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STUDY OF MICROSTRUCTURE AND TEXTURE DEVELOPMENT DURING
SUPERPLASTIC DEFORMATION OF STATICALLY AND DYNAMICALLY
RECRYSTALLIZING Al ALLOYS
Thesis · October 2014
DOI: 10.13140/RG.2.1.3461.3361
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STUDY OF MICROSTRUCTURE AND TEXTURE DEVELOPMENT
DURING SUPERPLASTIC DEFORMATION OF STATICALLY AND
DYNAMICALLY RECRYSTALLIZING Al ALLOYS
A THESIS
submitted by
ARUN BABU K
for the award of the degree
of
MASTER OF SCIENCE
DEPARTMENT OF METALLURGICAL AND MATERIALS ENGINEERING
INDIAN INSTITUTE OF TECHNOLOGY MADRAS.
OCTOBER 2014
Dedicated to My Parents
ii
THESIS CERTIFICATE
This is to certify that the thesis titled “STUDY OF MICROSTRUCTURE AND
TEXTURE DEVELOPMENT DURING SUPERPLASTIC DEFORMATION
OF
STATICALLY
AND
DYNAMICALLY
RECRYSTALLIZING
Al
ALLOYS’’ submitted by Arun Babu K, to the Indian Institute of Technology
Madras, Chennai for the award of the degree of Master of Science, is a bona fide
record of the research work done by him under our supervision. The contents of this
thesis, in full or in parts, have not been submitted to any other Institute or University
for the award of any degree or diploma.
Dr. V. Subramanya Sarma
Associate Professor
Dept. of Metallurgical and Materials Engineering.
IIT-Madras., 600 036
Date: 31st October 2014
Place: Chennai
iii
ACKNOWLEDGEMENTS
It has been a wonderful experience for me while doing my masters works that provided me
with the opportunity to come in touch with some of the nicest people with great abilities and
fortunate to find phenomenal support and encouragement from them. I sincerely acknowledge
all of them for their help and encouragement without which it would have been an up-hill
task to overcome the barriers on my way.
The first and foremost person I would like to thank is my research supervisor Dr. V.
Subramanya Sarma, Associate Professor, Department of Metallurgical and Materials
Engineering, Indian Institute of Technology, Madras, for his inspiring guidance, unstinted
support, persistent encouragement and timely help during my research career.
I would like to thank Dr. S. Sankaran, Associate Professor, Department of Metallurgical
and Materials Engineering, Indian Institute of Technology, Madras, for his timely help and
valuable advices for during my research career.
I express my sincere thanks to Dr. M. Kamaraj, Professor and Head of the Department; Dr.
S. S. Bhattacharya, Professor; Dr. Uday Chakkingal, Professor; Dr. P. Kesavan Nair,
Professor and Dr. N. V. Ravi Kumar, Associate Professor, for providing me with all the
facilities for the successful completion of the project. I express my sincere thanks to Graduate
Test Committee members, Dr. Uday Chakkingal, Professor and Dr. Srikanth Vedantam,
Associate Professor for their constructive criticism during the course of my master’s
program.
I would like to thank sincerely Dr. K. A. Padmanabhan, University Chair Professor of
University of Hyderabad for his valuable advices and suggestions for my research on
superplasticity.
I am grateful to Mr. C. K. Gopalakrishnan, staff, material testing facility, for his help and
support in carrying out high temperature tensile tests.
I am grateful for the help and co-operation received from all the supporting staff and
scientific officers in our department viz. Mrs. D. Kanchanamala, Mr. T. Rahavaiah, Mr. R.
Vardachari, and Mr. R. Parthiban, Mr. Veerabathiran for their assistance during my
experiments and characterization. I am thanking to Mr. V. Sampath and Mr. I. Prakash for
their administrative support.
i
I take immense pleasure in mentioning the names of my colleagues and my friends who
helped me during my course of stay at IIT Madras. My special thanks to C. N. Athreya for
timely help in my EBSD studies.
Last, but not the least, I am greatly thanks to my family for their continuous affection and
support throughout my research work. I also thank everyone who has supported me in my
M.S. either directly or indirectly.
Arun Babu K
October- 2014
ii
ABSTRACT
Keywords: Superplasticity, static recrystallization, Continuous recrystallization, Dynamic
grain growth, GBS, crystallographic slip
Superplasticity is the ability of the material to exhibit very high tensile elongation under
certain conditions such as high temperature and low stress and these materials are generally
referred to as superplastic materials. In aerospace and automotive industries superplasticity is
widely exploited in the forming of complex shapes and curved parts. Al alloys are widely
used in aerospace, automobile and marine applications due to their high strength to weight
ratio. Superplastic Al alloys are classified as: a) statically recrystallizing (e.g., AA 5083) and
continually recrystallizing (e.g., AA 2004). In the present work a systematic study of the
superplastic deformation characteristics of the statically and dynamically recrystallizing
alloys (AA 5456 and AA 2004) in the initial, transient and the steady state regions of
superplastic flow was carried out to distinguish between the operating mechanism(s) during
steady state, isotropic and anisotropic superplastic flow caused by grain shape and texture
anisotropy in the starting material
Also the optimum strain rate and temperature was
established for AA 5456 (a new variant of AA 5083 developed by Hydro Aluminum,
Germany) and the texture changes and microstructural changes were studied in the optimum
superplastic conditions. Aluminium alloy AA 5456 was exhibited discontinuous static
recrystallization following annealing at 550°C by particle stimulated nucleation. The
recrystallized microstructure was fine grained (~ 8µm) and was having a large fraction of
high angle boundaries. The optimal superplastic flow (highest % elongation) for AA 5456
was achieved at 550°C at a strain rate of 5×10-4 s-1. The alloy AA 5456 exhibited better
superplastic properties (higher elongation and lower stresses) in comparison to alloy AA
5083. The Initial rolling texture in AA 5456 was completely annihilated and a weak rotated
cube component was present following annealing at 550°C for 30 min. The weak rotated
cube component present in the annealed sample was modified following superplastic
deformation and after a strain of 250% (true strain ~ 1.25); a weak <100> fiber was observed.
With continued deformation, the fiber component was changed to a weak cube component.
Alloy AA 5456 exhibited dynamic grain growth during superplastic deformation. The grain
growth and growth rate is consistent with the model proposed by Caceres and Wilkinson
(1984).
iii
The optimal superplastic flow (highest percentage of elongation) for AA 2004 was
achieved at 457°C, at a strain rate of 3×10-3 s-1. It exhibited continuous dynamic
recrystallization during deformation and the initial banded microstructure changed to a fine
grained (~ 3 µm) microstructure, which was suitable for superplasticity after a true strain 0.7.
In AA 2004, during initial stages of superplastic deformation, the retention of {110} <
11̅2 > Brass’ texture component was observed and this is a stable orientation during
crystallographic slip under plane strain compression. However, the ODF values of major
texture components were decreased with increasing superplastic deformation. This is
believed to be due to the random rotation of grains. ‘Orientation diffusion’ model for texture
spread proposed by Engler et al. (2000) was validated for AA 2004 and the texture changes
following superplastic deformation are in good agreement with the predictions of the above
model.
iv
TABLE OF CONTENTS
ACKNOWLEDGEMENTS
i
ABSTRACT
iii
TABLE OF CONTENTS
v
LIST OF TABLES
viii
LIST OF FIGURES
ix
ABBREVIATIONS
xiv
NOTATIONS
xv
1 INTRODUCTION
1
2 LITERATURE REVIEW
3
2.1 Mechanisms of superplasticity………………………………………………….
5
2.1.1 Grain boundary sliding (GBS)..………………………………………….....
5
2.1.2 Diffusion creep……………………………………………………………...
9
2.1.3 Dislocation creep…………………………………………………………...
9
2.2 Recrystallization and superplasticity…………………………………………… 10
2.3 Texture and superplasticity……………………………………………………... 12
2.4 Grain growth during superplasticity……………………………………………
15
2.5 AA 5XXX alloys………………………………………………………………... 16
2.6 AA 2XXX alloys……………………………………………………………….. 17
3 OBJECTIVES AND SCOPE
18
3.1 Organization of thesis………………………………………………………….
4 EXPERIMENT DETAILS
19
20
4.1 Materials………………………………………………………………………… 20
4.2 Room temperature tensile testing…………………………………………….
20
4.3 High temperature tensile testing………………………………………………
20
v
4.4 Hardness………………………………………………………………………… 21
4.5 Scanning electron microscopy (SEM) studies………………………………….
21
4.6 Electron back scatter diffraction (EBSD) studies……………………………...
21
4.7 The overall experiment plan……………………………………………………
22
5 RESULTS AND DISCUSSIONS
23
5.1 Statically recrystallizing alloy, AA 5456………………………………………
23
5.1.1 Base metal characterization………………………………………………... 23

Composition…………………………………………………………….. 23

Micro hardness………………………………………………………….. 23
5.1.2 Room temperature mechanical properties of AA 5456……………………. 23
5.1.3 High temperature tensile behaviour of AA 5456…………………………... 24
5.1.4 Microstructure and texture evolution following annealing at 550°C……… 27

Texture evolution following annealing at 550°C……………………….. 29

Grain boundary mis-orientation distribution following annealing at
550°C….................................................................................................... 30
5.1.5 Microstructure and texture evolution following superplastic deformation at
optimal strain rate and temperature………………………………………..

31
Texture evolution following superplastic deformation at optimal strain
rate and temperature ……………………………………………………. 36

Grain boundary mis-orientation distribution following superplastic
deformation at optimal strain rate and temperature…………………….. 38
5.1.6 Microstructure and texture evolution following deformation at high strain
rate…………………………………………………………………………... 40
5.2 Dynamically recrystallizing alloy, AA 2004……………………………………. 42
5.2.1 Base metal characterization………………………………………………... 42

Composition…………………………………………………………….. 42
5.2.2 Room temperature mechanical properties of AA 2004……………………. 42
5.2.3 High temperature tensile behaviour of AA 2004…………………………... 43
5.2.4 Microstructure and texture evolution following annealing at 457°C……… 45

Texture evolution following annealing at 457°C……………………….. 47
vi

Grain boundary mis-orientation distribution following annealing at
457°C…………………………………………………………………… 49
5.2.5 Microstructure and texture evolution following superplastic deformation at
optimal strain rate and temperature…………………………………………. 50

Texture evolution following superplastic deformation at optimal strain
rate and temperature…………………………………….......................... 53

Grain boundary mis-orientation distribution following superplastic
deformation at optimal strain rate and temperature…………………….. 56
5.3 Orientation diffusion model…………………………………………………….
58
5.3.1 Validation in AA 2004……………………………………………………... 59
6 CONCLUSIONS
61
7 SCOPE FOR FURTHER RESEARCH
62
REFERENCES
63
LIST OF PAPERS BASED ON THESIS
69
CURRICULUM VITAE
70
GRADUATE TEST COMMITTEE
71
vii
LIST OF TABLES
5.1
Compositions of AA 5456 (present study) and AA 5083 (Pérez-Prado et al., 2001)
(in wt. %)…………………………………………………………………………...
23
5.2
Vickers hardness of AA 5456 in as-received and annealed condition……………...
23
5.3
Room temperature mechanical properties of AA 5456…………………………….
24
5.4
Estimation of constants……………………………………………………………..
34
5.5
Experimental and predicted dynamic growth model at 550°C at different strain
rates and strains using Eqn. 5.2……………………………………………………..
34
5.6
Values of kt and p at different temperatures and strain rates……………………….
35
5.7
Composition in wt. % of AA 2004…………………………………………………
42
5.8
Room temperature properties of AA 2004………………………………………….
43
viii
LIST OF FIGURES
2.1
Plot of log stress vs. log ε̇ during high temperature deformation…………………..
4
2.2
The variation of strain rate sensitivity (m) with elongation to failure (∆L⁄ L %),
for Zn-22%Al and Pb-62%Sn (Langdon, 2009)……………………………………
4
The variation of strain rate sensitivity (m) with strain rate ε̇ for Al- Cu eutectic
alloy (Davies et al., 1970)…………………………………………………………..
4
2.4
Rachinger sliding accommodated by (a) slip (b) Lifshitz sliding (Langdon, 2006)..
6
2.5
Illustration of grain switching mechanism, (a) Initial stage before superplastic
deformation, (b) intermediate stage, 4 grain boundaries are in contact, unstable (c)
stable after switching to a different configuration (Gifkins, 1978)………………...
6
‘a’ to ‘c’ the diffusion from inclined faces to horizontal faces; ‘c’ to ‘e’ the
restoration of equiaxed shape by grain boundary migration (Rust and Todd,
2011)………………………………………………………………………………..
7
2.3
2.6
2.7
EBSD map of AA 2004 shows the evolution of equi-axed structure from cold
rolled structure a) 0%, b) 10% c) 65%, d) 100% (Huang et al., 2012)…………….. 11
2.8
The mis-orientation distribution of AA 2004 for the evolution of equiaxed
structure from cold rolled structure (a) 0%, (b) 10%, (c) 65%, (d) 100% (Huang et
al., 2012).…………………………………….…………………………………….. 11
2.9
The illustration of GBS and random rotation of grains (Pérez-Prado et al., 2001)... 13
2.10
The deformation modes based on number of slip system, (a) incompatible for
deformation, (b) compatible for deformation, reorientation of grain, <111> fiber
or {100}<001> formed……………………………………….……………………. 13
2.11
Variation of the ODF, f (g) and orientation spread Ѱo obtained from the model
and the experimental ODF (Engler et al., 2000)…………………………………… 14
2.12
The demonstration of grain rotation and coalescence in grain coalescence model
(Li et al., 1997)……………………………………………………………………... 16
2.13
The dislocation glide and climb mechanism at grain boundaries (Li et al., 1997)… 16
4.1
Geometry of the room temperature tensile testing specimen prepared as per
ASTM E8 standards.………………………….…………..……............................... 20
4.2
(a) A view of high temperature tensile testing machine, and (b) drawing of the
sample for high temperature superplastic testing…………………....…………....... 21
4.3
Overall experimental plan for the current work……………………………………. 22
ix
5.1
Engineering stress vs. engineering strain curve of AA 5456 at room temperature... 24
5.2
Photograph of the specimens of AA 5456 before and after the superplastic
deformation at different strain rates at 550°C……………………………………… 24
5.3
Tensile stress-strain curves of AA 5456 at different strain rates at temperatures of
(a) 450°C, (b) 500°C, (c) 525°C and (d) 550°C (The stress jumps in the above
figure correspond to the strain rate jump tests that were carried out to evaluate the
strain rate sensitivity index ‘m’).…………………………………………………... 25
5.4
Plots of variation of (a) flow stress with temperature and strain rate, and (b)
percentage of elongation with strain rate at different temperatures in AA 5456.
(Data of AA 5083 from literature (Pérez-Prado et al., 2001) are also included for
comparison)………………………………………...………………………………. 26
5.5
Stress vs. strain curve at 550°C for 5×10-4 s-1 in AA 5456, the arrows indicating
the strains to which the intermediate superplastic tensile tests were carried out to
study microstructure development with strain.…………………………………….. 26
5.6
(a) BSE SEM image of AA 5456 ‘as-received’ condition, RD-ND section, (b)
EDAX spectrum from the coarse precipitate in AA 5456…………………………. 27
5.7
The EBSD IPF map of AA 5456 in (a) as-received condition (b) following
annealing at 550°C for 30 min. and (c) following annealing at 550°C for 5 h…….. 28
5.8
{100} pole figure of AA 5456 in (a) as-received condition (b) following
annealing at 550°C for 30 min. and (c) following annealing at 550°C for 5 h…….. 29
5.9
The ODF φ2 = 0° section of AA 5456 in (a) as-received condition (b) following
annealing at 550°C for 30 min. and (c) following annealing at 550°C for 5 h…….. 30
5.10
ODF values of dominant texture components present in as-received and annealed
conditions in AA 5456……………………………………………………………... 30
5.11
Mis-orientation distribution of AA 5456 in (a) as-received condition (b) following
annealing at 550°C for 30 min. and (c) following annealing at 550°C for 5 h…….. 31
5.12
EBSD IPF maps of AA 5456 (a) following annealing at 550°C for 30 min.; and
following superplastic deformation to different strains of (b)100% (0.6 true
strain), (c) 150% (0.9 true strain), (d) 250% (1.25 true strain), (e) 450% (1.7 true
strain), (f) 700% (2 true strain)…………………………………………………….. 32
5.13
(a) Grain growth in RD-TD and RD-ND sections following annealing at 550°C
for 30 min. and following superplastic deformation to different strains (b)
Variation of grain aspect ratio following superplastic deformation to different
strains; in AA 5456………………………………………………………………… 33
x
5.14
Variation of (a) grain size and (b) grain growth rate with strain rate following
superplastic deformation to a true strain of 1 at 550°C in AA 5456………………. 35
5.15
{100} pole figure of AA 5456 (a) following annealing at 550°C for 30 min.; and
following superplastic deformation to different strains of (b) 100% (0.6 true
strain), (c) 150% (0.9 true strain), (d) 250% (1.25 true strain), (e) 450% (1.7 true
strain), (f) 700% (2 true strain)…………………………………………………….. 36
5.16
ODF φ2 = 0° section of AA 5456 (a) following annealing at 550°C for 30 min.;
and superplastic deformation to different strains of (b) 100% (0.6 true strain), (c)
150% (0.9 true strain), (d) 250% (1.25 true strain), (e) 450% (1.7 true strain), (f)
700% (2 true strain)………………………………………………………………... 37
5.17
The variation of ODF value for cube, rotated cube and {110} < 11̅0 >
orientations with superplastic strain in AA 5456…………………………………... 38
5.18
The mis-orientation distributions of AA 5456 (a) following annealing at 550°C
for 30 min.; and following superplastic deformation to different strains of (b)
100% (0.6 true strain), (c) 150% (0.9 true strain), (d) 250% (1.25 true strain), (e)
450% (1.7 true strain), (f) 700% (2 true strain)……………………………………. 39
5.19
EBSD IPF map of AA 5456 following superplastic deformation to a strain of
200% (1 true strain), at 5×10-2 s-1 and at 550°C……………………………………. 40
5.20
{100} pole figure of AA 5456 following superplastic deformation to a strain of
200% (1 true strain), to different strain rates as (a) 5×10-4 s-1 (b) 10-2 s-1 (c) 5×10-2
s-1; at 550°C……………………………………………………………................... 40
5.21
ODF φ2 = 0° section of AA 5456 following superplastic deformation to a strain of
200% (1 true strain), at different strain rates of (a) 5×10-4 s-1 (b) 10-2 s-1 (c) 5×10-2
s-1, at 550°C………………………………………………………………………… 41
5.22
Mis-orientation distribution of AA 5456 following superplastic deformation to a
strain of 200% (1 true strain), at 5×10-2 s-1 and at 550°C………………………….. 42
5.23
Engineering stress vs. engineering strain curve of AA 2004 at room temperature... 43
5.24
Tensile stress-strain curves of AA 2004 at different strain rates at temperatures of
(a) 350°C, (b) 400°C, (c) 457°C. (The stress jumps in the above figure correspond
to the strain rate jump tests that were carried out to evaluate the strain rate
sensitivity index ‘m’).……………………………………………………………… 44
5.25
Plots of variation of (a) flow stress with temperature and strain rate, and (b)
percentage of elongation with temperature and strain rate in AA 2004………….... 44
xi
5.26
Stress vs. strain curve at 457°C for 3×10-3 s-1 in AA 2004, the arrows indicating
the strains to which the intermediate superplastic tensile tests were carried out to
study the microstructure development with strain…………………………………. 45
5.27
EBSD IPF maps of AA 2004 in (a) as-received, (b) following annealing at 457°C
for 30 min. and (c) following annealing at 457°C for 8 hours. The black regions
are corresponding to the regions with CI<0.1……………………………………… 46
5.28
Grain orientation spread maps of AA 2004 in (a) as-received, (b) following
annealing at 457°C for 30 min. and (c) following annealing at 457°C for 8 h…….. 46
5.29
{100} pole figure of AA 2004 in (a) as-received, (b) following annealing at
457°C for 30 min. (c) following annealing at 457°C for 8 h…………………......... 47
5.30
ODF at φ2 = 0° section of AA 2004 in (a) as-received, (b) following annealing at
457°C for 30 min. (c) following annealing at 457°C for 8 h…………………......... 48
5.31
Variation of the ODF values of the texture components {110} < 11̅√2 > and
{100}<011> following annealing at 457°C in AA 2004…………………………... 48
5.32
Mis-orientation distribution of AA 2004 in (a) as-received, (b) following
annealing at 457°C for 30 min., (c) following annealing at 457°C for 8 h and the
red lines correspond to the random Mackenzie distribution……………………….. 50
5.33
EBSD IPF maps of AA 2004 (a) following annealing at 457°C for 30 min.; and
following superplastic deformation to different strains of (b) 25% (0.22 true
strain), (c) 50% (0.4 true strain), (d) 100% (0.6 true strain, (e) 200% (1 true strain)
and (f) 500% (1.7 true strain)………………………………………………………. 51
5.34
Plots of variation of (a) micro-hardness (Hv) at room temperature and (b) grain
size in the as-received, following annealing at 457°C for 30 min. and following
superplastic deformation to different strains in AA 2004………………………….. 52
5.35
{100} pole figure of AA 2004 (a) following annealing at 457°C for 30 min., and
following superplastic deformation to different strains of (b) 25% (0.22 true
strain), (c) 50% (0.4 true strain), (d) 100% (0.69 true strain),(e) 300% (1.3 true
strain), (f) 500% (1.7 true strain)…………………………………………………... 53
5.36
ODF φ2 = 0° section of AA 2004 (a) following annealing at 457°C for 30 min.,
and following superplastic deformation to different strains of (b) 25% (0.22 true
strain), (c) 50% (0.4 true strain), (d) 100% (0.69 true strain), (e) 200% (1 true
strain), (f) 300% (1.3 true strain), (g) 500% (1.7 true strain)……………………… 54
5.37
Variation of ODF corresponding to {110} < 11̅√2 >, {110} < 11̅2 >
and {110} < 11̅1 > orientations following superplastic deformation to different
strains in AA 2004……………………………………………………………….
54
xii
5.38
Variation of orientation spread (Gaussian half width) during deformation
normalized with initial Gaussian half width for both Φ and φ1 directions following
superplastic deformation in AA 2004……………………………………………… 55
5.39
Mis-orientation distribution of AA 2004 (a) following annealing at 457°C for 30
min.; and following superplastic deformation to different strains of (b) 25% (0.22
true strain), (c) 50% (0.4 true strain), (d) 100% (0.69 true strain), (e) 200% (1 true
strain), (f) 500% (1.7 true strain)…………………………………………………... 57
5.40
Variation of SGBs, LABs and HABs fraction (including the symmetric variant of
majour texture components (45°-65° misoriented boundaries) (a) including SGB
fraction (b) excluding SGB fraction following superplastic deformation in AA
2004.……………………………………………………………………….............. 58
5.41
Comparison of experimental ODF of Φ section corresponding to φ2 = 0° and φ1 =
45° of AA 2004 following superplastic deformation to different strains and the
calculated ODF from orientation diffusion model…………………………………. 60
xiii
ABBREVIATIONS
ANN
Artificial Neural Network
BSE
Back Scattered Electron
EBSD
Electron Backscatter Diffraction
EDX
Energy Dispersive X-ray
FEG
Field Emission Gun
GBS
Grain Boundary Sliding
GOS
Grain Orientation Spread
HAB
High Angle Grain Boundary
IPF
Inverse Pole Figure
KAM
Kernel Average Misorientation
LAB
Low Angle Grain Boundary
ND
Normal Direction
ODF
Orientation Distribution Function
PSN
Particle Stimulated Nucleation
RD
Rolling Direction
SEM
Scanning Electron Microscopy
SGB
Sub-Grain Boundary
TD
Transverse Direction
TEM
Transmission Electron Microscopy
xiv
NOTATIONS
ε
True Strain
σ
True Stress in MPa
ε̇ dc
Strain Rate due to Diffusion Creep (s-1)
ε̇ g
Strain Rate due to Dislocation Creep (s-1)
ε̇ gbs
Strain Rate due to GBS (s-1)
b
Magnitude of Burgers Vector (m)
d
Grain Size (m)
D0
Frequency Factor (m2 s-1)
Dl
Lattice Diffusion Coefficient (m2 s-1)
E
Young’s Modulus (MPa)
G
Shear Modulus (MPa)
k
Boltzmann constant
m
Strain Rate Sensitivity
n
Stress Exponent
Ø
Change in Orientation in Degrees
p
Inverse Grain Size Exponent
Q
Activation Energy of Diffusion (J mol-1)
R
Universal Gas Constant (J mol-1K-1)
T
Absolute Temperature (K)
Tm
Melting Temperature (K)
Φ, φ2, φ1
Euler Angles
xv
CHAPTER 1
INTRODUCTION
In the present scenario, it is very important to reduce our energy consumption and improve
the energy efficiency. The energy efficiency can be improved by reducing the weight of the
components in automotive as well as aerospace sectors and this can be achieved by increasing
the use of lightweight materials. Al alloys are important candidates for various industrial,
automobile and marine applications due to their high strength to weight ratio. However, Al
alloys are not commonly used in various structural applications because of their lower
strength. High strength Al alloys are expensive and difficult to form. High strength Al alloys
are still widely used in various airframe components in spite of the high tooling cost as
complicated airframe components are very difficult to weld. It is therefore very important to
develop forming technologies, which can enhance the forming efficiency and produce
complicated parts in a single operation. Superplastic forming is an excellent substitute for
conventional forming technology and meets the above requirements. Superplasticity is the
ability of the material to exhibit very high tensile elongations under certain conditions such as
high temperature, low stress and these materials are generally referred to as superplastic
materials. In aerospace and automotive industry, superplasticity is widely exploited in the
production of complex shapes and curved parts (Barnes et al., 2007).
The superplastic properties of a material are controlled by the microstructure such as grain
size, second phase particles and forming parameters such as strain rate and temperature. For
achieving optimum superplastic flow, the grain size should be ≤ 10 µm and the operating
temperature should be ≥ 0.5 Tm; where Tm is melting temperature. During superplastic flow,
the grain growth needs to be controlled and this achieved by the second phase particles
(Langdon, 2009).
The commonly used superplastic Al alloys are AA 7075, AA 2004, and AA 5083 in order
of decreasing strength. Superplastic Al alloys can be classified as statically or dynamically
recrystallizing. Al alloys. AA 5083 is a statically recrystallizing alloy and an equiaxed
recrystallized grain size conducive to GBS is present right from the beginning of superplastic
deformation. Therefore, it is expected that there will hardly be any anisotropic superplastic
flow in this material. In contrast, in AA 2004 (commercially known as SUPRAL 100)
recrystallizes dynamically during superplastic deformation and anisotropic superplastic flow
is expected to be present often even up to 100% superplastic elongation in this alloy.
1
From a systematic study of the superplastic deformation characteristics of the statically
and dynamically recrystallizing alloys in the initial, transient and the steady state regions of
superplastic flow, it will be possible to clearly distinguish between the operating
mechanism(s) during steady state, isotropic and anisotropic superplastic flow caused by grain
shape and texture anisotropy in the starting material. This has obvious implications for the
superplastic forming because by understanding anisotropic superplastic flow better, there is a
possibility of improving the thickness uniformity in the formed products.
2
CHAPTER 2
LITERATURE REVIEW
Two types of superplasticity are generally observed in nature. They are environmental
superplasticity and structural superplasticity. The environmental superplasticity generally
occurs due to the allotropic phase transitions whereas the structural superplasticity occurs due
to the microstructural features. The favorable microstructural features for structural
superplasticity are (a) equiaxed and fine grains, (b) uniform distribution of second phase
particles and (c) microstructure, which is resistant to cavitation. To achieve superplastic
behavior, the deformation conditions are also very critical. The constitutive relation
connecting strain rate, temperature and stress is given by (Langdon et al., 1994),
𝜀̇ =
𝐴𝐷0 𝐺𝑏 𝑏 𝑝 𝜎 𝑛
𝑘𝑇
𝑄
(𝑑) (𝐺 ) exp (− 𝑅𝑇)
(2.1)
where ‘D0’ is frequency factor, ‘Q’ is activation energy of diffusion, ‘R’ is gas constant,
‘T’ is absolute temperature, ‘G’ is shear modulus, ‘d’ is grain size, ‘b’ is magnitude of
Burgers vector, ‘p’ is inverse grain size exponent, ‘n’ is stress exponent, ‘A’ is
microstructural and mechanism dependent constant, and ‘k’ is Boltzmann constant.
Most of the superplastic materials show sigmoidal variation of stress with strain rate, it
classified into three regions as shown in Fig. 2.1, and the slope of this curve is the strain rate
sensitivity ‘m’. Strain rate sensitivity (m) is an important parameter representing the ability of
the material to resist necking in the absence of cavitation. It is not the characteristic property
of the material but it is an “index of the rate effect on the flow stress corresponding to a
particular testing condition and microstructure” (Ghosh and Hamilton, 1982; Baudelet and
Suery, 1972). The region II shows maximum strain rate sensitivity of ~ 0.5 and maximum
superplastic elongation was observed in this region. The regions III and I exhibit values of m
~ 0.2 and in these regions conventional plasticity is observed. The variations of strain rate
sensitivity (m) with elongation to failure (∆L⁄ L %) or a number of different materials are
shown in Fig. 2.2, and it can be seen that that large elongations are achieved in materials with
higher values of strain rate sensitivity (m) (Baudelet and Suery, 1972). In region II, GBS is
the dominant mechanism and accounts for most of the strain, in region I (low strain rate and
low stress) diffusion creep is predominant and in region III (high strain rate and high stress)
power law dislocation creep is predominant.
3
Figure 2.1: Plot of log stress vs. log ε̇ during high temperature deformation.
Figure 2.2: The variation of strain rate sensitivity (m) with elongation to failure (∆L⁄ L %),
for Zn-22%Al and Pb-62%Sn (Langdon, 2009).
Figure 2.3: The variation of strain rate sensitivity (m) with strain rate ε̇ for Al- Cu eutectic
alloy (Davies et al., 1970).
4
Figure 2.3 shows the variation of ‘m’ with strain rate and temperature for Al-Cu eutectic
alloy which is clearly indicating that there is an optimal temperature and strain rate for
achieving the maximum value of strain rate sensitivity. The initial microstructure also plays
an important role in achieving superplastic properties. As fine grains are required for
achieving superplastic flow, the grain growth during deformation should be restricted and can
be controlled by introducing second phase particles or precipitates.
2.1 Mechanisms of superplasticity
Chokshi (1993) summarized the important microstructural evolution following superplastic
deformation and these are (a) equiaxed grain structure will be formed after few percent of
elongation, (b) at low strain rates strain enhances the grain growth, (c) grain boundary
migration and sliding occurs, (d) limited dislocation and diffusion activity is present, (e) 3
dimensional grain re-arrangement or rotation of grain takes place and (f) in-situ
recrystallization could occur and g) co-operative GBS may take place.
2.1.1
Grain boundary sliding (GBS)
Bell and Thompson (1962) reported the occurrence of GBS during the deformation of high
purity Zn bi-crystals. They found out that the influence of crystallographic slip changes the
shape of grain and reduces the rate of GBS. If the average aspect ratio of grains following
deformation is less as compared to the total strain, then GBS necessarily occurs during the
deformation and it can act as the dominant deformation mechanism (Gifkins., 1976; Caceres
and Wilkinson, 1984; Li et al., 1997). According to Ashby (1972), "sliding without the
change of the boundary structure is possible only if the boundary migrates". It is well
accepted that GBS can act as a dominant and rate-controlling mechanism during superplastic
deformation and its contribution to the total strain is significant (Beere et al., 1977,
Padmanabhan, 1977; Bhattacharya and Padmanabhan, 1989). However, during GBS, grain
compatibility is maintained by concurrent accommodation processes, which may involve
grain boundary migration, grain rotation, and diffusion or dislocation motion (Chokshi,
1993). Most of the models consider the above accommodation mechanisms in conjunction
with GBS. More than one mechanism can also operate simultaneously during superplastic
deformation.
It is widely believed that there are two mechanistically distinct types of GBS. In the first
type of GBS, the grains undergo relative displacement with each other while maintaining
their shape and this referred to as Rachinger sliding (Rachinger 1952). Grains are usually
5
irregularly shaped and therefore Rachinger sliding should necessarily be accommodated by
some intra-granular movement of dislocations within the adjacent grains as shown in Fig.
2.4a (Langdon, 2006). The second type of GBS occurs in Nabarro-Herring and Coble
diffusion creep and it refers to the boundary offsets that develop due to the stress-directed
diffusion of vacancies. This type of GBS is designated Lifshitz sliding (Lifshitz, 1963) is
illustrated in Fig. 2.4b (Langdon, 2006). Though the process of Rachinger and Lifshitz
sliding are mechanistically different, they produce similar offsets at grain boundaries.
Therefore, care must be taken to distinguish unambiguously between these two mechanisms.
It should be noted that the grains are retained their shape following Rachinger sliding while
the grains are elongated along tensile following Lifshitz sliding.
Figure 2.4: Rachinger sliding accommodated by (a) slip (b) Lifshitz sliding (Langdon, 2006).
The possible mechanisms for GBS are grain translation, grain rotation, neighbor grain
switching (see Fig. 2.5). Here, grains are considered as a ‘non-deforming rigid crystals’, and
the total strain is produced by the rearrangement of grains during deformation and this
analogous to Rachinger sliding as mentioned above.
Figure 2.5: Illustration of grain switching mechanism, (a) Initial stage before superplastic
deformation, (b) intermediate stage, 4 grain boundaries are in contact, unstable
(c) stable after switching to a different configuration (Gifkins, 1978).
6
In general, the total strain rate in a material subjected to stress at high temperatures will be
the sum of the strain rates contributed by each process i.e.
ε̇ total = ε̇ g + ε̇ gbs + ε̇ d c
(2.2)
where ε̇ g is the strain rate due to dislocation creep, ε̇ gbs is the strain rate due to GBS, ε̇ dc is
the strain rate due to diffusion creep. The dominant mechanism will vary according to the
strain rate and temperature, which in turn determines the ‘strain rate sensitivity index’.
Diffusion creep is the ‘stress-driven diffusional transport from one grain boundary to
another and occurs irrespective of the geometry of grains. The mechanisms proposed by
Herring (1950), Coble (1963), Ashby and Verrall (1973), Lee/Spingarn and Nix (1978), and
Hazzledine and Schneibel (1993) were generalized as accommodations processes by
diffusion creep. The grain boundary migration and grain neighbour switching can take place
with the help of diffusional flow. Lee-Spingarn and Nix (1978) proposed a mechanism in
which grains are elongated by diffusional transport during initial stages of the superplastic
deformation (Figs. 2.6a-c), and after a particular true strain (0.5) (Fig. 2.6c) the four grain
boundaries would meet (Fig. 2.6c) to form an unstable grain boundary structure and this
would lead to the restoration of the equiaxed shapes by grain boundary migration (shown in
Figs. 2.6d-e). Rust and Todd (2011) suggested that the ‘stress directed’ diffusional transport
and grain boundary migration leads to the grain boundary switching and restoration of
equiaxed shape following superplastic deformation and this mechanism can be responsible
for the high superplastic elongation observed in AA 5083.
Figure 2.6: ‘a’ to ‘c’ the diffusion from inclined faces to horizontal faces; ‘c’ to ‘e’ the
restoration of equiaxed shape by grain boundary migration (Rust and Todd,
2011).
The GBS could also be accommodated by dislocation creep (Chaudhari, 1967; Hayden et
al., 1967; Ball and Hutchinson, 1969; Valiev and Langdon, 1993; Li et al., 1996). If the
smooth GBS is obstructed by the precipitates or triple points, stress concentration present at
7
these points is relieved by the formation of dislocations and these dislocations generated
eventually pile up against opposite boundaries and it creates back stress, which prevents the
further sliding. The dislocation climb along with the GBS will reduce the dislocation pile up,
and the back stresses are reduced which facilitates further sliding. The dislocation activities
are confined to the region close to the grain boundaries and which in turn responsible for the
change in the shape of the grain following superplastic deformation (Gifkins, 1976; Mayo
and Nix, 1989; Padmanabhan and Davies, 1980). Some researchers proposed that the slip
could also take place along with diffusion during superplastic deformation (Kaibyshev, et al.,
1978; Sherby and Wadsworth, 1989). Li et al. (1996) observed dislocation activity in Al- Mg
alloys using transmission electron microscope and postulated that the occurrence of GBS in
the superplastic regime could be due to the movement of dislocations, and the observed
dislocation density would be more in larger grains at lower strain rates, and the fraction
dislocation free grains increases as strain rate increases. At high strain rate, the dislocation
cell structure also observed in Al-Mg alloy.
A random behaviour of grain rotation during GBS observed, and phenomenon is expected
to be responsible for the texture reduction following superplastic deformation. The presence
of the frictional forces at grain boundaries causes the grain rotation during GBS (Engler et al.,
2000). However, Bate et al. (2005, 2007) observed the persistence of orientation correlation
and retention of the banded structure even after considerable superplastic strain in
dynamically recrystallizing alloys which is not consistent with the ‘randomness’ of grain
rotation following superplastic deformation. Rust and Todd (2011) postulated that grain
rotation is not necessarily required for grain neighbor switching mechanism, and Lifshitz
sliding mechanism alone cannot generate sufficient superplastic strain. In addition, the intragranular straining was not observed following superplastic deformation. So grain
rearrangement and grain switching necessarily needed present for accommodating the
superplastic flow (Sotoudeh and Bate, 2010).
Zelin et al. (1991) postulated the possibility of co-operative GBS. Grains slide as a group
and it leads to the ‘sequential shear’ of grains and ‘grain boundary migration’ include grain
rotation, sliding, intra-granular deformation as well as ‘grain migration’ (Zelin et al., 1993).
Padmanabhan and Schlipf (1996) proposed a model in which the co-operative GBS is
coupled with diffusion, which could act as the rate controlling mechanism, occurs at both
regions IIa (lower stress range of region II) and I of superplastic flow. Shear bands observed
8
during optimal superplastic flow was reported in AA 2004 is an evidence for co-operative
GBS (Astanin et al., 1996).
2.1.2
Diffusion creep
At very low strain rates and very low stresses, diffusion creep can act as the dominant
deformation mechanism and it may result in the formation of elongated grain structure. If
grain elongation by diffusional transport is proportional to macroscopic strain, it is called as
Lifshitz sliding. Lifshitz siding can act as an accommodation mechanism for diffusion
process and its contribution to the total strain is believed to be negligible. Backofen (1968)
and Karim (1969) noted that Coble creep and GBS could be acting simultaneously, for that, a
threshold stress is required and the importance of threshold stress is more significant when
grain size is small (Edington et al., 1976). Raj and Ghosh (1981) proposed that strain rate
sensitivity (m) will be small if the GBS accommodated by diffusion creep. It is also observed
that some alloys like Al–Cu–Zr and the Al–Mg–Mn exhibit superplasticity only if the solute
concentration is higher. If the solute concentration is higher, it could enhance diffusion creep
and promote Lifshitz sliding which can act as an accommodation mechanism (Bate et al.,
2010). Bate et al. (2010) studied a statically recrystallizing Al alloy AA 5083 and postulated
that Lifshitz sliding is more predominant than grain neighbor switching mechanism in the
optimum superplastic regime.
2.1.3
Dislocation creep
Dislocations are generated during plastic deformation and the dislocation slip process
includes dislocation glide and dislocation climb during deformation (Chaudhari, 1967;
Hayden et al., 1967). The strain rate sensitivity for dislocation glide is ~0.3 and that for
dislocation climb is ~0.2.
For Al, the strain rate produced by dislocation creep when the creep rate is controlled by
dislocation climb is given as,
𝜀̇𝑔 =
1011 𝐷𝑙 𝜎 5
𝑏2
(𝐸 )
(2.3)
where Dl is the lattice diffusion coefficient (m2 s-1), σ is the stress (MPa), E is the
unrelaxed Young’s modulus (MPa), and b is the Burgers vector (m). If the dislocations are
dragged by solute atoms in dislocation creep, it is known as ‘solute drag creep’. The strain
rate sensitivity for dislocation climb controlled creep is 0.2 and in this regime, the maximum
9
elongation is achieved ~100%, whereas the solute drag creep has strain rate sensitivity of 0.2
to 0.3 and relatively higher percentage of elongation is obtained (Pérez-Prado et al., 2001).
2.2 Recrystallization and superplasticity
As was mentioned before, superplasticity is exhibited by materials with fine grain sizes (< 10
µm) which are stable (resist grain growth) during high temperature deformation. Most of the
commercial Al alloys are ‘pseudo-single-phase’ i.e., they consist of Al matrix with small
fraction of dispersed second phases. Superplastic microstructures in Al alloys are produced
by thermo mechanical processing (involving solution treatment, ageing treatment and cold or
warm working steps) based on the mechanism of PSN during static recrystallization or by
dynamic recrystallization in which fine grain size develops during the early stages of hot
deformation (Humphreys and Hatherly, 1995). In statically recrystallizing alloys (for e.g. AA
7075, AA 5083), the alloy is solution heat treated followed by over ageing treatment to
produce large intermetallic particles and subsequently cold/warm rolled and recrystallized to
produce fine grain size with nearly random texture. PSN occurs at the large overaged
particles during recrystallization annealing resulting in fine recrystallized grain size. The
grain growth during superplastic deformation is inhibited by large and small particles.
Fine grained microstructure in Al alloys is also obtained through dynamic recrystallization
phenomena. AA 2004, Al- 5%Ca-5%Zn are the examples for dynamically recrystallizing
alloys. AA 2004 is rapidly chilled during casting to avoid precipitation of ZrAl3 and
subsequently heat treated at 350°C to form fine dispersion of ZrAl3 and subjected to heavy
warm working and final cold reduction is given. The heavily rolled sheet is superplastic at a
temperature of ~ 460°C. The role of ZrAl3 particles is to inhibit recovery and suppress static
recrystallization of cold worked material and to prevent grain growth after dynamic
recrystallization. (Homogenization of microstructure was reported following annealing at
480°C for 8h. in AA 2004 (Pérez-Prado and McNelley, 1999)) The recrystallization during
superplastic deformation can be of continuous type; i.e. the deformation-induced
microstructures with a large fraction of HABs are produced by the short-range migration of
HABs accompanied by dislocation rearrangements during deformation. Different
mechanisms have been proposed to explain the superplastic behavior of dynamically
recrystallizing materials (Pérez-Prado et al., 2003). Initial banded microstructure with a large
fraction of low angle boundaries (LABs) is not favorable for superplastic deformation.
During dynamic recrystallization it is observed that the average mis-orientation angle is
10
increased and it was suggested that it can be due to the sub-grain growth (Nes et al., 1985),
sub-grain coalescence (Hornbogen., 1979), sub-grain rotation (Gudmundsson et al., 1991) or
accumulation of dislocations inside the grain (Tsuzaki et al., 1996). Some researchers
proposed that during initial stage of deformation, the crystallographic slip could act as the
dominant mechanism (Padmanabhan et al., 1991; Liu and Chakrabarti, 1996; Pérez-Prado et
al., 1998). It was also postulated that that the non-equiaxed or elongated grains would
become equiaxed due to the dislocation motion during crystallographic slip.
Figure 2.7: EBSD map of AA 2004 shows the evolution of equi-axed structure from cold
rolled structure a) 0%, b) 10% c) 65%, d) 100% (Huang et al., 2012).
Figure 2.8: The mis-orientation distribution of AA 2004 for the evolution of equiaxed
structure from cold rolled structure (a) 0%, (b) 10%, (c) 65%, (d) 100% (Huang
et al., 2012).
11
McNelley et al. (1997, 2002) studied AA 2004 and concluded that there is no evidence for
the formation and migration of HABs during recrystallization. In this alloy, it is observed that
‘a recovery dominated process in which continuous formation and migration 5° to 15° misoriented boundaries’ are observed. It is also found that there is no evidence for the conversion
of SGBs (2°-5°) to HABs whereas the 5° to 15° mis-oriented boundaries converted to high
angle boundaries during the deformation and this resulted in the breaking of the banded
structure during deformation (McNelley et al., 2004). However, Humphreys et al., 2005
observed the formation of HABs after dynamic recrystallization i.e., the fraction of HAB is
increased with increasing strain and after a true strain of 0.5 the material has a microstructure
suitable for superplasticity. Figure 2.7 is the EBSD map of AA 2004 at different strains,
which demonstrates the microstructure development during different stages of superplastic
deformation. From the microstructure, it is evident that HAB spacing is found to be increased
during superplastic deformation. The numbers of HABs per unit area are also decreased
during deformation. Figure 2.8 shows the mis-orientation distribution and it can be seen that
the increase in strain causes a reduction in LABs (Huang et al., 2004, 2012).
2.3 Texture and superplasticity
Texture analysis has been widely used to study the deformation mechanisms during
superplastic flow. Overall a reduction in texture following superplastic flow is observed
(Cutler et al., 1974; Melton et al., 1974; Matsuki et al., 1977; Engler et al., 2000) From
texture analysis, Padmanabhan et al. (2000) postulated that the observed texture annihilation
during superplastic flow is due to the random rotation of grains during superplastic
deformation and the formations of random orientations after random grain rotation followed
by slip are shown in Fig.2.9 (Pérez-Prado et al., 2001). Figure 2.10a shows the condition
when slip systems are <5 which is incompatible for deformation by slip. If slip systems are
≥5, dislocation creep can occur and it strengthens some texture components during
deformation depending on the slip system activated at that temperature. Melton et al. (1974)
observed the occurrence of crystallographic slip in Zn- 40(wt. %) Al eutectoid alloy even at
maximum strain rate sensitivity. He observed the stabilization of {110}<001> orientation
during deformation. Generally {100}<001> cube orientation will be developed in the
material under dislocation creep condition from the starting {100}<0vw> orientations. The
<111> fiber also formed by the reorientation of grains from random arrangement of grains
under dislocation creep conditions as shown in Fig. 2.10b (Pérez-Prado et al., 2001). The
formation of weak <111> fiber and weak {100}<001> components were reported in AA
12
5083 during uniaxial straining which is consistent with the dislocation activity along with
GBS.
Figure 2.9: The illustration of GBS and random rotation of grains (Pérez-Prado et al., 2001).
Figure 2.10: The deformation modes based on number of slip system, (a) incompatible for
deformation, (b) compatible for deformation, reorientation of grain, <111> fiber
or {100}<001> formed.
Dynamically recrystallizing materials are generally exhibiting crystallographic slip during
the initial stages of deformation and some texture components are developed due to the slip
process. Humphreys and Hatherly (1995) observed that {110} < 11̅2 > Brass’ component is
a stable texture component during plane strain compression and retention of brass component
at initial stages of deformation is analogous to plane strain compression of banded
microstructure during deformation In Al-5%Ca-5%Zn alloy, the major texture component in
as-received condition is {225}<554> and it changes to {110} < 11̅2 > Brass’ component
during initial stages of superplastic deformation which is consistent with texture evolution
during plane strain compression (Pérez-Prado et al., 2003). Bate et al. (2005) proposed a
13
model for the observed texture changes during superplastic deformation in which rate
sensitive slip and dynamic grain growth are considered and the effects of grain translation,
random grain rotation etc. are neglected. By using this model, the texture changes up to a true
strain of 1 were predicted correctly in AA 2004 (Bate et al., 2005).
Engler et al. (2000) proposed a model in which grains rotate randomly during GBS (due to
the effect of frictional forces) causing orientation spread and texture reduction. Grains are
assumed to be rigid and this texture change during superplastic deformation are treated as
analogous to the concentration changes during diffusion with time and is termed as
‘orientation diffusion’ during superplastic deformation. Consider the equation,
𝑗=
𝜕𝑓(𝑔) 𝜕𝑓(𝑔)
𝜕∅
;
𝜕(𝜀)
=𝑆
𝜕2 𝑓(𝑔)
𝜕∅2
(2.4)
∅ = changes in orientation ε = true strain, S is a constant.
If we are considering the two dimensional isotropic orientation diffusion coefficient and
uniaxial state of stress, the solution is obtained as,
𝑓(𝜀, ∅) =
𝐹
2
(4𝜋𝑆√𝜀)exp(−∅ ⁄𝑆𝜀 )
(2.5)
∅ = changes in orientation ε = true strain, S is a constant.
Figure 2.11: Variation of the ODF, f (g) and orientation spread Ѱo obtained from the model
and the experimental ODF (Engler et al., 2000).
Figure 2.11 shows the variation of the ODF f (g) and orientation spread Ѱo obtained from
the model for main texture component and the ODF value obtained from experiment at
14
different strains. The model was able to predict the texture changes with deformation
accurately.
2.4 Grain growth during superplasticity
Baudelet and Suery (1972) considered the superplastic deformation as non-strain hardening
process in which grain size remains constant and the flow stress is the function of initial
temperature, true strain and grain size. Later on, the deformation enhanced grain growth has
been reported following superplastic deformation (Caceres and Wilkinson, 1984; Chokshi.,
1993). Li et al. (1997) postulated that the low strain sensitivity produce sharp necks and to
prevent the material from necking due to low strain rate sensitivity, strain hardening and
grain growth take place that enhances the ductility with strain.
Experimental results indicate that the grain size increases linearly and this is described by
the relation,
∆d = k t ε ε̇ −p
(2.6)
where k t is the rate constant, and ∆d is the grain growth with respect to static annealing.
As strain rate decreases, the grain growth per unit strain increases and at high strain rate, the
grain growth was generally observed to be less (Caceres and Wilkinson, 1984). Dynamic
anisotropic grain growth were noticed in Al-Mg-Mn-Cu alloy during uniaxial tensile studies
(Li et al., 1997). Different theories are available account for this deformation enhanced grain
growth. It could be due to the creation of vacancies during GBS enhances diffusion rate as
well as grain growth rate , and the ‘strain induced’ grain growth increases the stability of
microstructure at low strain rates by reducing the possibility of necking (Clark and Alden,
1973; Senkov and Myshlyaev, 1986). Sato et al. (1991) postulated that the merging of grains
occurs during grain rotation and GBS could accelerate the dynamic grain growth. Caceres
and Wilkinson (1984) proposed that some damaged zones are created during GBS, these are
accommodated by grain rotation and grain switching, which in turn enhances the grain
growth following superplastic deformation. The plastic flow induced along the curvature to
the damaged edges can account for abnormal grain growth during GBS (Verma et al., 1996).
The enhanced grain growth in lateral or transverse direction can be explained either by grain
coalescence during grain rotation as shown in Fig. 2.12 or by dislocation climb at grain
boundaries as shown in Fig. 2.13 (Li et al., 1997).
15
Figure 2.12: The demonstration of grain rotation and coalescence in grain coalescence model
(Li et al., 1997).
Figure 2.13: The dislocation glide and climb mechanism at grain boundaries (Li et al., 1997).
This theory cannot predict the grain growth during continuous recrystallization. During
continuous recrystallization, the dislocations are recovered to SGBs, which increases the misorientation and the microstructure conductive to superplastic deformation after sufficient
straining. During this process, the coincidence-site-lattice boundaries (CSL) could be shifted
to higher values (Chokshi, 1993).
2.5 AA 5XXX alloys
AA 5XXX alloys are strain hardenable Al alloys, these possess moderate strength compared
with other Al alloys, and the mechanical properties can be enhanced by solid solution
strengthening. These alloys possess good corrosion resistance, good thermal conductivity,
good electrical conductivity (Gungor et al., 2014). Some of these alloys for e.g. AA 5083,
16
exhibit PSN during annealing. The fine Al6 (Mn, Cr) precipitates present in the material act
as the source for the PSN and this restricts grain growth following recrystallization. In AA
5XXX alloys, Mg is a major alloying addition that controls the mechanical properties. Si
improves fluidity; Mn and Cr improve the corrosion resistance. AA 5083 is a well-known
superplastic alloy widely used in marine, aerospace and automotive applications. AA 5456 is
a variant of AA 5083 with higher Fe and Mg content and lower Mn content.
2.6 AA 2XXX alloys
2XXX Al alloys are heat treatable alloys, possessing good mechanical properties and having
wide spectrum of applications especially in aerospace industry. Some of these alloys (for e.g.
AA 2004) exhibit continuous dynamic recrystallization during deformation. The Zr is added
to the material to form fine precipitate that can pin the grains to prevent static
recrystallization and grain growth. Al-Cu-Zr (Supral-100) is a commercial superplastic
aluminium alloy widely used in aerospace applications (Barnes et al., 2007).
17
CHAPTER 3
OBJECTIVES AND SCOPE
A systematic study of the superplastic deformation characteristics of the statically and
dynamically recrystallizing alloys in the initial, transient and the steady state regions of
superplastic flow, will enable to distinguish between the operating mechanism(s) during
steady state, isotropic and anisotropic superplastic flow caused by grain shape and texture
anisotropy in the starting material. This has obvious implications for the superplastic forming
because by understanding anisotropic superplastic flow better, there is a possibility of
improving the thickness uniformity in the formed products. AA 5456 recrystallize statically
and an equiaxed recrystallized grain size conducive to GBS act as the dominant and rate
controlling mechanism during steady state superplastic flow expected to present right from
the beginning of superplastic deformation. Therefore, it is expected that there will hardly be
any anisotropic superplastic flow in this material. However, some weak cube texture has been
reported during deformation due to the influence of crystallographic slip. In contrast, in AA
2004, it recrystallizes dynamically during superplastic deformation and anisotropic
superplastic flow is expected to present up to 100% of the superplastic elongation in this
alloy. After 100% deformation, the random grain rotation expected to present during
superplastic deformation and it is important to validate the ‘orientation diffusion’ model
proposed by Engler et al. (2000). It is imperative to study the texture development during
superplastic deformation to understand possible mechanisms during deformation, since there
is no consensus on definite mechanism that account for superplastic deformation for these
alloys.
The main objectives of the present investigation are:

To optimize the processing parameters (strain rate, temperature) to achieve maximum
superplastic elongation in AA 5456 (a new alloy variant developed by the Hydro
Aluminum GmbH, Germany).

Comparison of the superplastic property of AA 5456 with AA 5083 (a composition
from literature) and the modified AA 5083

Study of the texture and microstructure development during superplastic deformation
of aluminum alloys AA 5456 and AA 2004.

Validate the orientation diffusion model proposed by Engler et al. (2000) for AA 2004
to predict the nature of grain rotation during deformation.
18
3.1 Organization of thesis
A brief introduction of the current problem is outlined in Chapter 1. The literature review of
the current problem is presented in Chapter 2. The outlines the scope and objective of the
present work are included in Chapter 3. The experimental procedures and techniques used in
this work are explained in Chapter 4. The results and discussion of the investigation are
presented in Chapter 5. It is divided into two sections. The first section deals with the
statically recrystallizing aluminium alloy AA 5456 and the second section deals with the
dynamically recrystallizing aluminium alloy AA 2004. The summary and conclusions are
described in Chapter 6. Scopes for further research are described in the chapter 7.
19
CHAPTER 4
EXPERIMENT DETAILS
4.1 Materials
Al-Mg-Mn alloy AA 5456 and Al-Cu alloy AA 2004 were investigated in the current work.
Materials were received as rolled sheets with the sheet thickness of 2 mm for AA 5456 and 3
mm for AA 2004. AA 2004 were cladded with pure aluminium on both sides of the plate at a
thickness of 300 µm.
4.2Room temperature tensile testing
The tensile testing samples were prepared according to ASTM E8 standards using CNC
milling machine and the tests were carried out in an Instron® machine at 10-3 s-1 strain rate.
The dimensions of the tensile testing sample are shown in Fig. 4.1.
Figure 4.1: Geometry of the room temperature tensile testing specimen prepared as per
ASTM E8 standards.
4.3 High temperature tensile testing
High temperature tensile tests were carried out using a microprocessor controlled electromechanical screw driven tensile testing machine (Schenck Trebel ®) with a three zone tubular
furnace attached to it (Fig. 4.2a). The temperature in each zone was measured using a Pt-Rd S
type thermocouple coupled with a temperature controller (Eurotherm®) for maintaining
almost uniform temperature within ±3°C inside the furnace. The tensile testing specimens
were prepared by using CNC milling machine and its dimensions are shown in Fig. 4.2b.
20
(a)
(b)
Figure 4.2: (a) A view of high temperature tensile testing machine, and (b) drawing of the
sample for high temperature superplastic testing.
4.4 Hardness
Micro-hardness of the given alloys was measured using Wolpert Wilson Vickers Hardness
Tester® at a load of 100 g and dwell time of 10 s.
4.5 Scanning electron microscopy (SEM) studies
The samples were metallographically polished by using alumina paste of 1 µm and then with
colloidal silica of 0.05 µm to get smooth scratch-less surface. The specimens were examined
using Quanta 200 FEG equipped with BSE, EDX detectors, and with Quanta 400 FEG
equipped with BSE detector.
4.6 Electron back scatter diffraction (EBSD) studies
The samples were metallographically polished followed by diamond polishing up to 1 µm
and then it was subjected to electro-polishing. The electro-polishing was done in a solution of
perchloric acid, ethanol, butoxyethanol and water (8:73:10:9) at 22°C and 40V DC for 30 s.
AA 2004 was first polished using Struers® automatic polishing machine and then the sample
was electro-polished using the same parameters as mentioned above. High speed Hikari®
EBSD Camera was attached to FEI Inspect F SEM® was used to capture kikuchi patterns
used for EBSD scanning. The SEM was operated at 20 kV at a working distance of 16 mm.
The measurement was taken at 70° tilt and the tilt corrections were made by TSL-OIM®
software during scanning. The step size was used as 1 µm for recrystallized AA 5456, 0.15
µm for rolled AA 5456 and 0.3 µm for AA 2004.
21
4.7 The overall experiment plan
Figure 4.3 represents the sequence of experiments that have been conducted for the current
work.
Figure 4.3: Overall experimental plan for the current work.
22
CHAPTER 5
RESULTS AND DISCUSSIONS
5.1 Statically recrystallizing alloy, AA 5456
5.1.1

Base metal characterization
Composition
Composition of the alloy AA 5456 (in wt. %) (was determined using optical emission
spectrometry) is given in Table 5.1. For comparison, the composition of AA 5083 is also
included in the Table 5.1. AA 5456 is a variant of AA 5083 developed by Hydro Aluminum
GmbH, Germany. In comparison to AA 5083, AA 5456 has higher Mg and Fe content and
lower Mn content.
Table 5.1: Compositions of AA 5456 (present study) and AA 5083 (Pérez-Prado et al., 2001)
(in wt. %).
Elements
Si
Fe
Mn
Mg
Cr
Al

AA 5456
0.11
0.837
0.315
5.98
0.195
92.4
AA 5083 (Pérez-Prado et al.,
2001)
0.05
0.07
0.65
4.48
0.11
94.6
Micro hardness
The hardness of AA 5456 in as-received and following annealing at 550°C for 30 min. is
given in Table 5.2. It is clear that there was a significant reduction in hardness following
annealing, this could be attributed to recrystallization and this is further discussed in section
5.1.4.
Table 5.2: Vickers hardness of AA 5456 in as-received and annealed condition.
Condition
As-received
Annealed at 550°C for 30 min.
5.1.2
Hardness (HV 10)
120 ± 5
85 ± 5
Room temperature mechanical properties of AA 5456
The yield strength, ultimate tensile strength and the elongation to fracture are listed in Table
5.3 The room temperature stress strain curve for AA 5456 in as-received condition is shown
Fig. 5.1
23
Figure 5.1: Engineering stress vs. engineering strain curve of AA 5456 at room temperature.
Table 5.3: Room temperature mechanical properties of AA 5456.
(0.2%) Yield stress
Ultimate tensile strength (UTS)
Elongation to fracture
5.1.3
350 MPa
477 MPa
9.3%
High temperature tensile behaviour of AA 5456
To idetify the parameters for achieving the optimal superpalstic flow, the alloy AA 5456 was
tested at different temperatures (450°C, 500°C, 525°C, 550°C and 575°C) and different strain
rates (2×10-4 s-1, 5×10-4 s-1, 1×10-3 s-1, 5×10-3 s-1, 1×10-2 s-1 and 5×10-2 s-1). Figure 5.2 shows
the photograph of samples following tensile testing at 550°C at different strain rates.
Figure 5.2: Photograph of the specimens of AA 5456 before and after the superplastic
deformation at different strain rates at 550°C.
Figures 5.3a-d shows the representative tensile stress-strain curves for tests conducted at
different temperatures at different initial strain rates. The stress-strain curves show that there
is no steady state regime observed at strain rates ≥ 10-3 s-1, whereas the steady state can be
clearly identified at strain rates < 10-3 s-1. The strain rate sensitivity index (m) was evaluated
24
using strain rate jump tests (see the stress jumps in Figs. 5.3a-d) and the calculated values are
in the range of 0.5-0.6 at strain rates < 10-3 s-1 and 0.2 at strain rates > 10-2 s-1. It is well
established that a material exhibits better superplastic properties when the strain rate
sensitivity index is ≥ 0.4-0.5 (Padmanabhan and Davies, 1980). Consistent with the above
observation, in AA 5456 highest elongation of 50070% was obtained at 550°C at a strain
rate of 5 × 10-4 s-1 (Fig. 5.3 d).
Figure 5.3: Tensile stress-strain curves of AA 5456 at different strain rates at temperatures of
(a) 450°C, (b) 500°C, (c) 525°C and (d) 550°C. (The stress jumps in the above
figure correspond to the strain rate jump tests that were carried out to evaluate the
strain rate sensitivity index ‘m’).
Figure 5.4a shows the variations of flow stress (maximum or steady state stress taken from
Figs. 5.3a-d) with initial strain rate at different temperatures. The variation of percentage of
elongation with temperature and strain rate is shown in Fig. 5.4b. The data for AA 5083
tested at 535°C (data taken from literature, Pérez-Prado et al., 2001) is also included in the
plots for comparison. It is found that in AA 5456, the maximum elongation of 500±70% was
obtained at a temperature of 550°C and at a strain rate of 5×10-4 s-1. At higher and lower
temperatures and strain rates, the elongation was found to be reduced. From the above results,
the optimal temperature and strain rates were identified to be 550°C and 5×10-4 s-1
25
respectively. The flow stress (maximum/steady state) results also indicate that the stress
required for deformation of AA 5456 is lower in comparison with AA 5083 (Fig. 5.3a). The
elongations obtained in AA 5456 (at 525°C and 550°C) were much higher compared with
those reported for AA 5083 in the range of strain rates employed in this study (Fig.5.4b). It
can therefore be concluded that AA 5456 exhibits superior superplastic properties than AA
5083.
(b)
(a)
Figure 5.4: Plots of variation of (a) flow stress with temperature and strain rate, and (b)
percentage of elongation with strain rate at different temperatures in AA 5456.
(Data of AA 5083 from literature (Pérez-Prado et al., 2001) are also included for
comparison).
Figure 5.5: Stress vs. strain curve at 550°C for 5×10-4 s-1 in AA 5456, the arrows indicating
the strains to which the intermediate superplastic tensile tests were carried out to
study microstructure development with strain.
26
After establishing the optimal superplastic conditions, detailed microstructure and texture
studies using EBSD were carried out in the unsteady and steady state deformation regimes by
subjecting different samples to different superplastic strains (shown by arrows in Fig. 5.5)
and these results are discussed below.
5.1.4
Microstructure and texture evolution following annealing at 550°C
The microstructure of AA 5456 is consists of fine and coarse precipitates (Fig. 5.6a). Energy
dispersive spectroscopy in the SEM was used to identify the composition of coarse
precipitates (Fig. 5.6b). The relatively coarser precipitates were found to be the Al-Fe-Mn
type. The coarse precipitates play a critical role during recrystallization through particle
stimulated nucleation (PSN). PSN in Al, Fe, Cu and Ni alloys is very commonly observed to
occur at particles of diameter > 1 µm. Many commercial Al alloys contain large second-phase
particles and PSN is a very common recrystallization mechanism in these alloys. It has been
shown that if PSN is the only nucleation mechanism (which is often the case in materials with
a number of large second phase particles), the resulting recrystallization texture will be weak.
While the large particles act as the sites for recrystallization nuclei, the finer particles pin the
migrating grain boundaries through Zener pinning and suppress grain growth (Humphreys
and Hatherly, 1995). The observed microstructure and texture in AA 5456 after annealing can
be rationalized based on of PSN and Zener pining effects.
Figure 5.6: (a) BSE SEM image of AA 5456 ‘as-received’ condition, RD-ND section, (b)
EDAX spectrum from the coarse precipitate in AA 5456.
The inverse pole figure (IPF) maps obtained from the EBSD data analysis of ‘as-received’
and annealed (at 550°C for 30 min. and 5 h) samples are shown in Figs. 5.7a-c. The grain
boundary maps are also superimposed on the IPF maps. Different types of grain boundaries
are represented with different colour codes. The 2°-5° misoriented boundaries are highlighted
27
with yellow lines, 5°-15° misoriented boundaries are highlighted with red lines and 15°-60°
misoriented boundaries are highlighted with black lines. The EBSD scans in as-received
condition performed on RD-ND plane and for analyzing texture in the RD-TD plane; the data
are rotated by 90° about RD. The data points with confidence index (CI) values < 0.1 are not
considered for analysis and these are represented as black pixels in the IPF map. IPF maps
reveal a banded microstructure and the presence of strong texture in the as-received material.
Following annealing at 550°C for 30 min. the microstructure was equiaxed with
predominantly HABs and the texture is considerably weakened. The grain size following
annealing at 550°C for 30 min. was 9 ± 4 µm (Fig. 5.7b) and for 5h it was 13±4 µm (Fig.
5.7c), indicating a moderate grain growth following prolonged annealing for 5h.
T
D
TD
Figure 5.7: The EBSD IPF map of AA 5456 in (a) as-received condition (b) following
annealing at 550°C for 30 min. and (c) following annealing at 550°C for 5 h.
28

Texture evolution following annealing at 550°C
The as-received AA 5456 alloy shows strong ‘Brass, {110} < 11̅2}’, {110} < 11̅√2}
components. These are the typical texture components observed following cold rolling of
medium/high stacking fault energy (SFE) alloys. The above-mentioned rolling texture
components were weakened considerably and a weak cube and rotated cube texture
components were appeared following annealing at 550°C for 30 min. The calculated {100}
pole figures from EBSD data for as-received and annealed samples are shown in Figs. 5.8a-c.
The {100} pole figure of as-received material (Fig. 5.8a) confirms the presence of strong
texture with clear Brass, Copper components and is weakened following annealing at 550°C
for 30 min. (Fig. 5.8b) and 5 h (Fig. 5.8c) The orientation distribution functions (ODFs) were
calculated using series harmonic expansion method and the φ2 = 0 sections are plotted in
Figs. 5.9a-c. Quantitative analysis of texture components was performed from the ODF
analysis and this is shown in Fig. 5.10. Following annealing for 0.5h, the {110} <
11̅2}’, {110} < 11̅√2}
component decreased texture almost randomized, but following
annealing for 5h, the strength of cube component slightly increased. (Fig. 5.10).
Figure 5.8: {100} pole figure of AA 5456 in (a) as-received condition (b) following
annealing at 550°C for 30 min. and (c) following annealing at 550°C for 5 h.
29
(b)
(a)
(c)
Figure 5.9: The ODF φ2 = 0° section of AA 5456 in (a) as-received condition (b) following
annealing at 550°C for 30 min. and (c) following annealing at 550°C for 5 h.
Figure 5.10: ODF values of dominant texture components present in as-received and
annealed conditions in AA 5456.

Grain boundary mis-orientation distribution following annealing at 550°C
Figure 5.11 shows the mis-orientation angle distribution in the as-received condition and
following annealing at 550°C for 30 min. and 5 h. The as-received alloy has a large fraction
of LABs (Fig. 5.11a). Following annealing at 550°C the fraction of high angle boundaries
was increased (signifying recrystallization) and the mis-orientation distribution was very
close that obtained in randomly oriented poly-crystals, which are shown as dotted lines in
Figs. 5.11b and c.
30
(a)
(b)
(c)
Figure 5.11: Mis-orientation distribution of AA 5456 in (a) as-received condition (b)
following annealing at 550°C for 30 min. and (c) following annealing at 550°C
for 5 h.
5.1.5
Microstructure and texture evolution following superplastic deformation at
optimal strain rate and temperature
The EBSD IPF maps taken from samples subjected to different strains at the optimal
superplastic conditions are shown in Figs. 5.12(a-e). The grain boundaries are also marked on
Fig. 5.12 with the same color scheme employed in Fig. 5.7. It may be noted that the samples
were homogenized for 30 min. at the testing temperature before the commencement of the
tensile test. Figure 5.12a represents the microstructure at the beginning of superplastic
deformation at 550°C. Analysis of the microstructures following annealing and following
superplastic deformation at 550°C indicates that a significant grain growth and grain
elongation along the tensile loading direction was observed following superplastic
deformation (Figs. 5.12a-d). The grain size was evaluated following static annealing for 5 h
as this corresponds to the time taken to achieve the maximum elongation of 500% during
superplastic testing (Fig. 5.13a).
31
Figure 5.12: EBSD IPF maps of AA 5456 (a) following annealing at 550°C for 30 min.; and
following superplastic deformation to different strains of (b)100% (0.6 true
strain), (c) 150% (0.9 true strain), (d) 250% (1.25 true strain), (e) 450% (1.7 true
strain), (f) 700% (2 true strain).
32
From IPF maps, it is inferred that not all grains were elongated and a few grains were
retained their original shape following superplastic deformation. Normally relatively smaller
grains were more resistant to elongation that could be due to the influence of crystallographic
slip. Grain growth is diffusion controlled and it will cause an increase in the grain size in both
RD and TD directions (Fig. 5.13a). (It should be noted that the ratio of grain boundary area
surrounding a grain to its volume increases with a decrease in grain size. This also could be a
reason why grain boundary migration is less and the grain appears to have undergone less
shape change. Bringing in slip during superplastic deformation will make the picture more
complicated).
Figure 5.13: (a) Grain growth in RD-TD and RD-ND sections following annealing at 550°C
for 30 min. and following superplastic deformation to different strains (b)
Variation of grain aspect ratio following superplastic deformation to different
strains; in AA 5456.
Figure 5.13b shows the variation in grain aspect ratio with superplastic strain. It was
calculated by considering the grain as an ellipse and finding the ratio of intercept length along
major axis to the minor axis evaluated by using TSL OIM® (version 6.1) software. The
average aspect ratio of grains was found to be constant up to a true strain of 0.6, after which
there was a gradual increase up to the true strain 1.2 and thereafter it remained constant.
Constancy in the aspect ratio following superplastic deformation is consistent with the grain
switching mechanism (due to grain rotation caused by the presence of unbalanced shear
stresses at the boundary arising from GBS) during deformation. The spread in aspect ratio
also found to have increased following superplastic deformation up to a true strain of 1.2 and
after that, it was remained unchanged. This is indicating the possibility rearrangement of
grains and grain neighbor switching during deformation. Grain growth during deformation is
33
compared in both RD-TD section and RD-ND section and found that the grain growth
characteristics are similar in both the sections. (Superplastic deformation is isotropic).The
grain growth following superplastic deformation was found to be higher compared with grain
growth following static annealing. The dynamic grain growth Δd following superplastic
deformation under isothermal condition is a function of strain (ε) and strain rate (ε̇ ) (Caceres
and Wilkinson, 1984) as,
∆d = k t ε ε̇ −p
(5.1)
The Eqn. 5.1 is solved using measured grain growth (data shown in Table 5.4). The
solution gives values of p = 0.073, kt = 4 and the expression for estimating grain growth in
AA 5456 at 550°C is given as,
∆d = 4ε ε̇ −0.073
(5.2)
The grain growth following superplastic deformation is validated for different strain rates
at the same temperature (shown in Table 5.5). The calculated grain growth is consistent with
experimental result (using two fitting constants; but it can be said that the shape of the grain
growth curve, as predicted by Eqn. (5.1) is verified).
Table 5.4: Estimation of constants.
Condition,
550°C
True
strain 1
5×10-4 s-1
True
strain 1
5×10-2 s-1
Avg.
initial
grain
Size
(µm)
Avg. grain
size after
static
annealing
(µm)
9
9
12
9
Δda
(Static
grain
growth)
(µm)
3
0
Grain
Size
after
super
plastic
deform
ation
(µm)
Δdtotal
Overall
grain
growth
(µm)
22
13
14
5
Δd
Δdtotal -Δda
Dynamic
grain
growth
(µm)
kt
p
4
0.073
10
5
Table 5.5: Experimental and predicted dynamic growth model at 550°C at different strain
rates and strains using Eqn. 5.2.
Condition
Δd from experiment (µm)
True strain 2, 5×10-4 s-1, 550°C
True strain 1.5, 5×10-4 s-1, 550°C
True strain 1, 10-2 s-1, 550°C
14
10
5
34
Δd from Eqn. 5.2 (µm)
14
10.5
5.59
Table 5.6: Values of kt and p at different temperatures and strain rates.
Condition
True strain
1.7×10-4 s-1,
525°C
True strain
1.7×10-4 s-1,
500°C
True strain
1 5×10-4 s-1,
450°C
Avg.
initial
Grain
Size
(µm)
Avg.
grain
size
after
static
anneali
ng
(µm)
8
11
3
20
12
9
3
0.073
7
10
3
16
9
6
2.3
0.073
6
7
1
11
5
4
2.2
0.073
Δda
(Static
grain
growth)
(µm)
Grain
Size
after
super
plastic
deform
ation
(µm)
Δdtotal
(Overall
grain
growth)
(µm)
Δd
Δdtotal -Δda
(Dynamic grain
growth)
(µm)
kt
p
kt is dependent on temperature; the variation of kt with temperature is shown in Table 5.6.
Figure 5.14a denotes the variation of grain size with strain rate at a true strain of 1. It is noted
that that the dynamic grain growth is inversely proportional to the strain rate. The grain
growth rate (grain growth per unit time, which is normalized to initial grain size) is found to
be proportional to the strain rate (Fig. 5.14b). Caceres and Wilkinson (1984) proposed a rate
constant ~1. From Fig. 5.14b, the slope is calculated and it is nearly equal to 1 consistent with
the analysis of Caceres and Wilkinson (1984).
Figure 5.14: Variation of (a) grain size and (b) grain growth rate with strain rate following
superplastic deformation to a true strain of 1 at 550°C in AA 5456.
35

Texture evolution following superplastic deformation at optimal strain rate and
temperature
In the previous section, it was noted that the texture present in the as-received material was
weakened considerably following annealing at 550°C for 30 min. and a weak rotated cube
component {100}<011> was appeared which can be identified by {100} pole figure (Fig.
5.15a). In the present section, texture changes following superplastic deformation to different
strains are analyzed. The changes in texture with increasing superplastic strain are
represented in the form of {100} pole figures (Fig. 5.15a-e) and the φ2 = 0 ODF sections
(Figs. 5.16a-e). It is found that the rotated cube component present in the annealed sample
(Fig. 5.15a) was modified during superplastic deformation and after a strain of 250% (true
strain ~ 1.3); a weak <100> fiber was observed (Fig. 5.15d). With continued deformation, the
fiber component was changed to a weak cube component at the end of deformation (Fig.
5.15f).
(a)
(b)
(d)
(e)
(c)
(f)
Figure 5.15: {100} pole figure of AA 5456 (a) following annealing at 550°C for 30 min.; and
following superplastic deformation to different strains of (b) 100% (0.6 true
strain), (c) 150% (0.9 true strain), (d) 250% (1.25 true strain), (e) 450% (1.7 true
strain), (f) 700% (2 true strain).
36
(a)
(b)
(c)
(d)
(e)
(f)
Figure 5.16: ODF φ2 = 0° section of AA 5456 (a) following annealing at 550°C for 30
min.; and superplastic deformation to different strains of (b) 100% (0.6 true
strain), (c) 150% (0.9 true strain), (d) 250% (1.25 true strain), (e) 450% (1.7
true strain), (f) 700% (2 true strain).
The variation in the ODF values, f (g) for cube and rotated cube components during
superplastic deformation is shown in Fig. 5.17. The ODF values of rotated cube component
were decreased and the ODF values of cube component were increased after 250%
elongation (true strain of 1.25). The cube orientation and <111> fiber orientation are the most
stable orientations whereas the <110> orientation is the least stable orientation under
crystallographic slip for uniaxial loading condition (Pérez-Prado et al., 2001). Therefore, the
appearance of cube component and disappearance of rotated cube component could be due to
the effect of crystallographic slip during superplastic deformation and this was found to occur
at the final stages of the deformation after true strain 1.7. It is also observed that the
variations of other components were not consistent during deformation and this can be due to
the effect of random grain rotation. Pérez-Prado et al. (2001) postulated that the grains which
are {100}<0uv> oriented can be rotated to {100}<010> orientation during deformation under
dislocation creep condition. The evolution of weak cube component is not inconsistent with
the ‘random grain rotation’ theory, (which should only be regarded as an approximation
convenient for mathematical analysis, is a ‘random walk’ analogy) proposed by Engler et al.
(2000). Grains can rotate randomly during deformation except for a few grains that are
37
aligned in {100}<0uv> direction and these grains will rotate in order to achieve the most
stable orientation only under dislocation creep condition.
Figure 5.17:

The variation of ODF value for cube, rotated cube and {110} < 11̅0 >
orientations with superplastic strain in AA 5456.
Grain boundary mis-orientation distribution following superplastic deformation at
optimal strain rate and temperature
The evolution of grain boundary mis-orientation distribution (number fraction of boundaries)
during different stages of (transient and steady state) superplastic deformation is shown in
Figs. 5.18a-f. Following the recrystallization and annealing, the mis-orientation distribution
was very close to Mackenzie distribution that is the expected distribution in a randomly
oriented polycrystalline material (Fig. 5.18a). With increasing superplastic strain, LAB
number fraction was increased slightly (Fig. 5.18b) and it remained constant up to 250%
elongation (true strain ~ 1.3). At larger strains, a large increase in the LAB number fraction
was observed (Fig. 5.18e and f). This could be due to the fact that the LABs are more
resistant for GBS compared with HAB. With increasing superplastic strain, strain induced
grain growth occurs due to the migration of HABs and there is a reduction in the number
fraction of HABs. Since the LABs are unaffected, this leads to an increase in the LAB
fraction with increasing superplastic strain. In addition, the dislocations could be generated
during GBS as a non-rate controlling accommodation mechanism and the emitted
dislocations are accumulated in grains leading to an increase in the LAB fraction.
38
(a)
(b)
(d)
(c)
(f)
(e)
Figure 5.18: The mis-orientation distributions of AA 5456 (a) following annealing at 550°C
for 30 min.; and following superplastic deformation to different strains of (b)
100% (0.6 true strain), (c) 150% (0.9 true strain), (d) 250% (1.25 true strain), (e)
450% (1.7 true strain), (f) 700% (2 true strain).
39
5.1.6
Microstructure and texture evolution following deformation at high strain rate
Figure 5.19 shows the EBSD IPF map for the specimen deformed at a high strain rate (5×10-2
s-1), and at a temperature of 550°C, and maximum 200% (true strain 1) is obtained. Some
grains were observed to be elongated following deformation, and the grain elongation could
be stress-directed grain growth. (The effect diffusion / dislocation motion etc. could not be
deduced from a macro-measurement). The ‘strain rate sensitivity index’ was calculated using
strain rate jump test and it is nearly equal to 0.2, which is consistent with the presence of
dislocation creep during deformation.
Figure 5.19: EBSD IPF map of AA 5456 following superplastic deformation to a strain of
200% (1 true strain), at 5×10-2 s-1 and at 550°C.
(a)
(b)
(c)
Figure 5.20: {100} pole figure of AA 5456 following superplastic deformation to a strain of
200% (1 true strain), to different strain rates as (a) 5×10-4 s-1 (b) 10-2 s-1 (c) 5×
10-2 s-1; at 550°C.
40
(a)
(b)
(c)
Figure 5.21: ODF φ2 = 0° section of AA 5456 following superplastic deformation to a strain
of 200% (1 true strain), at different strain rates of (a) 5×10-4 s-1 (b) 10-2 s-1 (c)
5×10-2 s-1, at 550°C.
Figure 5.20, shows the {100} pole figure of the specimens deformed at different strain
rates of (a) 5×10-4 s-1, (b) 10-2 s-1, (c) 5×10-2 s-1 to a true strain of 1. Figure 5.21 shows the φ2
= 0° section of the ODF of the specimens deformed at different strain rates of (a) 5×10-4 s-1,
(b) 10-2 s-1, (c) 5×10-2 s-1 to a true strain of 1. From the pole figures and ODF, it is clear that
the intensity or ODF value of the cube component increased with the strain rate and it could
be due to the effect of Taylor slip (Taylor, 1938) during deformation. The stress exponent
following deformation at high stain rates is found to be close to 5 and this value is consistent
with dislocation creep. It was reported that certain texture components get stabilized by the
effect of dislocation creep (Engler et al., 2000; Pérez-Prado et al., 2001). For example, <111>
fiber orientation and ‘{100} <010> Cube’ orientations are the stable orientations under
dislocation creep (Pérez-Prado et al., 2001). The <111> fiber is weak but ‘{100}<010> cube’
orientation is stronger following deformation at high strain rate (5×10-2 s-1). It is probably
because {100}<010> slip system that leads to cube texture component and has equally
stressed eight slip systems has lower resistance to slip than the system that leads to the <111>
texture component caused by six equally stressed slip systems (both satisfy the Von Mises
criterion of a minimum of 5 independent slip systems). On the other hand at lower strain rates
(Figs. 5.20a and b), the stress exponent is found to be in the range of 2-3. The ODF value
corresponding to the cube orientation was not found to have increased. This is due to the fact
that the contribution of dislocation creep is believed to be to be negligible and GBS with
near-random grain rotation is expected to be present in that regime. Analysis of the misorientation distribution of samples deformed at high strain rates (Fig. 5.22) indicates that the
LAB fraction has increased during the deformation and this could be due to the ease of
41
accommodation of dislocation motion inside the grains by the motion of dislocations along
the LABs.
Figure 5.22: Mis-orientation distribution of AA 5456 following superplastic deformation to a
strain of 200% (1 true strain), at 5×10-2 s-1 and 550°C.
5.2 Dynamically recrystallizing alloy, AA 2004
5.2.1

Base metal characterization
Composition
The composition of AA 2004 was determined using optical emission spectroscopy and is
given in Table 5.7.
Table 5.7: Composition in wt. % of AA 2004.
Si
0.1
5.2.2
Fe
0.58
Cu
5.7
Mn
0.04
Zr
0.29
Zn
0.1
Al
Balance
Room temperature mechanical properties of AA 2004
The room temperature tensile properties are estimated using room tensile testing procedures
are described in Chapter 4. The yield strength, ultimate tensile strength and the elongation to
fracture are listed in Table 5.8.
The room temperature stress strain curve of AA 2004 in as-received condition is shown
Fig. 5.23.
42
Figure 5.23: Engineering stress vs. engineering strain curve of AA 2004 at room temperature.
Table 5.8: Room temperature properties of AA 2004.
Yield stress (0.2% offset)
Ultimate stress (UTS)
Maximum strain
5.2.3
225 MPa
277 MPa
8.07%
High temperature tensile behaviour of AA 2004
To idetify the parameters for achieving optimal superpalstic flow, the alloy AA 2004 was
tested at different temperatures (350°C, 400° C, 437°C, 457°C, 477°C) and different strain
rates (3×100-4 s-1, 3×10-3 s-1, 3×10-2 s-1). The stress-strain curves obtained at different
temperatures and strain rates are plotted in Fig. 5.24.The stress-strain curves indicate that
there is no steady state deformation regime in this material at all temperatures tested. The
variation of maximum stress with strain rate at different temperatures is shown in Fig. 5.25a.
The variation of the percentage of elongation with strain rate at different temperatures is
shown in Fig. 5.25b. The maximum stress decreases with increasing temperature (Fig. 5.23a)
and the maximum elongation of 500±50% in AA 2004 was obtained at a strain rate of 3×10-3
s-1 and at a temperature of 457°C, and this was taken as the optimum strain rate for further
studies. The microstructure and texture characterization was carried out in the transient
deformation regime by carrying out interrupted tests to different strains, which are indicated
by arrows in Fig. 5.26. It may be noted that the AA 2004 used in the present study is clad
with pure Al. From the comparison of elongations reported in the literature and the present
study it appears that the Al cladding has not affected the superplastic behavior significantly.
However, It was reported that the cladding of superplastic Al–Zn–Mg–Cu-Ni alloy with nonsuperplastic Al–1% Zn reduced the relative elongation by about a factor 1.5 (Portnoi et al.,
2010).
43
(b) 400°C
(a) 350°C
(c) 457°C
Figure 5.24: Tensile stress-strain curves of AA 2004 at different strain rates at temperatures
of (a) 350°C, (b) 400°C, (c) 457°C. (The stress jumps in the above figure
correspond to the strain rate jump tests that were carried out to evaluate the
strain rate sensitivity index ‘m’).
(b)
(a)
Figure 5.25: Plots of variation of (a) flow stress with temperature and strain rate, and (b)
percentage of elongation with temperature and strain rate in AA 2004.
44
Figure 5.26: Stress vs. strain curve at 457°C for 3×10-3 s-1 in AA 2004, the arrows indicating
the strains to which the intermediate superplastic tensile tests were carried out to
study microstructure development with strain.
5.2.4
Microstructure and texture evolution following annealing at 457°C
The optimum superplastic flow was observed at 457°C and this temperature was taken for the
annealing studies with a view to comparing microstructure evolution with and without strain.
The inverse pole figure (IPF) maps were obtained from the EBSD data analysis of AA 2004
samples in ‘as-received’ and following annealing at 457°C for 30 min. and 8 h are shown in
Figs. 5.27a-c. The annealing time of 30 min. is corresponding to the homogenizing time
before the start of superplastic deformation. The grain boundary maps are also superimposed
on the IPF maps. The 2°-5° misoriented boundaries are highlighted in yellow, 5°-15°
misoriented boundaries are highlighted in red and 15°-60° misoriented boundaries are
highlighted in black. The EBSD scans were performed on the RD-ND plane and for
analyzing texture in the RD-TD plane; the orientations were rotated by 90° about RD. The
data points with confidence index (CI) values < 0.1 were not considered for the analysis and
these are represented as black pixels in the IPF maps. The IPF maps in the as-received
material reveal a banded microstructure with the presence of a strong texture. Following
annealing at 457°C for 30 min., considerable fractions of grains were recrystallized and some
grains were exhibited abnormal growth (Fig. 5.27b). However, the banded microstructure
was still found to have persisted even after annealing. Following annealing for 8 h complete
recrystallization and grain growth was occurred and new orientations were developed (Fig.
5.27c).These recrystallized grains can be identified by analyzing the grain orientation spread
(GOS)* and if (GOS) < 1° are considered as recrystallized. Figures 5.28a-c shows the GOS
maps for the as-received, following annealing for 30 min. and 8 h respectively. Following
annealing for 30 min. ~ 40% grains have recrystallized based on the GOS < 1° criterion.
45
*GOS: “It is the average mis-orientation between all points inside the grain. I.e. it is the average deviation
between the orientation of each point in the grain and the average orientation for a grain” (Stuart et al., 2011).
Figure 5.27: EBSD IPF maps of AA 2004 in (a) as-received, (b) following annealing at
457°C for 30 min. and (c) following annealing at 457°C for 8 hours. The black
regions are corresponding to the regions with CI<0.1.
Figure 5.28: Grain orientation spread maps of AA 2004 in (a) as-received, (b) following
annealing at 457°C for 30 min. and (c) following annealing at 457°C for 8 h.
46
(Please note larger regions were investigated to study the recrystallization (Fig. 5.28) and
these regions are different from the maps shown in Fig. 5.27). The recrystallization is found
to be dependent on the orientation of grains. It is observed that the {110}<uv0> orientations
were more resistant to static recrystallization compared with other orientations, and only 25%
of grains with {110}<uv0> orientations were found to have recrystallized whereas almost 60
% grains away from {110}<uv0> orientations were found to have recrystallized following
annealing at 457°C for 30 min. The orientation development during annealing is discussed
below. Recrystallization is dependent on the stored energy, which is orientation dependent,
and further studies are required to establish a relation between dynamic recrystallization and
orientation of crystal.

Texture evolution following annealing at 457°C
The textures are represented in the form of {100} pole figures (Figs. 5.29a-c) and ODF (Figs.
5.30a-c). The alloy in the as-received state shows strong texture with dominant components
of
{110} < 11̅√2 >, {110} < 11̅2 > Brass’,
{110} < 11̅1 > orientations.
Following
annealing at 457°C for 30 min., the initial texture was retained indicating the occurrence of
recovery dominated continuous recrystallization. The material was found to be completely
recrystallized and the initial banded microstructure disappeared following annealing at 457°C
for 8 h. In addition, there was a significant change in the crystallographic texture following
annealing for 8h compared with the as-received condition is also observed.
(a)
(c)
(b)
Figure 5.29: {100} pole figure of AA 2004 in (a) as-received, (b) following annealing at
457°C for 30 min. (c) following annealing at 457°C for 8 h.
47
(a)
(b)
(c)
Figure 5.30: ODF at φ2 = 0° section of AA 2004 in (a) as-received, (b) following annealing
at 457°C for 30 min. (c) following annealing at 457°C for 8 h.
Figure 5.31: Variation of the ODF values of the texture components {110} < 11̅√2 > and
{100}<011> following annealing at 457°C in AA 2004.
The texture changes following annealing at 457°C are shown in Fig. 5.31 and it is
inferred that the microstructure was found to be textured even after prolonged annealing i.e.,
the texture randomization was not found to have occurred following annealing. I.e. the asreceived material was dominant with {110} < 11̅√2 > orientation and was modified to
{100}<011> orientation following annealing.
48

Grain boundary mis-orientation distribution following annealing at 457°C
It is possible to understand the mechanism of recrystallization from the analysis of misorientation distribution data. In discontinuous recrystallization, the nucleation and growth are
dominant and the mis-orientation distribution in recrystallized condition will be completely
different compared with the deformed state and a large fraction of HABs will be developed.
On the other hand, in continuous recrystallization, the fraction of HABs almost remains the
same, as it is a recovery-dominated process without significant nucleation and growth.
Consequently, the mis-orientation distribution after continuous recrystallization will be
almost similar to that of the deformed state (Humphreys and Hatherly, 1995). The misorientation distribution of as-received and annealed microstructure is shown in Figs. 5.32a-c.
The as-received material has a large fraction of sub-grain boundaries (SGBs) (mis-orientation
angles < 5°) and LABs (mis-orientation angles in the range of 5-15°). Following annealing at
457°C for 30 min. was resulted in a reduction in the SGB fraction and a slight increase in the
HAB fraction (Fig. 5.32b). Those HABs, mis-orientated by 45°-60°, are likely to be
symmetric variants of major texture components. These symmetric variants are the
orientations {110} < 11̅√2 > ; {110} < 11̅2 >, {110} < 11̅1 > etc. The decrease in the
SGB fraction following annealing could be due to the recovery process associated with the
annealing. The increase in HAB fraction observed during annealing could be due to the
overall reduction in the SGB fraction. It is observed that the material was completely
recrystallized following annealing at 457°C for 8 h and the mis-orientation distribution is
remained unchanged following complete recrystallization in comparison with the sample
annealed for 30 min. The random Mackenzie type distribution is also shown by red lines and
it is observed that the mis-orientation distribution is not close to random distribution even
after 8 h of annealing. It can therefore be inferred that the observed recrystallization during
static annealing is of continuous type.
49
(a)
(b)
(c)
Figure 5.32: Mis-orientation distribution of AA 2004 in (a) as-received, (b) following
annealing at 457°C for 30 min., (c) following annealing at 457°C for 8 h and the
red lines correspond to the random Mackenzie distribution.
5.2.5
Microstructure and texture evolution following superplastic deformation at
optimal strain rate and temperature
The EBSD IPF maps were taken from RD-ND sections of the samples subjected to different
superplastic strains at the optimal superplastic conditions are shown in Figs. 5.33a-f. For the
texture analysis in the RD-TD plane, the data were rotated about RD by 90°. In the IPF maps,
the 2° to 5° misoriented boundaries are highlighted by blue lines, 5° to 15° misoriented
boundaries are highlighted by red lines and 15°-180° misoriented boundaries are highlighted
by black lines. Figure 5.33a shows the IPF map of AA 2004, following annealing at 457°C
for 30 min., the material was recovered during annealing, and the SGBs present in the asreceived condition were reduced by this recovery.
50
Figure 5.33: EBSD IPF maps of AA 2004 (a) following annealing at 457°C for 30 min.; and
following superplastic deformation to different strains of (b) 25% (0.22 true
strain), (c) 50% (0.4 true strain), (d) 100% (0.6 true strain, (e) 200% (1 true
strain) and (f) 500% (1.7 true strain).
51
The initial microstructure following annealing was not suitable for superplasticity as there
was an elongated grains structure and low fraction of high angle boundaries was observed.
Up to a strain of 100% (true strain ~ 0.6), the deformation bands were found to be persistent
(Figs. 5.33b-d), However, with increasing strain the complete elimination of the banded
microstructure was resulted (Figs. 5.33e-f). This is consistent with the results reported by
Bate et al., 2007 and it was inferred that crystallographic slip act as a dominant mechanism
for deformation in the low strain regime (Bate et al., 2007).
Figures 5.34a and b show the variation of micro-hardness and grain sizes with
superplastic strain (measured on the RD-ND section of the sample) respectively. The
hardness of the as-rolled alloy was reduced from 80 Hv to 60 Hv following annealing at
457°C and for 30 min. and this could be attributed to the effect of recovery, partial
recrystallization and grain growth. Following superplastic deformation the hardness was
decreased linearly with strain to 40 Hv. The decrease in the hardness is attributed recovery,
recrystallization and grain growth following superplastic deformation. The average grain size,
on the other hand was decreased steadily up to a true strain of 0.5 and thereafter remained
constant (~5 µm) following superplastic deformation. The decrease in average grain size
during
the
initial
stages
of
the
deformation
is
attributed
to
the
dynamic
recovery/recrystallization during the stage and the microstructure is suitable for achieving
superplastic flow.
(b)
(a)
Figure 5.34: Plots of variation of (a) micro-hardness (Hv) at room temperature and (b) grain
size in the as-received, following annealing at 457°C for 30 min. and following
superplastic deformation to different strains in AA 2004.
52
 Texture evolution following superplastic deformation at optimal strain rate and
temperature
Figures 5.35a-g shows the {100} pole figures for the annealed and following superplastic
deformation to different strains. Pole figures clearly indicating the gradual weakening of
texture with increasing the superplastic elongation. To make quantitative analysis, the
orientation distribution function (ODF) was calculated from the EBSD data by using
‘Harmonic series expansion’ method up to L = 16 and HW = 5 imposing ‘cubic’ crystal
symmetry and ‘orthotropic’ sample symmetry. Since the intention was to study the texture
changes during superplastic deformation, the orthotropic sample symmetry which is
applicable for rolled sample was applied. The φ2 = 0˚ section of the ODF is plotted in Figs.
5.36a-f for the annealed and superplastically deformed samples. The ODF value decreased
after a particular strain and this phenomenon could be explained by ‘random grain rotation’
model proposed by Engler et al. (2000).
(a)
(b)
(d)
(f)
(c)
(g)
Figure 5.35: {100} pole figure of AA 2004 (a) following annealing at 457°C for 30 min., and
following superplastic deformation to different strains of (b) 25% (0.22 true
strain), (c) 50% (0.4 true strain), (d) 100% (0.69 true strain), (e) 300% (1.3 true
strain), (f) 500% (1.7 true strain).
53
(a)
(b)
(c)
(d)
(e)
(f)
(g)
Figure 5.36: ODF φ2 = 0° section of AA 2004 (a) following annealing at 457°C for 30 min.,
and following superplastic deformation to different strains of (b) 25% (0.22
true strain), (c) 50% (0.4 true strain), (d) 100% (0.69 true strain), (e) 200% (1
true strain), (f) 300% (1.3 true strain), (g) 500% (1.7 true strain).
Figure 5.37: Variation of ODF corresponding to {110} < 11̅√2 >, {{110} < 11̅2 >and
{110} < 11̅1 > orientations following superplastic deformation to different
strains in AA 2004.
54
Figure 5.37 shows the variation of the ODF value of different major texture components
corresponding to {110} < 11̅√2 >, {110} < 11̅2 > Brass’ and {110} < 11̅1 > orientations
following static annealing and following superplastic deformation. It is observed that the
ODF value of {110} < 11̅√2 > orientation was slightly decreased during annealing whereas
the ODF value of {110} < 11̅2 > Brass orientation was almost constant up to a true strain of
0.6 following superplastic deformation. Brass component is the most stable orientation under
plane strain condition and the deformation of banded microstructure is analogous to plane
strain compression (Pérez-Prado et al., 2003). Therefore, during the initial stages of
deformation, in the absence of GBS, crystallographic slip can acts as a dominant deformation
mechanism and this favour the development of Brass component in plane strain condition
(Bate et al., 2005). The ODF value of major texture components was found to have decreased
drastically after a true strain of 0.6 and this could be due to the effect of random grain
rotation during GBS (Engler et al., 2000).
Figure 5.38: Variation of orientation spread (Gaussian half width) during deformation
normalized with initial Gaussian half width for both Φ and φ1 directions
following superplastic deformation in AA 2004.
The normalized texture spread following superplastic deformation was calculated from
the Gaussian half widths of the ODF peaks for Φ when φ2 = 0°, φ1 = 45° and for φ1 when φ2
= 0°, Φ = 45° are plotted in Fig. 5.38. The spread increases with increase in strain and this is
an indicator of texture weakening/annihilation during superplastic deformation.
55

Grain boundary mis-orientation distribution following superplastic deformation at
optimal strain rate and temperature
Mis-orientation distributions following annealing and following superplastic deformation to
different strains are shown in Figs. 5.39a-f. The as-received material was having a large
fraction of SGBs and LABs (Fig. 5.32a). Following annealing at 457°C for 30 min. there was
a reduction in the SGB fraction and a slight increase in the HAB fraction is observed. (Fig.
5.39a). The decrease in the SGB fraction during annealing is most likely due to the recovery
process associated with annealing. The increase in the HAB fraction observed during
annealing could also be due to the overall reduction in the SGB fraction. The fraction of the
LABs and the fraction of the 45° to 60° misoriented boundaries were observed to be constant
up to a true strain of 0.69 (i.e., 50%) (Figs. 5.39a-d). It is expected that GBS and grain
rotation were not present during the initial stages of the deformation up to a true strain of
0.69. Further straining was resulted in the reduction in the fraction of LABs, because these
could be possibly converted to HABs during superplastic deformation. The SGB fractions
have increased after a true strain of ~0.4 (Figs. 5.39d-f) during GBS regime. This could be
due to fact that the SGBs are more resistant to GBS, only HABs undergo sliding during
deformation and eventually this leads to the accumulation of SGBs following superplastic
deformation (Figs. 5.39d-f). The fractions of LABs were reduced with increasing deformation
(due the slow conversion of LABs to HABs by strain at the very slow strain rate of GBS
possible at the SGBs) but the HAB fraction remained constant (most of the deformation is
concentrated in the HABs, which is also assisted by grain rotation) and this leads to an
overall increase in SGB fraction. Another possible reason for the increase in SGBs is the
accumulation of dislocations inside the grains during superplastic deformation (Pérez-Prado
et al., 2003). (It should be noted that experimental evidences are limited for the formation of
sub-grains and dislocation sub-structures following superplastic deformation. The popular
view is that dislocation motion, particularly of those emitted from the grain boundaries, can
act as an accommodation mechanism for GBS, and the emitted dislocations traverse the
grains and get absorbed at the opposite boundaries (Langdon, 2009)). Figure 5.40 represents
the variation of SGB, LAB and HAB fractions (45°-65° misoriented boundaries indicated
separately, most of the 45°-65° misoriented boundaries could be the symmetric variants of
majour texture components) following superplastic deformation with including SGBs (Fig.
5.40a) and without including SGBs (Fig. 5.40b). Due to the accumulation of SGB fraction in
the GBS regime, considering SGB fraction, it is not possible to understand the variation of
LABs and HABs fraction following superplastic deformation. By excluding the SGB
56
fraction, it is clear that LABs were converted to HABs following superplastic deformation.
The fraction of the 45°-65° misoriented boundaries were almost remained unchanged
following superplastic deformation (shown in Fig. 5.40 b).
(a)
(b)
(c)
(d)
(e)
(f)
Figure 5.39: Mis-orientation distribution of AA 2004 (a) following annealing at 457°C for 30
min.; and following superplastic deformation to different strains of (b) 25%
(0.22 true strain), (c) 50% (0.4 true strain), (d) 100% (0.69 true strain), (e) 200%
(1 true strain), (f) 500% (1.7 true strain).
57
(a)
(b)
Figure 5.40: Variation of SGBs, LABs and HABs fraction (including the symmetric variant
of majour texture components (45°-65° misoriented boundaries) (a) including
SGB fraction (b) excluding SGB fraction following superplastic deformation in
AA 2004.
5.3 Orientation diffusion model
Engler et al. (2000) proposed a model to predict the texture changes during superplastic
deformation termed as ‘orientation diffusion model’. According to this model, the grains are
considered to be rigid, the orientation spread and texture reduction is brought about by the
(near-) random rotation of grains during superplastic deformation. Since the GBS acts as the
fundamental atomistic mechanism for the superplastic deformation, the unbalanced shear
stresses caused by GBS along grain boundaries that slide at different rates depending on the
orientation of the boundary to the stress axis will give rise to grain rotation during GBS. The
grain rotation should be sufficiently fast to accommodate the degree of rotation needed by
GBS. In other words, consistent with experimental observations, superplastic deformation is
regarded as isotropic. The change in orientation is analogous to the concentration change
during diffusion and “time” in the diffusion equation gets replaced here by “strain”.
Therefore, Fick’s law of diffusion can be used for the present problem as follows.
𝑗=
𝜕𝑓(𝑔) 𝜕𝑓(𝑔)
𝜕∅
;
𝜕(𝜀)
=𝑆
𝜕2 𝑓(𝑔)
𝜕∅2
(5.3)
Ø = change in orientation ε = true strain, S is a constant
The solution for the texture peak broadening is obtained by considering a point source in
an infinite sphere, the infinite sphere corresponding to the orientation space for orientation
58
diffusion. If the orientation diffusion coefficient is isotropic and uniaxial state of stress
provided, the solution is obtained as,
𝑓(g, ε, Ø) =
5.3.1
F
(5.4)
2
(4πS√ε)exp(−Ø ⁄Sε)
Validation in AA 2004
Unlike in AA 5456, in AA 2004, a strong texture is present at the beginning of superplastic
deformation and orientation diffusion model can be validated using the experimental data.
Therefore experimental ODF values from φ2 section for φ2 = 0° and Φ = 45° are taken for the
verification of this model. (The ODF values from φ1 section when φ2 = 0° and Φ = 45° are not
considered, because the orientation change of φ1 when φ2 = 0° and Φ = 45° corresponds to
alpha (FCC rolling) fiber). Annealed condition cannot be described by ε = 0 because the ODF
will reach to infinity when ε = 0. In order to avoid that, the starting worked material is
assumed still to be somewhat strained following annealing (which actually is the case
because this material does not statically recrystallize) and the Eqn.5.4 can be modified as
𝑓(𝑔, 𝜀, Ø) =
𝐹
2
(4𝜋𝑆√(𝜀+𝜀0 ))exp(−Ø ⁄𝑆(𝜀+𝜀 ))
0
(5.5)
Ø = changes in orientation, ε = true strain, εo = strain when ε = 0 (insertion of εo in the
above equation would be an approximation based on mathematical convenience).
The constants F, S and εo are calculated using ‘Levenberq-Marquardt’ algorithm using
Matlab R2012a®, and the obtained values are F = 2.4, S = 0.04, εo = 0.1. (At this stage there
is no physical significance for the values of F, S and ε; curve fitting shows the texture
variation is according to the function that represents random grain rotation. Physical
significance of each of these constants is given in the paper of Engler et al. (2000)) The Eqn.
5.6 obtained by substituting F = 2.4, S = 0.04, εo = 0.1 to the Eqn. 5.5 becomes
𝑓(𝑔, 𝜀, Ø) =
2.4
2
(4𝜋0.04√(𝜀+0.1))𝑒𝑥𝑝(−Ø ⁄𝑆(𝜀+0.1))
(5.6)
The function f (g, ε, Ø) is plotted and compared with actual ODF values and this is shown
in Fig. 5.41. The thick solid lines correspond to the f (g, ε, Ø) and dotted lines are
corresponding to the ODF calculated from EBSD data. The texture changes during
deformation are in good agreement with the model and hence it can (near-) random grain
rotation does occur during superplastic deformation in the steady state.
59
Figure 5.41: Comparison of experimental ODF of Φ section corresponding to φ2 = 0° and φ1
= 45° of AA 2004 following superplastic deformation to different strains and the
calculated ODF from orientation diffusion model.
60
CHAPTER 6
CONCLUSIONS
From the investigation on the texture and microstructure development during superplastic
deformation of Al alloys AA 5456 and AA 2004; the following conclusions were arrived at:
1.
Al alloy AA 5456 exhibits discontinuous static recrystallization following annealing at
550°C by particle stimulated nucleation. The recrystallized microstructure is finegrained (~ 8 µm) and has a large fraction of high angle boundaries. The maximum
superplastic % elongation for AA 5456 is achieved at 550°C, at a strain rate of 5×10-4
s-1. The alloy AA 5456 exhibited better superplastic properties (higher elongation and
lower stresses) in comparison with alloy AA 5083.
2.
The initial rolling texture in AA 5456 is completely destroyed and a weak rotated cube
component is present following annealing at 550°C for 30 min. The weak rotated cube
component present in the annealed sample is modified during superplastic deformation
and after a strain of 250% (true strain ~ 1.25) a weak <100> fiber is observed. With
continued deformation, the fiber component changes to a weak cube component.
3.
Alloy AA 5456 exhibits dynamic grain growth during superplastic deformation. The
shapes of grain growth and growth rate experimental curves are consistent with those
predicted by the model of Caceres and Wilkinson (1984).
4.
The highest percentage superplastic elongation in AA 2004 is obtained at 457°C, at a
strain rate of 3×10-3 s-1. The alloy exhibits continuous dynamic recrystallization during
deformation and initial banded microstructure is changed into a fine grained (~ 3 µm)
microstructure, which is suitable for superplastic deformation after a true strain of 0.7.
5.
In AA 2004, during initial stages of superplastic deformation retention of ‘{110} <
11̅2 >’ texture component is observed and this is a stable orientation during
crystallographic slip under plane strain compression. However, the ODF values of
major texture components were found to be decreased with increasing superplastic
deformation. This is believed to be due to the (near-) random rotation of grains.
6.
The texture changes following superplastic deformation are in good agreement with the
‘Orientation diffusion model’ proposed by Engler et al. (2000).
61
CHAPTER 7
SCOPE FOR FURTHER RESEARCH
1. The LAB fraction found to be increased during straining at optimum strain rate
observed for both statically and dynamically recrystallizing aluminium alloys. It
could be due the accumulation of dislocation inside the grains which can be
identified using TEM.
2. Flow stresses, percentage of elongation etc. are connected with temperature, grain
size, dynamic grain growth etc. and it can be optimized using ANN, regression etc.
for achieving better superplastic properties.
3. The dynamic recrystallization during superplastic deformation is found to be
orientation dependent. Further studies are required in this area.
62
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68
LIST OF PAPERS BASED ON THESIS
1. Arunbabu, K., S. Sankaran, S. Jupp, V. Subramanya Sarma, K. A. Padmanabhan
(2014). On the implications of texture and microstructure changes during isotropic
and anisotropic superplastic deformation. International Conference on Textures in
Materials ICOTOM-17, Dresden, Germany, August 24-29, 2014 (Accepted for oral
presentation).
69
CURRICULUM VITAE
1. NAME
:
ARUN BABU K
2. DATE OF BIRTH
:
13/09/1987
3. EDUCATION QUALIFICATIONS
2010 B Tech
Institution
:
Specialization :
TKMCE Kollam
Mechanical Engineering
2014 Master of Science (By Research)
Institution
:
Indian Institute of Technology Madras, Chennai
Registration date
:
18-07-2011
Address
:
Arun Babu K
S/o Parameswaran K
Kizhakkiniyakathu Mana,
Klari (PO), Malappuram
Kerala-676501
70
GRADUATE TEST COMMITTEE
CHAIRPERSON:
Dr. M. Kamaraj
Professor and Head
Department of Metallurgical and Materials Engineering
GUIDE:
Dr. V. Subramanya Sarma
Associate Professor
Department of Metallurgical and Materials Engineering
MEMBERS:
Dr. Uday Chakkingal
Professor
Department of Metallurgical and Materials Engineering
Dr. Srikanth Vedantam
Associate Professor
Department of Engineering Design
71
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