High-Temperature Behaviour of Austenitic Alloys Mattias Calmunger

High-Temperature Behaviour of Austenitic Alloys Mattias Calmunger
Linköping Studies in Science and Technology, Thesis No. 1619
Licentiate Thesis
High-Temperature Behaviour of
Austenitic Alloys
-Influence of Temperature and Strain Rate on
Mechanical Properties and Microstructural Development
Mattias Calmunger
LIU-TEK-LIC-2013:53
Division of Engineering Materials
Department of Management and Engineering
Linköping University, SE-581 83, Linköping, Sweden
http://www.liu.se
Linköping, September 2013
Opponent: Professor Göran Engberg, Dalarna University, Falun, Sweden.
Date: 10:15, November 1, 2013
Room: ACAS, Linköping University
Cover:
Dynamic recrystallization in AISI 316L slow strain rate tensile tested at elevated
temperature.
Printed by:
LiU-Tryck, Linköping, Sweden, 2013
ISBN 978-91-7519-512-4
ISSN 0280-7971
Distributed by:
Linköping University
Department of Management and Engineering
SE-581 83, Linköping, Sweden
© 2013 Mattias Calmunger
This document was prepared with LATEX, October 2, 2013
Abstract
The global increase in energy consumption and the global warming from
greenhouse gas emission creates the need for more environmental friendly
energy production processes. Biomass power plants with higher efficiency
could generate more energy but also reduce the emission of greenhouse gases,
e.g. CO2 . Biomass is the largest global contributor to renewable energy and
offers no net contribution of CO2 to the atmosphere. One way to increase
the efficiency of the power plants is to increase temperature and pressure in
the boiler parts of the power plant.
The materials used for the future biomass power plants, with higher
temperature and pressure, require improved properties, such as higher yield
strength, creep strength and high-temperature corrosion resistance. Austenitic
stainless steels and nickel-base alloys have shown good mechanical and chemical properties at the operation temperatures of today’s biomass power plants.
However, the performance of austenitic stainless steels at the future elevated
temperatures is not fully understood.
The aim of this licentiate thesis is to increase our knowledge about the
mechanical performance of austenitic stainless steels at the demanding conditions of the new generation power plants. This is done by using slow
strain rate tensile deformation at elevated temperature and long term hightemperature ageing together with impact toughness testing. Microscopy is
used to investigate deformation, damage and fracture behaviours during slow
deformation and the long term influence of temperature on toughness in the
microstructure of these austenitic alloys. Results show that the main deformation mechanisms are planar dislocation deformations, such as planar
slip and slip bands. Intergranular fracture may occur due to precipitation
in grain boundaries both in tensile deformed and impact toughness tested
alloys. The shape and amount of σ-phase precipitates have been found to
strongly influence the fracture behaviour of some of the austenitic stainless
steels. In addition, ductility is affected differently by temperature depending
on alloy tested and dynamic strain ageing may not always lead to a lower
ductility.
iii
Acknowledgement
This research has been financially supported by AB Sandvik Materials Technology in Sandviken, Sweden and the Swedish Energy Agency through the
Research Consortium of Materials Technology for Thermal Energy Processes,
Grant No. KME-501, for which they are all greatly acknowledged.
Many thanks to my supervisors Sten Johansson, Guocai Chai and Johan
Moverare for their support, guidance and encouragement during this project.
I would also like to thank all my colleagues at the division of Engineering Materials for fruitful discussions and creating an enjoyable working environment.
The technical support from Annethe Billenius, Bo Skoog, Patrik Härnman,
Per Johansson and Peter Karlsson is greatly acknowledged. Agora Materiae and the Strategic Faculty Grant AFM (SFO-MAT-LiU#2009-00971) at
Linköping University are also acknowledged.
In addition, a collective acknowledgement goes out to my colleagues at
Sandvik Materials Technology, especially to Jan Högberg, Jerry Lindqvist
and Håkan Nylén for all their help and assistance.
Finally, I would like to thank my family and especially my dear wife and
son, for all their patient and support.
Mattias Calmunger
Linköping, September 2013
v
List of Papers
In this thesis, the following papers have been included:
I. M. Calmunger, G. Chai, S. Johansson and J. Moverare, Deformation
and damage behaviours of austenitic alloys in the dynamic strain ageing
regime, Submitted for publication.
II. M. Calmunger, G. Chai, S. Johansson and J. Moverare, Influence of
deformation rate on mechanical response of an AISI 316L austenitic
stainless steel, Accepted for presentation at THERMEC’2013, Las Vegas (USA), 2013, To appear in Advanced Materials Research or Materials Science Forum.
III. M. Calmunger, G. Chai, S. Johansson and J. Moverare, Mechanical
behaviours of Alloy 617 with varied strain rate at high temperatures,
Accepted for presentation at THERMEC’2013, Las Vegas (USA), 2013,
To appear in Advanced Materials Research or Materials Science Forum.
IV. M. Calmunger, G. Chai, S. Johansson and J. Moverare, Damage and
fracture behaviours in aged austenitic materials during high temperature
slow strain rate testing, Presented at MSMF7, Brno (Czech Republic),
2013, To appear in Key Engineering Materials.
V. M. Calmunger, R. L. Peng, G. Chai, S. Johansson and J. Moverare,
Advanced microstructure studies of an austenitic material using EBSD
in elevated temperature in-situ tensile testing in SEM, Presented at
MSMF7, Brno (Czech Republic), 2013. To appear in Key Engineering
Materials.
VI. M. Calmunger, G. Chai, S. Johansson and J. Moverare, Influence of
high temperature ageing on the toughness of advanced heat resistant
materials, Presented at ICF13, Beijing (China), 2013.
vii
Own contribution to the papers included:
In all papers above, I have been the main contributor of the microstructure
investigation, evaluation and manuscript writing. In addition, I have performed all the slow strain rate tensile testing in the listed papers and also
conducted the in-situ tensile testing. All other mechanical testing have been
performed at AB Sandvik Materials Technology in Sandviken, Sweden.
Papers not included in this thesis:
VII M. Calmunger, G. Chai, S. Johansson and J. Moverare, Damage and
fracture behaviours in advanced heat resistant materials during slow
strain rate test at high temperature, Presented at ICF13, Beijing (China),
2013.
VIII M. Lundberg, M. Calmunger and R. L. Peng, In-situ SEM/EBSD study
of deformation and fracture behaviour of flake cast iron, Presented at
ICF13, Beijing (China), 2013.
IX M. Calmunger, G. Chai, S. Johansson and J. Moverare, Influence of dynamic strain ageing on damage in austenitic stainless steels, Presented
at ECF19, Kazan (Russia), 2012.
viii
Contents
Abstract
iii
Acknowledgement
v
List of Papers
vii
Contents
ix
Abbreviation
xi
Part I
Background and Theory
xiii
1 Introduction
1.1 Introduction to the research project .
1.2 Background . . . . . . . . . . . . . .
1.3 Purpose of research and future work
1.4 Structure of the thesis . . . . . . . .
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1
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2 Austenitic alloys
2.1 Austenitic stainless steels . . .
2.1.1 Main alloying elements
2.1.2 Precipitation . . . . .
2.2 Nickel-base alloys . . . . . . .
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phenomena
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3 Microstructural mechanisms and
3.1 Deformation mechanisms . . . .
3.1.1 Dislocation movement .
3.1.2 Twinning . . . . . . . .
3.2 Softening phenomena . . . . . .
3.2.1 Dynamic recovery . . . .
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ix
3.3
3.2.2 Dynamic recrystallization . . . . . . . . . . . . . . . . 12
Dynamic strain ageing . . . . . . . . . . . . . . . . . . . . . . 12
4 Experimental and analytical methods
4.1 Material . . . . . . . . . . . . . . . .
4.2 Tensile deformation . . . . . . . . . .
4.3 Impact toughness testing . . . . . . .
4.4 Microscopy . . . . . . . . . . . . . .
4.4.1 Specimen preparation . . . . .
4.4.2 Scanning electron microscopy
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5 Summary of appended papers
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6 Conclusions
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Bibliography
29
Part II
35
Papers Included
Paper I: Deformation and damage behaviours of austenitic alloys in the dynamic strain ageing regime
39
Paper II: Influence of deformation rate on mechanical response
of an AISI 316L austenitic stainless steel
57
Paper III: Mechanical behaviours of Alloy 617 with varied strain
rate at high temperatures
65
Paper IV: Damage and fracture behaviours in aged austenitic
materials during high temperature slow strain rate testing 73
Paper V: Advanced microstructure studies of an austenitic material using EBSD in elevated temperature in-situ tensile
testing in SEM
79
Paper VI: Influence of high temperature ageing on the toughness of advanced heat resistant materials
85
x
Abbreviation
AUSC
BCC
BSE
DB
DRV
DRX
DSA
EBSD
ECCI
EDS
FCC
FEG
FIB
GAM
GB
IPF
LAGB
LCF
PLC
RT
SB
SEM
SF
SFE
SSRT
TB
TWIP
advanced ultra-super critical
body-centred cubic
backscattered electron
deformation band
dynamic recovery
dynamic recrystallization
dynamic strain ageing
electron backscatter diffraction
electron channeling contrast imaging
energy-dispersive system
face-centred cubic
field emission gun
focused ion beam
grain average misorientation
grain boundary
inverse pole figure
low angle grain boundary
low cycle fatigue
Portevin-Le Châtelier
room temperature
slip band
scanning electron microscopy
Schmid factor
stacking-fault energy
slow strain rate tensile testing
twin boundary
twinning induced plasticity
xi
Part I
Background and Theory
1
Introduction
1.1 Introduction to the research project
This licentiate thesis is a part of the ongoing research project Long term
high temperature behaviour of advanced heat resistant materials. The project
started with a M.Sc. thesis work [1] in the summer of 2011 and is carried
out in a strong collaboration between Linköping University and AB Sandvik Materials Technology in Sandviken, Sweden. It is financed through the
Research Consortium of Materials Technology for Thermal Energy Processes
(KME), grant no. KME-501. The purpose of research within KME is to
make thermal energy processes more effective and is financially supported by
both industries (60%) and the Swedish Energy Agency (40%).
1.2 Background
The research project is mainly concerned with two different groups of austenitic
alloys, austenitic stainless steels and nickel-base alloys. Some of them are
used in the biomass power plants of today and some are potential materials for the next generations biomass power plant [2]. Biomass is the largest
global contributor to renewable energy and has a great potential to expand
in the production of heat and electricity [3]. It is a sustainable fuel because
it gives no net contribution of CO2 to the atmosphere and it can be considered endless [3, 4]. However, the global increase in energy consumption
and the increase in emissions of greenhouse gases (e.g. CO2 ) causing global
warming, make needs for both an increase in energy production and a reduction of greenhouse gas emission [3, 5]. One way to accomplish both needs
is to increase the efficiency of biomass power plants, which could be reached
1
PART I.
BACKGROUND AND THEORY
by increasing temperature and pressure in the boiler sections [6]. Thus, the
requirement of more energy production is met and since biomass has no net
contribution of CO2 to the atmosphere less emissions of greenhouse gases is
the result.
Both groups of the austenitic alloys show good mechanical and chemical
properties at the operation temperatures of today’s biomass power plants [2].
However, the materials used for the future biomass power plants with higher
efficiency are required to display improved properties as higher yield strength,
creep strength and high-temperature corrosion resistance. The performance
of austenitic stainless steels at these elevated temperatures is not fully understood, but the nickel-base alloys are already operating at such conditions
in other applications and the group of nickel-base alloys is a possible option
[2, 7]. The nickel-base alloys are more expensive than the austenitic stainless
steels and the austenitic stainless steels are therefore an interesting option to
investigate as a material for the future biomass power plants. The austenitic
stainless steels are of main concern in this thesis and the nickel-base alloys
acting reference materials.
1.3 Purpose of research and future work
In the general goals of KME’s program period 2010-2013 it is stated that:
"The program will contribute to the conversion to a sustainable
energy system by development of more effective energy processes."
The purposes of this research project is to improve knowledge regarding
tensile deformation and cracking behaviour during very slow deformation and
the influence of long term ageing and tough environments, such as biomass
fuel, on structure integrity and safety for advanced heat resistant materials. Moreover, an evaluation of the degradation mechanism stress relaxation
cracking [8–11] is aimed to be achieved. In addition, the obtained knowledge will facilitate the development of new materials for the new generation
power plants with better performance with respect to temperature, pressure
and environmental conditions.
The aims of this licentiate thesis are to increase the knowledge concerning:
1. tensile deformation, damage and cracking behaviour during very slow
deformation
2
CHAPTER 1. INTRODUCTION
2. influence of long term ageing on toughness of austenitic alloys for the
future power plants.
Future work will cover the whole purpose of the research project and be
finalized in a Ph.D. thesis. It will partly consist of further evaluation of the
topics presented in this thesis. The long term ageing experiment will continue
and is planned to last for 30 000 hours, its influence on structure integrity and
safety will be evaluated using impact toughness and fracture toughness tests.
The stress relaxation cracking behaviour will be evaluated using mechanical
and thermal testing. Cyclic corrosion tests will be performed to investigate
how tough environment influence structure integrity and safety.
1.4 Structure of the thesis
This thesis consists of two parts:
• Part I: Background and Theory
• Part II: Papers Included
Part I, Background and Theory, cover an introduction of the research
project to the reader, also aims and future work are presented. Then information concerning the austenitic alloys and experimental details follows.
A summary of the included papers is provided the reader and finally the
conclusions of the thesis are given.
Part II, Papers Included, is based on six papers and describes the main
research that has been conducted in the project.
3
2
Austenitic alloys
This chapter provides general information about the austenitic alloys addressed in this thesis, austenitic stainless steels and nickel-base alloys. The
review will concentrate on the austenitic stainless steels.
2.1 Austenitic stainless steels
The main feature of stainless steels are their resistance to corrosion. They
also possess high ductility and toughness over a wide range of temperatures
and exhibit excellent high-temperature oxidation resistance [2, 12–15]. Stainless steels can be divided into five grades, ferritic, austenitic, martensitic,
dual and multiphase, and precipitation hardening. Four of them are based
on the characteristic crystallographic structure of the alloys in the grades,
ferritic, austenitic, martensitic and dual phase. The fifth grade, precipitation hardening is based on the type of heat treatment used rather than
microstructure [12–14]. Since this thesis only considers the chromium-nickel
alloyed austenitic stainless steel, the other grades will not be covered in the
review.
Austenitic stainless steels have a face-centred cubic (FCC) crystallographic
structure. They are the most commonly used and the grade containing the
largest number of alloys, of the stainless steel grades. Austenitic stainless
steels possess great corrosion resistance, good creep resistance and excellent
ductility, formability and toughness [12–14, 16, 17]. These materials exhibit
no ductile to brittle transition temperature, except for austenitic stainless
steels with very high content of nitrogen. In addition, they cannot be hardened by heat treatment but substantially hardened by cold work. Austenitic
stainless steels have relatively low yield strength but higher work hardening
5
PART I.
BACKGROUND AND THEORY
rates compared with other stainless steel grades [12, 14]. The alloys contains typically 16-26 wt.% Cr, 8-25 wt.% Ni and 0-6 wt.% Mo, and are fully
austenitic from well below RT to melting temperature [12, 14]. These alloys are non-ferromagnetic due to their crystallographic structure and have
greater heat capacity and thermal expansion, with lower thermal conductivity than other stainless steels grades [12].
2.1.1 Main alloying elements
Austenitic stainless steels are iron-based and the main alloying elements in
these materials are chromium, nickel, manganese, molybdenum, titanium,
niobium, carbon and nitrogen [12, 17].
Chromium is mainly added to obtain corrosion resistance, it reacts rapidly
with oxygen which creates a protective layer of chromium oxide on the surface. If the amount of chromium is 12 wt.% or more the oxide layer will selfrepair if it gets damaged, because of the rapid reaction between chromium
and oxygen [12].
Nickel stabilises the FCC structure in iron. Nickel increases the size of
the austenitic field, while nearly eliminating body-centred cubic (BCC) ferrite
structure from the iron-chromium-carbon alloys. Together with chromium it
produces high-temperature strength and scaling resistance.
Manganese forms austenite and can be used to replace nickel. Manganese
improves the solubility of nitrogen.
Molybdenum improves both the local and the general corrosion resistance. Molybdenum is a ferrite stabiliser and must therefore be balanced
with austenitic stabilisers to maintain the austenitic structure. It improves
the creep properties [12, 17].
Titanium stabilises the austenitic stainless steel against intergranular corrosion, if the carbon content is high. Titanium reacts more easily with carbon than chromium does, thus, titanium carbides are formed in preference
to chromium carbides and localised reduction of chromium is prevented. It
also greatly improves the creep strength [17].
Niobium creates carbides more easily than chromium and is therefore
used for intergranular corrosion resistance [12]. A too great amount niobium
may reduce the creep strength [16, 17].
Carbon additions stabilise the austenitic phase, but has a negative effect
on corrosion resistance, due to the formation of chromium carbides. If the
carbon content is below about 0.03 wt.%, the carbides do not form and the
steel is virtually all austenitic at room temperature [12].
Nitrogen additions stabilises the austenitic phase and strengthens the
material through solid solution [12, 18], leading to enhanced creep life and
6
CHAPTER 2. AUSTENITIC ALLOYS
low temperature yield strength. However, a too great amount of nitrogen
will reduce the creep life of austenitic stainless steels [17].
The stacking-fault energy (SFE) in austenitic stainless steels is influenced
by the alloying elements [19, 20]. Chromium decreases the SFE with increasing content in the austenitic stainless steels, at least within the range of 10-26
at.% chromium. Opposite to chromium, nickel increases the SFE in austenitic
stainless steels, at least within the range of 8-20 at.% nickel [20]. However,
the influence of cobalt, manganese and niobium on the SFE in austenitic
stainless steels depends on the amount of nickel. Cobalt decreases the SFE,
and the decrease is stronger in alloys with high nickel content. Manganese
decreases the SFE in alloys with <16 at.% nickel content. Niobium strongly
increases the SFE in alloys with low nickel content, but the increasing effect is weaker when increasing the nickel content [19]. In addition, the SFE
influence deformation mechanisms as twinning and dislocation movement,
twinning and dislocation mechanisms are further explained in section 3.1.
2.1.2 Precipitation
Austenitic stainless steels containing 18 wt.% chromium and 12 wt.% nickel
should be fully austenitic at high temperatures. However, the addition of
alloy elements may results in precipitation of carbides, nitrides and intermetallic phases. These precipitates may not be desirable since they can affect
the mechanical and corrosion properties. Their appearances depend on the
chemical compositions [12, 17, 21]. Only the most common precipitates in
austenitic stainless steels will be considered in this section.
M23 C6 is a carbide with FCC structure, usually containing chromium as
the main metallic element (M) but nickel, molybdenum and iron can be a
substitute for chromium. It nucleates very easily and therefore it appears
early in the precipitation process and it has been found after only 30 min at
750◦ C in an austenitic stainless steel. It can be located in grain boundaries
(GBs), twin boundaries (TBs) and intragranular sites. M23 C6 most often
appears in GBs and is often connected to intergranular corrosion [12, 13, 17].
σ-phase is an intermetallic phase with tetragonal structure, usually containing chromium and molybdenum [12, 13, 17, 21], but several other compositions have been reported [17]. It can appear after around 1000 hours
ageing at temperatures around 700◦ C. It can be located in GBs, TBs and
intragranularly. This brittle intermetallic phase affects mechanical properties
and influences the corrosion resistance by removing chromium and molybdenum from the matrix material [12, 17, 21].
Laves phase has a hexagonal structure, it contains mostly of iron and
molybdenum, a medium content of chromium and nickel and a small amount
7
PART I.
BACKGROUND AND THEORY
of Manganese and silicon. It precipitates intragranularly, often in small
amounts, after 1000 hours ageing at temperatures between 600◦ C and 800◦ C
[12, 17].
G phase is a silicide with FCC structure, it mostly contains of nickel
and it is titanium or niobium rich. It precipitates in GBs after less than 10
hours at temperatures between 700◦ C and 800◦ C and after longer time as
intragranular precipitates [12, 17].
Other precipitates that may form in austenitic stainless steel are carbonitrides such as NbC, NbN, TiC and TiN. These are carbides and nitrides with
FCC structure and contains of niobium or titanium and carbon or nitrogen
respectively [17].
2.2 Nickel-base alloys
Nickel-base alloys are widely used in high-temperature applications at temperatures between 650◦ C and 1100◦ C, since nickel is stable in the FCC structure from room temperature up to the melting temperature. Nickel-base
alloys possess great corrosion resistance, strength, creep and fatigue properties at elevated temperatures. They consist of many alloying elements,
but most of the nickel-base alloys have 10-20 wt.% chromium, up to 8 wt.%
aluminium and titanium, 5-10 wt.% cobalt and small amounts of boron, zirconium and carbon. Optional common additives are molybdenum, tungsten
and tantalum. The chromium content is enough to create a corrosion protective chromium oxide layer, at higher temperatures a corrosion protective
aluminium oxide layer is formed. Nickel-base alloys have an austenitic matrix, called γ. γ’ is a hardening precipitate that may improve the mechanical
properties at elevated temperatures due to the ordered FCC structure, it often contains of nickel, aluminium and titanium. Other precipitates that can
appear are M23 C6 and M6 C, which often occur at GBs [7, 22].
8
3
Microstructural mechanisms and
phenomena
In this chapter the main microstructural mechanisms and phenomena related
to the included papers are covered.
3.1 Deformation mechanisms
The main plastic deformation mechanisms in the investigated austenitic alloys can be divided into two types, dislocation and twin controlled deformations. Since they influence the mechanical behaviour and are different in
appearance, a review of these mechanisms is provided.
3.1.1 Dislocation movement
A dislocation is a lattice line defect and can be divided into two different
basic types, edge and screw dislocations. Edge dislocations have the Burgers
vector oriented normal to the dislocation line and screw dislocations has
the Burgers vector parallel to the dislocation line. Unlike edge dislocations,
screw dislocations don’t have a unique slip plane but several potential slip
planes. Thus, the screw dislocation possesses greater mobility than the edge
dislocation [23, 24].
There are two basic types of dislocation movements that may generate
plastic deformation, glide and climb movement. Glide occurs when the dislocation moves in the plane along the dislocation line and the Burgers vector
and climb occurs when the dislocation moves out of the plane perpendicular
to the Burgers vector. When many dislocations glide in the same slip plane
9
PART I.
BACKGROUND AND THEORY
it results in planar slip which is a common plastic deformation mechanism
in austenitic alloys treated in this thesis [15, 23, 25]. If there is a great
number of slip steps on closely spaced parallel slip planes, slip bands will be
formed [23, 26, 27], also common in the investigated alloys [28, 29]. Screw
dislocations can change from one slip plane to another, this is called crossslip, if many screw dislocations cross-slip it results in wavy slip [23–25]. The
stacking-fault energy (SFE) influences the cross-slip mechanism. When the
SFE is low, cross-slip is restricted so that barriers to dislocation movement
remain effective to higher stress levels than in the material of higher SFE.
Thus, when the SFE decreases the slip characters changes from wavy to planar mode [24]. In face-centre cubic (FCC) metals, as austenitic alloys, slip
appears generally between one of the four close-packed {111} planes and in
one of the three <110> directions. More than one slip system can be activate, which is called multi-directional slip. Activation of other slip systems
is rarely observed [23, 24]. To activate a slip system a critical shear stress is
required, this shear stress acting on a slip plane can be calculated as,
τ=
F
cos φ cos λ
A
(1)
Where τ is the resolved shear stress from the force F acting on crosssection area A, φ is the angle between F and the normal to the slip plane
and λ is the angle between F and the slip direction. The quantity of cosφ
cosλ is called the Schmid factor [23, 27].
Dislocation climb is dependent on diffusion and is for that reason thermally activated and temperature dependent, when atoms diffuse it enables
edge dislocations to move out of its original slip plane [23].
Temperature influences the energy that has to be provided for dislocations
to surmount the obstacles they encounter during slip. If the conditions are
sufficient, thermal vibrations of the crystal atoms may assist the dislocation
to surmount obstacles at lower values of applied stress than that required at
0 K. Thus, an increase in temperature, or a reduction in applied strain rate,
will reduce the flow stress [23].
3.1.2 Twinning
Twins may occur during different processes depending on origin. Annealing
twins nucleate during thermal processes [30], transformation twins come from
phase transformation and deformation twins nucleate from deformation [31–
33]. This review will concentrate on the later type.
10
CHAPTER 3. MICROSTRUCTURAL MECHANISMS AND PHENOMENA
Deformation twins are initiated by a certain shear stress, which is higher
than the stress needed for growth of an existing twin. The twinning process effects a rotation of the lattice such that the atom position in the twin
represent a mirror image of the atoms in the matrix material [24, 31, 32].
The angle of the boundary between matrix material and twin, called twin
boundary, get due to the mirror rotation a certain value, in austenitic alloys
often 60◦ [31, 32]. For FCC metals the critical twinning stress for initiation
of twins is slightly influenced by temperature and strain rate, where the critical twinning stress increases with increasing temperature and strain rate.
However, SFE have a larger influence on the critical twinning stress and it
increases with increasing SFE. Also the grain size influence the twinning, a
larger grain shows much greater twinning density than a smaller grain [32].
Formation of each deformation twin leads to a certain shear strain. This
will increase plasticity of the material if a large number of twins have been
formed [31, 32, 34], which is called twinning induced plasticity (TWIP) [34].
3.2 Softening phenomena
During deformation at elevated temperatures the softening phenomena dynamic recovery (DRV) and dynamic recrystallization (DRX) may occur, both
in hot working (strain rate range 1-100s−1 ) and slow creep deformation (strain
rates below 10−5 s−1 ). Signs of DRV and DRX both appear in the stress and
strain curve during flow stress and in the microstructure, see Fig. 1. Some of
the investigated alloys in this thesis have showed such signs of DRV and DRX,
see paper I, II and III. The phenomena affects the mechanical properties at
elevated temperatures and therefore some information is given.
3.2.1 Dynamic recovery
DRV is influenced by dislocation density, when the flow stress increases during the first stage of deformation and due to dislocation interaction and reproduction the rate of recovery increases when the dislocation density increases.
At this stage low angle grain boundaries (LAGB &5◦ ) and subgrains develop
in the microstructure [23, 35]. This process involves glide and climb of the
dislocations which form LAGB. Since climb involves diffusion a sufficient
temperature for thermal activation of diffusion of point defects is needed.
During DRV some dislocation annihilation occurs. DRV can be observed in
the stress-strain curve, a steady-state flow stress will be obtained due to a
dynamic equilibrium between the rates of recovery and hardening [35, 36].
In Fig. 1 (b) both DRV and DRX are active during the process of plastic
11
PART I.
BACKGROUND AND THEORY
flow [36].
3.2.2 Dynamic recrystallization
DRX may occur in alloys with low or medium SFE and initiates when a
critical strain has been reached at elevated temperature. The critical strain
for initiation of DRX not only depends on deformation rate and temperature
but also on chemical composition and initial grain size [23, 35–38]. During
DRX new grains originate at the old high angle grain boundaries but due
to the continuous deformation of the material, the dislocation density of the
new grains increases. This will reduce the driving force for further growth of
the new grains, and eventually the growth will stop. Since DRX originates
at existing high angle grain boundaries during straining, DRX may appear
at boundaries as deformation bands [35]. Thus, DRX shows recrystallized
grains in the microstructure often near or at high angle grain boundaries
[35, 37], Fig. 1 (a) shows recrystallized grains in AISI 316L after slow tensile
deformation. Other sign of DRX is the decreasing flow stress in the stress
and strain curve [35, 36], shown in Fig. 1 (b).
(a)
2 µm
(b)
Figure 1: Recrystallized grain structure from DRX (a) and a stress and strain
curve showing serrated yielding and the shape is due to DRV and DRX (b) in a
slow strain rate tensile tested AISI 316L at 650◦ C using a strain rate of 10−6 s−1
(based on paper II).
3.3 Dynamic strain ageing
Dynamic strain ageing (DSA) originates from interaction between solute
atoms and dislocations during plastic deformation. Under plastic flow dislo12
CHAPTER 3. MICROSTRUCTURAL MECHANISMS AND PHENOMENA
cations are gliding until they come across an obstacle where they are stationary until the obstacles are surmounted. When the dislocations are stationary
solute atoms can diffuse towards the dislocations which result in an increase
in the activation energy for re-activation and consequently also an increase
in the stress needed for overcoming the obstacle [39–44]. Thus, DSA is directly influenced by the deformation rate that affects the mobility of the
dislocations and the temperature that influences the diffusion rate of solute
atoms. At temperatures below 350◦ C carbon is responsible for DSA while
nitrogen and/or substitutional chromium atoms are responsible at higher
temperatures (400◦ C to 650◦ C) [18, 29]. It has been reported that mechanical properties like strength and ductility may be significantly changed due
to DSA [29, 45]. DSA is characterized by serrated yielding occurring in the
stress-strain curve, denoted as Portevin-Le Châtelier (PLC) effect or jerky
flow, Fig. 1 (b) show serrated yielding in AISI 316L. DSA can also lead to an
increase in flow stress, work hardening rate and most important a negative
strain rate sensitivity [46, 47] and DSA influence on ductility depends on the
alloy composition [15, 46].
The PLC effect is created by the pinning and unpinning of dislocations
and is recognized by serrated yielding in stress and strain curves [15, 48–50].
There are different types of PLC effects and they have designated appearance
[15, 48, 50]. Type A is considered as locking serrations, they abruptly rise and
then drop to a stress level below the general level. Type B is characterized
by small oscillations about the general level of the curve. Type C leads to
unlocking serration which is when the curve abrupt drops below the general
stress level. Type D is characterized by plateaus on the curve [50]. Serrated
yielding may also come from other mechanisms, e.g. twinning [24, 32].
13
4
Experimental and analytical methods
In this chapter, the conducted experimental and analytical methods are presented. The slow strain rate tensile testing (SSRT) using a strain rate of
10−6 s−1 and in-situ tensile tests were performed at Linköping University
(LiU) and most of the other mechanical tests were performed at AB Sandvik
Materials Technology (SMT) in Sandviken, Sweden. All ageing and solution
heat treatments have been done at SMT, except the ageing processes for the
specimens used in paper IV which were performed at LiU. The microscopy
analysis has mainly been performed at LiU but some of microstructural investigations have been made at SMT.
4.1 Material
All tested materials have been supplied by SMT in solution heat treatment
conditions according to table 1. Five austenitic stainless steels (AISI 304,
AISI 310, AISI 316L, Sanicro 25 and Sanicro 28,) and two nickel-base alloys
(Alloy 617 and Alloy 800HT) have been used in the conducted experiments.
The nominal chemical composition in wt% for each alloy is presented in table 2. All the specimens were manufactured at SMT after the heat treatment,
except the in-situ tensile specimens that were prepared at LiU.
4.2 Tensile deformation
Several uniaxial tensile tests have been performed within this project, different conditions as different temperatures and strain rates have been used.
From tensile testing many mechanical properties can be obtained, e.g. tensile
15
Alloy
AISI 304
AISI 310
AISI 316L
Sanicro 25
Sanicro 28
Alloy 617
Alloy 800HT
b balance
Alloy
1060
1050
1050
1250
1150
1175
1200
Temperature [ ◦ C]
15
10
10
10
15
20
15
Time [min]
Table 1: Solution heat treatments.
AISI 304
AISI 310
AISI 316L
Sanicro 25
Sanicro 28
Alloy 617
Alloy 800HT
C
0.35
0.55
0.4
0.25
0.43
0.04
0.71
Si
1.2
0.84
1.7
0.47
1.83
0.02
0.5
Mn
18.3
25.43
17.0
22.33
27.02
22.53
20.32
Cr
10.3
19.21
12.0
24,91
30.76
53.8
30.06
Ni
0.11
2.6
0.24
3.39
9.0
0.005
Mo
0.3
0.08
2.95
0.9
0.011
0.053
Cu
0.05
3.37
0.02
0.02
0.01
W
1.44
0.088
12.0
0.031
Co
0.01
0.52
0.02
0.01
Nb
0.07
0.04
0.236
0.047
0.005
0.013
N
0.001
0.005
0.003
0.46
0.52
Ti
0.031
0.0094
0.47
Al
0.046
0.054
0.017
0.048
V
Table 2: Nominal chemical composition in [wt%] of the austenitic alloys.
0.015
0.046
0.04
0.067
0.019
0.061
0.063
b
Fe
b
b
b
b
b
1.1
16
BACKGROUND AND THEORY
PART I.
CHAPTER 4. EXPERIMENTAL AND ANALYTICAL METHODS
strength, elongation to fracture, etc., using the stress and strain curve to illustrate those properties. In this thesis only engineering stress and strain curves
are considered. For the tensile testing a Roell-Korthaus and an Instron 5982
tensile test machine were used, the later is shown in Fig. 2. The machines
were equipped with an MTS 653 furnace and a Magtec PMA-12/2/VV7-1 extensometer and an Instron SF16 furnace and an Instron 7361C extensometer
respectively, both used in lab air environment. For the tensile tests roundbar specimens with a diameter of 5 mm and a gauge length of 50 mm were
used.
The tensile tests were carried out at different conditions. Strain rate
from 10−2 s−1 down to 10−6 s−1 and temperatures at 23◦ C referred to as room
temperature (RT), 400◦ C, 500◦ C, 600◦ C, 650◦ C and 700◦ C were used. The
SSRT was performed on the electromechanical tensile test machine showed
in Fig. 2.
Water cooled
grip
Furnace
Extensometer
Specimen with
thermocouples
Figure 2: Electromechanical tensile test machine used for SSRT equipped with
furnace, extensometer and water cooled grip.
The in-situ tensile testing was performed inside a HITACHI SU-70 FEGSEM scanning electron microscope (SEM) using a specially designed Gatan
microtest tensile test stage, Fig. 3 (a) show the stage that is tilted 70◦ for
optimal diffraction. The miniature tensile stage can produce a force of maximum 5kN. A small specimen showed in Fig. 3 (b) was used. The thickness
17
PART I.
BACKGROUND AND THEORY
of the specimens were ground down to less than 1 mm, then further preparation of one side of the specimen to enable the use of electron backscatter
diffraction (EBSD). The procedure is described in detail under section 4.4.1
later in this chapter.
(a)
Clamps/heaters
(b)
Specimen
Figure 3: Miniature tensile test stage (a) and a drawing of the small specimen
used for in-situ tensile test (b).
4.3 Impact toughness testing
The impact toughness tests were performed using the Charpy V method
according to ISO 14556 standard. Samples with a dimension of 10x10x55mm
and V type were used. The specimens were aged at 650◦ C and 700◦ C in air
environment for 1000 and 3000 hours before the toughness tests. The impact
toughness testing have been performed at room temperature, two to three
specimens at each ageing condition and non-aged specimens were tested.
4.4 Microscopy
4.4.1 Specimen preparation
To enable the use of electron channeling contrast imaging and electron backscatter diffraction, a careful specimen preparation must be performed. Both techniques are surface sensitive due to scattering of the backscattered electrons.
Thus, it is crucial to minimize the surface roughness before the microscopy.
Both electrolyte and mechanical polishing can accomplish the critical surface
preparation, in this project only the later one has been used according to the
following steps:
18
CHAPTER 4. EXPERIMENTAL AND ANALYTICAL METHODS
1. 500 SiC-paper (30 µm), 2 min
2. 1200 SiC-paper (15 µm), 2 min
3. 4000 SiC-paper (5 µm), 3 min
4. Silk cloth, diamond suspensions (3 µm), 5 min
5. Woven wool cloth, diamond suspensions (1 µm), lubricant, 10 min
6. Rayon-viscose fibres cloth, diamond suspensions (0,25 µm), lubricant,
15 min
7. Neoprene foam cloth, colloidal silica suspension (0,04 µm), 5 min
8. Neoprene foam cloth, water, 1 min
After each step, specimens and holder were cleaned carefully using water and
ethanol for step 1-2, and water + detergent, ethanol and ultrasonic cleaning
for step 3-9.
4.4.2 Scanning electron microscopy
The microstructural investigations were performed using different SEM related techniques like electron channeling contrast imaging (ECCI), electron
backscatter diffraction (EBSD) and energy-dispersive system (EDS).
To capture deformation, damage and even dislocation and twin structures
in the highly deformed alloys, the ECCI technique was used [51, 52]. In 1967
Coates [53] reported that the intensity of backscattered electrons is strongly
dependent on crystal orientation in the scanning electron microscope. The
same year, Booker et al. [54] suggested that the effect could be used to image
crystal defects near the surface of a bulk specimen using a SEM, since close
to the Bragg condition the backscattered electron intensity varies rapidly
with orientation. This forms the basis of the technique now called ECCI.
Thus, ECCI uses the interaction between backscattered electrons and the
crystal planes to generate contrast resulting in an image where local misorientation, defects and strain fields are shown as contrast variations [52–57].
Fig. 4 shows the sample position perpendicular to the incident electron beam
in the ECCI set up. Gutierrez-Urrutia et al. [52] studied dislocation structures and deformation twins using the ECCI and EBSD techniques in a SEM
and bright-field transmission electron microscopy of the same area in a tensile deformed TWIP steel. They found that ECCI could image dislocation
cell structures and mechanical twins of 30 nm thickness. Thus, ECCI is a
19
PART I.
BACKGROUND AND THEORY
powerful tool in characterizing highly deformed alloys. Fig. 5 shows dislocation structure and deformation as a plastic zone in front of a crack tip
and slip bands in an austenitic alloy, using the ECCI technique. Moreover,
the acceleration voltage influences the contrast in ECCI micrographs, where
decreasing acceleration voltage improves the contrast, according to GutierrezUrrutia et al. [52] there are two main possible reasons for this. First reason,
the backscatter yield increases with decreasing acceleration voltage, which
gives that the parts with higher intensity appear brighter while parts with
low intensity maintain dark. Second reason, the interaction volume decreases
with decreasing acceleration voltage, the lattice strain shows smaller variations in a smaller volume, thus diffraction conditions and contrast are better
defined. ECCI investigations were performed on a HITACHI SU-70 field
emission gun (FEG)-SEM and a Zeiss XB 1540 FEG-SEM, both equipped
with a solid state 4-quadrant backscattered electron (BSE) detector, using
10 kV acceleration voltage and working distances between 5 mm to 7 mm
[52, 58].
Incident electron beam
BSE detector
Backscattered
electrons
90
Specimen
Figure 4: ECCI set up with the incident electron beam 90◦ angled to the specimen, based on Gutierrez-Urrutia et al. [52].
The EBSD technique provides information about phases and crystallographic orientation, using Kikuchi patterns. As a simple analogy the grains
in a polycrystalline alloy act as a mirror reflecting it’s orientation to the
observers camera. Determination of characteristic crystallographic parameters as crystal plane spacing and angles between planes allow the phaseidentification. Crystallographic orientations are determined by identification
of the Kikuchi pattern [58–60]. To use EBSD the specimen has to be tilted
around 70◦ against the incident electron beam to provide the optimal diffraction [58, 60]. EBSD can be used to evaluate plastic deformation, because the
degree of deformation or damage can be expressed as local crystal reorienta20
CHAPTER 4. EXPERIMENTAL AND ANALYTICAL METHODS
(a)
(b)
Crack tip
100 nm
10 µm
Figure 5: Dislocations and dislocation ends in Sanicro 25 (a) and plastic zone in
front of a crack tip, slip bands and local strain concentrations at grain boundaries
in Alloy 617 (b).
tions of grains [60–62]. Also, the grain average misorientation (GAM) has a
linear relationship with plastic strain [60, 62]. Information about active slip
systems in the microstructure can be provided by the EBSD technique, using
obtained Euler angles and Equation 1 showed in section 3.1.1 [63]. EBSD investigations were performed in a HITACHI SU-70 FEG-SEM equipped with
an OXFORD EBSD detector and a 6500 F JEOL FEG-SEM equipped with a
TSL OIM EBSD system, both EBSD systems used the HKL software CHANNEL 5. The EBSD-maps were measured at 15 kV and 20 kV acceleration
voltage using a working distance of 12 mm up to 25 mm and step sizes of 0,1
µm up to 4 µm were used [58, 60].
The EDS technique provides chemical compositional information of the
specimen [58]. It is mainly used to collect information about precipitation
that occurs during the ageing and the long deformation process of SSRT. The
EDS investigations were performed using both a HITACHI SU-70 FEG-SEM
and a 6500 F JEOL FEG-SEM.
21
5
Summary of appended papers
Paper I
Deformation and damage behaviours of austenitic alloys in the
dynamic strain ageing regime
The purpose of this paper was to investigate the deformation and damage
mechanisms related to dynamic strain ageing (DSA) in three austenitic stainless steels (AISI 310, AISI 316L and Sanicro 25) and two nickel-base alloys
(Alloy 617 and Alloy 800HT). The materials were investigated by tensile
testing at different elevated temperatures (400◦ C to 700◦ C). One of the materials was low cycle fatigue (LCF) tested at 650◦ C using a strain range of
1,2%. Since DSA is present in different temperature ranges depending on
alloy composition, a scanning electron microscope (SEM) investigation was
mainly performed on the specimens tested at 650◦ C and 700◦ C where all five
materials displayed DSA.
As expected deformation in the DSA regime is planar slip and slip bands
in single and multi-direction. Local damage has been connected to interaction
between slip bands and/or interaction between twins and grain boundaries.
Ductility is affected differently by temperature depending on alloy tested,
DSA is not always related to a low ductility of the material, it is suggested
that DSA may introduce a phenomenon similar to twinning induced plasticity
(TWIP) that could increase the ductility.
23
PART I.
BACKGROUND AND THEORY
Paper II
Influence of deformation rate on mechanical response of an AISI
316L austenitic stainless steel
This paper investigates the high-temperature behaviour during uniaxial slow
strain rate tensile testing of an AISI 316L material, commonly used for components in power plants. Uniaxial slow strain rate tensile testing (SSRT)
has been performed at different temperatures up to 700◦ C using strain rates
down to 10−6 s−1 . An investigation of the microstructure was conducted on
the deformed and fractured materials using mainly the scanning electron microscopy methods, electron channeling contrast imaging (ECCI) and electron
backscatter diffraction (EBSD), to capture the microstructural mechanisms
coupled to the mechanical behaviours seen in engineering stress-strain curves.
It was found that DSA occurs in AISI 316L during tensile testing at
temperatures of 650◦ C and 700◦ C when using a strain rate of 10−6 s−1 . The
strength decreases with increasing temperature and decreasing strain rate.
Elongation on the other hand increases with decreasing strain rate applied
at the same temperature, this has been observed at both room temperature
(RT) and elevated temperature (up to 700◦ C). Dynamic recrystallization
(DRX) can occur during tensile deformation at elevated temperature. At
a low strain rate (10−6 s−1 ) DRX are more homogeneously spread in the
microstructure. The presence of dynamic recovery together with dynamic
recrystallization can be seen to effect the appearance of the stress-strain
curves at 650◦ C and 700◦ C when using a low strain rate.
Paper III
Mechanical behaviours of Alloy 617 with varied strain rates at
high temperatures
This study focuses on the deformation and damage mechanisms in Alloy 617
deformed using low strain rates (down to 10−6 s−1 ) at elevated temperatures
(650◦ C and 700◦ C) by uniaxial SSRT and microstructure evaluation using
scanning electron microscopy techniques as ECCI.
DSA can occur in Alloy 617 at temperatures between 650◦ C and 700◦ C
with strain rate from 10−2 s−1 down to 10−6 s−1 . TWIP is one of the mechanisms for a high elongation during DSA. Both strength and elongation increase with decrease of strain rate down to 10−4 s−1 , and then both decrease
with further decrease of strain rate. Micro and nano DRX can occur during
24
CHAPTER 5. SUMMARY OF APPENDED PAPERS
the tensile deformation with very low strain rates. Repeated DRX can lead
to the formation of damage in the material.
Paper IV
Damage and Fracture Behaviours in Aged Austenitic Materials
During High-Temperature Slow Strain Rate Testing
The aim of this study was to investigate damage and fracture mechanisms of
high-temperature long term aged austenitic materials (the austenitic stainless steel AISI 304 and the nickel-base alloy Alloy 617) during uniaxial
SSRT at RT and elevated temperature. The role of precipitation from hightemperature ageing and the long deformation process is evaluated by microscopy and coupled to the damage and fracture behaviour.
The investigation showed that SSRT caused intergranular cracking in
both high-temperature long term aged AISI 304 and Alloy 617 at both RT
and 700◦ C when using a strain rate of 10−6 s−1 . At RT the fracture is caused
by cracks initiated due to stress concentration formed by the precipitates
from the ageing process in the grain boundaries (GBs) for both alloys. Alloy
617 also exhibit crack initiation and propagation by slip band interaction
with the small GB precipitates. At 700◦ C the fracture is caused by GB
precipitates formed during both the ageing process and the tensile deformation. Elongation to fracture decreases for both the aged stainless steel and
the aged nickel-base alloy when a lower strain rate is used compared with a
higher strain rate at 700◦ C.
Paper V
Advanced Microstructure Studies of an Austenitic Material Using EBSD in Elevated Temperature In-Situ Tensile Testing in
SEM
In this study in-situ tensile testing was performed on Sanicro 25 at two different temperatures. An investigation of the influence of temperature on the
deformation behaviour was performed using the EBSD technique. Fracture
behaviour will also be discussed.
The analysis from in-situ tensile test at RT and 300◦ C in a SEM together
with EBSD in Sanicro 25 has shown that larger grains may tend to accumulate more local plastic strain for the same macroscopic strain values at
25
PART I.
BACKGROUND AND THEORY
both temperatures. Somewhat higher plastic strains at grain level can be
obtained at RT compared to elevated temperature for the same macroscopic
strain value. Cracks initiate and propagate along the slip system(s) with the
highest Schmid factor at RT.
Paper VI
Influence of High Temperature Ageing on the Toughness of Advanced Heat Resistant Materials
In this paper the influence of precipitation and growth of precipitates on
toughness due to different chemical compositions and high-temperature treatment are investigated. The experiments were conducted on two austenitic
stainless steels (AISI 304 and Sanicro 28) and one nickel-base alloy (Alloy
617). Impact toughness tests have been performed and the fracture surface
and cross-section have been investigated using microscopy. Thermocalc has
been used to predict possible precipitates.
From the study it was found that the fracture initiation and propagation in the aged austenitic stainless steel is very local. The initiation and
propagation of fracture behave differently in these materials due to different
chemical compositions affecting nucleation, growth and shape of precipitates.
The brittle σ-phase can appear in the austenitic stainless steel after 1000
hours at 650◦ C and then increases in amount. The amount and shape have
strong effect on the fracture behaviour, where needle shaped σ-phase which
mostly appear at high temperature (700◦ C) after longer ageing time (3000
hours) lead to a low impact toughness and brittle fractures both locally and
on a macro-level in the specimen. The nickel-base alloy show higher impact
toughness with increasing ageing temperature and time.
26
6
Conclusions
The presented research within this licentiate thesis deals with high-temperature
behaviour of austenitic alloys, five austenitic stainless steels and two nickelbase alloys, with focus on deformation, damage and fracture behaviour during
slow strain rate deformation and the influence of long term ageing.
It was found that the main deformation mechanisms in the austenitic
alloys are planar dislocation deformation, such as planar slip, slip bands in
single and multi-direction. Twinning has also been observed and related to
the dynamic strain ageing phenomenon, dynamic strain ageing occurred in
all tested materials but at different temperature ranges. The plastic performance of the austenitic alloys is influenced differently by temperature when
subjected to tensile deformation. Dynamic strain ageing may not always
lead to a lower ductility of the material, it is suggested that dynamic strain
ageing may introduce a phenomenon similar to twinning induced plasticity
that could increase the ductility. However, the ductility performance of some
of the tested austenitic alloys increased when subjected to a slow strain rate
tensile deformation process at elevated temperatures compared when using
a higher deformation rate at the same temperature. While others showed
an increase in ductility when decreasing to a certain deformation rate and
then the ductility performance decreased with further decreasing deformation rate. During these tests softening processes, such as dynamic recovery
and dynamic recrystallization appeared in some of the tested alloys, both in
stainless steel and nickel-base alloy. In aged conditions the tested austenitic
alloys showed only a decrease in ductility with decreasing deformation rate.
This is attributed to the formation of precipitates in the grain boundaries creating stress concentration that causes intergranular fracture. Also, in some
of these specimens dynamic recrystallization was observed.
The analysis from electron backscatter diffraction and in-situ tensile test
27
PART I.
BACKGROUND AND THEORY
in a scanning electron microscope at room temperature and elevated temperature in Sanicro 25 has shown that, larger grains may tend to accumulate
more local plastic strain for the same macroscopic strain values at both temperatures. Somewhat higher plastic strains at grain level can be obtained at
room temperature than at elevated temperature for the same macroscopic
strain value. Cracks initiate and propagate along the slip systems with the
highest Schmid factor at room temperature.
Local damage has been connected to interaction between slip bands and/or
interaction between twins and grain boundaries. Repeated dynamic recrystallization may also lead to the formation of damage.
It has also been found that the fracture initiation and propagation in
the aged austenitic stainless steel is very local during impact testing. The
initiation and propagation of fracture behave differently in the austenitic alloys due to different chemical compositions affecting nucleation, growth and
shape of the precipitates. The brittle σ-phase can appear in the austenitic
stainless steels AISI 304 and Sanicro 28 after 1000 hours at 650◦ C and then
increases in amount. The amount and shape have strong effect on the fracture behaviour, where needle shaped σ-phase which mostly appear at high
temperature (700◦ C) after longer ageing time (3000 hours) lead to a low
impact toughness and brittle fractures both locally and on a macro-level in
the specimen. However, the tested nickel-base alloy showed higher impact
toughness with increasing ageing temperature and time.
28
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34
Part II
Papers Included
The articles associated with this thesis have been removed for copyright
reasons. For more details about these see:
http://urn.kb.se/resolve?urn=urn:nbn:se:liu:diva-98242
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