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Microstructural evolution and mechanical properties of a copper-zirconium alloy processed by severe plastic deformation

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Microstructural evolution and mechanical properties of a copper-zirconium alloy processed by severe plastic deformation

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MICROSTRUCTURAL EVOLUTION AND MECHANICAL PROPERTIES OF A
COPPER-ZIRCONIUM ALLOY PROCESSED BY SEVERE PLASTIC
DEFORMATION

by
Jittraporn Wongsa-Ngam

A Dissertation Presented to the
FACULTY OF THE USC GRADUATE SCHOOL
UNIVERSITY OF SOUTHERN CALIFORNIA
In Partial Fulfillment of the
Requirements for the Degree
DOCTOR OF PHILOSOPHY
(MECHANICAL ENGINEERING)

May 2013

Copyright 2013 Jittraporn Wongsa-Ngam
ii

Acknowledgments
Firstly of all, I am deeply grateful to my advisor Professor Terence G. Langdon
for his encouragement, support and guidance throughout my Ph.D. study. This
dissertation would not have been possible without his invaluable advice. I hope that one
day I would become as good an advisor to my students as he has been to me. My thanks
also extend to his wife for her concern and kindness.
I am grateful to my dissertation committee, Professor Andrea M. Hodge and
Professor Edward Goo for their evaluation of my Ph.D. study. I am truly thankful to my
collaborators Dr. Youghao Zhao and Haiming Wen of the Department of Chemical
Engineering and Materials Science, University of California Davis for their outstanding
transmission electron microscope (TEM) work. I would also like to show my sincere
appreciation to Professor Steven R. Nutt, Professor John Platt and again Professor Andrea
M. Hodge for sharing the instrument of the material preparation.
Furthermore, I would like to take this opportunity to thank Dr. Megumi
Kawasaki, Dr. Roberto B. Figueiredo, Dr. Mahmood Shirooyeh, Dr. Vitor Sordi, Dr. Yi
Huang and my colleagues at USC for their instructive suggestions and discussions. I
would like to extend my emotional thanks to my friends for their friendships and
encouragement during this period. Their support and care helped me overcome setbacks
and stay focused on my graduate study.
iii

In addition, I am truly indebted and thankful to my country and the Royal Thai
Government for continuous funding me throughout my degree in the United States. I am
also grateful to all staff of Thai government for their advices and suggestions. Besides, I
extend my thanks to the staff of USC for their various forms of support during my study.
Most importantly, none of this would have been possible without the love of my
family. I would like to present this achievement as a gift to my grandma, dad and mom,
sisters, brother and the whole family members. It is their everlasting love, moral supports
and understanding that boosted my determination in pursuing this degree. I would like to
express my heart-felt gratitude to my family.
iv

Table of Contents
Acknowledgments ii

List of Tables vi

List of Figures vii

Abstract xii

1 Introduction 1

2 Literature Review 3
2.1 Severe plastic deformation (SPD) 3
2.2 Equal-channel angular pressing (ECAP) 5
2.2.1 Principle of ECAP 5
2.2.2 The fundamental parameters in processing by ECAP 9
2.2.2.1 Die angle  and  9
2.2.2.2 Pressing route 11
2.2.2.3 Pressing temperature 15
2.3 High-pressure torsion (HPT) 16
2.3.1 Principle of HPT 16
2.3.2 The fundamental parameters in processing by HPT 20
2.4 Microstructural characteristic after SPD 23
2.4.1 Grain refinement model in ECAP processing 24
2.4.2 Grain refinement model in HPT processing 27
2.5 Mechanical properties of materials processed by SPD 30
2.5.1 Strength and ductility 30
2.5.2 Superplasticity 34

3 Experimental Materials and Procedures 37
3.1 Experimental material 37
3.2 Methods of severe plastic deformation 37
3.2.1 ECAP processing 37
3.2.2 HPT processing 38
3.2.3 ECAP followed by HPT processing 39
3.3 Microstructural analysis 40
3.3.1 Optical microscope 40
3.3.2 Transmission electron microscope (TEM) 40
3.3.3 Electron-backscatter diffraction (EBSD) 41
3.4 Mechanical experiments 42
3.4.1 Microhardness measurements 42
v

3.4.1.1 Microhardness measurement of ECAP 42
3.4.1.2 Microhardness measurement of HPT 43
3.4.2 Tensile measurements 45

4. Experimental Results 47
4.1 ECAP experimental results 47
4.1.1 Microstructural observations after ECAP processing 47
4.1.1.1 TEM analysis for ECAP 47
4.1.1.2 EBSD analysis for ECAP 50
4.1.2 Microhardness measurement after ECAP processing 54
4.1.2.1 Hardness values on X plane 54
4.1.2.2 Hardness values on Y plane 58
4.1.3 Tensile properties of ECAP samples 63
4.2 HPT experimental results 66
4.2.1 Microstructural observation after HPT processing 66
4.2.1.1 TEM analysis for HPT 66
4.2.1.2 EBSD analysis for HPT 71
4.2.2 Microhardness measurements after HPT processing 75
4.2.3 Tensile properties after HPT processing 79
4.3 ECAP + HPT experimental results 84
4.3.1 Microhardness measurements after ECAP + HPT processing 84
4.3.2 Microstructural observations after ECAP + HPT processing 86

5. Discussion 91
5.1 ECAP discussion 91
5.1.1 Significant of the ECAP processing route 91
5.1.2 Microstructural evolution and grain refinement during
ECAP processing 93
5.1.3 A comparison of microhardness evolution on the X and Y plane
in ECAP processing 95
5.1.4 Hardness homogeneity after ECAP processing 100
5.2 HPT discussion 102
5.2.1 Hardness and microstructural evolution in HPT processing 102
5.2.2 Potential for achieving superplastic flow after HPT processing 107
5.3 Comparison of microstructures and mechanical properties in
different processing 110
5.3.1 Microstructures in different SPD techniques 110
5.3.2 Mechanical properties in different SPD techniques 113

6. Summary and conclusion 119

References 122

Alphabetized Bibliography 132
vi

List of Tables

Table 4.1 Microhardness values along the radii of Cu-0.1%Zr disks
processed by HPT. 77

Table 5.1 Average values of the microhardness and the 95% confidence
levels for the Cu-0.1% Zr alloy processed by ECAP through
one to eight passes. 98

Table 5.2 Microstructural parameter for the Cu-0.1% Zr alloy obtained
by TEM and EBSD 111

vii

List of Figures
Figure 2.1 The principle of ECAP through a die, showing the two angles
Φ and Ψ [19]. 7

Figure 2.2 Illustration of the X, Y and Z planes denote the transverse
plane, the flow plane and the longitudinal plane, respectively [20]. 8

Figure 2.3 Microstructures and SAED patterns when each sample is pressed
to different die angles with an imposed strain of ~4 [3, 19] . 10

Figure 2.4 Schematic illustration of four ECAP pressing routes [31]. 13

Figure 2.5 Shearing characteristics for four different processing routes [36]. 14

Figure 2.6 Grain size after ECA pressing versus pressing temperature for
pure Al, Al–3% Mg and Al–3% Mg–0.2% Sc [41]. 41

Figure 2.7 Schematic illustration of HPT processing [49]. 18

Figure 2.8 Principles of HPT for (a) unconstrained and (b, c) two types of
constrained condition [50]. 18

Figure 2.9 Parameters used in estimating the total strain in HPT [51]. 19

Figure 2.10 Parameters use in estimating the total strain in HPT [51]. 22

Figure 2.11 A model for grain refinement in the central region of the
billet in ECAP [69]. 25

Figure 2.12 The appearance of the microstructures on the Y plane after
four passes using routes A, B
C
and C [70]. 26

Figure 2.13 Schematic illustration of microstructural evolution with
straining along in pure Al. Thin double lines depicted
in region II represent low angle boundaries with some
extension of boundary width and thick lines in region III
represent high angle boundaries [71]. 28

Figure 2.14 Schematic illustration of microstructural evolution with
straining for grain refinement in pure Cu [55]. 29

viii

Figure 2.15 A comparison of yield strength and ductility for a 3004
aluminum alloy processed by cold rolling or ECAP [75]. 31

Figure 2.16 Tensile engineering stress-strain curves of (a) Cu tested at 22 C
and (b) Ti tested at 250 C at a strain rate of 10
-3
s
-1
[72, 76]. 33

Figure 3.1 Schematic illustrations showing the position of hardness
measurement on the longitudinal Y plane and cross-sectional
X plane. 43

Figure 3.2 Grid pattern of the Vickers hardness measurements on the HPT
disks. 44

Figure 3.3 Dimensions of tensile specimens [93]. 46

Figure 4.1 TEM images of the disks processed by ECAP for (a) one pass,
(b) two passes, (c) four passes and (d) eight passes. 48

Figure 4.2 Histograms of the grain size distributions measured at
(a) four passes and (b) eight passes: the individual average
grain sizes are indicated. 49

Figure 4.3 EBSD orientation images of the Cu-0.1% Zr disks processed
by ECAP for (a) four passes and (b) eight passes 52

Figure 4.4 Distributions of the number fraction of the misorientation
angles of the grain boundaries at the centers and edges
of the disks processed by ECAP through (a) four passes and
(b) eight passes. 53

Figure 4.5 Color-coded contour maps showing the distributions of the
Vickers microhardness on the cross-sectional planes after
ECAP for (a) one pass, (b) two passes, (c) four passes
and (d) eight passes. 56

Figure 4.6 Histograms of the Vickers Hardness (Hv) distribution of
Cu-0.1%Zr after ECAP (a) 1 pass, (b) 2 passes, (c) 4 passes
and (d) 8 passes on cross-sectional plane. 57

Figure 4.7 Color-coded contour maps showing the distributions of the
Vickers microhardness on the longitudinal planes after ECAP
for (a) one pass, (b) two passes, (c) four passes and (d) eight passes. 59

ix

Figure 4.8 The histogram of the Vickers hardness (Hv) distribution of
Cu-0.1%Zr after ECAP (a) 1 pass, (b) 2 passes, (c) 4 passes
and (d) 8 passes on longitudinal plane. 60

Figure 4.9 Individual values of Hv with regression lines recorded on the
longitudinal planes at 1.0 mm from the upper and lower
surfaces for the billets pressed through (a) one pass, (b) two
passes, (c) four passes and (d) eight passes. 62

Figure 4.10 Plot of engineering stress versus engineering strain for the
Cu-0.1%Zr alloy in an annealed specimen and after ECAP
processing through 4 passes and 8 passes at testing temperature
of 673 K with different strain rates. 64

Figure 4.11 Elongation to failure versus strain rate for the Cu-0.1%Zr
alloy in an annealed condition and after ECAP processing
through four and eight passes at testing temperature of 673 K 65

Figure 4.12 TEM images of the disks processed by HPT for (a) 1/4 turn,
(b) 1 turn, (c) 5 turns and (d) 10 turns: the left column
corresponds to the centers of the disks and the right column
corresponds to the edges of the disks. 68

Figure 4.13 Histograms of the grain size distributions measured at
(a) a peripheral region of the disk after 1/4 turn,
(b) a peripheral region of the disk after 1 turn,
(c) a central region of the disk after 5 turns (d) a peripheral
region of the disk after 5 turns (e) a central region of the
disk after 10 turns (f) a peripheral region of the disk after
10 turns: the individual average grain sizes are indicated. 70

Figure 4.14 EBSD orientation images of the Cu-0.1% Zr disks processed
by HPT for (a) 1/4 turn, (b) 5 turns and (c) 10 turns. 72

Figure 4.15 Distributions of the number fraction of the misorientation
angles of the grain boundaries at the centers and edges
of the disks processed by HPT through (a) a quarter turn
(b) five turns and (c) ten turns. 74

Figure 4.16 Values of the Vickers microhardness versus distance from
the center of the Cu-0.1%Zr disks after HPT processing
for various numbers of turns. 76

x

Figure 4.17 Color-coded contour maps showing the distributions of the
Vickers microhardness values over disks processed by
HPT at a present of 6.0 GPa for (a) 1/4 turn, (b) 1 turn,
(c) 5 turns and (d) 10 turns. 78

Figure 4.18 Plots of engineering stress versus engineering strain for the
Cu-0.1% Zr alloy in an as-received specimen and after HPT
processing through one, five and ten turns at an initial strain rate
of 1.0×10
-3
s
-1
and at temperatures of (a) 673 K and (b) 723 K. 81

Figure 4.19 Plots of engineering stress versus engineering strain for the
Cu-0.1% Zr alloy after (a) five turns and (b) ten turns at
temperature of 673 K with various initial strain rates. 82

Figure 4.20 Elongation to failure versus strain rate for the Cu-0.1% Zr alloy
in an as-received condition and after HPT processing through one,
five and ten turns and testing at temperatures of (a) 673 K and
(b) 723 K. 83

Figure 4.21 Color-coded contour maps and histograms of Vickers
microhardness distribution over disks processed by ECAP
(a), (b) 4p + 5 turns and (c), (d) 8p + 5 turns. 85

Figure 4.22 TEM figures of disks processed by ECAP 4 passes followed
by HPT 5 turns at (a) center and (b) edge of the disks. 87

Figure 4.23 Histograms of the grain size distributions at the centers and edges
of the disks processed by ECAP 4p + HPT 5 turns. 88

Figure 4.24 TEM figures of disks processed by ECAP 8 passes followed
by HPT 5 turns at (a) center and (b) edge of the disks. 89

Figure 4.25 Histograms of the grain size distributions at the centers
and edges of the disks processed by ECAP 8p + HPT 5 turns. 90

Figure 5.1 The average Vickers microhardness values measured on
the cross-sectional X plane and on the longitudinal Y plane
versus the number of ECAP passes. 97

Figure 5.2 EBSD orientation images of the Cu-0.1% Zr disks processed
by ECAP eight passes at different scanned positions of
(a) X plane and (b) Y plane combined with (c) corresponding
distributions of misorientation in each plane. 99

xi

Figure 5.3 Flow stress versus strain rate for the HPT specimen
subjected through 10 turns having tensile testing at
temperature of 723 K. 108

Figure 5.4 Color-coded contour maps showing the distributions
of the Vickers microhardness values (a) on the cross-sectional
plane after ECAP for 8 passes and (b) over the disk surface
after HPT for 10 turns: the hardness colors are shown in
the key on the right. 112

Figure 5.5 Color-coded contour maps showing the distributions
of the Vickers microhardness values processed by
(a) ECAP 8 passes (b) HPT 10 turns and (c) ECAP 8 passes
followed by HPT 10 turns. 115

Figure 5.6 Plot of engineering stress versus engineering strain at a testing
temperature of 673 K and an initial strain rate of 1.0 × 10
-3
s
-1

for the Cu-0.1% Zr alloy in the as-received condition, in the
annealed condition, after processing by ECAP for 8 passes
and after processing by HPT for 10 turns. 118

xii

Abstract
A copper alloy, Cu-0.1% Zr, has been processed at room temperature by different
techniques of severe plastic deformation (SPD), namely equal-channel angular pressing
(ECAP), high-pressure torsion (HPT) and a combination of both processing (ECAP +
HPT). The experiments were conducted to evaluate the microstructural evolution and
mechanical properties for each of the processed and their combination. A transmission
electron microscopy (TEM) and an electron backscatter diffraction (EBSD) techniques
were employed to measure the microstructural features, grain size distributions and the
distribution of the misorientation angles. The mechanical properties of the processed
samples were examined and compared both at a room temperature using microhardness
measurements and at an elevated temperature using tensile testing.
Using TEM and EBSD techniques, it is demonstrated that these three SPD
procedures have a potential for producing an ultrafine-grain structure containing
reasonably equiaxed grains with high-angle boundary misorientations. However,
microstructures are refined in different level depending on the processing operation. The
grain refinement mechanisms are primarily governed by dislocation activities.
Microhardness distribution of the strained samples shows that there is a non-
uniform of this distribution in the early stages of deformation where the lower hardness
values were measured near the bottom of samples for ECAP and at the central region for
HPT. This inhomogeneity is gradual decreased with increasing imposed strain and
xiii

ultimately the microhardness distribution is reasonably homogeneous when the sufficient
strain is subjected to the sample.
The tensile results demonstrate that the samples after SPD processing exhibit
superior mechanical properties with the combination of high strength and ductility
compared to the as-received materials where the maximum elongation to failure of
~240% at 723 K using a strain rate of 1.0 × 10
-4
s
-1
is achieved in a sample processed by
HPT. This elongation however does not fulfilled the requirements for true superplastic
flow where the measured elongation in tension should be at least 400%

1

1. Introduction
Ultrafine-grained materials are defined formally as polycrystalline materials
having grain sizes both within the submicrometer range from 0.1 to 1.0 µm and in the
nanometer range with grain sizes below 100 nm [1]. The fabrication of bulk ultrafine-
grained (UFG) materials has become important because of the potential for producing
materials having superior mechanical and physical properties. These UFG metals may be
produced through severe plastic deformation (SPD). The principle of SPD processing is
that the sample is subjected to very high strain but it retains approximately the same
dimensions. This contrasts with other procedures where generally the overall dimensions
of the sample are reduced during the strain operation. The SPD processing has become
attractive procedure for producing bulk fully-dense materials with exceptional refined
structure that cannot be achieved using more conventional metal-working processes.
Several different SPD processing techniques are now available [1, 2] but the two most
attractive procedures, and the ones receiving most attention at the present time, are equal-
channel angular pressing (ECAP) [3] and high-pressure torsion (HPT) [4].
In processing by ECAP, a sample is pressed through a die constrained within a
channel bent through a sharp angle [3]. Specific advantages of ECAP include the ease of
scaling to produce relatively large bulk solids [5] and the possibility of incorporating the
principles of ECAP into continuous production techniques involving rolling [6-8] or the
conform process [9]. In processing by HPT, a sample, generally in form of a thin disk, is
subjected to an applied pressure and concurrent torsional straining. The HPT is attractive
2

because it generally produces significantly smaller grains and a higher fraction of high-
angle grain boundaries [10].
Copper represents an ideal model material for studying the processes of
deformation and microstructure evolution, due to its low-cost, simple fcc structure,
medium stacking-fault energy (SFE), and the long history of research of this material
prepared by conventional techniques. The addition of a small amount of zirconium leads
to a significant increase in the recrystallization temperature and a potential for retaining
an ultrafine grain size at elevated temperature [11-14]. Furthermore, it is well established
that the incorporation of Zr in copper improved the fatigue properties [15]. Accordingly,
a Cu-0.1% Zr alloy was selected in this present investigation.
The purpose of the present work is to provide a systematic study in the relation
between processing, microstructure and mechanical properties of a Cu-0.1%Zr alloy.
Specifically, the objective of the present work is threefold: First, to apply three different
SPD processing to samples of a Cu-0.1%Zr; second, to evaluate the consequent
microstructural evolution with increasing imposed strain by using TEM and EBSD; third,
to provide information on the mechanical properties of the alloy both at room
temperature and at elevated temperature after each processing and investigate how
microstructure influences mechanical properties.

3

2. Literature Review
2.1 Severe plastic deformation (SPD)
Ultrafine-grained (UFG) materials are defined as polycrystalline materials having
very small grains with average grain size less than ~1 µm. The structures of UFG
materials divide into submicrometer (0.1-1µm) and nanometer (less than 100 nm) ranges.
For bulk UFG materials, there are additional requirements of fairly homogeneous and
equiaxed microstructures with a majority of the grain boundaries having high-angles of
misorientation. These UFG materials attract considerable attention over several decades
because they are expected to have superior mechanical and physical properties by
comparison with coarse-grain materials. Thus, interest in the processing of bulk UFG
materials has grown significantly. Two different approaches have been developed for
fabricating UFG materials: “bottom-up” and “top-down” approaches [16].
In the “bottom-up” method, the individual atoms or nanoparticles are used to
assemble bulk materials. Examples of the “bottom-up” approach include inert gas
condensation, high-energy ball milling and electro deposition. These techniques are
capable of producing materials with exceptional small grain sizes, even to the nanometer
level but they might introduce some contaminations and residual porosities during the
fabrication procedures. Furthermore, products from these techniques often have limited
sizes that not appropriate for large-scale structural application.
4

In the “top-down” approach, coarse-grained materials are processed to produce
UFG materials by applying heavy straining or shock loading. Since this approach
introduce severe plastic deformation (SPD) into the materials, the general term of SPD
processing is used to call the methods. SPD processing is defined as any method of metal
forming under an extensive hydrostatic pressure that may be used to impose a very high
strain to a bulk solid without the introduction and any significant change in the overall
dimensions of the sample and having the ability to produce exceptional grain refinement.
A unique character of SPD processing is that the shape of the sample is retained in SPD
processing by the use of special tool geometries which effectively prevent free flow of
the material and thereby produce a significant hydrostatic pressure. The presence of this
hydrostatic pressure is essential for achieving high strains and introducing the high
density of lattice defects and consequent grain refinement.
Numerous techniques for SPD processing are now available including
accumulative roll-bonding (ARB), cylinder covered compression (CCC), cyclic extrusion and
compression (CEC), constrained grooved pressing (CGP), equal-channel angular pressing
(ECAP), friction stir processing (FSP), high-pressure torsion (HPT), mechanical milling (MM),
multi-directional forging (MDF), repetitive corrugation and straightening (RCS) and twist
extrusion (TE). All of these procedures are capable of introducing large plastic straining.
However, the major methods already established for the fabrication of UFG materials are ECAP,
HPT, MDF and ARB. The principles of some SPD techniques were outlined in an earlier report
[1, 2]. Among all of various SPD techniques, special attention is given to the ECAP and
HPT techniques.
5

2.2 Equal-channel angular pressing (ECAP)
Equal-channel angular pressing (ECAP) is considered to be an attractive
processing technique for following reason. First, it is a relatively simple procedure that
can be performed on a wide range of materials. Second, this method has a potential for
achieving homogeneity through most of samples processed to a sufficiently high strain.
Third, it may be scaled-up for the pressing of large samples so that there is a potential to
apply to use in a wide range of structural application.
2.2.1 Principle of ECAP
Valiev and Langdon [3] reviewed this technique and summarized the recent
developments, effects to diverse processing parameters and resulting structure and
properties. The principle of the ECAP procedure is illustrated schematically in Fig. 2.1.
The ECAP die contains two intersecting channels having an identical cross-section. The
sample is machined to fit within the channels and pressed through the die using plunger.
As the sample passes through the abrupt angle of the two channels, an intense shear strain
is imposed into the sample. A basic characteristic of this technique is that the sample
cross section remains unchanged enabling repetitive processing, by which very high total
strain could be imposed on the samples.
When a sample is pressed through an ECAP die, the strain introduced into the
material is dependent on two angles associated with the construction of the die. These
angles are also illustrated in Fig. 2.1, where the angle  represents the angle of
6

intersection of the two channels and the angle Ψ represents the outer arc of curvature
where the two parts of the channel intersect. Calculations have shown [17] and
experimental results have confirmed [18] that the equivalent strain, , accumulated in
ECAP is given by the relationship











 


 





 



2 2
cosec
2 2
cot 2
3
N
N

2.1

where N is the total number of passage through the die. During repetitive pressing, the
shear strain is accumulated in the sample leading to a UFG structure.
Three separate orthogonal planes of ECAP processing are illustrated in Fig.2.2,
where X is the transverse plane perpendicular to the flow direction, Y is the flow plane
parallel to the side face at the point of exit from the die and Z is the longitudinal plane
parallel to the top surface at the point of exit from the die.
7

Figure 2.1 The principle of ECAP through a die, showing the two angles Φ and Ψ [19].

8

Figure 2.2 Illustration of the X, Y and Z planes denote the transverse plane, the flow
plane and the longitudinal plane, respectively [20].

9

2.2.2 The fundamental parameters in processing by ECAP
2.2.2.1 Die angle  and 
An internal angle  and an outer curvature  are two important parameters that
greatly impact imposed strain and thus microstructure of materials processed by ECAP.
The influence of the internal angle  on the microstructure attained by pressing was
investigated using a die having internal angles from 90 to 157.5° for pure aluminum [19].
The investigation demonstrated that an ultrafine microstructure of essentially equiaxed
grains, separated by high angle grain boundaries, is achieved most readily when a very
intense plastic strain is imposed on the sample in each passage through the die as when
using a die having  close to 90° as shown in Fig. 2.3. Despite the efficiency of the dies
with internal angle of  = 90 , it is important to recognize that the intense strain in each
passage might cause the difficulty for forming hard materials or lead to the cracks in the
materials having low ductility. In this case, using dies with internal angles are larger than
90 would help to successfully press in hard or low ductility materials. For example, an
experiment on commercial purity tungsten show that it was not feasible to press a sample
through a 90 die at a temperature of ~1273 K because of cracking in the sample but it
was successful in pressing a sample with die having  = 110 [21].
According to eq. 2.1 the angle of curvature  plays only a minor important in
determining the imposed strains by ECAP. However, this parameter has direct relation to
the development of the corner gap or “dead zone” which is formed at the outer corner
10

when the billets pass through the die. In this zone the billet no longer remains in contact
with the outer die walls. The model experiments [18, 22] and finite element analysis [23-
29] show the use of the sharper corner  = 0  leads to a corner gap or dead zone,
however, there is insignificant change in this gap with  angle varies from 0  up to ~
20.

Figure 2.3 Microstructures and SAED patterns when each sample is pressed to different
die angles with an imposed strain of ~4 [3, 19].

11

2.2.2.2 Pressing route
The ECAP processing has a capability for conducting repetitive passes to the
sample because there is no change in the cross-sectional dimension of the sample. The
use of repetitive pressing in ECAP allows the rotation of the sample in different ways
between consecutive pressings so that there is a possibility of introducing different slip
system in the sample because of the change of shearing plane [30]. The four basic
processing routes in ECAP are shown in Fig. 2.4 [31]: in route A the sample is pressed
repetitively without rotation, in route B
A
the sample is rotated by 90˚ in alternating
directions between consecutive passes, in route B
C
the sample is rotated by 90˚ in the
same sense between each pass and in route C the sample is rotated by 180˚ between
passes. The different shearing patterns in the sample lead to variation both in the
macrostructure distortions of the individual grain in polycrystalline matrices [32] and the
capability to develop a homogeneous and equiaxed ultrafine-grained microstructure [33-
35]. Furukawa et al. [36] examined the shearing characteristics associated with these four
processing routes in the X, Y and Z planes using a cubic element for up to a total of eight
ECAP passes. The results showed that the cubic element deformed under shearing in all
of the three planes when pressing through the die using route B
C
and the shape of the
cubic element was restored after 4N passes as shown in Fig. 2.5. It was thus concluded
that route B
C
is the optimum processing route to achieve a homogeneous microstructure.
This conclusion was based on investigations only concerning about ECAP using a die
with  = 90˚ and  = 0˚. Furukawa et al. [37] later expanded the investigation of the
influence of the routes on the shear pattern for ECAP under different die conditions  =
12

90 and 120˚. Route B
C
was again shown to be the optimum processing route. It was
concluded that the development of a uniform microstructure of equiaxed grains, separated
by high-angle grain boundaries, was favored using route B
C
because (i) shearing occurs
over large angular ranges on the three orthogonal planes within the sample, (ii) there is a
regular and periodic restoration of an equiaxed structure during consecutive pressing and
(iii) deformation occurs on each orthogonal plane. Detailed in a die angle  = 120, the
effectiveness of route B
C
was reduced and the effectiveness of route A was improved but
route B
C
remains the optimum procedure.
A recent report examined the effect to the route of ECAP on structure and
properties of oxygen-fee copper [38]. This sample was conducted by routes A, B
C
, and C
with the maximum number of pressed of N = 25. The results demonstrated that the
minimum size of structure elements, ~230 nm, was obtained with route B
C
, and the
maximum fraction of high-angle boundaries, 77%, was examined in route A. The
maximum UTS and YS values corresponded to route B
C
. The ductility decreased after the
first pass and increased after ten passes, especially route B
C
.

13

Figure 2.4 Schematic illustration of four ECAP pressing routes [31].

14

Figure 2.5 Shearing characteristics for four different processing routes [36].

15

2.2.2.3 Pressing temperature
Pressing temperature is an important factor for workability in ECAP procedure.
Typically, a certain material is more difficult to press at lower temperature, for example,
pure aluminum was able to be pressed up to an imposed strain of ~20.9 without fracture
at room temperature but the imposed strain decreases to ~13.6 when pressed at lower
temperature of 123 K using a 90  die [39]. In addition, in a hard material it is often
convenient to conduct ECAP processing at high temperature, For example CP titanium
billets had failure with serious cracks during ECAP processing at low temperatures at 298
K, 398 K and even 498 K, however, the success of pressing was achieved when billets
were pressed at higher temperatures of 548 K and 598 K [40].
In addition, the pressing temperature in ECAP also plays an important role in
material characteristics. Yamashita et al [41] investigated the influence of pressing
temperature on microstructural development in samples of pure Al, and Al-3%Mg alloy
and an Al-3%Mg-0.2%Sc alloy at pressing temperature from room temperature to 573 K.
Two important effects were concluded. First, the measured grain size tends to increase
with increasing pressing temperature as shown in Fig. 2.6. Second, the fraction of high-
angle grain boundaries decreased with increasing temperature. The inability to achieve
high angle boundaries at the higher pressing temperatures was attributed to the higher
rates of recovery which lead to dislocation annihilation rather than the absorption of
dislocations into the subgrain walls.
16

Several subsequent investigations confirmed the formation of larger grains or
subgrains at higher pressing temperatures [42-47] but the lower fraction of high-angle
grain boundary at higher pressing temperatures [42, 43, 46, 47]. These reviews
demonstrated that although high temperature facilitate workability in ECAP, optimum
ultrafine-grained structures will be achieved at lowest possible pressing temperature
where any significant cracks were not generated in the billet. By maintaining a low
pressing temperature, this ensures the potential of achieving both the smallest possible
grain sizes and highest fraction of high-angle boundaries.

Figure 2.6 Grain size after ECA pressing versus pressing temperature for pure Al, Al–3%
Mg and Al–3% Mg–0.2% Sc [41].

17

2.3 High-pressure torsion (HPT)
2.3.1 Principle of HPT
The principle of processing by HPT is based on the classic work of Prof.
Bridgman in 1943 where thin disk was subjected to compression and torsion to produce
large strains [48]. Fig. 2.7 illustrated the principle of modern HPT processing [49]. In
processing by HPT, the sample in the form of thin disk is located between upper and
lower anvils and subjected concurrently to a high pressure, P and torsional straining
which is imposed through rotation of the lower anvil at room temperature or an elevated
temperature. In practice there are two type of HPT processing as shown in Fig. 2.8;
unconstrained (a) and constrained (b, c) [50]. In unconstrained HPT, the disk is placed on
the lower anvil and then subjected to an applied pressure and torsional straining as shown
in Fig. 2.8(a). Lack of any constraints leads to free outward flow of materials under the
applied pressure and little or no back-pressure is introduced into the system. By contrast,
in constrained HPT as shown in Fig. 2.8(b), the specimen is prepared to fit into a
depression in the lower anvil and the load is applied so that, in principle at least, there is
no outward flow of material during processing. Therefore, the materials using this type of
HPT are conducted under a back-pressure condition. However, it is difficult to achieve a
fully constrained condition. In practice, there is some limited outward flow between
lower and upper anvils and the experiments are conducted under a quasi-constrained
condition as shown in Fig. 2.8(c). The schematic illustration of HPT processing in Fig.
2.7 is also the example of a quasi-constrained HPT experiment.
18

Figure 2.7 Schematic illustration of HPT processing [49].

Figure 2.8 Principles of HPT for (a) unconstrained and (b, c) two types of constrained
condition [50].

19

As described above, the sample between two anvils is subjected to a high pressure
and torsional strain by rotation of the lower anvils and surface friction forces deform the
sample by shear. Due to the specific geometric shape of the sample and facility, the main
volume of the material is strained under a quasi-hydrostatic compression. In order to
determine the strain values during processing the model in Fig. 2.9 was used in
estimating the total strain in HPT

Figure 2.9 Parameters used in estimating the total strain in HPT [51].
For an infinitely small rotation dθ and a displacement dl as depicted in Fig. 2.9 dl
= r dθ where r is the radius of the disk, and the incremental shear strain dγ is given by
[51]
h
r
h
l 

d d
d   2.2
20

where h is the disk thickness. By further assuming that the thickness of the disk is
independent of the rotation angle, θ, it follows from formal integration that, since θ =
2N, the shear strain γ is given by

h
r N 



2
, 2.3
where N is the number of revolutions. To compare the shear strain value during torsion
with a strain value during deformation by other schemes, the eq. 2.3 is usually converted
to the so-called equivalent strain,  according to the Mises criterion
h
r N
3
2 


 2.4
In practice, however, there is an earlier suggestion and simply specify the HPT
strain in terms of the numbers of turns imposed on the sample [52].
2.3.2 The fundamental parameters in processing by HPT
According to the operation of the HPT facility, there are three parameters that
need to assign before operation; an applied pressure, the numbers of turns and the speed
of rotation. These parameters, therefore, play an important role to the material
characteristics and corresponding properties.
The influence of the applied pressure on the microhardness was investigated for
Ni disks subjected to a certain torsional strain of 5 revolutions under different pressures
[51]. The experiments reported that the microhardness values are non-uniform across the
21

diameter where there are lower hardness values in the center of the disk. However, an
increase in the applied pressure from 1.0 to 9.0 GPa leads to an overall increase in
hardness, especially in the centers of the disks so that at a pressure of 9.0 GPa there is a
reasonably homogeneous distribution of hardness.
The influence of the number of revolutions is also an important parameter in
processing by HPT. Several experiments investigated the effect of the number of
rotations to microhardness evolutions. Generally, at the low imposed strains (the low
numbers of turns), there is a significant variation of microhardness values across the disk
diameter. This variation essentially depends upon the recovery rate of materials. Figure
2.10 illustrated the trends anticipated for two different types of materials in the early
stages of deformation by HPT [53]. For a material such as high purity aluminum where
recovery is rapid, the hardness is higher in the center of the disk and the microstructure is
more refined. By contrast, for a material where dynamic recovery occurs at a slower rate,
the hardness is lower in the center in the initial stages of torsional straining. However,
there is a gradual evolution of homogeneous microstructure after the higher numbers of
turns. The evolution of homogeneous in microhardness and microstructure with
increasing the numbers of turns, torsional strains, were confirmed in earlier reports for
pure Cu [54-61] and also for copper alloys [60-63] processed by HPT.
22

Figure 2.10 Parameters use in estimating the total strain in HPT [51].

Unlike the two previous parameters there are only limited reports on the influence
of rotation speed on HPT experiments [64-66]. The systematic investigation on the
influence of the rotation speed on the microstructural characteristics of magnesium AZ31
alloy was presented recently [64]. The experiments were conducted under a quasi-
constrained condition at a temperature of 453 K using an applied pressure of P = 6.0 GPa
and a total of N = 5 turns with different rotational speed of 0.5, 1 and 2 rpm. The
hardness values and the average grain sizes were correlated with the local strain rates by
using the imposed rotation rates and estimates of the strains within the disks. The results
demonstrate the hardness and the average grain size both exhibit a small dependence on
the local strain rate. In addition, the influence of rotation speed on slippage in HPT
processing was provided [66]. Slippage is defined explicitly as the difference between the
23

measured torsional rotation of the disk and the rotation imposed externally by the
processing facility. The experiments were conducted on three different metals; aluminum,
copper and iron, and two different rotation speeds. The results show the extent of
slippage varies between these three materials such that there is very little slippage in
aluminum, slightly more slippage in copper and significant slippage when using iron. For
all materials, the extent of slippage increases at faster rotational speeds.
2.4 Microstructural characteristic after SPD
Severe plastic deformation processing has a grain refinement by introducing large
numbers of dislocations that arrange themselves in low-energy configurations in the form
of low-angle grain boundaries and then evolve, with additional straining, into reasonably
homogeneous arrays of ultrafine grain separated by boundaries having high angles of
misorientation. Experiments have shown the grain size distributions in these materials
generally follow a log-normal statistical function [67]. In addition, very few lattice
dislocations are observed in the smaller grains, whereas high densities of lattice
dislocations are in the grains of intermediate sizes and subgrains are present in the larger
grains [68]. Although differences are expected in the various ultrafine-grain structures
generated by SPD techniques, in practice they are all produced through intense plastic
straining and this leads to a commonality in their substructural features.

24

2.4.1 Grain refinement model in ECAP processing
Xu et al [69] constructed a descriptive and feasible model of grain refinement in
pure Al during ECAP using route B
C
based on the mechanical shearing of grains [37].
This model is shown in Fig. 2.11 where the ultimate equiaxed grain size is denoted by d
and the parameter  denotes the estimate of fractions of boundaries having high-angle
misorientation in each substructure based on theoretical estimates [37]. The equiaxed
grains are formed when the elongated arrays of grains are subsequently sheared in other
directions. An important feature of this model is that the ultimate size of the equilibrium
equiaxed grains is dictated by the width of the subgrain bands produced in the initial pass,
whereas the misoreintation angle of grain boundaries increases gradually with an increase
of the numbers of ECAP passes.
The final equiaxed microstructure produced by ECAP is significantly related with
the shearing system differentiated by the pressing routes. The schematic appearance of
the microstructures in the materials on the Y plane after four passes using different routes
A, B
C
, and C is illustrated in Fig. 2.12 [70]. Even though the ultimate grain size was
achieved that may arise by a consistency of the same ECAP conditions such as pressing
number and the die angel of ECAP, less development of equiaxed grain structure can be
exhibited when the experiments are conducted using routes A and C. in term of a lower
value of the total angular angle. This model for grain refinement was developed
incorporating the major experimental observations. Calculations of the shearing patterns
for different processing routes lead to the conclusion that an equiaxed microstructure is
25

achieved most rapidly in ECAP when slip occurs on three orthogonal planes over a wide
range of angles, as in route B
C
.

Figure 2.11 A model for grain refinement in the central region of the billet in ECAP
[69].
26

Figure 2.12 The appearance of the microstructures on the Y plane after four passes using
routes A, B
C
and C [70].

27

2.4.2 Grain refinement model in HPT processing
Grain refining model in pure Al using HPT was discussed based on the change in
the hardness with imposed strain and a subsequent evolution of microstructure [71].
Since the hardness is related to the dislocation generation, thus, this model was focused
on dislocation behavior and the formation of high-angle boundaries as a function of
equivalent strain. As shown in Fig 2.13, the hardness variation with equivalent strain and
corresponding microstructure evolution was divided into three regions. In the first region
(I), the hardness increases with strain, the accumulation of dislocations and the formation
of subgrain boundaries occur. In the second region (II), the hardness decreases, some
dislocations are annihilated at subgrain boundary leads to increasing of the misorientation
angles. In the third region (III), the hardness remains constant, and there is a balance
between an increase in dislocation and absorption of dislocation at grain boundaries.
A similar model was reported in pure Cu using HPT [55], in Fig.2.14. Unlike
previous model [71], there is no hardness maximum in pure Cu and the steady-state
begins at equivalent strain of ~15 for Cu but ~6 for pure Al. These differences are due to
the dislocation mobility affected by SFE and straining temperature. The hardness
variation, therefore, was divided into two regions. In the first region (I), accumulation of
dislocation and formation of subgrain boundaries are introduced. As strain increases in
the second region (II) there is an increase in the misorientation angles and hardness
becomes constant because of a balance between dislocation generation and
recrystallization.
28

Figure 2.13 Schematic illustration of microstructural evolution with straining along in
pure Al. Thin double lines depicted in region II represent low angle
boundaries with some extension of boundary width and thick lines in region
III represent high angle boundaries [71].

29

Figure 2.14 Schematic illustration of microstructural evolution with straining for grain
refinement in pure Cu [55].

30

2.5 Mechanical properties of materials processed by SPD
2.5.1 Strength and ductility
As mentioned in the preceding section UFG materials prepared by SPD
processing usually have high dislocation densities, non-equilibrium grain boundaries, and
other structural features associated with intense plastic straining. The small grain sizes
and high defect densities give these UFG materials a much higher strength than their
coarse-grained counterparts. It is well known that plastic deformation induced by
conventional forming methods such as rolling, drawing or extrusion can significantly
increase the strength of metals but this increase is usually accompanied by a loss of
ductility. By contrast, some experimental results on materials processed by SPD
demonstrated extraordinary combinations of both high strength and high ductility [72-
74].
For example, experiments were conducted to compare the strength and ductility of
the 3004 aluminum alloy processed by ECAP and cold rolling (CR) [75]. The results as
shown in Fig. 2.15 demonstrated the yield strength for CR and ECAP exhibit a similar
variation with the equivalent strain where the yield strength increased with the increasing
equivalent strains. By contrast, there was the different trend in the ductility behavior for
two processing methods. After equivalent strain ~1, the elongation to failure (or ductility)
of both processing has rapidly decreased. Thereafter, the ductility of an ECAP alloy
maintain constant with increasing equivalent strain but in CR alloy the ductility continues
to decrease, but in a slower rate. In short, the improved ductility might be achieved
31

through ECAP by comparison with cold rolling. This improvement is attributed to the
development of the very fine structure of grain with high-angle grain boundaries of
samples after ECAP.

Figure 2.15 A comparison of yield strength and ductility for a 3004 aluminum alloy
processed by cold rolling or ECAP [75].

32

In addition, the interesting experiments for pure Cu processed using ECAP route
B
C
and for pure Ti processed using HPT [72, 76] were reported. All processes were
performed at room temperature and the resulting engineering stress-strain curve of a pure
Cu and a pure Ti were shown in Fig. 2.16(a) and (b), respectively. The Cu and Ti
specimens were tested in tension at room temperature and at 250 C, respectively. Figure
2.16 demonstrated that the initial coarse-grain Cu and Ti exhibited a typical mechanical
behavior of coarse-grained metal with a low strength but high strain hardening and
ductility. Moreover, a pure Cu with 60 % cold rolling and specimens with small SPD
strains (2 ECAP passes or 1 HPT revolution) had a significant increasing in the strength
but a dramatic reduction in ductility. However, at very large SPD strains (16 ECAP
passes and 5 HPT revolutions) exhibited significantly increased ductility and at the same
time further increase in the strength. This is contrary to the classical mechanical behavior
of metals deformed plastically. Typically, greater plastic deformation by conventional
techniques such as rolling, drawing or extrusion introduces greater strain hardening,
which in turn increases the strength, but decreases the ductility of the metal. The retaining
ductility in samples subjected to high strains SPD could be explained by the strain-rate
sensitivity, m. The strain-rate sensitivity of stress is defined as , where
is flow rate, is strain rate. The samples with high ductility were found to have higher
strain rate sensitivity because higher strain rate sensitivity renders the materials more
resistant to necking [77, 78]. For example, for Cu ECAP 16 passes had m = 0.14 while
Cu ECAP 2 passes had m = 0.06.
33

Figure 2.16 Tensile engineering stress-strain curves of (a) Cu tested at 22 C and (b) Ti
tested at 250 C at a strain rate of 10
-3
s
-1
[72, 76].

34

2.5.2 Superplasticity
The possibility of obtaining UFG metals with grain size of  1 µm has attractive
interests because of the potential for using these materials in the superplastic forming
industry for the production of complex curved parts. Superplastic forming is basically
achieved by using materials under conditions where they exhibit superplastic flow.
Superplasticity is defined as the ability of a material to pull out uniformly to a very high
elongation without the development of any incipient necking. The essential requirements
for achieving a superplastic forming capability are small grain sizes, typically less than
~10 µm, and high forming temperatures, typically above ~0.5T
m
where T
m
is the absolute
melting point of the material [79]. This high temperature is necessary because
superplastic flow is a diffusion-controlled process. Therefore, the thermal stability of
UFG materials becomes significantly important, at least the small grain size should be
reasonably stable at the high temperatures required for diffusion-controlled processes.
Furthermore, an additional requirement for superplasticity is the need to attain a high
volume fraction of grain boundaries having high angles of misorientation so that grain
boundary sliding can be achieved fairly easily at most of the interfaces. The constitutive
equation for conventional superplasticity is given by [38]
n p
G d kT
DG
A















b b
 2.5
where   is the steady strain rate, A is a dimensionless constant, D is the appropriate
diffusivity (lattice or grain boundary), G is the shear modulus, b is the Burgers vector, k
35

is the Boltzmann’s constant, T is the test temperature, d is the grain size, σ is the applied
stress, p is the grain size exponent and n is the exponent of the stress. Typically, in the
superplastic deformation regime p =2, n=2 and D=D
gb
where D
gb
is the coefficient for
grain boundary diffusion. Accordingly, it is anticipated that a decrease in the grain size
by one order of magnitude will lead to an increase in the optimal superplastic forming
rate by approximately two orders of magnitude. It can be shown also that this decrease in
grain size will lead to the advent of a superplastic forming capability at lower
temperatures than those generally associated with conventional superplastic flow. This
reduction in operating temperature is an attractive feature for the superplastic-forming
industry because of the problems associated with tool wear.
Accordingly, it is possible that the UFG material produced by SPD exhibits a
superplastic flow at high strain rate and/or low temperatures, where high strain rates refer
to the tensile testing of samples at rates at and above 10
-2
s
-1
and low temperatures refer
to tensile testing at homologous temperatures below 0.5T
m
. The occurrence of
superplasticity at high strain rates and low temperatures is an attractive feature for
superplastic forming operations because of the consequent reductions in forming time
and tool ware. For example, Valiev et al. [80] first reported the possibility of achieving
low-temperature superplasticity. In their work, an Al-Cu-Zr alloy were subjected to
severe plastic deformation to reduce the grain size to ~0.3µm and they tested the material
in tension to yield an elongation of ~250% at a temperature of 220 C corresponding to
~0.53T
m
. Subsequently, there have been reports of the occurrence of low-temperature
superplastic flow in UFG materials using different SPD techniques including HPT for an
36

Al alloy and Ni
3
Al [81], ECAP for a Mg-based alloy [82-85] and multiple forging for Ti
alloys [86]. Several works also demonstrated the superplastic ductilities at very high
strain rates [87-90]. In addition, the occurrence of superplastic flow after SPD techniques
in wide range of metals and metal alloys after ECAP [91] and HPT [4] was summarized
in earlier reports.

37

3. Experimental Materials and Procedures
3.1 Experimental material
The experiments were conducted using a commercial Cu151 alloy having a
composition, in wt. %, of Cu-0.1%Zr. The alloy was received from Olin Brass (East
Alton, Il, USA) in two types. The first one is in form of a plate with dimensions of 760 ×
250 × 15 mm. The plate then was machined into billets with a diameter of 10 mm and a
length of 70 mm. The other one is in form of a rolled strip with dimensions of 700 × 500
mm and a thickness of 1.5 mm. This strip was machined into disks with diameters of 10
mm and they were polished to final thicknesses of ~0.83 mm.
3.2 Methods of severe plastic deformation
The material was fabricated using severe plastic deformation techniques: ECAP,
HPT and ECAP followed by HPT.
3.2.1 ECAP processing
An ECAP facility is mainly composed of a die having two channels, a plunger
and a hydraulic press machine. The die and a plunger were made of a tool steel and were
heat treated to obtain a Rockwell hardness of ~55. The tolerance of the plunger was kept
extremely low to prevent materials from flowing in between the walls of the channel and
the plunger. Before pressing, the billets and plunger were well lubricated.
38

For processing by ECAP, the billets with a diameter of 10 mm and a length of
70mm were annealed at a temperature of 973 K for 1 h followed by furnace cooling. The
ECAP processing was conducted at room temperature using a solid die in which two
channels intersected at an internal angle of  = 110  and with an additional outer arc of
curvature of  = 20 . These geometries lead to an imposed strain of ~0.8 on each
separate pass through the die [17]. A set of billets was processed by ECAP for 1, 2, 4 and
8 which correspond to the total imposed strain up to ~6.4. All samples were processed
using route B
C
in which the sample is rotated by 90  in the same sense between each
consecutive pass [36]. This processing route was selected because it leads most
expeditiously to an array of equiaxed grains separated by an array of boundaries having
high-angles of misorientation [33-35].
3.2.2 HPT processing
The high pressure torsion was processed on Cu-Zr alloy disks with diameters of
~10 mm and thicknesses of ~0.83 mm. The HPT processing was conducted under quasi-
constrained conditions [92] using a facility consisting of upper and lower anvils having
central depressions with diameters of 10 mm and depths of 0.25mm. These anvils were
made from high-strength YXR3 tool steel and having nitride surfaces. Each disk was
placed in the central depression of the lower anvil. This lower anvil was brought into
position in which a pressure was applied on the disk and torsional strain was then
imposed by rotation of the lower anvil at constant speed of 1 rpm. All processing by HPT
was conducted at room temperature using an applied force of 470 kN which corresponds
39

to an imposed pressure, P, of 6.0 GPa with total strain through N = 1/4, 1/2, 1, 5 and 10
turns. The imposed pressure was determined by using following calculation
The applied force, F, exerted by the hydraulic system of HPT facility is calculated
from:

3.1
where P
h
is the hydraulic pressure, d
h
is the diameter of the hydraulic cylinder (250 mm
in this facility). This force then could be transferred to the pressure applied on the 10-
mm-diameter disk, P
d
using:

3.2

3.2.3 ECAP followed by HPT processing
ECAP billets after pressing through 4 and 8 passes were sliced in the direction
perpendicular to the longitudinal axis using a diamond blade sawing with a thickness of
1.5 and grinded to final thickness of ~0.83 mm. Disks are then subjected to HPT for 5
turns under an applied pressure of 6 GPa at speed of 1 rpm. Therefore, the combination
processing of ECAP + HPT was provided as ECAP 4 passes + HPT 5 turns and ECAP 8
passes + HPT 10 turns.

40

3.3 Microstructural analysis
3.3.1Optical microscope
An optical microscope was used to examine the microstructure of a Cu-Zr alloy
before processing. A conventional metallographic technique was used for the sample
preparation before the microscopy. The samples were molded in plastic resins for
grinding and polishing. The samples were ground with abrasive papers of silicone carbide
from 180 to 4000 grits subsequently to polishing with two ethanol solutions including
Al
2
O
3
powders of 0.3 and 0.05 µm on micro-cloths. The sample surface was finished with
chemical etching with a suitable etching reagent to observe the microstructure. The
chemical composition of the etching reagent for the Cu-Zr alloy was a solution of 1.0 ml
nitric acid with 3.0 ml hydrochloric acid. Linear intercept method was used to measure
grain sizes.
3.3.2 Transmission electron microscopy (TEM)
Specimens from three types of processing were prepared for examination by
transmission electron microscopy (TEM). For ECAP samples, TEM thin foils were
sectioned from the cross section (X plane) of the billets; in case of HPT and ECAP+HPT
samples, TEM foils were prepared from the center and the edge of the disks. TEM
specimens were obtained by mechanically grinding and dimpling from both sides to a
thickness of ~ 10 µm, and further thinning to electro transparency using a Gatan PIP 691
ion milling at a voltage of 4 kV. TEM was performed using a Philips CM12 microscope
41

operating at 100 kV for HPT samples and 120 kV for ECAP and ECAP+HPT samples.
Grain size measurements were conducted using an Olympus analySIS FIVE software
where the average grain size was obtained from at least 200 grains in ~50-100 TEM
images.
3.3.3 Electron-backscatter diffraction (EBSD)
Microstructures of samples after ECAP and HPT processing were also observed
using electron backscatter diffraction (EBSD) technique. The EBSD patterns were
observed on the screen with a charge coupled device (CCD) camera. Due to the
combination of good spatial and angular resolution in this technique, it provides unique
and quantitative information on the local preferred orientation of crystallite and the grain-
boundary misorientations.
For the billet after ECAP, a disk was cut from the central part of the billet in the
transverse plane or X-plane perpendicular to the pressing direction [20]. Then ECAP and
HPT disks were mounted for grinding and polishing. Grinding was started with 240 grit
SiC paper and continued through 1200 grit SiC. Polishing then is used to remove any
deformation introduced during grinding with 9, 6, 3 and 1 micron diamond suspension,
and ending with 0.04 micron colloidal silica suspension using the Vibromet machine. The
EBSD analysis was undertaken using a JSM 7001F scanning electron microscope (SEM)
operating at an accelerating voltage of 15 kV, a working distance of 15 mm, and a sample
tilt of 70 degree. The EBSD patterns were collected with TSL OIM software with the step
sizes varies from 0.04 µm to 0.10 µm, depending on the grain size. The grain boundaries
42

were analyzed using this software in which high-angle grain boundaries (HAGBs) are
defined as misorientations between adjacent measured points of more than 15˚. The
clean-up procedure was performed in an OIM-TSL analyzer and included grain dilation
(GD) and grain confidence index standardization. The total number of modified points
was not higher than 10% of the total points measured in the experiments.
3.4 Mechanical experiments
3.4.1 Microhardness measurements
Microhardness measurements were performed using a digital microhardness tester
FM-1e equipped with a Vickers diamond indenter using a load of 200 gf and a dwell time
of 10 s. The minimum distance between adjacent indentations and distance from
indentation to the edge of the specimen was designed based on ASTM E384 in which at
least 2.5 times the indentation size to be considered in order to avoid interaction between
the work-hardened regions and effects of the edge. The values of the Vickers
microhardness, Hv, were measured through different procedures for different processing;
ECAP and HPT.
3.4.1.1 Microhardness measurement of ECAP
Following ECAP processing, the billets were cut to reveal the cross-sectional
surfaces perpendicular to the pressing direction (X plane) and the longitudinal surface
along the pressing direction (Y plane) and then mounted for subsequent polishing to
mirror-like surfaces. Two different grid patterns of the indentation positions were used
43

for the measurements depending upon the planes as shown in Figure 3.1. First, the
individual values of microhardness were recorded over the surface in the X plane of each
billet with a spacing of 0.5 mm between each separated point. Second, hardness was
measured on a Y plane at positions of 0.5, 1.0, 2.0, 3.0, 4.0, 5.0, 6.0, 7.0, 8.0, 9.0 and 9.5
mm from the top of plane along a vertical direction and the measurements were continued
rectilinearly along X axis parallel to the pressing direction with an equal interval of 1 mm
for total 40 mm.

Figure 3.1 Schematic illustrations showing the position of hardness measurement on the
longitudinal Y plane and cross-sectional X plane.

3.4.1.2 Microhardness measurement of HPT
Following HPT, each processed disk was mounted and polished to have a mirror-
like surface quality and microhardness measurements were taken using two different
procedures as shown in Figure 3.2. First, the average microhardness values were
44

measured along a randomly selected diameter of each disk. These measurements were
taken at intervals of 0.3 mm and at every point the local value of Hv was taken as the
average of four separated hardness values recorded at uniformly separated points
displaced from the selected point by a distance of 0.15 mm. Using this procedure, it was
possible to achieve a high accuracy in the individual values of Hv and for every position
the error bars were estimated corresponding to the 95% confidence limits. Second, the
distribution of local hardness was measured over the total surface of each individual disk
by recording the individual values of Hv following a rectilinear grid pattern with an
incremental spacing of 0.3 mm. These individual datum points were then used to
construct color-coded contour maps to provide a simple visual display of the hardness
distributions within each disk.

Figure 3.2 Grid pattern of the Vickers hardness measurements on the HPT disks.
45

3.4.2 Tensile Measurements
In order to investigate the mechanical behavior at elevated temperature, tensile
experiments were conducted on the samples in the as-received condition, and after
processing. For ECAP samples, the tensile testing were carried out on X plane where
disks with diameter of 10 mm and thickness of ~0.75-0.8 mm, similar to the dimensions
of sample after HPT processing, were cut from the central part of billets. Thereafter,
tensile specimens were prepared from the disks in each condition. The dimension of
tensile specimens is presented in Fig. 3.3. This drawing was designed in order to avoid
any potential problems associated with microstructural inhomogeneities near the centers
of the disks in the early stage of HPT processing. In Fig. 3.3, two miniature tensile
specimens were cut from two off-center positions on either side of the disks using
electro-discharge machining (EDM). Earlier research showed that higher elongations are
generally achieved in specimens machined from the off-center positions in HPT
processing [93].
These tensile specimens were pulled to failure at temperature of 673 K and 723 K
using an Instron testing machine operating at constant rate of cross-head displacement
and with initial strain rates in range from 1×10
-4
to 1×10
-2
s
-1
. At least two specimens
were used for each condition to obtain consistent stress-strain curves.
The elevated temperature was achieved using a single-zone split furnace with a
vertical height of 60 cm and a proportional temperature controller. The temperature was
measured using a K-type thermocouple attached near to the surface of the specimen.
46

Usually it took approximately 30 minutes for the specimen in the furnace to achieve the
desired testing temperature and then held at that temperature for ~10 min in order to
reach thermal equilibrium prior to starting the test.

Figure 3.3 Dimensions of tensile specimens [93].

47

4. Experimental Results
4.1 ECAP experimental results
4.1.1 Microstructural observations after ECAP processing
4.1.1.1 TEM analysis for ECAP
Figure 4.1 presents transmission electron micrographs of the Cu-0.1% Zr after
ECAP processes in different passes. After one pass (Fig. 4.1 (a)), the microstructure
mainly consists of elongated dislocation cells or subgrains with a high density of
dislocations in their interior. These cells have average spacing of ~350 nm. Most of
boundaries are thick and wavy which attributed to their low-angle misorientation. After
two passes (Fig. 4.1 (b)), the microstructure has similar features to that of one pass. The
dominant structure is still characterized by elongated subgrains with majority of low-
angle misorienation and high density of dislocations. However, in some areas the
microstructure has evolved into equiaxed subgrain structures indicating the new slip
systems operated during the second pass. The microstructure of the specimen after four
passes is shown in Fig. 4.1(c). In this stage of deformation, microstructure has significant
change where grain structures become more equiaxed. These equiaxed grains are formed
when the elongated structures are subsequently sheared in other directions. Boundaries
are also better-defined as thinner and sharper possibly due to the increase of their
misorientation. However several grain interiors still contain a high density of
dislocations. The introduction of additional strain through eight passes in Fig. 4.1(d)
48

leads to higher grain refinement and more equiaxed grain structure with high-angle grain
boundaries.

Figure 4.1 TEM images of the disks processed by ECAP for (a) one pass, (b) two passes,
(c) four passes and (d) eight passes.

49

Figure 4.2 Histograms of the grain size distributions measured at (a) four passes and (b)
eight passes: the individual average grain sizes are indicated.

Figure 4.2 (a) and (b) show the statistical grain size distributions from TEM
images recorded in samples subjected to ECAP four passes and eight passes, respectively.
The measurements were taken from approximately 300 individual grains in each
distribution. It should be noted that in ECAP one pass and two passes the microstructures
were very poorly defined with a very high density of dislocation such that it was difficult
to take quantitative measurement. It can be seen from Fig. 4.2 that the distribution of
50

grain size for a Cu-Zr alloy measured using TEM was fitted by lognormal distribution
function. The grain sizes range between 200 and 800 nm after four passes and from 150
to 800 nm after eight passes. However, the distributions show different peaks leading to
different average grain size. The measurements gave an average grain size of ~440 nm
after pressing through four passes and this is larger than the average grain size of ~310
nm measured from eight passes.
4.1.1.2 EBSD analysis for ECAP
Crystal orientation images using EBSD were conducted for the ECAP samples of
4 passes and 8 passes as shown in Fig. 4.3 where the individual colors correspond to
different orientations of each grain described as a unit triangle. In this OIM image, and in
all subsequent OIM images, the low-angle grain boundaries defined by misorientation
angle of 2-15 are depicted by yellow lines and high-angle grain boundaries having
misorientation > 15  are denoted by black lines. Inspection of Fig. 4.3(a) reveals the
inhomogeneous microstructure of a sample deformed through 4 passes. Grain structure
consists of both original elongated shape grains and equiaxed grains. Many grains are
surrounded by low-angle grain boundaries. For this condition, and average grain size of
~410 nm was examined. This is consistent with the grain size reported by TEM as
discussed above. After eight passes in Fig. 4.3(b), the microstructure appears to be more
homogeneous and more equiaxed grains. Furthermore, the microstructure is more refined
where the average grain size is reported as ~350 nm. This is consistent with a report of a
gain size of ~400 nm, measured using TEM, for a Cu-0.18%Zr alloy processed by ECAP
51

for 6 passes at room temperature using a die with internals angle of  = 90  and  = 20 
[11]. Figure 4.4 presents the distributions of misorientation corresponding to the OIM
images of samples after ECAP (a) four passes and (b) eight passes. The data analysis
shows that the fraction of high-angle grain boundaries were measured as 32% after 4
passes and the higher fracture as 47% was observed from a sample subjected to 8 passes.
52

Figure 4.3 EBSD orientation images of the Cu-0.1% Zr disks processed by ECAP for (a)
four passes and (b) eight passes

53

Figure 4.4 Distributions of the number fraction of the misorientation angles of the grain
boundaries at the centers and edges of the disks processed by ECAP through
(a) four passes and (b) eight passes.

54

4.1.2 Microhardness measurements after ECAP processing
4.1.2.1 Hardness values on the X plane
After processing by ECAP through one, two, four and eight passes, the individual
values of Hv were recorded on the X plane to evaluate the development of hardness
homogeneity with increasing numbers of ECAP passes. These values are plotted in Fig.
4.5 in the form of color-coded contour maps after ECAP through (a) 1, (b) 2, (c) 4 and (d)
8 passes and they provide a visual display of the distributions of the local hardness across
the X planes. For all maps, the individual values of Hv are represented by a set of unique
colors corresponding to hardness values in the range from 100 to 160 in incremental steps
of 10: the significance of the different colors is denoted by the color key on the right in
Fig. 4.5. It should be noted that the Y and Z axes in Fig. 4.5 represent directions
perpendicular to the Y and Z planes of the ECAP billets, respectively, and the points (Y,
Z) = (0, 0) in the centers of each map correspond to the centers of the billets on the X
plane.
It is apparent from Fig. 4.5 that the microhardness values increase over the
surfaces of the X planes with increasing numbers of ECAP passes and the hardness
values become reasonably saturated at Hv  140 after 8 passes. Detailed inspection of
Fig. 4.5(a) shows there is an inhomogeneity in the hardness values in the lower area of
the X plane after a single pass where the values of hardness are Hv  110 and this
contrasts with an average value over the X plane of Hv  125. This result is similar to
data reported earlier for pure Al [94] and an Al-6061 alloy [69] processed by ECAP.
55

Thereafter, the average value of Hv gradually increases to ~135 after two passes and the
area having lower hardness at the lower edge of the billet is reduced to a width of ~1 mm
from the bottom surface as shown in Fig. 4.5(b). This area of lower hardness is further
reduced after four passes in Fig. 4.5(c) and it disappears to give excellent homogeneity
after eight passes in Fig. 4.5(d).
To clearly understand the Vickers hardness evolution and to evaluate the extent of
homogeneity quantitatively, histograms showing the number fraction of measurement
points were constructed for the samples subjected to ECAP with various numbers of
passes with the incremental segment of 5 Hv. These histograms of hardness distributions
are illustrated in Fig. 4.6 where the solid vertical lines represent the average of hardness
values in each pass. It is apparent that the average hardness values shift from left to right
as the number of passes increases. This means material exhibits more resistance or
hardness when subjected to higher strains (number of passes). Detailed inspection also
shows that the rate of increment of an average hardness value is obvious from 1 pass to 2
passes, thereafter, this rate slightly changes. The important point that can be concluded
from these histograms is the data distribution in each sample. It is obvious that the width
of the distribution tends to be narrower with increasing strain imposed. This
demonstrated the development of hardness homogeneity in a Cu-Zr alloy with increasing
the number of passes. In present measurement the best homogeneity achieved in a sample
subjected to 8 passes where a largest number fraction of ~80% lies into the range of
Vickers hardness of 140 to 145.
56

Figure 4.5 Color-coded contour maps showing the distributions of the Vickers
microhardness on the cross-sectional planes after ECAP for (a) one pass, (b)
two passes, (c) four passes and (d) eight passes.

57

Figure 4.6 Histograms of the Vickers Hardness (Hv) distribution of Cu-0.1%Zr after
ECAP (a) 1 pass, (b) 2 passes, (c) 4 passes and (d) 8 passes on cross-sectional
plane.

58

4.1.2.2 Hardness values on the Y plane
Figure 4.7 shows color-coded contour maps for the microhardness values
recorded over the Y planes of the samples after ECAP through (a) 1, (b) 2, (c) 4 and (d) 8
passes: as for the X planes in Fig. 4.5, the values of Hv are plotted in the range from 100
to 160 in incremental steps of 10 with the significance of the colors denoted by the key
on the right. These maps were constructed with the vertical axis in the Z direction where
Z = 0 denotes the center line of each billet along the X axis and the horizontal axis in the
X direction where X = 0 and 40 mm denote the rear and front positions of each billet as
illustrated in Fig. 3.1, respectively.
It is apparent from Fig. 4.7 that the microhardness values increase with increasing
numbers of passes up to 4 passes in Fig. 4.7(c) and the hardness is essentially saturated
after 8 passes in Fig. 4.7(d). In practice, there is a high degree of hardness inhomogeneity
after 1 pass in Fig. 4.7(a) and this continues through 4 passes in Fig. 4.7(c) whereas after
8 passes in Fig. 4.7(d) the sample shows a high level of homogeneity throughout the Y
plane. The lower hardness values near the lower edge of the billet in Fig. 4.7(a) are
consistent with the hardness data for the X plane in Fig. 4.5(a).

59

Figure 4.7 Color-coded contour maps showing the distributions of the Vickers
microhardness on the longitudinal planes after ECAP for (a) one pass, (b) two
passes, (c) four passes and (d) eight passes.
60

Figure 4.8 The histogram of the Vickers hardness (Hv) distribution of Cu-0.1%Zr after
ECAP (a) 1 pass, (b) 2 passes, (c) 4 passes and (d) 8 passes on longitudinal
plane.

61

Figure 4.8 presents the histogram of the Vickers hardness distributions of Cu-
0.1% Zr after ECAP through 1 pass to 8 passes as shown in Fig. 4.8(a)-(d), respectively,
where the solid line in each plot denotes the average hardness value. This plot again
provides valuable information on the evolution of hardness homogeneity. Inspection of
Fig. 4.8 leads to important conclusions. First, the average hardness value (solid line)
increases with imposed straining. Second, the narrower distributions of hardness are
observed on samples subjected to higher strains. These demonstrate the extent of
hardness homogeneity in the sample subjected to sufficient high strains. It is worth noting
that the histograms of all samples in longitudinal plane are similar to those of cross-
sectional plane.
In order to provide a more comprehensive display of the development of
homogeneity on the Y planes, an additional evaluation of the Hv values was performed
by plotting the individual microhardness values recorded along the pressing direction at
1.0 mm from both the top and bottom surfaces of the Y planes. These plots are shown in
Fig. 4.9 for the samples after (a) 1, (b) 2, (c) 4 and (d) 8 passes where the lower broken
line denotes the annealed condition without pressing and the two upper lines denote the
average values of Hv for these two positions, respectively.
Close inspection of Fig. 4.9 leads to several conclusions. First, the value of Hv
increases significantly in the first pass from a value of Hv  60 for the annealed condition
without processing to Hv  127 after 1 pass of ECAP. Thereafter, the increase is more
gradual to Hv  135 after 2 passes, ~140 after 4 passes and ~143 after 8 passes. Second,
62

Figure 4.9 Individual values of Hv with regression lines recorded on the longitudinal
planes at 1.0 mm from the upper and lower surfaces for the billets pressed
through (a) one pass, (b) two passes, (c) four passes and (d) eight passes.

there is a marked difference in the hardness values between the top and bottom regions of
the Y plane after 1 pass as is evident in Fig. 4.9(a). With increasing numbers of ECAP
passes, this difference in the hardness values between the upper and lower positions
becomes smaller and ultimately the top and bottom values are essentially identical after 8
passes as shown in Fig. 4.9(d). These results demonstrate the development of an
excellent hardness homogeneity in the Z direction on the Y plane of the processed
samples with increasing numbers of passes. Third, the individual Hv values are relatively
63

stable at the measurement points along longitudinal traverses after 2, 4 and 8 passes.
Therefore, when combined with the information in Figs 4.7- 4.9, it is concluded that there
is also excellent hardness homogeneity in the longitudinal or X direction on the Y plane.
4.1.3 Tensile properties of ECAP samples
The stress-strain curves are shown in Fig. 4.10 from tensile testing conducted at
temperature of 673 K at initial strain rates of (a) 1.0 × 10
-2
(b) 1.0 × 10
-3
and (c) 1.0×10
-4

s
-1
: each curve was plotted for the annealed condition and for specimens processed
through four passes and eight passes ECAP. At a strain rate of 1.0 × 10
-2
s
-1
in Fig.
4.10(a) the material after four passes exhibits higher yield stress but lower ductility
compared with annealed conditions. Thereafter, with increasing the number of passes to 8
passes the ductility increases and yield stress decreases, but still higher than an annealed
sample. A similar tendency was observed from different initial strain rates as shown in
Fig. 4.10(b) and (c). Detail inspection shows the material resistance decreases with
decreasing the tension strain rate. The effect of strain rate on elongation to failure of a
Cu-Zr alloy was presented in Fig. 4.11. The results show typical tensile properties where
the ductility of materials decreases with increasing the strain rates. In this experiment the
highest elongation to failure of an ECAP sample was measured as~100% at a strain rate
of 1.0 × 10
-4
s
-1
when processed through eight passes at 673 K.
64

Figure 4.10 Plot of engineering stress versus engineering strain for the Cu-0.1%Zr alloy
in an annealed specimen and after ECAP processing through 4 passes and 8
passes at testing temperature of 673 K with different strain rates.
65

Figure 4.11 Elongation to failure versus strain rate for the Cu-0.1%Zr alloy in an
annealed condition and after ECAP processing through four and eight passes
at testing temperature of 673 K

66

4.2 HPT experimental results
4.2.1 Microstructural observation after HPT processing
4.2.1.1 TEM analysis for HPT
The evolution of microstructure of the Cu-0.1% Zr alloy subjected to HPT
processing was examined using TEM. Representative microstructures examined by TEM
are shown in Fig. 4.12 for (a) 1/4 turn, (b) 1 turn (c) 5 turns and (d) 10 turns, where the
left column corresponds to the centers of the disks and the right column corresponds to
the peripheries of the disks, respectively.
Inspection shows the microstructures in the centers of the disks are similar after
1/4 and 1 turn in the left column of Fig. 4.12(a) and (b) where there are indistinct
(sub)grain boundaries but also a clear dislocation cell structure. Large numbers of
dislocations were observed within the grains for these two samples. At the edge of the
disk after 1/4 turns, dislocation cell structures are still dominant mixed with equiaxed
grains with sharp grain boundaries and high density of dislocations. In this stage of
deformation (sub)grains in the edge region are more refined than those in central region.
Continuous straining through 1 turn produced reasonably equiaxed grains with sharp
boundaries. Some larger grains still remained at the edge of the disk but the density of
dislocations in grains is reduced as shown in the right column in Fig. 4.12(b). After five
turns in Fig. 4.12(c) the microstructure becomes uniformly finer with equiaxed grains.
The grain boundaries become well-defined with sharp and straight grain boundaries in
67

both the central and peripheral regions of the disk. The microstructure after ten turns is
similar to five turns consisting of a reasonably homogeneous microstructure of equiaxed
grains. The SEAD patterns contain well-defined rings demonstrating the presence of
boundaries with high angles of misorientation.
Figure 4.13 shows histograms of the grain size distributions of Cu-0.1%Zr
samples measured from different number of turns at central and edge positions. In the
centers of the disks after 1/4 and 1 turn, the grains were poorly defined and it was not
possible to take quantitative measurements. At the edges of the disks after 1/4 and 1 turn,
individual grain sizes were recorded over a wide range from ~100 to ~800 nm and the
average grain sizes were measured as ~350-360 nm. Despite the similarity in these two
average sizes, there were different peak fractions of ~200 and ~300 nm for the disks
processed through 1/4 and 1 turn, respectively, as shown in Fig. 4.13 (a) and (b). This
difference arises because of the larger numbers of dislocations observed at most of the
grain boundaries at the edge of the disk after 1/4 turn as shown in the left column in Fig.
4.10 (a). After 5 turns of HPT, the range of grain sizes was reduced to 80-670 nm with a
peak fraction at ~250 nm in the central region and 40-460 nm with a peak fraction at
~100 nm in the peripheral region as shown in Fig. 6 (c) and (d). The reduction of range
indicates a more homogeneous microstructure. A smaller average grain size of ~180 nm
was recorded at the edge of the disk compared with an average size of ~310 nm at the
center of the disk after 5 turns. The grain size distribution after 10 turns is similar to that
of 5 turns. The average grain size measured as ~290 nm at the center and ~250 nm at the
edge.
68

Figure 4.12 continued on next page.

69

Figure 4.12 TEM images of the disks processed by HPT for (a) 1/4 turn, (b) 1 turn, (c) 5
turns and (d) 10 turns: the left column corresponds to the centers of the disks
and the right column corresponds to the edges of the disks
70

Figure 4.13 Histograms of the grain size distributions measured at (a) a peripheral region
of the disk after 1/4 turn, (b) a peripheral region of the disk after 1 turn, (c) a
central region of the disk after 5 turns (d) a peripheral region of the disk
after 5 turns (e) a central region of the disk after 10 turns (f) a peripheral
region of the disk after 10 turns: the individual average grain sizes are
indicated.
71

4.2.1.2 EBSD analysis for HPT
Figure 4.14 shows representative OIM images taken after (a) 1/4 turn (b) 5 turns
and (c) 10 turns, where the left column corresponds to the centers of the disks and the
right column corresponds to the peripheries: the individual colors correspond to different
orientations of each grain as depicted in the unit triangle. Figure 4.14(a) shows the
original coarse grains are visible in the central region after 1/4 turn. There are different
orientations within the grain and high fraction of low-angle misorientation due to
dislocation activities [95]. By contrast, grains are significantly refined after even a quarter
turn at the edge region but some large grains are still visible as shown in Fig. 4.14(a):
right column. After five turns, the microstructure at the center is significantly refined; the
original coarse grains are invisible as shown in left hand side of Fig. 4.14(b). By contrast,
at the edge region the grain refinement remains the same but an equiaxed grain structure
is more uniform over the scanned area. This can be confirmed by the grain size
distribution recorded by TEM where the range of grain sizes of a sample subjected to five
turns is narrower by comparison with the sample taken through a quarter turn as
presented in Fig. 4.13(a) and (d). Measurements gave average grain sizes of ~290 nm and
~240 nm at the center and edge of the disk processed through five turns. The introduction
of additional torsional straining to 10 turns leads to an insignificant change in grain
refinement. The microstructures of this condition in both regions are now essentially
homogeneous and the fine grains are reasonably equiaxed as shown in Fig. 4.14(c). The
average grain sizes were measured as ~270 nm and ~230 nm at center and edge of the
disk.
72

Figure 4.14 continued on next page.
73

Figure 4.14 EBSD orientation images of the Cu-0.1% Zr disks processed by HPT for (a)
1/4 turn, (b) 5 turns and (c) 10 turns.

Figure 4.15 shows the distributions of the misorientation angles of grain
boundaries at the centers and edges of the disks processed by HPT through (a) 1/4 turn
(b) 5 turns and (c) 10 turns: information is given also for the total percentage of high-
angle boundaries for each condition where the solid curves in each plot corresponding to
the statistical prediction for a set of random miorientations. It is apparent that the sample
processed through 1/4 turn exhibits a relatively low fraction (17%) of HAGBs in the
center of the disk and this fraction is significantly higher (51%) in the edge region. After
74

five and ten turns , as shown in Fig. 4.15(a) and (b), the distributions of the number
fraction of the misorientation angles have a similar tendency. Detailed inspection shows
the number fraction of high-angle grain boundaries is increased at the center and the peak
of low-angle is reduced significantly by comparison with those of a quarter turn whereas
this fraction of HAGBs is fairly increased.

Figure 4.15 Distributions of the number fraction of the misorientation angles of the grain
boundaries at the centers and edges of the disks processed by HPT through
(a) a quarter turn (b) five turns and (c) ten turns.
75

4.2.2 Microhardness measurements after HPT processing
The values of the Vickers microhardness, Hv are shown in Fig. 4.16 plotted as a
function of the position on the each disk after HPT processing: the lower dashed line
denotes the hardness value of Hv  95 in the as-received condition prior to processing.
These measurements demonstrate that the hardness at the edge of the disk after 1/4 turn
increases significantly with a value of Hv  160. By contrast, the hardness in the center of
the disk increases only from Hv  95 to Hv  110. Thereafter, there is little or no change
in the hardness at the edge of the disk even up to N = 10 turns but the hardness values
recorded in the centers of the disks gradually increase with increasing numbers of turns
and after five and ten turns there is essentially a hardness uniform with Hv  160 across
the diameter of the disk. This gradual development of hardness homogeneity is consistent
both with other experimental reports for materials processed by HPT [49, 55, 96-101] and
with strain gradient plasticity modeling of HPT process [102].
Detailed information on the values of Hv and the associated error bars are
summarized in Table 4.1 where the error bars denote the 95% confidence limits based on
the separate measurements records around each point. It is apparent from Table 4.1 that
the error bars tend to be higher in the central region of the disk where initially there are
significant inhomogeneities. However, the error bars in the central region decrease with
increasing torsional straining so that ultimately, after ten turns, the error bars are
reasonably consistent across the disk diameter.
76

Figure 4.16 Values of the Vickers microhardness versus distance from the center of the
Cu-0.1%Zr disks after HPT processing for various numbers of turns.

Using the microhardness values recorded following a rectilinear grid pattern,
color-coded contour maps were constructed for N = 1/4, 1, 5 and 10 turns as shown in
Fig. 4.17 where the X and Y labels denote two arbitrary and randomly selected
orthogonal axes on the HPT disks such that the center of each disk has a coordinate of
(0,0). For all maps, the individual values of Hv are presented as a set of unique colors that
cover hardness values from 110 to 170 in incremental steps of 10: the significance of
these colors is denoted by the color key on the right in Fig. 4.15.
77

In the very early stages of HPT processing though N = 1/4 turn in Fig. 4.17(a),
there is a relatively large area of lower hardness confined within the central region of the
disk and extending outwards through a diameter of ~ 4 mm. For this condition, the
hardness values reveal significant inhomogeneity. Increasing the straining to N = 1 turn
decreases the diameter of the central region of lower hardness to ~2 mm as shown in Fig.
4.17(b) and this central region is essentially removed after HPT through 5 and 10 turns as
shown in Fig. 4.17(c) and (d). Since the disks processed through 5 and 10 turns show
similar hardness distributions over the entire disk surfaces, these hardness measurements
suggest that a reasonably homogeneous hardness may be attained in the Cu-0.1% Zr alloy
by HPT processing at room temperature through N = 5 turns under a pressure of P = 6.0
GPa.
Table 4.1
Microhardness values along the radii of Cu-0.1%Zr disks processed by HPT.

Distance
from
center
(mm)
Hv
no HPT 1/4 turn 1/2 turn 1 turn 3 turns 5 turns 10 turns
0.0 95.4  1.7 116.8  4.7 124.6  5.1 131.0  8.0 143.1  4.8 156.2  4.5 157.0 ± 1.8
0.3 95.5  1.1 127.3  5.0 129.7  4.4 137.4  4.9 153.4  3.7 156.2  2.4 155.1 ± 2.4
0.6 95.5  1.4 133.8  4.2 136.2  2.0 145.1  3.3 156.8  2.4 155.0  2.1 155.9 ± 2.0
0.9 95.6  0.9 139.9  2.7 140.6  2.0 148.2  2.7 156.2  1.2 154.8  2.5 158.1 ± 1.9
1.2 96.1  1.2 148.0  3.0 144.3  1.4 154.0  1.6 156.9  0.4 155.6  2.0 159.4 ± 1.3
1.5 96.3  1.0 151.4  1.9 147.5  2.5 157.1  1.9 156.4  1.2 155.6  1.1 158.3 ± 1.2
1.8 94.8  1.1 152.0  2.3 149.8  1.2 160.1  1.5 157.0  2.3 157.9  1.3 158.4 ± 1.7
2.1 95.5  0.5 152.9  1.5 150.5  0.8 161.3  1.7 156.4  1.0 159.0  1.0 160.0 ± 1.3
2.4 95.8  1.0 156.5  2.0 153.1  1.8 161.2  1.5 156.7  0.8 159.0  1.3 160.5 ± 2.0
2.7 94.7  1.0 159.1  1.1 155.1  1.3 160.5  1.7 157.4  1.3 160.8  1.0 161.9 ± 1.3
3.0 95.3  0.8 158.3  1.7 155.2  0.5 160.5  1.1 158.0  1.1 161.0  1.6 163.3 ± 0.9
3.3 95.9  0.8 159.1  2.5 156.4  1.1 161.0  1.6 159.3  0.8 160.5  2.5 163.8 ± 0.6
3.6 96.5  0.9 161.4  2.7 158.6  1.5 159.9  2.0 160.9  1.6 161.4  1.3 163.4 ± 1.3
3.9 96.5  0.9 162.2  2.2 160.0  2.6 157.9  2.2 162.1  1.4 162.9  1.6 164.4 ± 0.5
4.2 96.1  1.4 163.8  2.0 159.5  3.3 157.4  1.2 163.1  1.0 163.2  1.4 163.8 ± 1.2
4.5 95.9  1.1 164.0  2.5 159.3  2.6 159.1  1.6 163.3  1.5 162.7  2.3 163.6 ± 1.1
78

Figure 4.17 Color-coded contour maps showing the distributions of the Vickers
microhardness values over disks processed by HPT at a present of 6.0 GPa
for (a) 1/4 turn, (b) 1 turn, (c) 5 turns and (d) 10 turns.

79

4.2.3 Tensile properties after HPT processing
Figure 4.18 shows representative stress-strain curves for the specimens after HPT
processing through 1, 5 and 10 turns and in the as-received condition without HPT: these
experiments were conducted using an initial strain rate of 1.0 × 10
-3
s
-1
at temperatures
of (a) 673 and (b) 723 K. The plots show the Cu-Zr alloy exhibits typical high
temperature behavior including little strain hardening and reasonable ductility. Detailed
inspection of Fig. 4.18 shows the maximum stresses recorded after HPT for 1 turn are
slightly lower than for the as-received material where this is attributed to the
microstructural inhomogeneities produced in the early stages of torsional straining.
Nevertheless, the yield strengths increase with increasing numbers of turns with the
highest values recorded after 5 and 10 turns. This is consistent with data reported for
several materials processed by HPT including Fe [103] and an Al-7034 alloy [104]. It is
also apparent that with an increment of testing temperature the strength tends to decrease
while the ductility increases for all conditions. In addition, the influence of initial strain
rate on tensile properties is presented in Fig. 4.19. The stress-strain curves in Fig. 4.19(a)
and (b) exhibit typical tensile properties where the tensile strength increases with strain
rate. However, these curves give an interesting point that there is identical tensile
properties of specimens processed through five turns and ten turns as shown in Fig.
4.17(a) and (b), respectively.
The results in Fig. 4.18 demonstrate that, at a strain rate of 1.0 × 10
-3
s
-1
, the
maximum elongations achieved in tension are ~130% at 673 K and ~160% at 723 K.
80

However, to provide a more comprehensive display of tensile data over a range of strain
rates, Fig. 4.20 plots the measured elongations to failure versus strain rate for samples in
the as-received condition and after processing through 1, 5 and 10 turns at (a) 673 and (b)
723 K. The results in Fig. 4.20 cover initial strain rates from 1.0 × 10
-4
to 1.0 × 10
-2
s
-1
,
respectively.
Inspection of Fig.4.20 leads to three important conclusions. First, the ductility
increases significantly after HPT for 5 and 10 turns whereas the samples processed
through only one turn show similar elongations to specimens without HPT. This lower
ductility in the samples after HPT for 1 turn is a direct consequence of the high degree of
microstructural inhomogeneity at this early stage of HPT processing. Second, after 5 and
10 turns the measured elongations to failure consistently increase with decreasing strain
rate and both processing conditions show almost identical results. The similarity between
the samples processed through 5 and 10 turns matches the hardness measurements shown
in Fig. 4.20 where essentially identical hardness values are recorded after processing
through 5 and 10 turns. Third, the highest elongations recorded in these experiments
were ~240% for the two specimens processed through 5 and 10 turns when testing at 723
K using a strain rate of 1.0 × 10
-4
s
-1
. This elongation does not fulfill the requirements for
true superplastic flow where the measured elongation in tension should be at least 400%

81

Figure 4.18 Plots of engineering stress versus engineering strain for the Cu-0.1% Zr
alloy in an as-received specimen and after HPT processing through one, five
and ten turns at an initial strain rate of 1.0×10
-3
s
-1
and at temperatures of (a)
673 K and (b) 723 K.

82

Figure 4.19 Plots of engineering stress versus engineering strain for the Cu-0.1% Zr
alloy after (a) five turns and (b) ten turns at temperature of 673 K with
various initial strain rates.
83

Figure 4.20 Elongation to failure versus strain rate for the Cu-0.1% Zr alloy in an as-
received condition and after HPT processing through one, five and ten turns
and testing at temperatures of (a) 673 K and (b) 723 K.

84

4.3 ECAP + HPT experimental results
4.3.1 Microhardness measurements after ECAP + HPT processing
The microhardness distributions on the cross-sectional plane for Cu-0.1%Zr after
combination processing are shown in Fig. 4.21 in the form of color-coded contour maps
and histogram of hardness distribution. In the color-coded contour maps the individual
values of Hv are represented by a set of unique colors corresponding to hardness values
in the rage from 110 to 170 in incremental steps of 10 Hv whereas the incremental steps
of 5 Hv were used to construct the histograms. It is apparent that the contour maps of
ECAP 4p + HPT 5 turns in Fig. 4.21(a) and ECAP 8p + HPT 5 turns in Fig. 4.1(c) are
essentially similar to each other. Detailed inspection of these maps show reasonably
homogeneous hardness distributions over the disks surface. This hardness homogeneity
may be confirmed by the corresponding histograms of hardness distribution as shown in
Fig.4.21 (b) and (d). The histogram information from both samples shows an excellent
uniform with high number fraction of ~60% in the peak in each condition. The average
values of Hv  157 and Hv 156 were measured in ECAP 4p + HPT 5 turns and ECAP 8p
+ HPT 5 turns samples, respectively
85

Figure 4.21 Color-coded contour maps and histograms of Vickers microhardness
distribution over disks processed by ECAP (a), (b) 4p + 5 turns and (c), (d)
8p + 5 turns.

86

4.3.2 Microstructural observations after ECAP + HPT processing
Figure 4.22 shows representative TEM images recorded (a) at the center of the
disk and (b) near the edge of the disk for a sample subjected to ECAP 4 passed followed
by HPT 5 turns. The microstructure in a central region is similar to that in an edge region
that mainly consists of reasonably equiaxed grains. Both sharp and wavy grain
boundaries can be observed in this sample. The corresponding histogram of grain size
distributions is illustrated in Fig. 4.23. Measurements gave grain sizes of ~300 nm in the
center region and ~290 nm near the edge.
With the same technique the microstructures after ECAP 8 passes followed by
HPT 5 turns were examined as shown in Fig. 4.24 for (a) center area and (b) edge area.
Again there is insignificant different between microstructure in center and edge regions
of the disk. The microstructure presents homogeneous equiaxed grains. The grain
boundaries appear to be better defined, more straight and shaper in this condition
compared to the former combination procedure as shown in Fig. 4.22. The individual
measurements of the grain sizes of the disk, as determined from TEM micrographs, are
presented as histograms in Fig. 4.25. The grain size distributions in both areas are almost
identical with measured average grain size of ~280 nm at center and ~270 nm at edge.

87

Figure 4.22 TEM figures of disks processed by ECAP 4 passes followed by HPT 5 turns
at (a) center and (b) edge of the disks.
88

Figure 4.23 Histograms of the grain size distributions at the centers and edges of the
disks processed by ECAP 4p + HPT 5 turns.

89

Figure 4.24 TEM figures of disks processed by ECAP 8 passes followed by HPT 5 turns
at (a) center and (b) edge of the disks.
90

Figure 4.25 Histograms of the grain size distributions at the centers and edges of the
disks processed by ECAP 8p + HPT 5 turns.

91

5. Discussion
5.1 ECAP discussion
5.1.1 Significance of the ECAP processing route
The present experiments were conducted using ECAP processing route B
C
where
the sample is rotated by 90  in the same sense between each pass. Other processing
routes commonly used in ECAP are route A where there is no rotation between passes,
route B
A
where samples are rotated by 90  in alternate directions between passes and
route C where samples are rotated by 180  between passes [36]. It is important,
therefore, to examine whether route B
C
represents the optimum condition.
The conclusion that route B
C
is the optimum route for most rapidly achieving an
array of equiaxed ultrafine grains was first put forward many years ago based on
observations using transmission electron microscopy (TEM) of high-purity aluminum
processed using an ECAP die with a channel angle of  = 90  [32, 34, 35].
Subsequently, in experiments on two aluminum alloys using a die with  = 120 , it was
proposed that processing using route A was more efficient than processing using route B
C

[105]. However, a later detailed analysis of these latter results showed that processing by
route A, although it led to a higher fraction of high-angle boundaries, produced grains
that were elongated rather than equiaxed [37].
92

More recently, there have been some attempts to clarify this issue through
extensive use of EBSD. In experiments on a Cu-8% Ag alloy with an ECAP die having
 = 90  it was concluded that route A was preferable rather than route B
C
because it
produced a higher fraction of high-angle boundaries and a smaller grain size [106].
However, the processing in these experiments was discontinued after 4 passes and it is
well established from other experiments that 4 passes are not sufficient to establish an
equilibrium structure [107, 108] Accordingly, more reliance can be placed on
experimental data where oxygen-free Cu was processed for up to 25 passes using a die
with a channel angle of  = 90  [38]. In these experiments it was found that route A
produced the highest fraction of high-angle boundaries after 25 passes but route B
C
gave
the smallest size for the structural elements. In addition, route B
C
gave the maximum
yield stress, the maximum ultimate tensile stress and the highest ductility after 10 passes
or more. Furthermore, route B
C
was most effective in producing equiaxed grains. Since
an equiaxed grain structure is an important requirement for achieving a superplastic
forming capability at elevated temperatures, these very recent results provide strong
support for using route B
C
as the optimum procedure.
Although most attempt to characterize the processing routes have been based on
the use of TEM or EBSD, there was an earlier experiment where samples of an Al-Mg-Sc
alloy were processed by ECAP with  = 90  using different routes and then the samples
were pulled in tension at a high temperature to check on the ability to produce
superplastic elongations [109]. These results also confirmed that a maximum
superplastic ductility was achieved when using route B
C
.
93

Finally, it is necessary to consider the reason for the optimum results achieved
when using route B
C
. An important characteristic of this route is that it produces the
largest angular ranges for slip on each of the three orthogonal planes in the ECAP billet
[70]. Additional advantages are that route B
C
restores the structural elements of the
initial grains every 4n passes where n is an integer and, in addition, the deformation
occurs on all three planes [36, 37].
5.1.2 Microstructural evolution and grain refinement during ECAP
processing
The TEM and EBSD analysis show that the microstructural features for a Cu-
0.1%Zr alloy subjected to 1-8 passes have a transition from elongated cell structures to
equiaxed (sub)grain structure. These analyses also inform a significant evolution of the
boundary misorientation during ECAP. Analysis by TEM shows that the grain boundaries
evolved from thick and wavy boundaries towards thinner and sharper as shown in the
Fig. 4.1. Furthermore, the qualitative investigation of the misorientation distributions in
EBSD reveals that the degree of misorientation among subgrains/grains increases with
increasing the number of passes. Dalla Torre et al. [110] reported that during the
evolution of the microstructure as the number of passes increases, the material is not only
subjected to a change in the misorientation relationship among cells or subgrains, but also
to a change from broader cell walls to narrower ones, which lead to equilibrium
boundaries as misorientation increases with strain. In their study the sharp boundaries
observed from TEM were referred as equilibrium boundaries while Valiev et al. [111],
94

defined the non-equilibrium boundaries as distorted grain boundaries that contain a high
density of extrinsic dislocations which are not needed to accommodate the misorientation
across the grain boundary. Similar results of the microstructural evolution during ECAP
were presented for Cu and Cu alloy [106, 112] and Al and Al alloy [108].
It is well-established that metals processed through ECAP processing lead to
significant grain refinement. The microstructure observation in this study presents an
attractive grain refinement in a Cu-0.1% Zr alloy after processed through eight passes of
ECAP. The grains were refined from ~30 µm in an annealed sample to ~310 nm in the
eight passes sample as shown in Fig. 4.2. Generally, grain refinement mechanism during
ECAP is controlled by various dislocation activities in metal with medium and high
stacking fault energy (SFE), e.g. Al, Cu, Fe, Ni. [33, 34, 110, 113-116]. A high density of
dislocations is introduced in materials subjected to severe plastic deformation such as an
ECAP method. The strong dislocation activities in the early stage of deformation results
in the formation of dislocation cell structures consisting of thick cell walls and low-
angles of misorientaions in the original coarse grains. As the strain increases, the
misorientations of the cell walls increase leading to the formation of high-angle grain
boundaries as well as smaller grains inside the original coarse grains. This is known as
dislocation subdivision mechanism [3, 33, 34, 68]. A recent investigation [117] proposed
that the grain refinement mechanism in ECAP processing is dependent upon the SFE of
the alloy. These authors reported that in materials with low SFE the grain refinement
mechanism was gradually transformed from dislocation subdivision to twin
fragmentation.
95

Moreover, earlier reports [69, 70] proposed a microstructural model for grain
refinement in ECAP based on an inter-relationship between the formation of subgrain
boundaries and slip system. Specifically, the original coarse grains contain bands of
elongated cells or subgrains in the first pass of ECAP and on subsequent passes these
bands evolve into an array of equiaxed grains separated by high-angle grain boundaries.
An important characteristic of this model is that the ultimate size of the equilibrium
equiaxed grains is dictated by the width of the subgrain bands introduced into the
material in the first pass through the die. These models are useful and provide a clearly
explanation for the microstructure observed in a Cu-0.1%Zr processed by ECAP.
Detailed inspection of TEM images in Fig. 4.1 reveals that the average width of those
elongated bands (~350 nm) in a single pass is similar to the size of equiaxed grains (~310
nm) after a total of eight passes.
5.1.3 A comparison of microhardness evolution on the X and Y planes in
ECAP processing
The results in the present experiments relate to hardness measurements taken on
two different sectional planes in billets processed by route B
C
. To provide a direct
comparison between the X and Y planes, Fig. 5.1 shows a plot of the average Vickers
microhardness values for the two different orientations after processing by ECAP for up
to 8 passes, where these average values of Hv were determined from all the measurement
points on each plane as used in the construction of Figs 4.5 and 4.6. The results confirm
there is excellent agreement in the hardness values on both the cross-sectional X plane
96

and the longitudinal Y plane for each pressing condition. Therefore, the present results
are consistent with an earlier report where the microhardness values increased with
increasing numbers of ECAP passes and there were similar hardness values on the three
different orthogonal planes, X, Y and Z, for the Al-6061 alloy after pressing up to 6
passes at 298 K [118]. By contrast, hardness measurements on an Al-7034 alloy showed
similar hardness values on the X, Y and Z planes but with the hardness decreasing up to 6
passes of ECAP at 473 K [118]. This decrease in hardness with processing by ECAP is
unusual and it was due to a transformation of the metastable  -phase during processing
by ECAP at the higher temperature of 473 K [119]. Nevertheless, all of these results are
consistent in showing the development of good hardness homogeneity on all sectional
planes after processing by ECAP. In addition, the results in Fig. 5.1 confirm that, as
indicated in Fig. 4.7, the Cu-0.1% Zr alloy exhibits a rapid increase in hardness, by a
factor of approximately two, after a single pass of ECAP and thereafter the values on
both the X and Y planes continue to increase slowly in each additional pass up to an
average value of Hv  143 after eight passes.
No error bars are included in Fig. 5.1 because the errors are small and the bars
become obscured by the individual datum points. Nevertheless, detailed information on
the average Hv values shown in Fig. 5.1 and the associated error bars at the 95%
confidence limit are summarized in Table 5.1 for the cross-sectional X plane and the
longitudinal Y plane in each sample after different numbers of ECAP passes. In practice,
these average values of Hv were calculated using two orthogonal axes on each plane: the
Y and Z axes on the X plane and the X and Z axes on the Y plane. Two additional
97

Figure 5.1 The average Vickers microhardness values measured on the cross-sectional X
plane and on the longitudinal Y plane versus the number of ECAP passes.

conclusions are apparent from Table 5.1. First, in the early stages of processing, such as
after one pass, the error bars tend to be larger in any directions on both the X and Y
planes but at higher numbers of passes, such as after eight passes, the error bars are
consistently smaller. The smaller error bars at larger numbers of passes confirms the
increasing evolution towards a homogeneity in hardness. Second, the values of the error
bars in the Z directions in both the X and Y planes are larger by comparison with those in
the Y direction in the X plane or the X direction in the Y plane and this is especially true
in the earliest stages of processing. The larger error bars in the Z direction are due to the
occurrence of lower hardness values near the bottoms of the billets along the pressing
98

directions. Nevertheless, the errors bars in the Z direction are reduced with increasing
numbers of passes which again indicates a gradual evolution in the Cu-0.1% Zr alloy
towards a hardness homogeneity throughout the billets.
Table 5.1
Average values of the microhardness and the 95% confidence levels for the Cu-0.1% Zr
alloy processed by ECAP through one to eight passes.

No. of passes
in ECAP
Average microhardness, Hv
X plane Y plane
Y direction Z direction X direction Z direction
1 125 ± 1.7 125 ± 2.3 127 ± 1.1 127 ± 2.9
2 134 ± 1.6 134 ± 1.9 134 ± 1.1 134 ± 2.2
4 141 ± 1.4 141 ± 1.6 140 ± 0.9 140 ± 2.1
8 144 ± 1.0 144 ± 1.1 143 ± 0.7 143 ± 1.8

In addition the EBSD technique was used to examine the microstructure of
samples after ECAP eight passes in X plane and Y plane. The measurements were
performed both to create the crystallographic microstructures and to provide the
distributions of grain boundary misorientations as shown in Fig. 5.2. In both planes (Fig.
5.2(a) and (b)) have identical microstructure consisting of equiaxed (sub)grains after
deforming through eight passes. The measured grain sizes of ~370 nm and ~310 nm was
reported in X plane and Y plane, respectively. A similar average grain size was measured
as ~ 310 nm in X plane using TEM analysis as presented earlier in Fig. 4.2(b). Moreover,
the misorientaion distribution in X plane is essentially consistent with that in Y plane as
depicted in Fig. 5.2(c) where the number fraction of HAGBs was calculated as ~47% in
X plane and ~48% in Y plane. This microstructure analysis of samples after eight passes
99

Figure 5.2 EBSD orientation images of the Cu-0.1% Zr disks processed by ECAP eight
passes at different scanned positions of (a) X plane and (b) Y plane combined
with (c) corresponding distributions of misorientation in each plane.

100

of ECAP confirms the identity of mechanical property such as Vickers hardness
measured from X plane and Y plane of a Cu-0.1%Zr alloy after ECAP processing.

5.1.4 Hardness homogeneity after ECAP processing
Three important conclusions may be drawn from the present experimental results.
First, the hardness measurements on the cross-sectional X plane and the
longitudinal Y planes after one pass show there is a narrow region of lower hardness
running in a band near the bottom surface of the billet. Similar effects were reported
earlier for an Al-6061 alloy [120] and an Al-1050 alloy [121] after ECAP for 1 and 2
passes and for pure Cu [122] and an Al-6082 alloy [123] processed through 1 pass of
ECAP. From the ECAP principles, the shear strain,
N
 , introduced in ECAP is given by
a relationship in eq. 2.1. This equation was initially derived theoretically [17] and
subsequently confirmed experimentally [18, 22]. It is important to note that in
conventional ECAP a billet is expected to fully fill the corner of the die at the intersection
of the two parts of the channel and therefore the microstructure should be uniform
throughout the billet. In practice, however, experiments have shown that a strain
inhomogeneity is formed near the lower surface of the ECAP billet as revealed when the
local strains are measured using grid pattern [22]. Furthermore, finite element simulation
indicates that the depth of any less-deformed zone at the lower surface is dependent on
the strain hardening properties such that a higher strain hardening rate leads to a lower
equivalent strain and to a greater depth for the less-deformed zone [122].
101

The non-uniform deformation in the lower regions of the billets is associated with
the development of a corner gap or dead zone which is formed at the outer corner when
the billets pass through the die in ECAP processing. Thus, the dead zone is a space where
the billet is no longer in contact with the die wall and this removes all constraints
imposed on the billet by the die wall. The formation of a dead zone has been widely
reported in ECAP both through experimental observation [18. 22, 29. 124], through finite
element modeling [23, 24, 26-28, 122, 124-135] and using 3D finite volume modeling
[136]. It was also reported for rotary-die ECAP [137]. The present results show lower
hardness values in the vicinity of the bottom surface of the billet in the early stages of
ECAP processing and this is consistent with these various analyses.
Second, another important result is that, although there is a variation in hardness
in the Z direction, the values of Hv taken along the X axes remain reasonably constant
under all processing conditions. This is consistent with earlier results both on aluminum
alloys [120, 121] and on a magnesium ZK60 alloy after eight passes [138]. A hardness
homogeneity in the X direction is attractive because it supports the feasibility of
developing a continuous processing operation for the production of metal sheets and
wires through the incorporation of ECAP and thus it confirms the potential for producing
materials by SPD processes for use in practical applications [16]. It is also consistent
with the feasibility of using the ECAP-Conform process for the processing of long rods
or wires [139].
102

Third, microhardness homogeneity was achieved throughout the billets after a
sufficiently large number of ECAP passes. For the Cu-0.1% Zr alloy, this critical
condition corresponds to pressing through at least eight passes and this gives good
microhardness homogeneity on both the cross-sectional and longitudinal planes.
Furthermore, the experimental results in Fig. 5.1 demonstrate there are similar amounts
of hardness evolution on both the X and Y planes in each separate pass. Earlier
experiments have demonstrated a direct correlation between microhardness
measurements and the internal microstructures observed in SPD materials processed
either by ECAP [49, 69] or HPT [49, 65, 100] and therefore it is reasonable to anticipate
that a microstructural homogeneity is achieved in the Cu-0.1% Zr alloy after processing
through eight passes.
5.2 HPT discussion
5.2.1 Hardness and microstructural evolution in HPT processing
The principle of HPT is relatively simple. When a disk is processed by HPT, the
equivalent von Mises strain imposed on the disk, ε
eq
, is given by a relationship of eq. 2.4.
It is apparent from eq. 2.4 that the imposed strain is a maximum around the edge of the
disk whereas it becomes equal to zero in the center of the disk where r = 0. In practice,
however, eq. 2.4 includes only the effect of torsional straining and there is also an
additional strain imposed by the compressive pressure, P. Recent experiments on an Al-
6061 aluminum alloy showed the initial applied pressure leads to a significant increase in
hardness across the disk diameter even in the absence of any torsional straining [141].
103

It is anticipated from eq. 2.4 that the microstructures and microhardness values
introduced by HPT will be extremely inhomogeneous. Nevertheless, early experiments
demonstrated a gradual evolution towards a homogeneous structure in disks of Ni
processed by HPT by increasing either the applied pressure, P, and/or the total numbers
of revolutions, N [51, 142]. Subsequently, numerous reports confirmed this gradual
evolution towards homogeneity [49, 50, 53, 65, 96, 143-150] and the evolution was
successfully modeled using strain gradient plasticity theory [102].
The results of this investigation show there is a gradual evolution toward hardness
homogeneity throughout the disks of the Cu-0.1% Zr alloy after five or more turns of
HPT processing. This is similar to earlier reports for pure Cu processed by HPT for up to
5 to 25 turns [54-61, 147] and also to reports for several Cu alloys [60-63]. The results
from Cu-Zn alloys [60, 61] suggest that the rate of hardness evolution is dependent upon
the stacking fault energy of the alloy and this is consistent also with other data suggesting
a significant influence of stacking fault energy on the development of an ultrafine-grained
structure in HPT processing [151-157] and in processing by ECAP [117, 158-160] .
The gradual evolution towards homogeneity with increasing numbers of turns is
visible for the Cu-0.1% Zr alloy both in the microhardness results and in the
microstructural analysis. A lower hardness was introduced in the central region of the
disk in the early stages of processing, equivalent to N = 1/4 turn in the present
experiments, and this was attributed to the reduced shear strain in the central area as
predicted by eq. (5.2). Conversely, at the peripheral region the imposed strain was high
and this led to increased values of Hv. Thereafter, the lower hardness in the central
104

region gradually swept outwards with increasing numbers of turns, as proposed earlier
[51], until it occupied the total area of the disk as is evident at N = 5 turns in Fig. 4.15(c).
Extending the torsional straining to N = 10 turns demonstrates complete homogenization
and a constant value of Hv throughout the disk surface as shown in Fig. 4.15(d).
Furthermore, these hardness values have negligible error bars as shown in the right
column of Table 4.1.
As discussed recently, the present results showing the variation of hardness with
the number of HPT turns, and therefore with the equivalent strain expressed by eq. (5.1),
depends in practice upon the nature of recovery in the material [53, 161]. Most of the
commercial purity metals and simple metallic alloys exhibit strain hardening in the initial
stages of HPT processing. These types of metals, such as commercial purity Al and Al
alloys [49, 50, 100, 148] and Cu and Cu alloys [54, 57, 60, 61, 147, 149] demonstrate
lower hardness values at the centers of the disks in the early stages of deformation but
with the hardness increasing with strain until ultimately there is a saturation value of Hv
throughout each disk. This suggests the hardness homogeneity is evolving in the absence
of any recovery due to the relatively low stacking fault energy. By contrast, a material
such as high-purity aluminum has a very high stacking fault energy and this leads to easy
cross-slip and rapid microstructural recovery [53] so that higher values of hardness are
then recorded in the centers of the disks in the early stages of straining [53, 71, 162, 163].
Moreover, very recent experiments showed that the situation is different for a material
such as the Zn-Al eutectoid alloy where all of the measured hardness values were lower
after HPT processing than the hardness values recorded in the annealed condition prior to
105

processing [65, 161]. This unique hardness distribution was attributed to the significant
reduction in the distribution of the Zn precipitates in the Al-rich grains due to the high
pressure applied to the disks during HPT processing [65, 161].
The present results on the Cu-0.1% Zr alloy permit a direct comparison between
the evolution of local hardness and the severely deformed microstructures in a series of
disks. It is observed in the microhardness measurements that there are regions of lower
hardness in the centers of the disks after processing at lower numbers of HPT revolutions
and these regions gradually disappear with increasing the numbers of HPT revolutions.
Using the microstructure observed by TEM and EBSD, these results are clearly explained
by the strong effect of grain size strengthening where larger grains were observed in the
centers at lower strains and thereafter smaller grains were evolved by dislocation
activities with increment of imposed strains (numbers of revolutions). By contrast, the
hardness near the edge of each HPT disk is similar while the measured grain sizes in Fig.
4.11 are different. It is worth noting that the hardness was averaged from all
microstructural characteristic not only the grain sizes. Dislocation density inside the
grains is another factor that will contribute to the hardness. Specifically, the
microstructure in the early stage of deformations i.e., a quarter and one turn, shows
mainly dislocation cell structures with moderate/high density of dislocations inside grains
with average grain sizes of ~350 nm. At high imposed strain levels, however, the grains
are refined with average grain size about 200 nm and low density of dislocations inside.
Therefore, it reasonable to anticipate that the change of the dislocation density balances
106

the contribution from the grain sizes resulting in constant hardness values near the edge
of the disks.
It is apparent that the hardness and microstructure evolution strongly depends on
the number of HPT revolutions. In the present experiments, reasonably homogeneous
hardness and homogeneous microstructures were observed after 5 turns and more. The
distributions of the grain boundary misorientation angles in Fig. 4.13 provide a direct
confirmation of the microstructural homogeneity, in addition to the hardness
homogeneity. There is an obvious difference in the fraction of high-angle grain
boundaries at the center and edge of the disk subjected to a quarter turn shown in Fig.
4.13(a), however, the fractions of high-angle boundaries are fairy similar in the center
and the periphery of the disks strained through 5 and 10 turns.
Furthermore, the HPT processing is especially effective not only in producing a
homogeneous microstructure throughout a disk of a Cu-0.1%Zr but also refining the
grains to submicrometer level. The present experiments give an average grain size of ~20
µm under as-received conditions, whereas the grain sizes achieved in the sample using
HPT are reported as ~300 nm at the center and ~250 nm at the edge. Therefore, using the
three separate procedures of microhardness measurements, TEM and EBSD, it is shown
that HPT may be used effectively to develop an exceptional grain refinement and
essentially homogeneous microstructure through the sample provided there is a sufficient
torsional strain (as measured in terms of N) under high pressure.

107

5.2.2 Potential for achieving superplastic flow after HPT processing
The tensile experiments were conducted at temperature of 673 K and 723 K
corresponding to ~0.50T
m
and ~0.53T
m
where T
m
is the absolute melting temperature of
materials. It should be noted that the addition of a small amount of Zr increases the
recrystallization temperature and inhibits grain growth [11] thereby leading to grain size
stability at elevated temperatures. The results in Fig. 4.18 show that the maximum tensile
elongation recorded in these tests was ~240% for specimens processed through 5 and 10
turns of HPT and tested at temperature of 723 K and initial strain rate of 1.0 × 10
-4
s
-1

whereas the maximum tensile elongation was recorded as ~100% for specimens
processed through 1 turns. These results can be explained using the microstructural
characteristics observed in each condition. Specifically, samples after five and ten turns
reveal a smaller grain size with higher fraction of high-angle misorientation boundaries
whereas a sample after one turn presents an inhomogeneous structure consisting of large
and small grains with a low fraction of high-angle boundaries. It is well-established that a
microstructure of equiaxed grains separated by high-angle boundaries is a prerequisite for
achieving high tensile ductility.
Superplasticity is defined as the ability of a material to pull out uniformly to a
very high elongation in tension without the development of any incipient necking. The
measured elongations in superplasticity are generally at least 400% and the measured
strain rate sensitivities are close to ~0.5. Based on this definition the maximum tensile
elongation of ~240% is not within the regime of superplasticity. In superplastic flow
alloys, the flow stress ( is related to the strain rate through the expression of
108

where m is strain rate sensitivity. The value of m is usually close to ~0.5 under
conditions of optimum superplasticity based on grain boundaries sliding [164]. In other
words, the strain rate sensitivity, m, is simply defined as . To check on the
strain rate sensitivity associated with this elongation, the measured flow stress was
plotted logarithmically against the initial strain rate for a specimen after HPT 10 turns as
shown in Fig. 5.3. It appears that the stain rate sensitivity is ~0.2 in this specimen which
is not close to 0.5 for superplastic flow.

Figure 5.3 Flow stress versus strain rate for the HPT specimen subjected through 10 turns
having tensile testing at temperature of 723 K.
109

The earlier experiments [11] studied the influence of Zn and Zr additions on
copper. The results demonstrated that the presence of Zn and Zr are both beneficial in
promoting the occurrence of superplastic ductilities. Zirconium is needed because it
increases the recrystallization temperature, inhibits grain growth and, therefore, serves to
retain a small grain size at elevated temperatures, and zinc is beneficial because it
introduces solute atoms into the matrix so that dislocation creep is inhibited and
superplastic flow can occur more easily. Therefore, the presence of a solid solution
alloying element is an important prerequisite for achieving good superplastic elongations
in tensile testing. This is confirmed by a recent experiment on a Zn-22%Al eutectoid
alloy processed by HPT [165] where the tensile elongation of ~1800% was reported.
At the present, however, there are only a limited number of reports of true
superplastic flow in materials processed by HPT. Moreover, most elongations are not
large by comparison with materials processed by ECAP where the grain size are larger
but the reported elongations to failure may exceed 1000% [91] or exceed 3000%[166]. A
possible explanation for this dichotomy may lie in the very small thicknesses within the
gauge lengths of samples processed by HPT. In present experiments, the gauge widths
were only 1 mm and the gauge thicknesses were 0.8 mm after HPT processing. It is well
known that the measured elongations in tensile testing are dependent upon the
dimensions of the samples [167, 168]. Specifically, tensile specimens having very small
cross-sections are generally less ductile than larger specimens.
110

5.3 Comparison of microstructures and mechanical properties in
different processing
5.3.1 Microstructures in different SPD techniques
Detailed inspection in Table 5.2 as well as the microstructure images presented in
Chapter 4 demonstrated the interesting conclusions in microstructures developed by
different SPD processing. It is apparent that the microstructures in the samples after these
three different processing procedures are significantly refined and contain an array of
grains that are reasonably equiaxed; it should be noted that the average grain size of ~20
µm was measured from an as-received sample. However, the microstructures of the Cu-
Zr alloy are refined to different level depending upon the SPD procedures. Detailed
measurements gave an average grain size of ~ 330 nm after ECAP 8 passes, thereafter,
with the combination of ECAP 8 passes + HPT 5 turns the grains are more refined to the
average grain size of ~280 nm in the center region and ~270 nm near the edge area.
Similar average grain sizes are measured in processing by only HPT 10 turns with an
average grain size of ~280 nm and ~240 nm at the center and edge of the disk,
respectively; the values are the average values from TEM and EBSD methods as
presented in Table 5.2. Therefore, three important views can be summarized from these
results. First, it is confirmed that the SPD techniques including ECAP, HPT and
ECAP+HPT are effective processing in grain refinement for this Cu-Zr alloy. Second,
there is a potential to improve the grain refinement in ECAP processing by subjected to
further straining using a different path of deformation as ECAP followed by HPT. Third,
111

HPT processing can produce a Cu-Zr alloy having smaller grain sizes than in processing
by ECAP when the processing is conducted at room temperature. This is consistent with
earlier results showing that HPT is generally more effective than ECAP in producing
exceptionally smaller grain sizes [10, 115, 169, 170] including recent results on a pure
copper [54] and a copper alloy [171].

Table 5.2 Microstructural parameter for the Cu-0.1% Zr alloy obtained by TEM and
EBSD
Cu-0.1%Zr
ECAP
8 passes
HPT
10 turns
ECAP 8p +
HPT 5 turns
center edge center edge
TEM Grain size (nm) 310 290 250 280 270
EBSD Grain size (nm) 350 270 230 - -
Fraction of HAGBs (%) 47 55 65 - -

In addition, the EBSD images in Fig. 5.4 reveal the microstructure produced by
HPT appears generally to be more homogeneous and there is a larger fraction of equiaxed
grains having boundaries with high-angle misorientation (Fig. 5.2(b)). By contrast, some
of the grains are separated by LAGBs after processing by ECAP (Fig. 5.2(a)). Table 5.2
shows the fraction of HAGBs as ~47% after ECAP, ~55 after HPT in the center and
~65% after HPT in the edge region. These measurements show HPT is also more
effective in producing boundaries having high angles of misorientation.

112

Figure 5.4 Color-coded contour maps showing the distributions of the Vickers
microhardness values (a) on the cross-sectional plane after ECAP for 8 passes
and (b) over the disk surface after HPT for 10 turns: the hardness colors are
shown in the key on the right.

113

5.3.2 Mechanical properties in different SPD techniques
The influence of the imposed strain on the hardness evolution in Cu-Zr samples
after processing by ECAP and HPT are observed in sections 4.1.2 and 4.2.2, respectively.
The results reveal that the values of Hv gradually increase with increasing imposed strain
and therefore with the numbers of passes in ECAP and the numbers of turns in HPT.
Ultimately, the hardness becomes reasonably saturated when a sufficient high strain is
attained. Moreover, the values of Hv in the combination processing are also measured (in
section 4.3.1) to evaluate whether it is possible to improve the properties if a sample is
subjected to higher plastic deformation by using different processing.
In order to provide a visual comparison of the levels of hardness saturation
achieved in ECAP, HPT and ECP+HPT processing, the color-coded contour maps are
plotted in the same range as shown in the Fig. 5.5. The contour maps show that
reasonably homogeneous hardness distributions are attained over the disk surfaces of the
Cu-Zr alloy after both ECAP and HPT, thereby confirming the processing conditions for
this alloy of ECAP for 8 passes and HPT for 10 turns are sufficient to produce high level
of hardness homogeneity for both conditions as demonstrated in Fig. 5.5(a) and (b). In
addition, there is a potential to improve the hardness of this alloy after deformed through
ECAP 8 passes by adding higher strains using different processing such as HPT as
depicted as ECAP 8p + HPT 5 turns in Fig. 5.5(c). A detailed evaluation of the
microhardness measurements shows that the average saturation values vary from Hv 
143 after ECAP to Hv 156 after ECAP + HPT and Hv  164 after HPT. By comparison,
114

the microhardness value in the as-received condition was Hv  95. Thus, this processing
at room temperature leads to a significant improving in the hardness value with increase
by factors of approximately 1.5, 1.6 and 1.7 for ECAP, ECAP + HPT and HPT,
respectively. This different in hardness for the Cu-Zr alloy is primarily attributed to: first,
the grain size; second, the defects in materials; and third, the history of materials. The
first reason is well-known as the grain size strengthening or Hall-Petch relationship [172,
173] where the relation between the yield stress and grain size is described by

5.1
where σ
y
is the yield stress, σ
o
is a materials constant for the starting stress for dislocation
movement, k
y
is the strengthening coefficient depending on each material, and d is the
average grain diameter. In the absence of appreciable strain hardening, there is a linear
correlation between the yield strength and hardness of the materials of the form [174]

5.2
From eq. 5.1 and 5.2, it is apparent that the material becomes stronger if the grain size is
reduced. The average grain size and fraction of HAGBs of each processing observed by
TEM and EBSD are summarized in Table 5.2. It is clear that the lower hardness in ECAP
is attributed to the larger grain size compared to HPT and combination processing.
Despite the similar average grain size observed in ECAP + HPT and HPT the average
hardness value in ECAP + HPT is slightly lower than in HPT. This might be explained by
the history of a sample. Prior to deformation in ECAP at room temperature the sample
115

was heat treatment at 973 K for 1 hour which leads to a lower hardness in processing of
ECAP and ECAP + HPT.

Figure 5.5 Color-coded contour maps showing the distributions of the Vickers
microhardness values processed by (a) ECAP 8 passes (b) HPT 10 turns and
(c) ECAP 8 passes followed by HPT 10 turns.

116

Figure 5.6 shows representative stress-strain curves for the specimens after
ECAP and HPT processing and for specimens in the as-received and annealed
conditions. Testing was conducted using an initial strain rate, , of 1.0 × 10
-3
s
-1
at a
temperature of 673 K corresponding to an homologous temperature of ~0.5T
m
where T
m
is the absolute melting temperature of the material. This testing condition was selected
for two reasons. First, the strain rate of 1.0 × 10
-3
s
-1
is similar to the operating strain
rate for material fabrication in conventional industries. Second, it was shown in an
earlier investigation using a Cu-0.18% Zr alloy processed by ECAP that the grains of
this alloy remain ultrafine and essentially constant up to a temperature of 773 K [11].
The plot shows that the heat treatment at 973 K for 1 hour leads to a lower yield
stress but an increase in the elongation to failure compared to the as-received condition.
However, the Cu-Zr alloy processed by ECAP for 8 passes exhibits a higher strength and
an elongation which is higher than in the as-received condition but similar to the annealed
condition. For the specimen processed through 10 turns by HPT, there is a significant
improvement in both the tensile strength and the ductility with a yield stress that is about
two times higher and an elongation that is about 1.7 times higher than in the annealed
condition.
Generally, the materials produced by SPD techniques usually have special
structural characteristics such as high dislocation densities, non-equilibrium grain
boundaries and other structural features associated with intense plastic straining [175,
176]. These defects produced by SPD processing lead also to a high resistance to plastic
deformation as confirmed by the improved yield strengths after ECAP and HPT in these
117

experiments. Typically, materials processed by SPD procedures exhibit high strength but
relatively low ductility. However, the results from tensile testing at a high temperature in
this study, as illustrated in Fig. 5.6, demonstrate that the Cu-Zr alloy has both higher
strength and increased ductility after HPT processing compared with the material in the
as-received condition. This result is similar to an earlier report where pure Ti was
subjected to HPT processing through 5 turns and then tested in tension at 523 K [72]. A
recent review summarized the key strategies now available for achieving both high
strength and good tensile ductility in ultrafine-grained and nanocrystalline materials
[177]. In the present experiments, the unusual combination of both high strength and
good ductility in the Cu-Zr alloy after processing by HPT is attributed to the exceptional
grain stability and to the high fraction of HAGBs. The grain stability in this alloy is
achieved through the addition of a small amount of Zr and it should be noted that this
addition increases the recrystallization temperature and also inhibits grain growth thereby
producing excellent grain stability at elevated temperatures [11]. Thus, grain stability
and a high fraction of HAGBs contribute directly to the high tensile ductility achieved in
this alloy at a temperature of 673 K.

118

Figure 5.6 Plot of engineering stress versus engineering strain at a testing temperature of
673 K and an initial strain rate of 1.0 × 10
-3
s
-1
for the Cu-0.1% Zr alloy in the
as-received condition, in the annealed condition, after processing by ECAP for
8 passes and after processing by HPT for 10 turns.

119

6. Summary and Conclusions
In the present research, a Cu-0.1%Zr alloy was processed by different techniques
of severe plastic deformation (SPD): equal-channel angular pressing (ECAP), high-
pressure torsion (HPT) and a combination of both processing (ECAP + HPT). In order to
undertake a meaningful comparison between processing, all procedures were conducted
at room temperature. Prior to processing by ECAP the as-received samples were annealed
for 1 h at 973 K due to the difficulty to process at room temperature. For ECAP the
sample was processed up to 8 passes using a solid die with a channel angle of 110  and
route B
C
and for HPT the sample was processed through up to 10 turns under an applied
pressure of 6.0 GPa. For the combination processing, samples were prepared from ECAP
4 passes and 8 passes and then subjected to HPT for five turns. The experiments were
conducted to evaluate the microstructural evolution and mechanical properties for each of
the processed and their combination. Transmission electron microscopy (TEM) and an
electron backscatter diffraction (EBSD) technique was employed to measure the
microstructural features, grain size distributions and the distribution of the misorientation
angles. The mechanical properties of the processed samples were examined and
compared both at a room temperature using microhardness measurements and at an
elevated temperature using tensile testing. The summary of results and discussion are
described as follows.
1) A Cu-0.1% Zr processed through ECAP processing leads to significant grain
refinement. This processing reduced the grain size from an initial value of ~30 µm in
120

annealed condition to a value of ~330 nm after 8 passes. The microstructure features
observed from TEM and EBSD demonstrated that there is a transition from elongated cell
structure to equiaxed grain structure where the ultimate size of the equiaxed grains is
approximately equal to the width of the subgrain bands introduced into the material in the
first pass of deformation.
2) The Vickers hardness measurements on X and Y planes of billets after ECAP
revealed the presences of the lower hardness value adjacent to the lower surfaces of the
billets in the early stages of processing due to a corner gap or dead zone. Then there is a
gradual evolution to a homogeneous hardness distribution throughout the billet after 8
passes. The experiments also show that the evolution towards homogeneity occurs at a
similar rate on both the X and Y planes. In addition, the tensile testing at elevated
temperature demonstrated a significant improvement of the tensile strength with
admissible penalization on ductility in specimens after ECAP processing.
3) The HPT processing provides an exceptional grain refinement. The
measurements gave an average grain size of ~280 nm in the central region and ~240 nm
near the edge of the disk where the average grain size of ~20 µm was measured from the
as-received sample. The microstructures after five and ten turns are reasonably uniform
consisting of equiaxed grains with a high fraction of high-angle grain boundaries.
4) The measured hardness values are low in the central region of the disks in the
early stages of processing but there is a high level of hardness homogeneity after 5 and 10
turns. The hardness values are matched by microstructural observations showing a
121

gradual evolution into a well-defined equiaxed grain structure. The tensile testing shows
that the samples after HPT processing exhibit superior mechanical properties over the as-
received samples with higher strength and higher ductility. These combination properties
are attributed to the grain stability and the high fraction of high-angle grain boundaries.
5) The experimental results demonstrated there is a potential for improving the
microstructure through the combination processing where the grains can be refined more
after ECAP 4 passes and 8 passes by subsequently subjecting to HPT 5 turns; denoted as
ECAP 4 passes + HPT 5 turns and ECAP 8 passes + HPT 5 turns. The grain size is
reduced from ~430 nm after ECAP 4 passes to ~300 nm after ECAP 4 passes + HPT 5
turns and the grain size is reduced from ~330 nm after ECAP 8 passes to ~290 nm after
ECAP 8 passes + HPT 5 turns. The reduction of grain sizes in this combination
processing leads to the increment of hardness values.
6) Although the combination processing exhibits an improvement in grain
refinement and hardness compared to ECAP but the average grain size in combination
processing is similar to HPT 5 turns and 10 turns. Moreover, the hardness of combination
processing tends to be lower than that of only HPT processing. This might be due to the
heat treatment before ECAP processing.
7) For a comparison between ECAP processing and HPT processing, it is
demonstrated that ultrafine grains are achieved after processing by both ECAP and HPT
but the mean grain size is smaller and the fraction of high-angle grain boundaries is
higher after HPT
122

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Abstract (if available)

Abstract A copper alloy, Cu-0.1% Zr, has been processed at room temperature by different techniques of severe plastic deformation (SPD), namely equal-channel angular pressing (ECAP), high-pressure torsion (HPT) and a combination of both processing (ECAP + HPT). The experiments were conducted to evaluate the microstructural evolution and mechanical properties for each of the processed and their combination. A transmission electron microscopy (TEM) and an electron backscatter diffraction (EBSD) techniques were employed to measure the microstructural features, grain size distributions and the distribution of the misorientation angles. The mechanical properties of the processed samples were examined and compared both at a room temperature using microhardness measurements and at an elevated temperature using tensile testing. ❧ Using TEM and EBSD techniques, it is demonstrated that these three SPD procedures have a potential for producing an ultrafine-grain structure containing reasonably equiaxed grains with high-angle boundary misorientations. However, microstructures are refined in different level depending on the processing operation. The grain refinement mechanisms are primarily governed by dislocation activities. ❧ Microhardness distribution of the strained samples shows that there is a non-uniform of this distribution in the early stages of deformation where the lower hardness values were measured near the bottom of samples for ECAP and at the central region for HPT. This inhomogeneity is gradual decreased with increasing imposed strain and ultimately the microhardness distribution is reasonably homogeneous when the sufficient strain is subjected to the sample. ❧ The tensile results demonstrate that the samples after SPD processing exhibit superior mechanical properties with the combination of high strength and ductility compared to the as-received materials where the maximum elongation to failure of ~240% at 723 K using a strain rate of 1.0 × 10⁻⁴ s⁻¹ is achieved in a sample processed by HPT. This elongation however does not fulfilled the requirements for true superplastic flow where the measured elongation in tension should be at least 400%.

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Asset Metadata

Creator Wongsa-Ngam, Jittraporn (author)

Core Title Microstructural evolution and mechanical properties of a copper-zirconium alloy processed by severe plastic deformation

School Viterbi School of Engineering

Degree Doctor of Philosophy

Degree Program Mechanical Engineering

Publication Date 03/22/2013

Defense Date 12/10/2012

Publisher University of Southern California (original), University of Southern California. Libraries (digital)

Tag copper alloy,mechanical properties,microstructure,OAI-PMH Harvest,severe plastic deformation

Language English

Contributor Electronically uploaded by the author (provenance)

Advisor Langdon, Terence G. (committee chair), Goo, Edward K. (committee member), Hodge, Andrea M. (committee member)

Creator Email iamojeed@gmail.com,wongsang@usc.edu

Permanent Link (DOI) https://doi.org/10.25549/usctheses-c3-226747

Unique identifier UC11292781

Identifier usctheses-c3-226747 (legacy record id)

Legacy Identifier etd-WongsaNgam-1479.pdf

Dmrecord 226747

Document Type Dissertation

Rights Wongsa-Ngam, Jittraporn

Type texts

Source University of Southern California (contributing entity), University of Southern California Dissertations and Theses (collection)

Access Conditions The author retains rights to his/her dissertation, thesis or other graduate work according to U.S. copyright law. Electronic access is being provided by the USC Libraries in agreement with the a...

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