2.

EXTRUSION OF 2024 ALLOY SECTIONS

BY

JAYA SUBRAMANIYAN

A Thesis submitted for PhD of London University

John Percy Group

Department of Materials Imperial College London SW7 2BP.

March 1989 TO MY PARENTS,

MICHAEL CLODE &

RAGHUNATHAN 1

Abstract

Extrusion of 2024 sections

The Aluminium Association alloys designated 2XXX have extensive application in the aerospace and transport industries. Their hot working range is limited because their high temperature flow properties are relatively strong and the alloys are prone to incipient melting at higher extrusion temperatures.

Much work has been accomplished on the production of limit diagrams for axi-symmetric rod extrusion which allows the prediction of pressure limits, surface acceptability structure and properties. The work is, however, not generally applicable to the production of shaped extrusions which constitute the bulk of output in the industrial environment.

The project seeks to relate both direct and indirect modes axi-symmetric and shaped sections by considering the geometrical variables involved (e.g. the aspect ratio and circumscribing circle diameter) in order to produce the pressure and surface limitation loci on the limit diagram. The project also involved a detailed investigation of the structure and substructure at all stages of processing from the as cast to the final heat treated extrudates. Also studied was the variation of structure within the direct and indirect deformation zones and across and along the extruded sections. Such work allowed the introduction of the structure loci on the limit diagram. Four different sections (round, square, T-shape and U-shape) were investigated both by the direct and indirect route.

The pressure, surface quality, recrystallisation (particularly in the outer band) and the structural loci were quantified in terms of the material constants contained in the high temperature constitutive equation and/or the temperature compensated strain rate. Conclusions are presented regarding the possibility of structure and property control in the extrusion process using the results from the the experimental matrix designed to establish the interaction of the process variables and strength, ductility, fracture toughness etc. 2

ACKNOWLEDGMENTS

The author would like to express his sincere thanks to Professor Terry Sheppard for his initiation of the project and his supervision, guidance and encouragement throughout the course of this work.

The author would also like to express his appreciation to Alcan International Ltd., for their financial and material support experimental facilities and useful discussions and technical advice notably Dr. A. J. Bryant and Dr. Nick Parson. I am also grateful to the Overseas Reasearch Scholarship Fund for providing partial financial assistance.

Sincere thanks are also expressed to present and past members of the John Percy Research Group and in particular Drs. Henry McShane, Mike Clode, N. Raghunathan for their useful discussions, suggestions, patience and willingness to help. Also I would like to thank Alison Mew and Dr. Efklidis loannidis for their help with the corrections and Apple Macintosh respectively. Sincere thanks to The JPDC for making every Friday afternoon eventful.

Technical assistance from Mr. Alec Neve, Mr Melvyn Andrews, and Mr. Roy Baxter and his colleagues in the workshop are also greatly appreciated.

Thank also to Ian and Alison Metcalfe and the lads at Garden Hall for making my stay there extremely enjoyable. I would also like to thank Mr. and Mrs Ludden for all their encouragement and help throughout this PhD programme.

And finally to my parents for all their love and support. 3

CONTENTS

Abstract 1 Acknowledgements 2

Contents 3 List of figures 7

List of tables 11

Chapter 1 Introduction

1.0 Introduction 1 4

Chapter 2 Literature Survey

2.1 The Al-Cu-Mg system 1 7 2 .2 The effect of alloying elements 1 9

2 .3 Preheat treatment 2 0

2.3.1 Homogenisation 2 0 2 .3 .2 Solution treatment 2 2

2.3.3 The mechanisms of age-hardening 2 2 2.3.4 The effect of age hardening 2 5

2 .4 Hot working 2 7 2.4.1 Empirical relationships 2 7

2.4.2 Structural aspects of hot working 2 9

2 .4 .3 Restoration processes 3 0 2 .4 .4 Static restoration processes 31

2 .5 Substructure strengthening 31 2 .6 Extrusion 3 3 2.6.1 Direct extrusion 3 3 2 .6 .2 Indirect extrusion 3 6

2 .6 .3 Limit diagrams 3 7 2 .6 .4 Structure produced by extrusion 3 9 4

2 .6 .5 Shaped extrusion 4 0

Chapter 3 Theory

3.1 Introduction 4 3

3.2.1 Torsion analysis 4 3 3 .3 Extrusion analysis 4 4 3.3.1 Temperature rise 4 4 3 .3 .2 Strain rate 4 4 3 .4 Evaluation of the billet container friction in direct extrusion 4 5

Chapter 4 Experimental

4.1 Introduction 4 8 4 .2 The extrusion press 4 8 4.2.1 Tooling 5 0

4.3 Material 5 0

4 .4 Examination of extrudates 5 3

4.4.1 Heat treatment of extrudates 5 3 4 .4 .2 Hardness tests 5 3

4 .4 .3 Optical microscopy 5 4 4 .4 .4 Electron microscopy 5 5 4 .4 .5 Tensile testing 5 5

4 .4 .6 Fracture toughness tests 5 6

4 .5 Torsion testing 5 6

4 .6 Partially extruded billets 5 8 4 .7 Surface quality 5 8

Results and Discussion

Chapter 5 Torsion Analysis

5.1 Torsion data analysis 61

5.2 Torque/twist curves 61 5 .3 Flow stress characteristics 6 3 5

5 .4 Constitutive equations 6 7 5.5 Torsion structures 6 9 5 .6 Conclusion 7 0

Chapter 6 Extrusion Analysis

6.1 Extrusion data analysis 7 2 6 .2 Load-displacement curves 7 2 6 .3 Effect of extrusion parameter on the pressure to extrude 7 4 6.3.1 Extrusion temperature 7 4 6 .3 .2 Extrusion ratio 7 9 6.4 The variation of peak pressure with Z 8 2

6 .5 Variation of AP with Ln (Zp) 8 6

6 .6 Evaluation of the friction conditions 8 9 6 .7 General pressure equation 9 5

6 .8 The effect of section shape on the pressure 1 01 6.8.1 Effect of section shape on Peak pressure 1 01

6 .8 .2 Effect of section shape on APeak pressure 1 0 6

6 .9 Geometrical considerations 1 0 6

6 .1 0 Surface quality of extrudates 1 1 3 6.11 Conclusion 1 2 0

Chapter 7 Structural Analysis

7.1 The evolution of structure and substructure 1 23

7.1.1 As-cast structure 1 2 3 7 .1 .2 Homogenised structure 1 2 5

7 .2 Structural variations with extrusion conditions 1 2 8

7.2.1 Extrudate structures 128 7 .2 .2 Substructural variation with extrusion conditions 1 3 6 7 .2 .3 Effect of process conditions on the steady state substructure 1 4 2

7 .2 .4 Variation of substructures across and along extrudates 1 45

7 .3 Material flow during steady state extrusion 1 4 6 7 .4 Development of substructure during extrusion 1 5 3

7 .5 The effect of solution treatment on extrudate microstructures 1 61 7 .6 The effect of ageing on the microstructure 171 7 .7 Conclusion 1 7 4

Chapter 8 Room Temperature Properties

8.1 Room temperature properties of extrudates 1 7 7 8.1.1 Effect of extrusion conditions on the hardness properties 1 7 7 8 .1 .2 Variation of hardness along extrudate length 1 7 9

8 .1 .3 Variation of hardness across the extrudate 1 8 2 8 .1 .4 Effect of solution treatment on the hardness 1 8 2

8 .2 Ageing characteristics of alloy 2024 extrudates 1 8 4

8 .3 Tensile properties of 2024 alloy extrudates 1 9 0 8.3.1 Stress-strain characteristics of alloy 2024 1 90 8 .3 .2 T1-Tem per 1 93

8 .3 .3 T6-Temper 1 99 8.3.4 T5-Tem per 201

8 .3 .5 The effect of a preageing stretch on the tensile properties 201

8 .4 Fracture toughness properties of alloy 2024 2 0 2

8 .5 Limit diagrams 2 0 7 8 .6 Conclusion 2 1 3

Chapter 9 Conclusions

9.1 Major conclusions 2 1 5 9.2 Recommendation for further work 2 1 7

Appendices

1 Extrusion data 2 2 0

2 Ageing data 2 3 2

3 Tensile data 2 3 5 4 Short rod fracture toughness data 2 3 7

5 Experimental errors 2 3 8 6 Modification to calculations involving j x 2 4 0

References 2 4 3 7

List of figures

Fig. 2.1 Phase diagram of aluminium- 24

Fig. 2.2 Principles of direct and indirect extrusion 34 Fig. 2.3 Idealised load-displacement curve for direct and indirect extrusion 35 Fig. 2.4 Extrusion ratio vs extrusion temperature limit diagram 37

Fig. 4.1 General layout of press and extrusion tooling 49 Fig. 4.2 The dies used for rod extrusion 51

Fig. 4.3 The dies used for shaped extrusion 52

Fig. 4.4 Short rod fracture toughness test sample 57

Fig. 5.1 Schematic torque vs twist curves 62 Fig. 5.2 Stress-strain curves for torsion (strain rate =16.0/s) 64 Fig. 5.3 Stress-strain curves for torsion (strain rate =5.0/s) 65 Fig. 5.4 Stress-strain curves for torsion (strain rate =0.16/s) 66

Fig. 6.1 Typical load-displacement trace for direct extrusion of alloy 2024 at high and low temperatures 73

Fig. 6.2 Typical load-displacement trace for direct and indirect extrusion of alloy 2024 73

Fig. 6.3 Peak pressure vs initial billet temperature for direct extrusion 75

Fig. 6.4 Peak pressure vs initial billet temperature for indirect extrusion 76 Fig. 6.5 Peak pressure vs Ln R for direct extrusion 80 Fig. 6.6 Peak pressure vs Ln R for indirect extrusion 81 Fig. 6.7 Peak pressure vs Ln (Z) for direct and indirect extrusion 84 Fig. 6.8 Peak pressure vs Ln (Z/A) for direct and indirect extrusion 85

Fig. 6.9 Tangential construction for AP evaluation 87

Fig. 6.10 A Peak Pressure vs Ln Z for extrusion 88

Fig. 6.11 Peak pressure vs billet length for direct extrusion alloy 2024 90

Fig. 6.12 Peak pressure vs predicted pressure calculated using equation 6.12 for both modes of extrusion 98

Fig. 6.13 Effect of die geometry on load displacement curves for both modes of extrusion 102 8

Fig. 6.14 Graph of peak pressure vs Ln Zj for shaped sections (direct mode) 103

Fig. 6.15 Graph of peak pressure vs Ln Zj for shaped sections (indirect mode) 104

Fig. 6.16 Graph of APeak pressure vs Ln Zj for shaped sections (direct mode) 107

Fig. 6.17 Graph of APeak pressure vs Ln Zj for shaped sections (indirect mode) 108

Fig. 6.18 Graph of peak pressure vs Ln Zj (direct mode) 111

Fig. 6.19 Graph of peak pressure vs Ln Zj (indirect mode) 112

Fig. 6.20 Predicted pressure from general pressure equation vs experimental extrusion pressure for shaped sections 114

Fig. 6.21 Typical examples of extrudate surfaces of alloy 2024 116

Fig. 6.22 Ln Zj vs initial billet temperature for surface conditions - direct and

indirect extrusion of alloy 2024 119

Fig. 7.1 As cast microstructure 124

Fig. 7.2 Homogenised material microstructure 126 Fig. 7.3 Heat treatment profile given prior to extrusion 127 Fig. 7.4 Extrudate microstructure 129 Fig. 7.5 Shaped microstructure 131 Fig. 7.6 Volume % recrystallised vs initial billet temperature for direct extrudates 132

Fig. 7.7 Volume % recrystallised vs initial billet temperature for direct extrudates 132

Fig. 7.8 Recrystallised layer thickness vs Ln Zj for direct shapes 134

Fig. 7.9 Recrystallised layer thickness vs Ln Zj for indirect shapes 134

Fig. 7.10 Comparison of substructures in transverse plane for direct and indirect extrusion of alloy 2024 137

Fig. 7.11 Comparison of substructures in transverse plane for direct and indirect extrusion of alloy 2024 rod section 140

Fig. 7.12 Comparison of substructures in transverse plane for direct and indirect extrusion of alloy 2024 U-section 141

Fig. 7.13 Influence of strain on subgrain size 142

Fig. 7.14 Inverse subgrain size (d'1) vs Ln Zc relationship 143

Fig. 7.15 Inverse subgrain size (d"1) vs Ln Zj relationship 143 9

Fig. 7.16 Macrosections of partially extruded billets of alloy 2024 at 425°C - direct extrusion 146 Fig. 7.17 Macrosections of partially extruded billets of alloy 2024 at 300°C -

direct extrusion 149

Fig. 7.18 Macrosections of partially extruded billets of alloy 2024 at 300°C - indirect extrusion 150 Fig. 7.19 Metal flow in T-section 152

Rg. 7.20 Position extrusion cycle stopped for structural development 154 Fig. 7.21 Location of TEM specimens for structural development studies 155

Fig. 7.22 Development of substructure along flow line during steady state direct extrusion 157

Fig. 7.23 Development of substructure along flow line during steady state indirect extrusion 159

Fig. 7.24 Effect of solutionising on the structure of alloy 2024 extrudates 162 Fig. 7.25 Effect of heat treatment on the substructure of alloy 2024 extrudates 165 Fig. 7.26 Recrystallised grain size vs initial billet temperature for solutionised direct extrudates 166

Fig. 7.27 Recrystallised grain size vs initial billet temperature for solutionised indirect extrudates 166 Fig. 7.28 Recrystallised grain size vs subgrain size for solution direct and indirect extrudates of alloy 2024 168

Fig. 7.29 Volume % recrystallised vs initial billet temperature for solutionised extrudates 169

Fig. 7.30 Volume % rextd vs Ln Z\ for direct shapes after solutionising 170

Fig. 7.31 Volume % rextd vs Ln Zj for indirect shapes after solutionising 170

Fig. 7.32 Mechanism of grain growth in extrudates 164

Fig. 8.1 Hardness variation versus initial billet temperature 178 Fig. 8.2 Hardness variation versus extrusion ratio 180

Fig. 8.3 Variation of hardness along extrudate length 181 Fig. 8.4 Variation in microhardness across extrudate 181 Fig. 8.5 Hardness variation with extrusion temperature after solutionising 183 Fig. 8.6 Ageing characteristics at 120°C of solutionised 2024 alloy extrudates 185

Fig. 8.7 Ageing characteristics at 160°C of solutionised 2024 alloy extrudates 186 Fig. 8.8 Ageing characteristics at 180°C of solutionised 2024 alloy extrudates 186 10

Ageing characteristics at 120°C of as-extruded 2024 alloy extrudates 188 Ageing characteristics at 160°C of as-extruded 2024 alloy extrudates 188 Ageing characteristics at 180°C of as-extruded 2024 alloy extrudates 189 Stress-strain curves for tensile test of alloy 2024 extrudates 191

Proof stress vs Ln Zj - direct extrusion of alloy 2024 194

Proof stress vs Ln Zj - indirect extrusion of alloy 2024 194

Ultimate tensile strength vs Ln Zj - direct extrusion of alloy 2024 195

Ultimate tensile strength vs Ln Zj - indirect extrusion of alloy 2024 195

Percentage elongation vs Ln Zj - direct extrusion of alloy 2024 196

Percentage elongation vs Ln Zj - direct extrusion of alloy 2024 196

Variation of tensile properties with extrusion ratio for direct and indirect extrudates of alloy 2024 198

Fracture toughness vs Ln Zj for direct extrusion of alloy 2024 203

Fracture toughness vs Ln Zj for indirect extrusion of alloy 2024 203

Fracture surfaces of alloy 2024 extrudates T1, T5 and T6 conditions 206 Limit diagram for direct and indirect extrusion of round sections in alloy 2024 209 Limit diagram for direct extrusion of shaped sections in alloy 2024 211 Limit diagram for indirect extrusion of shaped sections in alloy 2024 212 11

List of tables

2.1 Typical composition of commercial Al-Cu-Mg alloys 18 2.2 Possible equilibrium compounds in alloy 2024 21 2.3 Transition phases in alloy 2024 21

4.1 Composition of cast material investigated 50 4.2 Classification of heat treatments employed 54

5.1 The hot working constants 67

6.1 Temperature dependency of peak pressure in alloy 2024 77 6.2 Temperature dependency of the flow stress 78

6.3 Effect of extrusion ratio on the peak pressure 82

6.4 The dependence of peak pressure on Ln Zp 83

6.5 The dependence of peak pressure on Ln ( Zp /A) 86

6.6 The dependence of peak height AP on Ln Zp 87

6.7 Peak pressure vs billet length regression data for direct extrusion 89

6.8 Coeff of friction jll obtained from direct and indirect extrusion data 92 6.9 Extrusion constants in the general pressure equation 97 6.10 Extrusion constants in general pressure equation 6.14 99

6.11 Extrusion constants in general pressure equation 6.16 100 6.12 Extrusion constants in equation 6.3 for shaped sections 105

6.13 Extrusion constants for shaped sections 106

6.14 Correction factors for non circular sections 110 6.15 New values of m and fi 94

7.1 The variation in volume % recrystallised during direct and indirect extrusion of alloy 2024 in the F-temper 135

7.2 The variation in volume % recrystallised along direct and indirect extrudates of alloy 2024 in the F-temper 136 7.3 Temperature compensated subgrain size relationships for extrusion and torsion 144 7.4 Subgrain size relationships for extrusion without temperature rise compensation 145 12

7.5 Variation of substructure size across and along the extrudate 146

7.6 Subgrain size measurements within the steady state direct extrusion

deformation zone; Tj =400°C 158

7.7 Subgrain size measurements within the steady state indirect extrusion

deformation zone; Tj = 400°C 160

7.8 The variation in volume % recrystallised after solutionising in direct

and indirect extrudates 169 7.9 Comparison between F, T5 and T6 subgrain dimensions for both modes of extrusion 172

8.1 Hardness values of alloy 2024 extrudates in both modes at different extrusion temperatures 178 8.2 Hardness values of extrudates for both modes at different extrusion

ratios 180 8.3 Variation of strain hardening exponent of alloy 2024 alloy extrudates 192 1.0 Introduction 14

Introduction

Aluminium and its alloys in general have always possessed the advantage of good working characteristics in all the conventional metallurgical processes such as rolling and drawing. Aluminium alloys may be conventionally grouped into two classes, the heat treatable and non-heat treatable. The non heat treatable alloys include all the various grades of pure aluminium and all other alloys in which strength is developed largely by solid solution hardening from the annealed temper. These alloys include the 1XXX, 3XXX, 4XXX, and 5XXX series. The heat treatable alloys generally belong to the 2XXX series (aluminium - copper - ), the 6XXX series (aluminium - magnesium - ) and the 7XXX series (aluminium - - magnesium - copper). All three of these systems obtain their strength primarily by a precipitation treatment which requires that the elements contributing to hardening must be in supersaturated solid solution before the sequence.

One of the major applications of heat treatable aluminium alloys is in airframe construction. The excellent weight/strength relationship and good durability characteristics of aluminium alloys make them an ideal choice for transport applications, particularly in the aircraft industry. They emerged during the 1914-1918 period in the form of , and during the 1930's began to dominate the airframe as a structural material, notably with the introduction of the archetypal DC3 twin engine transport aircraft. By 1940, the stronger zinc containing alloys were being introduced, the available design strength having being doubled from the earlier alloys. A rather serious problem that arose from stress in these alloys led the aircraft industry back to the use of copper alloys such as 2024 for critical structural components. These latter alloys in improved versions have retained a strong position in airframe design up to the present time. The wing stringer for example in the A300, A320 and A340 Airbus are typical examples.

Over the last fifty years, the degree of control of their properties has been increasing as new insights and test techniques have become available. Today, the final properties of particular interest are strength, toughness, crack growth 15

resistance, exfoliation behaviour and stress corrosion resistance. Of the two major systems, 7XXX series alloys are chosen for their high strength, while 2XXX series alloys are generally selected for fatigue sensitive applications and where higher service temperature may be encountered.

Today extrusion is one of the principal metal forming processes, and despite its relative newness has achieved special importance. The main advantage of extrusion is that complex sections, both solid and hollow, can be manufactured in a single operation, meeting stringent geometric, property and cosmetic goals. The relative importance of these goals may vary with the product application, but there is a distinct advantage gained over other forming processes where laborious and expensive finishing operations are necessary when surface finish is critical. In some cases, extrusion is the only manufacturing route which is commercially viable for the net shape production of components to high dimensional tolerance, aircraft window frames being a classic example. An additional advantage gained by using extrusion processing is that different microstructures may be obtained by varying process parameters, such that optimisation of processing conditions leads to improved mechanical properties.

It is in the area of net shape extrusion and microstructural control that this thesis was initiated. The effect of process parameters on the hot working characteristics of a commercial Al-Cu-Mg alloy, AA2024, has been studied. Also included in the investigation is the effect of extrusion mode on the as-extruded and heat treated product. 2.0 Literature Survey 17

2.1 The Al-Cu-Mg system

Al-Cu-Mg systems have been known for over half a century since their discovery by Wilm5 . He observed that the hardness of aluminium alloys containing small amounts of copper, magnesium, silicon and , increased with time at room temperature after having been quenched from a temperature near the solidus. The first plausible explanation for this "ageing process" was put forward by Merica et al4 who postulated that age-hardening occurred in alloys in which the solid solubility increased with increasing temperature, thereby enabling a new phase to form at a lower temperature by precipitation from an initially super saturated solid solution (s.s.s.s.). Since then a range of commercial Al-Cu-Mg alloys has been developed, of which 2024 and 2014 are the most widely used.

The alloy 2014 was developed utilising the effect of silicon to produce an Al-Cu-Mg alloy that is more susceptible to artificial ageing and provides a higher level of strength than the original Duralumin alloy (2017). Alloy 2024 was developed in the 1930's as a higher strength natural ageing alloy to replace 2017. The response of alloy 2024 to artificial ageing is accelerated by strain hardening of the heat treated material and is available in a number of artificially and naturally aged tempers that reflect several degrees of strain hardening5 .

Both these alloys found widespread applications in the aerospace industry until the introduction of the higher strength Al-Zn-Mg-Cu alloys some 30 years ago.

Today, the 2000 series alloys are selected where fatigue, service temperature or ease of fabrication dictate.

Improved versions of the 2XXX series alloys, such as X2048, 2214, 2124 have recently been introduced for improved fracture toughness properties6 . For specific applications, other alloys have been developed, such as 2618 (RR58- Concorde alloy) for elevated service temperature or 2219 for weldability7 .

Typical compositions of some commercial Al-Cu-Mg alloys are given in table

2 . 1 . Table 2.1 Typical composition of commercial Al-Cu-Mg Alloys

Alloy Al Cu Si Mn Mg Fe Ni Ti Zr

2014 Rem 4.4 0.9 0.8 0.5 --- -

2017 Rem 4.0 - - 0.5 - - - -

2024 Rem 4.5 - 0.6 1..5 0.2 - --

2214 Rem 4.1 0.7 0.8 0.5 0.05 - - -

2124 Rem 4.1 - 0.6 1.4 0.05 - - -

2048 Rem 3.1 - - 1.3 0.09 - - -

2618 Rem 2.3 - - 1.5 1.1 1.0 0.07 -

2219 Rem 6.3 0.3 _ 0.06 0.15

oo 19

2.2 The effect of alloying elements

Copper has been the most common and probably the most important alloying element since the beginning of the aluminium industry. It has appreciable solubility and a strengthening effect on the parent metal, the strength increasing with increasing copper content up to a maximum of 6%. Magnesium is used in combination with copper to accelerate and increase the age-hardening at room temperature. The addition of silicon improves the artificial ageing response and has significant strengthening effects at higher temperatures but at the expense of ductility8.

It has been reported that when the Cu:Mg ratio is above 2 and the Mg:Si ratio is about 1.7, the equilibrium phases present are CuMgAl 2 (S phase), Mg 2Si and CuAl2

(8 phase)9.

Iron in the alloy forms compounds such as CuFeAlg and Cu 2 FeAl 7 . It has a beneficial strengthening effect at higher temperatures but causes embrittlement and hence must be kept to a minimum'10. , too, has a marked strengthening effect and this is due to its solubility and the formation of several intermetallic compounds8 *11. During the casting of the alloy, most of the manganese remains in solution but during homogenisation it combines with iron to form intermetallic compounds such as (CuFeMn )3 and (CuFeMnJAly which in turn reduce the embrittling effect of iron7 *9 *12 .

The effect of manganese in solution is reported to retard the onset of recrystallisation whereas the manganese rich intermetallics, which are associated with the sub-boundary dislocation network, inhibit grain growth7 *8 *1 2 . The size and spacing of these particles, therefore, have an important effect on the mechanical properties of the material. Iron and manganese, being transitional elements do not influence the age hardening mechanism8.

The other minor additions include zinc and . Zinc increases the strength by entering solid solution while titanium acts as a grain refiner8 . 20

Possible equilibrium compounds that may occur in alloy 2024 are given in table

2 . 2 .

2.3 Preheat treatment

During the processing of 2XXX series alloys, several heat treatments are applied. These usually consist of an initial homogenisation prior to the hot working process (extrusion), followed by solution treatment and age hardening to give the required strength and ductility.

2.3.1 Homogenisation

The solidification rate during casting is rarely slow enough for equilibrium to be maintained, and compositional gradients (coring) exist within the dendrites. In addition, the last region to freeze may be enriched in non equilibrium, low melting point interdendritic eutectic or constituent phases which enhance cracking during subsequent working operations8 . Homogenisation treatments are normally given to reduce segregation, remove the low melting phases and thus improve workability.

Homogenisation also serves to precipitate dispersoid forming elements such as , manganese and iron so that they can perform their role of grain control during subsequent processing. The elements are often trapped in solid solution during solidification because of their low diffusion coefficients in the solid state and rapid cooling of the ingot. The treatment allows precipitation under controlled conditions and results in a more uniform precipitate distribution than can be obtained during solidification18.

The rate of heating to homogenisation temperature is important; relatively slow heating rates are necessary in order to dissolve low melting point eutectics and this prevents incipient melting from occurring at higher temperatures. The conventional homogenisation treatment for alloy 2024 is reported to be 24 hours at

5 0 0 °C 8 *1 1 *1 4 . 21

Table 2.2 Possible equilibrium compounds in alloy AA2024

Compound Structure Stable as solid °C

Al Cubic 660

CuAl2 Tetragonal 591

CuMgAl2 Orthorhombic 518

CuMg4Als Cubic 495

Mg2Si Cubic 1102

CuFeMnAIg Orthorhombic 654

Table 2.3 Transition Phases in alloy AA2024

Phase S tr u c tu r e a -A I F.C.C.

0"- C uAI2 Tetragonal

0'- CuAl2 Tetragonal

0- CuAl2 Tetragonal

S'-CuMgAl2 Orthorhombic 22

2.3.2 Solution treatment

Solution heat treatment is usually performed after hot working with the purpose of obtaining in solid solution, the maximum practical concentration of the hardening solutes such as copper, magnesium and silicon. It is most effective near the solidus or eutectic temperature, where maximum solubility exists and diffusion rates are rapid. However, care must be taken to avoid incipient melting of low temperature eutectics, and grain boundary phases15. Such melting results in quench cracks and loss in ductility.

The optimum solution treatment for alloy 2024 is reported to be between 490°C and 500°C for 20 minutes to 2 hours depending on the product thickness16*17. The material is rapidly quenched following solution heat treatment in order to maintain the maximum degree of supersaturation of solute elements and vacancies for subsequent ageing treatments. The importance of vacancies as heterogeneous nucleation sites is demonstrated by the presence of precipitate free zones adjacent to grain boundaries; where the vacancies in these zones have been absorbed by the grain boundaries18.

The highest strength and most often the greatest corrosion resistance is obtained with the fastest quenching rates. Water is the most widely used and efficient medium available. It has been reported that during quenching, dislocations appear as helices, loops or tangles generated as a result of quenching strains and condensation of excess vacancies5.

2.3.3 The mechanisms of aae hardening

The basis of age hardening is the formation during the decomposition of the super saturated solid solution (s.s.s.s.) of one or more metastable transition phases prior to the equilibrium precipitate. These precipitates develop sequentially, with increasing time at temperature between room temperature and the solvus. In the binary Al-Cu system, the three stages are identified as; 23

SSSS -*• GPI GPII(0") 0* -► 0(C u AI2)

GP1 phase is a zone structure and forms immediately at room temperature as plate like copper rich regions on the (100) planes of the aluminium rich matrix19. They are reported to be 30-50A in diameter and 100A apart. On raising the temperature above 100°C, the GPI zones are replaced by the GPU or 0" zones.

Structurally, they are similar to GPI zones except bigger. They also lie on the (100) plane and are coherent in the aluminium matrix20*21. On further raising the temperature, the 0 ” grows and the coherency with the matrix breaks and this leads to the formation of dislocation rings. It is at these rings that the semi coherent 0 ' nucleate and grow at the expense of the 0 n. The semi coherent 0 ' platelets are approximately 250A thick and lie parallel to the (100) plane. The final stage is the transformation of the 0 ' into the non coherent equilibrium 0 (C uA ^). The 0 nucleates at the 0'/matrix boundaries and consumes the 0 ' platelets at which it is nucleated. All three precipitate phases are reported to have tetragonal structures22.

The phases have been tabulated in table 2.3 and the Al-Cu phase diagram is shown in fig 2.1.

In commercial Al-Cu-Mg alloys the presence of magnesium and silicon modify the precipitation sequence depending on the Mg:Cu and Mg:Si ratio9. The other additives, iron and manganese, do not affect the age hardening directly, but affect the overall response by "tying" up the hardening solute via compound formation23. Silcock25 investigated the Al-Cu-Mg systems with Cu:Mg ratios of 2.2:1 and 7:1. At a ratio of 2.2:1, the ageing sequence is;

SSSS -► GPB -»■ S’(CuMgAI2 ) S

The formation of the GPB zone is not fully understood but it is known to be stable at temperatures up to 260°C with smaller lattice strains than the GP1 zone23' 26.

The transitional phase S"(GPB) is analogous to the 0 " of the Al-Cu system, and the 6 25

semi coherent S' (CuMgA^) consists of platelets coherent in the (021) aluminium matrix plane. The equilibrium S phase is orthorhombic and incoherent with the m atrix.

In alloy 2024, there is a possibility that either the Al-Cu or Al-Cu-Mg type transformation may occur. Bonfield and Datta28»27 linked the presence of silicon to the formation of GPB zones (Al-Cu-Mg-Si) as well as GP1 zones, during the initial stages of ageing at temperatures above 150°C. On further ageing the GP1 zones were succeeded by 0 " and 0 ' precipitates, while the GPB zones were persistent throughout the ageing sequence. Peak hardness was associated with partially coherent 0 ' and the subsequent limited overageing at these temperatures was controlled by the slow growth of these particles. They also suggested that at ageing temperatures between room temperature and 130°C, the transformation of GPB zones to S' precipitates may be favoured27.

2.3.4 The effect of aae-hardenina

The main objective of age hardening is to increase the strength. This increase in strength is, however, accompanied by a loss in ductility and fracture toughness28*29.

The microstructural features which affect the mechanical properties, in particular the fracture toughness, are the second phase particles and grain structure. Several workers have reviewed the effect of these particles according to their size7*30;

1) coarse particles formed during casting; 1p.m-10p.rn in diameter

2) intermediate particles formed during the homogenisation treatment;

0.05p.m-0.5p.m

3)ageing particles; 0.01-0.05p.m 26

Coarse particles are deleterious to fracture toughness since they crack easily promoting void growth and coalescence ahead of the crack tip3 0 . Furthermore, preferred orientation of these particles in the extrusion direction reduces the toughness in the short transverse direction30. Intermediate particles are not injurious to toughness but may influence it by retarding recrystallisation and grain growth thus controlling the grain size. Toughness and strength may be improved by a reduction in grain size7 *30 .

The fracture toughness decreases with increase in tensile strength on ageing31 and this has been attributed to a reduction in work needed to link voids prior to fracture. The loss of toughness is connected with the fine precipitate particles that form during ageing but the exact mechanism is not fully understood29*32. A decrease in toughness is also observed between the naturally and artificially aged conditions7 . This is partly due to the increase in strength and in part due to the detrimental effect of S' and 9 ' precipitates. It is reported that 2024 and 2014 exhibited higher toughness in the underaged condition than in the overaged condition7 .

Ageing has a significant effect on the corrosion resistance of aluminium alloys9. During ageing those areas of the structure which are in a more advanced stage of ageing are electronegative to the rest of the grain and so corrosion concentrates there. Maximum susceptibility to stress corrosion cracking, inter-granular corrosion and exfoliation are all greatest in the underaged condition, and least in the overaged condition28*33*34. Spiedal28 attributed this to the interaction of dislocations and precipitates during ageing. When the GP zones or precipitates are small, they are cut by the glide dislocations during plastic deformation. This is the condition at peak strength. After this, however, deformation tends to be on the active slip planes which is detrimental to stress corrosion resistance and fatigue strength.

Fine38 suggested that a duplex structure of small closely spaced particles to give high strength and large particles to distribute the plastic deformation, improved resistance to stress corrosion cracking and fatigue cracking. 27

2.4 Hot working

Hot working refers to deformation of a metal at elevated temperatures, usually greater than 0.6Tm (where T m is the melting point in degrees Kelvin). At elevated temperatures, the stress required to deform a metal at a given strain decreases due to the increased ease of dislocation movement. Furthermore, the rate of work hardening decreases as high temperature restoration processes become effective. It is these restoration processes that have been studied in investigations into hot-workng^7,38<

The approaches used in hot working analysis have been classified into two general areas of investigation

1) Empirical relationships: Equations relating flow stress, strain rate and temperature.

2) Structural analysis: structural changes which occur during hot working can aid the interpretation of empirical equations.

2.4.1 Empirical relationships

in order to investigate hot working, several techniques which can simulate high temperatures and strain rates associated with hot working operations are employed. These include the tensile test, compression test and torsion testing, all of which have been reviewed extensively by several authors16,39,40. The following relationships have been determined using data derived from these tests.

Two relationships can be applied to hot working data. For low stress at constant temperature;

e= A icnl...... 2.1

where A-| and n-j are temperature independent constants. For high stresses at 28

constant temperature;

£=A2exp(pa)...... 2.2 where A 2 and p are also temperature independent constants. Equation 2.1 and 2.2 are similar to those for creep, and from the work by Garafola41, these two equations can be combined into the form;

8 = A3(sinh(aa))n...... 2.3

At low stress levels (a a < 0.5) equations 2.3 reduces to a power form while at high stresses (a o > 1.2) it gives an exponential form.

The similarity between steady state creep and steady state hot working led Sellars and Tegart42 to propose a relationship of the form;

8 = A(sinh(aa))n exp(-AH/RT»...... 2.4

where A, a and n are temperature independent constants and AH is the activation energy.

Equation 2.4 can be rewritten in terms of Z, the Zener Hollomon parameter, the temperature compensated strain rate. Z = 8exp(AH/RT) = A(sinh(ac))n...... 2.5

The equation can now be applied to hot working data over a wide range of strain rates and temperatures.

Despite the accuracy of the Zener Hollomon equation in predicting the deformation response of aluminium alloys in the hot working regime, it does not relate the deformation behaviour to the varieties of metallurgical microstructure. To overcome this, Anand79 proposed a simplified approach to the notion of a mechanical 29

state equation by formulating a phenomenological theory with a single, non-zero, positive valued scalar (s) to represent the internal variables. In addition the kinetic equation used by Anand was that proposed by Lee and Zaverl80:

a = S.( (8/A).exp(AH/RT))m...... 2.6

where m is the material conatant

This equation was recently applied with great success by Sample and Lalli81 at the Alcoa Technical Centre for compression of commercial purity aluminium at hot working temperatures. They presented a relationship between the developed micro-hardness and the state parameter, which displayed good predictive capabilities. However, it is likely that such a relationship would be difficult to determine in alloys where contributions to hardness originates from numerous sources (e.g. solution hardening and grain size effects)

2.4.2 Structural aspects of hot working

The empirical relationship obtained earlier indicates that the flow stress is independent of strain in the plastic regime. This is further reinforced by high temperature stress strain curves, in which there is an initial rapid rise in stress followed by a steady state region. Hence, some high temperature restoration process must be operating to overcome the effect of work hardening. The type of restoration process will depend very much upon the material and in particular its stacking fault energy (S.F.E.). This is a measure of the ease with which dislocations can divide into partial dislocations, a situation which is energetically more favourable for cross slip motion. The higher the stacking fault energy the stronger is its binding energy and the closer together are the partials. Before dislocation climb can occur, the partials must come together under the action of the applied force, thus climb and hence recovery is easier in high stacking fault energy metals4 3 .

In high S.F.E. materials such as aluminium and its alloys, the steady state structure is characterised by subgrains of low misorientation, whilst in low S.F.E. materials, a recrystallised grain structure is found44-45. The difference in 30

structures shows that a different mechanism operates in each system.

2.4.3 Restoration processes

In high stacking fault energy materials, the dominant restoration mechanism is dynamic recovery. As deformation begins, the flow stress rises rapidly to a maximum due to the generation and multiplication of dislocations46. These dislocations are initially in tangles but soon arrange themselves to form rough cell-like structures by intersecting each other4 7 *4 6 . However, thermal activation assists the dislocations to cross slip and climb and either annihilate by mutual interaction, or rearrange by polygonization, into well defined walls. The increasing influence of dislocation annihilation and movement causes a reduction in the rate of work hardening until a steady state stress is reached, the rate of work hardening is in equilibrium with the rate of annihilation4 9 .

Dynamic recovery leads to the formation of sub grains. These are within the deformed original grains and have a low misorientation of approximately 1-2° between them. When the steady state flow stress is reached, the sub grains are equiaxed and have a steady state size indicating that they were continuously reforming during hot working. This process is termed "Repolygonisaton".

The mean sub grain size increases with increase in temperature and with decrease in strain rate37-*42 . The size can be controlled through the temperature compensated strain rate, Z, as the sub grain size produced during deformation is always the same at a given Z value. Hence by varying the strain rate and temperature the optimum operational and structural properties can be obtained.

In low stacking fault energy materials, however, two restoration mechanisms are available. At small strains, dynamic recovery is dominant, while at high strains dynamic recrystallisation occurs42. In these materials, the dislocation increases to higher levels than in materials that dynamically recover, until local differences in density are high enough to permit the nucleation of recrystallisation during deformation. Recrystallisation leads to the elimination of large numbers of dislocations by the migration of high angle boundaries. 31

Dynamic recrystallisation does not seem to have been found in commercially pure aluminium or in Al-Cu alloys, dynamic recovery being the only reported mechanism16,50-52.

2.4.4 Static restoration processes

If a metal that contains a substructure is heated, the stored energy gives rise to static restoration processes, which are of two types; static recovery and static recrystallisation.

Static recovery involves individual dislocation motions within existing grains, and although these dislocations undergo considerable rearrangement, the grain structure remains almost unaltered and the preferred orientation remains unchanged.

Static recrystallisation eliminates large numbers of dislocations simultaneously as a result of the motion of high angled boundaries, from a point of nucieation. If the material had a preferred orientation, the recrystallised material is also likely to have one, usually different but related.

If a metal reduces its stored energy through static recovery, then recrystallisation is less likely to occur. Therefore in materials where dislocation climb and polygonisation is relatively easy, the possibility of recrystallisation is reduced. Recrystallisation is enhanced, however, when second phase particles which accelerate the nucieation process are present.

2.5 Substructure strengthening

Under hot working conditions, the sub-grain size can be related to the steady state flow stress by the equation; 32

a » a 0 + kd*m...... 2.6

where a 0 , k, and m are constants and d is the sub grain diameter. During high temperature deformation, a 0 is approximately 0 and m ranges from 1 to 1.582.

Although the strengthening mechanism is not fully understood, it is thought to be due to sub boundaries opposing dislocation motion.

An alternative relationship involving the hyperbolic sine function was proposed by Jonas et al3 8 ;

d_1 = a + b. Ln(sinh(aa))...... 2.7

or

d_1 = p(sinh(aa))q ...... 2.8

where a, b, p and q are constants.

As the compensated strain rate is a function of hyperbolic sine, it can be introduced into the equation such that;

d*m = a + b.Ln(Z)...... 2.9

This enables the process variables, temperature and strain rate, to control the sub grain size and hence the hot worked materials' room temperature properties. 33

2.6 Extrusion

Extrusion is a process in which a billet is converted into a continuous length of uniform cross section by forcing it to flow under high pressure through a die orifice, which is so shaped as to impart the required form. Most metals are processed at elevated temperatures to reduce the influence of work hardening on the load and to prevent the occurrence of a brittle, cold worked structure in the extrudate.

Figure 2.2 illustrates the essential principle of the process, and, at the same time demonstrates the difference between direct and indirect extrusion.

In the case of direct extrusion, a billet in the container is forced by a ram through a die to give the product. For indirect extrusion, the container, which is sealed with a dummy block, moves together with the billet towards the die which is situated at the end of a hollow stem. The essential difference between these methods is that for indirect extrusion, there is no motion between the billet and the container, and hence minimal friction.

2.6.1 Direct extrusion

A typical load-ram displacement curve for direct extrusion is shown in fig 2.3. It can be divided into 3 stages. In stage I, the load increases rapidly as the billet upsets to fill the container. Stage II is referred to as steady state flow and shows a decrease in load with ram travel since the functional force generated by the movement of the billet relative to the container is reduced as the billet length decreases. Finally, at stage III, a rapid pressure increase is observed as the pressure pad impinges on the deformation zone. For this reason, a discard is left in the container and is ejected before the next extrusion.

The prominent peak in the load trace was proposed to be associated with the formation of the dead metal zone (D.M.Z.) by Avitzur et al3 3 , while Johnson et al34 linked it with frictional effects. Ziegler33 disproved both theories with the evidence of peak pressure in indirect extrusions, where friction is minimal and no dead metal Direct Indirect

M ------e x tr u d a te ------► u

Figure. 2.2 Principles of extrusion LOAD i 23 da la cre o extrusion for curve load Ideal 2-3 Fig 36

zone forms. The only acceptable theory is that put forward by Castle and Sheppard5 6 , who associated the peak with the energy required to establish a quasi static deformation zone. They proposed that the formation of the zone required an excess of dislocation sources leading to a greater than equilibrium dislocation density. This excess of dislocation production necessary to establish the steady state process would also imply a non equilibrium flow stress resulting in the requirements of excess pressure.

2.6.2 Indirect extrusion

A typical load versus ram displacement curve for indirect extrusion is also shown in fig 2.3. It, too, can be divided into 3 stages. During stage 1, the billet is upset and the load rises to a maximum. Stage II produces a slight decrease in load from the peak since the billet/container friction has been eliminated; there is no reduction in load with decrease in billet length. Finally, as stage III is reached, the load rises rapidly. The peak pressure, which also exists in indirect extrusion, is probably due to similar reasons as those for direct extrusion. The necessary pressure for indirect extrusion is considerably reduced and has been reported to be 60-70% of that of direct extrusion under the same conditions10.

Indirect extrusion without a shell without lubrication has many advantages over direct extrusion. The most important of these are as follows: 1. Greater productivity due to higher extrusion speed at a lower billet temperature, and the use of larger and longer billets. 2. Uniform grain structure and mechanical properties over the entire rod length. 3. Closer and more uniform dimensional tolerances.

Certain prerequisites are essential to achieve these advantages, the most important among which are: 1. Stricter requirements on the casting quality. Only flaw free, turned or scalped billets produce a good surface finish. 2. It must be possible to clear away any remaining extrusion film (shell) from the container after each pressing, and to vent after insertion of the billet. 37

From the above points it can be seen that indirect extrusion has certain limitations. If these are exceeded the quality of the product and the profitability of the process will suffer. The most important limitations are:

1. Friction between billet and die basically limits the extrusion speed, since the deformation heat from the zone ahead of the die must be dissipated into the surrounding metal and this takes time. Too high a surface temperature leads to hot shortness. 2. The billet temperature cannot be chosen too low since this can lead to undesirable changes in structure due to partial or total recrystallisation. This also limits the extrusion speed.

3. The billet length is limited not only by mechanical factors, but also by the fact that as the billet length increases, it clearly becomes more difficult to achieve good venting of the container, giving a much higher risk of producing surface defects and

blisters. 4. Solution treatment and quenching at the press are more difficult because of the length of the hollow ram, and thus the greater distance between tool and cooling section. Thus, quench-sensitive alloys cannot all be processed, or at least not in all dimensions. Moreover, for effective quenching at the press high extrusion temperatures are needed, and this also limits the extrusion speed. 5. The hollow ram restricts the useful cross-section, and thus the simple possibilities for multiple extrusion without having to use a composite, segmented

hollow ram.

To carry out indirect extrusion in the optimum way from the economic and metallurgical standpoints, many parameters that vary from alloy to alloy must be taken into consideration. It should be said, however, that direct extrusion was the subject of intensive development work with respect to the optimisation of alloy composition, heat treatment at high temperatures, extrusion temperatures, extrusion speeds and other details of the process, for many years. The relatively new method of indirect extrusion is a technique employing lower extrusion temperatures. It is to be expected that as well as further improvements to the method, there may well be important developments in the material used, that will lead to even better

results in indirect extrusion 38

2.6.3 Limit diagrams

In an attempt to define the limiting conditions under which an extrusion can be produced, Hirst and Ursell ®7 developed a graphical method of representing certain factors. Figure 2.4 shows a type of limit diagram where the extrusion ratio, ER, is plotted against the initial billet temperature. The limiting line on the left of the diagram represents the maximum pressure achievable by the press. Above this line, i.e. at low temperatures and/or high extrusion ratios, extrusion will not be possible as the required pressure for "break through" is not available; below the line, however, extrusion is possible. On the right hand side of the diagram another limiting line is produced, this represents melting of the extrudate during extrusion. Thus extrusions produced with extrusion ratios and preheat temperatures above this line will be unsatisfactory due to melting, while those below it will be satisfactory. The location of these lines are dependent on the maximum available pressure and on the speed of extrusion, thus, a limit diagram is dependent on many factors and applicable to only very specific conditions

Extrusion Temperature

Fig 2.4. Extrusion Ratio vs Temperature Limit Diagram

Since the pioneering work of Hirst and Ursell the usefulness of limit diagrams has been extended considerably. Numerous workers have shown how the diagrams 39

can be used to display structural information and how these metallurgical data can be applied to process control5 8 *59 '6 0 . It has been shown earlier that a relationship between the sub-grain size and the temperature compensated strain rate can be formed. This allows the microstructural variables to be incorporated into the extrusion limit diagram. For example, if lines of constant subgrain size are plotted, it is seen that constant subgrain size could be obtained by using numerous temperature-extrusion ratio combinations. It has been shown by Sheppard 61 that these substructural lines have the additional advantage of being representative of the room temperature properties due to their relation to subgrain size (and by extension 'Z'). In conclusion, results have shown that there are two principle disadvantages to this representation of structural information in heat-treatable alloys. Firstly, it appears that the relationship between process conditions (and hence subsequent structure and mechanical properties) is not so simple. Secondly, the structure limit diagrams and limit diagrams make no allowance for inhomogeneity in deformation and therefore give no information on the variation of structure across the extruded section (this is especially critical for shaped extrusion)

2.6.4 Structure produced by extrusion

Extrusion produces a marked directionality in the structure of metals that dynamically recover, due to the very high strains that are encountered. The original grains are elongated and particles aligned in the extrusion direction. However, the strain produced, is not uniform across the extrudate, with surface layers undergoing considerably more deformation at a higher strain rate in passing the face of the dead metal zone , and along the die lands. The higher strains in the surface may lead to a higher dislocation density which can often promote static recrystallisation. This leads to an annulus of coarse recrystallised grains in the extrudate.

The grain's directionality in the extrusion of aluminium, due to preferred orientation is [100] and [111], 75% being of the type [111]. The texture produced increases the strength of the extrudate in uniaxial tension from the [ 1 1 1 ] directional grains. Strength of the extrudate is decreased, however, when a recrystallised annulus occurs. 40

2.6.5 Shaped extrusion

The advantage of extrusion as a manufacturing process is that complex sections can be formed in one operation. Therefore it is important to understand the effect of section shape on extrusion load.

There has yet been no great study of the effect of the section shape upon the load. Several authors 62'64 have noted that the deformation resistance is dependent upon section shape, this becoming larger as the complexity of the section shape increases. They attributed this to the inhomogeneous nature of the material flow under such conditions. Honauer 62 suggested that the resistance to deformation could be evaluated by defining a shape factor such that:-

Kw= Kf fp fR...... 2.10

Kw » deformation resistance

Kf =* external friction

fp = shape factor

fp = friction factor

The friction factor is given a value of unity for a round bar. Kienzle and Gruning65 also suggested using a rod product as a base and demonstrated that as much as a 28% increase in pressure was required to extrude shaped sections in 99.5% aluminium than was necessary to accomplish the extrusion of a smooth rod. It is thought that this increase is in some way related to the geometry of the shape being extruded. This train of thought led to the proposal of a peripheral ratio, which is defined as the periphery of the section shape to the periphery of a rod of equal area. This method has a limited use as it does not account for the variation of shape with a constant periphery ratio.

Farag102 developed Johnsons 102 proposal using a concept of an effective extrusion ratio to predict the increase in extrusion pressure of shaped sections. The effective extrusion ratio is obtained by summing up the extrusion ratios over the 41

different meridian planes and then averaging these.

Sheppard and Wood6** proposed an alternative technique to allow for section shape. They found that when the generalised hot working formula was plotted logarithmically, a series of parallel straight lines were obtained, each corresponding to a different section shape. To normalise these points into a single line, a modified extrusion ratio was defined, which was then incorporated into an adapted form of the standard extrusion equation. Peripheral ratio was used to define the modified extrusion ratio (or shape correction factor) as follows:

R' = QR...... 2.11

where;

« = (Pe/Pr)0'5 = f° -5 pe = circumference of the shaped section pr = circumference of the rod section

R = extrusion ratio

By extension, the relationship between the extrusion ratio and Zenner-Holloman parameter was modified to the following:

LnR’ - a + b.LnZ...... 2.12

where; a,b - constants Ln R' * shape correction factor. 3.0 Theory 43

U __ Introduction

The analysis of extrusion through empirical and theoretical equations has been extensively reviewed by other workers16,17,66 ancj hence will not be repeated here.

One of the most important factors that must be known before any understanding of the behaviour of the metal during extrusion can be gained is its flow stress characteristics at high temperatures. There are several methods which include compression, torsion and tensile testing. The torsion test is however by far the most suitable method to obtain hot working, flow stress data to model the extrusion process.

3.2. Torsion analysis

It was shown earlier in section 2.4.1 that the flow stress behaviour in hot working can be expressed mathematically using the constitutive equation providing the four constants are known. The following theory was used to derive these from an experimental torsion data matrix.

To analyse the test accurately, it is necessary to allow for the heat generated within the specimen due to the work input. For this, a finite difference model was used allowing for heat conduction and convection.

The torque generated by the deformation of a torsion specimen can be written as;

M = a r2 d r .3.1

where r is the specimen radius. The flow stress substitution can be made by rewriting the constitutive equation and using the sinh’ 1 (x) expansion to give: 44

o = (1/a) Ln[(Z/A)1/n + {(Z/A)2711 + 1}1/2] ...... 3.2

Z contains the mean equivalent surface strain rate which can also be substituted,

6 = 2itr0/(V3 L)...... 3.3

L = gauge length

0 = twist rate in rev/secs.

The torque is now related to the twist rate and temperature by the four constants. This was solved by minimising the difference between the torque calculated in this way and the measured torque66*67.

3.3 Extrusion analysis

3.3.1 Temperature rise

The temperature rise during extrusion has been estimated using an integral profile model first developed by Sheppard and Wood 66 and later modified to accommodate for use in indirect extrusion16. Patterson 16 also introduced a modification to the temperature differential term to allow for the heat losses to the tooling during extrusion since the original model assumed a constant linear temperature of 300°C as opposed to the variable container temperature in the present work. In the results sections the temperatures and values of Ln Z quoted refer to the initial extrusion conditions, unless otherwise stated.

3.3.2 Strain rate

Obtaining a single value for the strain rate during extrusion is difficult as there 45

is a variation from point to point in the deformation zone. In the present work a single triangle axisymmetric upper bound computer model developed at Imperial College was used to obtain a mean strain rate. This allowed for energy dissipation due to velocity discontinuities, circumferential straining and container billet friction. For direct extrusion the mean strain rate is comparable to the modified Felthams strain rate evaluated by Tutcher and Sheppard®9, whilst the indirect strain rate is higher at equivalent extrusion ratios and ram speeds due to smaller deformation zone size. Further details are outlined by Vierod 1 7 and Patterson 1 ®.

3.4 Evaluation of the billet container friction in direct extrusion

During direct extrusion the force required to overcome the friction between the billet and container results in a marked increase in the overall pressure to extrude.

The friction may be defined in one of two ways:-

1 ) the friction coefficient, p, where;

X = jiP ...... 3.4

X is the tangential stress acting opposite the motion and P is the pressure between the bodies.

2) a constant friction factor, m, where;

X - m(a0 /V3)...... 3.5 a 0 is the flow stress of the plastic material and m ranges in value 0 < m < 1 .

The coefficient of friction p can be experimentally evaluated by one of two

methods: 46

1) Measuring the pressure to extrude two billets of different lengths L-j and l _2 under identical extrusion conditions, where \i is given by5 7 :

n = ~ 3 _ . Ln(Pi/P2)...... 3.6 40-i-L2)

where D b is the billet diameter.

2) The rate of change of pressure with billet length can be used to evaluate \l using the relationship77*78:-

Dq dP ...... 3.7 4P dx

where P is the ram pressure and x the axial distance along the container.

The friction factor m can be calculated using the following relationship derived from the single triangle upper bound analysis 16 *71

m - V3 (Pi - P2 ) R b ...... 3.8 2 a (Li - L2)

where R b is the billet radius.

The accuracy of the friction factor m is largely dependent upon the value of stress a chosen to represent the billet material at container interface, where theoretically the strain rate is zero. A mean strain rate derived from the adjacent layers must therefore be used to determine a value of flow stress, so that m can only be considered as an approximate guide to the friction conditions. 47

4.0 Experimental Procedure 48

4.1 Introduction

The following chapter describes the techniques developed and briefly outlines the experimental procedures used to carry out the investigation, from the starting material, through process monitoring, to final product properties.

4.2 The extrusion press

Extrusion was performed on a 5 MN press operating with tooling set up for direct and indirect extrusions. Extrusion ratios varied between 20:1 and 50:1, ram speeds between 3 and 15 mm/s, and initial billet temperature between 250 and 480°C . The billets were 75mm diameter x 95mm long and were heated in an induction heater or in an air circulating furnace.

The load was measured by a Mayes load cell situated directly above the ram, the output from the cell being recorded on a Labmaster. Output from a pressure transducer situated at the inlet to the main cylinder was also recorded in order to check load measurements. Ram speeds and displacement were measured by a rectilinear potentiometer fixed between the moving crossheads and the press bolster which transmitted to the Labmaster.

The container was hydraulically lowered into position and the ram removed to its highest point. Two semi circular rings were placed on top of the container to prevent any damage to the main ram.

The hot billet was transferred from the induction heater into the container. A pressure pad was dropped on top of the billet. The ram was then lowered, initially under a fast approach and then at a predetermined speed during the extrusion cycle.

The ram, followed by the container, was then raised allowing the extrudate to be cut and pushed into the quench tank. The discard was then removed by raising the container, and pushing it out slowly using a tight fitting scrapper pad in front of the main ram. Fig. 4.1 General layout of the extrusion press and the direct and indirect tooling 7777 1 |d w 1 j Hv II Off* j 50

The procedure for indirect extrusion was essentially the same except that the 75mm ram was removed from the main ram and immediately prior to extrusion the container was raised such that the die assembly at the top of the mandrel was positioned in the bottom 50mm of the container. Upon transferring the preheated billet, the dummy block was placed in the container and the extrusion cycle was initiated as for the direct extrusion. When the main ram hit the dummy block both the billet and container were pushed down onto the mandrel and moved simultaneously at the predetermined speed during extrusion.

The layout of the extrusion press and the direct and indirect extrusion toolings are shown in fig. 4.1.

4.2.1 Tooling

The geometries of the dies used for rod extrusion are shown in plan and section view in fig 4.2. The size of the die orifices used for other shaped extrusions are shown by plan view only in fig 4.3. A serious limitation on the ability to extrude shapes in the indirect mode was the maximum allowable circumscribed circle diameter (C.C.D.). This restricted the minimum reduction ratio to 40:1 to avoid back up of aluminium in the tooling stem as a result of flexing of the emergent extrudate.

4.3 Material

The material used for the current research programme was supplied by Alcan Labs, Banbury, In the form of semi continuous logs of 86mm diameter. The quoted composition of the cast material is given below in table 4.1.

Cu Mg Si Mn Fe Zn Ti Al

4.66 1.35 0.08 0.69 0.19 0.02 0.009 Balance

Table 4.1 Composition of cast alloy AA2024 (wt %) 51

ii i p>- CD o> T CM

I

Extrusion a b Ratio

20:1 16.32 20.41 30:1 13.69 17.78 40:1 11.54 15.63 50:1 10.32 14.41

(all measurements in mm)

Fia 4.2 The dies used for rod extrusion i 43 is o sae extrusion shaped for Dies 4.3 Fig

10.57 <4 ------10.57 ► 52 53

The billets were homogenised prior to extrusion at 500°C for 24 hours and furnace cooled. The homogenised billets were then machined to a diameter of 73mm and cut into billets of required length.

4.4 Examination of extrudatea

Specimens for mechanical testing, for heat treatment and for optical and electron microscopy were cut from a position one third along the extrudate in order to ensure steady state conditions. Specimens were also taken along the length of the extrudate to determine the range of properties. The final 60cm of the extrusion was never used since this region remained unquenched.

4.4.1 Heat treatment of extrudatea

The typical heat treatment given to alloy 2024 was a solution treatment at 500°C for 1 hour in an air circulating furnace, followed by quenching to room temperature and subsequent age hardening in an air circulating oven. Hardening versus time characteristics were determined for solution treated and press quenched extrudates by artificially ageing at temperatures of 120, 160, and 180°C for up to 720 hours. Standard treatments of 8 hours at 180°C or 18 hours at 160°C were adopted for most applications. The classification of the various heat treatments is given below in table 4.2.

4.4.2 Hardness taste

Hardness tests were carried out on a number of extrudates in the as extruded, solution treated, naturally aged and artificially aged conditions. Transverse specimens cut from the extrudates were tested at the edge, mid radius and centre using a standard Vickers hardness machine fitted with a 10Kg weight. 54

Table 4.2 Classification of heat treatments used

Description of Heat Treatment Aluminium Association Temper Designation

As-extruded, quenched F As-extruded, quenched, natural age T1 Solution treated, quenched, natural age T4 As-extruded, quenched, artificial age T5 Solution treated, quenched, artificial age T6 As-extruded, quenched, stretched (1%), artificial age T51 Solution treated, quenched, stretched (1%), artificial age T510

4.4.3 Optical microscopy

Both longitudinal and transverse sections were cut from the extrudate and prepared for optical microscopy. The sections were mounted in Metaserv S.W. resin and ground on 220-1200 grade silicon carbide paper and polished using 6-1 micron diamond compound spray.

The etchants used to show grain structure and precipitate distribution were:

Grain structure (Barkers Reagent) 46 mis HBF

7 grams Boric Add

970 mis H 2O

With the spedmen as the anode and using a stainless steel cathode, a potential of 20 Volts was applied with a maximum current density of 0.2 A/cm. The specimen was then observed under crossed polarisers.

Precipitate distribution (Kellers Reagent) 2ml HF 3ml HCI 55

20ml H N 03

175 ml H20

The specimen was immersed for 10-15 seconds and washed in a stream of water and finally blown dry.

4.4.4 Electron microscopy

Transverse and longitudinal sections of 2-3 mm thickness were cut from the extrudates and ground to a thickness of 0.25 - 0.3 mm on the silicon carbide paper. 3 mm discs were punched from these sections and electropolished in a commercial Struers jet thinner using a solution of 30% nitric acid in methanol. The solution was

maintained at -30°C and a potential of 13 volts was applied between the discs and solution. The specimens were then examined in a high resolution Philips EM 301 100KV transmission electron microscope, (TEM), and in a JEOL Temscan ( 100KV) with quantitative EDAX analysis.

Scanning electron microscopy (SEM) was used to examine both structural details and topographical features of the extrudate surface. Structural examinations were made in a JEOL T200 (25KV) SEM using a backscattered electron imaging (BEI) detector. Samples for the SEM structural study were polished using 0.25pm diamond compound spray and examined unetched. In both imaging modes an EDAX qualitative analysis probe was available to assist identification of second phases.

4.4.5 Tensile testing

Standard No. 14 Hounsfield specimens were machined from extrudates in the as extruded, solutionised and artificially aged conditions. Tests were conducted on an Instron universal testing machine at a cross head speed of 0.5 mm/min. Elongation values were measured from the load-time traces and proof stress values quoted are for a 0.2% offset.

As the type of fracture of the tensile specimen was considered to be important, 56

the fracture surfaces were examined under the scanning microscope.

4J4^6__Fcacture toughness tests

The fracture toughness test used was developed by ALCOA and called the "Terratek Short Rod Fracture Toughness Test"70. The specimens to be tested were machined from the longitudinal section of the round bar extrusion (R«20:1, v - 6.9 mm/s) to the specifications shown in fig 4.4. Every effort was made to obtain specimens at the same location as those for the tensile tests. The fracture toughness was then calculated using the following equation;

kv -A.F*B-(3/2) ...... 4.1

A * calibration constant - 22 F = maximum Load B « specimen diameter

All tests were carried out at room temperature on an Instron Universal Testing Machine using a cross head speed of 0.2 mm/min. The fracture surface was also examined under the scanning microscope.

4.5 Torsion teating

Specimens for torsion testing were machined from 10:1 rod extrudates. The testing machine was capable of surface shear strain rates of up to 50s'1 on the specimen geometry used. Tests were performed over a range of shear strain-rates from 0.05s*1 to 50s*1 and over a temperature range of 350 to 500°C.

An induction coil around the gauge section heated the specimen to the test temperature using a Eurotherm programmable controller. The heating rate for all the tests was set at 100°C per minute. After testing to failure, the specimen, enclosed in a sliding, sealed container, was quenched with a water jet to achieve rapid cooling. 57

F1g.4.4 Short rod fracture toughness test sample 58

The torque generated was detected by a lever arm operating a load cell, and the rotational displacement measured by a transducer, the outputs being amplified and subsequently recorded using a programme written by Clode66 for an IBM Labmaster PC.

4.6 Partially extruded billets

To examine the flow patterns during steady state direct and indirect extrusion, billets were partially extruded to a position of quasi-static deformation.. The partially extruded billet was then rapidly removed from the container and water quenched. The billets were then sectioned along the longitudinal axis, polished and etched to observe the flow pattern. However, it was found that due to the very fine grain size no flow patterns could be easily distinguished.

To overcome this difficulty, super purity Aluminium billets were partially extruded, sectioned and polished and gave a superior representation of the flow pattern. Disc specimens of 3mm diameter were then punched from locations of interest in the original alloy to be observed under the TEM.

4.7 Surfaw quality

The surface quality Is of prime importance in many industrial extrusion processes. Extrudates were assessed for quality and categorised as follows

A - Good surface finish along the entire length B - Surface cracking along part of extrudate length

C - Surface cracking along entire extrudate length. Results & Discussion 5.0 Torsion Analysis 61

5.1 Torsion data analysis

The analysis of any extrusion process requires information about the hot working characteristics of the material to be extruded. To evaluate these characteristics, a small scale testing technique with a known stress system such as hot torsion testing is employed.

The torsion test has the advantage that high strains can be applied at elevated temperatures and the results can be used to model the extrusion process. The test can also be used to assess a material's hot workability by measuring the strain to failure.

In this section, data from the hot torsion tests is analysed using several programs developed at Imperial College16,6 6 ,67 # j h 0 (jata jS then presented in the form of constitutive equations. The Zenner-Holloman equation is subsequently used in the following sections to analyse the extrusion process.

5.2 Torque/Twlat curve*

The torque/twist curves from the matrix of test conditions produced two characteristic sets of curves as shown in fig 5.1.

The first, which is typical of tests at lower strain rate conditions (< 5s*1) over the entire temperature range is shown in fig 5.1 A. It shows an initial rapid rise in torque corresponding to the interval of microstrain deformation, where the strain rate rises from zero to that applicable to the test. The gradient then falls off as recovery mechanisms driven by the deformation energy take an increasing effect. The peak torque corresponds to the steady state situation where dislocation generation and annihilation processes are balanced. It was also observed that the interval of work hardening becomes more prominent with reduction in temperature.

The second type of curve, as shown by fig. 5.1 B is typical of tests at low temperatures and high strains rates (high Z). It shows a similar initial rise but a slower decrease. The initial stage of microstrain deformation is followed by a decreasing work hardening rate prior to a fall in torque and frequent failure. The CU RVE TV 63

steady decrease in torque after peak is due to the conversion of energy input to heat, resulting in a reduced flow stress. The drop in torque can be attributed to the near adiabatic heating due to high strain rate deformation.

5.3 EIow stress characteristics

The torque-twist records were converted into equivalent stress-strain curves using the effective radius analysis. This method of conversion is detailed by Clode68 . The curves obtained by this method are shown in figs. 5.2-S.4.

The stress flow curves have a similar shape to the torque-twist curves reflecting the general characteristic of restoration by dynamic recovery. They also show the expected trend of increasing stress with decreasing temperature at constant strain rate, and with increasing strain rate at constant temperature. The strain to peak decreases with increase in temperature and corresponds to the twist to peak torque. It is worth noting that the stress data at higher strains is liable to increasing error due to the temperature rise effect which occurs during testing. To overcome this, peak values of stress which correspond to the onset of the steady state region are taken. The drop in the flow curve after peak stress is probably due to the temperature rise but it has been reported that other mechanisms such as dynamic recrystallisation and precipitate coarsening may also be the cause3 8 .

There are several factors which affect the ease of dynamic recovery during deformation and therefore the flow stress characteristics. The two most obvious are the process parameters of temperature and strain rate. These effects have been well documented3 7 *38 where the relationship between them and the flow stress can be accurately described by either the general hot working equation or the general stress-strain relationships as shown in chapters 2&3. An increase in temperature reduces the flow stress as the thermally activated cross slip and climb increase the rate of recovery. An increase in strain rate increases the dislocation generation rate and effectively reduces the time for the recovery process resulting in a less recovered substructure and higher flow stress. The two conflicting parameters can be combined into a single parameter, Z, the temperature compensated strain rate. Previous workers have shown that it can be used to predict a specific substructure 140

Equivalent Strain

Fig. 5.2 Stress-Strain curves for torsion of alloy AA2024 Fig. 5.3 Stress-Strain curves for torsion of alloy AA2024 LfxON Fig. 5.4 Stress-Strain curves for torsion of alloy AA2024

ON O n 67

or flow stress for any combination of temperature and strain rate16’3 8 ’3 9 *71.

5.4 Constitutive equations

The flow stress behaviour with temperature and strain rate can be concisely described by the constitutive equation once the four constants have been established. This equation is generally arranged in the form which gives a temperature

compensated strain rate parameter, Z;

z = Eexp(AH/RT) = A(sinh(aa))n ...... 2.5

where;

AH = activation energy of the rate controlling mechanism a,n ,A = hot working material constants

R = Universal gas constant

The four material constants, AH, a,n, and A in the equation were determined using a program developed at Imperial College by Wright8 7, Vierod18 and Clode88based on the theory in chapters 2&3. The values given below in table 5.1 are effectively an average over the range of test conditions used.

Units AA2024 AA2014 (ref 17)

AH KJ/mole 146880 1 4 4 4 0 8

a m2/M N 0 .0 1 6 0 .0 1 5

n - 4 .2 7 5 .2 7

LnA . 1 9.6 24.11

Table 5.1 The hot working constants 68

The activation energy for creep of aluminium has been quoted as 150KJ mole'1 whilst that for lattice self diffusion has been given as 142KJ mole-1 for polycrystalline aluminium72*73. These figures confirm that the hot deformation of aluminium is a thermally activated process, where AH is equal to the activation energy of the rate controlling mechanism. A value of 146KJ mole-1 found for alloy

2024, indicates that similar processes are occurring in this material. This confirms that dynamic recovery is the overridding softening mechanism operative during deformation. Therefore it can be said that the magnitude of AH represents the size of the barrier preventing atomic rearrangement. However, it must be noted that other mechanisms may be occurring concurrently whose energy barriers are lower than AH.

The difficulty in interpreting the rate controlling mechanism for the values of

AH is that the remaining constants, A, a and n do not have a simple physical interpretation. However, an attempt has been made in comparing the results established for AA2024 with that of other workers.

Jonas et al defined n as a measure of the strain rate sensitivity of the material3 7 . It is reported to be in the range of 2-6 and to decrease with increase in solid solution alloying3 7 . A value of n=3 obtained from high temperature creep of solid solution alloys has been associated with the mechanism of solute atmosphere drag and viscous glide73*74. A value of between 4 and 5 has also been reported for pure metals75. Hence, the value of n computed for the alloy of 4.27 is consistent with this theory and can be associated with a climb controlled creep process.

The commonly used definition of a is that it is the reciprocal stress at which the temperature compensated strain rate changes from power to exponential stress dependence37. Hence it is dependent on the alloying additions and material purity.

The value of a determined in the analysis (0.016) agrees well with data obtained by earlier workers for a similar alloy16*17. 69

Jonas et al considered A to be a structure factor at high stress levels, being proportional to the density of activatable sites3 7 . As such it should be dependent upon alloying additions and material purity, the values of A here being consistent with many reported values for both aluminium and its alloys16’17'6 6 '7 1 .

In conclusion, the use of the computer programs to determine the material constants is a significant advance in the analysis of the hot working characteristics of the material. The laborious and error producing method previously used has been superseded by assessment of material constants directly from a matrix of torque-twist rate test temperature data. It also demonstrates the accuracy of Clode's66 assessment of the effective radius analysis.

5.5 Torsion structures

All the quenched steady state structures consisted of well formed subgrains, typical of a material that undergoes dynamic recovery during hot deformation. Some static recrystallisation was observed but this was not extensive, due to the rapidity of the quench. The structures will be discussed further in chapter 7 where they are compared with extruded substructures. 70

5.6 Conclusions

1. The flow stress data obtained from the torsion tests has been correlated over a range of strain rates and can be described by the hot working equation;

Z = eexp(AH/RT) = A(sinh(aa))n

2. The values of the constants in the above equation were found to be consistent with results of other workers and the activation energies obtained indicate that the hot deformation is thermally activated and diffusion controlled. 6.0 Extrusion Analysis 72

6.1 Extrusion data analysis

The following data was extracted from each extrusion run for analysis;

1 ) Extrusion ratio

2 ) Ram speed

3 ) Initial billet temperature 4 ) Extrusion pressure

5 ) Billet length

A temperature range of 300°C to 500°C, exit speeds of 300 to 500 mms*1 and extrusion ratios from 20:1 to 50:1 were utilised. The process matrix was repeated for direct and indirect extrusion of shapes. The strain rate and temperature rise during extrusion have been calculated using a computer programme developed by several workers and outlined by Cooper76 and Vierod17.

A full listing of all the above data is given in appendix 1.

6.2 Load-Displacement curves

Typical experimental load-displacement curves for direct and indirect extrusion are shown schematically in figures 6.1 & 6.2. All the curves show the initial stages of extrusion to be characterised by a rapid increase in load as the billet is upset to fill the container. During this stage a very small amount of extrusion takes place, reported to be of the order of the extrudate diameter in length6 6 . Once the peak is reached there is a steady fall in load as the billet length decreases. The extent of this fall in load is seen to decrease as the initial billet temperature increases. This trend is probably due to the increased heat generation in low temperature extrusions, which causes a rapidly decreasing flow stress and resultant lower extrusion loads.

The indirect process shows a much flatter and lower curve than the direct extrusion for identical process conditions. This is primarily due to the removal of the container wall/billet friction in this mode. It has also been reported that the 73

Fig 6-1 Typical load curves for direct extrusion

indirect extrusion 74

larger deformation zone in the direct mode may require more pressure to produce a larger number of excess dislocations to accommodate deformation66. In the indirect mode the initial rise to peak is followed by either a steady load at low temperature , or a slowly increasing load at higher temperatures. This interesting feature is probably the consequence of temperature losses in the billet, which are either cancelled out at low temperatures by the heat generation, or produce an increased load at higher temperatures. These trends point to the decreased heat generation in the indirect process, due to the absence of friction at the container/billet interface, and to the increased heat losses which result from the elevated position of the container on top of the mandrel.

6.3 Effect of extrusion parameters on the pressure to extrude

When evaluating the two extrusion modes, it is necessary to compare the relationships that exist between the extrusion parameters. Hence, the dependence of the peak pressure on the temperature, ram speed and extrusion ratio will be discussed.

6.3.1 Extrusion temperature

The effect of the initial billet temperature on the peak pressure is shown in figures 6.3 & 6.4 for both modes of extrusion. The data was plotted for a given ram speed and extrusion ratio. The statistical data for the plots is given in table 6.1.

As expected, a general trend of increasing extrusion pressure with decreasing initial billet temperature (as the flow stress increases) was found, but the relationship was not considered linear. Although previous workers have reported linear39, reciprocal51 and exponential relationships78, the use of a power relationship of the form;

P = ATn...... 6.1 was favoured as an excellent correlation was obtained for the data plotted. Vierod17 iue . Pa pesr v iiil ilt eprtr fr iet extrusion direct for temperature billet initial vs pressure Peak 6.3 Figure 5 Q. Peak Pressure CO 1100 500 0 0 800 700 600 T(K) T (K) T 75 76

Figure 6.4 Peak pressure vs initial billet temperature for indirect extrusion 77

and Clode66 also favoured this relationship and argued that earlier workers had employed constant container temperatures leading to errors in temperature calculations. It is now common practice to maintain the container at 50°C below the billet temperature, hence producing a more uniform heat balance at each extrusion temperature.

P = aTn

Extrusion D irect/ Ram a n CC Ratio Indirect Speed E8

20 direct 3.00 20.41 -2.303 0.99 20 direct 6.90 3.78 -2.024 0.99 20 indirect 3.00 131.41 -2.618 0.99 20 indirect 6.90 17.29 -2.284 0.99

30 direct 3.00 7.53 -2.137 0.99 30 direct 6.90 2.23 -1.931 0.99 30 indirect 3.00 11.44 -2.229 0.99 30 indirect 6.90 6.43 -2.1 18 0.99

50 direct 3.00 1.71 -1 .888 0.99 50 direct 6.90 1.23 -1.831 0.99 50 indirect 3.00 5.37 -2.089 0.99 50 indirect 6.90 3.72 -2.026 0.99

Table 6.1 Temperature dependency of peak pressure in alloy 2024

The similarity of the results in table 6.1 & 6.2 imply that the pressure can be described in a similar manner to the flow stress characteristics; this is consistent with the basic hypothesis that extrusion is a thermally activated hot working process similar to hot torsion testing. However, the total peak pressure is not a simple function of the steady state flow stress. The absence of a direct correlation demonstrates the difficulty in predicting the peak pressure from the flow stress characteristics. This may be partly explained by the non-homogeneous nature of the deformation during extrusion which necessitates defining a mean temperature and strain rate and that the work done against friction and the incremental increase in 78

pressure

a = aTn

Extrusion Direct/ Ram a n cc Ratio Indirect Speed E7 (mm/s)

20 Direct 6 .9 3.46 3.08 0.993 30 Direct 6 .9 1.87 2.78 0.995 50 Direct 6 .9 1.31 2.62 0.995

20 Indirect 6 .9 2.70 2.96 0.994 30 Indirect 6 .9 1.51 2.69 0.995 50 Indirect 6.9 1.09 2.54 0.996

Table 6.2 Temperature dependency of the flow stress

associated with the development of the steady state deformation zone must also be allowed for. The pressure dependencies shown in table 6.1 will therefore also be dependent on the effect of process conditions on the friction pressure and peak height, which will be considered in the following sections. As one might expect, the peak pressures obtained at higher ram speeds are greater than those at the lower ram speeds, and the rate of reduction of the peak pressure is also greater due to the

increased temperature rise.

A similar set of data is shown for indirect extrusion. It is noted that in all cases the load was lower as expected. At lower ram speeds, the indirect peak pressures are approximately 83MPa lower than the direct peak pressures over the whole temperature range. At higher ram speeds (6.9 mm/s), the difference is 125MPa at

300°C , but only 95MPa at 450°C. This difference is probably caused by the fact that as the temperature is increased, the heat losses through the hollow stem of the indirect tooling become more significant than the heat losses through the die assembly of the direct tooling. This behaviour indicates that the benefit of indirect extrusion, as regards to pressure, is likely to be more pronounced at higher speeds and lower temperatures. 79

The usefulness of this kind of data interpretation is that for a given press capacity, the minimum temperature at which extrusion is possible may be determined for a given reduction ratio and ram speed.

6.3.2 Extrusion ratio

The effect of extrusion ratio on the peak pressure is shown in figures 6.5 & 6.6 for direct and indirect extrusion at a given ram speed and temperature (which is assumed to be constant). This assumption is not strictly true, as was shown by

Vierod17, but is compensated for by maintaining the container at 50°C below the billet temperature. The results show an increase in the peak pressure with increase in extrusion ratio and the statistical data in table 6.3 indicate a linear relationship in the form of;

P= A + B.lnR...... 6.2 with good correlation at all temperatures. This is in accordance with data presented by previous workers 1 6,1 7,66,71

The value of B has been suggested as a measure of the pressure sensitivity to the magnitude of deformation by Clode®®. Hence the lower values of B in indirect extrusion can be explained by the nature of material flow which is more homogeneous and has a lower degree of redundant deformation.

The constant A decreases with temperature and hence appears to represent the pressure dependence upon temperature. The constant 'A' values are also smaller at lower ram speeds and this demonstrates that the pressure increases with both decreasing temperature and increasing ram speed.

Although both modes of extrusion show similar trends in the dependence of peak pressure on the extrusion ratio, the difference between them reduces as the initial billet temperature increases. This reduction in pressure difference between Peak Pressure (MPa) Peak Pressure (MPa) iue . Pa pesr v L R o drc extrusion direct for R Ln vs pressure Peak 6.5 Figure LnR 80 81

© 2a

w© 3 (0 M © afc.

© © a

a Q.

a © CL

Figure 6.6 Peak pressure vs Ln R for indirect extrusion 82

extrusion modes is a direct result of the increased thermal energy allowing the dislocations to cross slip and climb more readily, reducing dislocation storage and the redundant work prevalent in the direct extrusion mode.

P = A + B LnR

Mode Temp Ram Speed AB CC C mm/s

Direct 4 5 0 6.9 2 9 3 1 09 0.99 Direct 4 0 0 6.9 3 9 4 1 1 0 0.98 Direct 300 6.9 648 115 0.98

Direct 4 5 0 3 2 5 1 6 7 0.99 Direct 4 0 0 3 121 1 6 9 0.99 Direct 3 0 0 3 3 9 8 1 6 7 0.99

Indirect 4 5 0 6.9 2 0 7 1 0 2 0.98 Indirect 4 0 0 6.9 3 5 6 8 6 0.95 Indirect 3 0 0 6.9 361 1 0 2 0.98

Indirect 4 5 0 3 - 5 3 1 5 9 0.99 Indirect 4 0 0 3 3 9 1 60 0.99 Indirect 3 0 0 3 2 6 4 1 66 0.91

Table 6.3 Effect of extrusion ratio on the peak pressure

6.4 The variation of peak pressure with Z

The separate relationship generated for the variation of peak pressure with

temperature at different ram speeds can be overcome by considering the variation in terms of the Ln Z parameter. The temperature compensated strain rate at peak, Zp, incorporates both strain rate and temperature, which should enable the parameters to be related by one relationship. 83

The pressure could then be related to Ln Zp by the simple equation;

P = A + B.Ln(Zp)...... 6.3

where A and B are constants.

The results obtained are shown statistically and schematically in table 6.4 and figure 6.7 respectively.

P = A + B.LnZp

Mode AB cc

Direct -905.8 59.8 0.98

Indirect -954.2 56.8 0.97

Table 6.4 The dependence of peak pressure on Ln

V ierod17 also suggested modifying the equation to include the hot working constant A to deal with different alloy compositions. The modified equation is then;

P = A' + B'.Ln(Zp/A)...... 6.4

The results obtained for this analysis are shown in figure 6.8 and tabulated in table 6.5.

The high correlation obtained throughout the analysis indicates that both relationships are applicable to all temperatures and strain rates considered. One advantage of this concept is that it gives the freedom to allow for the variation in peak pressure with reduction ratio and even alloy variations. iue . Pa pesr v L Z o drc ad niet extrusion indirect and direct for Z Ln vs pressure Peak 6.7 Figure Peak Pressure (MPa) Peak Pressure (MPa) 84 85

a.a

a o a.

Ln (Z/A)

Figure 6.8 Peak pressure vs Ln(Z/A) for direct and indirect extrusion 86

P = A* + B\Ln(Zp/A)

Mode A B cc

Direct 265.9 59.8 0.98

Indirect 158.4 56.8 0.97

Table 6.5 The dependence of peak pressure on LnfZg/Al

6.5 Variation of AP with Log

Castle and Sheppard^® proposed that the excess pressure (AP) for the initiation of steady state extrusion was associated with the energy required to establish a quasi static deformation zone. The formation of the zone requires considerable dislocation generation and movement, and the pressure associated with this may therefore be expected to vary with temperature, strain rate, reduction ratio and alloy composition. Since the concept of Z is founded in hot working theory, it is reasonable to assume that the size of the peak is also related to log Zp. The values of AP have been measured using the tangential construction shown in figure 6.9. Previous workers have shown that the magnitude of AP could be directly related to Ln Z at peak in a similar fashion to the peak pressure in the previous section16,17,66,71 #

The results of using this relationship are shown in table 6.6 and plotted in figure

6.10 for both modes of extrusion.

The figure shows that for a given value of Z, the peak pressure obtained for direct extrusion is greater than that for indirect extrusion, which can be attributed to the different features of the two processes. It is evident that the size of the peak increases with increased extrusion pressure but it was found that the ratio of AP to extrusion pressure was greater for direct extrusion and so the noted difference is not ih h wr dn i oecmn te iltcnanr rcin A similar A friction. billet/container the overcoming in done work the with f n nrmna pa drn idrc etuin mle ta A i nt associated not is AP that implies extrusion indirect during peak incremental an of oey u t te rae pesrs soitd ih h drc poes Te existence The process. direct the with associated pressures greater the to due solely

Pressure Table 6.6 The dependenceof peakheight APon LnfZjJ Fia6.9 Peak height AP tangentialconstruction Direct Indirect Mode -203.8 -225.8 AP= A B.Ln(Zp) + B A 10.32 9.28 CC 0.97 0.96 87

88

(0 2CL © 3 0) 0)

a a

Xmo 3 « M © a. a © Q. <

Figure 6.10 A Peak pressure vs Ln Z for extrusion 89

observation in alloy 2014 led Vierod17 to postulate that a smaller AP for indirect extrusion was related to the geometrically smaller deformation zone, which required fewer dislocations to accommodate the imposed strain gradients and hence the energy required to establish the deformation zone was smaller.

6.6 Evaluation of the friction conditions

The results in the previous sections have dealt with extrusions with a constant billet length, 95mm being the usual length. In this section the effect of the billet length on the peak pressure will be studied. The variation of peak pressure with the billet length is shown in figure 6.11 for the alloy extruded at identical strain rate conditions at two different extrusion ratios and temperatures. A linear rise in peak pressure with billet length is noticed and this is confirmed by the regression data shown in table 6.7. Previous workers have reported a similar observation above a minimum billet length equal to the size of the deformation zone, below which the pressure increases7 1 *7 7 . The results obtained for this alloy do not reveal this feature and indicate that the deformation zone and flow conditions are similar for all the billet lengths considered. The values of p. obtained using equation 3.6 are also shown in table 6.7.

P = a + c.L

Extrusion Temp a C cc m Ratio K

20:01 698 1.61 418.8 0.99 0.058 0.911 20:01 623 1.67 589 0.99 0.047 0.669

40:1 698 1.65 531.6 0.99 0.045 0.715 40:1 623 1.67 730.8 0.99 0.037 0.567

Table 6.7 Peak pressure vs billet length regression data for direct extrusion 1000 □ ER=20:1,T=698K ♦ ER=20:1,T=623K

900 ER=40:1,T=698K ER=40:1,T=623K

800

700

600

500

T 400 60 1

Billet Length (mm)

Figure 6.11 Peak pressure vs billet length for direct extrusion 91

An alternative method for evaluating p. can be derived from the differences between the direct and indirect extrusion pressure for identical conditions. Assuming that the decrease in pressure associated with indirect extrusion can be related to the removal of the billet container friction i.e.

Pf=Pe -P i ...... 6.5

where;

Pj = indirect extrusion pressure

Pe = direct extrusion pressure.

The pressure required to overcome friction, Pf, can be defined by the equation;

P f= P t[exp(4nL/D)-l]...... 6.6 where;

Pf = Total pressure

L = Length of the billet/container shear zone

D = Billet diameter

A constant value of p. implies that Pf increases with decreasing temperature as Pf the total pressure increases. Hence substituting equation 6.5 into 6.6 gives;

|i = Ln[(2 Pe-P j)/ Pe] • D /4L ...... 6.7

The results of this analysis are shown in table 6.8 and indicate that, for both reduction ratios and at all temperatures, a fairly constant value of p. is obtained, similar to the values shown in table 6.7.

The trend in the values indicates that the difference in pressure between direct and indirect extrusion increases with decreasing temperature, due to the increase in pressure required to overcome friction during direct extrusion. This is shown more clearly in figures 6.3 & 6.4, which show the variation of peak pressure with initial 92

billet temperature for direct and indirect extrusion at 20:1. Paterson16 and Vierod17 reported a similar trend and this clearly demonstrates the benefit to be gained from indirect extrusion in terms of the reduction in peak load with decreasing initial billet temperature.

Extrusion Ram Extrusion Friction Ratio Speed Temp Coeff

mm/s K it

20:01 3 723 0.0745 20:01 3 673 0.0654 20:01 3 623 0.0576 20:01 3 573 0.0509

20:01 6.9 723 0.0663 20:01 6.9 673 0.0595 20:01 6.9 623 0.0538 20:01 6.9 573 0.0463

30:1 3 723 0.0674 30:1 3 673 0.0601 30:1 3 623 0.0533 30:1 3 573 0.0644

30:1 6.9 723 0.0605 30:1 6.9 673 0.0566 30:1 6.9 623 0.0513 30:1 6.9 573 ’ 0.0464

Table 6.8 Coefficient of friction p. obtained from direct and indirect extrusion data

The value of p. calculated implies that sliding friction conditions prevails in the container, yet experimental observation indicate that the condition inside the container is one of complete sticking friction. For sticking friction conditions, p. should exceed 0.5; the values obtained here are less than 0.1. This implies that the assumptions made in calculating p. are invalid. It is assumed that the excess pressure required in direct extrusion over indirect extrusion is only due to friction. It is 93

known that the flow, and hence the associated redundant deformation in both direct and indirect extrusion are different. The calculation of the Coulomb friction, p., according to equation 3.4, involves a shear stress and the normal pressure, P. The shear stress in this equation is assumed to be purely deviatoric and ignores the hydrostatic component of the shear stress. Hence, these assumptions could account for the anomalously low value of p..

An alternative method of defining the friction conditions is to evaluate the friction factor, "m". The values of m are shown in table 6.7 and have been calculated using equation 3.8 and the strain rate predicted from theory. The average value for m (m=0.8) obtained from these calculations shows that conditions close to full sticking friction (m=1.0) are experienced.

The fact that m is not exactly 1 indicates that the friction conditions in the container may not be 100% sticking and that there may be localised areas where the interface may be slipping and resulting in a stick-slip friction situation. Additionally, in calculating m using equation 3.8, the flow stress used is probably not representative of the true flow stress at the interface. To calculate the flow stress at the interface, o*|, it is first necessary to calculate the temperature at the billet/container interface. This is done by employing a simple heat transfer equation;

O b - T|)/ u r t S) = ( « s/ <*A|)0'5...... 6.22 where;

T b = Billet temperature

Tg = Container temperature

T| = Billet/container temperature

Q£s= thermal capacity of steel=143 MJ/ Kg3 qsa|= thermal capacity of aluminium=478 MJ/Kg3 94

The billet/container interface temperature calculated, however, ignores the temperature rise. Hence, to obtain the true interface temperature, equation 6.23 is employed.

(P2- P^nD ^La- L-|)/4 = p(L2-L,) x Cp AT...... 6.23

where;

P 2 and Pi are the steady state pressures for billet lengths L2 and Li

D = billet diameter p = density of aluminium

Cp = specific heat capacity of aluminium x = thickness of interface

AT = temperature rise

Having obtained the true interface temperature, Z is calculated. In the calculation of Z, the strain rate value utilised was calculated assuming the shear of a continuous 100 pm, thick layer of aluminium.

The new flow stresses predicted at the interface, with the resulting values of m, are shown in table 6.15. 1- > ER

m m/s K M N /m 2 M N /m 2

20:1 6.9 698 143 0.911 132 0.987 0.058 0.18

40:1 6.9 698 160 0.715 138 0.865 0.045 0.38

Table 6.15 New values of m and u. 95

In all instances, the flow stress at the interface, c r (, was lower and correspondingly, m was higher. It is worth noting that for m to be exactly 1, the continuous layer is calculated to be approximately 70 pm for an extrusion ratio of 20:1 and 50 pm for an extrusion ratio of 40:1. This indicates that the layer thickness decreases with increasing extrusion ratio.

Sticking friction conditions are operative when p, the coefficient of friction, exceeds 0.5. The average value of p obtained in the work was only approaching 0.05.

The low value of p obtained can be partly explained by the assumptions made in deriving equation 3.6. In the process of derivation of the equation, it is assumed that the flow stress at the interface is the same as in the bulk material. This assumption, as was shown earlier, is invalid. Hence < r x is replaced by o*| and the new equation for p derived in Appendix 6 is; (p2- pi)p

^ " ( L 2 - L l)4 °,

The values of p obtained using this approach are also shown in table 6.15. It is interesting to note that although these values do not exceed 0.5, they are approximately 6 times larger than previously calculated values of p. Hence further refinement in the equation is necessary.

In conclusion, although the value of m obtained in the revised method was satisfactory, further work must be carried out to obtain an equation which reflects the differences in flow stress at the interface and in the bulk material. Likewise p, the coefficient of friction, needs extensive investigation as it is clear that the assumptions made in obtaining the equation were invalid.

6.7 General pressure equation

The total extrusion pressure has been shown to be dependent on the strain rate, extrusion ratio and extrusion temperature in the previous sections. The general 96

pressure equation relating these parameters has been predicted by mathematical models based on slip line field theory and upper bound solutions to be in the form;

P/G = A + B.LnR...... 6.8

where; P = the extrusion pressure,

a = mean equivalent flow stress,

R = extrusion ratio A,B = extrusion constants

Assuming the value of G equates to the steady state flow stress given by the general hot working equation;

a = 1/a Ln[(Z/A )1/n + ((Z/A)27" + 1 )0-5] ...... 6.9

then, since the two terms within the brackets are approximately equal, this equation reduces to;

P = [0.7/a + 1/an.Ln(Z/A)] ...... 6.10

substitution of equation 6.9 into equation 6.10 gives a new equation of the form;

P = [0.7/a + 1/an.Ln(Z/A)]. [A + B.LnR] ...... 6.11

The similarity between this equation and the empirical relationships established earlier indicate that the total pressure may be more conveniently described by an equation of the form;

P = 1/an [a + b.LnR + c.Ln(Zp/A)] ...... 6.12

where a, b and c are extrusion constants. 97

To assess the validity of this equation, a multiple regression analysis on all the extrusion data was carried out. The results obtained are presented in table 6.9 along with the correlation coefficient.

The high correlation for both extrusion modes (>0.96) indicates that the derived

equation is applicable over the range of extrusion conditions considered. The equation has incorporated dependencies upon temperature, strain rate and extrusion ratio and hence defines the extrusion pressures under a given set of conditions. This is clearly demonstrated in figure 6.12, which shows the experimental pressures against predicted pressures.

P = 1/an [a + b.LnR + c.Ln(Zp/A)]

MODE a b CCC

DIRECT 11.4 2.15 4.0 0.97

INDIRECT 5.3 1.79 3.8 0.96

Table 6.9 Extrusion constants in the general pressure equation

The constants a, b and c reveal some very interesting features. Clode66 associated b.LnR to the useful work done, as Ln R is a measure of the true strain suffered by the material, 'a' was the measure of redundant work, which can been seen to be higher for the direct mode of extrusion. The final term c.Ln(Z/A), introduces the hot working constants into the equation, effectively allowing the material properties for flow stress characteristics to be deduced.

The principal limitation of the equation is that it applies to billets of specific dimensions and is therefore of little use as a general pressure equation. The following discussion will therefore consider the inclusion of a billet dimension term and the relationship of the individual constants to the extrusion process. 98

Predicted extrusion pressure (MPa)

a a.

Q. X HI

Predicted extrusion pressure (MPa)

Figure 6.12 Peak pressure vs predicted pressure calculated using equation 6.12 for both modes of extrusion 99

It was shown earlier that equation 6.8 could be written as;

P /a = a' + b'.LnR + c'L ...... 6.13

where; L is the billet length remaining in the container and a, b and c are constants.

Hence equation 6.13 may be written in the form;

P = 1/an [a + b.LnR + c.Ln(Zp/A) + d.(L/D e)] ...... 6.14

where L represents the length of the shear zone in contact with the container wall and D q the billet diameter.

The validity of this form of equation to demonstrate the extrusion pressure was assessed using a multiple regression analysis on the variable billet length extrusion data. The constants determined are listed in table 6.10 together with the correlation coefficient. The correlation coefficient is still high, but it does show a slight decrease due to the inclusion of the billet term.

a b c d cc

0 .3 2 .1 5 4.0 8 .4 5 0 .9 4

Table 6.10 Extrusion constants In general pressure equation 6.14

The question arises as to whether the linear increase in pressure with billet length shown in fig. 6.11 and equation 6.14 can be simply allowed for by introducing the separate d 17 D b term. It was shown earlier that a friction pressure, Pf, may be related to the total extrusion pressure, P\, by: 100

Pf = Pt [e x p (4 p L /D )-l] ...... 6.6

which can also be written as;

Pf = [A + B.Ln(Z/A)] . [exp(4pl_/D) - 1 ] ...... 6.15

which implies that the friction pressure is a function of the process conditions and may therefore be incorporated into the equation by;

P = 1/an [a + b.LnR + c.Ln(Z/A) + d.Ln(Z/A). exp(4^L/D) -1 ] ...... 6.16

The validity of this equation was assessed by carrying out a multiple regression analysis on the extrusion data, using a constant value of \l= 0.058. The statistical results obtained are shown in table 6.11 and they indicate a decrease in coefficient correlation when compared to the correlations achieved earlier.

a b c d cc

5 .0 6 2 .1 5 4.0 7.25 0.81

Table 6.11 Extrusion constants in general pressure equation 6.16

In conclusion, it must be noted that the inclusion of the billet length term into the general pressure equation should be taken as a guide, rather than an absolute measure of the pressure increase, due to the simplifying assumptions made. This is reflected in the decrease in correlations in these equations and may be related to the assumptions made in the present analysis. Firstly, it is assumed that the equations will apply to all other ratios and billet diameters. Secondly, it is assumed that the work required to overcome friction during direct extrusion is a simple function of the contact area and steady state flow stress and not the shear stress within the shear 101

zone. It may, however, be possible to obtain a more accurate equation by considering the effect of billet length and billet diameter at different extrusion ratios and Z conditions. A study of this is recommended in the "further work" section.

6.8 The effect of section shape on the pressure.

It is known from practical experience that an additional increment of pressure is required to extrude a section of complex design, over that of an equivalent circular cross section. Additionally, the extrusion speed attained is lowered while production generally is more difficult. This section deals with these phenomena and investigates the influence of cross-seclion on extrusion pressure. Round, square, T-shape and U-shape forms, as shown in figure 4.3, were investigated.

6.8.1 Effect of section shape on peak pressure

The effect of die geometry on the shape of the load-displacement curve is shown in fig. 6.13 for the direct extrusion mode; analogous curves for the indirectly extruded sections have not been represented, as the trend observed is similar. The curves have been reproduced for a nominal ram speed of 3mm/s, an extrusion ratio of 40:1 and a temperature of 400°C. It is interesting to note that the peak pressure increment as well as the steady state extrusion pressure are increased with increasing complexity in extruded section. This is accounted for by the increase in complexity of the deformation zone. Thus the average stress needed to continue extrusion will vary, depending upon the extrusion geometry; as the peak load is virtually constant under otherwise almost identical initial conditions, there will be a consequent variation in peak pressure.

The values of peak pressure obtained for indirect extrusion show a similar trend to those obtained in direct extrusion; they are however lower. This latter phenomenon has been fully explained in section 6.3

When the peak pressures are plotted against Ln Z as shown in figures 6.14 & i..3 fet f i goer o la dslcmn tae f iet extrusion direct of trace displacement load on geometry die of Effect Fig.6.13 (M P a) ipaeet (mm) Displacement O to i. .4 rp o Pa pesr v L (i fr hpd etos Drc Mode) (Direct sections shaped for (Zi) Ln vs pressure Peak of Graph 6.14 Fig. Peak Pressure (MPa) 1100

1000

900

800

700

600

500 26 28 30 32 3 Ln (Zi) >h of Peak pressure vs Ln (Zi) for shaped sections (Indirect I 105

6.15, a linear relationship of the form:

P = A+B.LnZp...... 6.3

is apparent. Although not all the experimental points were added due to overcrowding, it can be concluded that a linear relationship of this form does exist, consistent with work done by Wood6 1. The figures indicate that over the range of LnZ considered, the functions relating P and LnZ for each section are more less co-linear.

The relationships established as a result of linear regression analyses are given in table 6.12. In all cases excellent correlations were obtained.

Shape Mode AB CC

Round Direct -873 6 2 .4 0 .9 8 Square Direct -909 6 4 .7 0 .9 8 T-shape Direct -926 6 9 .4 0 .9 9 U-shape Direct -990 6 5 .9 0 .9 9

Round Indirect -1260 68.1 0 .9 8 Square Indirect -1288 72.1 0 .9 9 T-shape Indirect -1208 6 7 .6 0 .9 9 U-shape Indirect -1135 6 5 .5 0 .9 9

Table 6.12 Extrusion constants in equation 6.3 for shaped sections

The extrapolated values of the extrusion constant A show an increase with section complexity. This suggests that the amount of redundant work increases in shaped extrusions. However, Sheppard and Wood68 pointed out that the value of Z which is calculated from the averaged process parameters must vary with section complexity, due to local distortions which occur at the die entry, producing large variations in Z. It would therefore be useful to correct the value of Z to allow for the required pressure increment. This will be dealt with in section 6.9 106

6.8.2 Effect of section shape on APeak pressure

A similar comparison between AP and Ln Z was also made. Figures 6.16 & 6.17 reveal the plots for both modes of extrusion. The scatter in these plots is greater and this is reflected by the lower correlation in table 6.13. This is not surprising, however, when the accuracy with which the peak height can be determined is considered. However, the plots made do show a linear tendency and are approximately parallel.

AP = A + B.LnZ

Shape Mode A B cc.

Round Indirect -736 2 8 .1 5 0 .9 7 Square Indirect -803 3 1 .2 0 .9 5 U-shape Indirect -391 1 8 .3 0 .9 8 T-shape Indirect -534 2 2 .4 0 .9 9

Round Direct -947 3 7 .4 2 0 .9 9 Square Direct -941 3 7 .5 8 0 .9 9 U-shape Direct -936 3 8.5 1 0 .9 8 T-shape Direct -901 3 6 .5 5 0 .9 9

Iahla.6i13-Extruslon constants for.shaped sections

This is surprising because extrusion does not commence until the load has neared peak and full ram speed is certainly not attained until just after peak load. In addition, maximum pressure incorporates a friction component, which has always been considered to be independent of both the strain rate and deformation zone shape. The implication is that peak load is independent of all variables except the initial billet temperature and the extrusion ratio (or modified extrusion ratio for shaped sections). These observations would therefore indicate that the pressure increment,

AP, is linked with the development of the deformation zone. 300

200

100

0

of APeak pressure vs Ln (Zi) for shaped sections (Direct Me i..7 rp o Aek rsue s n Z) o sae scin (niet Mode) (Indirect sections shaped for (Zi) Ln vs pressure APeak of Graph Fig.6.17

APeak Pressure (MPa) oo o 109

6.9 Geometrical considerations

One of the major problems encountered when attempting to correlate results from different workers is in defining the complexity of the shape. One method utilised to compensate for the effect of differing shape has been the concept of peripheral ratio (r). This parameter is defined as the ratio of the extrudate perimeter, pe , to the circumference of a rod, pr, having the same reduction ratio.

The peripheral ratio and section geometry of the dies used in the present study are given in fig 4.3. Wood used the peripheral ratio to obtain a modified extrusion ratio defined by:

R’ = QR...... 6.17

where;

0 = (pe/p r)0-5 = r 0-5

R = extrusion ratio

Wood then used this to obtain a correction factor to modify the temperature compensated strain rate Z through an equation of the form:

LnR' = a + b.LnZ...... 6.18

where; a,b = constants Ln R' = shape correction factor.

In the present work it was found that the best correlation of the data was obtained when the peripheral ratio was applied directly to the equation, resulting in the modified form given by:

P = A + B.Ln(rZ)...... 6.19

These curves clearly show that the extrusion pressure necessary to form 110

complex sections increases with increasing peripheral ratio, this mainly linear pressure increase being considerably less for indirect extrusion. In this equation the extrusion constants A and B are derived from rod extrusion results, i.e. r=1.

In his work, Wood51 used a fixed container temperature of 300°C, whereas in the current work the container temperature was varied such that it was 50°C below the initial billet temperature. This difference in temperature control would cause a variation in the peak pressures, due to a change in heat flow during extrusion. This difference is highlighted by figures 6.18 and 6.19, which show extrusion pressures predicted by equations 6.18 and 6.19 respectively. The former was compiled using Wood's correction (r0-5) and shows considerable scatter, whereas the latter plot used r=1 and gives an excellent prediction of the required extrusion pressure for shaped products. The difference in the magnitude of the two correction factors is shown in table 6.14 below.

Shape r p0.5 correction Wood's correction factor factor

Round 1 1 Square 1 .1 3 1 .0 6 T-shape 1.31 1 .1 4 U-shape 1 .6 4 1 .2 8

Table 6.14 Correction factors for non circular sections

The greater pressures required for the extrusion of sections, compared with those necessary for the extrusion of a symmetrical round bar, are principally due to the extra forces required to overcome the additional shear resistance arising from the asymmetrical flow. From the increase in deformation resistance, it follows that the force required to overcome friction between billet and container also increases. Ill

Graph 6.18 Graph of peak pressure vs Ln(Zi) (direct mode) rp 61 Gah f ek rsue s nZ) idrc mode) (indirect Ln(Zi) vs pressure peak of Graph 6.19 Graph Peak Pressure (MPa) Peak Pressure (MPa) Peak Pressure (MPa) 113

Inclusion of peripheral ratio (T ) in the general pressure equation extends its usefulness, since it takes into account section geometry. This results in the following modified equation;

P = 1/an(A' +B’lnR + C’Ln (rZ/A)...... 6.20

where the constants are found for rod extrusion (r= 1 ), as described earlier. Figure

6.20 illustrates the accuracy of this approach for both extrusion modes, high correlations being obtained in both cases.

It is evident that the use of a peripheral ratio to modify 'Z' in shaped extrusion is simple and effective. However, it requires further extensive research to confirm these findings, by including a wider range of peripheral ratios and varying design while maintaining constant peripheral ratio. It is also envisaged that this treatment might be extended to include multi-hole dies. The incremental internal shear arising from extrusion of sections is quantitatively very difficult to determine. A study of metal flow will later be used in an attempt to explain the increase in pressure requirements for different shapes.

6.10 Surface quality of extrudates

During the extrusion of aluminium alloys, especially those for high strength in aerospace applications, the attainment of good mechanical properties is one of prime importance. However, these extrudates are often rendered useless by the onset of deleterious surface defects; these can provide sites for crack initiation, leading to reduced mechanical properties and lower resistance to degradation by the environment. Hence a good product surface finish is essential if any extrusion process is to be commercially successful.

In the present work, the necessity to machine the billets prior to extrusion resulted in a smooth billet surface, which eliminated many of the sources of poor surface finish. However, defects common to all extrudates were die lines and pick up. i..0 rdce pesr fo mdfe gnrl rsue equation pressure general modified from pressure Predicted Fig.6.20

Predicted peak pressure (MPa) Predicted peak pressure (MPa) s xeietl xrso pesr frsae sections. shaped for pressure extrusion experimental vs pressure (MPa) pressure pressure (MPa) pressure 114

115

The appearance of these defects showed no significant dependence on either strain rate or temperature, but was found to increase dramatically if the dies used were not cleaned before use. Therefore, particular care was taken to ensure that all dies were clean, preheated to the container temperature and of good dimensional accuracy. It is believed that the trends to be described are for a generally consistent die condition.

The surface finish was categorised into one of three grades:

A - no evidence of cracking

B - cracking commences at some some distance along the extrudate C - cracking occurred along the entire length of the extrudate.

Typical examples of the surfaces found midway through extrusion of 2024 alloy are shown in fig 6.21 for both direct and indirect process. The extrudates were extruded through a 20:1 die, at a ram speed of 6.9 mm/s, over the range of extrusion temperatures. It can be seen that for direct extrusion, that the severity of the surface cracking increased with extrusion temperature. A similar trend is also apparent for the indirect extrudates, except that cracking is not as severe as that for direct extrusion. This behaviour may be attributed to lower temperature rises associated with the indirect process and suggests that extrusion without the onset of surface cracking is possible over a wider range of extrusion conditions with the indirect process.

Figure 6.21 also shows examples of the types of surface found along shaped sections extruded through a 40:1 die, at a ram speed of 3.0 mm/s and an extrusion temperature of 350°C. Again a similar observation of the differences between direct and indirect extrusion was made. Figure 6.21 d shows examples of the type of defects which are obtained; the nature and origin of these defects will now be discussed.

During extrusion the surface layer of the extrudate heats up more than the central core. In many cases, the eutectic temperature is exceeded and regions of eutectic composition begin to melt i.e. incipient melting. Even if no actual melting occurs, these regions become so weak that their flow stress is approximately zero. Al-Cu-Mg and Al-Cu-Zn-Mg alloys are especially susceptible to hot shortness as they contain low melting point eutectics. It was observed that incipient melting was Fig. 6.21 Surface quality of extrudates a) Direct extrusion, ER=20:1, v=6.9mm/s, decreasing from top to bottom b) Indirect extrusion, ER=20:1, v=6.9mm/s, Tg decreasing from top to bottom c) Shaped extrusion, ER=40:1, v=3.0mm/s d) Poor surface quality in shaped extrusions

117

always preceded to some extent by 'speed cracking'. This type of surface cracking is caused by a reduction in flow stress, due to the higher temperatures at this location. The occurrence of speed cracking becomes more likely as the strain rate (or exit velocity) is increased; hence the term speed cracking. This can be explained by considering the temperature rise at the extrudate surface, as a result of friction between the metal and the die lands. This considerable heat source has a decisive influence on the extrusion speed and of course the maximum initial billet temperature. The hot shortness phenomenon is enhanced by the sudden loss of ductility, as the billet material moves into the die throat, due to the release of the hydrostatic component of stress. This explains the commonly observed inclination of the surface cracks at approximately 45° towards the rear of the extrudate.

The danger of hot shortness increases when extruding complex sections, especially when these include sharp corners. At such locations the susceptibility to crack formation is enhanced by two factors. The first is that the matrix at a corner point will receive almost double the frictional force, exerted by the perpendicular die lands. The second reason is that the volume of material receives a heat input from two perpendicular die lands. Therefore the combination of increased temperature (and therefore reduced flow stress) and increased frictional forces tends to tear the matrix, leading to crack initiation at this location. This is consistent with the assumption made by Lange83, that the temperature rise at the edge of a section was approximately twice that at the mid point of the sides. Closer examination of the shaped sections revealed that the cracking was primarily confined to the corner locations of the section, and consisted of a network of incipiently melted second phase particles. This confirms the hypothesis that the susceptibility to hot shortness is greater at the corner section of an extrudate.

It has been recognised84'85*88 that the process parameters causing the onset of hot shortness can be predicted using an equation of the form :

Ln Z< a T b ...... 6.21

where a and b are constants and T is the initial billet temperature in degrees Kelvin. This equation was applied to both direct and indirect extrusion of round sections as 118

shown in figs 6.22. The statistical data is given in below. The high correlations indicate that it is also applicable to this alloy.

Direct LnZ| < 5246.T'0-810 cc 0.991

Indirect LnZj < 20008.T'1 009 cc 0.985

It can be concluded from this analysis that use of the indirect process does indeed extend the range at which extrusion is possible without the onset of surface cracking.

However, although it would appear that the indirect extrusion produces better surfaces than the direct process, it should be remembered that the sources of other defects have been removed by machining, which is the normal procedure for the commercial production of this alloy87. Although the equation is useful in determining the initial extrusion condition resulting in cracking, the exact conditions cannot be properly defined due to the uncertainty of the temperature and

stress conditions at corner locations in complex sections.

Recent work has shown that it may be possible to reduce the onset of surface cracking by effectively increasing the strength of the material at the edge of the extrudate through solid solution strengthening8 8 . This was achieved by solution soaking the extrusion billets at 500°C for 30 minutes before furnace cooling to the extrusion temperature. It was assumed that sufficient solute had been retained in solution to increase the peak pressure by 5% and reduce the surface cracking at an extrusion temperature at 400°C. It is predicted that a similar increase in extrudability will be achieved in the extrusion of shaped sections using this

technique. i 62 Gah f n Z) s eprtr fr ufc quality surface for temperature vs (Zi) Ln of Graph 6.22 Fig Ln Zi Ln Zi considerations 120

6.11 Conclusions

1. The dependence of peak pressure on Ln R is linear whilst a power relationship provides the best fit for the variation of peak pressure with initial billet

temperature.

2. Use of indirect extrusion reduces the extrusion pressure requirement. This effect is particularly significant at low extrusion temperatures and high ram speeds.

3. The excess pressure required to initiate extrusion can be related to the initial

Zenner-Holloman parameter by a linear relationship. AP during indirect extrusion

is smaller. Both these observations are consistent with the theory that the peak is associated with the establishment of the quasi static deformation zone, the formation of which requires an excess of dislocations.

4. The pressure for direct extrusion increases with billet length due to increase in the friction between the billet and container. The friction conditions may be described by either a friction factor, 'm', or a friction coefficient, 'jj.' but neither method is completely satisfactory.

5. The value of the temperature compensated strain rate, Z, must be corrected when considering the die geometry. The load encountered during extrusion is a function of the load required for rod extrusion and the correction factor. The necessity for this correction factor arises from inaccuracies in calculating the strain rate and deformation during shaped extrusion. An additional consideration is the use of mean strain rates and temperatures calculated from rod extrusion in the calculation of the Zenner-Holloman parameter. The correction factors can be related to the geometric variables but further work is required in this area. As the complexity of section shape is increased, the magnitude of the excess load required to initiate extrusion

increases.

6. A general pressure equation incorporating the hot working constants has been established for a constant billet length. Inclusion of a billet dimension has met with only a limited success due to the difficulty in defining the friction conditions. 121

7. The flow patterns for indirect extrusion exhibit no dead metal zone. The deformation field for section shapes is severely distorted by the asymmetry of the section.

8. The surface condition of 2024 extrudates is dependent upon the strain rate and temperature. Expressions have been derived which predict the onset of surface cracking by relating the Z parameter and the initial temperature. The onset of surface cracking is less evident for indirect extrusion and thus higher ram speed is possible in indirect extrusion compared to direct extrusion at the same temperature. 7.0 Structural Analysis 123

7.1 The evolution of structure and substructure

The microstructure is an important factor in the overall analysis of the extrusion process. Hence the evolution of the structure and substructure has been investigated at all stages of processing, from the as-cast state to the fully heat treated extrudates. The development of microstructure is similar in shaped extrusions unless otherwise stated.

7.1.1 As-cast structure

The structures shown in fig. 7.1 a,b are optical micrographs of the as-cast alloy taken from the centre of the ingot, but are typical of the structure throughout. The micrographs show that the solidification mechanism was dendritic, with the major alloying elements segregated to the interdendritic regions; these have been etched up black in fig. 7.1a. The cast material also reveals considerable coring, labelled X on fig. 7.1b. This is a common feature in all cast alloys due to the non-equilibrium conditions prevailing during solidification. The grain size in this alloy was

approximately 90jim.

A scanning electron micrograph of the cast material in the back-scattered mode is shown in fig. 7.1c. A higher magnification of the region marked Y in fig. 7.1c is shown by fig. 7.1 d. Both micrographs reveal the interdendritic regions to contain lighter and darker regions. The lighter regions were found by EDAX analysis to contain predominantly copper and magnesium while the darker ones were found to be rich in magnesium, silicon and iron. The available literature5'16'17'86'92

indicates that these phases are CuAl 2, CuMgAl2 and CuFeMnAlg. This was confirmed in the present case by X-ray diffraction analysis. The darker regions have been identified by previous workers5 '8 6 '92 to consist of two separate eutectics with compositions of a + (FeMnSiCu)AI and a + 0 + C ^M gsSisA ls-

The relatively low ductility of the cast material, caused by the presence of these intergranular and interdendritic second phase constituents, accounts for the inferior hot-workability of as-cast materials. In order to improve this workability it is Fig. 7.1 As-cast material microstructure a,b) optical c,d) scanning electron micrographs

125

necessary to perform a homogenisation treatment prior to extrusion.

7.1.2 Homogenised Structure

Previous workers have shown that a standard homogenisation treatment of 24 hours at 500°C is suitable for most commercial direct chill-cast aluminium-copper-magnesium alloys5*15»17. Differential Scanning Calorimetry tests on the as-cast material revealed a melting peak at about 511°C. Homogenisation was therefore limited to 500°C, with a 10°C safety margin. Heating to temperatures beyond 500°C risks the onset of eutectic melting which can severely reduce the properties of the product, whilst prolonged heating has been reported to produce no real significant change in the overall structure15. The billets were furnace cooled.

An optical micrograph of the material homogenised for 24 hours at 500°C is shown in fig. 7.2a. The heavy coring and dendritic structure characteristic of the as-cast alloy has been completely eliminated. It should also be noted that no substantial grain growth occurred during the homogenisation treatment; the approximate grain size was found to be 95jim.

The structure after homogenisation for 4 hours at 500°C is shown in the transmission mode in fig. 7.2b. It reveals needle-like structures surrounded by a region containing very fine precipitates. It is apparent from the morphology of the needle-like structures that they have a habit plane relationship to the matrix of the aluminium solid solution. After 24 hours homogenisation, however, the matrix is more uniformly interlaced with the needle-like precipitates shown by fig. 7.2c. The structure was also examined in the backscattered mode as shown in fig. 7.2d. Both figs. 7.2c and 7.2d reveal that the segregation has been completely broken down after 24 hours of homogenisation and the soluble components have been distributed uniformly throughout the structure by diffusion. The structure contains a homogeneous distribution of small particles with only the isolated occurrence of large particles. Homogenisation takes much of the copper and magnesium into solution; these reprecipitate on subsequent furnace cooling as CuMgAl 2 with the possible formation of some CuAl 2. The needle-like structure shown in fig 7.2c was Fig. 7.2 Homogenised material microstructure a) optical, 24 hour homogenisation b) TEM, 4 hour homogenisation c) TEM, 24 hour homogenisation d) SEM, 24 hour homogenisation e) SEM, structure just prior to extrusion > 127

indeed identified as CuMgAl 2 and adopts a Widmanstatten structure on slow cooling.

All the iron containing phases undergo transformation to A ly C ^ F e , FeMnAIg,

MnAIg, possibly accompanied by other minor phases. Simultaneously, manganese is precipitated from solid solution as the complex C ^ M n g A ^ o dispersoid. The smaller, lighter particles were analysed to be generally FeMnAIg and MnAlgf while the darker particles were found to be Mg 2Si. The presence of Mg 2Si implies that the increase in copper taken into solution during homogenisation results in silicon combining preferentially with magnesium.

The microstructure just prior to extrusion is shown in fig. 7.2e. It was obtained by induction heating a homogenised billet to the process temperature, followed by rapid quenching, as shown by the heating profile in fig. 7.3 below. The striking difference between this microstructure and the homogenised structure is the amount of precipitate dissolution which has occurred during the short induction heating period.

Fig. 7.3 Heat treatment given to material prior to extrusion 128

7.2 Structural variation with extrusion conditions

In the following section the effect of process conditions and heat treatments on the structural and substructural features will be discussed. The structures obtained for the as-extruded material, as well as the heat-treated extrudates were examined.

7.2.1 Extrudates structures

Fig. 7.4. shows examples of the microstructures observed in selected longitudinal sections taken from extrudates obtained under a variety of extrusion conditions.

The as-quenched microstructures could be classified into two types; the first consisted of a core of fibrous grains surrounded by a more dislocated structure at the periphery, examples of which is shown in figs. 7.4 a.b; fig. 7.4 a shows a region near the periphery of the extrudate, fig. 7.4 b the centre. The original homogenised grains, elongated in the extrusion direction, give rise to the fibrous core of the structure, with the periphery originating from the heavily sheared regions outlining the deformation zone. This structure is typical of those produced in extrusion at temperatures below 350°C. The second type of structure is shown in fig. 7.4 c and consists of an inner core of fibrous grains surrounded by a peripheral zone of recrystallised grains, an appearance which is characteristic of high temperature extrudates. The formation of this annulus of recrystallised grains can be attributed to the fact that the metal at the surface of the outgoing extrudate has undergone a greater degree of shear in entering the deformation zone than the material at any other point. Thus the surface also tends to be hotter than any other part of the product and is therefore more likely to be recrystallised.

The structure observed in the indirect mode, as shown in fig. 7.4 d, was similar except that the depth of the peripheral region of recrystallisation was reduced and the structure was more fibrous; this can be attributed to the smaller temperature rises and less shear involved with this process.

It was also observed that the depth of this recrystallised surface region increased Fig. 7.4 Microstructure of longitudinal sections of extrudates

a) extrudate periphery, direct mode, = 300°C

b) extrudate centre, direct mode, Tg = 300°C c) extrudate periphery, direct mode, = 450°C d) extrudate periphery, indirect mode, = 450°C

130

quite markedly at some of the corners in the non-circular extruded sections. It was found that the depth of these regions was invariably greater in proximity to an internal acute angle in the die orifice. This is shown in fig. 7.5 a for a T-section. This effect can again be attributed to localised variations in the mode of deformation and the temperature. The massive shear occurring when there is an abrupt change in section results in a locally increased temperature rise, which in turn leads to an increase in recrystallisation (be it dynamic or static) and explains the difference in thickness of the recrystallised annulus at sharp section changes in the non-circular extrudates.

In practice, the structures arising as a result of extrusion are not comprised solely of one of these basic elements, but are often a mixture of the types of structures described. The extent of mixing of the basic structures is dependent upon the initial billet temperature and the degree of deformation produced during the

extrusion process.

The effect of the extrusion ratio on the volume percent of material recrystallised per unit length at 4 different extrusion temperatures for both direct and indirect extrusion is shown in figs. 7.6 and 7.7. The equation used to determine this volume fraction is shown below:

Vol % = (l-[(R-d)/R]2) .1 0 0 ...... 7.1

R = Extrudate Radius d = Recrystallised depth

The measurements were taken at a constant distance along the extrudate corresponding to 0.3L. Fig. 7.5 Microstructures observed in the transverse plane of press quenched shaped extrudates a) T-section, as extruded b) T-section, solution soaked at 500°C for 0.5 hours c) U-section, solution soaked at 500°C for 0.5 hours

132

Extrusion temperature (C)

Fig. 7.6 Volume % recrystallised vs initial billet temperature for direct extrudates

Extrusion temperature (C)

Fig. 7.7 Vol. % recrystallised vs temperature for indirect extrudates 133

At the lower extrusion ratio of 20:1, the recrystallised grain size was normally found to be limited to the width of the recrystallised band, while at higher extrusion ratios finer, less elongated grains were found. On increasing the extrusion ratio to 30:1, the higher strains involved increase the driving force for the nucleation of recrystallisation. Increasing the ram speed, on the other hand, reduces the quench delay but increases the temperature rise. A combination of these factors helps to increase the fraction of recrystallised material with increase in extrusion ratio and ram speed, although the recrystallised grain size was found to decrease.

For shaped sections, estimation of the depth of the recrystallised layer was more difficult, since the layer was no longer uniform (as described earlier). The structures used for this measurement consisted of a fairly well-defined recrystallised region surrounding a recovered structure. Due to the constraint on the maximum indirect extrusion ratio described earlier, only one extrusion ratio was used in these determinations and hence only a plot of recrystallised layer thickness against Ln Zj was made, as shown in figs. 7.8 and 7.9. In general, a decrease in extrusion temperature and an increase in ram speed and extrusion ratio were found to increase the volume fraction recrystallised for Zj < 30; additionally the layer was thicker for direct extrudates and more complex sections due to the larger temperature rise as explained earlier. It is also worth noting that at high values of Zj

(Zj > 30) the depth of the recrystallised layer decreases. This is due to a combination of two factors. The first is the method of calculation of Z: at low Z, use of the initial billet temperature is more accurate because the ram speed is low and therefore the temperature rise is smaller; when Z is high, the ram speed is correspondingly high so that the instantaneous temperature rises are greater, causing the value of Zj calculated to be artificially raised. The second is due to the quench delay: at high ram speeds, the quench delay is decreased, inhibiting extensive recrystallisation.

The effect of indirect extrusion on the volume percentage recrystallised is represented graphically in fig. 7.7 and 7.9. There is a large reduction of percentage recrystallisation for both extrusion ratios investigated when the indirect process is used, the same general trend with temperature and ram speed being observed as in direct extrusion. This reduction in the percentage recrystallisation has been i. . Vrain f ersalsto lyr hcns wt L (i fr hpd sections shaped for (Zi) Ln with thickness layer recrystallisation of Variation 7.9 Fig. i. . Vrain f ersalsto lyr hcns wt L (i fr hpd sections shaped for (Zi) Ln with thickness layer recrystallisation of Variation 7.8 Fig.

Rexn Layer Thickness (nm) l _ CC H >* © © CO © X © o c co E a. in the as extruded condition for indirect extrusionindirect for condition extruded as the in in the as extruded condition for direct extrusion direct for condition extruded as the in 134

135

observed by previous workers71*93'94'99 and can be attributed to the more uniform nature of flow and lower temperature rises experienced during indirect extrusion.

Table 7.1 shows the volume percentage recrystallised related to Ln Zj. Although reasonable correlations were obtained, such relationships can only be regarded as an approximation, since the quench delay varies with the strain rate and, in addition, is virtually doubled during indirect extrusion due to the difference in tooling arrangement. For this reason, the percentage recrystallisation in equivalent indirect extrusions may be even smaller than shown by the results in table 7.1. Since a duplex structure is reported to be deleterious to fracture toughness and corrosion resistance, the reduction in peripheral recrystallisation is an important feature of the indirect process.

Volume % = a + b.LnZj

Extrusion Extrusion a b cc Mode Ratio

Direct 20:01 17.62 -0.54 0.82

Direct 30:01 50.83 -1.53 0.91

indirect 20:01 8.03 -0.25 0.96

Indirect 30:01 16.74 -0.5 0.96

Table 7.1 The variation in volume % recrystallised during direct and indirect extrusion of alloy 2024 in the T1 temper

A study of the structure along the length of the 20:1 extrudate revealed an interesting feature as shown in table 7.2. In low temperature extrusions, little or no variation was observed in the fibrous nature of the grains along the entire length of the extrudate. However, in extrusions above 400°C, the depth of the peripheral region of recrystallisation increased by about 9% along the length. This may be 136

related to the increase in the temperature associated with deformation at higher temperatures, producing a greater proportion of recrystallised grains. This increase was also observed by other workers and has been attributed to a decrease in effectiveness of the recrystallisation inhibitors at higher temperatures16 »17«71. However, this is debatable as recrystallisation inhibitors are generally insoluble even at high temperatures.

Distance Down T =:300C T = 450C Extrudate

Direct Indirecl Direct Indirect

0.3 0.8 0.2 4.5 1.8

0.6 1.2 0.7 8.2 2.5

0.9 2.3 1.1 13.5 3.7

Table 7.2 The variation in volume % recrystallised along direct and indirect extrudates of alloy 2024 in the T1 temper

7.2.2 Substructural variations with extrusion conditions

Typical substructures observed in the transverse plane of the extrudate for direct and indirect extrusion are shown in fig. 7.10. All the microstructures exhibited a subgrain structure characteristic of dynamic recovery and the lack of contrast between the subgrains is an indication of the low misorientation that exists. It has been reported16,17,50,84 that f0r alloy 2014 this misorientation varies between 0.17° and 1.13°. Evidence of dynamic recrystallisation was not observed in any of the specimens examined; this may be attributable to the presence of fine intermediate precipitates (<1p.m) which inhibit recrystallisation100. It is interesting to note that with increasing extrusion temperature there was a decrease in the density of particles, indicating a greater amount of alloying element in solution during extrusion at higher temperature. This may be beneficial to the Fig. 7.10 Substructures observed in the transverse plane of press quenched direct and indirect rod extrudates of alloy 2024 (ER =20:1) a) Direct, Ln Zj = 25.59 b) Direct, Ln Zj = 27.41 c) Direct, Ln Zj = 31.99 d) Indirect, Ln Zj = 25.87 e) Indirect, Ln Zj = 27.68 f) Indirect, Ln Zj = 32.26

138

age-hardening properties of the extrudate in the as-extruded condition. At low extrusion temperatures the subgrains are fairly well defined and show a relatively high internal dislocation density. When correct contrast conditions were chosen, bands of dislocations were visible in the structure, although these are not evident in the micrographs presented. The subgrains are equiaxed in the transverse plane of the extrudate but are elongated in the extrusion direction (i.e. the longitudinal plane). The particles are also aligned in the extrusion direction and are arranged uniformly throughout the structure. Pinning of the sub-boundaries is evident, particularly by larger particles, but the subgrain size was not found to be dependent upon the interparticle spacing. Since it is apparent from their alignment that the majority of particles are present as precipitates during extrusion, the interaction between some of the particles and the subboundaries must have been overcome by the applied stress. This evidence indicates that some recovery has occurred at this temperature, but the process is incomplete.

At high extrusion temperatures the subgrains increase in size and improve their perfection, showing reduced boundary widths and internal dislocation density. This is a result of the increase in thermal activation, permitting easier cross slip and climb of dislocations. The subgrains are also more equiaxed in the longitudinal plane, which indicates that the increase in thermal activation and reduction in precipitate density produces an increase in the mobility of the subboundaries, allowing polygonisation to become more complete.

The substructures obtained from indirect extrusion show the same characteristics as those observed for direct extrusion, the main difference between the two structures being the indirect subgrain size, which is larger. This may be attributed to the arrangement of the indirect tooling, which is such that the delay time between die exit and quenching is greater. Clearly, the growth of the subgrains during this time is more significant at high temperatures than at low temperatures.

An important feature of the substructure is the presence of helical dislocations, which are not a direct result of the deformation, but are formed during quenching; the proportion of these dislocations tends to increase with increasing extrusion temperature. Examples of these can be seen clearly in fig. 7.10a. They are caused by 139

the interaction of quenched vacancies with screw or mixed dislocations during the quenching process. The vacancies condense onto the screw dislocations causing the dislocation line to become curved, forming one turn of a spiral. Growth occurs as other vacancies are attracted to different points along the dislocation line. The helical dislocations observed in these specimens tended to be locked by precipitates, which prevented them from straightening out, and are generally orientated in [200] directions. Their presence is beneficial during the ageing process as they act as nucleation sites for precipitation.90*91

Figs. 7.11 and 7.12 show the substructures for round and T sections at 40:1. The shaped extrusion showed similar features to rod extrusions such as decreasing subgrain size, helical dislocation density and precipitate density, with decreasing temperature but showed one anomaly. It is interesting to note that the 40:1 round extrudate shows a coarser substructure than the 20:1 round bar extrudate at a corresponding Z\. Though in general it could be expected that the subgrain size will decrease with increasing extrusion ratio, this effect may in fact be limited to the lower extrusion ratios (extrusion ratio up to 20:1) with the substructural refinement remaining constant beyond this point, as represented schematically in fig. 7.13. However, as extrusion ratio increases still further an additional influence begins to predominate; this additional effect is due to the temperature rise which increases in magnitude and increases the temperature in the deformation zone. This increased temperature and the strain effect will result in a slightly coarser structure than expected at higher extrusion ratios.

The subgrains were also found to increase in size with complexity of shape as shown by figs. 7.11 and 7.12. Although in this particular instance the difference was not substantial it could again be attributed to the increase in heat generation with increase in complexity of shape for constant extrusion ratio. Similar effects were observed in indirect extrusion.

The dependence of the subgrain size on the process variables for extrusion and torsion testing will be discussed in the following section. Fig. 7.11 Substructures observed in the transverse plane of press quenched direct and indirect rod extrudates of alloy 2024 (ER =40:1) a) Direct, Ln Zj = 25.81 b) Direct, Ln Zj = 28.95 c) Direct, Ln Zj = 32.20 d) Indirect, Ln Zj = 25.81 e) Indirect, Ln Zj = 28.49 f) Indirect, Ln Zj = 30.59

Fig. 7.12 Substructures observed in the transverse plane of press quenched direct and indirect T-section extrudates of alloy 2024 (ER =40:1) a) Direct, Ln Zj = 27.01 b) Direct, Ln Zj = 29.05 c) Direct, Ln Zj = 30.44 d) Indirect, Ln Zj = 26.83 e) Indirect, Ln Zj = 29.05 f) Indirect, Ln Zj = 30.57 e 142

Strain

Fig. 7.13 Influence of strain on subgrain size

7.2.3 The effect of process conditions on the steady state substructure

In section 7.2.2 it was shown that the subgrain size was dependent upon the

processing conditions, it would therefore be reasonable to assume that a relationship between the sub-grain diameter and the temperature compensated strain rate (Z) should exist. The expression most commonly used for the relationship is of the form:

d'm = a + b.LnZ...... 7.2

where a, b, and m are constants. Although the value of m=1 is generally considered to give the best overall correlations, recent work89 has shown that equally good correlations can be achieved with 0.3 < m < 1.25. However, in the present work a value of m=1 was used and was found to result in the best correlations.

Subgrain sizes, produced by varying process conditions during both extrusion and hot torsion tests, were measured using the mean linear intercept method, to test the validity of this relationship. The results from the torsion tests represent a wide range of temperatures and strain rates, whilst those from extrusion are for ratios of 20, 30 and 50:1. The value of the temperature compensated strain rate was corrected for the temperature rise encountered during extrusion using the integral Fig. 7.14 Inverse subgrain size vs Ln Zc for extrusion and torsion

Fig. 7.15 Inverse subgrain size vs Ln Zi for extrusion 144

profile method described by Cooper76.

The results obtained are plotted as reciprocal subgrain size against Ln Zc in figs.

7.14 for both extrusion and torsion tests (where Ln Zc is the corrected temperature compensated strain rate). Although there is some scatter in the data, a linear relationship was found to exist for all cases; the statistical correlation data for these plots are given in table 7.3

d-1 = a.LnZc - b

Data a b corr

Torsion 0.1033 2.0781 0.943

Direct Extrusion 0.0913 1.7675 0.944

Indirect Extrusion 0.0943 1.8474 0.931

Combined Extrusion 0.0923 1.7936 0.946

Table 7.3 Temperature compensated subgrain size relationships for extrusion and torsion

The relationship of d~1 to Ln Zj was also examined, where Zj is the value of Z referring to the initial extrusion conditions; the schematic and statistical data are given in fig. 7.15 and table 7.4 respectively. This demonstrates the justification of using Zc rather than Zj, as the former shows an improvement in correlation.

It can be seen that, in general, the subgrain size increases with decreasing Z, which is associated with a decreasing strain rate and increasing temperature. The structural implications of this may be explained by remembering that at lower strains rates there is an increase in the time available for dislocation rearrangement and at higher temperature, increased thermal activation appreciably increases dislocation mobility, thus resulting in larger subgrain sizes. 145

d’1 = a.LnZj - b

Data a b corr

Direct Extrusion 0.0378 0.5778 0.941

Indirect Extrusion 0.0426 0.6457 0.924

Table 7.4 Subgrain size relationships for extrusion without temperature compensation

The similarity between the torsion and extrusion substructures over a similar range of Z conditions confirms the use of the hot torsion test to model the flow stress characteristics of a material during extrusion and demonstrates the possibility of using as a reference sets of torsion substructures obtained for a particular alloy to predict the final substructures in extrusion, with the aid of the hot working constants established from torque-twist data.

7.2.4 Variation of substructure across and along extrudates

The variation of the substructure across the extrudate has been investigated by measuring the subgrain sizes of specimens taken from the centre, mid-radius and edge of the extrudate. The results are given in table 7.5 for direct and indirect extrusion at 400°C and for an extrusion ratio of 20:1. The variation across the extrudate can be seen to be relatively small, although there is a tendency for the subgrain size to increase from the centre to the edge. This is consistent with the temperature distribution expected in the extrudate due to the inhomogeneous nature of the deformation. The structure at the edge of the extrudate which undergoes greater deformation during extrusion experiences a greater increase in temperature and thus produces larger subgrains. However, it should be noted that the temperature distribution in the extrudate is not particularly significant at this temperature. 146

Distance down Extrusion Centre of Mid-radius of Edge of extrudate mode extrudate extrudate extrudate

0 .3 Direct 2.13 2.17 2.28

0.6 Direct 2.18 2.21 2.31

0.9 Direct 2.22 2.24 2.45

0.3 Indirect 1.68 1.74 1.87

0.6 Indirect 1.77 1.85 2.03

0.9 Indirect 1.91 1.92 2.12

Table 7.5 Variation of substructure size in pjn across and along the extrudate

The variation of the substructure along the extrudate has been investigated by measuring the subgrain sizes in specimens taken at specified intervals along the extrudate. The results are also given in table 7.5 for both direct and indirect extrusion at an initial temperature of 400°C and an extrusion ratio of 20:1. In direct extrusion a progressive increase in subgrain size along the extrudate was seen which may be related through Z to the rise in temperature which occurs as extrusion proceeds. The indirect extrudate also shows an increase in subgrain size along its' length, but the variation is not as significant and is consistent with the observation that lower temperature rises are experienced for indirect extrusion.

7.3 Material flow characteristics during steady state extrusion

The development of material flow during direct and indirect extrusion has been the focus of much interest in aluminium alloys in general, but has been limited to axisymmetric sections and most often rod extrusions. The influence of section geometry on the quasistatic deformation zone is difficult to establish as, unlike in axisymmetric rod extrusion where the circumferential strain is essentially zero, 147

the introduction of the third dimensional strain requires more careful interpretation. This is especially so near the die throat where the extrudate surface and structural features are being generated. One of the techniques most commonly used to study metal flow is the gridded billet technique, but as this relies on planar flow, it was unsuitable for use here.

In the present work, the variation in flow between different shapes has been established using super purity aluminium, extruded at 300°C and 425°C at an extrusion ratio of 40:1. Due to its large grain size and ease of macroetching, pure aluminium is particularly suited to deformation behaviour studies of this type. The differences in the flow patterns attained during steady state extrusion are shown in figs. 7.16-7.18. Details of the experimental procedure are given in section 4.6.

Fig. 7.16a and 7.17a show the steady state flow pattern developed during direct extrusion of a circular section at 300°C and 425°C. It is evident that the temperature has very little effect on the development of the deformation zone. The primary deformation is characterised by an elliptical region of intense shear

extending from the rear of the billet container interface to the die throat. For reasons not fully understood, there is a secondary zone of intense shear along the billet axis indicating that there is a back flow of material along the billet axis in the container. This may however be possible, if for some reason the pressure pad has a recess at the centre and during the initial upsetting stage a small back flow developed at the billet rear; this might then be preserved throughout the extrusion stage. This

necessitates that the material adjacent to the billet axis flows towards the die throat and results in the strange shape of the deformation pattern in contrast to the normally observed pattern during dry extrusion. The rate of travel of material at the billet centre is increasing from the rear of the billet towards the die throat, progressively augmenting the restrained flow of metal near the container wall. Also evident is a tertiary flow in the regions adjacent to the die shoulders, the reason for which is also not known. In some respects, this flow regime bears a certain resemblance to flow during indirect extrusion.

The effect on the deformation patterns of changing the section geometry from circular to square is shown in figures 7.16b and 7.17b. The planar section shown Fig. 7.16 Macrosections of partially extruded billets of super purity aluminium, Direct mode, ER = 40:1, Tj = 425°C.

(sectioned as shown in fig. 4.3) a) Rod b) Square shape c) T-shape d) U-shape

Fig. 7.17 Macrosections of partially extruded billets of super purity aluminium, Direct mode, ER = 40:1, Tj = 300°C.

(sectioned as shown in fig. 4.3) a) Rod b) Square shape c ) T-shape d) U-shape

Fig. 7.18 Macrosections of partially extruded billets of super purity aluminium, Indirect mode, ER = 40:1, T,- = 425°C.

(sectioned as shown in fig. 4.3) a) Rod b) Square shape c) T-shape d) U-shape

151

was taken along the diagonal of the square. The flow pattern resembles a typical direct extrusion flow with the fast moving central zone delineated by a slow moving restrained periphery, resulting in a funnel shaped flow.

In many respects one would expect the flow in a square section to be the same as the flow in circular sections, mainly because there is not a drastic change in the symmetry of the section between a square and a circle. Nevertheless, in the present case, the flow during circular section extrusion being radically different from the normally expected axisymmetric flow, it is difficult to visualise to what extent the non circular sections affect the flow stability during extrusion.

The flow pattern corresponding to the extrusion of T sections is shown in fig. 7.16c and 7.17c for the two temperatures stated, in the direct mode. It can be seen that the flow patterns are different at the two temperatures, in contrast to those obtained in rod and square sections. Although to a certain extent this variation can be explained by the sectioning method used for flow visualisation (i.e. a slight deviation from the required line in taking the section will result in a different area of the flow pattern being seen), the major changes in the deformation patterns indicate that, to some degree, the variations in temperature profile across the section during low temperature extrusion may contribute to changes in the flow patterns with decrease in extrusion temperature. The plane section and the expected metal flow patterns are shown below in fig. 7.19. In fig. 7.16c, though the general flow pattern remains similar to that in rod extrusion, a certain amount of asymmetry about the billet axis can be envisaged, especially in the regions close to the die shoulders (dead metal zone). This can be explained on the basis that, if the flow is expected to be radial from the container wall, at least on the plane of projection parallel to the die plane, the distance ab and cd over which the metal flows are different, even though die design is based on a circumscribed circle diameter. This can to some extent cause the asymmetry observed in the present instance. Due to limitations on the experimental data available, it is difficult to comment any further on the effect of die geometry on flow behaviour. 152

Fig. 7.19 Metal flow expected in a T-section

It is not unreasonable to expect a certain amount of asymmetry and complexity as the flow of the extruded sections becomes more and more asymmetric. On changing the geometry from T to U, much of the deformation pattern and the variations in the flow with temperature are maintained. This indicates that, in terms of complexity of flow during extrusion, T and U are similar.

The shape of deformation zone in the indirect mode is radically different from that seen in the direct mode and now extends in a circular manner from the die/container interface to the die mouth. The size of the deformation zone is reduced, as predicted by the upper bound solution, where a dead metal zone can only be considered to exist at the billet/die interface.

The flow of material proceeds from the sides of the billet to the centre, where it is extruded through the die. A much larger percentage of the billet remains undeformed before reaching the deformation zone and there is no evidence of any 153

billet/container friction. The absence of the heavily sheared region which existed in the direct mode and the change in flow pattern produces an extrudate with a much more uniform structure across its section. The extrudate structure is again essentially fibrous with the grains elongated in the extrusion direction; the heavily sheared peripheral zone seen in directly extruded material is no longer apparent.

In a similar fashion to the direct process, all section shapes produce deformation zone patterns with the fundamental features of a simple rod extrusion. However, circumferential flow again results in subtle differences. This is most evident for the

U-section. Also evident is the change in the penetration height of the deformation zone with circumscribed circle diameter position.

An important feature of the indirect process is that defects from the surface of the billet will be transferred to the surface of the extrudate, due to the nature of the flow during extrusion, in contrast to direct extrusion where the surface of the extrudate originates in the regions of intense shear. However, scalping of the as-cast billet before extrusion removes the main source of such defects, which were, therefore, not a serious problem in the present work.

It is evident from the plates that there is a significant difference in the flow characteristics during steady state direct and indirect extrusion. It should be noted, however, that super pure aluminium was used for these flow visualisation studies making it is difficult to draw any conclusions regarding the flow behaviour of 2024 alloy from these observations, as the flow behaviour of much stiffer alloys such as 2024 and 7075 can be considerably different. The difference in substructures within the deformation zone and extrudate will now be considered.

7.4 Development of substructure during extrusion

Having described the flow patterns of the material during extrusion, the next step is to consider the development of the substructure during the process.

The evolution of the substructure during the direct and indirect extrusion of 154

2024 round sections has been investigated by interrupting individual extrusions at the position marked X on the load-displacement curve shown in fig. 7.20

Extrusion was performed at a reduction of 40:1 at 400°C and at a ram speed of

3mm/s for a round section. Specimens for transmission electron microscopy examination were removed as described in the section 4.4.4 from positions as shown in fig. 7.21. Although shapes were also investigated the differences observed in the substructural development between these extrusions and round bar extrusions were very small.

Fig. 7.20 Position extrusion cycle stopped for structural development study Direct partials

Indirect partial

Fig. 7.21 Location of the TEM specimens for structural study 156

The development of the substructure along the flow line during steady state deformation is shown in fig. 7.22. Micrograph 7.22a corresponds to position A in fig. 7.21a and represents a position just inside the deformation zone. The substructure has already begun to form with a fairly low dislocation density; the large second phase particles are still randomly distributed. The substructure at this stage contains mainly dislocation tangles and poorly formed walls with low internal dislocation density between the walls. Humphreys100 has shown that large lattice rotations occur at coarse second phase particles and the requirement of high dislocation at these sites is beneficial to subgrain wall formation.

Micrograph 7.22b corresponds to position B in fig. 7.21 a and represents a position midway through the deformation zone. The structure consists of well defined subgrains with low internal dislocation densities. This is indicative of the substructure being fully developed between positions A and B. Further changes in substructure beyond B can only occur by a combination of dynamic recovery and polygonisation.

Micrograph 7.22c corresponds to position C in fig. 7.21a and represents a position near the exit of the deformation zone, adjacent to the die mouth. The subgrains are well defined and equiaxed. The particles have been aligned in the extrusion direction but although some sub-boundary pinning has occurred, the subgrain size bears no relation to the interparticle spacing.

Micrograph 7.22d shows that the structure in the extrudate is not too different from that obtained at position C, adjacent to the die mouth. The structure in the dead metal zone has also been examined and is shown in fig. 7.22e. Although the dislocation density is high, there is no evidence of substructure or alignment of particles, which confirms that no significant deformation occurs within this zone.

The subgrain sizes measured for each of these locations is given in table. 7.6. The subgrain size decreases as the material passes through the deformation zone, but shows a slight increase in the extrudate. Assuming that the subgrain size is related to Z as shown in the equations in table 7.3, one would expect a decrease in subgrain size Fig. 7.22 Development of substructure along the flow line during steady state direct extrusion. (ER =40:1, v = 3mm/s, = 400°C)

158

from position A to C due to the increase in the strain gradient and hence strain rate. This will of course be influenced greatly by any temperature variation and thus it would appear that the expected decrease in this case is somewhat balanced by a small rise in temperature along the flow line.

Position Longitudinal Transverse

A 3.10 2.91

B 2.37 2.20

C 2.49 2.25

D 2.55 2.30

E 2.66 2.40

Table 7.6 Subgrain size measurements within steady state direct extrusion deformation zone for positions shown in fig. 7.21a (all values in jim)

The variation in structure during steady state indirect extrusion is shown in fig.

7.23. The change in flow pattern shown earlier, indicates that material in the deformation zone now flows parallel to the die face. Taking this into consideration, specimens were taken from partially extruded billets at the positions shown in fig.

7.21b.

Micrograph 7.23a corresponds to the position at the edge of the deformation zone,

A in fig. 7.21b, and shows that the subgrain structure is fairly well defined but shows a relatively high internal dislocation density. The subboundary walls formed by the interactions of the mobile dislocations are clearly associated with larger, 0.5 to 1.0 p.m particles, which are beginning to be aligned in the extrusion direction.

Micrograph 7.23b shows the structure at the edge of the deformation zone, near the billet container interface corresponding to position B in fig. 7.21b. The subgrains are smaller and more well developed than those at A. Fig. 7.23 Development of substructure along the flow line during steady state indirect extrusion. (ER =40:1, v = 3mm/s, = 400°C)

160

Micrograph 7.23c corresponds to position C in fig. 7.21b, midway through the deformation zone. The subgrains are well developed with a low internal dislocation density and all the particles are aligned parallel to the die face. Micrograph 7.23d shows that a similar substructure exists adjacent to the die mouth, at D in fig.

7.21b.

The substructure of the emerging extrudate is shown in fig. 7.23e. It is evident that the subgrain structures have similar characteristics to those observed in the direct case.

The subgrain sizes at each of the positions are shown in table 7.7. From A to C there is a marginal decrease in subgrain size due to the increase in strain and strain rate as the material enters the deformation zone. From B to C to D there is no real change in the subgrain size, indicating that only a small change in strain rate and temperature occurs along the flow line, similar to that observed in direct extrusion between B and C in fig. 7.22. The increase in subgrain size from D to the extrudate substructure at position E again indicates that a large temperature rise occurs as material flows round and through the die. The effect is likely to be accentuated during direct extrusion, due to the higher strain and strain rate gradients.

Position Longitudinal Transverse

A 2.71 2.51

B 2.44 2.23

C 2.47 2.19

D 2.39 2.07

E 2.51 2.21

Table 7.7 Subgrain size measurements within steady state indirect extrusion deformation zone for positions shown in fig. 7.21b

(all values in p.m) 161

It may be concluded that during steady state direct and indirect extrusion, the variations in substructure within the deformation zones are not significant, considering the changes in the flow characteristics shown in fig 7.16-7.18. It is,

however, evident from the subgrain size measurements that the final product

structure and hence properties are largely dependent on the temperature and strain

rate distribution adjacent to the die mouth, since the steady state structure changes in accordance with the prevailing Z conditions. During direct extrusion, the less recovered material in the periphery of the extrudate derived from the shear zone outlining the dead metal zone results in a greater variation in substructure across the extrudate, although the predicted increase in strain rate within the deformation zone is smaller.

7.5 The effect of solution treatment on the structure of extrudates

Having discussed the variation of the structure in the as-extruded condition, attention will now be focussed on the heat treatments commonly used subsequent to extrusion and their effect on the structure.

All precipitation hardening aluminium alloys are given a solution treatment after extrusion, irrespective of whether the extrudates are press quenched or not, in order to nullify the effects of any variation in properties due to processing. The recommended solution treatment for hot worked alloy 2024 is 20 minutes to 2 hours at 500°C (depending on section size), followed by water quenching to obtain a supersaturated solution16. The effect of this solutionising on the extrudates is shown in fig 7.24. Fully recrystallised structures were found in all extrudates below an extrusion temperature of 350°C. Below this temperature the fibrous structure observed in the central region of the extrudate was no longer retained, static recrystallisation having occurred to form new strain free grains. The recrystallised grains are elongated in the extrusion direction, their size increasing with increasing extrusion temperature. Typical examples of longitudinal and transverse sections are shown in figs. 7.24a,b, taken from a 350°C extrudate, showing full recrystallisation with an average grain size of 0.3 mm in the transverse direction. Above this temperature the extrudates exhibited a fibrous core similar to the as-quenched specimens surrounded by a recrystallised grain structure as shown in figs. 7.24 e,f. Fig. 7.24 Microstructure of extrudates after 500°C solutionising followed by water quench a,b) 350°C extrudate,0.5 hour soak, longitudinal and transverse c,d) 350°C extrudate,1 hour soak, longitudinal and transverse e) 450°C extrudate periphery, 1hour soak, longitudinal f) 450°C extrudate core, 1hour soak, longitudinal

163

The elongated nature of the recrystallised grains in the longitudinal direction is due to pinning of the boundaries by second phase particles. This is shown schematically in fig. 7.32a for low temperature extrudates, where the density of second phase particles is high. For high temperature extrudates however, the propensity for subgrain coalescence to be predominant is high, as shown schematically in fig. 7.32b,c. Additionally, in high temperature extrudates, the number of second phase particles is reduced, hence grain growth is not impeded in either the transverse or longitudinal direction resulting in the equiaxed nature of the grains as shown in figs. 7.4 c,d earlier.

The difference observed after solutionising is probably due to the intermediate particles shown in fig. 7.10; these have been reported17 to restrict the nucleation and growth of statically recrystallised grains in the more highly recovered substructures, causing the fibrous structure to be retained. This phenomenon is known as the "press effect". Continued soaking seems to have had little effect, as shown in figs. 7.24c,d and hence the optimum solutionising time would be between 1/2 and 1 hour at 500°C.

The use of electron microscopy to further study the effects of solutionising revealed some very interesting features, as shown in fig. 7.25. The low temperature extrudates, as shown in fig. 7.25a, were found to be fully recrystallised with the particles aligned in the extrusion direction. There was no evidence of subgrains but helical dislocations were observed as shown in fig. 7.25b These dislocations were almost certainly formed during quenching. It has been reported that high concentration of dislocations of this type are beneficial for the ageing properties as they act as nucleation sites for precipitation90'91.

In high temperature extrusions, however, a retained substructure was observed, as shown by fig. 7.25c. The amount of retained substructure is found to increase with increasing extrusion temperature. This behaviour can be related to the stored deformation energy present after extrusion. This will be higher at low extrusion temperatures because recovery is not as complete as that at higher extrusion temperatures, and so the driving force for recrystallisation during solutionising will be more significant. The structure of the low temperature extrudate will 164

(a)

------► extrusion direction

(b)

extrusion direction

(C)

Second phase particles

Fig. 7.32 Mechanism of grain growth in high and low temperature extrudates Fig. 7.25 Heat treated microstructures of alloy 2024 extrudates a) solutionised, 1hour and water quenched, T e =300°C b) helical dislocations present after solutionising, T e =300°C c) solutionised, 1hour and water quenched, T e =450°C d) Peak aged T6 substructure, 18 hours at 160°C, T e =350°C e) Peak aged T6 substructure, 18 hours at 160°C, T e =450°C f) Peak aged substructure showing a Precipitate Free Zone (PFZ) g) Peak aged T5 substructure, 18 hours at 160°C, T e =450°C

0.5

□ ed g e ♦ c e n tre

unrecrystallised 0 .0 -t— 250 350 450 Initial billet temperature (°C)

Fig. 7.26 Recrystallised grain size vs initial billet temperature for solutionised 2024 direct extrudates

Fig. 7.27 Recrystallised grain size vs initial billet temperature for solutionised 2024 indirect extrudates 167

therefore also contain a greater number of regions of high strain which act as nucleation sites for recrystallisation which is shown by the decrease in recrystallised grain size with decreasing extrusion temperature.

The variation in recrystallised grain size at the centre and edge of the 2024 direct and indirect solutionised extrudates is shown in figs 7.26 and 7.27. It is evident that the grain size decreases from the centre to the edge of the extrudate, the heavily strained periphery providing a larger number of nucleation sites for static recrystallisation. It is also interesting to note that this was also observed in the indirect extrudates, despite the more uniform flow characteristics which one would expect to lead to a uniform structure. In addition, it was found that the grain size in the solutionised direct extrudates was much smaller than in corresponding indirect extrudates; this can be associated with the higher strains characteristic of the direct process, producing more nucleation sites; in addition, as observed previously, the indirect structure is more recovered than the direct structure. The reduction in strain energy associated with an increase in recovery with extrusion temperature reduces the driving force for recrystallisation so that the grain size increases with temperature as shown in both figs. 7.26 and 7.27.

It has previously been reported by Sheppard et al101 that the recrystallised grain size can be directly related to the hot worked subgrain size. Using data for recrystallised grain size obtained by direct measurement and the subgrain dimensions predicted by using the variables obtained in table 7.3, the recrystallised grain size has been plotted as a function of the subgrain size as shown in fig. 7.28. A direct correlation between the subgrain size (d), and the recrystallised grain size (D), is obvious and indicates that recrystallisation occurs primarily by subgrain coalescence or rotation mechanism. A specific number of subgrain coalescences are required to form recrystallisation nuclei, having a high angle boundary.

The results in fig. 7.28 show that for a given subgrain size the recrystallised grains are larger in the indirect extrudates, consistent with the decrease in the degree of peripheral recrystallisation due to the decrease in stored strain energy. Additionally, there is a possibility of subgrain coalescence due to the longer press quench delay during indirect extrusion thus reducing the energy available for 168

nucleation.

The effect of extrusion mode and temperature on the volume percentage of material recrystallised was calculated using equation 7.1 and is given in table 7.8 and shown in fig. 7.29. The figure shows that the percentage recrystallisation increases with decreasing extrusion temperature, this is due to the increased flow stress and hence the driving force for static recrystallisation.

Subgrain size Oim)

Fig.7.28 Recrystallised grain size vs subgrain size for solutionised direct and indirect extrudates of alloy 2024 Extrusion temperature (C)

Fig. 7.29 Volume % recrystallised vs Initial billet temperature for solutionised extrudates

Vol % = a + b.LnZj

Extrusion Extrusion a b cc Mode Ratio

Direct 20:1 -518.9 20.81 0.954

Indirect 20:1 -158.8 6.41 0.935

Table 7.8 The variation in volume % recrystallised after solutionising in direct and indirect extrudates of alloy 2024 100

3 “ 80 “

60 -

40 -

o E solutionised 3 20 - O direct mode > ER*40:1

“T“ 25 26 27 28 29 30 31

Ln(ZI)

Fig. 7.30 Variation of volume fraction recrystallised with Ln (Zi) for shaped sections after solutionising for direct extrusion

£ ■o X o

c o «o

® E 3 O >

Fig. 7.31 Variation of volume fraction recrystallised with Ln (Zi) for shaped sections after solutionising for indirect extrusion 171

If the amount of recrystallisation in shaped extrusions is considered, it can be seen by referring to fig. 7.30 and 7.31 that initially the volume of recrystallised material increases with Z, subsequently tailing off after Z > 30. This phenomenon can be explained in the same terms for solutionised material as for the as-extruded material, detailed earlier in section 7.2.1. Again, the benefits of indirect extrusion can be seen in the lower volume of recrystallised material obtained. Full recrystallisation is shown in fig. 7.15 for T and U-shaped sections; the trend in the grain size under different extrusion conditions was found to be the same as for the rod extrusions at 20:1, described earlier.

7.6 The effect of ageing on the microstructure, of the extrudates

The effect of age-hardening on the structure is shown in fig. 7.25d-g for material which had been solutionised and aged for 18 hours at 160°C. The feature common to all extrudates, irrespective of the extrusion temperature, is the presence of peak aged precipitates and coarse dispersoids. Apart from the proportion of retained substructure, the extrusion temperature had no apparent effect on the structures observed. The chequered patterns observed in fig.7.25d can be attributed to the elastic coherency strains associated with the 0" precipitates.

In the extrudate produced at 450°C, as shown by fig. 7.25e, it can be clearly seen that substructure has been retained after solutionising and ageing. The misorientation between the subgrains has been reported17 to be between 0.1° and

3°. In order to ascertain whether any change in the substructure occurred during the ageing treatment, transverse subgrain sizes in all the T1, T5 and T6 temper conditions were compared, as shown in table 7.9; subgrain sizes for the T1 temper were predicted using equation 7.1. These results indicated that some subgrain growth and coalescence may have occurred during the solution treatment. It is clear that this coalescence would not have been sufficient to initiate recrystallisation, although more definite conclusions cannot be drawn due to errors inherent in the

measurements. 172

ER = 20:1 T=450C, v=6.9mm/s

Temper Direct mode Indirect mode pm pm

T1 2.56 2.37 T5 2.58 2.36 T6 2.66 2.45

Table 7.9 Comparison between T1, T5 and T6 subgrain dimensions for both modes of extrusion

The overall ageing characteristics have been reported86 to be little affected by the presence of retained structures, although as the micrograph shows, some preferential precipitation occurs at the subgrain boundaries. This may have an undesirable effect on the corrosion and fracture properties which are very much dependent on the nature of the precipitate distribution and solute concentration at grain boundaries. A higher magnification fig. 7.25f shows that a precipitate free zone approximately 0.05 pm wide is associated with the subgrain boundary, which probably formed due to the depletion of solute from the surrounding matrix caused by preferential precipitation at the subboundary. It could be seen in this micrograph that within the grains there was a fine distribution of spherical precipitates approximately 0.1-0.2 pm in diameter. The presence of this fine precipitate distribution indicates that GPB type zones are present in this material, in agreement with the findings of previous workers2 6 *2 7 .

Examples of the microstructures from the press quenched and aged extrudates processed at 450°C are shown in fig. 7.25g. The possible increase in solid solution content prior to ageing might be expected to improve the ageing response and hence precipitate distribution, which is partly confirmed by the micrographs. If this is the case, it implies that if the age hardening in the T5 temper can be increased by a presolution soak treatment, then the incremental increase in strength which may be derived from the as-quenched substructures may produce extrudates with equivalent 173

properties to those of T6 extrudates. An example of the subgrain size found in a T5 extrudate is also shown in table 7.9 and reveals some interesting facts when compared to values at the T1 and T6 tempers. The results imply that little or no subgrain growth has occured during the course of the ageing treatment, effectively counter-balancing the possible reduction in age hardening due to the absence of a solution treatment by providing greater substructural strengthening. 174

7.7 Conclusions

1. The structure of 2024 alloy In the as cast condition consists of dendrites interspersed with eutectics.

2. The extruded rod exhibited a fibrous core at all extrusion temperatures with a strained periphery or recrystallised grains at the periphery. The extent of this recrystallisation increases with extrusion ratio, initial billet temperature and ram speed. In shaped extrusion the extent of recrystallisation was found to increase with increase in complexity of shape. Also the recrystallised layer was deeper where there was a reentrant corner in the section due to the increased shear occurring in

these regions.

3. Investigation of the structural development of alloy AA2024 shows that dynamic recovery is the operative high temperature restoration mechanism. Increased temperature and decreased strain rate increases the size and perfection of the

subgrains. The presence of a large amount of precipitate in 2024 alloy pins the sub-boundaries and thus the subgrain size is smaller. The variation in substructure within the steady state direct and indirect deformation zone results in a more uniform substructure across and along the indirect extrudate. The structures observed in the torsion samples have in general been consistent with those found in extrusion. The subgrain size'd' may be related to the process conditions by:

d-1 = a.LnZc +b

4. The structure in the extrudate after solutionising and quenching is dependent upon

the initial extrusion temperature. With low extrusion temperatures, complete recrystallisation occurs during solutionising, but with high extrusion temperatures

the substructure is retained, the proportion of which increases with increasing extrusion temperature. The solution treated extrudates all showed varying degrees of recrystallisation; those extruded below 350°C were fully recrystallised. Indirect extrusion reduces the volume percent recrystallised. The increase in recrystallised grain size with extrusion temperature can also be directly related to the increase in 175

subgrain size. The effect of ageing the solution treated material is to produce GP zones and transition precipitates, hence improving the strength.

5. The effect of artificial ageing on the solutionised extrudates is to produce transition precipitates. In the low Z partially recrystallised extrudates the presence of a retained substructure provides sites for preferential precipitation. Artificial ageing of the high extrusion temperature press quenched extrudates results in a pronounced ageing reaction. 176

8.0 Room Temperature Properties 177

8.1 Room temperature properties of the extrudates

It is evident from the previous chapter that the choice of extrusion process and process conditions had a marked influence on the structure of the material and hence the properties. In this chapter, the relationship between the processing conditions and the properties in the T1, T5 and T6 temper have been investigated for both direct and indirect extrudates. Tests were carried out on 20:1 rod-shaped extrusions unless otherwise mentioned.

8.1.1 Effect of extrusion condition on the hardness properties in the as extruded condition

The variation of hardness with extrusion temperature and extrusion mode is shown in fig 8.1. The hardness measurements were taken at the centre, mid radius and periphery of the extrudate to determine the variation in properties across the section. The results tabulated in table 8.1 report an average of the readings, since

little or no variation was found in hardness across the extrudate. There is no apparent difference between the values obtained for direct and indirect extrusion. The alloy does, however, show an increase in hardness with increasing extrusion temperature.

The hardness values of any press quenched extrudate are in general due to contributions from the following factors: the substructure of the extrudate, formation of precipitates on natural ageing and solid solution hardening due to the elements (e.g. in this material mainly due to copper and magnesium) present. The current observations indicate that the contributions due to the former two factors are more significant in the material, as the extrudates when tested were in the naturally aged condition. Hence, the increase in hardness is likely to be due to a combination of substructural strengthening and precipitation hardening associated with the formation of GP1 and GPB zones at room temperature, which act as obstacles to dislocation motion2 6 *2 8 . Furthermore, with increase in extrusion temperature, the increasing amount of solute taken into solution will reprecipitate during subsequent ageing of this extrudate. Hence it would seem likely that the mechanical properties of as-extruded 2024 alloy would improve with higher extrusion temperatures. 178

150

9 140 xw 0)01 ® 130

■o1 _ £(0 kco 120 xo o > 110

100 200 300 400 500

Temperature (C)

Fig 8.1 Hardness variation versus initial billet temperature

Run Extrusion Extrusion Hardness Hardness Code Mode Temp°C (as extruded) (after solutionising)

JS007 Direct 450 142 100

JS008 Direct 400 131 90

JS009 Direct 350 112 89.6

JS010 Direct 300 103 91.6

JS011 Indirect 450 141 97

JS012 Indirect 400 133 90

JS013 Indirect 350 110 89

JS014 Indirect 300 105 92

Table 8.1 Hardness values of alloy AA 2024 extrudates 179

The effect of increasing extrusion ratio on the hardness is shown in fig. 8.2 and table 8.2. Again, the difference between direct and indirect extrusion is insignificant. There is, however, an increase in hardness with increasing extrusion ratio. This effect is marginal at low temperatures but prominent at high temperatures. This increase is probably related to the greater temperature rises at higher extrusion ratios, which assists the solutionising effect.

8.1.2 Variation of hardness along extrudate length

The variation of the structure along the length of the extrudate was also discussed earlier in section 7.2.2 and it was proposed that the differences could be related to the rise in temperature in the extrudate during extrusion. The effect of this temperature variation on the hardness of the extrudate was investigated by taking hardness along the length of the extrudate. The results are shown in fig. 8.3 for direct and indirect extrusion at an extrusion ratio of 20:1 and initial billet temperatures of 450°C and 300°C.

Although the variation of the hardness along the length of the extrudate could not be described as significant, a slight increase in hardness along the length was apparent. This increase is comparatively larger in the lower temperature extrudates than in the higher temperature material. This is consistent with the concept of an increased amount of solute due to the increase in the deformation zone temperature towards the end of an extrusion cycle caused by the temperature rise. As observed earlier, the hardness was greater for higher initial billet temperatures and extrusion ratios. However, it would appear from these results that it is not the instantaneous increases in temperature taking place during extrusion, but the initial extrusion conditions which have the greatest influence on the hardness in the extrudate.

It is interesting to note that at the end of each indirect extrudate the hardness showed a sharp decrease, tending to confirm that due to the configuration of the indirect tooling there is a considerable delay before the extrudate ends are quenched. 180

Extrusion Ratio Extrusion Extrusion Mode Temperature 20:01 30:1 40:1 50:1 (C)

Direct 450 1 4 2 1 4 8 151 1 5 5

Direct 300 1 0 3 1 0 7 1 1 0 1 1 6

Indirect 450 141 1 4 9 1 5 3 1 5 7

Indirect 300 1 0 5 1 0 5 1 1 3 1 1 5

Table 8.2 Hardness values of extrudates for both modes at different extrusion ratios

z z > co CO 9 C V i m £a k.co o 2C o >

Fig 8.2 Hardness variation versus extrusion ratio Fig.8.3 Variation of hardness along extrudatelength Fig. 8.4Variation of microhardness across extrudate Vickers hardness (VHN) 150 120 130 - ■ Indirect (450C) Direct (450C) Direct (300C) Distance along extrudate Distance

182

8.1.3 Variation of hardness across the extrudate

It was mentioned earlier that there was little or no variation across the extrudate section. However, as the Vickers hardness machine takes values over a relatively large area in relation to the extrudate geometry, it was decided to take measurements across the section of selected samples using a microhardness tester. The results are shown in fig. 8.4 and the results are comparable to those obtained from the Vickers hardness machine.

It was noted in section 7.2.2 that the subgrain size increased from the centre to the edge of the extrudate. Hence it could be expected a corresponding decrease in microhardness to be shown, but it was evident that the microhardness of the as-extruded material increases from the centre to the edge of the extrudate for all extrusions conditions, particularly at low extrusion temperatures. This contradictory observation can be attributed to a number of factors. For instance, it was observed that although the subgrain size increased at the edge of the extrudate, these subgrains had a higher internal dislocation density than those in the centre of the material, an effect which was especially evident at low extrusion temperatures; this may increase the microhardness. This may be attributed to the decrease in quench rate from the edge to the centre of the extrudate.

8.1.4 Effect of solution treatment on the hardness

It was shown in section 7.3 how the extrudate structures are modified during solutionising. In order to determine the hardness of the alloy in this condition, measurements were taken immediately on quenching after solution treatment. The results are given in table 8.1 and are also shown in fig. 8.5 revealing that the hardness is approximately constant for both direct and indirect extrudates at all temperatures. The hardness after solutionising is also considerably lower than the as extruded hardness at all extrusion temperatures as shown in table 8.1.This drop in hardness also decreases with the decrease in the extrusion temperature.

This loss of the hardness can be explained on the basis of the phenomena taking place during solutionising of the extrudate. Solutionising is a high temperature heat treatment; in this instance temperatures used were in the range of 500°C. During 183

solutionising, two physical processes occur. All soluble elements are taken back into solution, while at the same time the deformation substructure undergoes large scale static recovery and/or static recrystallisation. As a result, the hardness of the extrudate will drop. On the other hand, as the solute elements go back into solution, a certain amount of solid solution strengthening can contribute to increase the hardness. The net hardness value will be a combination of these two effects; the drop in hardness due to the loss of deformation substructure by static recovery and/or recrystallisation can be expected to be much higher then the increase due to solid solution hardening, the net effect therefore being one of lower hardness in the solutionised condition. Additionally, fig. 8.5 shows that the trend in the hardness with increasing extrusion temperature is retained after solutionising. This effect can be easily explained, in that only static recovery occurs in the high temperature extrudates during solutionising, while in the low temperature extrudates large scale static recrystallisation removes the deformation substructure. This necessarily means that in the high temperature extrudates a certain amount of substructure will be retained after solutionising, which was indeed observed; this can result in higher hardness than low temperature extrudates.

Temperature (C)

Fig. 8.5 Hardness variation after solutionising

It has been reported that the hardness of the solution treated material can be 184

greatly improved by naturally ageing for 28 days or more16. However, due to the length of this natural ageing period, alternative ageing treatments have been developed.

8.2 Ageing characteristics of 2024 alloy extrudates

The improvement in the mechanical properties of the extrudates by artificial ageing at elevated temperature was discussed in section 2.3.3, but it would be useful to summarise the mechanisms of strengthening associated with precipitation hardening. In aluminium-copper alloys these are: a ) hardening due to coherency strains b ) chemical hardening c ) modulus hardening d) dispersion hardening

In an alloy such as AA 2024 the first three mechanisms are associated with zone formation, while dispersion hardening is linked with the formation of 6 and 6' precipitates. Since it is likely that for a given precipitate hardening treatment there will be a combination of zones and precipitates, it is necessary to determine the optimum ageing temperatures and times in order to achieve the maximum possible benefit from these strengthening mechanisms.

The artificial ageing characteristics of the alloy were determined for 4 different extrusion temperatures at a constant extrusion ratio (20:1) in both modes by ageing them in the as-extruded condition or after solutionising for half an hour at 500°C.

The ageing treatments were conducted at 120°C, 160°C and 180°C. The specimens were removed from the oven at set time intervals for hardness measurements to be made. The hardness measurements obtained in the solution treated condition and aged condition are given in full in Appendix 2, and the results obtained for direct extrusion at temperatures of 300°C, 350°C, 400°C and 450°C are shown in figs.

8.6-8.8 at ageing temperatures of 120°C, 160°C and 180°C respectively. It should be noted that identical characteristics were observed for the indirect extrudates. 185

The general trend observed in the evolution of hardness with ageing time showed the hardness increasing with time to peak and then dropping. However, fig. 8.6 shows that peak hardness is not achieved even after 200 hours ageing at a temperature of 120°C. Previous workers have reported that for alloy 2014 the peak is obtained after 800 hours, whilst for ageing at 100°C, peak hardness is only

attained after about 2000 hours2 6 *8 6 . The slow initial ageing response indicates that the early stages of ageing are dominated by the formation of GP1 or GPB zones

and the nucleation of 6" or S" precipitates does not occur until the later stages. The ageing response is markedly improved at 160°C, as shown in fig. 8.7, where the peak hardness is achieved after 18 hours. At 180°C, as shown in fig. 8.8, the peak for the alloy is attained after 8 hours. Increasing the temperature has accelerated the

Fig. 8.6 Ageing characteristics of solutionised 2024 direct extrudates at 120°C 186

Fig. 8.7 Ageing characteristics off solutionised 2024 direct extrudates at160°C

Fig. 8.8 Ageing characteristics of solutionised 2024 direct extrudates at180°C 187

transition of GP1 zones to 0" and it is likely that some nucleation of the 0' precipitate has occurred. The decrease in hardness after the peak is relatively slow and can be attributed to the further nucleation of 0' precipitates which subsequently coarsen, resulting in the loss of coherency with the matrix. The effect of ageing at 180°C is shown in fig. 8.8, emphasising the increase in rate of transition with temperature; peak hardness was obtained after only approximately 8 hours. The particle growth has accelerated at this temperature and so softening subsequent to the peak is more significant. Overageing is associated with the continuing transformation of 0" to 0' and the growth of 0*. This is followed by a decrease in hardness which results from the annihilation of the zones and the formation of large precipitates.

The time taken to attain and the magnitude of the maximum hardness are observed to decrease with increasing ageing temperature. The effect of the extrusion temperature on the hardness of the solution treated and aged material is only minimal, although it would appear that there is a tendency for a small increase in hardness (about 10Hv) with increasing extrusion temperature to be shown for both direct and indirect extrudates. This increase in hardness at high temperatures can result from the contribution of the retained substructure. There is no noticeable difference between the hardness values for the corresponding direct and indirect extrudates.

The hardness measurements obtained after ageing this alloy in the as extruded conditions are also given in full in Appendix 2, the results for direct extrusions extruded over a range of temperatures are shown in figs. 8.9-8.11 for ageing temperatures of 120°C, 160°C and 180°C respectively. Again identical characteristics were obtained for the indirect extrudates.

The ageing curves obtained were similar to those obtained from the solutionised extrudates. However, the results show a wide dependence on extrusion condition, maximum hardness only being achieved at higher initial billet temperatures, and the values generally being lower than those obtained for solutionised materials. At lower temperatures a progressive decrease in strength was observed. Fig.8.10 Ageing characteristics of as extruded2024 direct extrudates Hardness (Hv) 3 Hardness (Hv) i.8.9 Ageing characteristicsof as extruded2024 direct extrudates at160°C at120°C 188 189

Fig. 8.11 Ageing characteristics of as extruded 2024 direct extrudates at180°C

The wide variations observed in the ageing behaviour of the extrudates without solution treatment originates from the prior thermomechanical history. In the present case the billets were homogenised and furnace cooled prior to extrusion. During the homogenisation treatment, the segregated second phases are taken into solution, but the furnace cooling treatment allows reprecipitation of the solute elements in a finer form. Such material, when subsequently reheated in the induction heater prior to extrusion, will only undergo a limited amount of solutionising. As this depends on the preheat temperature, the billets which have been subjected to a higher reheat temperature will have a higher solute concentration available for precipitation on subsequent ageing, giving rise to the observed variation in the ageing behaviour of the extrudates with extrusion temperature.

It may therefore be concluded that by using the "press quenching" and ageing treatment described above, it is possible to produce a vast improvement in the hardness of the extrudate. This phenomenon has obvious benefits but economic considerations must also be made when combining the solution heat treatment with the hot working process in a commercial environment. Although this programme was not verified in the present research, it has been reported that billets held at the 190

required temperature in the furnace for 30 minutes prior to extrusion produced equally as good hardness values16*17. Thus at high extrusion temperatures most of the hardening elements, i.e. copper and magnesium, will be in solid solution and may be retained in solid solution in the extrudate by press quenching. Therefore, it is perfectly feasible to age harden a product which has been extruded at high temperatures and quenched at the press. However, it should be remembered that other factors such as surface quality may impose limitations on the use of high extrusion temperatures.

8.3 Tensile tests of 2024 alloy extrudates

The tensile testing programme was carried out in order to establish the effect of extrusion conditions on the properties in the T1, T5, T6, T51 and T510 temper conditions. Again the effect of both modes of extrusion was established. It must be noted however that before the tensile tests were made, the specimens were unavoidably naturally aged for 2 weeks.

The tensile specimens were all machined from sections of the extrudates in the steady state region corresponding to 0.3 to 0.5L which showed no signs of surface cracking. The presence of even the finest cracks can be deleterious to the tensile properties. The tests were conducted at room temperature and a full list of the results is given in Appendix 3

8.3.1 Stress strain characteristics of_allav 2024

Typical stress strain curves obtained from testing alloy 2024 are shown in fig. 8.12 for material that has been extruded at 450°C and 300°C in the T1 temper and 450°C in the T6 temper. The figures show an initial approximately linear rise in stress followed by an extended plastic region, with the stress increasing to a maximum value at which necking occurs. The stress then steadily decreases until fracture occurs. In addition, the stress strain curves for the material in the T1 temper contained serrated flow behaviour which was absent in the T6 temper extrudate. The appearance of instability in the curve at some strain subsequent to yielding 191

has been recognised by many workers investigating aluminium-magnesium alloys3 9 *95-97. This phenomenon is known as serrated yielding or the Portevin-Le Chateiier effect and is manifested by the appearance of jerks in the plastic region of the stress strain curve. Cottrell93 suggested a mechanism whereby moving dislocations are repeatedly locked by solute atoms. When a sufficient stress is applied, the dislocations break away and the load drops. The solute atoms, in this instance magnesium, then diffuse through the matrix to catch up the dislocations and repin them; the process is then repeated. The magnesium atoms, due to their large size and correspondingly slow diffusion rate, would seem to render this mechanism unlikely. However, the increased vacancy concentration resulting from the deformation will permit an enhanced rate of diffusion and make the mechanism feasible. Tutcher39 observed that an increase in the strain rate of the test caused the serrations to appear at lower strains and stresses.

Fig. 8.12 Stress strain curves for alloy 2024

Finally in the T6 peak aged material, serrated flow is not observed. This is consistent with the fact that most of the copper and magnesium come out of solution to precipitate as the CuMgA^-based S phase, both homogeneously in the matrix, as 192

well as heterogeneously at the low angle subgrain boundaries.

The strain hardening exponents determined for this alloy, when extruded over a range of extrusion temperatures are given in table 8.3. It is apparent that the strain hardening exponent of the extrudate increases with increasing extrusion temperature; this may be attributed to the greater amount of zone formation during natural ageing of the as-extruded material which is in turn dependent on the amount of supersaturation, which thus increases the number of zones/dislocation interactions upon straining. This reasoning is supported by the observations that the strain hardening exponent of the peak aged T6 extrudate is comparable to the material extruded at 450°C, and only shows a limited decrease with extrusion temperature due to the presence of 0" precipitates.

g = Aem

Extrusion 300C 350C 400C 450C Temperature

Direct A 124.8 136.2 126.1 142.1 m 0.370 0.382 0.491 0.512

Indirect A 135.7 133.9 147.7 146.7 m 0.377 0.418 0.423 0.492

T6 Temper A 146.6 149.9 154.5 155.6 m 0.483 0.512 0.544 0.571

Table 8.3 Variation of strain hardening exponent of 2024 alloy extrudates

The relationship of m to Ln (Zj) for the as extruded material in the T1 temper was also examined since a linear relationship was reported by Sheppard52 and Paterson16. The results obtained are given below: 193

Direct m = -0.025Ln( Zj) +1.152 cc 0.881

Indirect m =-0.016Ln(Zj) + 0.895 oc 0.873

and are comparable to those obtained from previous work by Paterson16 and Vierod17 for alloy 2014. The low correlations obtained in here are probably due to insufficient data points.

8.3.2 Tensile properties of_allov 2024 m . temper)

The effect of extrusion conditions on the T1-temper tensile properties are shown in figs. 8.13-8.18 for extrusion ratios of 20:1 in both the direct and indirect mode.

The extrusion conditions have been plotted as a function of Ln Z\ since, as shown earlier, this parameter can be directly related to the microstructure of the extrudates. It is evident from the data that the proof stress decreases from approximately 400 MPa to 330 MPa as Ln Z\ increases, and also that there is no significant variation between the strengths in the direct and indirect mode. The values of ultimate tensile stress also decrease from 470 MPa to 420 MPa in the same range of Ln Z\. These variations in the mechanical properties are consistent with the hardness variations discussed earlier in section 8.1.1.

It was suggested earlier that the mechanical properties of the extrudates depend upon a balance between substructure, solute strengthening, and precipitation strengthening. The subgrain size in the extrudate was shown earlier to decrease with increasing in Ln Z\ and it would therefore be expected that the influence of substructural strengthening is greatest at high Ln Zj (low extrusion temperature).

However, at low temperatures, since the concentration of copper and magnesium in solution in aluminium is limited, the influence of the solute strengthening and the precipitation hardening will be minimal. At high extrusion temperatures (low Ln

Zj) the opposite applies and it can be seen from these figures that solute hardening and precipitation hardening were the more dominant hardening mechanism.

The effect of extrusion conditions on the T1-temper tensile properties will i. .4 ro Srs v L(i fr niet xrdts (ER=20:1) extrudates indirect for Ln(Zi) vs Stress Proof 8.14 Fig. Proof Stress (MPa) i. .3 ro Srs v L(i fr iet xrdts (ER=20:1) extrudates direct for Ln(Zi) vs Stress Proof 8.13 Fig. Temper e p m e -T 1 T ■ - mper p em 5-T T er p • em T6-T □ 194

U.T.S. (MPa) U.T.S. (MPa) i. .5 .. v L(i fr iet xrdts 04 (ER=20:1) 2024 extrudates direct for Ln(Zi) vs U.T.S 8.15 Fig. i. .6 .. v L(i fr niet xrdts 04 (ER-20:1) 2024 extrudates indirect for Ln(Zi) vs U.T.S 8.16 Fig. 4 6 8 0 2 34 32 30 28 26 24 Z n L Temper e p m e -T 1 T er p em ■ 5-T T er p • em 6-T T □

195 Elongation (%) i. .7 Eogto v L(i fr iet xrdts (ER=20:1) extrudates direct for Ln(Zi) vs Elongation % 8.17 Fig. i. .8 Eogto v L(i fr niet xrdts (ER=20:1) extrudates indirect for Ln(Zi) vs Elongation % 8.18 Fig. Temper e p m e -T 1 T er p em ■ -T 5 T r e p ♦ m e 6-T T H

196 197

therefore depend on whether the strengthening mechanism in the extrudate is a function of substructure strengthening (as is the case at high Ln Zj) or the amount of solute (as at low Ln Zj). However, at high Ln Zj, the maximum increase in strength will be limited by the press capacity, whilst at low Ln Zj the increase is limited by the requirements of an adequate surface quality, which in the present work has been associated with the onset of surface cracking. However, because the samples have been naturally aged, the increase in strength shown in the figs. 8.13-8.16 indicates that natural ageing is a more effective strengthening mechanism than either substructural or solution strengthening; therefore extrusion at the lowest Ln Zj (i.e. highest temperature) will result in the highest T1 -temper strength for press quenched extrudates because the high super saturation in this condition promotes maximum natural ageing. It is important to remember, however, that the overall strength at low Ln Zj will still have the substructural strength contribution. This is shown by the values of proof stress and ultimate tensile strength at Ln Zi < 25, which are higher than the values quoted9 for the T4 temper (the condition which corresponds to the maximum increase in strength due to natural ageing after solutionising and quenching; proof stress - 290 MPa, ultimate tensile strength - 425 MPa) strengths of 2014. Therefore, control of the subgrain structure is still an important consideration even though the solute is the principal source of strengthening, and an optimum combination of these factors must be sought.

It was apparent from the results that the tensile strengths of the indirect extrudates were similar to those of the direct extrudates over the entire range of Ln

Zj studied. Although this is consistent with the hardness values, in the light of the variations observed in the substructure, it could be expected that the indirect tensile properties would be lower. Although this inconsistency cannot be accounted for d irectly, a possible explanation m ay be that the uniform nature of the flow in indirect extrusion produces more uniform properties across the extrudate section and the lower tensile strength of the direct extrudate reflects its greater anisotropy.

The effects of peripheral grain growth on the tensile properties can be assessed if the tensile strength is plotted against extrusion ratio. Generally, the strength, as shown in fig. 8.19, is observed to increase with extrusion ratio initially marginally and then to decrease. This can be attributed to the increased degree of substructural 198

strengthening occurring as the strain increases. For an initial billet temperature of 350°C, a maximum strength is reached at an extrusion ratio of 30:1, while extrusion ratios of 40:1 and 50:1 show a decrease in strength.

Fig. 8.19 Variation of tensile properties with extrusion ratio for direct and indirect extrudates of alloy 2024

it was earlier shown that at these high extrusion ratios a significant degree of peripheral recrystallisation and grain growth has occurred. Material from this region is retained in the tensile test sample, since very little of the extrudate exterior is removed during specimen manufacture at high extrusion ratios. This is therefore the probable cause of the fall off in strength at high extrusion ratios; theoretically it could be expected that the strength would have otherwise continued to increase with increasing extrusion ratio. Again, a similar pattern was seen for indirect extrusion with increase in extrusion ratio. Although the strengths obtained were slightly higher in the lower extrusion ratio, the difference in the strengths between direct and indirect becomes insignificant at higher extrusion ratios. This can be attributed to the fact that the temperature rise associated with indirect extrusion is lower at lower extrusion ratios than in direct extrusion, and any loss of 199

strength due to substructural coarsening is therefore minimised in indirect extrusion. However, at higher extrusion ratios, in addition to the increased temperature rise, a greater volume of recrystallised material is present. This results in a considerable loss of strength, and accounts for differences in the strength between direct and indirect extrusion. These results are not however conclusive, since tests were only carried out on material from one extrusion temperature.

8.3.2 Tensile properties of 2024 allOY^fT6 temper)

The variation of the tensile properties of 2024 alloy extrudates in the fully heat treated condition has been investigated by testing specimens which had been extruded over the range temperatures and peak aged at 160°C for 18 hours.

The results obtained are listed in detail in Appendix 3 and are shown graphically in figs. 8.13-8.18 for both direct and indirect extrusion. Since the effect of the solution treatment is to remove the differences in solute supersaturation between the different extrudates and reduce the effect of substructural strengthening in the extrudates one would expect the tensile properties to be approximately equal for materials having different thermomechanical history. However, it can be seen that there is a considerable amount of variation in the proof stress and ultimate tensile strength with increase in Ln Z\ and it is apparent that the values for the direct extrudates are not very much different to those from the indirect extrudates. The data agrees with the hardness data in that the strength decreases with increase in the

Ln Zj values. Substructural examination did not indicate any significant differences in the hardening precipitates (size or density) observed in the materials extruded at different temperatures. Hence this decrease has to be due to the loss of deformation substructure on solutionising, either by static recovery or static recrystallisation, as already explained in the earlier chapter on hardness evolution during ageing. At high Ln Z; all the extrudates undergo complete recrystallisation. Hence there may be some contribution to strength from grain boundary strengthening according to the

Hall-Petch equation: 200

G = a 1 + kD'0-5 ...... 8.1 where; D = recrystallised grain size

The equation would predict a trend of increasing stress with decreasing grain size, however, it is well known that in aluminium alloys grain boundary strengthening plays only a minor role in the overall strengthening of the material. Therefore the major strengthening in this alloy in the T6 condition would be due to precipitation hardening and only to a minor extent due to the retained substructure in the high temperature extrudates.

It is interesting to note that the data obtained in the course of this investigation

indicate that the mechanical properties in the T6 condition are superior to those

obtained in the T1 and T5 conditions under all conditions. This is in contrast to an

earlier report by Paterson16 that at low Ln Zj, the T6 strength levels are lower

than the T1 strength levels in an AA2014 alloy. This discrepancy is possibly due to the fact that the alloy AA2014 may have recrystallised on solutionising under all

conditions of Ln Zj, resulting in the complete loss of the retained substructure

contribution; in the present case, where AA2024 is under investigation, this

phenomenon is restricted to high Ln Zj. Additionally, the development of higher

mechanical properties in the T6 condition over the T1 and T5 tempers can be

attributed to more complete precipitation hardening in the T6 condition than in the

other tempers.

The ductility of all the T6 and T5 extrudates is lower than those of the T1 temper. The ductility is clearly related to the nucleation and growth of voids, which is

assisted by the presence of precipitate particles, making the precipitate size and

distribution important. Thus in the T1 temper, though zones and coarse constituent particles are present, the nucleation, growth and linkage of the voids is delayed to a larger strain than in the T6 and T5 tempers, indicating that the tensile ductility is more a function of precipitate hardening and particle distribution than the deformation microstructure of the extrudate. 201

8.3,3 Tensile properties of alloy 2024 - ( T5 temper)

The alternative to a full solution soak and ageing treatment is to use the billet preheat cycle as the solution treatment, and age the press quenched material artificially immediately after extrusion. This process is feasible if economy rather than property optimisation is the dominant consideration.

The tensile properties of the T5 extrudates given in figs. 8.13-8.16 show a general increase in ultimate tensile strength and proof stress with increasing extrusion temperature, principally due to the solutionising effect of the preheat temperature on the emergent extrudate. The results show a wide dependence on the extrusion conditions, maximum ultimate tensile strength and proof stress being achieved at 450°C. At lower extrusion temperatures a progressive decrease in strength is observed, reflecting the decrease in solid solubility with temperature.

In conclusion, the T5 tensile strength of the extrudates considered does not quite match the equivalent T6 extrudates at low Z; conditions due, to the greater degree of age-hardening and substructure strengthening associated with these latter extrudates. However, at the low Z\ conditions below the specific values mentioned, the T5 strengths are greater than those obtained in the fully recrystallised T6 extrudates at high Z\ and the naturally aged T1 extrudates at the lowest Zj conditions.

As for the T1 and T6 temper, the variation in peripheral recrystallisation with extrusion conditions shown in the as extruded material in section must also be considered when choosing the optimum process conditions, since for all the extrudates the increase in T5 strength is accompanied by an increase in the recrystallised depth, although it should be noted this did not increase during the ageing treatment.

8.3.4 The effect of a preaaeina stretch

In the absence of quenched-in defects in aluminium-copper-magnesium based alloys, a preageing stretch has the effect of accelerating the nucleation of the S and

0" phases by providing nucleation sites in the form of dislocation defect structures. 202

In the present investigation the extruded material was solution treated, water quenched, stretched (1%) and then artificially aged at 160°C for 18 hours. In a further experiment, the extrudate was stretched (1%) and then peak aged as before without a prior solutionising treatment.

The results obtained are shown in figs 8.13 and 8.15 for the direct mode only. It is interesting to note that solutionising in this instance has fulfilled its role and there is very little difference in the ultimate tensile strength and proof stress for the different extrusion conditions. However, the press quenched and aged samples did show a wide dependence on the extrusion temperature, again reflecting the solutionising effect at high temperature extrusions. Although the values obtained by the use of this additional process were increased significantly over the T6 conditions, the benefit may not be sufficient to render this treatment of the material economical.

8.4 Fracture toughness properties of alloy 2024

It is accepted that the fracture toughness of metals generally decreases as the yield strength is raised. For high strength aluminium alloys, the increase in strength on ageing can result in a decrease in toughness, to the extent that alloys such as those in the 2XXX series are often used in the overaged condition to enhance the fracture toughness properties^1. Several review papers have listed the important metallurgical factors which affect the toughness in aluminium alloys, but have largely ignored the effect of process conditions7»30,31 # |n the present work, therefore, the effect of process conditions on the fracture toughness of alloy 2024 in the T1, T5 and T6 tempers has been investigated.

The specimens for the fracture tests were machined from material which had been extruded over a range of extrusion temperatures for a given extrusion ratio; again it must be noted that specimens were unavoidably naturally aged for two weeks. At least three tests were conducted in order to confirm their reproducibility. The fracture toughness has been evaluated using the equation described in section 4.4.6; all the raw data are listed in Appendix 4.

The effect of the initial extrusion conditions on fracture toughness in the T1, T5 and T6 temper are shown in figs. 8.20 and 8.21 for direct and indirect extrudates. It i. .1 hr rd rcue ogns v L (i fr niet extrudates indirect for (Zi) Ln vs toughness fracture rod Short 8.21 Fig. Fracture Toughness (MPa) Fracture Toughness (MPa) i. 20 Sot o fatr tuhes s n Z) o drc extrudates direct for (Zi) Ln vs toughness fracture rod Short 0 .2 8 Fig. 203 204

is apparent from these graphs that the extrudates show a marked increase in fracture toughness at low Z\ as indicated by the curves. These results are somewhat surprising since the extrusion conditions also correspond to the highest proof stresses and highest strain hardening coefficients, deformation characteristics which are normally associated with a decrease in fracture toughness3 1 . If the above trend is correct, then extruding at low Ln Zj conditions to retain the substructure during subsequent heat treatment not only increases the tensile strength but also increases the fracture toughness in the transverse plane.

The values of fracture toughness obtained are of a similar order to those obtained by previous workers for a similar alloy5*^» 16,17,28; \\ \s worth noting that the technique used by previous workers was not the same as that used in the present investigation, but was based on measurements of crack opening displacement in three point bend tests. It was apparent that the best fracture toughness was produced in material which had been extruded at high temperature (i.e. low Ln Zj ). The rapid decrease in fracture toughness with increasing Ln Zj is contrary to the expected trend with the yield strength. It is evident, therefore, that the fracture toughness must be associated with other factors such as the microstructural influence on the crack path.

For a 7000 series alloy a large decrease in the plane stress fracture toughness from an unrecrystallised product to a recrystallised material has been reported, due to a change from the transgranular to intergranular fracture mode7 . In this case however, the fracture surfaces were relatively flat and smooth, and showed no obvious signs of intergranular cracking. This can perhaps be expected, since the fracture path in the transverse plane is normal to the grain boundaries, whilst intergranular fracture is more likely to occur in the longitudinal plane. Preferential precipitation at the grain boundaries and precipitate free zones are instrumental in reducing the grain boundary cohesive strength, thus promoting intergranular fracture. The presence of aligned particles or stringers in the longitudinal direction also reduces the initiation fracture properties by increasing the amount of particle matrix decohesion ahead of the crack tip.

In order to establish the fracture mode in the high Ln Zj and low Ln Zj extrudates, 205

the fracture surfaces were examined in the centre of the specimen. The scanning electron micrographs are shown in fig. 8.22 for the extrudates in all 3 tempers at high and low Ln Zj. Micrographs 8.22 a,b show that in the T6 temper, the fracture surfaces have a characteristic dimpled appearance associated with a ductile or fibrous mode of fracture. Micrograph 8.22b shows that for a fully recrystallised, high Ln Z\ extrudate the large dimples have an average spacing of 18 tim and are associated with inclusions of 3-4 p.m in diameter, similar in size to the as cast particles shown earlier. The spacing of the smaller dimples, 2-3 urn, is similar to the separation of the intermediate sized particles shown in fig. 7.23e. The observations are consistent with the fracture mode reported by several workers for commercial age hardenable aluminium alloys. Cracking is initiated within the plastic zone ahead of the crack tip by decohesion of the particle matrix interface at the large inclusions; the voids grow to a critical size until the strain is sufficient to cause decohesion at the intermediate particles so that smaller secondary dimples are formed in the ligaments between the larger voids. The fracture surfaces from the low

Ln Zj extrudates shown in the micrographs indicate that a similar fracture mechanism occurs in the presence of a substructure, although it is noticeable that the larger dimples are larger and fewer in number. Assuming that the initiation stage of fracture in the high Ln Zj and low Ln Zj extrudates involves particle matrix decohesion at the larger inclusions, the increased void size indicates that void growth continues to a higher strain, and hence stress, before decohesion of the ligament between the void occurs. The substructure strengthening may effectively be increasing the fracture toughness by increasing the energy required for decohesion of the ligaments. The unstable crack growth in the high Ln Zj extrudates shown in fig.

8.22b is therefore due to rapid decohesion of the inter-void ligaments associated with a recrystallised matrix. Therefore, although the decrease in strain hardening coefficient shown in table 8.3 may result in an increase in strain and stress localisation at the crack tip and hence increase in the particle matrix decohesion, the increase in energy required to cause decohesion in the substructure strengthened matrix results in an equivalent or greater fracture toughness when compared to the high Ln Zj extrudates. Since the subgrain structure is relatively equiaxed, it is likely that the fracture toughness in the transverse longitudinal direction will also be enhanced. Fig. 8.22 Fractured surfaces of extrudates a) T6-temper, Ln Z\ = 25.59 b) T6-temper, Ln Z\ = 29.52 c) T5-temper, Ln Z\ = 25.59 d) T5-temper, Ln Z\ = 29.52 e) T1-temper, Ln Zj = 25.59 f) T1 -temper, Ln Z\ = 29.52

207

The fracture toughnesses of T5 and T1 extrudates are also shown in figs. 8.20 & 8.21. There is no appreciable difference in the fracture toughness values between T5 and T 1 tempers. The fracture toughness values, which follow the same trend with Ln

Zj as the fracture toughness values of the T6 temper material, are in fact considerably lower under all conditions. The fracture surface of the T1 and T5 extrudates are shown in fig. 8.22 c-f. The fracture surfaces show larger voids compared to the T6 micrographs. This can be rationalised on the basis of the microstructure of the extrudates. In addition to the fine ageing precipitates, T5 and T1 microstructures contain a larger number of coarser aluminium-copper-magnesium particles which are remnants of the homogenisation and furnace cooling treatment. The dimpled nature of the fractured surface indicates that nucleation and growth of the voids occurs in front of the advancing crack. Since it is easier to nucleate voids in a material that contains a larger volume fraction of coarser particles than a material with the same volume fraction of fine particles, the fracture initiation toughness will be lower in a material containing coarse particles.

In the present case, the T1, T5 and T6 extrudates conformed to this microstructural variation, hence the large difference in the fracture toughness between T5 and T6 under all conditions. The size and separation of the dimples in the micrographs indicate that ageing precipitates play no rdle in the fracture mechanism, but influence the fracture toughness by changing the matrix deformation characteristics.

8.5 Limit diagrams

Extrusion limit diagrams are a practical way of expressing the operating limits for a given press configuration and alloy. The maximum press capacity available for extrusion in the present case is 1130 MPa on a 3 inch billet. The limit lines for round extrusion for alloy 2024 are shown in fig 8.23 for the direct and indirect extrusion mode. In this instance, the ram speed was 6.9 mm/s. The diagrams contain two limits, one on the right hand side and one on the left. The limit line on the left hand side is the limiting pressure, while that on the right is the surface quality line. Aluminium, its alloys and, in the present instance, 2024 has a decreasing flow stress with increasing temperature; thus it is possible to extrude at high ratios with increasing extrusion temperature. However, at higher temperatures, if the extrusion ratio is increased above a certain limit, surface deterioration due to 208

surface cracking commences. As shown in the limit diagram for indirect extrusion in fig 8.23, the limiting pressure line is much higher than for direct extrusion; this variation can be accounted for, in that the contribution to the pressure due to the friction is much lower in indirect extrusion, compared to direct extrusion. This allows the material to be extruded at much higher ratios for a given temperature in indirect extrusion. The surface cracking line is approximately constant for both direct and indirect extrusion. This is not surprising since the mechanism for surface cracking is not much different in either mode.

Aluminium and its alloys have comparatively high strain rate sensitivity, hence an increase in ram speed lowers the pressure limit line, due to the increase in flow stress and hence load, although ram speed has a more pronounced effect on the surface cracking limits which may be attributed to the higher temperature rise and stresses generated at the die exit. The expected direction of shift in the limiting pressure line and surface cracking lines with increase in ram speed is also shown in

fig. 8.23.

Sheppard^2 has shown that limit lines may also be constructed to represent

specific structures and properties, which are a function of the process conditions. The results in earlier sections have shown that for heat treatable alloys such as 2024 the situation is not so simple, since the relationships between the process conditions and structures and properties are also dependent on the final temper. However, the most important structural feature in relation to the as extruded material is the amount of peripheral recrystallisation and the scale of substructure. It was shown in section 7.2.1 that the percentage recrystallisation could be related to the extrusion conditions. The line obtained from table 7.1 represents a relationship between peripheral recrystallisation and initial billet temperature at a fixed ratio and ram speed. Using equations developed in this way, the limiting lines representing 2.5 and 5.0% recrystallisation in the T1 condition have been shown in fig. 8.23 The limiting lines show that as the percentage recrystallisation increases with ram speed, the extrusion range is reduced, which is accentuated by the larger temperature rises during direct extrusion. It should be noted, as discussed earlier, that the arrangement of the indirect tooling almost doubles the quench delay, during which static recovery and recrystallisation take place, so the limiting lines for indirect extrusion may actually be larger than are shown in fig. 8.23. 209

FigB*23Limits for extruding alloy 2024 round section 210

From the relationship between Ln Z and d‘ 1, the reciprocal of the subgrain size, it has proved possible to plot subgrain size on the limit diagram, as included on fig. 8.23. However, since the overriding strengthening mechanism for alloys such as 7075 and 2024 is precipitation hardening and not substructural strengthening the practical use of such lines is limited.

The limit diagrams for different shapes in direct and indirect mode are shown in fig. 8.24 and 8.25 respectively. The type of limit diagram shown here is different from the conventional limit diagram, in that the ram speed instead of extrusion ratio is plotted along the y-axis, while the x-axis remains the initial temperature. This is because the die assembly limited the minimum extrusion ratio that could be obtained in the available press configuration. It can be seen that the window for acceptable extrusions is expanded in the indirect mode compared to the direct mode, for simple symmetric sections such as round and square sections, while the difference is only marginal for the complex sections like T and U sections. Also evident from the diagrams is, that with increase in complexity of section geometry, the extrusion

limit field is restricted more in the direct mode compared to the indirect mode.

In conclusion, the indirect mode clearly offers the greatest range of extrusion conditions over a range of processing conditions, due to lower loads at low extrusion temperatures, while at high temperatures benefit is obtained due to the more uniform nature of flow and thus from the smaller temperature rises reducing the likelihood of surface cracking. It should be noted that the above results are relevant to a specific set of preprocessing conditions such as homogenisation, preheating etc. Although it is expected that the extrusion limit windows are likely to be altered, depending on the preprocessing, the exact nature of this change has to be further investigated. 211

Temperature’c Fig8*24 Limit diagram for direct extrusion of shapes 10

250 500 350 400 450 500 Temperature *c Fig 8 25 Limit diagram for indirecf extrusion of shapes 213

8.6 Conclusions

1. Hardness testing of the extrudate generally shows a similar trend to that found in tensile testing and can be related to the initial extrusion conditions.

2. The tensile properties of 2024 alloy in the as extruded condition is related to the substructure strengthening but the influence of solid solution strengthening is far greater at high extrusion temperatures.

3. The tensile properties of the solutionised and aged extrudates are superior to those that are press quenched and aged. In the solutionised and aged extrudates, the effect of the extrusion temperature on the strength is minimal. In the press quenched and aged extrudates however the strength increases with increasing extrusion temperature.

4. The fracture toughness of the heat treated 2024 extrudates increases with increasing extrusion temperature due to the tendency for retention of substructure. This results in a change in fracture mode from an intergranular mode in the recrystallised structure to a transgranular mode in the recovered structure.

5. Basic limit diagrams showing the limiting conditions of pressure and surface finish have been produced for 2024 alloy. These show that the range of extrusion conditions for this alloy is expanded by the use of the indirect process. 9.0 Conclusions 215

1. The value of the temperature compensated strain rate, Z, must be corrected when considering the die geometry. The load encountered during extrusion is a function of the load required for rod extrusion and the correction factor. The necessity for this correction factor arises from inaccuracies in calculating the strain rate and deformation during shaped extrusion. An additional consideration is the use of mean strain rates and temperatures calculated from rod extrusion in the calculation of the Zenner-Holloman parameter. The correction factors can be related to the geometric variables but further work is required in this area. As the complexity of section shape is increased, the magnitude of the excess load required to initiate extrusion increases.

2. The extruded rod exhibited a fibrous core at all extrusion temperatures with a strained periphery or recrystallised grains at the periphery. The extent of this recrystailisation increases with extrusion ratio, initial billet temperature and ram speed. In shaped extrusion the extent of recrystailisation was found to increase with increase in complexity of shape. Also the recrystallised layer was deeper where there was a reentrant corner in the section due to the increased shear occurring in these regions.

3. Investigation of the structural development of alloy AA2024 shows that dynamic recovery is the operative high temperature restoration mechanism. Increased temperature and decreased strain rate increases the size and perfection of the subgrains. The presence of a large amount of precipitate in 2024 alloy pins the sub-boundaries and thus the subgrain size is smaller. The variation in substructure within the steady state direct and indirect deformation zone results in a more uniform substructure across and along the indirect extrudate. The structures observed in the torsion samples have in general been consistent with those found in extrusion.The subgrain size'd' may be related to the process conditions by: d'1 = a.LnZc +b

4. The structure in the extrudate after solutionising and quenching is dependent upon the initial extrusion temperature. With low extrusion temperatures, complete recrystailisation occurs during solutionising, but with high extrusion temperatures the substructure is retained, the proportion of which increases with increasing 216

extrusion temperature. The solution treated extrudates all showed varying degrees of recrystallisation; those extruded below 350°C were fully recrystallised. Indirect extrusion reduces the volume percent recrystallised. The increase in recrystallised grain size with extrusion temperature can also be directly related to the increase in subgrain size. The effect of ageing the solution treated material is to produce GP zones and transition precipitates, hence improving the strength.

5. The effect of artificial ageing on the solutionised extrudates is to produce transition precipitates. In the low Z partially recrystallised extrudates the presence of a retained substructure provides sites for preferential precipitation. Artificial ageing of the high extrusion temperature press quenched extrudates results in a pronounced ageing reaction.

6. The tensile properties of 2024 alloy in the as extruded condition is related to the substructure strengthening but the influence of solid solution strengthening is far greater at high extrusion temperatures.

7. The tensile properties of the solutionised and aged extrudates are superior to those that are press quenched and aged. In the solutionised and aged extrudates, the effect of the extrusion temperature on the strength is minimal. In the press quenched and aged extrudates however the strength increases with increasing extrusion temperature.

8. The fracture toughness of the heat treated 2024 extrudates increases with increasing extrusion temperature. This trend can be attributed to the difference in the distribution and density of ageing precipitates in the extrudates. In the low Zj conditions i.e. at high extrusion temperatures, the greater density of these precipitates accounts for the increased density of ligaments observed at fracture. 217

Recommendations for further work

During the course of the present investigation, the products studied were of relatively simple section, i.e. low peripheral ratio. The analytical solutions used in extrusion analysis were developed for rod extrusion and were, in the present work, also applied to shaped extrusion, with the application of appropriate correction factors. Obviously, a change in section affects the metal flow, strain rate, stress and temperature distribution in the deformation zone; the effect of the section change on these parameters will not be reflected in the results of the calculations. Although this does not greatly affect the results for simple sections, since mean values for the parameters listed are used in the theory in any case, increasing complexity of section, and therefore flow, may lead to discrepancies. Thus it may be recommended that the flow of material during complex shape extrusion may be studied; no detailed experimental work has yet been independently reported on the extrusion of high peripheral ratio shapes. A useful future study might include the investigation of gridded billets placed at various positions relative to the die orifice. Furthermore, in the present work, only one extrusion ratio was studied in detail for the indirect % mode, due to structural limitations on the tooling design; evidently, further work should include an investigation of a wider range of extrusion ratios for the shapes considered.

Material flow studies are one aspect of extrusion characterisation which have received little attention, especially in the extrusion of complex shapes. In the material flow studies made in this investigation, super purity aluminium was used for ease of macroetching. However, its flow characteristics are different to those of AA2024, and thus a flow study using this alloy is to be advised. Problems with the examination of the flow will have to be overcome if such an investigation were to be successful. The introduction of a series of radial pins of a different alloy which possesses similar deformation properties, but is more susceptible to macroetching may provide one solution to this difficulty, the introduction of a radioactive tracer element another.

One of the principle observations made in the present work is that press quenching can be effectively used in industry to improve the energy economics of the products by eliminating an extra solutionising treatment after extrusion. However, in the present work the observations made indicate that strength developed in the 218

extrudate produced by this route is far inferior to that developed after traditional heat treatment. The thermal history of the former extrudate does not guarantee complete solutionising of the solute elements prior to heat treatment. This is primarily because the material was homogenised and then furnace cooled and then induction heated in a very short time prior to extrusion. In future work it would be interesting to see if preheating for a considerably longer time prior to extrusion helps to improve the strength in the T5 or T1 condition. Appendix Table of extrusion data

Extrusion Extrusion Extrusion Direct/ Extrusion Ram Strain Ln (Zi) Ln (Zi/A) Peak load Ctxte Shape Ratio indirect Temperature Speed Rate Load /Kelvin) (mm/s) / s * * - 1 ) /Tons)

JS450R7D/20 Round 20:01 Direct 723 7 3.188 25.59 5.99 258 JS400R7D/20 Round 20:01 Direct 673 7 3.188 27.41 7.81 301 JS350R7D/20 Round 20:01 Direct 623 7 3.188 29.52 9.92 352 JS300R7D/20 Round 20:01 Direct 573 7 3.188 31.99 12.39 413

JS450R7D/30 Round 30:1 Direct 723 7 6.725 26.34 6.74 280 JS400R7D/30 Round 30:1 Direct 673 7 6.725 28.16 8.56 324 JS350R7D/30 Round 30:1 Direct 623 7 6.725 30.26 10.66 376 JS300R7D/30 Round 30:1 Direct 573 7 6.725 32.74 13.14 438

JS450R7D/50 Round 50:1 Direct 723 7 11.178 26.85 7.25 301 JS400R7D/50 Round 50:1 Direct 673 7 11.178 28.66 9.06 343 JS350R7D/50 Round . 50:1 Direct 623 7 11.178 30.77 11.17 395 JS300R7D/50 Round 50:1 Direct 573 7 11.178 33.24 13.64 457

JS450R3D/20 Round 20:01 Direct 723 3 0.719 24.10 4.50 222 JS400R3D/20 Round 20:01 Direct 673 3 0.719 25.92 6.32 264 JS350R3D/20 Round 20:01 Direct 623 3 0.719 28.03 8.43 316 JS300R3D/20 Round 20:01 Direct 573 3 0.719 30.50 10.90 378

JS450R3D/30 Round 30:1 Direct 723 3 1.460 24.81 5.21 244 JS400R3D/30 Round 30:1 Direct 673 3 1.460 26.63 7.03 287 JS350R3D/30 Round 30:1 Direct 623 3 1.460 28.73 9.13 338 JS300R3D/30 Round 30:1 Direct 573 3 1.460 31.21 11.61 401

JS450R3D/50 Round 50:1 Direct 723 3 5.913 26.21 6.61 285 JS400R3D/50 Round 50:1 Direct 673 3 5.913 28.03 8.43 328 JS350R3D/50 Round 50:1 Direct 623 3 5.913 30.13 10.53 379 JS300R3D/50 Round 50:1 Direct 573 3 5.913 32.61 13.01 442 to o

Appendix 1. Extrusion data for alloy 2024 Table of extrusion data

Extrusion Peak APeak AT Ln (Zc) Steady State Rexn Laver Rexn Layer Volume fraction Cods Pressure Pressure Pressure Thickness Thickness Rextd after (MPa) (MPa) (MPa) (as extruded) (solutionised) solutionisinq (nm ) (um )

JS450R7D/20 617 39.70 64.19 23.40 519.51 JS400R7D/20 718.4 54.80 80.09 24.34 597.17 JS350R7D/20 841.9 80.00 99.46 25.24 685.63 JS300R7D/20 986.3 104.30 122.09 26.08 793.71

JS450R7D/30 669 47.50 72.34 23.89 559.29 JS400R7D/30 775 62.70 88.97 24.79 641.00 JS350R7D/30 898 87.70 108.25 25.67 729.19 JS300R7D/30 1048 111.80 131.77 26.46 842.49

JS450R7D/50 719 53.00 80.18 24.16 599.33 JS400R7D/50 820 67.90 96.02 25.07 676.81 JS350R7D/50 944 93.00 115.46 25.92 765.81 JS300R7D/50 1093 119.00 138.82 26.70 876.50

JS450R3D/20 530 24.00 50.55 22.34 455.35 JS400R3D/20 631 39.50 66.39 23.32 532.29 JS350R3D/20 755 64.50 85.83 24.26 621.38 JS300R3D/20 904 88.30 109.19 25.11 734.05

JS450R3D/30 584 31.50 59.02 22.78 497.19 JS400R3D/30 685.5 47.00 74.93 23.74 574.59 JS350R3D/30 808 71.70 94.14 24.65 662.60 JS300R3D/30 959 95.70 117.81 25.47 776.88

JS450R3D/50 682 46.40 74.38 23.70 571.98 JS400R3D/50 783.5 61.30 90.30 24.62 649.91 JS350R3D/50 906.5 86.40 109.58 25.49 738.01 JS300R3D/50 1057 110.50 133.18 26.27 851.76

Appendix 1. Extrusion data for alloy 2024 Table of extrusion data

Extrusion Extrusion Extrusion Direct/ Extrusion Ram Strain Ln (Zi) Ln tZi/A) Peak load Coda Shape Ratio Indirect Temperature Speed Rate Load (Kelvin) (m m/s) ( s * * - 1 ) (Tons)

JS450R7I/20 Round 20:01 Indirect 723 6.9 4.182 25.87 6.27 213 JS400R7I/20 Round 20:01 Indirect 673 6.9 4.182 27.68 8.08 253 JS350R7I/20 Round 20:01 Indirect 623 6.9 4.182 29.79 10.19 302 JS300R7I/20 Round 20:01 Indirect 573 6.9 4.182 32.26 12.66 362

JS450R7I/30 Round 30:1 Indirect 723 6.9 9.006 26.63 7.03 235 JS400R7I/30 Round 30:1 Indirect 673 6.9 9.006 28.45 8.85 276 JS350R7I/30 Round 30:1 Indirect 623 6.9 9.006 30.55 10.95 325 JS300R7I/30 Round 30:1 Indirect 573 6.9 9.006 33.03 13.43 385

JS450R7I/50 Round 50:1 Indirect 723 6.9 15.085 27.15 7.55 253 JS400R7I/50 Round 50:1 Indirect 673 6.9 15.085 28.96 9.36 286 JS350R7I/50 Round 50:1 Indirect 623 6.9 15.085 31.07 11.47 342 JS300R7I/50 Round 50:1 Indirect 573 6.9 15.085 33.54 13.94 402

JS450R3I/20 Round 20:01 Indirect 723 3 0.942 24.37 4.77 178 JS400R3I/20 Round 20:01 Indirect 673 3 0.942 26.19 6.59 218 JS350R3I/20 Round 20:01 Indirect 623 3 0.942 28.30 8.70 267 JS300R3I/20 Round 20:01 Indirect 573 3 0.942 30.77 11.17 327

JS450R3I/30 Round 30:1 Indirect 723 3 2.064 25.16 5.56 201 JS400R3I/30 Round 30:1 Indirect 673 3 2.064 26.97 7.37 241 JS350R3I/30 Round 30:1 Indirect 623 3 2.064 29.08 9.48 290 JS300R3I/30 Round 30:1 Indirect 573 3 2.064 30.83 11.23 333

JS450R3I/50 Round 50:1 Indirect 723 3 8.143 26.53 6.93 238 JS400R3I/50 Round 50:1 Indirect 673 3 8.143 28.35 8.75 279 JS350R3I/50 Round 50:1 Indirect 623 3 8.143 30.45 10.85 328

JS300R3I/50 Round 50:1 Indirect 573 3 8.143 32.93 13.33 387 222

Appendix 1. Extrusion data for alloy 2024 Table of extrusion data

Extrusion Peak APeak AT Ln (Zc) Steady State Rexn Layer Rexn Layer Volume fraction Code Pressure Pressure Pressure Thickness Thickness Rextd after (MPa) (MPa) (MPa) (as extruded) (solutionised) soiutionisina (um ) (um )

JS450R7I/20 509 36.10 47.26 24.21 425.56 JS400R7I/20 605 53.20 62.31 25.23 496.56 JS350R7I/20 722 73.00 80.66 26.22 584.04 JS300R7I/20 865 96.30 103.08 27.12 691.75

JS450R7I/30 562 43.25 55.57 24.71 466.82 JS400R7I/30 659 60.50 70.78 25.70 538.59 JS350R7I/30 776 80.20 89.12 26.66 626.15 JS300R7I/30 919 103.50 111.54 27.54 733.87

JS450R7I/50 605 48.50 62.31 25.01 500.79 JS400R7I/50 684 65.20 74.70 26.08 556.86 JS350R7I/50 818 87.10 95.71 26.93 657.74 JS300R7I/50 960 106.40 117.97 27.80 768.15

JS450R3I/20 426 23.00 34.24 23.15 362.66 JS400R3I/20 522 38.10 49.30 24.21 435.46 JS350R3I/20 640 57.20 67.80 25.23 524.46 JS300R3I/20 782 80.30 90.06 26.18 631.46

JS450R3I/30 481 30.40 42.87 23.65 405.49 JS400R3I/30 577 45.20 57.92 24.68 478.57 JS350R3I/30 694 68.30 76.26 25.68 563.07 JS300R3I/30 796 80.90 92.26 26.87 643.52

JS450R3I/50 570 43.30 56.83 24.57 473.98 JS400R3I/50 666 58.50 71.88 25.56 546.69 JS350R3I/50 783 81.20 90.22 26.52 631.55 JS300R3I/50 926 100.60 112.64 27.40 742.78 u>to

Appendix 1. Extrusion data for alloy 2024 Table of extrusion data

Extrusion Extrusion Extrusion Direct/ Extrusion Ram Strain Ln (Zi) Ln (Zi/A) Peak load Cbde Shape Ratio Indirect Temperature Speed Rate Lead (Kelvin) (m m/s) ( s * * - 1 ) (Tons)

JS450O3D Round 40:1 Direct 723 3 3.967 25.81 6.21 303 JS390O5D Round 40:1 Direct 663 5 6.667 28.54 8.94 394 JS37506D Round 40:1 Direct 648 6 8.017 29.34 9.74 400 JS34506D Round 40:1 Direct 618 6 8.017 30.67 11.07 429 JS360O7D Round 40:1 Direct 633 7 9.364 30.15 10.55 423 JS350O3D Round 40:1 Direct 623 3 3.967 29.73 10.13 403 JS300R3D Round 40:1 Direct 573 3 3.967 32.20 12.60 472 JS325R5D Round 40:1 Direct 598 5 6.667 31.44 11.84 455 JS330R6D Round 40:1 Direct 603 6 8.017 31.38 11.78 452 JS380R5D Round 40:1 Direct 653 5 6.667 28.95 9.35 390

JS425U3D U-shape 40:1 Direct 698 3 3.967 26.69 7.09 356 JS380U5D U-shape 40:1 Direct 653 5 6.667 28.95 9.35 427 JS350U5D U-shape 40:1 Direct 623 5 6.667 30.25 10.65 471 JS360U6D U-shape 40:1 Direct 633 6 8.017 29.99 10.39 442 JS350U3D U-shape 40:1 Direct 623 3 3.967 29.73 10.13 453 JS450U3D U-shape 40:1 Direct 723 3 3.967 25.8 1 6.21 336 JS315U3D U-shape 40:1 Direct 588 3 3.967 31.42 11.82 ex JS315U4D U-shape 40:1 Direct 588 4 5.317 31.72 12.12 ex JS400U3D U-shape 40:1 Direct 673 3 3.967 27.63 8.03 414 JS340U5D U-shape 40:1 Direct 613 5 6.667 30.72 11.12 ex

JS330S4D Sauare 40:1 Direct 603 4 5.317 30.97 11.37 450 JS350S6D Square 40:1 Direct 623 6 8.017 30.44 10.84 446 JS375S6D Sauare 40:1 Direct 648 6 8.017 29.34 9.74 421 JS425S3D Square 40:1 Direct 698 3 3.967 26.69 7.09 338 JS390S5D Square 40:1 Direct 663 5 6.667 28.54 8.94 391 JS370S5D Square 40:1 Direct 643 5 6.667 29.37 9.77 420 JS365S7D Square 40:1 Direct 638 7 9.367 29.93 10.33 426 to

Appendix 1. Extrusion data for alloy 2024 Table of extrusion data

Extrusion Peak APeak AT Ln (Zc) Steady State Rexn Laver Rexn Laver Volume fraction Gxfe Pressure Pressure Pressure Thickness Thickness Rextd after (MPa) (MPa) (MPa) (as extruded) (solutionised) solutionising (um ) (um )

JS450O3D 723.00 80.81 23.11 340 360 11.5 JS390O5D 942.00 121.00 115.15 24.23 755.32 520 880 27.22 JS37506D 958.00 150.00 117.66 24.76 743.36 650 1120 33.96 JS34506D 1027.00 200.00 128.48 25.30 760.84 750 1520 44.45 JS360O7D 1012.00 182.00 126.12 25.08 763.60 800 1300 38.79 JS350O3D 965.00 166.00 118.75 24.77 735.08 720 1200 36.13 JS300R3D 1130.00 250.00 144.62 25.45 809.60 --- JS325R5D 1088.00 221.00 138.04 25.40 797.64 --- JS330R6D 1082.00 217.00 137.10 25.46 795.80 - -- JS380R5D 934.00 136.00 113.89 24.55 734.16 450 950 25.65

JS425U3D 853.00 88.00 101.19 23.17 703.80 550 900 49.66 JS380U5D 1021.00 184.00 127.53 24.12 770.04 880 1420 78.36 JS350U5D 1126.00 232.00 144.00 24.45 822.48 1050 1780 98.22 JS360U6D 1057.00 216.00 133.18 24.70 773.72 1010 1740 96.01 JS350U3D 1084.00 205.00 137.41 24.15 808.68 925 1620 89.4 JS450U3D 805.00 60.00 93.67 22.73 685.40 500 700 38.6 JS315U3D ex ex --- JS315U4D ex ex - - - JS400U3D 990.00 128.00 122.67 23.20 793.04 600 1000 60.11 JS340U5D ex ex 1000 1720 94.23

JS330S4D 1077.00 223.00 136.31 25.08 785.68 805 1520 57.76 JS350S6D 1066.00 206.00 134.59 24.94 791.20 850 1600 60.8 JS375S6D 1008.00 156.00 125.50 24.51 783.84 740 1220 46.36 JS425S3D 808.00 64.00 94.14 23.39 684.48 420 560 21.28 JS390S5D 936.00 128.00 114.21 24.26 743.36 590 1060 40.28 JS370S5D 1006.00 161.00 125.18 24.48 777.40 760 1220 46.36 JS365S7D 1020.00 182.00 127.38 24.89 770.96 810 1440 54.72 t o Lh

Appendix 1. Extrusion data for alloy 2024 Table of extrusion data

Extrusion Extrusion Extrusion Direct/ Extrusion Ram Strain Ln (Zi) Ln (Zi/A) Peak load Cods Shape Ratio Indirect Temperature Speed Rate Load (Kelvin) (mm/s) (s**-1) (Tons)

JS300S3D Square 40:1 Direct 573 3 3.967 32.20 12.60 ex JS400S4D Square 40:1 Direct 673 4 5.317 27.92 8.32 375 JS350S5D Square 40:1 Direct 623 5 6.667 30.25 10.65 438

JS400T3D T-shape 40:1 Direct 673 3 3.967 27.63 8.03 372 JS350T3D T-shape 40:1 Direct 623 3 3.967 29.74 10.14 431 JS350T7D T-shape 40:1 Direct 623 7 9.367 30.59 10.99 458 JS400T7D T-shape 40:1 Direct 673 7 9.367 28.49 8.89 387 JS450T3D T-shape 40:1 Direct 723 3 3.967 25.81 6.21 326 JS450T7D T-shape 40:1 Direct 723 7 9.367 26.67 7.07 350 JS450T5D T-shape 40:1 Direct 723 5 6.667 26.33 6.73 340

JS38505I Round 40:1 indirect 658 5 9.048 29.05 9.45 287 JS36007I Round 40:1 Indirect 633 7 12.608 30.44 10.84 327 JS32507I Round 40:1 Indirect 598 7 12.608 32.08 12.48 386 JS34009I Round 40:1 Indirect 613 9 16.168 31.60 12.00 378 JS31008I Round 40:1 Indirect 583 8 14.388 32.9,' 13.37 422 JS42503I Round 40:1 Indirect 698 3 5.488 27.01 7.41 257 JS40005I Round 40:1 Indirect 673 3 9.048 28.45 8.85 282 JS300O7I Round 40:1 Indirect 573 7 12.608 33.37 13.77 421 JS32508I Round 40:1 Indirect 598 8 14.388 32.21 12.61 386 JS42504I Round 40:1 Indirect 698 4 7.268 27.30 7.70 248

JS350U7I U-Shape 40:1 Indirect 623 7 12.608 30.89 11.29 397 JS400U5I U-Shape 40:1 Indirect 673 5 9.048 28.45 8.85 321 JS415U3I U-Shape 40:1 Indirect 688 3 5.488 27.38 7.78 280 JS375U6I U-Shape 40:1 Indirect 648 6 10.828 29.64 10.04 359 JS360U6I U-Shape 40:1 Indirect 633 6 10.828 30.29 10.69 382 JS350U5I U-Shape 40:1 Indirect 623 5 9.048 30.58 10.96 371

Appendix 1. Extrusion data for alloy 2024 Table of extrusion data

Extrusion Peak APeak AT Ln (Zc) Steady State Rexn Layer Rexn Laver Volume fraction Cocfe Pressure Pressure Pressure Thickness Thickness Rextd after (MPa) (MPa) (MPa) (as extruded) (solutionised) solutionisina (nm) (um )

JS300S3D ex ex - -- JS400S4D 897.00 109.00 108.09 23.94 724.96 500 950 30.25 JS350S5D 1048.00 195.00 131.77 24.85 784.76 820 1500 54.66

JS400T3D 891.00 110.00 107.15 23.68 718.52 550 960 42.29 JS350T3D 1032.00 183.00 129.26 24.42 781.08 900 1480 65.19 JS350T7D 1096.00 212.00 139.29 24.95 813.28 1000 1700 74.89 JS400T7D 926.00 143.00 112.64 24.37 720.36 700 1160 51.1 JS450T3D 780.00 41.00 89.75 22.85 679.88 415 520 22.91 JS450T7D 837.00 70.00 98.69 23.45 705.64 480 760 33.48 JS450T5D 814.00 62.00 95.08 23.21 691.84 420 660 29.07

JS38505I 685 78.10 74.85 26.03 558.35 200 840 26.05 JS36007I 781 130.00 89.91 26.64 598.92 340 1100 33.67 JS32507I 922 156.00 112.02 26.98 704.72 140 1020 31.19 JS34009I 903 164.00 109.03 26.84 679.88 200 1040 31.75 JS31008I 1008 226.40 125.50 27.11 719.07 100 1000 30.63 JS42503I 614 39.00 63.72 24.68 529.00 60 ' 520 16.51 JS40005I 674 85.90 73.13 25.62 541.05 160 740 23.14 JS30007I 1010 200.00 125.81 27.31 745.20 JS32508I 925 166.00 112.48 27.10 698.28 JS42504I 596 31.00 60.90 25.05 519.80

JS350U7I 950.3 176.00 116.45 26.01 712.36 400 1260 69.52 JS400U5I 767 124.90 87.71 25.13 590.73 340 1100 60.69 JS415U3I 665.6 110.10 71.81 24.71 511.06 240 880 48.55 JS375U6I 857.7 156.20 101.93 25.58 645.38 540 1200 66.21 JS360U6I 912.7 179.60 110.55 25.75 674.45 600 1280 70.63 JS350U5I 886.4 171.70 106.43 26.03 657.52 700 1320 72.83 N> "j

Appendix 1. Extrusion data for alloy 2024 Table of extrusion data

Extrusion Extrusion Extrusion Direct/ Extrusion Ram Strain Ln (Zi) Ln (Zi/A) Peak load Code Shape Ratio Indirect Temperature Speed Rate Load (Kelvin) (mm/s) ( s * * - 1 ) (Tons)

JS335U3I U-Shape 40:1 Indirect 608 3 5.488 30.76 11.16 392 JS315U3I U-Shape 40:1 Indirect 588 3 5.488 31.75 12.15 408 JS290U3I U-Shape 40:1 Indirect 563 3 5.488 33.08 13.48 457 J*380U5I U-Shape 40:1 Indirect 653 3 5.488 28.76 9.16 306 J*365U7I U-Shape 40:1 Indirect 638 7 12.608 30.22 10.62 350 J*390U4I U-Shape 40:1 Indirect 663 4 7.268 28.63 9.03 300

JS360S7I Sauare 40:1 l Indirect 633 7 12.608 30.44 10.84 355 JS410S3I Square 40:1 Indirect 683 3 5.488 27.57 7.97 276 JS420S3I Square 40:1 Indirect 693 3 5.488 27.19 7.59 260 JS390S5I Square 40:1 Indirect 663 5 9.048 28.85 9.25 299 JS345S5I Square 40:1 Indirect 618 5 9.048 30.79 11.19 363 JS310S4I Square 40:1 Indirect 583 4 7.268 32.29 12.69 410 JS295S3I Square 40:1 Indirect 568 3 5.488 32.80 13.20 413 JS335S6I Square 40:1 Indirect 608 6 10.828 33.09 13.49 400 JS365S5I Square 40:1 Indirect 638 5 9.048 31.44 11.84 345 JS350S7I Square 40:1 Indirect 623 7 12.608 29.89 10.29 375 JS295S4I Square 40:1 Indirect 568 4 7.268 30.89 11.29 424 J*375S6I Square 40:1 Indirect 648 6 10.828 29.64 10.04 352

JS475T3I T-Shape 40:1 Indirect 748 3 5.488 25.32 5.72 224 JS460T3I T-Shape 40:1 Indirect 733 3 5.488 25.80 6.20 234 JS450T3I T-Shape 40:1 Indirect 723 3 5.488 26.14 6.54 244 JS350T3I T-Shape 40:1 Indirect 623 3 5.488 30.06 10.46 345 JS335T3I T-Shape 40:1 Indirect 608 3 5.488 30.78 11.16 376 JS335T8I T-Shape 40:1 Indirect 608 3 14.388 31.72 12.12 394 JS360T8I T-Shape 40:1 Indirect 633 8 14.388 30.57 10.97 359 JS400T5I T-Shape 40:1 Indirect 673 5 9.048 28.45 8.85 29 JS415T3I T-Shape 40:1 Indirect 688 3 5.488 27.38 7.78 268 to ooto

Appendix 1. Extrusion data for alloy 2024 Table of extrusion data

Extrusion Peak APeak AT Ln (Zc) Steady State Rexn Laver Rexn Laver Volume fraction Cbcte Pressure Pressure Pressure Thickness Thickness Rextd after (MPa) (MPa) (MPa) (as extruded) (solutionised) solutionisinq

JS335U3I 935 169.50 114.05 25.74 704.26 460 1300 71.73 JS315U3I 974.8 187.40 120.30 26.18 724.41 240 1060 64 JS290U3I 1092 214.00 138.66 26.33 807.76 60 1200 66.21 J*380U5I 730 131.00 81.91 25.44 551.08 J*365U7I 840 166.00 99.16 26.14 620.08 J*390U4I 720 133.00 80.34 25.46 540.04

JS360S7I 858.2 150.00 102.01 26.20 651.54 400 1140 38 JS410S3I 669.4 53.00 72.41 24.84 567.09 640 23 JS420S3I 636 40.00 67.17 24.71 548.32 160 580 20.74 JS390S5I 724.4 93.70 81.04 25.66 580.24 200 880 30.5 JS345S5I 877.3 164.00 105.00 26.24 656.24 320 1100 37.3 JS310S4I 990 210.80 122.68 26.54 716.86 100 1060 36 JS295S3I 997 171.80 123.77 26.74 759.18 60 1040 35.48 JS335S6I 1023 242.00 127.85 25.93 718.52 1020 34.87 JS365S5I 966 210.80 118.91 25.14 694.78 240 1080 36.67 JS350S7I 834 131.20 98.22 26.66 646.58 300 . 1100 37.29 JS295S4I 906 164.00 109.51 27.59 682.64 300 1 1100 37.29 J*375S6I 845 122.00 99.94 25.65 665.16

JS475T3I 535.2 33.00 51.37 23.64 462.02 560 24.67 JS460T3I 559.1 44.00 55.12 23.94 473.89 200 580 25.55 JS450T3I 583 52.00 58.87 24.11 488.52 200 640 28.19 JS350T3I 825 141.00 96.80 25.88 629.28 520 1200 52.86 JS335T3I 898 151.00 108.25 25.96 687.24 380 1200 52.86 JS335T8I 941.4 177.00 115.05 26.67 703.25 240 1120 49.34 JS360T8I 857.7 154.30 101.93 26.34 647.13 600 1040 54.62 JS400T5I 702 104.00 77.52 25.47 550.16 360 920 40.53 JS415T3I 640.3 78.90 67.85 24.84 516.49 240 700 30.83 to VO

Appendix 1. Extrusion data for alloy 2024 Table of extrusion data

Extrusion Extrusion Extrusion Direct/ Extrusion Ram Strain Ln (Zi) Ln (Zi/A) Peak load Qxfe Shape Ratio Indirect Temperature Speed Rate Load (Kelvin) (mm/s) (s**-1) (Tons)

JS430T3I T-Shape 40:1 Indirect 703 3 5.488 26.83 7.23 270 JS385T5I T-Shape 40:1 Indirect 658 5 9.048 29.05 9.45 297 JS350T5I T-Shape 40:1 Indirect 623 5 9.048 30.56 10.96 373 JS365T7I T-Shape 40:1 Indirect 638 7 12.608 30.22 10.62 367

Appendix 1. Extrusion data for alloy 2024 Table of extrusion data

Extrusion Peak APeak AT Ln (Zc) Steady State Rexn Layer Rexn Layer Volume fraction Qxfe Pressure Pressure Pressure Thickness Thickness Rextd after (MPa) (MPa) (MPa) (as extruded) (solutionised) solutionisina (nm ) (um )

JS430T3I 645.1 66.70 68.60 24.37 532.13 220 620 27.31 JS385T5I 870.9 117.60 93.59 30.06 691.50 400 1040 45.82 JS350T5I 891.2 150.50 107.19 26.00 681.44 600 1240 54.62 JS365T7I 876.9 142.70 104.94 25.94 675.46 540 1200 52.86

Appendix 1. Extrusion data for alloy 2024 Run Extrusion Extr. Extrudate Ageing Hardness Data H v (1 0 ) Code Mode Temp °C Condition Temp °( after no of hours

1 1 0 2 0 7 2 1 0 0 2 0 0

JS450R7D/20 Direct 450 Solutionised 120 105 115 117 117 119 127

JS400R7D/20 Direct 400 Solutionised 120 97 108 112 112 114 123

JS350R7D/20 Direct 350 Solutionised 120 99 105 113 113 115 125

JS300R7D/20 Direct 300 Solutionised 120 95 107 1 11 111 112 122

Run Extrusion Extr. Extrudate Ageing Hardness Data H v ( 1 0 ) Code Mode Temp °C Condition Temp °C after no of hours

0 .1 1 4 8 1 6 4 8 7 2

JS450R7D/20 Direct 450 Solutionised 160 100 120 133 139 151 136 132

JS400R7D/20 Direct 400 Solutionised 160 90 111 126 134 148 128 124

JS350R7D/20 Direct 350 Solutionised 160 89 110 123 128 147 133 124

JS300R7D/20 Direct 300 Solutionised 160 91 112 120 124 140 135 131

t o t o

Appendix 2. Ageing data for alloy 2024 Run Extrusion Extr. Extrudate Ageing Hardness Data H v (1 0 ) Code Mode Temp °C Condition Temp °C after no of hours

0 .1 1 4 8 2 6 4 8 7 2

JS450R7D/20 Direct 450 Solutionised 180 100 122 131 136 144 136 130

JS400R7D/20 Direct 400 Solutionised 180 90 113 126 133 130 130 123

JS350R7D/20 Direct 350 Solutionised 180 89 113 120 142 133 133 126

JS300R7D/20 Direct 300 Solutionised 180 91 114 116 135 136 135 127

Run Extrusion Extr. Extrudate Ageing Hardness Data Hv(1 0) Code Mode Temp °C Condition Temp °C after no of hours

1 1 0 4 8 1 0 0 2 0 0

JS450R7D/20 Direct 450 As extruded 120 102 113 115 117 125

JS400R7D/20 Direct 400 As extruded 120 95 106 108 111 121

JS350R7D/20 Direct 350 As extruded 120 95 103 105 110 122

JS300R7D/20 Direct 300 As extruded 120 92 101 103 108 115 233

Appendix 2. Ageing data fo r alloy 2024 Run Extrusion Extr. Extrudate Ageing Hardness Data H v (1 0 ) Code Mode Temp °C Condition Temp °C after no of hours

1 4 1 8 4 8 1 0 0

JS450R7D/20 Direct 450 As extruded 160 115 124 140 126 120

JS400R7D/20 Direct 400 As extruded 160 112 120 128 120 118

JS350R7D/20 Direct 350 As extruded 160 104 110 118 115 113

JS300R7D/20 Direct 300 As extruded 160 90 100 104 102 100

Run Extrusion Extr. Extrudate Ageing Hardness Data H v ( 1 0 ) Code Mode Temp °C Condition Temp °C after no of hours

1 6 8 1 2 2 4 1 0 0

JS450R7D/20 Direct 450 As extruded 180 118 130 135 131 122 115

JS400R7D/20 Direct 400 As extruded 180 116 122 124 125 121 118

JS350R7D/20 Direct 350 As extruded 180 106 110 111 113 112 107

JS300R7D/20 Direct 300 As extruded 180 93 98 100 102 100 98 234

Appendix 2. Ageing data for alloy 2024 Extrusion Extrusion Extrusion Direct/ Extrusion Ram Ln (Zi) Temper Cross-sec Proof Proof Code Shape Ratio Indirect Temp. Speed Desianation Area Load Stress (Kelvin) (mm/s) (sq. mm) .....(KN) _ (MPa)

JS450R7I/20 Round 20:01 Indirect 723 6.9 25.87 T6 22.99 9.63 468.91 JS450R7I/20 Round 20:01 Indirect 723 6.9 25.87 T5 23.11 9.63 451.70 JS450R7I/20 Round 20:01 Indirect 723 6.9 25.87 F 23.16 8.50 407.01 JS400R7I/20 Round 20:01 Indirect 673 6.9 27.68 T6 22.99 8.74 451.17 JS400R7I/20 Round 20:01 Indirect 673 6.9 27.68 T5 23.20 8.43 398.37 JS400R7I/20 Round 20:01 Indirect 673 6.9 27.68 F 23.16 7.79 366.36 JS350R7I/20 Round 20:01 Indirect 623 6.9 29.79 T6 23.24 8.69 440.92 JS350R7I/20 Round 20:01 Indirect 623 6.9 29.79 T5 23.20 8.25 390.60 JS350R7I/20 Round 20:01 Indirect 623 6.9 29.79 F 22.99 7.88 372.75 JS300R7I/20 Round 20:01 Indirect 573 6.9 32.26 T6 23.20 8.13 425.43 JS300R7I/20 Round 20:01 Indirect 573 6.9 32.26 T5 23.24 7.94 349.65 JS300R7I/20 Round 20:01 Indirect 573 6.9 32.26 F 23.16 7.50 353.83

JS450R7D/20 Round 20:01 Direct 723 6.9 25.59 T510 23.16 13.20 571.30 JS450R7D/20 Round 20:01 Direct 723 6.9 25.59 T51 23.16 12.30 532.10 JS450R7D/20 Round 20:01 Direct 723 6.9 25.59 T6 23.11 9.16 446.37 JS450R7D/20 Round 20:01 Direct 723 6.9 25.59 T5 23.16 9.12 428.78 JS450R7D/20 Round 20:01 Direct 723 6.9 25.59 F 22.99 8.10 382.33 JS400R7D/20 Round 20:01 Direct 673 6.9 27.41 T510 22.99 12.80 557.10 JS400R7D/20 Round 20:01 Direct 673 6.9 27.41 T51 23.16 11.65 503.50 JS400R7D/20 Round 20:01 Direct 673 6.9 27.41 T6 22.99 8.80 440.78 JS400R7D/20 Round 20:01 Direct 673 6.9 27.41 T5 23.20 8.46 399.66 JS400R7D/20 Round 20:01 Direct 673 6.9 27.41 F 23.16 7.73 363.78 JS350R7D/20 Round 20:01 Direct 623 6.9 29.52 T6 22.99 8.30 422.03 JS350R7D/20 Round 20:01 Direct 623 6.9 29.52 T5 22.99 7.89 378.19 JS350R7D/20 Round 20:01 Direct 623 6.9 29.52 F 23.11 7.57 357.56 JS300R7D/20 Round 20:01 Direct 573 6.9 31.99 T510 23.16 11.90 515.20 JS300R7D/20 Round 20:01 Direct 573 6.9 31.99 T51 22.99 10.15 441.00 JS300R7D/20 Round 20:01 Direct 573 6.9 31.99 T6 23.20 7.66 410.17 JS300R7D/20 Round 20:01 Direct 573 6.9 31.99 T5 23.20 7.53 359.57 JS300R7D/20 Round 20:01 Direct 573 6.9 31.99 F 23.16 7.15 338.72 235

Appendix 3. Tensile tests data for alloy 2024 Extrusion Temper U.T.S. U.T.S. Percentage (%) Final Percentage (%) Code Desiqnation Load Elonaation Cross-sectional Reduction in _ (KN) (MPa) Area Cross-sectional (sa. mm) Area

JS450R7I/20 T6 15.11 606.20 17.82 19.17 16.60 JS450R7I/20 T5 14.68 562.22 18.59 18.71 19.04 JS450R7I/20 F 13.75 488.70 22.44 17.32 25.22 JS400R7I/20 T6 13.49 535.77 18.94 19.12 16.83 JS400R7I/20 T5 12.95 485.19 19.20 17.98 22.50 JS400R7I/20 F 13.11 461.07 24.40 16.55 28.54 JS350R7I/20 T6 13.36 523.87 20.02 18.97 18.37 JS350R7I/20 T5 12.71 472.84 19.38 17.05 26.51 JS350R7I/20 F 13.67 461.40 26.19 15.38 33.10 JS300R7I/20 T6 11.99 465.80 19.58 16.30 29.74 JS300R7I/20 T5 11.48 420.98 19.58 15.80 32.01 JS300R7I/20 F 12.89 425.60 26.51 23.08 33.72

JS450R7D/20 T510 15.14 655.50 11.60 20.94 9.60 JS450R7D/20 T51 14.10 607.40 14.90 20.42 11.80 JS450R7D/20 T6 14.53 598.73 17.72 19.12 17.27 JS450R7D/20 T5 13.95 550.33 18.63 18.72 19.17 JS450R7D/20 F 13.18 463.29 22.61 17.32 24.66 JS400R7D/20 T510 14.98 651.60 12.90 20.18 12.20 JS400R7D/20 T51 13.40 577.10 15.70 19.94 13.90 JS400R7D/20 T6 13.47 555.91 18.78 18.98 17.44 JS400R7D/20 T5 12.85 505.88 19.16 18.00 22.41 JS400R7D/20 F 13.09 455.20 24.60 16.44 29.00 JS350R7D/20 T6 12.78 525.90 20.40 18.76 18.40 JS350R7D/20 T5 12.12 491.19 19.61 17.05 25.84 JS350R7D/20 F 13.05 440.69 26.02 15.53 32.80 JS300R7D/20 T510 14.98 646.60 15.50 18.89 18.40 JS300R7D/20 T51 12.20 531.20 20.00 18.01 21.70 JS300R7D/20 T6 11.99 496.80 19.56 16.28 29.83 JS300R7D/20 T5 11.56 450.28 19.52 15.76 32.07 236 JS300R7D/20 F 12.92 424.86 26.43 15.33 33.81

Appendix 3. Tensile tests data for alloy 2024 Extrusion Extrusion Extrusion Ram Direct/ Ln (Zj) Extrusion Temper Fracture Code Shape Ratio Speed Indirect Temp. Designation Touqhness (mm/s) (Kelvin) (MPa)

JS450R7I/20 Round 20:01 6.90 Indirect 25.87 723 F 18.40 JS450R7I/20 Round 20:01 6.90 Indirect 25.87 723 T5 18.45 JS450R7I/20 Round 20:01 6.90 Indirect 25.87 723 T6 22.96 JS400R7I/20 Round 20:01 6.90 Indirect 27.68 673 F 20.17 JS400R7I/20 Round 20:01 6.90 Indirect 27.68 673 T5 20.41 JS400R7I/20 Round 20:01 6.90 Indirect 27.68 673 T6 25.60 JS350R7I/20 Round 20:01 6.90 Indirect 29.79 623 F 21.52 JS350R7I/20 Round 20:01 6.90 Indirect 29.79 623 T5 21.71 JS350R7I/20 Round 20:01 6.90 Indirect 29.79 623 T6 26.32 JS300R7I/20 Round 20:01 6.90 Indirect 32.26 573 F 21.81 JS300R7I/20 Round 20:01 6.90 Indirect 32.26 573 T5 21.92 JS300R7I/20 Round 20:01 6.90 Indirect 32.26 573 T6 27.28

JS450R7D/20 Round 20:01 6.90 Direct 25.59 723 F 18.35 JS450R7D/20 Round 20:01 6.90 Direct 25.59 723 T5 18.45 JS450R7D/20 Round 20:01 6.90 Direct 25.59 723 T6 22.88 JS400R7D/20 Round 20:01 6.90 Direct 27.41 673 F 20.07 JS400R7D/20 Round 20:01 6.90 Direct 27.41 673 T5 20.29 JS400R7D/20 Round 20:01 6.90 Direct 27.41 673 T6 25.36 JS350R7D/20 Round 20:01 6.90 Direct 29.52 623 F 21.43 JS350R7D/20 Round 20:01 6.90 Direct 29.52 623 T5 21.64 JS350R7D/20 Round 20:01 6.90 Direct 29.52 623 T6 26.11 JS300R7D/20 Round 20:01 6.90 Direct 31.99 573 F 21.74 JS300R7D/20 Round 20:01 6.90 Direct 31.99 573 T5 21.93 JS300R7D/20 Round 20:01 6.90 Direct 31.99 573 T6 27.11

u>to

Appendix 4. Short rod fracture toughness data of alloy 2024 238

Appendix 5

Experimental errors

During the experimental programme, repeat tests were made whenever possible to ensure that no 'freak' value affected the final result.

In extrusion runs, due to the limited material, extrusions were not repeated. Hence the results should be treated with care. Extrusion being a complex process with numerous variables is prone to errors. The two most important sources of errors are the ram speed and initial billet temperature. The initial billet temperature is known to an accuracy of 1°C and coupled with the short transfer time to the container minimises any possible errors. The second source, the ram speed is calculated with the aid of a stopwatch and marker and hence human error is possible. The errors involved in the load calculations were 4% and that of the ram speed was approximately 6%.

The accuracy of the hot working constants derived from the torsion tests cannot be easily assessed as it depends on the experimental results such as the torque-twist curves, heating rate and temperature rise. However, in calculating the torque and flow stress values there was an error of 7%. The temperature rise model itself could not be experimentally verified during extrusion but was considered realistic after comparison with the structure and properties of the extrudates.

The subgrain size measurements were made using the Mean Linear Intercept Method. Errors can arise due the inherent wedge shape of the transmission electron microscope specimens. The relative error using this technique for the measurements was approximately 9%. The recrystallised layer thickness measurements although revealing the trend of the thickness with extrusion temperature and shape should be treated cautiously. As mentioned in chapter 7, the layer thickness in the rod extrudates was uniform and the average measurements were made easily. In the case of the shape extrudates however, there was a variation in thickness around the extrudate. The variation between the thickest and thinnest layers was approximately 9%. For industrial purposes however, the maximum thickness may be more important. 239

In tensile tests and short rod fracture toughness tests, 3 tests were carried out and the mean value was taken. The errors involved here were 6% and 12% respectively. In hardness testing, where each value quoted was based on an average of 10 readings, the error involved was 4%. 240

Appendix 6

Modification to the equation calculating u. the coefficient of friction

V -°~x v

^ (°x+ dax) IcD^°x dx = 0

but at the interface crx=CT|, therefore;

2 TCP - rcDjiGjdx = 0 4 K + dox) - rearranging;

4 no, d

Integrating these terms between the limits;

7 f 4 n c V '

-fd

P 2 - P 1 = 4 hct|(L2-L 1)=A P

where;

4na.(L -L ) P 2 = ct(0.47 + 1.2LnR) + ------^

4na.(L -L ) P, = d (0.47 + 1.2LnR) + ------^

Hence;

'4*10, P - P = AL 2 1 V D rearranging for \i;

(p2- pi)p ^ " ( L2- Li)4a,

Terms used

P2 and P-| are steady state pressures for billet lengths L2 and 1^

D = billet diameter

AL = change in billet length

crf= flow stress of bulk material

In the calculation of the strain rate;

v_ x where; v = ram speed x = aluminium layer thickness at the billet/container interface (this thickness was initially assumed to be 100 pm thick) 243

References

1) Proceedings of Inti. Aircraft Industries, Torino, Oct. 1976

2) Proceedings of Inti. Aircraft Industries, Torino, Oct. 1976

3) A. Wilm, 'Aluminium' (1936),18,336

4) P. D. Merica et al, Sci. Paper (US Bureau Of Standards) (1919), 347

5) K. R. Horn, Aluminium, Vol. 1, (1967), A.S.M. Metals Park, Ohio

6) A. Perrone, Alumino (1976), 45,512

7) D. S. Thompson, 'Met. Trans.' (1975), 6A, 671

8) M. V. Lancker, Metallurgy of Aluminium Albys, (1967), Chapman Hall, London

9) L. F. Mondolfo, Aluminium Alloys, Vol 1 (1976) Butterworth

10) M. V. Hyatt, Alumino (1977), 46,81

11) R. B. Ross, 'Metallic Materials Specificartion Handbook' (1972) SPON, London

12) H. Westengen et al, Proc. Ist.Riso Inti Symp. Met and Mat Sci, (1980), Riso

13) J. Blade et al, 'Met. Tech.' (1976), 3,380

14) Aluminium Assoc., Aluminium Standards and Data (1979)

15) P. C. Varley, The Technology of Al and its Albys'. (1970), Butterworth, London

16) S. J. Patterson, Ph. D Thesis, University of London, 1981

17) R.P. Vierod, Ph. D Thesis, University of London, 1983

18) I. Kovacs et al 'Acta Met' (1972), 20,975

19) A. Guinler, Compt. Rend.' (1938), 206,1641

20) R. Nicholson and J. Nutting, 'Phil. Mag.' (1958), 38,531

21) V. A. Phillips, 'Acta Met’, (1973), 21,219

22) J. W. Martin, 'Micromechanisms in Particle Hardened Alloys', (1980),

Camridge University Press

23) H. K. Hardy, 'J. I. M.', (1954), 83,17

24) R. Aekert, 'J. I. M.', (1967), 95, 204 244

25) J. M. Silcock, 'J. I. M.' (1960), 89, 203

26) W. Bonfield and P. K. Datta, J. Mat Sci\ (1977), 12,1050

27) W. Bonfield and P. K. Datta, J. Mat Sci\ (1976), 11,1661

28) M. O. Spiedal, 'Met. Trans.’ (1975), 6A, 631

29) G. Hahn and A. Rosenfield, 'Met. Trans' (1975), 6A, 653

30) G. Hahn and A. Rosenfield, 'Met. Trans' (1975), 6A, 653

31) G. G. Garet et al 'Met. T rans' (1978), 9A, 1187

32) R. Van Stone and J. Psioda, 'Met. Trans.' (1975), 6A, 668

33) K. Sugimoto et al, 'Corrosion Science', (1975), 15,709

34) W. Bell and H. Campbell, 'J. I. M.', (1960/61), 89,464

35) E. Liddiar and J. Whittaker, 'J. I. M .\ (1960/61), 89,423

36) M. Fine, 'Met. Trans.', (1975), 6A, 625

37) J. Jonas et al, 'Met. Rev.' (1969), 14,1

38) H. McQueen and J. Jonas, 'Treatise on Mat Sci\ Vol 6, p393, Academic Press,

1975

39) M. G. Thutcher, Ph.D. Thesis, University of London, 1979

40) H. J. McQueen and J. Jonas, 'Hot Workability Testing Techniques'

41) F. Garafola, 'Trans Met. Soc. AIME', (1963), p227

42) C.M. Sellars & W.J. McTegart, Int. Met. Rev.,1972,17,1

43) P. R. Swan, 'Electron Microscopy and Strength of Metals', (1963), Interscience

44) T. Sheppard and D. Raybould, 'J. I. M.', (1973), 101,65

45) J. Jonas, H. McQueen and W. Wong, 'Deformation under Hot Working Conditions',

I. S. I. Publications, (1968), 108,49

46) O. Sherby et al, 'Met Trans", (1977), 25,897

47) H. McQueen and J. E. Hockett, 'Met Trans', (1970), 1,2296

48) A. Korbel et al, 'Met Tech', (1979), 6, 391 245

49) H. Stuwe, 'Deformation under Hot Working conditions', (1968), J. S. I., 108, P1

50) A.F. Castle, PhD Thesis, University of London, 1974

51) E.P. Wood, PhD Thesis, University of London, 1978

52) T. Sheppard, 'Modern Extrusion Symposium'. (1980), Birmingham, Met. Assoc.

53) B. Avitzur et al, Trans AIME', (1965), 8 7 8 ,1 ,5 7

54) W. Johnson et al, 'The mechanics of metal extrusion', (1962), Manchester

University Press

55) W. Ziegler et al, 'Metaiikunde', (1973), 63,224

56) A. Castle and T. Sheppard,' Metals Tech.', (1976), 63,224

57) S. Hirst & D.H. Ursell, 'Metal Treatment', (1958), 25,409

58) K.A. Ashcroft, G.S. Lawson, 'J.I.M.', 1960-61,89,369

59) B.J. Meadows, M J. Cutler, 'J.I.M.', 1969, 97,321

60) T. Sheppard, D. Raybould, 'J.I.M.', 1973,101,73

61) T. Sheppard, 'Proc 2nd Inti Al. Ext, Tech. Sem.', 1979,1,331

62) K. Laue, H. Homauer, 'Z. Metallkd', 1956,47,117

63) T.M. Watkins, K. Ashcroft, J. McKenzie, 'J. Proc Tech Eng Manuf.', 1958, p155

64) E.P. Wood, T. Sheppard, 'Metal Tech', 1980,2,58

65) O. Kienzle, K. Gruning, 'Metaiikunde', 1962,16,276

66) M.P. Clode, PhD Thesis, University of London, 1987

67) D.S. Wright, PhD Thesis, University of London,1978

68) T. Sheppard, E.P. Wood, 'Proc 17th M.T.D.R., 1976, McMillan, London, p411

69) T. Sheppard, M.G. Tutcher, 'Met. Tech.', 1980,12,488

70) S.J. Cieslak, 'ALCOA short rod fracture toughness indicator test', report no. GA

17684.1, ALCOA Engineering Divison, 1985

71) P.A. Tunnicliffe, PhD Thesis, University of London, 1979 246

72) F. Garafola, 'Fundamentals of creep & creep rupture in metals', (1965),

McMillan, New York

73) H.Oikawa et al.'J.Metal Science', 1975,9,209

74) K. Murty et al, 'Acta Met.', 1962,20,1009

75) R.W. Cahn(Editor), 'Physical Metallurgy', North Holland, Amsterdam, p821

76) P. Ccooper, PhD Thesis, University of London,

77) W.A. Wong, J.J. Jonas, 'Trans Met Soc AIME', 1968,242,2271

78) C.E. Parson, R.N. Parkins, 'The extrusion of metals', 1961, Chapman & Hall,

London

79) L. Anand, T rans ASME', 1982,104,12

80) D. Lee, F. Zaverl, 'Acta Met', 1978,26,177

81) U. M. Sample, L.A. Lall, 'Mat Sci & Tech', 1987,3,28

82) J. Cotner, W.J. McG Tegart,' J.I.M.', 1969, 97,73

83) G. Lange, 'Z. Metallkd', 1971,67,571

84) M.G. Tutcher, T. Sheppard, 'Met Tech', 1980,7,488

85) T. Sheppard, 'Met Tech', 1981,8,130

86) T. Sheppard, S.J. Patterson, 'Met Tech', 1982,9,274

87) S. J. Patterson, PhD Thesis, University of London, 1981

88) S. J. Patterson, PhD Thesis, University of London, 1981

89) M. A. Zaidi, PhD Thesis, University of London, 1980

90) R. E. Smallman, 'Modern Physical Metallurgy', 1970, Butterworth, London

91) R. Nicholson, G. Thomas and J. Nutting, 'J.I.M.', 1958-59,87,429

92) A. Munitz, 'Met. Trans.', 1980,11,1863

93) J. Eulitz, G. Scharf, 'Aluminium', 1975,50,10,639

94) Hiroaki Yokobayashi, 'Proc 2nd Inti Extrusion Tech. Seminar', 1977, Vol. 1,

115 247

95) V. A. Phillips, A.J. May, 'J.I.M.', 1952-53, 81,265

96) A. T. Thomas, 'Acta Met.', 1966,14,1363

97) B. J. Brindley, P.J. Worthington, 'Acta Met.', 1969,17,1357

98) A. Cottrell, 'Phil. Mag.', 1958,44, 829

99) Hiromasa Hayashi, 'Light Metal Age', 1975, Feb., p8

100 ) F.J. Humphreys, 'Acta Met.', 1977, 25,1323

101) T. Sheppard, M.A. Zaidi, 'Met. Tech. 1982,9,368

102) M. Farag, C. M. Sellars, 'J.I.M.', 1973,101,137

103) W. Johnson, H. Kudo, 'The mechanics of metal extrusion', 1962, Man. Univ Press