MODIFICATION OF MECHANICAL PROPERTIES OF 6351 Al-Mg-Si … 2021. 1. 2. · Fig. 2.14 The phase...
Transcript of MODIFICATION OF MECHANICAL PROPERTIES OF 6351 Al-Mg-Si … 2021. 1. 2. · Fig. 2.14 The phase...
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MODIFICATION OF MECHANICAL PROPERTIES OF
6351 Al-Mg-Si ALLOY BY AGING HEAT TREATMENT
By
Ahmed Yehia Ahmed Abd El-Rahman
A Thesis Submitted to the
Faculty of Engineering at Cairo University
in Partial Fulfillment of the
Requirements for the Degree of
MASTER OF SCIENCE
In
Metallurgical Engineering
FACULTY OF ENGINEERING, CAIRO UNIVERSITY
GIZA, EGYPT
2015
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MODIFICATION OF MECHANICAL PROPERTIES OF
6351 Al-Mg-Si ALLOY BY AGING HEAT TREATMENT
By
Ahmed Yehia Ahmed Abd El-Rahman
A Thesis Submitted to the
Faculty of Engineering at Cairo University
in Partial Fulfillment of the
Requirements for the Degree of
MASTER OF SCIENCE
In
Metallurgical Engineering
Under the Supervision of
Prof. Dr. Mohamed Mamdouh Ibrahim Prof. Dr. El-Sayed Mahmoud El-Banna
Professor of Metallurgy
Mining, Petroleum and Metallurgical
Department
Faculty of Engineering, Cairo University
Professor of Metallurgy
Mining, Petroleum and Metallurgical
Department
Faculty of Engineering, Cairo University
Prof. Dr. Taher Ahmed El-Bitar
Head of Plastic Deformation Department
Central Metallurgical R&D Institute (CMRDI)
FACULTY OF ENGINEERING, CAIRO UNIVERSITY
GIZA, EGYPT
2015
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MODIFICATION OF MECHANICAL PROPERTIES OF
6351 Al-Mg-Si ALLOY BY AGING HEAT TREATMENT
By
Ahmed Yehia Ahmed Abd El-Rahman
A Thesis Submitted to the
Faculty of Engineering at Cairo University
in Partial Fulfillment of the
Requirements for the Degree of
MASTER OF SCIENCE
In
Metallurgical Engineering
Approved by the
Examining Committee
____________________________
Prof. Dr. Mohamed Mamdouh Ibrahim, Thesis Main Advisor
____________________________
Prof. Dr. El-Sayed Mahmoud El-Banna, Member
____________________________
Prof. Dr. Taher Ahmed El-Bitar, Member Central Metallurgical R&D Institute (CMRDI)
___________________________
Prof. Dr. Abd El-Hamid Ahmed Hussein, Internal Examiner
___________________________
Prof. Dr. Mohamed Abd El-WahabWaly, External Examiner Central Metallurgical R&D Institute (CMRDI)
FACULTY OF ENGINEERING, CAIRO UNIVERSITY
GIZA, EGYPT
2015
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Engineer’s Name: Ahmed Yehia Ahmed Abd El-Rahman
Date of Birth: 4/3/1989
Nationality: Egyptian
E-mail: [email protected]
Phone: 01004438769
Address: El Qlubia, El Khanka, El Qalag
Registration Date: 1/10/2011
Awarding Date: …./…./……..
Degree: Master of Science
Department: Metallurgy Departement
Supervisors: Prof. Mohamed Mamduoh Ibrahim
Prof. Elsayed Mahmoud Elbanna
Prof. Taher Ahmed El-Bitar
Examiners: Prof. Mohamed Abd El-WahabWaly (External examiner)
Central Metallurgical R&D Institute (CMRDI)
Prof. Abdel Hamid Ahmed Hussein (Internal examiner)
Prof. Mohamed Mamdouh Ibrahim(Thesis main advisor)
Prof. Elsayed Mahmoud Elbanna (Member)
Prof. Taher Ahmed El-Bitar (Member)
Central Metallurgical R&D Institute (CMRDI)
Title of Thesis:
MODIFICATION OF MECHANICAL PROPERTIES OF 6351
Al-Mg-Si ALLOY BY AGING HEAT TREATMENT
Key Words:
Artificial Aging; Natural Aging; Pre-aging; XRD; SEM; EDAX
Summary:
The present study is dealing with modification of mechanical properties of Al-Mg-Si alloy
6351 by age hardening. The study investigates the effect of aging temperature, time, natural
aging and pre-aging on artificial aging behavior in terms of mechanical properties and
fractography examination. Artificial aging after solution treatment-water quenched resulted
in a sharp increase in both ultimate tensile strength UTS and yield stress YS, can lead with a
decrease in total elongation. As the time of aging increase the strength increase slightly till
reaches peak strength after that it starts to decrease with increasing time of aging. Better
mechanical properties are observed at lower aging temperature. Natural aging at room
temperature (25 ±3oC) after solution treated-water quenched resulted in a mild increase in
tensile properties with a slight drop in total elongation, natural aging for 170 h and for 1000
h after solution treatment followed by artificial aging of this alloy at 160oC, shifted the time
to reach peak strength to shorter aging time (8- 4 h respectively) in comparison to peak-aged
condition (160oC for 18 h). Pre-aging at 100
oC for various times before artificially aging at
160oC for 18 h was investigated. It was found that the pre-aging for 10 min followed by
artificially peak aging led to slight increase in ultimate tensile strength and yield stress YS
associated with a reasonable total elongation.
ere
mailto:[email protected]
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I
AKNOWLEDMENT
First and foremost, I have to thank my research supervisors, Prof. Mohamed Mamdouh
Ibrahim, Prof. El-Sayed Mahmoud El-Banna and Prof. Taher Ahmed Al-Bitar. Without their
assistance and dedicated involvement in every step throughout the process, this thesis
would have never been accomplished. I would like to thank you very much for your
support and understanding over these past two years.
I would also like to show gratitude to my committee, including Prof. Mohamed
Mamduoh Ibrahim and Prof. El-Sayed Mahmoud El-Banna were my third-year professor in
metallurgy department at faculty of engineering, Cairo University. Their teaching style and
wide knowledge for different topics made a strong impression on me and I have always
carried positive memories of their classes with me. I discussed early versions of this work
with them. They raised many precious points in our discussion and I hope that I have
managed to address several of them here. Working with Prof. Mohamed Mamduoh Ibrahim
and Prof. El-Sayed Mahmoud El-Banna were an extraordinary experience. Much of the
analysis presented in Section IV and V is owed to my time at physical metallurgy classes I
had been through in the undergraduate level and in the postgraduate level.
I am very grateful to Prof. Taher Ahmed Al-Bitar at the Central Metallurgical Research and
Development Institute (CMRDI) kindly assisted me in my recent work, present all available
methods to accomplish my work and his experience to get a very useful suggestion and
discussion and he was very patient with my knowledge gaps in the area.
I must also thank two colleagues at the Department of Mohamed Hafez and Mustafa Ahmed
Othman, for giving me the retreat to have this thesis rushed to the printer. I would also like to
present a great thankful to Eng. Almosilhy at CMRDI for his helpful in my present work. I do
not forget Mr. Tarek a technician at CMRDI and Mechanical Testing Laboratory staff for
their efforts in preparation and testing my specimen.
Most importantly, none of this could have happened without my family. My father, my
mother and my wife, who offered me encouragement through everything limited devotion to
correspondence. Every time I was ready to quit, you did not let me and I am forever grateful.
This dissertation stands as a testament to your unconditional love and encouragement.
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Dedication
I dedicate this thesis to my parents, my brother and sisters, my wife their love give
me forces to perform this work.
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III
TABLE OF CONTENTS
Page
ACKNOWLEDGMENT………………………………………………………...... I
DEDICATION…………………………………………………………………….. II
TABLE OF CONTENTS…………………………………………………............. III
LIST OF FIGURES………………………………………………………............. V
LIST OF TABLES ………………………………….............................................. XI
ABSTRACT…………………………………………............................................. XII
CHAPTER 1: INTRODUCTION……………………………………………….. 1
CHAPTER 2: LITERATURE SURVEY……………………………………….. 3
2.1 Aluminum…………………………………………………………………… 3
2.1.1 History of Aluminum……………………………………………………. 3
2.1.2 Application……………………………………………………………… 4
2.1.3 Alloy Types………………………………………………………............ 4
2.2 Strength of Metals……………………………………………………………… 6
2.2.1 Dislocations……………............................................................................ 6
2.2.2 Slip………………………………………………………………………. 6
2.2.3 Particle coherency……………………………………………………….. 7
2.2.4 Solute solution hardening……………………………………………….. 8
2.2.5 Precipitation hardening …………………………………………………. 9
2.2.5.1 Precipitation hardening mechanism……………………………… 9
2.2.5.1.1 Cutting versus bowing…………………………………. 10
2.2.5.1.2 Shearing mechanisms of particle strengthening………... 11
2.2.5.1.2.1 Chemical hardening………………...................... 11
2.2.5.1.2.2 Stacking fault hardening……............................... 12
2.2.5.1.2.3 Modulus hardening……………………………... 12
2.2.5.1.2.4 Coherency hardening…………………………… 12
2.2.5.1.2.5 Order hardening………………………………… 13
2.2.5.1.2.6 Dispersion hardening…………………………… 13
2.2.5.1.3 Orowan bowing or bypass mechanism…………............. 14
2.2.5.2 Precipitation hardening in aluminum alloys……………………… 14
2.3 Heat treatment of Aluminum alloys……………………………………………. 17
2.3.1 Solute solubility………………………………………………………….. 19
2.3.2 The usual heat treatment procedure for aluminum……………………….. 19
2.3.2.1 Solution heat treatment (SHT)……………………………………. 20
2.3.2.2 Room temperature storage. (RT-storage)………………………… 21
2.3.2.3 Artificial aging (AA)……………………………………………... 21
2.4 The Al-Mg-Si (6xxx) alloy system…………………………………………. 21
2.4.1 Precipitation Hardening on Al-Mg-Si alloys………………………….. 22
2.4.1.1 Pseudo-binary Al-Mg2Si………………………………………… 22
2.4.1.2 Precipitation sequence………………………………………… 22
2.5 Factors Affecting the Precipitation Hardening in Al-Mg-Si alloys…………. 26
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IV
2.5.1 Solution heat treatment……………………………………………............. 26
2.5.2 Aging condition……………………………………………………………. 27
2.5.2.1 Time-Temperature variation………………………………............. 27
2.5.2.2 Two-step aging………………………………………..................... 27
2.5.3 Chemical compositions……………………………………………………. 29
CHAPTER 3: MATERIALS AND EXPRIMENTAL TECHNIQUE…………. 32
3.1 Materials……………………………………………………………………….. 32
3.2 Heat-treatment………………………………………………………………….. 32
3.3 Tensile Test…………………………………………………………………….. 34
3.4 Hardness test …………………………………………………………………… 36
3.5 XRD Analysis ……………………...................................................................... 37
3.6 Microstructure Examination……………………………………………………. 38
3.7 Fractographic Examination (SEM)……………………………………………... 38
3.8 Energy Dispersive X-rays Analysis (EDAX)………………………….............. 39
CHAPTER 4: RESULTS AND DISCUSSION………………………………….. 40
4.1 Effect of Artificial Aging on Tensile Properties……………………………….. 40
4.2 Factors Affecting the Artificial Aging…………………………………………. 52
4.2.1 Natural Aging……………………………………………………………. 52
4.2.1.1 The Influence of Natural Aging Duration on Mechanical
Properties……………………………………………….
52
4.2.1.2 Effect of natural aging time on artificial aging…………………… 59
4.2.2 Pre-aging………………………………….................................................. 67
4.2.2.1 Effect of pre-aging time on artificial peak aging condition………………… 67
4.3 Microstructure Examination and XRD Analysis ………………………………. 72
4.4 Scanning Electron Microscope (SEM) with Energy Dispersive X-rays
Analysis (EDAX)…………………………………………………………..
79
4.5 Fracture behavior………………………………………………...……………... 84
CHAPTER 5: CONCLUSIONS………………………………………………….. 87
REFERENCES……………………………………………………………............. 89
ARABIC SUMMARY ……………………………………………………............ أ
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V
LIST OF FIGURES
Page
Fig. 2.1 AA Designation of wrought Aluminum and its alloys.
5
Fig. 2.2 Illustrations of a line dislocation (a) and a screw dislocation (b). In
the case of the line dislocation, Burgers vector can be seen to lie in
the same plane as the plane 1 → 5, while it lies perpendicular to it
in the case of the screw dislocation.
7
Fig. 2.3 Figure (a) shows a fully coherent particle, figure (b) a coherent
particle, figure (c) a partially coherent particle and figure (d) a non-
coherent particle dispersed in the surrounding matrix.
8
Fig. 2.4 Figure (a) shows a schematic drawing of an atom dispersed in the
surrounding matrix which demands more space than the matrix
atoms. Figure (b) shows a schematic drawing of an atom which
requires less space than the surrounding matrix. Both can be seen to
cause coherency strain.
9
Fig. 2.5 A dislocation held up by a random array of slip-plane obstacles.
10
Fig. 2.6 A dislocation motion through strong and weak obstacles.
10
Fig. 2.7 Variation of yield strength with aging time for typically age-
hardening alloys with two different volume fractions of
precipitates.
11
Fig. 2.8 Schematic representation of the shape change accompanying the
movement of a dislocation through a GP zone.
12
Fig. 2.9 View of edge dislocation penetrating an ordered particle.
13
Fig. 2.10 Shown the precipitation sequence in Al-Mg-Si from the
supersaturated solid solution.
16
Fig. 2.11 GP zones in Al-Cu, Al-Zn and Al-Mg-Si.
17
Fig. 2.12 Coherency in a cubic lattice; [001] section of GP zone.
17
Fig. 2.13 The temper designation scheme of aluminum alloy.
18
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VI
Fig. 2.14 The phase diagram of silicon and aluminum. Theα phase to the left
is silicon fully dissolved in aluminum while the phase to the lower
right is a combination of the α-phase and solid silicon. The
horizontal line at 577oC is the solidus line. All phases above this
line except for the α-phase consists partly or fully of a liquid state.
19
Fig. 2.15 Schematic drawing of the heat treatment procedure. TRT, TAA and
TSHT denote room temperature (RT), temperature during artificial
aging (AA) and temperature during solution heat treatment (SHT)
respectively. The symbols tRT, tAA and tSHT denote the times for the
three steps. The vertical slopes in the temperature indicate assumed
instantaneous changes in temperature as the sample goes from one
treatment to another.
20
Fig. 2.16 Pseudo-binary diagram of Al-Mg2Si.
22
Fig. 2.17 Pictures of the β" precipitate taken with conventional TEM. (a)
shows the original picture, while (b) shows a filtered version. The
precipitate eyes can be seen as small rings, and denote the unit cell
centers.
24
Fig. 2.18 Picture of the β' precipitate taken with conventional TEM. The unit
cell can be observed to be hexagonal with lattice parameters a = b =
7.05o A.
25
Fig. 2.19 Picture of the B‟ precipitate taken with conventional TEM. The
precipitate eyes can be seen as hexagonal rings, and denote the unit
cell centers. The unit cell can be observed to be hexagonal with
lattice parameters a = b = 10.4˚ A.
25
Fig. 2.20 Al-Mg2Si-Two step aging.
28
Fig. 3.1 Heat-treatment furnace
33
Fig. 3.2 Heat-Treatment process
33
Fig. 3.3 Age hardening sequence of Aluminum alloys
34
Fig. 3.4 Tensile Test Specimen according to ASME E8
35
Fig. 3.5 Universal tensile testing machine
35
Fig. 3.6 Hardness Machine test
36
Fig. 3.7 XRD Machine 37
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VII
Fig. 3.8 Optical Microscope
38
Fig. 3.9 Scanning Electron Microscope
39
Fig. 4.1 Effect of artificial aging on tensile strength for Al-alloy 6351.
40
Fig. 4.2 Effect of artificial aging on 0.2% offset yield stress for Al-alloy
6351
41
Fig. 4.3 Effect of artificial aging on hardness for Al-alloy 6351
42
Fig. 4.4 Effect of artificial aging on total elongation for Al-alloy 6351
42
Fig. 4.5 True stress-true strain curve of the received Al-alloy 6351.
44
Fig. 4.6 True stress-true strain curve of solution treatment-water quenched
of Al-alloy 6351.
45
Fig. 4.7 True stress-true strain of artificially aging Al-alloy 6351 at 160oC
for 4 h.
46
Fig. 4.8 True stress-true strain of artificially aging Al-alloy 6351 at 160oC
for 18 h.
47
Fig. 4.9 True stress-true strain of artificially aging Al-alloy 6351 at 160oC
for 24 h.
48
Fig. 4.10 True stress-true strain curves of Al-alloy 6351 for solution treated-
water quenched, the received conditions in comparison with
various artificially aged conditions
49
Fig. 4.11 Change in 0.2%yield strength, MPa of Al-alloy 6351 due to the
effect of natural aging for various times.
54
Fig. 4.12 Change in ultimate tensile strength, MPa of Al-alloy 6351 due to
the effect of natural aging for various times.
54
Fig. 4.13 Change in hardness, HV of Al-alloy 6351 due to the effect of
natural aging for various times.
55
Fig. 4.14 Change in total elongation, % of Al-alloy 6351 due to the effect of
natural aging for various times.
55
Fig. 4.15 True stress-true strain of natural aging of Al-alloy 6351 at room
temperature for 170 h.
56
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VIII
Fig. 4.16 True stress-true strain of natural aging of Al-alloy 6351 at room
temperature for 1000 h.
57
Fig. 4.17 True stress-true strain curves of Al-alloy 6351 for naturally aged
condition in comparison with solution treated-water quenched and
peak-aging conditions.
58
Fig. 4.18 The effect of natural aging for 170 h followed by artificial aging at
160oC for various times on ultimate tensile strength, Mpa of Al-
alloy 6351.
60
Fig. 4.19 The effect of natural aging for 170 h followed by artificial aging at
160oC for various times on 0.2% offset yield stress, MPa of Al-
alloy 6351.
60
Fig. 4.20 The effect of natural aging for 1000 h followed by artificial aging at
160oC for various times on ultimate tensile strength, MPa of Al-
alloy 6351.
61
Fig. 4.21 The effect of natural aging for 1000 h followed by artificial aging at
160oC for various times on 0.2% offset yield stress, MPa of Al-
alloy 6351.
61
Fig. 4.22 The effect of natural aging for 170 h followed by artificial aging at
160oC for various times on hardness, HV of Al-alloy 6351.
62
Fig. 4.23 The effect of natural aging for 1000 h followed by artificial aging at
160oC for various times on hardness, HV of Al- alloy 6351.
62
Fig. 4.24 The effect of natural aging for 170 h followed by artificial aging at
160oC for various times on total elongation, % of Al-alloy 6351.
63
Fig. 4.25 The effect of natural aging for 1000 h followed by artificial aging at
160oC for various times on total elongation, % of Al-alloy 6351.
63
Fig. 4.26 True stress-true strain of natural aging of Al-alloy 6351 at room
temperature for 170 h followed by artificial aging for 8 h at 160oC.
65
Fig. 4.27 True stress-true strain of natural aging of Al-alloy 6351 at room
temperature for 1000 h followed by artificial aging for 4 h at 160oC
66
Fig. 4.28 Change in tensile properties difference of Al-alloy 6351 due to the
effect of pre-aging at 100oC on the artificial peak aging (160
oC for
18 h).
68
Fig. 4.29 Change in elongation difference of Al-alloy 6351 due to the effect
of pre-aging at 100oC on the artificial peak aging (160
oC for 18 h).
68
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Fig. 4.30 Change in hardness difference of Al-alloy 6351 due to the effect of
pre-aging at 100oC on the artificial peak aging (160
oC for 18 h).
69
Fig. 4.31 True stress-true strain of pre-aging of Al-alloy 6351 at 100oC for 10
min followed by artificial aging for 18 h at 160oC.
70
Fig. 4.32 True stress-true strain curves to illustrate the effect of natural aging
and pre-aging on artificial peak aging.
72
Fig. 4.33 Microstructure of the as received specimen at magnification
73
Fig. 4.34 Microstructure of the as quenched specimen (540oC for 45 min).
74
Fig. 4.35 Microstructure of 6351 Al-Mg-Si alloy treated at 540oC for 45 min
then artificially aged at 160oC for 4 h (under-aging condition).
74
Fig. 4.36 Microstructure of 6351 Al-Mg-Si alloy treated at 540oC for 45 min
then artificially aged at 160oC for 18 h (peak-aging condition).
75
Fig. 4.37 Microstructure of 6351 Al-Mg-Si alloy treated at 540oC for 45 min
then artificially aged at 160oC for 24 h (over-aging condition).
75
Fig. 4.38 Show XRD analysis of the as received specimen.
77
Fig. 4.39 Show XRD analysis of solution treatment water-quenched.
77
Fig. 4.40 Show XRD analysis of under-aging specimen.
78
Fig. 4.41 Show XRD analysis of peak-aged specimen.
78
Fig. 4.42 Show XRD analysis of over-aging specimen.
79
Fig. 4.43 SEM microstructure with EDAX of solution heat treated specimen
80
Fig. 4.44 SEM microstructure with EDAX of under-aged condition
81
Fig. 4.45 SEM microstructure with EDAX of under-aged condition
82
Fig. 4.46 SEM microstructure with EDAX of under-aged condition
83
Fig. 4.47 Fracture surface of solution treated-water quenched condition.
84
Fig. 4.48 Fracture surface of under-aged condition.
85
Fig. 4.49 Fracture surface of peak-aged condition.
85
Fig. 4.50 Fracture surface of over -aged condition.
86
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LIST OF TABLES
Page
Table 2.1 Strengthening methods for aluminum metal.
3
Table 2.2 AA Designation of cast aluminum and its alloys.
5
Table 2.3 Overview of the precipitate phases U1, U2 and B‟ (A, B and C).
26
Table 3.1 Chemical composition of Al-alloy 6351 used in the present work
32
Table 4.1 Strain hardening exponent and strengthening coefficient of
solution treated-water quenched alloy and artificially aged alloy.
43
Table 4.2 Strain hardening exponent and strengthening coefficient of
solution treated-water quenched alloy, artificially peak-aged alloy
and the effect of natural and pre-aging on artificial aging.
69
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Abstract
Modification of mechanical properties of Al-Mg-Si alloy 6351 by age hardening involve
studying the effect of aging temperature, time, natural aging and pre-aging on artificial aging
behavior in terms of mechanical properties (ultimate tensile strength, yield stress and
elongation), hardness and fractography examination.
Artificial aging after solution treatment-water quenched resulted in a sharp increase in both
ultimate tensile strength UTS and yield stress YS related to a sharp decrease in total
elongation with respect to solution treatment-water quenched only. As the time of aging
increase the strength of the investigated material increase slightly till reaches peak strength
after that it starts to decrease with increasing time of aging. As the aging temperature
decreases the precipitation of secondary solute rich phases takes place in the more number of
intermediate stages. The intermediate phases strain the matrix during their precipitation to
enhance the mechanical properties, so better mechanical properties are observed at lower
aging temperature.
Natural aging at room temperature (25 ±3oC) after solution treatment-water quenched resulted
in a slight increase in tensile properties with a slight drop in total elongation, natural aging for
170 hours and for 1000 hours after solution treatment followed by artificial aging of this alloy
at 160oC, shifted the time to reach peak strength to shorter aging time (8- 4 hours
respectively) in comparison to peak-aged condition (160oC for 18 hours).
Pre-aging at 100oC for various times after solution treatment then artificially aging at 160
oC
for 18 hours (peak-aged condition) was investigated. It was found that the pre-aging for 10
min followed by artificially peak aging at 160oC for 18 h led to slight increase in ultimate
tensile strength UTS with a higher increase in yield stress YS associated with a reasonable
total elongation.
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CHAPTER 1: INTRODUCTION
Al-Mg-Si Wrought alloys (6xxx series aluminum alloys) are generally used for structural
engineering applications in aerospace and automotive industries, and in civil engineering
owing to their strength to weight ratio, good formability, reasonable weldability, good
corrosion resistance, and lower cost. Al 6351 is identified for its light weight (ρ= 2.7g/cm3)
and good corrosion resistance to air, water, oils and many chemicals. Electrical and thermal
conductivity is 4 times greater than steels. Its chemical compositions are Si (0.93), Fe (0.36),
Cu (0.1), Mn (0.57), Mg (0.55), Zn (0.134), Ti (0.014) and remaining Al. It has higher
strength among the 6000 series alloys. Alloy 6351 is also known as a structural alloy, in plate
form and commonly used for machining. However relatively a new alloy the higher strength
of 6351 has replaced 6061 alloy in numerous applications. Mechanical properties can be
easily achieved at tension tests, with great precision. Thus, alloy such as 6351 have
considerably more silicon than magnesium or other elements, but find themselves in the form
Mg2Si series β phase. The AA 6351 aluminum alloy is used in manufacturing owing to its
strength, bearing capacity, reasonable workability and weldability. It is also used in
construction of boats, columns, chimney, rods, pipes, tubes, automobiles, bridges. Al
(6351 H30) series alloy can be also used in structural and general engineering objects such as
rail & road transport automobiles, bridges, cranes, roof trusses, rivets and so on with a good
surface finishing. Also it was observed from research that for the wrought aluminum alloy
AA6351-T6 show the lowest and most stable strain amplitude.
The main advantages of Al 6351 have some important performance characteristics that
make them very attractive for aircraft structures, namely light unit weight, simply one
third that of steel, strength compared to other aluminum alloys, good corrosion resistance,
with a negligible corrosion even in the presence of rain and other extreme conditions, high
toughness and resistance to low-ductility fracture at very low temperatures, and without any
ductile-to-brittle transition and excellent fabricability. These performance characteristics
make available advantages over conventional aircraft design, fabrication and creation of
aerospace structures like light weight and comparable strength enables the use of a higher
ratio of live load to dead load, superior corrosion resistance eliminates the need to paint the
aluminum components except may be for aesthetic purposes resulting in lower maintenance
costs, superior low-temperature toughness eliminates concerns about brittle fracture even in
the most severe freezing weather, ease of extrusion enables the design of more weight-
efficient beam and component cross sections, placing the metal where it is most needed
within a structural shape or assembly including providing for interior stiffeners and for joints
and the combination of light weight and ease of fabrication.
Si and Mg considered the main alloying element in 6xxx series, these elements are partially
dissolved in α-Al matrix and then present in the form of intermetallic phases depending on
composition and solidification condition. In the technical 6xxx aluminum alloys contents of
Si and Mg are in the range of 0.5-1.2wt%, usually with a Si/Mg ratio more than one. In
addition the intentional additions, transition metals like Fe and Mn are always present. If Si
content exceeds the amount that is required to form Mg2Si phase, the remaining Si is present
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in other phases, like AlFeSi and AlFeSiMn particles. A large number of wrought Al-Mg-Si
alloys contain an excess of Si, above that required to form the Mg2Si (β) phase, in order to
improve the age hardening response. In Al-Mg-Si addition of Mn is generally used to
decrease the grain size, restrain recrystallization and increase the strength as finely
precipitated intermetallics modifies the shape of plate-like iron phases which reduces their
embrittling effect. The combination of manganese with Fe, Si, and Al also formsα-
Alx(Fe,Mn)ySiz phase that acts as nucleation sites for Mg2Sicrystals, which eventually
influences the alloys behavior.
For these alloys, the accepted precipitation sequence starting from a supersaturated solid
solution is separate clusters of Si and Mg atoms, co-clusters containing Mg and Si atoms,
(spherical) GP zones, (needle-like) metastable β” phase, (rod-like) metastableβ‟ phase, Si
precipitates, and (platelets) of equilibrium βphase. The β” precipitates are considered the most
effective phase to give the main contribution to strength and hence they are mostly
responsible for the peak age hardening effect. The medium strength Al-Mg-Si aluminum
alloys are commonly processed by extrusion.
It is well known that heat treating variables in addition to the final aging time and temperature
can have a marked effect on the hardening response of heat-treatable aluminum alloys.
Variables are: delay time between the solution heat treating and aging concept of natural
aging, rate of heating to the aging temperature, and aging at an intermediate temperature prior
to final aging (pre-aging). Generally, natural aging and pre-aging treatments are beneficial;
they support fine, uniform precipitate dispersions and high strength. The situation appears to
be more complicated in the Al-Mg-Si system due to the fact that the precipitation reactions in
this alloys system are very sensitive the alloys compositions and the alloy history.
The objective of the current work is to study the influence of several heat treatments on the
mechanical properties of Al-alloy 6351; particular attentions were given to the
following points:
1- The effect of time and temperature variation on the artificial aging behavior of the alloy in
terms of hardness (HV), tensile properties and fractography.
2- The variation of time on natural aging behavior of the alloy in terms of hardness (HV),
tensile properties.
3- Natural aging before artificial aging has an important effect on the behavior on the alloy in
terms tensile properties.
4- The influence of pre-aging on the artificial aging behavior of the alloy in terms tensile
properties.
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CHAPTER 2: Literature survey
2.1 Aluminum and Its alloys
2.1.1 History of Aluminum
Aluminum (Al) is the third most common element in the earth‟s crust, but was not discovered
as an atomic element until the discovery of bauxite in 1821 in Les Baux, [1]. than to exist in
nature in its pure form it is found as aluminum oxide Al2O3 in different minerals with the
reddish stone Bauxite as the most common. It was first produced in its pure form in the late
1820‟s & remained an exclusive metal far more expensive than gold until the late 1800‟s. A
known story is that the Emperor of Germany, Napoleon III, one time invited to a banquet
where the emperor‟s relatives & the most honored guests where given the privilege of eating
from aluminum plates while the guests of lower ranks had to manage with gold. The age
when pure aluminum was a precious metal ended in 1886 with the discovery that pure
aluminum could be produced industrially from Al2O3 by electrolysis. Although the methods
from then are slightly changed, electrolysis still remains the principal process for producing
pure aluminum. Today, however they have the possibilities of producing far more waste
amounts of it.
Aluminum in its pure form is normally very soft and has none or few practical applications.
Adding small amounts of other elements to the liquid metal, in order to make an alloy where
its strength strongly increased. The principle alloying additions to aluminum are copper,
manganese, silicon, magnesium, and zinc; other alloying elements are also added in smaller
amount for grain refinement and to develop special properties. So there is a wide variety of
aluminum alloy. Nowadays the hardness of a typical aluminum alloy actually scales like ∼10 compared to the hardness of pure aluminum, and make it to one of the most common
materials utilized in daily life. In order to take advantage of its low density, aluminum has to
be strengthened by one or more of the following mechanisms. Table 2.1 showed four
completely different strengthening mechanisms that are used to strength aluminum alloys.
Table 2.1 Strengthening Methods for Aluminum Alloys
Mechanism Description Dislocation barrier
Strain
hardening
Plastic deformation, or work hardening, of metals
increases the dislocation density. Dense
dislocation 'tangles' can form, obstructing the
movement of other dislocation.
Other dislocation
Solute
hardening
Alloy elements such as Mg, Mn and Cu can 'pin'
dislocation, thereby strengthening the material. Solute atoms
Precipitation
hardening
Small, finally dispersed precipitates can
significantly increase the strength of aluminum
alloy.
Precipitates
Grain size
hardening
Reducing the grain size increases the alloy
strength according to the Hall-Petch relationship. Grain boundary
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2.1.2 Applications
Aluminum is what‟s called a lightweight metal with a density of 2700 kg/m3 in comparison
with steel which has a density of 7800 kg/m3 [2]. Although it doesn‟t have the same strength
as steel it has a higher strength-to-weight ratio which makes it appropriate for several
lightweight applications in i.e. Cars and airplanes. In addition to the high strength to weight
ratio aluminum in the form of Al-alloys has many other excellent properties, including high
electrical and thermal conductivity, high resistance to corrosion, and no ductile to brittle
transformation at low temperatures, easy shapeability and low energy amounts needed for
recycling. Only 5% of the energy required making it, Al-alloys are greatly used in different
articles such as packaging like in beverage cans [2].
However, despite of its benefits, Al-alloys possess weaknesses that confine their areas of
application. Their low fatigue limit, low hardness compared with steel and a melting point of
only ≈ 660oC make them unsuitable for several applications. For example certain parts of
automotive need to be strong to withstand high forces, and therefore need strength higher than
obtained by Al-alloys. Improving today‟s Al-alloys to be able to overcome some of the
mentioned weaknesses can be of excellent industrial importance. It allows Al-alloys to
substitute steels in a higher number of applications that means great environmental
advantages could be achieved.
Al-Mg-Si alloys are commonly used as medium strength structural alloys in many
applications, such as construction or automotive industry due to their favorable
formability, weldability, corrosion resistance and so on [3].
2.1.3 Alloy types
When dealing with alloys general one refers to all possible mixings of aluminum with
different elements. Since there are many different alloys and a system for classifying them is
needed. Aluminum alloys can most roughly be divided into the two groups wrought and
casting alloys, dependent on the way they are fabricated. According to the two groups, the
alloys have their own designation system that sorts them into different subcategories. They
are organized by using the category yxxx for wrought alloys and yxx.x for casting alloys.
Designed for wrought alloys y denotes the main group of alloying elements and the remaining
numbers xxx denote the modifications and amount of alloying elements. The identical applies for the casting alloys only that here the last digit stands for the product form.
In addition to the numbering system, all aluminum alloys also can be divided into to two
groups influenced by whether they are heat treatable or non-heat treatable. By heat treatable
one means that the alloy can be exposed to elevated temperatures for various times to alter
their particular atomic structure. Complete overviews of the different types of alloys found in
table 2.2 that illustrate the meaning of cast alloy and figure 2.1 that also illustrate the meaning
of wrought alloy.
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Fig. 2.1 AA Designation of wrought Aluminum and its alloys
Table 2.2: AA Designation of cast aluminum and its alloys
Definition of Casting Alloy Groups
Aluminum, 99.00% and greater 1xx.x
Aluminum alloys grouped by major alloying elements
Copper (Cu) 2xx.x
Silicon (Si), with added copper and/or magnesium 3xx.x
Silicon (Si) 4xx.x
Magnesium (Mg) 5xx.x
Zinc (Zn) 7xx.x
Tin (Sn) 8xx.x
Other elements 9xx.x
Unused series 6xx.x
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2.2 Strength of metals
Assume that you want to calculate the strength of a metal from an atomistic viewpoint; a
reasonable approach would be to combine the crystal structure of the metal with inter-atomic
bonding energies and then summarize to get an estimate of the bulk strength. The predicted
strength is between 103and 104 times higher than the actual strength of the metal [2]. How
come it so? How can the strength of the metal be so much smaller than the one calculated
from its atomic bonding? To understand this, it required to understand the concepts of
dislocations and slip.
2.2.1 Dislocations
A dislocation is taken as a line defect or imperfection in an otherwise ideal crystal.
Dislocations understood to be one-dimensional and really exist in two forms; line (edge)
dislocations and screw dislocations.
Line dislocations: A line (edge) dislocation exists when a crystallographic half-plane can be
introduced into or removed from the crystal structure, followed by re-bonding of the atoms
towards the termination interface on this plane. A schematic drawing of a line dislocation can
be shown in figure 2.2a where the lower a part of the central upper half plane is what defines
the dislocation. If you go into equal numbers of atomic distances in a very loop round this
dislocation, you will find yourself in an atomic position not the same as the one you started at.
The vector from the end point to the starting point is called „Burgers vector‟ and is denoted as
b. A line dislocation can be defined by this particular Burgers vector because it lies in the
same plane as the path of propagation throughout the dislocation [2]. A visualization of this
looping is seen in figure 2.2a. Starting in position 1 before traveling throughout the
dislocation by taking one step in every direction will lead you to position 5. To accomplish
the loop, you need to take one extra step to the right which defines the burgers vector.
Screw dislocation: A screw dislocation could be visualized by an ideal crystal that have been
sliced halfway though and then ‟screwed‟ to move the atomic bonding one crystal spacing.
Basically the screwing is really as shearing of each side of the cut in opposite directions. In
that case, the “burgers vector” is not in the plane of propagation as with the line (edge)
dislocation, but perpendicular to it [2]. These can be seen in figure 2.2b where this vector
from point 5 to point 1 lies perpendicular to the plane of propagation.
2.2.2 Slip
Dislocations will not stationary, but may undertake the process called slip. In case of line
(edge) dislocations, the process happens in „the direction of burgers vector‟ and it is in
„perpendicular direction to burgers vector‟ in case of screw dislocations. The direction of
motion is usually known that the slip direction, together with the slip-plane formed from the
dislocation itself and burgers vector, where the total process called slip system.
Slip can be easily visualized throughout the motion of a line dislocation. For the dislocation to
able to jump a single atomic spacing in the direction of burgers vector, only one particular
column of atomic bonds need to be broken at any one time. Following the breaking of the
bonds, the dislocation is transferred to the neighboring column wherever new bonds are
produced at the time rather than at the same time. It is usually this simple fact that explained
why metals are not as strong evidently from their own inter-atomic bonding energies.
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Fig.2.2. Illustrations of a line dislocation (a) and a screw dislocation (b). In the case of
the line dislocation, Burgers vector can be seen to lie in the same plane as the
plane 1 → 5, while it lies perpendicular to it in the case of the screw dislocation
[4].
The local stress has to exceed so-called Peierls–Nabaro stress τ given by the relation (2.1) [2];
for slip to happen,
τ = c · exp (−k d / b) (2.1)
Where k and c are constants for materials, d is the inter-planar distance between two
neighboring slip planes and b is the magnitude of burgers vector. The latter is important to be
aware when discussing interference with dislocation movements.
2.2.3 Particle coherency
To understand later sections regarding precipitation hardening, it is necessary to know the
concepts associated with coherency. Coherency could be understood by considering a particle
of one phase dispersed inside a matrix of another phase. Its fit with the host matrix might be
described through what is defined as coherency. The degree of coherency divided into four
groups, according to how well the dispersed phase fits in [4].
Fully coherent: The dispersed particle is considered to be fully coherent if it fits perfectly
with the host matrix in terms of crystal structure and lattice parameter. In other words, the
atoms within the particle fills already existing lattice points within the host matrix (figure
2.3a).
Coherent: The dispersed particle is said to be coherent if it fits perfectly into the host matrix
in addition to a small variation in lattice parameter. This difference in lattice parameter causes
a so-called coherency strain in the host matrix to induce the particle to fit in (figure 2.3b).
Partially coherent: The dispersed particle is considered to be partially coherent if it has
interfaces with different coherency. This can be seen in (figure 2.3c) wherever there is fully
coherency between the planes in the y-direction whereas there‟s coherency between the
planes in the x-direction.
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Incoherent: The dispersed particle is said to be incoherent if it does not fit with the host
matrix at all. The host matrix can thus be unstrained for the reason that crystal structure of the
dispersed phase is so different from the particular host lattice, that a coherency is
unobtainable even through coherency strain (figure 2.3d).
Fig.2.3 Figure (a) shows a fully coherent particle, figure (b) a coherent particle, figure
(c) a partially coherent particle and figure (d) a non-coherent particle dispersed
in the surrounding matrix [4].
2.2.4 Solute solution hardening
Hardening effects because of precipitation might not only be caused by Nano-sized
precipitates, but also by individual alloying elements being dissolved within the matrix. As
the alloying elements are of different chemical character compared to the matrix, they are
going to cause local expansion or contraction of the lattice, resulting in coherency strain [5].
The particular coherency strain effect is visualized in figure 2.4, showing two completely
different atoms dispersed in a host lattice.
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Fig.2.4 Figure (a) shows a schematic drawing of an atom dispersed in the surrounding
matrix which demands more space than the matrix atoms. Figure (b) shows a
schematic drawing of an atom which requires less space than the surrounding
matrix. Both can be seen to cause coherency strain.
2.2.5 Precipitation hardening
The strength of a metal could be increased through increasing its resistance against slip. In the
case of nonferrous metals as aluminum, this is done through the process called precipitation
hardening wherever a large amount of Nano-sized precipitates are introduced into the metal
that helps the metal stand up to dislocation motion. This interference process between these
precipitates and the dislocation motion could be described through different mechanisms,
coherency strain hardening, chemical hardening, stacking-fault hardening, order hardening,
modulus hardening and dispersion hardening [6].
2.2.5.1 Precipitation hardening mechanisms
Most alloys rely on precipitation hardening in one form or another to accomplish high
strengths and the central concept is that the strength of a ductile material is governed by
dislocation flow past obstacles. To understand the relationship between microstructure and
strength, we need to get into the subject of hardening mechanisms. Therefor strength can be
designed by controlling the density and the nature of the obstacles to dislocation motion.
When a glide dislocation incurs one of numerous obstacles as shown in Fig 2.5 it must be bent
to some angle υc (0 ≤ υc ≥ π) before it can move on where angle υc is measure of the strength
of the obstacles [7]: Weak obstacles can be overcome with very slight bending (υc ≈ π) while
strong obstacles cannot be overcome unless the dislocation practically double on itself (υc ≈
0) as shown in Fig 2.6.
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Fig. 2.5 a dislocation held up by a random array of slip-plane obstacles [7].
The following equation is given:
(2.2)
The given equation expresses the shear stress that required to beak the obstacles when the
dislocation is held in equilibrium where G is shear modulus, b is burger‟s vector and L' is
Mean intercept length of precipitates. At a critical stress the dislocation breaks the obstacles
and advances to other obstacle depending on the size of the obstacles and interaction between
dislocation and obstacles (critical break angle ).
Fig.2.6 A dislocation motion through strong and weak obstacles [7].
2.2.5.1.1Cutting versus bowing
Second phase particles act within two distinct ways to retard the dislocation motion, the
particle either might be cut by the dislocations or the particles resist cutting and the
dislocations are forced to bypass them [8]. At small sizes or soft particles the dislocation cut
or deforms through the particles, there are six properties of particles which affect the ease
with which they are often sheared, they called strengthening mechanisms. The summation of
these mechanisms leads to an increase in strength with increasing the particle size till reaches
a point where the cutting of the particle becomes very hard, and instead the dislocations find
ways of moving around the particles [8]. When the particles become very strong or coarse it
does not break even at ≈ 0, then the dislocations reach an unstable (Frank-Read) configuration and slip occurs by dislocation multiplication, leaving a small dislocation loop
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(Orowan loop) around the unbreakable particle. The stress to accomplish this obtained from
equation (2.2) by putting ≈ 0 which called „Orowan bowing stress‟ [7]. Large particles mean fewer particles, large particles interspacing and lower flow stresses are obtained, as
shown in fig 2.7[9].
Fig. 2.7 Variation of yield strength with aging time for typically age-hardening alloys
with two different volume fractions of precipitate [9].
2.2.5.1.2 by shearing mechanisms of particle strengthening
To obtain and estimate the strengthening in the case of particle that are cut through by a glide
dislocation, there are a number of possible source for this shear strengthening. They are as
follow:
2.2.5.1.2.1Chemical hardening
The hardening caused by the stress required to force a dislocation through the precipitate itself
referred to as cutting. If the precipitate is coherent with the matrix, the dislocation could move
by the same slip mechanism as in the matrix. However, as the dislocation moves though the
precipitate, the precipitate will for the case of a line dislocation, increasing in size due to the
introduction of the extra-half plane, as the precipitate is inhomogeneous in comparison to the
rest of the matrix. Both these events will as well as additional effects result in a hardening due
to the extra energy required to inflict them [6].
Cutting through a particle with a dislocation displaces one half relative to the other by b
(burger‟s vector), as shown in Fig. 2.8.
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Fig. 2.8 Schematic representation of the shape change accompanying the movement of a
dislocation through a GP zone [10].
2.2.5.1.2.2 Stacking fault hardening
For precipitates that have stacking-fault energies significantly different from the matrix, the
interaction between the dislocation and the particles can be dominated by the local variation
of fault width when glide dislocations enter the particles. A large difference in stacking fault
energy between particle and matrix, i.e. Ag in Al, increases flow stress because of the
different separation of partial dislocations in the two phases [8]. In order to operate this
mechanism, the particle must have a structure which gives ride to extended dislocations.
2.2.5.1.2.3 Modulus Hardening
A large difference in elastic modulus results in image forces when a dislocation in the matrix
approaches a particle. Considering, i.e. the difference between silver, Ag particles (nearly the
same shear modulus) and iron, Fe particles (much higher shear modulus) in aluminum. Think
of modulus hardening as being caused by a temporary increase in dislocation line energy
whereas it resides among a particle [10].
2.2.5.1.2.4 Coherency hardening
Coherency strain hardening is a hardening mechanism that results from the coherency strain
fields produced by precipitates within the matrix. The strain fields are generally produced as
the precipitates are not fully coherent with the matrix, but obtain coherency through bending
and stretching of the surrounding matrix as shown in figure 2.3b. The hardness is obtained
though the altering of crystallographic structure such that the Peierls -Nabaro stress (2.1)
increases as the dislocation moves closer to the precipitate. This causes the precipitate to be
able to repulse the dislocation. The latter has consequences as the precipitates could also aid
dislocation motion by repulsing them in their motion direction. If maximum strength is to be
required, the density of precipitates must therefore not be too high [6].
Differences in density between particle and the matrix give rise to elastic stresses near the
particle. This has been analyzed based on the elastic stresses that exist in the matrix adjacent
to a particle that a different lattice parameter than the matrix. This mechanism can be applied
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to the early stages of precipitation, i.e. strengthening by „GP zones‟ and very fine secondary
phases [10].
2.2.5.1.2.5 Order hardening
The hardening due to ordering depends on the product of the anti-phase-boundary energy
(APBE) and the area swept by a dislocation in a particle. Passage of a dislocation through an
ordered particle, i.e. Ni3Al in super-alloys, results in a disordered lattice and the creation of
anti-phase boundaries. Generally, low values of the APBE not only predict slight increase in
hardness, but also the result which the dislocations can move through the particles
independently of one another.
This may be understood from Fig. 2.9, in which the particular crystal structure is cubic and
has composition AB.
In (a) the dislocation has not yet entered the particle, in (b) it is partially entered through the
particle and the slip result in the formation of an anti-phase boundary (A-A and B-B bonds)
across the slip plane. After the dislocation exited the particle, the ant-phase boundary
occupies the whole of the slip plane area of the particles. This mechanism is more important
for Ni-based super alloys [10].
Fig. 2.9 View of edge dislocation penetrating an ordered particle [10].
2.2.5.1.2.6 Dispersion hardening
Hardening obtained from larger incoherent precipitates called dispersoids. If the dispersoids
are totally incoherent with the matrix, the dislocation may no longer pass through them
through cutting as with coherent precipitates, but have to find alternative mechanisms to pass,
the hardness is thus obtained by the stress required for the dislocation to pass the dispersoid
by any alternative mechanisms [6].
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2.2.5.1.3 Orowan bowing or bypass mechanism
Increasing aging times or aging temperatures, precipitates come to be incoherent and
dislocations are no longer able to cut through them. Rather, they must by-pass these
precipitates by one of a number of possible mechanisms. These mechanisms include bowing,
climb and cross-slip. One of the important features of dispersion hardened materials is the
homogenous nature of slip. This feature has important consequences in terms of mechanical
properties; the process of particle by-passing is called “Orowan bowing mechanism”. The
Orowan shear stress require to bowing a dislocation between two precipitate particles is
directly proportional to burger‟s vector and inversely proportional to the particle separation L'
as given by:
τ = Gb/L'(2.3)
The generation of dislocation loops around the particles results as a result of the Orowan
bowing mechanism. As subsequent dislocations pass, dense tangles involving dislocation
form resulting to a high rate of work hardening [10]. Most theories of strengthening with
second-phase particles derive from idealized spherical particles. However, particle shape
could be important, at equal volume fraction, rods and plates strengthen about twice as much
as spherical particles [8].
2.2.5.2 Precipitation Hardening in Aluminum Alloys
The most important methods for strengthening alloys, specifically nonferrous alloys, utilizes
the solid state reactions referred to (precipitation or age hardening).
The history of precipitation hardening of aluminum alloys goes back to 1906 when A. Wilm
[11] discovered that quenched from a high temperature nearly ~ 550°C in a cold water, Al-
Cu-Mg alloy initially increased in hardness as it was spent at room temperature; the alloy
hardened with age, which led to the phenomenon being known as “age hardening”, Wilm
examined his samples within an optical microscope, but not able to detect any structural
change as the hardness increased. At 1919 Mercia, Waltenberg and Scott [12] supposed that
in their study of an Al-Cu alloy, they also observed that the hardness increase after quenching.
They provided that the solid solubility of copper in aluminum decreases with decreasing
temperature and this led them to propose that the hardening with age after quenching was
caused by copper atoms precipitating out as particles from supersaturated solid solution
(SSSS).
In a review paper published in 1932, Mercia [13] recommended that “age hardening in Al-Cu
alloys resulted from the assembly of copper atoms into a random array of small clusters
“knots” which interfere with slip when grains are generally deformed”. In 1938 Mercia‟s
“knots” was provided by the historic work of Guinier [14] and Preston [15] who,
independently, interpreted features in diffuse x-ray scattering from aged aluminum alloys as
evidence for clustering of atoms into very small zones; since classified as Guinier-Preston
zones, or GP Zones. Direct observation of the precipitated GP zones did not occur until the
transmission electron microscopy (TEM) was developed. For the first time, the transmission
electron microscope provided an investigation technique with enough resolution to reveal the
very small precipitate particles (GP zones) responsible for age hardening.
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Aluminum alloys may be hardened (or strengthened) by heat treatment is complete solute
solid solubility at high temperature but only very limited solute solid solubility at room
temperature. The required heat treatment to increase strength of aluminum alloy is explained
in three steps process:
First, Solution heat treatment: dissolution of soluble phases,
Followed by quenching: development of supersaturated solid solution and
Finally, age hardening: precipitation of solute atoms either at room temperature (natural
aging) or at elevated temperature (artificial aging).
Fig. 2.10 shows the precipitation sequence in Al-Mg-Si from the supersaturated solid solution
as example in Al-alloys.
It was found that the rate and the degree of hardening increase if an alloy is aged at an
elevated temperature, say up to 200°C; this was termed artificial aging as distinct from aging
at room temperature. For some alloys (for example, Al-Mg2Si) there may be important
differences in detail between the metallurgical processes that occur at different temperatures
and times, significantly within the sequence of phase transformations that present the
precipitation sequence; that is, the manner in which solute clusters (zones) grow and change
in shape and crystal structure [13, 14].
Strengthening by age hardening involves the formation of coherent clusters of solute atoms,
that is, the solutes atoms have collected into a cluster still have the same crystal structure as
the solvent phase. This causes a lot of strain because of a mismatch in size between the
solvent and solute atoms. The cluster stabilizes, because dislocation has a tendency to reduce
the strain. The alloy is said to be strengthened and hardening when dislocations are sheared by
the coherent solute clusters. Consequently, higher strength by obstructing and retarding the
movement of dislocations may be because of the presence of the precipitate particles, and
more importantly the strain fields in the matrix were surrounding the coherent particles.
However, a dislocation can circumvent the particles only by bowing into a roughly
semicircular shape between them under the action of the applied shear stress if the precipitates
are semi-coherent, incoherent or incapable of reducing strain behavior because they are too
strong,. The characteristic that determines whether a precipitate phase is coherent or non-
coherent, is that the closeness of match between the atomic spacing on the lattice of the
matrix and on that of the precipitate [17, 9].
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Fig.2.10 shown the precipitation sequence in Al-Mg-Si from the supersaturated solid
solution
For understanding of how GP zones harden aluminum alloy is the fact that the GP zones
consist of clusters of solute atoms that are said to be coherent with the aluminum lattice. For
Al-Cu, as showed in Fig. 2.11 a, the copper atoms assemble in singles atoms layers on (100)
plane, which creates a distortion, in this case a contraction, of the lattice (remember, Cu atoms
are smaller in comparison with Al atoms). Nonetheless, continuity of the crystallographic
planes is maintained; the platelets of copper are fully coherent with the aluminum lattice. GP
zones as in Al-Zn are also fully coherent, see Fig.2.11 b. Here, the zones are approximately
spherical in shape and, because Zn atoms are slightly smaller in comparison with Al atoms,
the distortion is again a contraction of the lattice. However, the zones are fully coherent again.
In Al-Mg-Si, GP zones are only semi-coherent, Fig.2.11 and 2.12. The needle-shaped (or rod-
shaped) zones are coherent with the matrix along their length, which can along an aluminum
matrix direction. Detailed Electron Microscopy with a Transmission Electron
Microscope [14] has shown that, these zones have a hexagonal structure [18] with the close-
packed planes parallel to the cube planes of the aluminum matrix and coherent with it. There
is considerable mismatch in crystal structures perpendicular to the major axis of the needle-
shaped zone, associated with the cylindrical interface between the needle and the surrounding
matrix where the matrix within the neighborhoods of the cylindrical interface expands to
accommodate the mismatch.
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Fig. 2.11 GP zones in Al-Cu, Al-Zn and Al-Mg-Si [16].
Fig.2.12 Coherency in a cubic lattice; [001] section of GP zone in Al-Mg-Si [18].
2.3 Heat Treatment of Aluminum Alloys
In fact, the properties of aluminum alloy are not given entirely by the atomic composition of
the alloys. This has already been mentioned by the fact that the two major types of aluminum
alloys are defined by the way they are fabricated. In order to give the aluminum alloys a
desire set of mechanical properties, the alloys undergo different treatments to reshape their
atomic. The different possible treatments will be summarized in five major groups denoted by
the symbols F, O, H, Wand T wherever the temper designation scheme is shown in Fig.2.13.
The five major treatments had the meaning of as-fabricated, annealed, cold-worked, solution-
treated and age-hardened, respectively. Solution treatment may in some cases be included as a
part of the age-hardening, and a common term used in this case to include both is “heat-
treating”.
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A heat treatment is thereby a treatment wherever the alloy is kept at different temperatures for
various times. The hardness enhance that age-hardenable alloys obtain during heat treatment
was in the late nineteen hundreds discovered to be caused by Nano size particles known as
“precipitates”. There are different precipitates with different morphologies, but they can
commonly be interpreted as particles that jam the matrix in such a way that slip becomes
more difficult. Slip was described previously as the movement of a dislocation, and imped the
dislocation motion will make the alloy very harden. The types of precipitates that are created
depend on the temperatures utilized in the heat treatment and the corresponding storage times,
and they can be represented in a temperature-wise succession known as the precipitation
sequence. In such a sequence, the precipitates formed at the beginning of the process at the
lowest temperatures for shortest times and subsequently formed at the highest temperatures at
the end.
Fig. 2.13 the temper designation scheme of aluminum alloy.
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2.3.1 Solute solubility
In order to understand the reason for performing a heat treatment, first we should know the
concept of solute solubility. There are limited amounts of alloying elements that can be added
and dissolved before the solution splits into two separate phases. Figure 2.14illustrate the
phase diagram of aluminum and silicon wherever the α-phase denotes fully dissolved silicon
in aluminum and also can be observed that the amount of silicon that may be dissolved in
aluminum before pure silicon starts to split is strongly temperature dependent. Investigation
of the phase diagram, it noticed that the maximum solid solubility of silicon in aluminum
about at 2% is found at 577oC. As shown in figure 2.14, the solubility of Silicon in Aluminum
varies with temperature. If 2 % of Si is completely dissolved in the host Al-matrix at 577oC, a
lowering of the temperature will result in a phase separation. Provided that, lowering of
temperature quickly, a supersaturated solid solution would be the result where SSSS is an
unstable/metastable phase and the driving force for aggregation of Si atoms is very large.
Fig.2.14 The phase diagram of Magnesium silicide and aluminum. The α phase to the
left is silicon fully dissolved in aluminum while the phase to the lower right is a
combination of the α-phase and solid silicon. The horizontal line at 595oC is the
solidus line. All phases above this line except for the α-phase consists partly or
fully of a liquid state.
2.3.2 The usual heat treatment procedure for aluminum
For producing desired properties of aluminum alloys, a heat treatment could be performed on
them to alter their atomic structure. It carried out by kept alloys at different temperatures for
various times, and take care that the transition time from one temperature to another is as
short as possible. The traditionally heat treatment is divided into three parts, namely solution
heat treatment (SHT, room temperature storage (RT-storage) and artificial aging (AA). A
schematic diagram for explaining this procedure can be shown in figure 2.15. Different heat
treatments are usually referred to by the abbreviation TX, where X is often a number and T
denotes that the alloy is susceptible to age hardening.
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2.3.2.1 Solution heat treatment (SHT)
When an alloy is solution heat treated, it is heated to a high temperature (500∼577oC for aluminum) where it is hold for a time tSHT which this time can vary from 30 minutes to several
hours. The temperature needs be chosen such that dissolve all solute elements, but without
any transition to liquid state (below solvus line). The purposes of solution heat treatment are:
1. In order to dissolve all phases consisting of solute elements in the aluminum matrix so that
the solute elements are homogeneously spread out where this is a good starting point for
constructing new phases.
2. To introduce vacancies within the Al matrix. The density Cv of vacancies present in a metal
will increase exponentially with the temperature, and the vacancy concentration is explicitly
given by [19]:
(2.4)
Where Ef is the energy required introducing one vacancy into the system, kB is Boltzmann‟s
constant and T is the absolute temperature in Kelvin. The diffusion of substitutional solute
atoms is dependent on vacancies, and vacancy diffusion is many orders of magnitude larger
than the so-called “self-diffusion” [20]. The process is actually for that reason required to
form clusters and led to growth of precipitates.
In order to obtain a super saturated solid solution, after solution treatment the alloy is quickly
cooled to room temperature, the process known as quenching. In this case the state of the
system is then no longer stable, and it will undergo phase separation to lower its energy to
achieve the stability. After quenching, the treatment enters the next step (Phase) which is
called room temperature storage.
Fig.2.15 Schematic drawing of the heat treatment procedure. TRT, TAA and TSHT denote
room temperature (RT), temperature during artificial aging (AA) and
temperature during solution heat treatment (SHT) respectively. The symbols
tRT, tAA and tSHT denote the times for the three steps. The vertical slopes in the
temperature indicate assumed instantaneous changes in temperature as the
sample goes from one treatment to another.
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2.3.2.2 Room temperature storage. (RT-storage)
The storage of the alloy at room temperature, the diffusion processes of solute atoms often
have enough energy to proceed, and then aggregate either favorably or not. The solutes spread
along the matrix forming phases, and the time of storage affects greatly on this process. In
principle, the RT-storage step could go on till equilibrium is reached, but diffusion at this
temperature is too slow process and it would take an infinitely long time [21].
2.3.2.3 Artificial aging (AA)
In this the treatment, the storage at elevated temperatures may create large precipitate
particles. The temperatures TAA and time tAA for this process is depending on which
precipitate phases are desired. AA treatment for Al-alloys is typically performed at
temperatures in the range 160-260oC, but the exact temperatures and times are dependent on
the alloy composition and solute atoms content. Once the desired precipitates are obtained,
the alloy is quenched and then ready for use.
2.4 The Al-Mg-Si (6xxx) alloy system
Al-Mg-Si alloys (6xxx) alloys are considered the most commercially used Al alloys these
days. they can be used in everything from the transport industry to the consumer industry, due
to their good corrosion, welding properties, high strength to weight ratio and low cost.
Particularly they are used as automobile body sheets, and before they are used, the car body
sheets treated by process namely paint-baked cycle at 180oC, is a temperature at which the
peak hardness of these particular alloys [21].
6xxx series alloys contain silicon and magnesium approximately in the proportions
required for formation of (Mg2Si) compound magnesium silicide, making them heat
treatable. Although not as strong as 2xxx and 7xxx alloys, 6xxx series alloys behave good
formability, weldability, machinability, corrosion resistance, and medium strength. Alloys in
this heat-treatable group could possibly be formed in the T4 temper (solution heat
treated but not precipitation heat treated) in addition to strengthened after forming to
full T6 properties by precipitation hardening heat treatment.
Al-Mg2Si alloys can be divided into three groups. The first group, the total amount of
magnesium and silicon does not exceed 1.5%; the elements are in a nearly balanced ratio;
typical alloy of this group is 6063 alloy. This alloy is widely used for extruded architectural
sections. It nominally contains 1.1% Mg2Si. The second group nominally contains 1.5% or
more of magnesium, silicon and other addition elements such as .3% Cu, which increase
strength in the T6 temper. Elements such as manganese, chromium, and zirconium are used
for controlling grain structure. Alloys of this group such as 6061 alloy achieve strength higher
than in the first group in the T6 temper by about 70 MPa. The third group contain an amount
of Mg2Si overlapping the first two but with excess silicon. An excess of .2% Si increase the
strength of alloy containing .8% Mg2Si by about 70 MPa (10 KSi). Increasing the amounts of
excess silicon is less beneficial. Excess magnesium, however, is of beneficial only at law
Mg2Si contents because magnesium lower the solubility of Mg2Si. In excess silicon alloys,
segregation of silicon to grain boundaries causes grain-boundaries fracture in recrystallized
structures. Additions of manganese, chromium or zirconium counteract the effect of silicon by
preventing recrystallization during heat treatment. Addition of lead and bismuth to an alloy of
this group improve machinability. Common alloys of this group are 6009, 6010, and 6351
alloys [9].
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2.4.1 Precipitation Hardening in Al-Mg-Si alloys
2.4.1.1 Pseudo-binary Al-Mg2Si
Al-Mg-Si alloy is a ternary system. Engineering Al-Mg-Si alloys are based on the pseudo-
binary composition Al-Mg2Si %. Fig.2.16. the equilibrium precipitate in the Al-Mg-Si is
Mg2Si which known as a balanced compositions contain magnesium and silicon in same
atomic ratio of 2:1 as the equilibrium precipitate. In terms of Wt%, this translates to the
ratio1.73:1.
Fig. 2.16 pseudo-binary diagram of Al-Mg2Si
2.4.1.2 Phase co-exist and precipitation sequence
For a balanced alloy, the precipitation sequence is specifically as follow:
Embryo clusters →needle-shaped GP zones β” →intermediate β‟→ β (Mg2Si)
The expression of “embryo cluster” is introduced into this sequence. The recent work in this
field by Murayama et al [22] who studied the pre-precipitation stages of Al-0.70Mg-0.33Si
and Al-0.65Mg-0.70Si alloys by using Atom Probe Field Ion Microscopy (APFIM) and High
Transmission Electron Microscopy (HTEM) claim to have detected the separation of Mg and
Si clusters atoms. They were incapable of detect either separate clusters or co-clusters in a
High Resolution Transmission Electron Microscope. The smallest clusters that can be
detected within the TEM are needle like-shaped zones that grow in length and rather more
slowly in diameter, with increasing aging time. They proposed the following precipitation
sequence:
Separate Mg and Si clusters →co-clusters of Mg and Si →small equiaxed precipitation → β''
precipitates → β' precipitates → β (Mg2Si)
The effect of aging treatment on mechanical properties and precipitation behavior in Al-Mg-
Si alloy (0.95%Mg, 1.55%Si and 0.1%Zr) were studied by Kang et al [23]. The results
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indicate that the precipitation sequence of Al-Mg-Si alloy with excess Si content is proposed
to be:
SSSS → independent clusters of Si and Mg atoms, co-clusters of Si and Mg atoms → GP
zones → Si rich phase → β'' phase → β' phase → Si precipitates→ β (Mg2Si)
Studies carried out on Al-alloy 6082 confirm that the precipitation sequence that in generally
accepted is the following:
SSSS → atomic → clusters → GP zones → β'' → β' → β.
Some authors of these studies consider GP zones as GP-1 zones while β'' particles are referred
to GP-2 zones. It has been shown that Mg atoms from clusters in the as-quenched stage and
eventually from co-clusters with Si. The atomic ratio of Mg: Si atoms in the Mg-Si co-clusters
are chosen to be 1: 1. The equiaxed zones observed by artificial aging for 3 h at 175 have a
higher Mg: Si ratio of 1.6: 1. Increasing artificial aging suggests that the atom ratio of Mg: Si
approaches the equilibrium value of 2: 1 [24].
Other studies showed that the hardness obtained for age-hardenable alloys after heat treatment
is caused by the strain-field surroundings of Nano-sized particles known as precipitates and
the precipitation sequence for 6xxx alloys studied has been reported as follow:
SSSS → AC → GP zones → β''→ β', U1, U, B' → β/Si
Where SSSS referred to super saturated solid solution, AC is atomic clusters and GP zones
standing for Guinier-Preston zones. The other symbols denote the respective precipitate
phases; with the uttermost right phase β (Mg2Si) that called the equilibrium phase. Phases on
the right of the sequence are larger phases which they are produced at higher temperatures
and longer times than those to the left.
a) Atomic clusters
Each two solute atoms, which distribute homogeneously, start to cluster with each other to
form precipitates. A sophisticated technique like Atom Probe Tomography (APT) is used to
observe this precipitates in order to prove the presence of clusters. The solute clusters in the
precipitation sequence begins from the step where two solute atoms are next to each other and
still progress until the cluster begins to grow large. The coherency between the clusters and
the Al matrix deteriorate the contrast, which makes it difficult to be observed by TEM [25].
b) GP-zones
The GP-Zone is formed due to the continuous growth of clusters because of the random
distribution of solutes. The pre-β” precipitate is the predominant evolved phase among several
differently evolved phases from GP-Zones [21]. Coherency effects of GP-Zones make it
possible to investigate with HRTEM because of its large size compared with clusters.
Marioara et al [26] discovered that needle-like GP-zones in the 6082 Al alloy were less
coherent with the matrix than β”. Three dimensional atom probes (3DAP) studies by
Murayama and Hono [25] have shown that GP-zones in the same alloy system have equal
amount of both Mg and Si approximately 1. The GP-Zone usually defines a small particle
with little coherency with the matrix.
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c) The β" precipitate
The β” precipitate or some author‟s called it the GP-II zone which considered the main
hardening phase in 6xxx-alloys [27]. This phase can be created when the alloy artificially
aged at temperature in between 125oC and220
oC [21] as the temperature increase i.e. 250
oC
and more the β"-phase will start to dissolve and or transform [29]. For a long time the
composition of the β" phase was believed to be Mg2Siafter the composition of the equilibrium
phase β. In 1996 Edwards et. al. [30] showed that the Mg/Si ratio was closer to 1 using the
APT investigations while Andersen et. al. [28] in 1997 found that the composition of β" phase
to be Mg5Si6. Finally, the most likely composition of β" phase that was founded by Hasting et
al [31] using APT and DFT techniques is Mg5Al2Si4 which have Mg-rich, and not Si-rich
according to Andersen et al suggestion.
The β” precipitate has needle shape morphology, fully coherent with the Al-matrix along the
b-axis and semi-coherent along the two other axes and is elongated along the direction
of the aluminum lattice with size nearly ∼ (4x4x50 nm) [28]. The β" precipitate has monoclinic crystal structure with a = 1.516 nm, b =0.405 nm, c = 0.674 and β = 105.3
o as
shown in figure 2.17, and it is ordered relative to the host aluminum lattice in such a way that
(001)Al|| (010)Β", [310]Al||[001] Β" and [230]Al||[100]β". the angle between the β" a-vector and
[010]Al is 33.69oand therefore the angle between the β" c-vector and [100]Al is 18.43
o [28].
d) The β' precipitate
Increasing the aging time or aging temperature, β" phases will start to dissolve or transform
and a new phase will create known as β' [29]. Which is bigger than β" phases and have
dimensions nearly∼ (10x10x500 nm) in compared to∼4x4x50 nm for β'' precipitate. It has a hexagonal unit cell with a = 0.705 nm and c = 0.405 nm, and the latter coinciding with the
4.05 ˚A lattice parameter of fcc aluminum making it fully coherency with the Al. Fig.
2.18 show the hexagonal unit cell of the β' precipitate. The unit cell of β' doesn‟t have a
required orientation in the aluminum (001) plane and may be observed with many different
orientations unlike β" [32].
Fig.2.17 Pictures of the β" precipitate taken with conventional TEM. (a) shows the
original picture, while (b) shows a filtered version. The precipitate eyes can be
seen as small rings, and denote the unit cell centers [28].
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Fig.2.18 Picture of the β' precipitate taken with conventional TEM. The unit cell can be
observed to be hexagonal with lattice parameters a = b = 7.05o A [32].
e) The B', U1 and U2 precipitates
The B‟, U1 and U2 precipitates or also known A, B and C which are coexist with β‟. U1 are
Si-rich and belongs to space group P3m1 which have a hexagonal rod-shaped, semi-coherent
phase which is often found on dislocations, while U2 have orthorhombic with space group
Pnma [33]. Table 2.3 gives more information about their crystal structure and Fig.2.19 shows a
conventional TEM-picture of the B‟-phase.
Fig.2.19 Picture of the B’ precipitate taken with conventional TEM. The precipitate eyes
can be seen as hexagonal rings, and denote the unit cell centers. The unit
cell can be observed to be hexagonal with lattice parameters a= b= 10.4 ˚A
[32].
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Table 2.3 Overview of the precipitate phases U1, U2 and B’ (A, B and C) [29].
f) The equilibrium phase β
If the heat treatment of a 6xxx-alloy are conducted at high temperature for long times, all
solute within the precipitate phases will finally promote in the formation of the equilibrium
phase β. The crystal structure ofβ phase is fcc type like Ca2F with a lattice parameter equals
0.639and its stoichiometric composition is Mg2Si [34]. It was believed that all the hardening
phases had the same composition (Mg2Si) and this belief is changed by Andersen et al in the
late of the last century [28]. The β phase is very large with dimensions ∼µm and predominant
in influence compared to the other precipitate phases in 6xxx an alloy.
2.5 Factors Affecting the Precipitation Hardening in Al-Mg-Si alloys
2.5.1 Solution Heat Treatment Solution Heat Treatment includes heating the alloy to a temperature which below the solvus
line of the alloy in order to avoid partial melting. In case of Al-Mg-Si alloy the temperature
ranged from 500 to 577o
C for enough time till all solute atoms are dissolved followed by
rapid cooling (water-quenched) to obtain a super saturation solid solution (SSSS). Prolong
heat treatment will cause a migration of Mg atoms to t