111 pole figure of the orientation distribution of 64 rapidly growing grains within the Al-layer of SSSS-ARB showing retained rolling texture but with a notable absence of the Brass texture component, as indicated in the reference pole figure in (b). orientations from the ideal orientations of the -fibre (shown in Fig. 5.23b). The origin of this texture during annealing is discussed in section 8.3.1.2.
5.4.3 Microstructure and texture of Al- and Al(Sc) layers after annealing for 6
h at 350 °C
The foregoing analysis (section 5.4.2) showed that the Al(Sc) layers do not undergo substantial microstructural change during annealing for short times at 350 °C. In contrast, certain grains within the Al layers coarsen rapidly during the early stages of annealing. These grains also have orientations within the orientation spread of the deformation texture although there is a notable absence of Brass texture components.
The as-rolled material was further annealed for 6 h at 350 °C; representative ICC micrographs of the three cross-sections (ND-RD, TD-RD and ND-TD) are shown in
Fig. 5.24. Both the ND-RD and ND-TD sections in Fig. 5.24 reveal the generation of a very coarse grained microstructure in the Al layers where grains are of the size of the layer thickness. It is also evident in the TD-RD section that these grains are pancake-shaped. In contrast, the ND-RD and ND-TD sections show that the Al(Sc) layers are only slightly recovered exhibiting a fine (0.4- ) " # $ structure which is largely similar to the as-deformed microstructure.
These cross-sections also reveal shear band-induced waviness of the layers, but this is very minor in SSSS-ARB compared to Aged-ARB, as described in chapter 6.
Overall, extended annealing at 350 °C has generated a microstructure containing alternating coarse-grained and recovered layers that align parallel to RD. Figure 5.25 is a three-dimensional reconstructed micrograph of such a structure clearly revealing the pancake-like growth of grains in Al layers, by restricting the growth along ND and RD directions, and without these grains invading the Al(Sc) layers. It is pertinent to note that the widths of the different layers are similar after annealing, a
61
Figure 5.24. ICC micrographs showing the microstructure of SSSS-ARB after annealing for 6 h at 350 °C: (a) ND-RD section, (b) ND-TD section and (c) RD- TD section.
Figure 5.25. A reconstructed three-dimensional ICC micrograph showing the unique lamellar microstructure generated after annealing for 6 h at 350 °C. change from the ~4:3 ratio in the as-deformed state. Hence, the coarse grains within the Al layers appear to migrate slightly into the Al(Sc) layers.
Figure 5.26 shows an EBSD micrograph of the ND-RD section showing that many of the coarse grains in the Al layers are of full layer thickness. The layers labelled
{123}<634> and {142}<211> within the Al(Sc) layers do not contain HAGBs except at a single deformation band in {123}<643>. The layer labelled {110}<112> in Fig.
5.26 contains changes among complementary texture components at an average spacing of ~7 μm. The orientations of 128 coarse grains in the Al layers were determined by EBSD and given in the 111 pole figure in Fig. 5.27a. The rolling texture is largely retained after extended annealing, although a comparison with Fig.
5.27b shows that the {110}<112> (Brass) texture component is negligible. In summary, extended annealing of SSSS-ARB at 350 °C has generated a unique microstructure consisting of alternating layers of coarse grains (Al layers) and a recovered substructure (Al(Sc) layers); the latter are likely to be stabilized by the presence of scandium in the microstructure. The next section examines, in further detail, the role of scandium on the creation of this unique microstructure.
5.4.4 Al/Al(Sc) interface behaviour during annealing at 350 °C
The unique lamellar microstructure generated by annealing of SSSS-ARB is a result of unimpeded grain coarsening in the Al layers and the retention of the deformation substructure in the Al(Sc) layers. It is clear that scandium in solid solution in the aluminium matrix prior to ARB has a substantial influence on the deformation microstructure with significant differences in substructure development and texture evident between the Al- and Al(Sc)-layers (section 5.3). Furthermore, annealing also resulted in substantial retardation of normal processes of recovery and recrystallization in the Al(Sc) layers. In order to clarify the influence of scandium on
62 (1 3 2)[6 -4 3] (1 2 3)[-6 -3 4]
{1 1 0}<1 1 2>
( 1 4 2)[2 -1 1]
Figure 5.26. EBSD micrograph of SSSS-ARB after annealing for 6 h at 350 °C (ND-RD section) where LAGBs (5-15°) and HAGBs (> 15o) are denoted by black and blue lines, respectively.
Figure 5.27. (a) EBSD-generated 111 pole figure of 128 coarse grains in Al layers of SSSS-ARB showing the dominance of Copper and S3 orientations, as indicated in the reference pole figure in (b). microstructural development during annealing within both the Al(Sc) layers and at the Al/Al(Sc) interfacial regions, site-specific TEM foils were prepared using FIB
(section 4.4.3.1).
Figure 5.28a shows a scanning TEM (STEM) micrograph of the Al/Al(Sc) interface region after annealing for 6 h at 350 °C; a very coarse grain has formed in the Al layer at the left and the fine substructure associated with Al(Sc) is retained on the right of the interface. Figure 5.28b shows a schematic reconstruction of the interface structure revealing the very fine-grained structure within the Al(Sc) layer where the grain size remains submicron. The figure also reveals significant localised curvature at the interface as a result of what appears to be localised pinning by fine particles
(circled regions). Figure 5.29 is a bright field TEM micrograph of the same region shown in Fig. 5.28a (the same locations are circled in the micrographs). The substructure within the Al(Sc) layer is equiaxed with size 0.6 ) %re is a dispersion of nanosized and coherent Al3Sc particles within this layer. At the
Al/Al(Sc) interface, there is substantial boundary pinning by the Al3Sc dispersoids, as shown in Fig. 5.30, which is the white circled region shown in Fig. 5.29.
Since the Al(Sc) alloy sheet was initially prepared in the supersaturated condition prior to ARB, these Al3Sc dispersoids are known to form in this (Al-0.3%Sc) alloy after ~1000 s at 350 °C (Seidman et al., 2002). Hence, annealing of SSSS-ARB results in concomitant recovery and precipitation within the deformation substructure during annealing at this temperature; such an effect on microstructural development is discussed in detail in chapter 8. Within the coarse grains of the Al layer, there is also evidence of precipitation that is probably associated with boundary migration of these grains into the Al(Sc) layer in the early stages of annealing, that is, prior to the nucleation of the Al3Sc dispersoids. However, this migration may also have occurred after the formation of the Al3Sc dispersoids, whereby to prevent growth, it requires
63 Figure 5.28. (a) STEM image of the TEM foil prepared by site-specific FIB across an Al/Al(Sc) interface in SSSS-ARB (ND-RD section) after annealing for 6 h at 350 °C. The white and black circles indicate Al3Sc particle pinning the grains at the interface. (b) Schematic representation of the interface region shown in (a) illustrating the Al and Al(Sc) layers where a recovered substructure of ~1 m diameter is generated in the Al(Sc) layer. Figure 5.29. Bright field TEM micrograph of the area in Fig. 5.28a showing the fineness of the retained substructure in Al(Sc) and interfacial pinning by Al3Sc particles.
Figure 5.30. (a) TEM micrograph of the black-circled area in both Figs 5.28 and 5.29 demonstrating the particle/boundary interactions at the Al/Al(Sc) interface. (b) Schematic representation of (a) displaying the interface between Al-layer and Al(Sc)-layer highlighting the precipitate projection associated with the particle- interface boundary interaction area (circled). an optimum particle density which may not be present at the boundary, although these dispersoids are abundant in the interior of the Al(Sc) layers.
5.5 Summary
A high purity Al alloy and a supersaturated Al-0.3 wt.% Sc alloy were accumulative roll bonded at 200 °C to generate 0.5 mm gauge sheet consisting of 32 alternating layers of Al and Al(Sc). The deformation structure of Al and Al(Sc) layers contain lamellar bands aligned parallel to RD. The bands in the Al(Sc) layers are more refined than those in the Al layers, but they contain fewer high angle grain boundaries (HAGBs). Extensive dynamic recovery was observed in the lamellar bands of the Al layers. Subsequent annealing at 350 C eventually generated alternating layers of coarse grains (Al layers) and a recovered substructure (Al(Sc) layers). Within the Al layers, annealing resulted in rapid coarsening of certain grains within the deformation microstructure with the resulting coarsening process not significantly altering the rolling texture (-fibre) thereby indicating that static restoration did not occur by conventional nucleation and growth processes characteristic of discontinuous recrystallization.
In contrast to the rapid grain coarsening of the Al layers, the Al(Sc) layers were only slightly recovered, which generated a stable submicron grain structure even after extended annealing. It was observed that extensive boundary pinning by nanosized
Al3Sc particles impeded recovery and recrystallization in the Al(Sc) layers and these particles also severely impeded the propagation of coarse grains from within the Al layers into these layers. This ultimately generated a lamellar composite structure consisting of alternating coarse grained layers and slightly recovered layers. A more detailed discussion of the deformation and annealing behaviour of this Al/Al(Sc) lamellar composite is given in chapter 8.
64 Chapter 6 ______RESULTS II Deformation and Annealing Behaviour of Al/Al(Sc) Lamellar Composite (32 Alternating Layers: Aged-ARB)
6.1 Introduction
Chapter 5 described the development of a multilayered composite consisting of alternating layers of Al and Al(Sc) (termed SSSS-ARB). Based on the initial heat treatments, the hardness of the Al and Al(Sc) sheets prior to ARB were comparable with the resultant rolling procedures generating a multilayered structure (c.f. Fig.
5.25) in which the layers were aligned parallel to RD even after rolling to very high strains. The present chapter is focused on the microstructure and texture of a multilayered Al/Al(Sc) composite using the same base alloys as those given in chapter 5. However, the supersaturated Al(Sc) alloy was subsequently artificially aged at moderate temperature prior to ARB that generated a considerably higher starting hardness than the supersaturated alloy (see e.g. Fig. 6.2). The principal aim of this chapter is to investigate the roll bonding behaviour of these aluminium sheets where the additional heat treatment of the Al(Sc) has generated two Al alloys with a
65 significantly different starting hardness and microstructure. The resultant Al/Al(Sc) composite generated by ARB is hereafter termed Aged-ARB.
Similar to SSSS-ARB, the microstructure and texture of Aged-ARB during various stages of processing were investigated mainly by ICC imaging and EBSD in the
DualBeam platform and TEM using site-specific foils produced by FIB milling.
Again, most of the analyses were carried out on ND-RD sections of the fabricated material. Since a large shear strain was also generated in the surface and subsurface regions of Aged-ARB during rolling, all investigations were limited to mid-thickness regions of the as-rolled sheet.
6.2 Starting Microstructure of Al(Sc) Alloy
In this part of the work, the Al-0.3 wt.% Sc alloy (Al(Sc)) was initially heat treated for 48h at 640 °C followed by cold water quenching. The material was then artificially aged for 4h at 350 °C in order to generate a uniform dispersion of Al3Sc dispersoids in a coarse-grained matrix (Aged-ARB) prior to roll bonding. A more complete description of material preparation is given in section 4.2.2. The initial hardness of the Al(Sc) alloy in the SSSS- and Aged-ARB condition was 29 HVN and
78 HVN, respectively. The latter exhibits a much higher starting hardness due to the generation of the Al3Sc dispersoids throughout the aluminium matrix. Figure 6.1a shows a dark field TEM micrograph of the artificially aged microstructure of Aged-
ARB showing coherent, spherical Al3Sc particles of average diameter less than 10 nm. The selected area electron diffraction (SAED) pattern in Fig. 6.1b shows large and small diffraction spots that correspond to the Al matrix and coherent Al3Sc particles, respectively. Figure 6.1c shows a high resolution TEM micrograph of a typical particle in Fig. 6.1a demonstrating its spherical shape and full coherency with the aluminium lattice.
66 Figure 6.1. (a) Dark field TEM micrograph and (b) accompanying selected area diffraction pattern (<001> zone axis) of the artificially aged Al-0.3%Sc alloy, showing a uniform dispersion of coherent Al3Sc dispersoids. (b) High resolution TEM micrograph of a ~5 nm diameter particle showing the full coherency with the aluminium matrix. 6.3 Deformation Microstructures
6.3.1 Hardness of individual layers during ARB
The effect of scandium in the form of a dispersion of nanosized Al3Sc dispersoids on the relative hardening behaviour as a function of true strain of the Al and Al(Sc) layers during roll bonding is shown in Fig. 6.2 with the data for SSSS-ARB shown for comparative purposes. Similar to SSSS-ARB, hardness measurements were conducted after each ARB pass and the hardness impressions were measured by
SEM for improved accuracy.
Irrespective of the heat treatment method for generating a given starting microstructure in the Al(Sc) layers (Fig. 4.1), the hardness of the Al layers in both materials increased from 30 HVN to 44 HVN after two ARB cycles ( = 1.38) with a plateau in hardness thereafter resulting in an overall hardness increase of ~14 HVN.
During ARB, the hardness of SSSS-ARB increased from 29 to 68 HVN at = 4.2
(D'()*+% -ARB increased in hardness from 78 to 88 HVN at = 2.0
(D'()*, -% %% %. -
ARB is consistent with the work of Apps et al. (2005) on a similar alloy after equal channel angular pressing (ECAP). These workers argued that a starting microstructure containing a dispersion of fine Al3Sc dispersoids results in a homogenized microstructure after ECAP.
6.3.2 Microstructural development within the Al and Al(Sc) layers
The through-thickness microstructure of roll bonded Aged-ARB is shown in Fig. 6.3 presenting a representative section of the 32 alternating layers of Al and Al(Sc). The principal difference between Aged-ARB (Fig. 6.2) and SSSS-ARB sheets is the significant waviness of the layers in the former. This behaviour appears to be a result of the formation of shear bands inclined at ±17±3 to RD throughout the sample
67 Figure 6.2. Vickers hardness (HVN-10 mg) as a function of true strain for the Al and Al(Sc) layers measured after each ARB cycle in Aged-ARB. Data for SSSS- ARB (Fig. 5.1) is included for comparative purposes. thickness during rolling, as indicated in Fig. 6.3. Along the shear bands, the thickness of the Al layers remains uniform, but they often converge due to localized narrowing of the Al(Sc) layers consistent with multiple necking. Similar to SSSS-ARB, the interfacial regions contain a discontinuous distribution of oxide fragments (arrowed in Fig. 6.3). While the total true strain of the roll bonded Al/Al(Sc) sheets was essentially the same for both SSSS- and Aged-ARB, the different rates of thinning of both the Al and Al(Sc) layers resulted in larger true strain within the Al layers; that is
= 4.63 in SSSS-ARB and = 4.83 in Aged-ARB.
Figure 6.4a is an ICC micrograph showing seven alternating Al and Al(Sc) layers and the extensive shear banding through the layers. The layers in Aged-ARB are more extensively undulated compared with SSSS-ARB (see e.g. Fig. 5.2). Figure
6.4 also shows that he Al layers have a sharper grey scale contrast aligned along RD whereas the Al(Sc) layers are more diffusely shadowed structures. The lamellar structure within the Al layers can be seen in the higher resolution ICC micrograph in
Fig. 6.4b. In this image, they are aligned along the shear bands, as bands with an inclination of 17o ±3 to RD and are accommodating roughly 3-4 bands of width in a
2 m scale bar length. The LBs were measured to be ~0.7-1.0 μm in thickness which is slightly greater than that generated in SSSS-ARB.
The deformation substructure within the Al(Sc) layers was investigated in more detail by bright field TEM with a representative region of microstructure shown in
Fig. 6.5. This figure demonstrates that the LBs in these layers are separated by dislocation walls (dark lines) and the band thickness is ~0.2-0.3 μm, which is comparable to the thickness of LBs measured in the Al(Sc)-layers of SSSS-ARB.
The planes and their associated zone axes of five consecutive LBs in Fig. 6.5a were measured by CBED in the TEM. The beam axis was close to [011] zone axis and there is also a little TD rotation (i.e. TD // beam axis) between the patterns. This
68 Figure 6.3. Overview ion channeling contrast (ICC) micrograph of Aged-ARB (ND-RD section) showing several alternating Al and Al(Sc) layers. The small boxes are described in later figures. Oxide debris at the Al/Al(Sc) interface and shear bands are also labelled. Figure 6.4. (a) Magnified ICC micrograph (ND-RD section) of Aged-ARB showing seven alternating Al and Al(Sc)-layers. Two shear bands are also labelled. (b) Higher resolution ICC micrograph (ND-RD section) of an Al-layer showing lamellar bands (LBs). Figure 6.5. (a) Bright field TEM micrograph (ND-RD section) of a typical Al(Sc)- layer in Aged-ARB showing high density of internal dislocations and lamellar band boundaries. (b) Five typical indexed CBED patterns from five corresponding LBs in (a) showing very small change of orientations between them. indicates that the misorientations between adjacent LBs in this Al(Sc) layer is quite small, that is, they are LAGBs. This is consistent with the structure of the Al(Sc) layers in SSSS-ARB.
6.454B Deformation Textures
6.4.171B General texture development
The deformation textures, as measured by EBSD, were carried out on the Al layers by area scanning (4 m whole layer thickness) throughout the sheet thickness, as shown by the boxes in Fig. 6.3. Figure 6.6 shows the set of EBSD micrographs of these regions. Similar to SSSS-ARB, the data for the first two surface layers were discarded due to the intense shearing generated at the sheet surface by the unlubricated rolls.
The acquired data from the twelve interior layers were combined into a single data file and plotted as a 111 pole figure and 2 sections of Euler space in Fig. 6.7 and 6.8, respectively. The pole figure shows that the orientation distribution of Aged-ARB in the as-rolled condition appears to be a distorted -fibre, which is different to that produced in SSSS-ARB for a comparable rolling strain (see e.g. Fig. 5.9). This distortion is revealed more clearly in the ODF of Fig. 6.8 although the classic FCC - fibre is still evident whereby the orientations range from {011}<112>,
{123}<634>{112}<110 to {001}<100>). It can also be seen that {001}<110> and
{110}<100> texture components are also present. The most notable features of the
ODF in Fig. 6.8 are: (a) the high intensities of Brass, S3 and Copper orientations, and
(b) the generation of two partial fibre textures, whereby one is extracted from Brass to Rotated-Cube which is clear in the 2= 0° section and another is S3 extracted from
2= 10–to–35° sections, which is much higher than that generated in SSSS-ARB
(Fig. 5.10). A further comparison of these figures shows that the maximum intensity
69 Figure 6.6. Combined EBSD micrograph of twelve interior Al-layers in Aged- ARB, as shown in Figure 6.1 (some scans are labelled for demonstration purpose). Figure 6.7. EBSD-generated 111 pole figure of the twelve combined scans of the Al-layers shown in Fig. 6.6.
Figure 6.8. 2 sections of Euler space (Bunge notation) showing the orientation distribution of the Al layers in Aged-ARB. Orientations of interest are labelled as B (Brass{011}<112>), S3 ({123}<634>), Cu (Copper {112}<110>), C (Cube {001}<100>), RC (Rotated Cube {001}<110>) and G (Goss {110}<100>). for Aged-ARB and SSSS-ARB is 12.5× and 16.3× random, respectively. The possible reasons for the difference in texture development within the Al layers of these materials are discussed in section 8.2.4.
6.4.2 Effect of shear banding on local texture development
As noted in section 6.3.2, the microstructure of the Al layers consists of a strip-like substructure uniformly distributed and largely aligned parallel to RD and, hence, they a consistent with the deformation features generated in SSSS-ARB. However, Fig.
6.4 clearly demonstrates that the lamellar bands are traversed by gross shear bands.
The difference in the orientation distribution in the ODFs for SSSS- and Aged-ARB, respectively, highlights the likely effect of shear banding on texture development (a detailed discussion is given in section 8.2.4.2). The process by which a shear band alters the orientations of the lamellar bands in an Al layer can only be seen from local orientation measurements of a newly-formed shear band.
Figure 6.9 shows an ICC image and corresponding EBSD micrograph of an Al layer in Aged-ARB showing lamellar bands sheared by an operative shear band (marked on the micrographs). The lamellar bands, whose boundaries are parallel to RD, are rotated ~15 from RD in the sheared region (marked in Fig. 6.9b). The distribution of orientations in the un-sheared and sheared regions (enclosed by lines in Figs 6.9a-b) is shown by the 111 pole figures in Fig. 6.10. There is clear indication of a ~15 rotation about TD in the sheared region. Figure 6.11 shows a series of 111 pole figures generated by plotting individual orientations at ~1 m intervals along a lamellar band into the sheared region where misorientations of 9-20 about TD are evident between the matrix ‘M’ and the sheared segment ‘S’.
70 Figure 6.9. (a) ICC micrograph of Aged-ARB (ND-RD section) showing a freshly deformed shear band in which the lamellar band boundaries have been inclined ~15° to RD. (b) EBSD micrograph of region in (a) showing the HAGB’s in black lines.
Figure 6.10. <111> pole figures of non-sheared and sheared regions (bounded by grey lines), as shown in Fig. 6.9. In the sheared region, a ~15° rotation is generated about TD. Figure 6.11. Series of 111 pole figures showing the crystallographic rotation of the lamellar bands during a traverse from the outside (M, matrix) to the inside of the sheared segments (S) of the EBSD micrograph in Fig. 6.9. The magnitude of the rotation in each case is marked, and varies from 9-20 about TD. 6.5 Microstructural Evolution During Annealing
6.5.1 Microstructure and texture of Al and Al(Sc) layers after annealing for 3 min at 350 °C
A detailed investigation of the static restoration behaviour of SSSS-ARB was given in section 5.4. Similar to this material, Aged-ARB sheet was annealed for 3 min at
350 C for investigating the early stages of static restoration. Figure 6.12 shows a notable change in the microstructure after annealing for 3 min where a ~12 μm diameter grain (arrowed white) has grown along the line of an operating shear band within the Al layer. However, most of the subgrains in the Al layers are less than 3
μm in diameter and the microstructure of the neighbouring Al(Sc) layers show no sign of rapid grain coarsening.
6.5.2 Microstructure and texture of Al and Al(Sc) layers after annealing for 6 h at 350 °C
The microstructure of Aged-ARB after annealing for 6 h at 350 C is given by the
ICC micrograph in Fig. 6.13. The Al layers contain coarse grains consuming the layer thickness bounded by Al(Sc) layers that have not undergone extensive static restoration. This generated a similar microstructure as the annealed SSSS-ARB but the more pronounced waviness of the layers in Aged-ARB is clearly visible.
Similar to SSSS-ARB, the microstructure within the Al(Sc) layers after extensive annealing is similar to that observed directly after deformation whereby the deformation substructure persists after extensive annealing and, while the cells are now equiaxed of size ~0.4 μm, they still contain a high density of dislocations (Fig.
6.14). It can also be seen in Fig. 6.13 that shear bands are still present in the Al(Sc) layers. It is pertinent to note that Al(Sc) in Aged-ARB contained fine Al3Sc dispersoids prior to rolling whereas Al(Sc) in SSSS-ARB precipitated dispersoids after the first few seconds of annealing at 350 C. Regardless of their stage of
71 Figure 6.12. ICC micrograph (ND-RD section) of the partly recrystallized Al- layer in Aged- ARB showing a rapidly growing grain (arrowed) within a shear band (after annealing for 3 min at 350 °C).
Figure 6.13. ICC micrograph (ND-RD section) of Aged-ARB showing grain growth to full layer thickness within the Al-layers (after annealing for 6 h at 350 °C). Figure 6.14. Bright field TEM micrograph (ND-RD section) of the deformation substructure after annealing for 6 h at 350 °C within a typical Al(Sc) layer showing a moderately recovered deformation substructure. formation, the Al3Sc dispersoids in both Al(Sc) materials are clearly retarding normal recovery and/or recrystallization processes.
Figure 6.15 shows a representative EBSD micrograph of the extensively annealed material showing coarse-grained Al layers and recovered Al(Sc) layers. Again, the waviness of the layers is clearly evident in the micrograph. The various types of grain boundary in the microstructure is also shown whereby LAGBs (5-15°) and
HAGBs (>15°) are shown as black and blue lines, respectively. An EBSD analysis of the coarse grains in the microstructure was carried out with Fig. 6.16 showing a
111 pole figure of 130 coarse grains.
The texture is not random with the -fibre texture components being largely retained although there is a larger spread in orientations compared to annealed SSSS-ARB
(Fig. 5.27). Therefore, a similar process of static restoration is operating in both
SSSS- and Aged-ARB with the additional orientation spread in the latter likely to be caused by shear-band induced waviness of the alternating layers. Such an effect is discussed in detail in sections 8.2.4.2 and 8.3.1.3.
Since the Al(Sc) layers are only partially recovered, they are expected to exhibit the same types of orientations as the as-deformed material. Figure 6.17 shows a series of
111 pole figures of typical regions within four different Al(Sc) layers where it can be seen that layers 1 and 2 contain texture components ranging from Cu ((121) 111 ) to S3 ((123) 6 34 ) and S2 ((132) 4 2 1 ) to S3 ((321) 3 4 6 ), respectively. It can also be seen that layer 3 contains Brass texture components whereas layer 4 shows a single S1 component ((241) 112 ).
72 Figure 6.15. EBSD micrograph of Aged-ARB (ND-RD section) after annealing for 6 h at 350 °C showing the curvature of the Al and Al(Sc) layers. Black lines indicate HAGBs (>15°).
Figure 6.16. 111 pole figure showing individual orientations of coarse grains within the Al layers of Aged-ARB after annealing for 6 h at 350 C, revealing some degree of randomness of the resultant texture. Figure 6.17. 111 pole figures of individual Al(Sc)-layers after annealing for 6 h at 350 C showing texture components as: (a) Al(Sc)-1 with two orientations; (b) Al(Sc)-2 with two orientations; (c) Al(Sc)-3 with single orientation, and (d) Al(Sc)-4 with single orientation. 6.6 Summary
A high purity Al alloy and an artificially aged Al-0.3 wt.% Sc alloy were accumulative roll bonded at 200 °C to generate 0.5 mm gauge sheet consisting of 32 alternating layers of Al and Al(Sc). The deformation microstructure of Al and Al(Sc) layers in this material was considerably different to that observed in SSSS-ARB although the Al layers contain lamellar bands aligned largely parallel to RD but with some deviation due to the generation of undulating layers of Al and Al(Sc). The marked waviness of the alternating layers was argued to be a result of extensive and large-scale shear banding throughout the sheet thickness. This also resulted in notable differences in deformation textures of SSSS- and Aged ARB.
Extended annealing of Aged-ARB at 350 C generated alternating layers of coarse grains (Al layers) and a recovered substructure (Al(Sc) layers) and the substantial waviness of the layers was inherited from the as-deformed material. Similar to SSSS-
ARB, rapid coarsening of certain grains within the deformation microstructure within the Al layers generated a texture containing a spread in rolling texture components
(-fibre) but with a much larger spread compared with that generated in SSSS-ARB.
Again, the Al(Sc) layers were resistant to recrystallization and a stable submicron grain structure containing shear bands was observed after extended annealing. The stabilization of the microstructure against recrystallization within the Al(Sc) layers was argued to be a result of the nanosized Al3Sc dispersoids that were generated during the heat treatment prior to ARB (Fig. 6.1).
73 Chapter 7 ______RESULTS III Microstructure and Mechanical Properties of Al/Al(Sc) Lamellar Composite (64 Alternating Layers: SSSS-ARB-HT)
7.1 Introduction
Chapter 5 and 6 described the microstructural development in an Al/Al(Sc) multilayered composite using five ARB cycles to generate 32 alternating Al/Al(Sc) layers. The influence of initial heat treatment of the alloys on microstructural development and general stability of the layers during rolling was examined in detail, as well as the influence of the effect of low temperature annealing on the final microstructure. The present chapter investigates the same heat treated material as that described in chapter 5, but with an additional rolling pass to generate 64 alternating layers and with deformation carried out at the higher temperature of 350
°C. This material is hereafter termed SSSS-ARB-HT. The as-rolled material was also used for an investigation of the influence of microstructural change during annealing on mechanical behaviour. Here, annealing was carried out for 3 min at
250 to 350 C and the microstructure and texture were investigated by EBSD.
74 7.2 Microstructural Development During Deformation and Annealing
7.2.1 General observations
The multilayered composite produced by six consecutive rolling cycles generated 64 alternating Al and Al(Sc) layers in 0.5 mm sheet thickness and a total true strain of
~4.9. A low magnification RD-ND section ICC micrograph of the material is given in Fig. 7.1 showing twenty alternating layers from the mid-thickness of the sheet.
Unlike SSSS-ARB (chapter 5), the layers are not aligned parallel to RD but were undulated in a manner similar to Aged-ARB (chapter 6). EBSD was carried out in the mid-thickness of SSSS-ARB-HT where Fig. 7.2a shows that the Al and Al(Sc) layers are elongated but substantial distortion of the individual layers is evident.
Furthermore, both type of layer was highly fragmented with a large fraction of
HAGBs generated within the Al(Sc) layers and a coarser recovered substructure within the Al layers. This demonstrates the significant influence of the distribution of scandium in the microstructure on the deformation substructure, as discussed in section 8.2.2.
Annealing was carried out for 3 min at temperatures ranging from 250 to 350 C and resulted in substantial microstructural changes mainly in the Al layers (Fig. 7.2).
Fig. 7.2b-c shows the evolution of the microstructure indicating rapid recrystallization in the Al layers and a recovered substructure in the Al(Sc) layers.
After 3 min at 300 C (Fig. 7.2c), the Al layers are completely recrystallized consuming the entire layer thickness although some regions of deformation microstructure persist (see e.g. Fig. 7.15).
75
Shear band
Figure 7.1. Overview ion channeling contrast (ICC) micrograph of SSSS-ARB-HT (ND-RD section) showing several alternating Al and Al(Sc) layers and shear banding.
Figure 7.2. EBSD micrographs of SSSS-ARB-HT (ND-RD section): (a) as- deformed sample, and after annealing for 3 min at (b) 250 °C; (c) 300 °C and (d) 350 °C (ND-RD sections; RD is horizontal).
7.2.2 Deformation textures within individual deformed layers
Figure 7.3 is an EBSD micrograph of the as-deformed material showing fourteen alternating layers of Al and Al(Sc) (labelled). These layers were more difficult to identify compared with SSSS-ARB. Based on the gradual colour changes observed in Fig. 7.3, the Al layers appear to show a gentle change in orientation over large distances although a few HAGBs can be seen aligned approximately parallel to RD.
However, these features are of a lower frequency compared with the lamellar band boundaries (LBs). In this context, the Al layers in this material are different to SSSS-
ARB, whereby the latter contains a larger number of LBs. Similar to SSSS-ARB, the
Al(Sc) layers in SSSS-ARB-HT contain a fine, partly recovered substructure.
The average thickness of the Al and Al(Sc) layers were measured to be 8.7 m and
6.2 m, respectively in the as-deformed SSSS-ARB-HT. However, while the Al layers are thicker on average, there is more deviation in thickness in the Al(Sc) layers where some regions are thicker than the Al layers. Hence, it is clear that substantial instability in the form of necking occurs in the Al(Sc) layers and some areas actually show contact between the Al layers, thereby indicating substantial non-uniform plastic flow in the material. This behaviour was not observed in SSSS-ARB (chapter
5) but is similar to that observed in Aged-ARB (chapter 6)
The texture of the combined layers in Fig. 7.3 is shown in the 111 pole figure in Fig.
7.4a. The classic !-fibre deformation texture, as expected in high SFE, FCC alloys, was generated in SSSS-ARB-HT although the texture is not as uniform as that generated in conventional aluminium after cold rolling to high strain (Hansen et.al,
2004). Figures 7.3b and 7.3c show the texture of the individual Al and Al(Sc) layers respectively, where it can be seen that the former displays a larger spread in orientations compared to the Al(Sc) layers; the latter also generates a more regular !- fibre deformation texture. This indicates that the Al layers are either not deforming
76
Figure 7.3. High magnification EBSD micrograph of as-deformed SSSS-ARB-HT (ND-RD section) showing alternating Al/Al(Sc) layers (labeled).
Figure 7.4. 111 pole figures generated from EBSD data of as-deformed SSSS-ARB-
HT showing the texture of: (a) Al and Al(Sc) combined layers with intensity contours; (b) Al layers (raw EBSD data), and (c) Al(Sc) layers (raw EBSD data)
(inset shows main FCC rolling texture components).
in plane strain conditions or the layers are continually recrystallizing during the various rolling cycles at 350 °C thereby randomizing the texture. This aspect of deformation texture development in the Al layers is discussed in section 8.3.2.
Figures 7.5 and 7.6 shows a series of 111 pole figures of the specific layers labelled in Fig. 7.3. The 111 pole figure in Fig. 7.5a indicates that Al-1 is a highly fragmented layer and is very unlike that expected in conventional cold rolled aluminium thereby indicating that the layers have recrystallized to some extent. A1-
2 shows a major texture component centred at ]81111)[1144( , termed the Taylor component (Randle and Engler, 2000; Humphreys and Hatherly, 2004), although there is a large spread in orientations of 35-45°. Both Al-4 and Al-5 have generated rather complex textures where a given layer has split into two major texture components: ]211)[241( to ]643)[321( in Al-4 and ]010)[101( to ]211)[241( in Al-
5, although their spreads are continuous from one to the other. The spread in orientations in Al-4 and Al-5 is ~25° and ~35°, respectively. It can also be seen that
Al-4 has generated a complementary texture component to Al-2. Finally, Al-6 shows a texture closer to the classic !-fibre deformation texture. These observations are consistent with recrystallization occurring during multi-pass hot rolling to produce the multilayered composite.
Figure 7.6 shows the texture of various Al(Sc) layers shown in Fig. 7.3, although every layer is not given due to difficulties locating them in the micrograph (i.e.
Al(Sc)-2, -4 and -5). The deformation textures of the individual Al(Sc) layers are somewhat different to that generated in the Al layers where it can be seen that the distribution of orientations are more comparable to those comprising the classic !- fibre (see inset in Fig. 7.5). Nevertheless, there are subtle differences in the Al(Sc) layers with high intensities at some specific orientations, such as the complementary
Copper components dominant in Al(Sc)-3 (Fig. 7.6b) and the S1-components in both
77
Figure 7.5. 111 pole figures of the deformed Al layers in SSSS-ARB-HT, labeled in Figure 7.3, showing the texture of: (a) Al-1; (b) Al-2 with superimposed copper texture variant; (c) Al-4 with superimposed S1 and S3 components; (d) Al-5 with superimposed Goss ]010)[101( and S1 (]211)[241 components, and (e) Al-6 (inset shows main FCC rolling texture components).
Figure 7.6. 111 pole figures of the deformed Al(Sc) layers in SSSS-ARB-HT, labeled in Figure 7.3, showing the texture of: (a) Al(Sc)-1; (b) Al(Sc)-3 with superimposed copper texture component; (c) Al(Sc)-6 with superimposed
S1 (]121)[214 component; (d) Al(Sc)-7 with superimposed
S1 (]211)[241 components.
Al(Sc)-6 (Fig. 7.6c) and Al(Sc)-7 (Fig. 7.6d). It is clear that the specific textures in these individual Al(Sc) layers effectively generates the standard !-fibre deformation texture.
Figure 7.7 and 7.8 shows typical misorientation profiles parallel to ND within the Al and Al(Sc) layers, respectively. Since the textures are different in a given layer, the misorientation profiles also vary between layers. Figure 7.7 shows a misorientation profile across the width of Al-2, -3, -5 in (a-c) respectively, where they exhibit an
HAGB spacing of ~1.1, ~0.75, ~1.7μm to give an average of ~1.2 μm. However, Fig.
7.8 shows that this HAGB spacing in the Al(Sc) layers is considerably lower (~0.79
μm in Al(Sc)-1, -3, -7). In some cases the width of a single Al-spacing of HAGB is much larger in a given Al layer, where it was found for Al-5, that the HAGB spacing is up to 3 ) / 0 -% % 1 % % indicates that the material has recrystallized prior to the final rolling pass such that these partly/fully recrystallized grains are compressed to a low strain in the final deformation pass.
7.2.3 Texture development within individual layers during annealing
7.2.3.1 Microstructure after annealing at 250 C
An enlarged EBSD micrograph of Fig. 7.2b showing the as-deformed microstructure following annealing for 3 min at 250 C is given in Fig. 7.9. There are four and three layers, respectively, of Al and Al(Sc) shown in the figure (labelled). Both types of layer have undergone static recovery, but to a significantly different extent. For example, grains of up to ~5μm in diameter are evident in the Al layers whereas the
Al(Sc) layers are only slightly recovered and resemble that of the deformation microstructure (Fig. 7.3). Figure 7.9 also demonstrates clearly the non-uniformity in
78
Figure 7.7. Typical misorientation profiles within the Al layers in the as-deformed SSSS-ARB-HT: (a) Al-2, (b) Al-3 and (c) Al-5 (parallel to ND).
Figure 7.8. Typical misorientation profiles within the Al(Sc) layers in the as- deformed SSSS-ARB-HT: (a) Al-Sc-1, (b) Al-Sc-3 and (c) Al-Sc-7 (parallel to ND).
Figure 7.9. EBSD micrograph of SSSS-ARB-HT annealed for 3 min at 250 °C showing four Al layers and three Al(Sc) layers revealing the differential thinning and bending of the various layers after rolling.
thickness of the various layers generated by ARB, which is particularly noticeable between Al and Al(Sc).
The distribution of all orientations in the EBSD micrograph of Fig. 7.9 is shown as the 111 pole figure in Fig. 7.10a where a distorted !–fibre texture is evident. The distribution of orientations in the specific Al and Al(Sc) layers are given in Fig.
7.10b and 7.10c, respectively, showing the similar distortion of the !–fibre in the Al layers and a more standard !–fibre in the Al(Sc) layers. Figure 7.11 shows the texture within each of the labelled Al layers in Fig. 7.9 where it can be seen that Al-1
(Fig. 7.11a) has two major texture components centred close to ]111)[112( and
]112)[011( . Al-2 and Al-3 show a significant deviation from the standard !-fibre showing no specific texture components whereas Al-4 is comprised of two S components close to ]211)[132( and ]124)[132( . The distribution of orientations of the individual Al(Sc) layers of Fig. 7.9 are given as 111 pole figures in Fig. 7.12, where it can be seen that Al(Sc)-1 consists of two S1 components ]112)[142( and
]112)[124( ,whereas the other two layers have a reasonably well-developed !-fibre texture.
Therefore, in the Al layers the orientations are either concentrated at a few components of !-fibre or distributed as a distorted !-fibre, as observed in the deformed state (Fig. 7.5). Likewise, the Al(Sc) layers show a distorted !-fibre after both deformation and annealing (at 250 °C) except for the example in Fig. 7.12a, which has generated a complimentary texture component.
The misorientations within the Al and Al(Sc) layers (parallel to ND) are given in Fig.
7.13 and 7.14, respectively. Figure 7.13 shows typical misorientation profiles across the width of Al-1-to-3 whereby they have an average HAGB spacing of ~1.6, ~2.0 and ~1.5μm, respectively (average of ~1.7 μm). In contrast, Fig. 7.14 shows that this
79
Figure 7.10. 111 pole figures generated from EBSD data of the sample given in Figure 7.9 showing the texture of: (a) Al and Al(Sc) combined layers with intensity contours; (b) Al layers only, and (c) Al(Sc) layers only (inset shows main FCC rolling texture components).
Figure 7.11. 111 pole figures of the deformed Al layers labeled in Figure 7.9 showing the texture of: (a) Al-1 with superimposed Brass and Copper components; (b) Al-2; (c) Al-3, and (d) Al-4 with superimposed S2 ]211)[132( and (]124)[132 components. \
Figure 7.12. 111 pole figures of the deformed Al(Sc) layers labeled in Figure 7.9 showing the texture of: (a) Al(Sc)-1 with superimposed S1 ]112)[124( and (]112)[142 components; (b) Al(Sc)-2, and (c) Al(Sc)-3 (inset shows main FCC rolling texture components).
Figure 7.13. Typical misorientation profiles within the Al layers in as-deformed and annealed SSSS-ARB-HT (3 min at 250 °C): (a) Al-1, (b) Al-2 and (c) Al-3 (parallel to ND).
Figure 7.14. Typical misorientation profiles within the Al(Sc) layers in as- deformed and annealed SSSS-ARB-HT (3 min at 250 °C): (a) Al-Sc-1, (b) Al-Sc-2 and (c) Al-Sc-3 (parallel to ND). spacing is considerably less (~0.9 μm) in Al(Sc). Nevertheless, in some regions the width of a single Al-spacing of HAGB is much larger in a given Al layer, where it was found for Al-2 % % '34 1 1 ) / 0 ,
Interestingly, a similar scenario occurs for the Al-Sc-1 (Fig. 7.14) and even exceeds that for the Al-2 situation where an HAGB spacing greater than 4 ) is evident.
7.2.3.2 Microstructure after annealing at 300 C
The EBSD micrograph of Fig. 7.2c is shown again in Fig. 7.15, where Fig. 7.15a and
7.15b, respectively, reveal the HAGBs and LAGBs. It is clear that annealing at this temperature has resulted in substantial static restoration of the Al layers with some grains have growing to full layer thickness although some areas are only recovered, as shown clearly in the EBSD image quality (IQ) map in Fig. 7.15c. It can be seen in
Fig. 7.15c that grain A and B are beginning to grow within these recovering regions and these grains are bounded by HAGBs. Overall, recrystallization of the deformation microstructure in the Al layers appears to be similar to that occurring in the same material after roll bonding to five passes (chapter 5). In contrast, the Al(Sc) layers show no sign of recrystallization and only limited recovery. The misorientation profile parallel to ND for a typical Al(Sc) layer shows that the HAGB spacing is ~0.55 μm (see Fig. 7.16), which is only marginally larger than that observed in the as-deformed state.
The overall texture of the area given in Fig. 7.15, together with the texture of the specific Al and Al(Sc) layers, is shown as a series of 111 pole figures in Fig. 7.17.
The global texture shown in Fig. 7.17a is a distorted !-fibre and the other layers show similar behaviour as the lower annealing temperature. However, substantial grain coarsening has occurred at this higher temperature and Fig. 7.17c shows the orientation distribution of the largest grains given in Fig. 7.13. There is a considerable spread in orientations although it appears that the grains contain
80
Figure 7.15. EBSD micrographs of a typical area of microstructure in SSSS-ARB- HT after annealing for 3 min at 300 °C showing: (a) alternating recovered (Al(Sc) and partly recrystallized (Al) layers highlighting HAGBs; (b) same area as (a) but highlighting LAGBs, and (c) Image Quality (IQ) map revealing the distribution of the various types of boundary (arrows indicate unrecrystallized areas within the recrystallized Al layers).
Figure 7.16. Representative misorientation profiles of the sample given in Fig. 7.15 within: (a) unrecrystallized region of an Al layer, and (b) recovered Al(Sc) layer (parallel to ND).
Figure 7.17. 111 pole figures of the sample given in Figure 7.15 showing the texture of: (a) all layers; (b) Al(Sc) layers only, (c) recrystallized grains in Al layers, and (d) unrecrystallized regions in Al layers.
orientations within the spread of orientations of the deformed layers. Figure 7.18 shows the orientations of the large grains in various Al layers of Fig. 7.15.
Unfortunately, the insufficient number of grains per layer makes it difficult to make firm conclusions about their origin and mechanism of growth within the deformation substructure. Figure 7.19 shows the texture within the individual Al layers where only limited recovery is observed. The textures shown in this figure are comparable to texture development in the Al(Sc) layers after ARB and subsequent annealing at the lower temperature.
Finally, the texture within specific Al(Sc) layers of Fig. 7.15 are shown in Fig. 7.20.
The general observations are similar to both SSSS- and Aged-ARB, thereby indicating that very little structural change within the deformation microstructure is possible during annealing at 300 °C.
7.2.3.3 Microstructure after annealing at 350 C
Annealing at 350 C resulted in significant changes to the Al layers and, again, no major changes to the Al(Sc) layers although some subgrain growth has occurred in the latter. This is demonstrated in Figs 7.21a and 7.21b showing full thickness grains and some recovered regions in the Al layers (see e.g. IQ map in Fig. 7.21c), whereas only static recovery has occurred within the Al(Sc) layers (Fig. 7.21d). Figure 7.22 shows a typical misorientation profile through an Al(Sc) layer parallel to ND % % %5 1 6) 1 773
300 C.
Similar to the previous processing conditions, Fig. 7.23a shows the distribution of orientations within the entire area of Fig. 7.21. Of particular note is the retained deformation texture in the Al(Sc) layers (Fig. 7.23c) and the essentially random texture in the Al layers (Fig. 7.23b), which is different to the retained rolling texture
81
Figure 7.18. 111 pole figures of the sample given in Figure 7.15 showing the orientation distribution of the recrystallized grains in: (a) Al-1 (b), Al-2 and (c) Al- 3.
Figure 7.19. 111 pole figures of the Al layers in Figure 7.15 showing the texture of: (a) unrecrystallized regions only; (b) Al-1 with superimposed Taylor, S3 (]463)[213 and S1 (]121)[214 components; (c) Al-2 with superimposed S3 (]643)[321 component; (d) Al-3 with superimposed S1 (,]643)[321 S2 (]124)[132 and S3 (]346)[132 components, and (e) Al-4 with superimposed Brass component (inset shows main FCC rolling texture components.
Figure 7.20. 111 pole figures of Al(Sc) layers in Figure 7.15 showing the texture of: (a) Al(Sc)-1; (b) Al(Sc)-2 with superimposed S1: ]121)[214( and (]211)[241 components; (d) Al(Sc)-3 with superimposed S1:(]112)[124 and Copper components, and (d) Al(Sc)-4.
Figure 7.21. EBSD micrographs of a typical area of microstructure in SSSS-ARB- HT after annealing for 3 min at 350 °C showing: (a) alternating recovered (Al(Sc) and fully recrystallized (Al) layers highlighting HAGBs; (b) same area as (a) but highlighting LAGBs, and (c-d) IQ maps revealing the distribution of the various types of boundary in the regions shown in (b).
Figure 7.22. Typical misorientation profile within a recovered Al(Sc) layer of the sample given in Figure 7.21 (parallel to ND).
Figure 7.23. 111 pole figures of the sample given in Figure 7.21 showing the texture of: (a) Al and Al(Sc) combined layers; (b) Al layers only, and (c) Al(Sc) layers only (inset shows main FCC rolling texture components).
developed in the Al layers in both SSSS- and Aged-ARB. To generate better statistics on the orientations of the coarse grains within the Al layers after annealing at 350 C, a series of area scans were carried out in the ND-RD section. Figure 7.24 shows the montage 5 % 5 94.: 1% %,) taken at 400 μm intervals and a scan height that covers ~20 consecutive layers in the mid-thickness region. This analysis generated ~130 orientations for the through- thickness grains within the Al layers. The 111 pole figure in Fig. 7.24b shows that such grains have orientations that are randomly distributed. The different texture that develops in the Al layers in SSSS-ARB-HT compared with both SSSS- and Aged-
ARB is discussed in more detail in section 8.3.2. Finally, Fig. 7.25a shows the orientations within all Al(Sc) layers in Fig. 7.21 with Fig. 7.25b-e showing the texture of the individual Al(Sc) layers. Again, the layers are unrecrystallized and possess similar textures as that observed for the previous processing conditions.
Table 7.1 provides a summary of the general evolution of microstructure and texture in SSSS-ARB-HT during various stages of processing.
Table 7.1. Summary of microstructure and texture development within the samples after various stages of processing.
General microstructure of Al and Major texture components in Al and Processing Al(Sc) layers (ND-RD, along ND) Al(Sc) layers (average) conditions Al layers Al(Sc) layers Al layers Al(Sc) layers Complex texture - Close to typical As- subgrain subgrain large spread in FCC -fibre deformed structure structure orientations texture ARB + Complex texture - Close to typical annealing recovered recovered large spread in FCC -fibre at 250 C substructure substructure orientations texture Mixed ARB + structure – Coarse grains, Close to typical annealing coarse grains recovered randomly FCC -fibre at 300 C and recovered substructure distributed texture substructure Coarse ARB + Coarse grains, Close to typical through- annealing recovered randomly FCC -fibre thickness at 350 C substructure distributed texture grains
82
Figure 7.24. A schematic diagram of the layered structure of Al-layers (coloured) and Al(Sc) layers (white) in SSSS-ARB-HT after annealing for 3 min at 350 °C (here, nine EBSD scan areas of 10 m width are superimposed). An orientation data point for each recrystallized grain is given in the corresponding pole figure in (b).
Figure 7.25. 111 pole figures of the Al(Sc) layers in Figure 7.21 showing the texture of: (a) Al(Sc) combined layers with intensity contours; (b) Al(Sc)-2 with superimposed Copper and two S1 components at: ]211)[241( and (;]112)[142 (c) Al(Sc)-3 with superimposed Copper, Brass and S1 (]112)[142 components; (d) Al(Sc)-4 with superimposed S2 (]412)[132 and two Brass components at: ]211)[110( and (]112)[011 , and (e) Al(Sc)-6 with superimposed two S2 at: ]412)[321( and (]421)[231 components (inset shows main FCC rolling texture components).
7.3 Mechanical Properties
The mechanical properties of as-rolled and annealed SSSS-ARB-HT were investigated by tensile testing using ASTM standardized samples (Fig. 4.3). This material was compared with SSSS-ARB following annealing for 6 h at 350 °C and two conventionally rolled Al and Al-0.3%Sc samples. Figure 7.26a shows stress- strain curves for SSSS-ARB-HT after the various stages of processing. It can be seen that the as-rolled material undergoes early plastic instability where there is a peak in stress soon after yielding that decreases with increasing strain to generate a low elongation to failure. With increasing annealing temperature, it can be seen that the extent of uniform elongation is considerably increased, although the maximum tensile stress is lowered.
Figure 7.26b compares the tensile test data of SSSS-ARB-HT with the other materials. It can be seen, for the as-rolled composites, that SSSS-ARB has a higher maximum tensile stress compared with SSSS-ARB-HT as well as a greater uniform elongation before failure. This difference in flow behaviour is likely to be associated with the propensity for flow localization during ARB whereby SSSS-ARB consists of alternating Al and Al(Sc) layers largely parallel to RD and very little evidence of shear banding (Fig. 5.2). In contrast, it can be seen in Fig. 7.1 that ARB-processed
SSSS-ARB-HT shows extensive shear banding similar to that produced in Aged-
ARB (Fig. 6.3). While the microstructure of SSSS-ARB-HT undergoes continual recrystallization during the multi-stage rolling operations, it is unlikely that these structural changes would influence the tensile flow behaviour.
A more plausible explanation of the early onset of plastic instability in SSSS-ARB-
HT during tensile straining (Fig. 7.26b) is the fact that this material already contains numerous microstructural instabilities generated during roll bonding. Hence, tensile straining will simply result in strain localization along the shear bands soon after
83 Figure 7.26. Engineering stress-strain curves for: (a) SSSS-ARB-HT in the as-rolled condition and after annealing for 3 min at 250 to 350 °C; and (b) the same data in (a) including additional data for SSSS-ARB (ARB-32) together with as-rolled conventional Al and Al(Sc) alloys.
yielding thereby generating the flow curves shown in Fig. 7.26b. It is clear from Fig.
7.26b that the tensile flow behaviour of the composite materials falls within the limits of the monolithic Al and Al(Sc) alloys where these materials have the lowest and highest tensile stresses, respectively. Figure 7.27 shows the effect of material type and processing route on both yield stress (0.2% proof stress) and elongation to failure showing that SSSS-ARB-HT has an inferior combination of properties. The influence of annealing on the improvement in elongation to failure in SSSS-ARB-HT is given in Fig. 7.28. Section 8.4 discusses the influence of roll bonding conditions during the production of the Al/Al(Sc) composites described in chapters 5-7 on tensile flow behaviour and other mechanical properties.
7.4 Summary
A high purity Al alloy and a supersaturated Al-0.3wt.% Sc alloy were accumulative roll bonded at 350 °C to generate 0.5 mm gauge sheet consisting of 64 alternating layers of Al and Al(Sc). The deformation microstructure of the Al layers in this material was considerably different to both SSSS-ARB and Aged-ARB. It was found that the higher rolling temperature and the large strain (50%) generated during a given rolling pass was sufficient to considerably alter the microstructure and texture of the Al layers. These layers are expected to undergo extensive dynamic recovery during the first rolling pass and reheating for the second rolling pass resulting in static recovery or partial/complete recrystallization of the Al layers.
These deformation and dynamic/static restoration processes are then repeated in further rolling bonding cycles. These continual changes to the microstructure of the
Al layers during multi-stage rolling at 350 °C also generated a complex texture in
SSSS-ARB-HT compared with the well-developed -fibre texture in both SSSS- and
Aged-ARB.
84 180 Al-Sc 165
150 ARB-32 135
Yield 120 Stress (MPa) 105 90
75 2500C ARB-64 Al 60 3000C 3500C 45 0 5 10 15 20 25 30 35 40 % Elongation Figure 7.27. Yield stress (0.2% proof stress) as a function of total elongation to failure for SSSS-ARB-HT after annealing for 3 min at a given temperature together with data for as-rolled SSSS-ARB (ARB-32) and Al and Al(Sc) alloys.
14
12
10
8
6 % Ductility % 4
2
0 200 250 300 350 400 Temperature (C)
Figure 7.28. Total elongation to failure in SSSS-ARB-HT after annealing for 3 min at 250 to 350 °C.
However, the deformation microstructure of the Al(Sc) layers in SSSS-ARB-HT was similar to that observed in both SSSS- and Aged-ARB. Though, unlike SSSS-ARB, the higher processing temperature is expected to result in precipitation of a fine dispersion of Al3Sc within the Al(Sc) layers after the first few roll bonding cycles, thereby impeding static/dynamic restoration during the subsequent roll bonding operations. This retention of the deformation microstructure throughout the roll bonding schedule resulted in both the generation of the classic FCC rolling texture in these layers.
The microstructural development within the Al layers in SSSS-ARB-HT during annealing at 350 °C was consistent with normal discontinuous recrystallization occurring to generate coarse-grained layers. However, the final texture within these layers was more random compared with the retained deformation texture in both
SSSS- and Aged-ARB due to recrystallization probably occurring in a discontinuous manner. All materials retained the deformation substructure in the Al(Sc) layers due to particle pinning effects and this generated a final material containing alternating coarse-and fine-grained layers.
A notable feature of SSSS-ARB-HT after ARB, and similar to Aged-ARB, was the considerable waviness of the Al/Al(Sc) layers caused by extensive shear banding.
This plastic instability in the as-processed microstructure is the likely cause of the poor flow behaviour during tensile testing whereby plastic instability occurred soon after yielding and resulted in a low strain to failure. However, further annealing resulted in a slight improvement in the ductility of the material. A more detailed discussion of the deformation and annealing behaviour of this material is given in chapter 8.
85 Chapter 8 ______DISCUSSION
8.1 Introduction
Multilayered aluminium alloy composites have been produced by accumulative roll bonding (ARB) to generate a thin-gauge sheet product comprising alternating layers of commercial purity Al and Al-0.3% Sc alloy. Chapters 5 to 7 described the development of microstructure and texture of these multilayered composite materials whereby both the control of the starting microstructure of the Al(Sc) alloy and roll bonding temperature had a significant influence on the development of microstructure and texture during both deformation and subsequent annealing and the general integrity of the multilayered structures.
Chapter 5 demonstrated that roll bonding of Al(Sc) in the supersaturated condition with Al resulted in uniform deformation and generated reasonably parallel Al/Al(Sc) layers. In contrast, chapter 6 demonstrated that artificial aging of the Al(Sc) sheet prior to ARB had a marked influence on the integrity of the layers during rolling whereby large-scale shear banding caused marked buckling of the layered
86 microstructure. This is seen clearly in the ICC micrographs in Fig. 8.1 showing parallel layers of material in SSSS-ARB and highly distorted layers in Aged-ARB.
Finally, chapter 7 explored the roll bonding behaviour of SSSS-ARB at a higher temperature (350 °C) and higher strain (6 rolling cycles to produce 64 layers). It was shown that roll bonding at this temperature resulted in repeated recrystallization of the Al layers during each reheating and rolling stage and generated a different final deformation microstructure and texture in the Al layers, since ARB of both materials at 200 °C resulted in moderate recovery in the Al layers without recrystallization.
Scandium, either in solution or in the form of a dispersion of fine Al3Sc particles in the Al(Sc) layers, significantly retarded static restoration during subsequent annealing whereby, for all materials, the deformation substructure was retained in the
Al(Sc) layers whereas the higher purity Al layers were rapidly replaced by coarse grains at the scale of the layer thickness. This ultimately generated alternating layers of coarse grains (Al) and a partially recovered substructure (Al(Sc)). Furthermore, the general shape of the layers produced by ARB was inherited in the annealed materials; this was due to substantial pinning of the deformation substructure in the
Al(Sc) layers by both solute and particle pinning effects. Table 8.1 provides a general summary of the development of microstructure and texture of the three materials after various stages of processing.
The following sections discuss the general evolution of the microstructure and texture of these materials after roll bonding and during subsequent annealing with particular emphasis on the fundamental nature of shear band formation, texture development during processing and particle-pinning effects for generating the unique multilayered microstructures shown in Fig. 8.1. Finally, the effect of roll bonding and annealing on the microstructure and the mechanical behaviour of SSSS-ARB-HT is discussed in detail in sections 8.3.2 and 8.4.
87 (a) (b)
Figure 8.1. ICC micrographs showing the comparison of the lamellar structures generated in (a) SSSS-ARB and (b) Aged-ARB due to the different initial heat treatments. Table 8.1: Summary of microstructural development during ARB and subsequent annealing.
SSSS-ARB Aged-ARB SSSS-ARB-HT Rolling conditions Roll bonding temperature 200 °C 200 °C 350 °C ARB cycles (layers) 5 (32) 5 (32) 6 (64) True strain 4.15 4.20 4.95 Microstructures prior to ARB Al layers Coarse grain size. Coarse grain size. Coarse grain size. Coarse grain size. Coarse grain size. Coarse grain size. Al(Sc) layers Supersaturated solid Aged to generate a Supersaturated solid solution. uniform dispersion of solution. fine Al3Sc particles. Microstructures after ARB Parallel, alternating Considerable large- Moderate degree of layers of Al and Al(Sc). scale shear banding shear banding General Some shear banding. through many layers through many layers. resulting in major perturbation. Severely strained, Severely strained, Evidence of recovery containing LBs containing LBs and recrystallization (recovery evident but (recovery evident but during each ARB cycle no recrystallization). no recrystallization). – final microstructure Al layers resembling typical low strain deformation substructure. Severely strained, Severely strained, Severely strained, containing very fine containing very fine containing very fine LBs. Very little LBs. Very little LBs. Very little recovery evident. recovery evident. recovery evident. Al(Sc) layers No Al3Sc particles Particles present. Particles present observed. indicating precipitation during ARB. Microstructures after annealing at 350°C Parallel, alternating Major perturbations Moderate layers of Al and Al(Sc) of the layers inherited perturbations of the still evident. from the as-rolled layers inherited from General Well-defined state. the as-rolled state. boundary between Al Well-defined Well-defined and Al(Sc) layers. boundary between Al boundary between Al and Al(Sc) layers. and Al(Sc) layers. Rapid recrystallization Rapid recrystallization Rapid Al layers to full layer thickness. to full layer thickness. recrystallization to Retained rolling Retained rolling full layer thickness. texture texture. Randomized texture. Moderate recovery Moderate recovery Moderate recovery after extended after extended after extended Al(Sc) layers annealing. Al3Sc annealing. Al3Sc annealing. Al3Sc particle pinning of particle pinning of particle pinning of substructure. substructure. substructure.
88 8.2 Deformation and Stability of Multilayered Structures
8.2.1 Microstructural development in the Al and Al(Sc) layers during ARB
The notable microstructural features observed in the layers of both SSSS- and Aged-
ARB are lamellar bands (LBs), shear bands (SBs) and deformation bands (DBs); these are high strain structural elements typical of medium to high stacking fault energy (SFE) metals such as Cu (Duggan et al., 1993), LC steel (Ueji et al., 2004),
ARB-Ti (Terada et al., 2007), Al-Mg alloys (Hurley and Humphreys, 2003), Ni
(Hughes and Hansen, 2000), and Al (Humphreys and Hatherly, 2004). It is pertinent to note that SBs and DBs were observed only in the Al(Sc) layers and these were observed in the Al/Al(Sc) composite deformed at both 200 °C (chapter 6) and 350 °C
(chapter 7).
The formation of SBs is common at intermediate to high strains ( > 0.7), and they often form in the harder layers of an alternating hard and soft layered structure. This was demonstrated by Yeung and Duggan (Yeung and Duggan, 1987) in a layered
Al/Cu composite, where flow instability becomes inevitable in the Cu layers to match the strain generated in the softer Al layers. The large initial grain size in the
Al(Sc) sheet also promoted the formation of shear bands and deformation bands.
Shear bands were generally confined within the Al(Sc)-layers in SSSS-ARB, whereas they extended through several Al and Al(Sc) layers in Aged-ARB and
SSSS-ARB-HT.
Over the past few decades, several investigators have established a framework for the evolution of microstructures during cold deformation (Hughes and Hansen, 1997;
Winther et al., 2000; Hansen et al., 2001). This framework was based on the subdivision of grains by deformation-induced dislocation boundaries which at medium strain levels are separated into two groups: (i) geometrically necessary boundaries (GNBs) separating crystallites that deform by different selections of slip
89 systems and/or different strain and (ii) incidental dislocation boundaries (IDBs) formed by the trapping of glide dislocations (Bay et al., 1989; Hughes and Hansen,
1991; Bay et al., 1992). These workers concluded that, with increasing strain, the misorientation angle across the two types of boundary increases and, accordingly, the spacing between boundaries, decreases.
At larger strain levels this evolution leads to structures composed of dislocation boundaries having a wide range of misorientations and spacings in the submicron range that, by further increase in strain, there is a dominant tendency for the dislocation boundaries to reorient from a typical cell block structure into a structure termed a lamellar structure. In the typical cell block structure, the GNBs include microbands (MBs) and single dense dislocation walls (DDWs) that surround blocks of equiaxed cells. In a typical lamellar structure at large strain, the lamellar boundaries (LBs) sandwich thin layers of cells and subgrains oriented along the material flow direction, that is, the structural evolution brings the cell block boundaries into a macroscopic inclination almost parallel to the rolling plane forming
LBs that, on average, consist of only one row of cells.
The main microstructural similarity between the Al and Al(Sc) layers in both SSSS- and Aged-ARB is that the LBs are separated by DDWs (Figs 5.4-5.8 and Figs 6.4-
6.6). This type of structure is found in many metals deformed to true strains greater than ~3 including rolled Ni (Hughes and Hansen, 2000), Al (Jazaeri and Humphreys,
2004a) and Al alloys produced by ARB (Kamikawa et al., 2006). However, there are two notable differences in the structure of LBs in Al and Al(Sc) in both SSSS- and
Aged-ARB: (i) the degree of refinement whereby the LBs in Al are greater than those in Al(Sc), and (ii) the misorientation between contiguous volumes of LBs, which is generally high in Al and low in Al(Sc) (Figs 5.6-5.15 and Figs 6.5-6.11).
The refining effect is expected to arise during the formation stage, governed by the
90 dislocation generation rate and mobility, and/or due to differential dynamic recovery in the Al and Al(Sc) layers. In alloys in the supersaturated condition, mobile dislocations can be pinned by solute atoms with the extent of pinning dependent on solute concentration and the size ratio between the solute and solvent atoms
(Zakharov, 1997). Although the difference in atomic size between Al and Sc is small, workers have shown that the stress-strain curve in an Al-4Mg-0.4Sc alloy was serrated, which was attributed to solute drag on gliding dislocations (Zakharov, 1997;
Woo et al., 2003). Likewise, scandium in solution may contribute to enhanced work hardening in the Al(Sc) layers. Hence, a large portion of hardening within the Al(Sc) layers in SSSS-ARB (as demonstrated in Fig. 6.2) may arise due to the presence of solute atoms, as discussed below.
Fleischer (1964) derived the following relationship for the critical stress required to free a pinned dislocation from a solute atmosphere:
W bc )( 2/1 (8.1) crit ab 0 where W is the energy required to displace a dislocation through a region of solute of concentration c0 , and b is the Burgers vector of the dislocation.
The mechanism is similar to that of the movement of jogs and, therefore, a is the distance a jog moves to reach the top of a barrier. By assuming a ~ b , and noting that / bW is at least an order of magnitude smaller than the line tension of a dislocation, ~Gb 2 , where G is the shear modulus (Hirth and Lothe, 1982), / bW is
2 2/1 related to Gb by a factor B. The short range interaction bc0 )( is the spacing between the pinning points of a dislocation, that is, the distance, l , between two solute atoms. Hence, Eq. (8.1) can be rewritten as:
91 Gb (8.2) crit Bl
Equation (8.2) is of the same form as Taylor’s work hardening relation (Hirth and
Lothe, 1982):
Gb (8.3) wh 2 r where r is the separation distance between dislocations.
The differences between Eqs. (8.2) and (8.3) are the values of coefficients and the substitution of solute spacing for dislocation separation. This implies that the pinning force induced by solute atoms is significant enough to hinder dislocation movement in the matrix. As a consequence, additional dislocations are required to maintain deformation. The restricted movement of dislocations also results in a refined substructure. It is relevant to note that a high dislocation density does not necessarily mean the generation of densely spaced dislocation walls of LBs.
The stages of wall formation from a low strain cell block structure have been reviewed recently (Hansen and Jensen, 1999; Hughes and Hansen, 2000; Jazaeri and
Humphreys, 2004a). The degree of dynamic and post-dynamic recovery also determines the extent of structural refinement of a deformed metal. Recovery in Al occurs readily due to its high SFE (Humphreys and Hatherly, 2004), but a possible reduction in SFE due to scandium will reduce the rate of recovery. While such an effect of solute atoms on SFE is well-documented for Cu, the effect of scandium in
Al remains unresolved.
The influence of scandium in solution on structural refinement in the Al(Sc) layers is comparable to that found in other work. For example, the mean separation of LBs in
92 high purity aluminium after six-cycle ARB was ~0.8 m (Kamikawa et al., 2006) which is at least three times thicker than that measured in solute-containing Al sheets i.e. AA 3XXX, AA5XXX and AA8XXX Al alloys (Slamova et al., 2005; Slamova et al., 2006; Worswick et al., 2006; Cieslar et al., 2007; Halim et al., 2007; Kang et al.,
2007; Pirgazi et al., 2008b).
In the roll-bonded materials it was found that the area fraction of HAGBs in the
Al(Sc) layers was lower than in the Al layers for both SSSS- and Aged ARB. This is in general agreement with the work of Apps et al. (2005) where it was shown that Al-
Sc alloys are highly resistant to fragmentation because of the strain homogenizing capability of the fine Al3Sc particle dispersion in the matrix. However, the Al(Sc) sheets in SSSS-ARB were deformed in the supersaturated condition with precipitation unlikely happening since the cumulative heating period of all ARB passes at 200 °C was only 25 min which, according to Seidman et al. (2002), is too short a time for precipitation in the Al-Sc system. Therefore, solute Sc presumably also plays a vital role in fragmentation of the substructure.
8.2.2 Effect of scandium on differential hardening of the Al and Al(Sc) layers in SSSS- and Aged-ARB
The overall effect of scandium on hardening during ARB for both SSSS- and Aged-
ARB was shown in Fig. 6.2. The hardening response of the Al layers in both materials during ARB was similar: there is a plateau in hardness (~44 HVN) after the second cycle ( = 1.38) resulting in a hardness increase of ~14 HVN. The hardness of the Al layers after ARB is lower than values found in the literature (Park et al.,
2004; Valiev and Langdon, 2006). This is likely to be a result of dynamic recovery during rolling and static recovery during reheating to a reasonably high homologous temperature ( TT m 5.0~/ ) between ARB cycles.
93 The hardness of the Al(Sc) layers reached a plateau of ~68 and ~88 HVN for SSSS- and Aged-ARB, respectively. Hardening of Al(Sc) in the supersaturated condition is a result of interactions of solute Sc atoms with dislocations; this also enhances structural refinement and suppresses recovery, as discussed in section 8.2.1. In the artificially aged condition, Al(Sc) exhibits a high starting hardness (78 HVN), due to the generation of nanosized, coherent Al3Sc dispersoids, as shown in Fig. 6.1 and by other workers (Jazaeri and Humphreys, 2004a). Following ARB, the moderate increase in the hardness of the Al(Sc) layers (D'()*, % % of Prangnell et al. (2004) on a similar alloy after ECAP where two single-phased Al- alloys (Al-0.13%Mg and Al-0.2%Sc) were studied. These researchers argued that a starting microstructure containing a dispersion of fine Al3Sc dispersoids results in a homogenized microstructure after ECAP.
Using Fig. 6.2, initial hardening rate of Al, Aged Al(Sc) and SSSS Al(Sc) layers was computed to be 0.66, 0.32 and 0.90, respectively (measured from the slope of initial rising part of curves). The high hardening rate of the Al(Sc) layers in SSSS-ARB may be explained by solute-dislocation interactions, as discussed elsewhere
(Seidman et al., 2002; Apps et al., 2005; Ferry et al., 2005). However, from the description given by Prangnell et al. (2004), there should not be pronounced hardening in the particle-containing Al(Sc) layers of Aged-ARB. Then again, the following factors are still a potential source of some hardening in these aged alloys.
The main one is the distribution of particles, which is not always homogenous, leaving particle-free pockets of material either as low-solute Al or Al containing scandium in solution. Such regions undergo deformation but are not greatly influenced by the nano-sized particles. However, since artificial ageing of Al-0.3%Sc generates a dispersion of closely-spaced Al3Sc dispersoids (see e.g. Fig. 6.1), this effect on hardening should be minor.
94 8.2.3 Origin of large-scale shear banding in Aged-ARB during ARB
The formation of large-scale shear bands in Aged-ARB compared with SSSS-ARB
(Fig. 8.1) appears to be a result of sequential necking in the harder Al(Sc) layers and cooperative flow in the softer Al layers. This implies that shear bands form in the
Al(Sc) layers, although their evolution is attributed to the larger difference in flow properties of the Al and Al(Sc) layers (as indicated by the hardness results of Fig.
6.2), since shear banding was negligible in SSSS-ARB. The origin of shear banding may be explained as a response to flow instability that appears in single-phase alloys as the onset of necking during tensile testing, that is, when (Hutchinson, 1974):
d (8.4) d where and are true stress and true strain.
The same is applicable for plane strain condition (e.g. rolling) that can be re-arranged and given as:
1 d 0 (8.5) d
The important parameters influencing the instability condition were incorporated in a single equation by Dillamore et al. (1979) and outlined in a crystallographic model showing the coincidence of the angle of shear banding to the direction of geometrical softening. These workers expressed such an instability condition as:
1 d 1 $$ mnmn d# dm $$ 0 (8.6) d # d ! d where n and m are the strain hardening and strain rate exponents, ! the dislocation density, and MT the Taylor factor defined as "s"% "s is the total shear strain on
95 all active slip systems and is the normal strain. In their analysis, various terms are disregarded except for Taylor hardening and, therefore, it is exclusively applicable to shear banding of a crystallographic nature (Dillamore et al., 1979; Van Houtte, 1979;
Yeung and Duggan, 1987).
It is well-known that a layered microstructure can result in shear banding due to the difference in flow properties of neighbouring layers, i.e. this is a prerequisite feature for triggering the shear banding process, as reported for layered polymers (Crosky et al., 2006), ceramics (termed ‘sausaging’) (Zeimetz et al., 1995) and clad metals
(Yeung and Duggan, 1987). In such situations, this behaviour may be regarded as a necking-related fabrication issue. For a layered composite containing both hard and soft layers, the condition of instability for shear banding has been given as (Yeung and Duggan, 1987):
d d 1 2 (8.7) d d 21
where 1 and 2 is the plane strain flow stress of the hard and soft layers, respectively.
Using Eq. (8.7), a plot of both Al Sc)( Al // dddd and ( Al Sc)( Al ) as a function of true strain, , is expected to highlight the critical strain for instability by the intersection of the curves. Figure 8.2 shows these plots for both SSSS- and Aged-
ARB, respectively. The instantaneous values of were calculated by multiplying the average hardness of the Al and Al(Sc) layers after each ARB cycle using the following conversion factors: tensile strength MPa VHNH )(807.9)( (Dieter,
1986) and plane strain flow stress 3/2 (Rowe, 1968). According to Fig.
8.2a, plastic instability in the SSSS-ARB sheet is expected to occur at = 0.6.
However, shear banding was minimal even at = 4.15 in this material. Most
96 Figure 8.2. Plot of ScAl )( Al // dddd and ( ScAl )( Al ) as a function of true strain for both the Al and Al(Sc) layers in (a) SSSS-ARB and (b) Aged-ARB. importantly, Fig. 8.2b shows that Aged-ARB should be stable against shear banding, since there is no point of intersection on the graph. This behaviour is understandable from the hardness-strain curves in Fig. 6.2 showing a minor hardening rate of both the Aged Al(Sc) and Al layers. Therefore, Eq. (8.7) fails to identify the onset of plastic instability of a multilayered alloy composite composed of materials of comparable hardening rates.
Rolling a layered composite (consisting of hard and soft layers) creates strain incompatibility at the layer interfaces during deformation. This is because the rolling mill induces the same magnitude of load to both the layers separated by a bonded interface, across which the material in the softer layer flows to a larger strain than the harder layer. Thus, the soft layer pulls the harder layer resulting in an in-plane shear force at the interface (schematically shown in Fig. 8.3a). The magnitude of this stress ( Shear ) might play an important role in explaining the onset of shear banding if it exceeds the yield strength in pure shear ( y Al Sc)( 3/ ), where y Al Sc)( is the yield stress in uniaxial tension of the harder layers, that is:
y Al Sc)( shear (8.8) 3
Since shear banding is a generic phenomenon in layered structures of a thickness length scale down to a few microns, it is pertinent to note that measuring the flow properties of individual layers is often difficult because of the fineness of their features. For instance, the flow properties of 0.5-2μm thick individual layers of microbands in SFE alloys (Quadir and Duggan, 2006) and 4-10 μm wide twin bands in low SFE alloys (Duggan et al., 1978; Hirsch et al., 1988) involve large errors, although shear banding is common in them. In regard to the current composite materials, it is possible to calculate shear using the equations developed by Semiatin and Piehler (1979); these are applicable to frictionless plane strain deformation
97 (assuming perfect bonding at the interfaces) of layered clad structures. Hence, the yield condition of the Al(Sc) layers using Von-Mises yield criterion is (Semiatin and
Piehler, 1979):
$ shear Al Sc)( (8.9)
where Al Sc)( is the flow stress of an individual Al(Sc) layer and the flow stress of the entire composite, as calculated using the following rule of mixtures (Semiatin and
Piehler, 1979):
$ AA AlAlAl sc Al Sc)()( (8.10)
where AAl , AAl Sc)( , Al and Al Sc)( are cross-sectional areas and flow strengths of the relevant Al and Al(Sc) layers, respectively.
These data can be extracted from the hardness values (using the previous conversion factors) and the instantaneous thickness of the layers after each ARB cycle. Figure 8.3b shows a plot of the in-plane shear stress, shear as a function of the hardness ratio of the Al and Al(Sc) layers at a given strain (after each ARB cycle). It can be seen in Fig. 8.3b that the acting in-plane shear stress ( shear ) on the harder Al(Sc) layers increases with increasing hardness ratio and the lower and upper data sets are for SSSS- and Aged-ARB, respectively. It is interesting to note that the tensile strength of an artificially aged Al-0.3%Sc alloy in pure shear is 138 MPa (Toropova et al., 1998); this value is far exceeded in the data set of Aged-ARB (Fig. 8.3b).
However, the data set of SSSS-ARB does not extend far beyond the value measured in the SSSS condition (Polmear, 2006). The data point at the top of this plot is from the single cycle ARB sample. This implies that shear banding in Aged-ARB commences very early in the process. The information shown in Fig. 8.3b provides a basis for understanding why shear localization is profuse only in the Aged-ARB.
98 Figure 8.3. (a) Schematic diagram showing the operation of an in-plane shear stress at the boundary between the Al and Al(Sc) layers in Aged-ARB, thereby creating a tensile drag force on Al(Sc). (b) Magnitude of the in-plane shear stress as a function of the hardness ratio of Al and Al(Sc) layers in both SSSS-ARB and Aged-ARB. The foregoing analysis might also be useful in the choice of candidate materials suitable for ARB, and where shear localization is to be avoided. It may also be possible to extend the analysis to quantitatively predict the extent of instability from the magnitude of stress beyond the critical value. The present analysis provides a conceptual understanding of the correlation between the acting stress (at the interface) and mode of deformation; this behaviour also clearly affects the subsequent recrystallization textures that form (Fig. 6.16).
8.2.4 Texture development of the Al and Al(Sc) layers during ARB
8.2.4.1 Absence of large-scale shear banding (SSSS-ARB)
The rolling texture of the Al layers in SSSS-ARB, shown by the 111 pole figure in
Fig. 5.9 and the ODF in Fig. 5.10, is typical of that developed in heavily cold rolled aluminium alloys (Grewen and Huber, 1978). It contains texture components distributed along the -fibre from {112}<111> (Copper) to {110}<112> (Brass) through {123}<634> (S3). As observed for the present material, and discussed by
(Hirsch and Lucke, 1988a,b), the {110}<001> (Goss) component was not observed as this component is only common at low to moderate rolling strains and is not sustained after high strains.
The rolling texture of the Al layers of the SSSS-ARB is similar to that observed by other investigators of ARB of Al alloys (Tsuji, 2006). It is also clear that, while the rolling situation in ARB is different to conventional rolling, similar rolling textures are generated. Nevertheless, rolling in the absence of lubrication can generate a shear texture in the subsurface region of ARB sheet with the core of the sheet generating a conventional rolling texture (Kamikawa et al., 2006).
99 8.2.4.2 Effect of large-scale shear banding (Aged-ARB)
The orientation distribution in as-rolled Aged-ARB was found to be a distorted - fibre with variations in peak intensities and extension along a new fibre from B to
RC (Fig. 6.8). This can be seen clearly from the intensity plot across the B and Cu orientations in Fig. 8.4. Across B, plots along =0-90 and 1=0-90 , in 2 =0 are shown in Figs 8.4a and 8.4b, respectively. Compared with SSSS-ARB, the heavily sheared Aged-ARB material showed a large intensity reduction in B from 9.8 to 3.6 random and some increment in the G component and along the fibre to RC. The spread of orientations from Cu cannot be seen in a single 2 section, although a spread from 2=45 to 2=30 (SSSS-ARB) and 2=15 (Aged-ARB) is clear in Fig.
8.4c and 8.4d, respectively. Therefore, for clarity 1=90 sections are shown in the insets of these figures. An intensity plot along O-P and M-N (across Cu) shows a reduction in intensity from 16 random in the SSSS-ARB to 12 random in the
Aged-ARB. By comparing Fig. 8.4a-b with 8.4c-d, the orientation spread of Cu is also less than Brass.
During single-pass, lubrication-free rolling of metals to high strain (>50%), the considerable friction between the sheet and rolls generates a texture gradient within a narrow region near both surfaces of the as-rolled sheet (Holscher et al., 1994; Huh et al., 1998). Due to a combination of complex shearing and bonding damage at the sheet surface of the present materials, the texture of the first two outermost Al layers
(plus two Al(Sc) layers) were discarded during the analysis of overall texture development. This is justified since a parallel investigation of ARB textures of Al showed that the first four layers after roll bonding contain a substantial shear gradient
(as measured by the deflection of an embedded through-thickness scratch) (Lee et al., 2002b; Kamikawa et al., 2004; Kamikawa et al., 2007a). In contrast, there was only minor shearing in the core of the as-processed material (excluding the four layers adjacent to each surface after ARB).
100 Figure 8.4. Intensity distribution across the B and Cu orientations, in SSSS-ARB and Aged-ARB, respectively. The intensity distribution of B is shown in the 2=0 section (along (a) 1 and (b) angles) and Cu in the 1=90 section (along (c) 2 and (d) angles). The shear strain distribution during ARB was also measured in recent studies by observing the deflections of a straight pin introduced before rolling (Lee et al.,
2002b). A parabolic through-thickness strain distribution was measured after the first cycle, which was eliminated following several further cycles. Such an effect on microstructural development was also monitored by studying the cell size and misorientation distribution within the deformation substructure; this was found to remain in gradient up to a few ARB cycles, but was not sustained thereafter. This phenomenon was explained by the complexity of ARB processing whereby cutting, stacking and re-bonding alternates in a cyclic manner. Thus, one surface of every layer comes in contact with the rolls and, thereafter, becomes embedded between layers in the next cycle. The through-thickness shear effect on texture development during ARB has not been thoroughly investigated for Al alloys. An additional complexity arising in the present work is the hard-soft combinations of material, which induces additional in-plane shear, as described in section 8.2.3. This effect might also influence texture development during ARB.
The spread in orientations from B to RC in Aged-ARB, as shown in the 2=0 section of Euler space in Fig. 6.8 is a well-known shear deformation phenomenon
(Morii et al., 1985; Engler et al., 2001). However, other shear components
({111}<112>, {111}<110> and {110}<100>) common in rolled FCC metals are not strongly present ( Choi and Lee, 1997; Choi et al., 1997; Huh et al., 1998; Lee et al.,
2002a,b). The {100}<110> shear component was also predicted by simulation in two distinct ways:
(i) using the relaxed Taylor framework (Honneff and Mecking, 1981), and
(ii) exchanging of the rolling and shear textures between the FCC and BCC systems
governed by slip systems (Holscher et al., 1994).
101 The primary effect of producing shear texture components is to reduce the intensities of the standard rolling components and, in combination with shear banding, the resultant effect is to produce a distorted low intensity distribution in the 2 sections of Euler space (c.f. Fig. 6.8). A 15 deviation in orientations is illustrated in Fig.
6.10 in Aged-ARB, whereas a larger magnitude of up to 20 is also evident in individual lamellar bands (Fig. 6.11). A comparable magnitude of the crystallographic rotation and turning of the lamellar band boundaries in the sheared region can readily be attributed to a rigid body rotation (Rowe, 1968). This rotation occurred in both directions about TD. Figure 8.5 show 111 pole figures of as-rolled
SSSS- and Aged-ARB with Fig. 8.5c showing an intensity plot along A-B, that is, across the common <111> pole of all main rolling components. From Fig. 8.5c it is also clear that the intensity is reduced in the heavily shear banded Aged-ARB, while maintaining the distribution symmetric about <111>, which is relevant to the fact that shear bands are inclined in both positive and negative angles to RD.
Considering the B and Cu orientations, extensive shear banding in Aged-ARB generated a lower reduction in intensity of the Cu texture component as compared with B (cf. Fig. 8.4). For Al-Mn (3XXX) alloys, high rolling reductions results a large spread about the Brass texture component, as reported by (Engler et al., 1996) and confirmed by (Todayama and Inagaki, 2005). However, the Cu orientation was found to be relatively more stable (Kocks and Chandra, 1982), particularly when shear banding was common. Shear banding (Wagner et al., 1995) is argued to account for the stability of this component due to the rotation of the TD-rotated Cu components (e.g. Taylor: {4 4 11}<11 11 8>) back to Cu. Examples of TD rotations between the Cu and Taylor components are shown in Fig. 6.11. For the Brass component, such a phenomenon does not occur and, hence, this component experiences a greater reduction in intensity, as shown in Figs 8.4a-b. A large rotation was also noted for the Brass matrix orientation in Fig. 6.11. Oscillation around stable
102 Figure 8.5. 111 pole figure of as-rolled (a) SSSS-ARB and (b) Aged-ARB. (c) Intensity plot (× random) along AB in (a) and (b) showing a symmetric reduction in texture intensity due to shear banding. end orientations (such as the Cu component in aluminium), was recently also demonstrated in steel by (Shen and Duggan, 2007), and explained by the notion described as “geometric orientation stability” by (Kocks and Chandra, 1982).
8.3 Annealing of As-Deformed Al/(Al(Sc) Composites
In chapters 5-7, it was demonstrated that annealing of the as-deformed Al/Al(Sc) multilayered composite, in its various forms, resulted in rapid grain coarsening in the
Al layers and partial recovery in the Al(Sc) layers. The latter can be attributed to
Zener pinning effects due to nanosized Al3Sc dispersoids that were either present during rolling (Aged-ARB) or by precipitating out into the recovering substructure during subsequent annealing (SSSS-ARB and SSSS-ARB-HT). Ultimately, annealing generated a material consisting of alternating layers of coarse- grained/recovered material, as shown in Figs 5.26, 6.15 and 7.2.
8.3.1 Annealing behaviour following ARB at low homologous temperature
At low homologous temperatures (200 °C), multi-pass ARB coupled with reheating between passes did not result in major static/dynamic restoration of the Al and al(Sc) layers in either SSSS- (section 5.2.2) or Aged-ARB (section 6.3.2). Nevertheless, the elevated rolling temperature did result in considerable softening of the Al layers compared with alloys roll-bonded at lower temperatures due to a combination of static and dynamic recovery. The Al(Sc) layers were highly resistant to recovery and generated a considerably finer microstructure compared with the Al layers (section
8.2.1).
During subsequent annealing of both SSSS- and Aged-ARB, the Al layers were rapidly transformed from a fine-grained deformation substructure to a coarse-grained structure that, by first sight, resembles normal discontinuous recrystallization
103 (Humphreys and Hatherly, 2004). However, the general coarsening behaviour of the deformation substructure in conjunction with texture evolution in the Al layers, as shown in Figs 5.16-5.24 and Figs 6.12-6.17, generated a retained rolling texture although the Brass component was largely eliminated. The retained rolling texture was less developed in Aged-ARB, as shown, for comparative purposes, in Fig. 8.6.
Despite the larger spread in texture in Aged-ARB, the results indicate that normal discontinuous recrystallization is not solely operative but, rather, the Al layers are undergoing continuous recrystallization (Humphreys and Hatherly, 2004).
8.3.1.1 Continuous and discontinuous recrystallization of the Al layers
In contrast to the Al(Sc) layers, the Al layers were unstable during annealing with two major changes occurring: (i) the lamellar structure rapidly collapses into an array of spheroidized grains (up to 5 m in diameter) separated by LAGBs, and (ii) the misorientation profile in a given band changes from a gradual to sharp configuration over a repeated distance of 3-5 m (parallel to RD) resulting in an increase in misorientation spread (Fig. 8.7). Prangnell et al. (2004) showed, by both experimental investigations on severely deformed Al-Mg alloys and Monte Carlo-
Potts simulations, that changes in the microstructure during continuous recrystallization depend largely on the newly-formed boundary spacing of lamellar bands. If the boundary spacing is low (low aspect ratio blocks), then these blocks readily turn into small spheroids. However, for a very large spacing (high aspect ratio blocks), surface tension leads to a slow rate of spheroidization and the lamellar structure persists.
It is known that the microstructural changes that occur when a highly deformed material undergoes continuous recrystallization are rather small, and occur incrementally as the annealing temperature is increased (Humphreys and Hatherly, 2004). Here, it is not expected to be any significant change in the fraction of HAGBs, but the deformed
104 Figure 8.6. Distribution of orientations of the coarse grains within the Al layers after annealing for 6h at 350 C: (a) key -fibre components; (b) SSSS-ARB; (c) Aged-ARB, and (d) SSSS-ARB-HT. Figure 8.7. Misorientation profile along a lamellar band formed in the Al layer (SSSS-ARB) shown in Fig. 5.11 and along a recovered lamellar band (after annealing for 3 min at 350 C) comprising freshly formed LAGB’s, as indicated in the ICC micrograph of Fig. 5.20a. grains/subgrains become more equiaxed and slightly larger. Figure 8.8 shows a series of schematic diagrams illustrating the effect of annealing on a highly deformed microstructure comprising of lamellar HAGBs, together with intersecting boundaries which are mainly LAGBs. On annealing of the microstructure shown in Fig. 8.8a, the energy of the system may be lowered by localized boundary migration occurring in two stages: (i) collapse of the lamellar microstructure (Figs 8.8b and 8.8c) – this is due to the surface tension at the node points such as A in Fig. 8.8a, where the boundaries (of energy R) aligned in the rolling plane, are pulled by the boundaries (of energy N) aligned in the normal direction, as shown in Fig. 8.8b.
The process eventually generates an equiaxed microstructure by (ii) spheroidization and growth – when nodes A and A' touch, node switching will occur, and two new nodes
A1 and A2 will form and be pulled apart by the boundary tensions as shown in Fig. 8.8c.
Further spheroidization and growth occurs due to boundary tensions, as shown in Fig.
8.8d, leading to a more equiaxed grain structure. The initial coarsening of the
SMG/UFG microstructure in the Al layers is consistent with this mode of microstructural evolution up to the stage where rapid growth of certain grains occurs to consume the fine-grained matrix.
Jazaeri and Humphreys (2004a,b) showed that either continuous or discontinuous recrystallization occurs depending on the relative fraction of high and low angle boundaries in the deformation microstructure, with continuous recrystallization prevailing for HAGB fractions greater than 0.7; this annealing process also retains the rolling texture (Prangnell et al., 2004). Figure 8.9 shows the results of Jazaeri and Humphreys (2004a,b) for heavily cold rolled AA8006 Al showing the effect of both strain and starting grain size on the transition from discontinuous to continuous recrystallization. It can be seen that small starting grain sizes and large strains promote continuous recrystallization.
105 (a) (b)
(c) (d)
Figure 8.8. Schematic diagram showing continuous recrystallization of a highly deformed lamellar microstructure: (a) Initial structure; (b) collapse of the lamellar boundaries; (c) spheroidization beginning by Y-junction migration, and (d) further spheroidization and growth, after Humphreys and Hatherly (2004). Figure 8.9. The effects of initial grain size and strain on the transition from discontinuous to continuous recrystallization in AA8006, after Jazaeri and Humphreys (2002). It is pertinent to note that the strain generated in the Al layers of the present materials by ARB ( = 4.2) is on the borderline of the transition from continuous to discontinuous grain coarsening in heavily cold rolled commercial purity aluminium
(Jazaeri and Humphreys, 2004b). Although the deformation microstructure in the Al layers of the present material contained a low fraction of LAGBs (Fig. 8.10a) there was a significant increase in the fraction of these boundaries in the early stages of annealing (Fig. 8.10b). This rapid change in boundary character is probably sufficient to eventually induce some discontinuous recrystallization in the Al layers.
A very similar misorientation distribution was also shown by Prangnell and co- workers (Zahid et al., 2009) in Al-Mg alloy after ECAP and recovery. In their work, recovery at 100 °C generated a similar misorientation distribution as in the rolled
SSSS-ARB (Fig. 8.10a); this is rational since dynamic recovery is inevitable during rolling at 200 °C. In the recovery annealed sample for 3 min at 350 °C (Fig. 8.10b) an increment in LAGB is clearly obvious and this is similar to the distribution shown by Zahid et al. (2009) after annealing at 220 °C.
8.3.1.2 Elimination of Brass texture component in the Al layers during annealing
While the Brass texture component dominates the deformation microstructure within the Al layers in both SSSS- and Aged-ARB (cf. Figs 5.10-5.12 and 6.8), there was a gradual elimination of this component during annealing (Fig. 8.6b), which is particularly evident in SSSS-ARB. It is known that discontinuous recrystallization can result in the preferred growth of grains of minor texture components, and these components dominate the final texture (Humphreys and Hatherly, 2004). The retained deformation texture in the annealed Al layers indicates that continuous recrystallization is occurring, but the eventual rapid coarsening of certain grains within the deformation substructure of the Al layers (Fig. 5.21) and elimination of the Brass component indicates that discontinuous recrystallization is also important.
106 Figure 8.10. Misorientation distribution (5-63 ) in an Al layer of SSSS-ARB after: (a) ARB deformation processing and (b) annealing for 3 min at 350 C showing an increase in the fraction of 5-25 misoriented boundaries. The orientation-dependent stored energy of the microstructure arguably plays a primary role in the orientation selectivity to generate the textures shown in Fig. 8.6 will now be addressed. Using neutron diffraction, the stored energies of various texture components in 88% cold rolled AA3104 Al after light annealing have been determined by Rajmohan and Szpunar (1999) and given in Table 8.2. It is significant that the {100}<001> (Cube) orientation, which has been reported as the recrystallization texture in many investigations (Humphreys and Hatherly, 2004), has the highest stored energy. Low carbon steel provides a similar example, where the
-fibre components are known as soft components since they possess lower stored energy but, nevertheless, they are eliminated during growth of higher energy -fibre recrystallized grains (Humphreys and Hatherly, 2004). These cases point out the necessity for considering local variations, and not average values, in stored energy of the deformed microstructure, along with the temporal change during annealing.
Table 8.2: Stored energy values of selected orientations in 88% rolled and stress-relieved (15 min at 200 °C) AA3104 sheet (Rajmohan and Szpunar, 1999).
Stored energy, J (g-atom)1 Orientation After deformation After stress relief Cube {001}<100> 13.6 6.4 Copper {112}<111> 8.2 5.0 S {123}<634> 4.7 2.5 Brass {110}<112> 3.2 1.2
Gottstein and co-workers (Sebald and Gottstein, 2002) recently developed a time- resolved, but space-averaged, statistical “ReNuc” model for recrystallization texture evolution, in which nucleation is assumed to be distributed on grain boundaries, shear bands and transition bands between two orientations. The growth rates of
107 nuclei depend on grain boundary velocity, vGB, which is a product of boundary mobility, m, and local driving force, p:
4 GB D Rx D Rx gpggmggv D )(),(),( , (8.11)
where boundary mobility D ggm Rx ),( is a misorientation function between deformed
(gD) and recrystallized (gRx) regions of the microstructure, and the driving force
gp D )( is calculated using the Taylor criterion of orientation-dependent stored energy arising from interactions of operating slip systems (Sebald and Gottstein, 2002)
Some notable outcomes of the computer simulations i.e. “ReNuc” model are described by Crumbach et al. (2006b). It was found, for nucleation occurring exclusively at grain boundaries, that the resultant nuclei spectrum is dominated by
Copper and S3 orientations, while nucleation at transition bands results mainly in nuclei of the Cube orientation. The Brass component appears as a low intensity component when orientation dependent recovery (ODR) is assumed, but is a dominant component otherwise. Therefore, consideration of ODR may be essential for texture prediction. Briefly, orientations with a large number of activated slip systems are assumed to recover the fastest (Sebald and Gottstein, 2002; Crumbach et al., 2006a). However, no experimental evidence in support of their mechanism was provided. They also highlighted that the available experimental data are inadequate in terms of spatial and temporal resolution. Therefore, these workers could only compare with measured annealing textures to justify their explanations.
In the Al layers of the present materials, the presence of a large fraction of HAGBs in the deformed and recovered LBs is sufficient to create a situation where grain boundary nucleation becomes dominant. In this context, the “ReNuc” model successfully explains the elimination of Brass component in the Al layers if both grain boundary nucleation and orientation-dependent recovery are operative. As
108 shown in Fig. 5.27a, the texture of the Al layers after annealing for 6 h at 350 °C contains the same orientations as the rapidly growing grains observed after 3 min at this temperature (Fig. 5.23a).
The foregoing explanation has been recently reconsidered by Zahid et al. (2009) where it was argued that the Brass orientation is a low energy texture component and also remains at a low level after recovery. Thus, it is more likely for the Brass orientation to recrystallise and consume the neighbouring substructure, thereby becoming a prominent component in the recrystallization texture. Zahid et al. (2009) also highlighted that a more accurate investigation of the statistical substructural parameters is required to further study and compare the evolution of the Brass orientation relative to other orientations. Their explanation could be tested by conducting an investigation to show the spatial distribution of substructure within the ribbons/LBs of a particular orientation and their subsequent alteration during the early stages of recovery.
It may happen that a large block of substructure recovers at high energy texture components and recrystallise in advance of lower energy components. One notable case of this type of proces % 5 % % " =,,,>): ? =,,>@: grains. This concept can reconcile the stored energy and recrystallization argument.
8.3.1.3 Influence of shear banding on recrystallization textures in Aged-ARB
Figure 8.6c shows that there is a weakening of the recrystallization texture in the Al layers for the material that generated profuse shear banding during processing (Aged-
ARB). It is well-known that shear bands are preferred sites for recrystallization during annealing and these deformation features can also strongly affect the recrystallization texture (Humphreys and Hatherly, 2004). Shear bands may, or may
109 not, be crystallographic. For the former (i.e. those that operate within grains of particular orientations), the material within a given shear band will rotate to particular orientations. Hence, recrystallization will generate grains with orientations found within the orientation spread of the shear band, that is: Goss, Q ({013}<231>)
(Engler and Lucke, 1992) and some other -fibre ({111}//ND) components dominate the recrystallization texture (Engler et al., 2001). These components are also correlated with the -fibre rolling textures by systematic rotations. For the case where shear bands are not crystallographic, as observed in Al-Mg alloys (Koizumi et al., 2000), high strain rolling can generate a very high density of intersecting shear bands of different sign. This results in extreme perturbation of the lattice towards scattered orientations with a consequent randomization of the recrystallization texture (Koizumi et al., 2000).
The shear bands in Aged-ARB (Fig. 8.1b) are not the same as the aforementioned shear bands. They are examples of macro-scale shear bands that generally alter the orientations of a bundle of lamellar bands in the Al layers (irrespective of orientation) by a simple rotation (up to ±20 ) about TD. Therefore, the resultant weakening of the recrystallization texture (Fig. 8.6c) is probably due to nucleation within shear banded segments (as seen in Fig. 6.12) of the lamellar bands, since these segments are rotated 9-20 from -fibre (Fig 6.11). The same magnitude of orientation spread in the recrystallization textures is indeed expected to appear in a pole figure towards a random distribution, as shown in Fig. 8.5c. In contrast, substantial continuous recrystallization of the lamellar bands in the Al layers in
SSSS-ARB occurs by random orientation selectivity, thereby leading to a retained rolling texture after annealing (Fig. 8.6b), although the <111> pole distribution in
Fig. 8.6c is not completely random; such a large spread of orientations is difficult to generate in commercial purity aluminium and, therefore, might improve sheet formability.
110 8.3.2 Annealing behaviour following ARB at high homologous temperature
8.3.2.1 Microstructural development in Al layers during ARB in the presence of concomitant recovery and recrystallization
Chapter 7 described the production of a 64 layered Al/Al(Sc) composite material
(termed SSSS-ARB-HT). The major difference between this material and those described in chapters 5 and 6 was the higher roll bonding temperature of 350 °C (Tdef
~ 0.67Tm). This high rolling temperature and the large strain (50%), generated per rolling pass, is sufficient to considerably alter the microstructure and texture of the
Al layers (cf. section 7.2). In contrast, the high processing temperature is also expected to result in precipitation of a fine dispersion of Al3Sc within the Al(Sc) layers after the first few roll bonding cycles, thereby impeding static/dynamic restoration during the subsequent roll bonding operations.
This retention of the deformation microstructure throughout the roll bonding schedule resulted in both the generation of the classic FCC rolling texture in these layers (Fig. 7.4) and extended annealing up to 350 °C was not sufficient to alter the deformation microstructure and texture due to particle pinning effects (Figs. 7.10,
7.17 and 7.23). Returning to the Al layers, repeated heating to 350 °C for each roll bonding cycle is expected to recrystallize the Al layers either after each cycle or once sufficient strain is generated within these layers.
For example, the Al layers in SSSS-ARB-HT will undergo extensive dynamic recovery during the first rolling pass and reheating for the second rolling pass will result in static recovery or partial/complete recrystallization of the Al layers. These deformation and dynamic/static restoration processes are then repeated in further rolling bonding cycles. These continual changes to the microstructure of the Al layers generates a complex texture that, after a given rolling pass, may be a combination of the deformation texture and that generated by recrystallization. This
111 type of behaviour is evident in the deformation microstructure of the fabricated material, whereby the microstructure (Fig. 7.3) and texture (Fig. 7.5) of the central
Al layer (Al-4) is consistent with that produced in coarse-grained aluminium after rolling to moderate strains (Humphreys and Hatherly, 2004).
8.3.2.2 Development of microstructure and texture during annealing
Due to lack of recrystallization after roll bonding at 200 °C, the very high strains generated in both the Al layers in SSSS- and Aged-ARB generated a largely retained rolling texture on subsequent annealing (Figs 8.6b and c). In contrast, roll bonding at
350 °C results in a complex combination of deformation and annealing for each rolling cycle. The resultant microstructure after annealing at 350 °C (Fig. 7.21) shows the same characteristic features as that observed in SSSS- and Aged-ARB whereby recrystallizing grains have grown to complete layer thickness. However, the texture of these layers is different, as shown in Fig. 8.6d, due to normal discontinuous static recrystallization occurring during annealing.
8.3.3 Effect of Al3Sc dispersoids on structural stability of the Al(Sc) layers
The Al-0.3%Sc alloy used in this thesis was heat treated to generate either a uniform dispersion of nanosized Al3Sc dispersoids prior to rolling, as shown in Fig. 6.1
(Aged-ARB) or roll-bonded with Al in the supersaturated condition (SSSS-ARB and
SSSS-ARB-HT). For SSSS-ARB, subsequent annealing at 350 °C resulted in rapid precipitation of these dispersoids within the deformation substructure whereas, for
SSSS-ARB-HT, this precipitation process probably commenced during the multi- pass rolling stages at 350 °C, with subsequent annealing likely to complete the precipitation reaction. Therefore, regardless of the processing route, all materials are expected to contain Al3Sc dispersoids within the Al(Sc) layers and this is argued to result in the remarkable structural stability of these layers and the generation of the multilayered recrystallized/recovered microstructure shown in Figs 5.24-5.26, 6.13,
112 6.15, 7.2 and 7.21.
8.3.3.1 Kinetics of coarsening of Al3Sc dispersoids
The coarsening behaviour of Al3Sc dispersoids in a range of coarse- and fine-grained
Al-Sc alloys has been studied in detail by several research groups. In general, it was found that a dispersion of nanosized, coherent Al3Sc particles is highly resistant to coarsening up to 500 °C. The following provides an analysis of the expected coarsening behaviour of Al3Sc dispersoids in an Al-0.3%Sc alloy for gaining insight into the stability of the microstructure within the Al(Sc) layers and the generation of the stable Al/Al(Sc) interfaces in the multilayered composites. The kinetics of particle coarsening can be understood by the use of the Lifshitz-Slyozov-Wagner
(LSW) theory for conditions where lattice diffusion is the rate controlling step in the coarsening process (Lifshitz and Slyozov, 1961, Wagner, 1961):
33 )1(8 Vcc m Dt rr o 2 (8.12) RT cc )(9
where ro is the initial particle radius either prior to annealing or once precipitation is complete during annealing, r is the average particle radius at time, t, c and c is the equilibrium concentration of Sc in the Al matrix and Al3Sc, respectively, Vm is the
-3 3 partial molar volume of Sc in Al3Sc (~ 1.035 x 10 m /mol (Seidman et al., 2002)), D is a diffusivity term, is the interfacial free energy between the Al matrix and Al3Sc and RTis the grain or subgrain radius and temperature, respectively.
There have been several studies of particle coarsening in Al-Sc alloys where -values ranging from 0.041-0.226 J/m2 have been reported (Novotny and Ardell, 2001;
Seidman et al., 2002; Iwamura and Miura, 2004). While a direct measurement of has not been reported, = 0.2 J/m2 appears to be the most reasonable value for particle coarsening in Al-Sc alloys for the case where nucleation is completed and
113 when the particle size is greater than a few nanometres (Robson et al., 2003).
The diffusion coefficient, D, in Eq. (8.12) may be given by the following Arrhenius- type relationship (Fujikawa, 1997):
o exp( QDD b RT )/ (8.13)
-4 2 Tracer diffusion experiments (Fujikawa, 1997) have shown that Do = 5.31 x 10 m /s and Qb = 173 kJ/mol, the activation energy for bulk diffusion of Sc atoms in Al.
The solid solubility, c , of Sc in the -Al matrix is given as (Jo and Fujikawa, 1993):
c exp( RS )/ exp( '4' H RT)/ (8.14) where 'S/R is a constant (= 6.57) and 'H is the enthalpy change (= 62.8 kJ/mol) for dissolving Sc in the aluminium matrix. Equation (8.14) is shown to closely fit the available experimental data of the solvus boundary in Al-Sc alloys (Drits et al.,
1973).
Equation (8.12) has been shown to provide a good fit of the experimental data for
Al3Sc particle coarsening in: (i) a coarse-grained Al-0.2%Sc alloy at temperatures in the range 400-490 °C (Iwamura and Miura, 2004), and (ii) an ECAP-generated submicron grained Al-0.2%Sc alloy at temperatures in the range 350-500 °C (Ferry et al., 2005). Both studies also showed that third order particle coarsening kinetics
3 ( tr ) is valid, thereby indicating that the Al3Sc particles coarsen at a rate controlled by the bulk diffusion of scandium in the Al matrix. Based on the good fit of recent experimental data with Eq. (8.12), Fig. 8.11 shows the computed average particle radius in the Al-0.3Sc alloy as a function of time for annealing temperatures 9 ranging from 300 to 500 °C. To generate the data in Fig. 8.11, ro 105.2 m, c =
0.126 at.% and c = 25 at.% was used. It is clear that particle coarsening within the
114 Figure 8.11. Computed average Al3Sc particle radius (Eq. 8.12) as a function of annealing time for temperatures ranging from 300 to 490 °C. Al(Sc) layers is expected to be negligible, even after extended annealing at 350 °C.
8.3.3.2 Conditions for retardation of recrystallization in the Al(Sc) layers
It has been argued that the presence of small, closely-spaced particles affects the development of a cell or subgrain structure during deformation, thereby resulting in cells or subgrains with low misorientations. Indeed, this is the case for the Al(Sc) layers in Aged-ARB. Nevertheless, it is well-known that nucleation of recrystallization occurs preferentially at microstructural heterogeneities associated with large strain gradients such as grain boundaries, deformation bands and shear bands and deformation zones at large particles (Humphreys and Hatherly, 2004).
These types of heterogeneities are present in the roll-bonded composite with particular prevalence of shear banding in Aged-ARB (Figs 6.3-6.5). Hence, their effectiveness as potential nucleation sites is clearly reduced in the Al(Sc) layers due to the boundary pinning effects of the Al3Sc dispersoids.
Retardation of recrystallization in particle-containing alloys depends strongly on the particle size and the inter-particle spacing (Doherty and Martin, 1962; Humphreys and Hatherly, 2004). Compared with single-phase alloys, recrystallization can be completely inhibited by a dispersion of closely-spaced, fine particles (Humphreys and Hatherly, 2004). A number of experimental investigations have shown that the transition from accelerated to retarded recrystallization is primarily a function of the dispersion parameter (f/r) where f is the particle volume fraction and r the particle radius, that is, for f/r > 2) -1 (Humphreys and Hatherly, 2004). Assuming that nucleation of recrystallization occurs at microstructural heterogeneities containing a " % 5 % % % % substructure to form a viable nucleus for recrystallization has been shown to occur when (Humphreys and Hatherly, 2004):
115 f 4 * > f <( 0.05) (8.15a) r 3) m f 3/1 * f >(> 0.05) (8.15b) r ) m
-1 where * 7"! = 3 (both and ! are constant values), )m ~ 15 *7 ) and %m = high angle boundary of misorientation.
For alloys containing a low volume fraction of particles, nucleation is expected to be suppressed if f/r > ,7) -1 (Orsund et al. 1989), which is close to the experimental observations noted earlier. Conversely, for a nucleus that has formed in a particle- containing alloy, the grain will want to expand due to the driving pressure for growth
(PD) (due to dislocation density effects) but this expansion is offset by the Zener pinning pressure (PZ) due to the presence of particles and the retarding pressure due to boundary curvature of the newly-formed nucleus (PC). The net driving pressure for recrystallization (P) is given by the following energy balance (Humphreys and
Hatherly, 2004):
3 f 2 PPPP ! Gb 2 - - (8.16) CZD 2 Rr where ! is dislocation density, G is shear modulus and b is dislocation burger vector.
Based on the available experimental and theoretical evidence, Zener pinning plays a major role in retarding the nucleation stage, as given by Eq. (8.15). This suppression of nucleation in the Al(Sc) layers is evident in the EBSD micrographs of Figs 5.26 and 6.15 where no nucleation was observed after 6 h at 350 °C.
Assuming that the principal criterion for retarding recrystallization in microstructural heterogeneities is the suppression of nucleation (i.e. Eq. 8.15), the change in magnitude of the dispersion parameter as a function of annealing parameters
116 provides a platform for understanding the stability of the deformation substructure in the Al(Sc) layers. Unfortunately, the particle volume fraction, f, is a difficult parameter to determine experimentally in heavily deformed microstructures. The equilibrium volume fraction, feq, can therefore be used as an estimate which is determined by the lever rule (Andersen and Grong, 1995):
o cc feq (8.17) cc
where co is the initial solute concentration and ca is given by Eq. (8.14).
By combining Eqs (8.12) and (8.17), the effect of annealing time for a range of temperatures on the dispersion parameter is given in Fig. 8.12a. It can be seen that this parameter remains well-above the critcal range for the transition from retarded to accelerated recrystallization at annealing temperatures below 500 °C. This simple analysis indicates that the Al(Sc) layers are expected to remain unrecrystallized even after very long times at 350 °C. Based on the expected annealing behaviour in both the particle-free Al and particle-containing Al(Sc) layers, Fig. 8.12b is a schematic diagram showing the variation in dispersion parameter at the Al/Al(Sc) interface after extended annealing at 350 °C, whereby a distinct boundary between the recrystallized/unrecrystallized regions is both expected and observed, as shown in
Fig. 5.29 (shown again in Fig. 8.12b).
8.3.3.3 Particle-controlled subgrain coarsening in the Al(Sc) layers
Recent work on the annealing behaviour of ECAP-deformed Al-0.2%Sc alloy in the supersaturated condition, Ferry et al. (2005) showed that the deformation substructure following straining to an effective true strain of 9.2 was highly resistant to coarsening at 350 °C and no recrystallization was observed after 50h at this temperature, as shown in Fig. 8.13a. The present material is more highly alloyed
117 (a)
(b)
Figure 8.12. (a) Effect of anealing time for a range of annealing temperatures on the dispersion parameter, f/r, calculated using Eqs (8.12) and (8.17). (b) Schematic diagram showing the variation in dispersion parameter at the Al/Al(Sc) interface (Fig. 5.29) after extended annealing at 350 °C. thereby indicating further stability due to a higher volume fraction of coherent, Al3Sc particles. The accompanying TEM micrograph in Fig. 8.13b shows nanosized, coherent particles throughout the Al(Sc) microstructure, together with particle pinning at the Al/Al(Sc) interface. The microstructure is partially recovered with an average cell size of ~400 nm. In the same study, Ferry et al. (2005) demonstrated that the rate of Al3Sc particle coarsening in their as-deformed Al-Sc alloy was controlled by bulk diffusion of scandium atoms in the Al matrix, and argued that the cell size in the deformation microstructure is controlled by the rate of growth of particles since a linear correlation was found between cell size and particle size.
Similar to other particle-containing alloys (Manohar et al., 1998), the cell size in the as-deformed Al(Sc) layers in the present work should be controlled by the dispersion parameter (f/r) via the particle limiting grain size or Zener limit (Smith, 1948).
Zener predicted that the grain coarsening is inhibited completely when the grain size reaches a critical maximum grain radius (Rc) given by:
4r R (8.18) c 3 f
Equation (8.18) is often expressed in the general form:
r KR (3.3) zc f m
Experimental grain size data for a range of particle-containing alloys (for initial grain sizes greater than 0.1m) was found to closely follow Eq. (3.3) with Kz = 1/6 and m =
1 for low volume fractions of particles (f < 0.05) (Fig. 3.9). It is relevant to note that
Doherty (1982) has shown that the Zener pinning pressure for alloys containing coherent particles is z /6 rfP which generates Eq. (8.19) with Kz = 1/6 and m = 1. Using Eq. (3.3), the cell size expected in the Al(Sc) layers after 6 h at 350 °C (f =
118 Figure 8.13. (b) Effect of annealing time at 350 °C on subgrain radius in an ECAP-deformed Al-0.2Sc alloy, after Ferry et al. (2005). (b) Nanosized, coherent particles dispersed throughout the Al(Sc) microstructure showing boundary pinning (encircled). 0.011 from Eq. (8.17) and r = 10 nm from Eq. (8.12)) is in the range 150 to 600 nm, depending on the the constant, Kz, used in the calculation These values are comparable to the observed cell size in the Al(Sc) layers after extended annealing,
Figs. 5.28a-5.29.
8.4 Strength and Ductility of Al/Al(Sc) Composites
A major aim in the structure/property control in alloys is the development of a material that exhibits both high strength (yield etc.) and elongation to failure. There is a common relationship between strength and ductility in most alloys whereby an increase in strength of the alloy results in a concomitant decrease in ductility thereby generating a characteristic ‘banana’ curve. In recent years, several workers have exploited various processing routes for improving the ductility of a material without reducing its strength; this enabled a move into new region of property space for the material with the overall outcome being an increase in the so-called tensile toughness of the alloy.
A notable example of ductility enhancement is the work of Wang et al. (2002) who rolled copper at -196 °C followed by annealing for 3 min at 200 °C. This process generated a bimodal grain size distribution containing ultrafine grains (80–200 nm) and ~25% volume fraction of coarser grains (1–) -% % % % excellent combination of strength and ductility is the result of: (i) multi-axial stress states in the confined grains, (ii) twinning in the larger grains, and (iii) preferential accommodation of strain in the larger grains. Based on this discovery, several workers have demonstrated an increase in ductility without substantial loss in strength for a range of materials produced by various processing routes (see e.g.
Koch (2003). Figure 8.14, taken from Koch (2003), shows the normalized yield strength as a function of elongation to failure for a range of ultrafine grain materials
119 Figure 8.14. Normalized yield stress as a function of elongation to failure for a range of ultrafine grain metals showing enhancement in elongation to failure, after Koch (2003). (grain sizes of 100–500 nm) produced soon after the work of Wang et al. (2002).
With the exception of several electrodeposited Ni samples, all materials exhibit increased yield strength along with good ductility.
In the recent work of Huang et al. (2008c), nanostructured Al and IF steel were produced by ARB where both metals showed a similar structural size scale (200 nm), a formation of a high fraction of HAGBs, the presence of more than 10% of boundaries with misorientation angles less than 2° and a density of interior dislocations. The two nanostructured metals showed response to low-temperature annealing treatments which increases the maximum strength but decreased the elongation to failure. However, by introducing a cold rolling step following ARB, an improvement in ductility was achieved. These workers suggested that deformation rather than annealing should be considered carefully when optimizing the properties of nanostructured metals processed by plastic deformation to high strains.
Since the Al/Al(Sc) composites developed in this thesis contain coarse-grained Al layers and recovered Al(Sc) layers after ARB and subsequent annealing, it is useful to explore the strength/ductility relationships in these materials and compare the results with conventional Al alloys. Figure 8.15 shows the relationship between yield strength and elongation to failure for SSSS-ARB-HT and SSSS-ARB composites together with monolithic Al and Al(Sc) alloys and a range of commercial grade Al alloys. It is clear that the classic inverse strength/ductility relationship exists in all materials although the composites exhibit slightly superior ductility over the monolithic alloys. The SSSS-ARB composite annealed for 6 h at 350 °C shows slightly better ductility than SSSS-ARB-HT. Nevertheless, the exaggerated increase in ductility (for a given strength) reported in several recent researches is not produced in the Al/Al(Sc) composites. Further extension into new regions of strength/ductility property space for the present materials may be achieved by
120 Developed Alloys Vs other AA 55
Commercial Alloys
45 Developed AA
Al & AlS c- Expe r ime nte d
35
25 % Elongation 15
5
-5 95 195 295 395 495 -5 Yield Stress (MPa)
Figure 8.15. Elongation to failure as a function of yield stress in SSSS-ARB-HT and extensively annealed SSSS-ARB compared with a range of commercial aluminium alloys and the two conventionally rolled Al and Al(Sc) alloys. intermittent annealing strategies that generates a moderately recovered substructure within the Al layers and a retained deformation microstructure in the Al(Sc) layers or following the work of Huang et al. (2008c) by incorporating a final rolling stage.
The present work is concentrated on the development of microstructure on properties such as yield strength and tensile ductility. While these parameters are not improved dramatically, it is possible that other important properties such as fatigue life and fracture toughness, particularly in bending, may be improved by roll bonding dissimilar metal combinations. In recent work, Cepeda-Jimènez et al. (2009) produced a multilayered composite comprising of alternating layers of Al and Al-Zn-
Mg-Cu alloy by hot rolling. This material showed a considerable increase in Charpy impact toughness compared with the as-received Al-Zn-Mg-Cu alloy. They attributed the improved toughness to crack deflection at interfaces prone to delamination during impact testing. The work of Cepeda-Jimènez et al. (2009) demonstrates that other useful properties such as impact toughness may indeed be improved substantially in multilayered composites produced by accumulative roll bonding.
8.5 Control of Annealing Texture and Formability
The influence of crystallographic texture on sheet formability is well known, with superior formability generally characterized in the practical sense by a high Lankford parameter, R, and planar anisotropy, 'R, close to unity (Humphreys and Hatherly,
2004). Optimizing R in cubic metals usually requires a strong -fibre ({111} || ND) texture (Humphreys and Hatherly, 2004), which is relatively easy to achieve in steels but yet to be realized in aluminium. Therefore, Al alloys require a texture trade-off approach known as “texture balancing” whereby a combination of certain orientations is required to nullify the adverse effect of other orientations. For
121 example, the Cube orientation alone is undesirable, but can be beneficial in the presence of -fibre components. This has been recently demonstrated by Hu et al.
(1998) using the continuum mechanics of textured polycrystals framework proposed by Montheillet et al. (1985).
They showed that the Brass component contributes in-plane anisotropy ('R) that can be nullified by the Cube component. A method for achieving such a mixed texture has been demonstrated by Jazaeri and Humphreys (2004a,b), whereby the Cube component is dominant on annealing after rolling to true strains below 2.6, but the - fibre orientations gradually take over with increasing strain and dominate for strains greater than 3.9; this change in texture of the annealed material was a result of the transition from discontinuous to continuous recrystallization. In the present thesis,
ARB and annealing of SSSS-ARB did not significantly change the deformation texture. For a material comprising only of Al layers, the resulting texture may be favourable in the context of improving the formability of aluminium.
122 Chapter 9 ______CONCLUDING SUMMARY
Multilayered aluminium alloy composites have been produced by accumulative roll bonding (ARB) to generate thin-gauge sheet product comprising alternating layers of commercial purity Al and an Al-0.3% Sc alloy. Three different Al/Al(Sc) combinations were produced by varying the processing parameters. In one set of fabricated materials, two varieties of roll bonded sheet were produced by rolling at
200 C whereby the Al(Sc) alloy was heat treated in two different conditions prior to rolling. The first material, termed SSSS-ARB, Al(Sc) was heat treated to generate a supersaturated solid solution and subsequently roll bonded with Al. For the second material, termed Aged-ARB, roll bonding was carried out with Al(Sc) in the artificially aged condition. Here, the initial heat treatment generated a uniform dispersion of nanosized Al3Sc dispersoids. Following five ARB cycles, 0.5 mm gauge sheet consisting of 32 alternating layers of Al and Al(Sc) were produced. To generate the third material, termed SSSS-ARB-HT, roll bonding was carried out with
Al(Sc) in the supersaturated condition at the higher temperature of 350 C. This
123 material was rolled in six cycles to a higher total strain and generated 0.5 mm gauge sheet consisting of 64 alternating Al/Al(Sc) layers. The roll-bonded Al/Al(Sc) composites were subsequently annealed for investigating the influence of heat treatment and rolling parameters on static restoration processes. Annealing was carried out for various times at a temperature of up to 350 C.
A detailed study of microstructural evolution of the three materials during roll bonding and subsequent annealing were carried out using backscatter electron (BSE) imaging in the FEGSEM, ion channeling contrast (ICC) imaging in the DualBeam
Platform and transmission electron microscopy (TEM) using bright field and dark field imaging modes. The crystallographic nature of the microstructure of all materials, in both the deformed annealed states, were investigated using electron backscatter diffraction (EBSD) in the FEGSEM and both selected area electron diffraction (SAED ) and convergent beam electron diffraction (CBED) in the TEM.
Finally, certain mechanical properties of the materials were investigated by hardness testing and tensile testing.
It was found that the initial microstructure and roll bonding temperature played an important role in the development of the microstructure and texture in both the as- deformed and annealed states. The major conclusions of the study are now described:
The Deformed State
Depending on the rolling and initial heat treatment condition of the Al(Sc) layers, roll bonding generated a certain deformation microstructure in each of the Al and
Al(Sc) layers and also affected the uniformity of the alternating Al/Al(Sc) layers.
When Al was roll bonded with Al(Sc) in the supersaturated condition and at a low deformation temperature (200 °C), deformation of the individual layers was largely
124 uniform and generated parallel Al/Al(Sc) multilayers. However, artificial aging of the Al(Sc) sheet prior to rolling at the same temperature had a marked influence on the integrity of the layers whereby large-scale shear banding resulted in marked buckling of the layered microstructure.
For the Al/Al(Sc) combination roll bonded at a higher temperature (350 °C) and to a larger strain, large-scale shear banding was also evident despite the Al(Sc) alloy being in the supersaturated condition. The reasons for such behaviour were attributed to the changes in microstructure within the Al and Al(Sc) layers during multi-pass rolling at the higher temperature. Here, repeated recrystallization of the Al layers occurred during each reheating and rolling stage to maintain relatively soft layers. In contrast, the higher processing temperature resulted in precipitation of a dispersion of fine Al3Sc particles within the Al(Sc) layers after the first few roll bonding cycles, thereby altering the work hardening behaviour of these layers and generating conditions similar to that observed in Aged-ARB.
The effect of the initial heat treatment condition of Al(Sc) and the rolling temperature had a significant influence on the deformation microstructure and texture of the material. For SSSS-ARB, the interior microstructure of Al and Al(Sc) layers consisted of lamellar bands (LB) aligned parallel to the rolling direction (RD).
Compared with the Al layers, LBs in Al(Sc) layers were more refined with the average thickness being 0.15 m and 0.4 m, respectively. This significant structural refinement was explained by the interaction of solute atoms with dislocations during deformation. Here, extensive dynamic recovery of the LBs readily occurred in the Al layers and these bands were separated by high angle boundaries. In the Al(Sc) layers, however, recovery was retarded and the generated LBs separated largely by low angle boundaries. The influence of artificial aging of Al(Sc) also had an impact on the deformation microstructure resulting in more uniform deformation.
125 For both SSSS- and Aged-ARB in the as-rolled state, the Al layers exhibited orientations (within a given Al layer) distributed along the FCC -fibre. However, shear banding resulted in a spread of orientations about the transverse direction (TD) in the highly sheared segments of magnitude equal to the inclination of the lamellar bands with respect to RD. Hence, shear banding in Aged-ARB reduced both the intensity and sharpness of the deformation texture and produced a larger intensity reduction of the Brass texture component compared with Copper. This was explained by a relatively higher orientation stability of Copper than Brass under the influence of shear banding.
For SSSS-ARB-HT, the deformation microstructure and texture within the Al layers was considerably different to both SSSS-ARB and Aged-ARB due to the higher rolling temperature. Here, the Al layers undergo extensive dynamic recovery during the first rolling pass and reheating for the second rolling pass resulting in static recovery or partial/complete recrystallization of the layers; these processes were then repeated in further roll bonding cycles. These continual changes to the microstructure of the Al layers during multi-stage rolling at 350 °C also generated a complex texture compared with the -fibre texture after rolling at 200 °C. The Al(Sc) layers developed the classic -fibre texture due to precipitation of Al3Sc dispersoids within these layers after the first few roll bonding cycles, thereby impeding static/dynamic restoration during the subsequent roll bonding operations.
Both Aged-ARB and SSSS-ARB-HT developed large-scale shear bands during roll bonding. The origin of shear banding was investigated in detail for both SSSS- and
Aged-ARB and was attributed to the differential hardening behaviour of Al and
Al(Sc) layers. During ARB processing of either material, the hardening behaviour of the Al layers was similar and increased from ~30 to ~44 HVN whereas the SSSS
Al(Sc) and aged Al(Sc) layers increased from 29 to 68 HVN and 78 to 88 HVN,
126 respectively. The rate of hardening of the layers therefore increased in the following order: aged Al(Sc), Al and SSSS Al(Sc). Such differential flow behaviour between the adjacent Al and Al(Sc) layers is expected to induce an in-plane shear stress at their interface. This is because the rolling mill induces the same magnitude of load to both the layers separated by a bonded interface, across which the material in the softer layer flows to a larger strain than the harder layer. Thus, the soft layer pulls the harder layer resulting in an in-plane shear force at the interface. Using a simple flow model, it was shown that plastic instability in the Al(Sc) layers occurs when the in- plane shear stress becomes considerably greater than the yield strength (in pure shear condition) of Al(Sc).
The Annealed State
The three composite materials were annealed for up to 6 h at 350 °C. This extended annealing generated alternating layers of coarse grains (Al layers) and a recovered substructure (Al(Sc) layers) with the substantial waviness of the layers in both Aged-
ARB and SSSS-ARB-HT being inherited from the as-deformed material. It was found that scandium, either in solution or in the form of Al3Sc particles, significantly impeded static recrystallization within the Al(Sc) layers during annealing and these particles also severely impeded the growth of coarse grains from the Al layers into
Al(Sc) layers. For all ARB materials, the deformation substructure was retained in the Al(Sc) layers whereas the deformation substructure within the high purity Al layers was replaced by coarse grains at the scale of the layer thickness. Such differences in static restoration behaviour generated composite materials containing alternating coarse-and fine-grained layers.
At the low rolling temperature, dynamic recovery in the Al layers was the sole restoration process and generated the classic -fibre texture. On annealing, there was rapid coarsening of certain grains within the deformation microstructure within the
127 Al layers to generate a texture containing a spread in rolling texture components (- fibre) although the spread in orientations was larger in Aged-ARB due to shear banding caused by rolling. The retention of the rolling texture (-fibre) within the Al layers in both SSSS- and Aged-ARB after annealing indicates that static restoration did not happen by conventional nucleation and growth processes characteristic of discontinuous recrystallization but occurred by continuous recrystallization, consistent with that generated in many other severely strained alloys.
There was a major difference in the texture of the Al layers when roll bonding was carried out at elevated temperature due to repeated recrystallization during this processing stage. The microstructural development within the Al layers in SSSS-
ARB-HT during subsequent annealing at 350 °C was consistent with normal discontinuous recrystallization occurring to generate the coarse-grained layers.
Hence, the resultant recrystallization texture in Al layers was more random compared with the retained deformation texture in both SSSS- and Aged-ARB.
An interesting observation in both SSSS- and Aged-ARB was the decrease in intensity of the Brass texture components from within the -fibre rolling texture during annealing at 350 C. This was found to occur in the early stages of annealing whereby grains exhibiting other texture components within the -fibre grow rapidly and subsequently eliminate the Brass components. In the very early stage of annealing, extended recovery occurred whereby LAGBs formed along the lamellar bands and the lamellar nature of microstructure collapsed by relaxation of the boundaries. This was followed by spheroidization initiated boundary migration with grains containing the Copper and S3 texture components growing at the expense of grains of the Brass orientation. The orientation selectivity that eliminated the Brass texture components in the Al layers was explained by an orientation dependent
128 recovery (ODR) model of recrystallization, in which nucleation is assumed to occur on HAGBs between the -fibre components in the LBs.
Mechanical Behaviour
The mechanical properties of as-rolled and annealed SSSS-ARB-HT were investigated by tensile testing. This material was compared with SSSS-ARB following annealing for 6 h at 350 °C and two conventionally rolled Al and Al-0.3Sc samples. It was found that as-rolled SSSS-ARB-HT was prone to early plastic instability during tensile straining by which there was a peak in stress soon after yielding that decreases with increasing strain to generate a low elongation to failure.
It was found that the extent of uniform elongation was considerably increased with increasing annealing temperature, although the maximum tensile stress was lowered.
The early onset of plastic instability in SSSS-ARB-HT was argued to be a result the numerous microstructural instabilities generated during roll bonding in the form of gross shear bands. Hence, tensile straining will result in strain localization along the shear bands soon after yielding.
It was found that the composite materials conform to the classic inverse strength/ductility relationship although they exhibit slightly superior ductility over their monolithic alloy counterparts. Hence, there was no significant improvement in ductility (for a given strength), as reported in several recent researches on a range of materials. For the present materials, an extension into new regions of strength/ductility property space may be achieved by intermittent annealing strategies that generates a moderately recovered substructure within the Al layers and a retained deformation microstructure in the Al(Sc) layers or by incorporating a final rolling stage in the materials. While both the yield strength and tensile ductility were not improved considerably in these Al/Al(Sc) composites, it is possible that other important properties such as fatigue life and fracture toughness may be improved.
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