modification of mechanical properties of 6351 al-mg-si...
TRANSCRIPT
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
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
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
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
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.
II
Dedication
I dedicate this thesis to my parents, my brother and sisters, my wife their love give
me forces to perform this work.
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
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
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
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
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
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
IX
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
X
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
XI
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.
1
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
2
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.
3
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
4
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.
5
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
6
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.
7
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.
8
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.
9
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.
11
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
11
(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.
12
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
13
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].
14
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.
15
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].
16
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 <100> 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.
17
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”.
18
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.
19
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.
21
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.
21
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].
22
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
23
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.
24
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 <100>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 = 105.3o 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 <001>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].
25
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].
26
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 the surface [35].
Increasing solutionizing temperature increase the strength of the alloy where the best range in
between 540-550o
C was founded by Dorward et al [36]. Mechanical properties are more
sensitive to quenching rate, alloys have low Mg2Si content showed low quenching rate
sensitivity but the sensitivity increases as Mg, Si, Mn, Cu and Cr contents are increased [35].
More studies were done on Al-alloy 6082 using Small-Angle Neutron Scattering (SANS)
technique in order to show the effect of quenching rate on the (Mg/Si) precipitates final size
after aging. It was found that low quenching rate leads to create (Mg/Si)precipitates with a
course size ∼30-200nm and precipitates size are quenching rate dependent while high
quenching rates leads to(Mg/Si) precipitates with a fine size∼ 2-30nm and precipitates size
depends on the aging condition [37]. Loss of the quenched in-vacancies and the grain
boundary precipitation can be occurred as a result of slow cooling rate as oil quenching which
cause depletion solute atoms from matrix. Samples quenched in oil behave dominant
intergranular fractures in contrast to that were quenched in water which behave a
27
transgranular fracture [38]. In order to reduce warping, quenching directly to aging
temperature may be better than water quenching [21].
The final mechanical properties are also affected by the alloy grain size. The factors affecting
grain growth during SHT on Al-Mg-Si alloy were studied by Zhuang et al and found that
controlling the grain size was achieved by large dispersoid like Fe-rich phase which acts as a
favorable sites for the nucleation of new grains and the fine dispersoid like Mn-phase, Zr-
phase or Cr-phase which pin grain boundary migration [39].
2.5.2 Aging Condition
2.5.2.1 Time-Temperature Variation
Mechanical properties of the Al-alloy was affected by the time and temperature variation
which also play an important role in the precipitation hardening process of Al-alloy showed
by many experimental works [40]. Vacancies assisted diffusion mechanism and the formation
of high volume fraction of Guinier Preston (GP) zones which disturb the regularity in the
lattice result in initial increase in tensile properties and hardness. Increasing the aging time to
an over-aging condition, at which the individual particles increase in size, but the number of
particles decreases and this cause an increase in inter-particle spacing, therefore the
mechanical properties decrease. Aging of 6063 Al-alloy between 8 and 10 h at 175o C is the
most suitable combination of time and temperature achieving maximum tensile strength, yield
strength and hardness to the alloy. Another study of aging on the same alloy at two
temperatures 160 and 250o C for various times showed that the alloy attains its peak hardness
after 64 h at 160o C and 2.5 h at 250
o C; the peak strength is much higher at 160
o C [41].
Aging at the higher temperature produce earlier but somewhat lower peak hardness compared
with aging at lower temperature and also accelerate the over-aging condition. Artificial aging
of alloy 6061 at 175o C and 200
o C showed that the time to peak aging at 175
o C was 8 h
while this time considers an over-aging condition at200o C [42].
2.5.2.2 Two-step aging
Three decades ago, it was reported that the number of density and the size of the precipitation
products in the peak hardness condition are highly sensitive to the pre-aging condition in two-
step aging, and they proposed a kinetic model to explain their microstructure observation
under various heat treatment conditions. Until recently, this two-step aging process of Al-
alloy was mostly of scientific interest. However, renewed interest in this process has been
stimulated by the possibility of using these alloys for automobile body sheet, where age
hardening is carried out during the paint-bake cycle [43]. The two step aging process consists
of scheme like in Fig. 2.20.
28
Fig. 2.20 Al-Mg2Si-Two step aging
A benefit of two-step aging was explained by Ber's [44] which aimed to obtain the same
strength level with lowering the aging time. He observed that for Al-alloy 6063 and Al-alloy
6061 used in his studies aging the alloys at 165o C for 1 h and subsequent aging at 220 for 0.5
h gives approximately the same strength level as the alloy aged with the standard long aging
time.
Jacobs [45] and Pashley et al [46] study the mechanism of two-step aging whereas pre-aging
is carried out at room temperature. Their studies explain the influence of clustering during the
delay in the first step of aging at room temperature on the stability of clusters when the
temperature is suddenly raised to the second step aging of artificial aging at a higher
temperature. They found that for a very short delay for a few minutes, all of very small
clusters formed at room temperature dissolve when the aging temperature is suddenly raised,
and re-nucleation occurs at the second step artificial aging. For slightly longer delays of a few
hours, a percentage of clusters formed at room temperature survive where the sudden increase
in aging temperature; the cluster density is lower than that which results from a very short
delay.
More studies carried out to select the proper pre-aging temperatures for an Al-1.4%Mg2Si-
0.3%Si alloy and final aging at 175oC, and showed that pre-aging at a temperature
lowerthan70oC will give a negative effect on mechanical strength while pre-aging at a
temperature higher than 70o C will give a positive effect on mechanical strength [47]. Pre-
aging studies carried out on Al-Mg-Si alloy (0.65%Mg and 0.7%Si) proposed that pre-aging
at a temperature higher than 70o C increases the density of the precipitate dispersion formed
after subsequent artificial aging at 175o C because of GP zones formed during pre-aging are
large enough to serve and act as heterogeneous nucleation sites for β'' precipitates. However,
the co-clusters that formed during natural aging reduce the hardening response at 175o C
because they revert at the artificial aging temperature [43]. Muryama et al [48] also explain
the two-step aging process, and found that the density of the GP zones is significant affected
by the pre-aging condition. Pre-aging at temperatures higher than 70o C increase the number
density of GP zones in artificially aged alloy, but natural aging suppresses the precipitation
kinetics of the GP zones in artificial aging. This part can be summarized that the GP-Zones
formed in the pre-aging condition grow in the subsequent artificial aging process, but the co-
clusters formed by natural aging are completely reverted.
29
2.5.3 Chemical Composition
The presence of selected trace element additions or micro-alloying such as Mg, Cu, Si, Mn,
Fe, Cr, Zr and other alloying element considered the most important factor affecting on the
process and/or kinetics of precipitation hardening.
a) Silicon Addition
Mg and Si consider the two major alloying elements in Al-Mg-Si alloys, which form the
equilibrium phase β or Mg2Si of approximate ratio 1.73:1, whose one of its metastable forms
is reasonable for hardening. The maximum solid solubility of Mg2Si in α-solid solution is
1.85% at 585oC. It decreases with decreasing the temperature. Excess Si up to some level
increases the strength because it enhances precipitation [17]. Mondolfo studied Al-Mg2Siand
found that excess Si accelerates precipitation and hardening [35]. It is not believed that the
excess Si alter the precipitation sequence, crystal structure and lattice parameters of the
different metastable precursors, but rather promotes the formation of additional phases which
do not contribute to hardening significantly. The existence of excess Si changes the
composition and density of metastable β" phase; it modifies the Mg/Si ratio in the clusters
zones and β" precipitates and improves strength by changing their size, number, density and
distribution. Moreover, the rate of strengthening increases until the overall Mg and Si ratio in
the alloy is close to approximately 0.4%. The hardening precipitates with reduce Mg to Si
ratio becomes less stable with aging and cause a decrease in strength during over-aging, and
as a result, stability in strength beyond the peak aging condition is somewhat reduced [49]. It
is known that excess silicon, i.e. silicon above that required for producing Si/Mg ratio of the
Mg2Si phase, improves the mechanical properties of Al-Mg-Si alloys. Previously this was
attributed to more silicon clusters that nucleated a denser dispersion of β" precipitates.
However this effect may be at least partly due to a greater volume fraction of β" precipitates
in the "excess" silicon alloys [50].
In order to identify the proper amount of excess silicon, an investigation [51] was made on
two Al- Mg2Si alloys containing 0.3 and 1.0 Si in excess. The first one reaches peak hardness
faster than the second i.e. when excess silicon reaches certain amount, Si atoms tends to form
Si particles preferentially, which reduce both Si content in solid solution and vacancies,
affecting the β" precipitations. Excess Si in a level of 0.4-0.5% in alloy containing 0.9 %Mg
is benefit but more is harmful especially on toughness because it segregates on grain
boundaries promoting the intergranular fracture were founded by Dorward and bouvier [36].
They also found that the peak strength was enhanced by 10-15 MPa for each 0.1 % excess Si
with corresponding decreases in the elongation about 0.25 %.
Another studies carried out to show the effect of excess Si on precipitates composition, it
carried out on the solute clusters, GP zones and β" precipitate within the Al-0.65% Mg-0.70%
Si (Si-excess) and the Al-0.70% Mg-0.33%Si (balance) alloys after pre-aging (70oC for16h)
or artificial aging (175oC for 10min). It was found that the Si excess is the reason that the
spherical GP zones and the needle-shaped β" precipitates contain more Si than those in the
balanced alloy otherwise the excess Si has to form its own precipitates or clusters [48].
31
b) Mg Additions
Excess Mg will reduce the response of the alloy to heat treatment especially at higher content
of Mg2Si [35]. The investigation done on AA6061 confirms that the excess Mg has a negative
impact on precipitation hardening because it lowers the solubility of Mg2Si in Al-Mg-Si
alloys [36].
c) Mn, Cd, Ag and Cr Additions
Additions of Mn or Cr at level of 0.4-0.7% and 0.3% respectively reduce the harmful effect
associated with excess Si, and also found that Mn and Cr reduce grain boundary precipitation,
thus reducing embitterment and susptability to intergranular corrosion. It was also found that
Cd and Ag retard GP zones formation where they reduce effect of natural aging, and
accelerate intermediate phase formation, also elimination of retrogression results [35].
d) Fe Additions
Mondolfo proposed that Fe and Zn do not have appreciable effect on precipitation [35],
whereas Tanihata et al found that increasing Fe in the range 0-0.3% in AA6063 reduces the
peak hardness reasonably [52]. Fe additions play an important role during SHT (solution heat
treatment) because it forms large dispersions which act as nucleation sites for the new grains
at the SHT temperature (grain refining activator) [39].
e) Cu Addition
The addition of copper increases the peak hardness and yield strength during aging of Al-Mg-
Si alloys. Copper will concentrate in the precipitates and increase the volume fraction of it.
First investigation of the strengthening due to copper addition was caused by the additional S'
and Q' precipitation while another type of precipitates reported in Al-Mg-Si-Cu alloys is the
quaternary Q phase with a lath morphology formed at later aging stages [53].It was found that
the addition of Cu changes the precipitation sequence and phases coexist.
With the alloy free from Copper alloy at UA (under-aged condition) and PA (peak-aged)
condition, the alloy contains only one precipitating phase identified as β" with a monoclinic
crystal structure and Mg/Si ratio increase from 0.8 to 1.01% with increasing aging time while
in alloy containing copper two phase coexist β" (monoclinic crystal structure) and Q
(hexagonal crystal structure) [54]. It was found that Cu induces the formation of Q and its
precursor metastable phases and has a beneficial effect on the kinetics of artificial aging, for
the alloy containing 0.07% Cu, the precipitation sequence may be GP zones → needlelike β”
→ rod-like β‟ + lath-like Q‟ → β+Si. On the otherhand, the precipitation sequence in the alloy
containing 0.91% Cu may be GP zones → needlelike β” →lathlike Q‟ →Q + Si [55].
The addition of Cu result in improved tensile properties and this appears to be mainly due to
refinement in the precipitate dispersion; this result was pointed out by Ringer and Hono [43].
Recent Three Dimensions Atom Probe (3DAP) results indicate that Cu is incorporated
exclusively in the β" precipitate phase while it is not a constituent of the clusters or GP zones.
Another studies have reported the presence of the Q' and Q phases (Al5Cu2Mg8Si6), these
precipitates are observed only after prolonged aging time at elevated temperatures, and do not
contribute to age hardening during the usual industrial heat treatment [43].
31
Mondolfo [35] found in his investigation that presence of Cu reduces the effect of delay in
natural aging Al-Mg-Si alloys.
f) Mg2Si Content
Increasing Mg2Si content (0.77, 1.19 and 1.68) in Al-Mg-Si alloy reduces the aging time to
reach maximum hardness at the same temperature (500, 200 and 100 min); this result was
pointed by Mahota and Takeda et al. Alloy with Mg2Si content (1.19 and 1.68) shows a
different precipitation sequence than that have Mg2Si content (0.77). TEM observations show
that Mg2Si content affects the precipitates formed during aging; at higher content needle like
shaped precipitates accompanying high strain contrast were observed while in the lower
contents larger precipitates (rod shape) were the predominant [56].
It was found that the alloys possess higher solute content of Mg2Si in the alloy produce a
higher density of needle like shaped precipitates and consequently the mechanical properties
of the alloy increases [57].It was found that the peak strength increased by 5 MPa per 0.1%
Mg2Si content [36].
32
CHAPTER 3: MATERIALS AND EXPERIMENTAL
WORK In this chapter, experimental techniques used for modification mechanical properties of Al-
Mg-Si 6351 alloy by precipitation hardening, several steps had been made so as to improve
mechanical properties. The material used in this work is Al-Mg-Si 6351 alloy.
3.1 Materials
The alloy used in the present study is a grade Al-alloy. Its chemical composition is given in
table 3.1, the chemical composition emphasizes that the material under analysis is Al-alloy
6000 near the specifications of grade 6351 (AlSiMg0.5Mn). The as received condition was
extruded seamless tube (20.6 mm outer diameter and 6.7 mm wall thickness) and its
mechanical properties in an extruded condition referred to a temper T54. (Property limits in
an extruded condition according to ASTM standard specimen, T54 temper; tensile strength
(min) 207 MPa; 0.2% yield strength (min), 138 MPa; elongation (min), 10%).
Table 3.1 Chemical composition of Al-alloy 6351 used in the present work
Element Si Fe Cu Mg Mn Cr Ti Al
Received
alloy 0.918 0.395 0.0319 0.735 0.511 0.0211 0.0258 97.35
Standard
composition
ASTM
0.7-1.3 0.5
max
0.1
max
0.4-0.8 0.4-0.8 ----- 0.2
max
bal
3.2 Heat-treatment
The specimens were heat-treated in a chamber furnace, whose temperature reached to 3000oC
and controlled by ± 5oC as shown in figure 3.1. Heat -treatment was conducted at Central
Metallurgical Research and Development Institute (CMRDI). Solid solution heat treatment
was firstly conducted at 540oC for 45 min in to obtain α–super saturated solid solution then
followed by water quenching. After solution treatment specimens were divided into 3 groups;
the first group was specialized to study the artificial aging behavior of the alloy by varying the
aging temperature from 160oC to 260
oC and aging time from 0.5 to 32 h, second group was
prepared to study the effect of natural aging and the final group was conducted to study the
effect of pre-aging on artificial aging in terms of the tensile properties of the alloy. The heat
treatment processes which were performed in this study are summarized in figure 3.2.
33
Fig. 3.1 Heat treatment furnace
Fig. 3.2 Heat treatment process
34
In order to achieve optimal mechanical properties, three steps including solution, quenching
and aging are generally used in heat-treatable:
(1) SHT: this is where the material is held at 540oC for 45 min, so that all the elements are
taken into solution, resulting one single phase.
(2) Quenching: this is when the material is rapidly cooled from the SHT temperature to room
temperature so as to -freeze- this super-saturated state within the material at room
temperature, giving a microstructure condition known as „Super Saturated Solid Solution‟
(SSSS).
(3) Ageing: age-hardening is the final stage in the development of the properties of heat
treatable alloys which controlled the decomposition of the SSSS to form finely dispersed
precipitates. Some alloys undergo aging at room temperature (natural aging), but most require
heating at a certain temperature for a time interval (artificial aging).
Fig. 3.3 Age hardening sequence of Aluminum alloys
3.3 Tensile Test
Standard plate tensile specimens were prepared with 12.35 mm width, 6.7 mm thickness and
50.0 mm gauge length. The tensile specimens were machined from extruded tube of
investigated alloy by cutting machine according to ASME (E8), as shown in fig 3.4. Coolant
was used during machining and cutting. Suitable strips were taken from the main tube in the
longitudinal direction for making tensile test samples. The tensile tests were carried out to
fracture at room temperature using tensile machine of type (UH-X Japan) as shown in figure
3.5. Tensile test was conducted at Cairo University, Faculty of Engineering (Mechanical
Testing Lab MTL). The tensile test machine has loading range from 0 to 10 ton. According to
35
ASTM specification the cross-head speed 5 mm/min was used to investigate the tensile
properties of the alloy. The machine was equipped with a chart recorder for the stress-strain
curves, which synchronized with the crosshead speed.
Fig. 3.4 Tensile test specimen according to ASME E8
Fig. 3.5 Universal tensile testing machine
36
3.4 Hardness test
The Vickers hardness number was measured by using hardness machine of type
(Zwick/ Roell ZHU260 Standard, Germany and 115/230V-50/60Hz), as shown in figure.
Hardness test was conducted at Cairo University, Faculty of Engineering (Mechanical Testing
Lab MTL). Load of 5 kg was applied and the times of loading 15 seconds 4 indentations were
taken on surface of specimen.
The calculated Vickers hardness according to the following equation:
(
)
W = Weight in (kg)
d = Average diameter of indenter (mm)
Fig. 3.6 Hardness Test Machine
37
3.5 XRD Analysis
X-rays diffraction (XRD) is a very powerful technique used to investigate crystal structure of
a material as well as its chemical compounds from their crystalline structure detected. In this
work, XRD instrument of type (X‟Pert PRO PANanalytical) was used. The in-plane
diffraction technique was used so as to determine the crystal structure of the film. XRD
analysis was conducted at Central Metallurgical Research and Development Institute
(CMRDI). Diffracted x-rays are detected and analyzed using computer software. Results of
this analysis are displayed on the screen as a graph between beam intensity (counts) versus
angle of incidence of x-rays beam (2θ). Before specimens‟ spectroscopy, anodized face is cut
into dimensions to fit machine‟s stage dimensions. The XRD instrument is shown in figure
3.7.
Fig. 3.7 XRD Machine
38
3.6 Microstructure Examination
After finishing all heat treatment conditions, various techniques were applied for
metallographic preparation and examination. The specimen was ground and then polished into
a smooth finishing by using abrasive papers of grades 400, 600, 800, 1000 and 1200 followed
by polishing using fine alumina powders (50g / 500 ml H2O). Finally the specimens were
etched in Keller solution containing 0.5% HF in 50ml H2O. After that, the specimen was dried
well and then microscopic examination, optical microscope of type Carl Ziess Baujahr and
100/240V-50/60Hz was used as shown in figure. Optical Microscope investigation was
conducted at Central Metallurgical Research and Development Institute (CMRDI)
Fig. 3.8 Optical Microscope
3.7 Fractographic Examination Scanning electron Microscopy (SEM) is a technique used to investigate surface
topography, surface morphology as well as elemental analysis and compounds on the
surface (relative amounts of them). Fracture surface for some selected specimens was
examined by using Scanning Electron Microscope of type (Jeol-KSM 5410, Japan and of 30
Kv) as shown in figure 3.9. In this work, the SEM analysis was done at SEM laboratory in
Egyptian Mineral Resources Association using Quanta FEG 250 instrument as shown in
Figure 3.9.
39
Fig. 3.9 Scanning Electron Microscope (Quanta FEG 250)
3.8 Energy Dispersive X-rays Analysis
Energy dispersive x-rays analysis is a technique that depends on the characteristic x -rays
emitted by an element‟s atoms when being irradiated by a high-energy beam. This
technique is used for investigating the presence and quantity of chemical elements exist in
analyzed sample. The elemental analysis was done in SEM laboratory in Egyptian
Mineral Resources Association using the same instrument used in scanning electron
microscopy section, Quanta FEG 250.
41
CHAPTER 4: RESULTS AND DISCUSSION
4.1 Effect of Artificial Aging on Tensile Properties
The artificial aging behavior of Al-Mg-Si alloy 6351, in which was solution treated at 540oC
for 45 min followed by water-quenching and then immediately aged without any delay at
room temperature at 160oC, 175
oC, 200
oC and 260
oC as is illustrated in figures 4.1, 4.2 and
4.3 in terms of hardness, yield stress and ultimate tensile strength respectively versus artificial
aging time. The peak strength values were reached after approximately 18 h at 160oC for
hardness, yield stress (YS) and ultimate tensile strength (UTS). The corresponding peaks of
hardness and strength at previous temperature higher than 160oC were reached after shorter
times as shown in figures. The maximum peak hardness (108 HV) and peak strength
(305MPa) for UTS were achieved by aging at 160oC for 18 hours, while the maximum peak
strength (279 MPa) for YS was achieved by aging at 260oC for 1 hour as shown in figure 4.2.
The yield stress is more sensitive to both changing in aging temperature and aging time than
the ultimate tensile strength and hardness as shown in figures 4.1 and 4.3.
Fig 4.1 Effect of artificial aging on tensile strength for Al-alloy 6351
100
150
200
250
300
350
400
0 5 10 15 20 25 30
Ult
imate
ten
sil
stre
ngth
, M
pa
Artificial aging time, h
AA at 261ᵒC AA at 211ᵒC AA at 175ᵒC AA at 161ᵒC
41
Aging times lower than the peak aging time represents the under-aging condition, while those
higher represent the over-aging condition. Aging at temperatures higher than 160oC for times
higher than the peak aging time leads to quick over-aging which is manifested by rapid
decrease in hardness and ultimate tensile strength as observed at 260oC. It was found that
over- aging occurred after one hour at 260oC while it occurred after 18 hours at 160
oC. It was
found that, as the aging temperature increase (260oC) the peak strength reached at lower time
1 hours but little lower in value than that obtained in case of lower aging temperature (160ᵒC
at 18 hours).
Fig 4.2 Effect of artificial aging on 0.2% offset yield stress for Al-alloy 6351
Figure 4.4 illustrates the effect of artificial aging time and aging temperature on the total
elongation. At the same aging temperature as the aging time increase the total elongation
decrease in the early stage of aging but for longer aging times they showed slight change with
time reaching a saturation value.
100
150
200
250
300
0 5 10 15 20 25 30
0.2
% o
ffse
t yie
ld s
tres
s, M
pa
Artificial aging time, h
AA at 261ᵒC
AA at 211ᵒC
AA at 175ᵒC
AA at 161ᵒC
42
Fig 4.3 Effect of artificial aging on hardness for Al-alloy 6351
Fig 4.4 Effect of artificial aging on total elongation for Al-alloy 6351
50
70
90
110
130
0 5 10 15 20 25 30
Hard
nes
s, H
V
Artificial aging time, h
AA at 261ᵒC
AA at 211ᵒC
AA at 175ᵒC
AA at 161ᵒC
0
10
20
30
0 5 10 15 20 25 30
Tota
l el
on
gati
on
, %
Artificial aging time, h
AA at 261ᵒC
AA at 211ᵒC
AA at 175ᵒC
AA at 161ᵒC
43
In general as the aging temperature increase the total elongation decreased; this was occurred
from the early stage of aging. The saturated value of total elongation after aging at 160oC was
about 25% while it was about 20% after aging at 260oC. It was observed that for each aging
temperature the minimum elongation is corresponding to the peak strength, as shown in
figures 4.2 and 4.4.
It is clearly observed that changing the aging temperature or changing aging time causes the
same effect on total elongation. If the total elongation divided into uniform and necking
elongation, the decrease of the total elongation by increasing the aging time or aging
temperature is mainly caused by the decrease in the uniform elongation.
Figures 4.5 – 4.9 show the true stress-true strain of the as received alloy, the solution treated-
water quenched condition, and the artificially aged conditions with the calculation of strain
hardening exponent for each condition. The artificial aging for 4 hours at 160oC represents the
under-aged condition while artificial aging at 160oC for 18 hours represents the peak aged
condition finally aging at 160oC for 24 hours represents the over-aged condition. All true
stress-true strain curves were drawn up to the true maximum tensile strength. Generally all
true stress-true strain curves show parabolic hardening after yielding. Artificially aged alloy at
160oC raises the true stress-true strain curve level to much higher stresses compared to
solution treated-water quenched only. Strain hardening exponent for every condition is
calculated in the range from true yielding stress to true maximum tensile strength according
the following equation:
σ = K
ln(σ) = ln(K) + n ln(ε)
Where;
σ is true stress, MPa ε is true strain,
K is strengthening coefficient, MPa intersect part and
n is strain hardening exponent, slope
Conditions are summarized in table 4.1 as follow:
Table 4.1 Strain hardening exponent and strengthening coefficient of solution treated-
water quenched alloy and artificially aged alloy.
Heat treatment
Strain hardening exponent Strengthening coefficient
As received condition 0.1389 299.7
Solution treated-water quenched 0.2022 364.5
Aging at 160oC for 4 h
0.1752
413.6
Aging at 160oC for 18 h
0.0801
395.1
Aging at 160oC for 24 h
0.0982
393.2
44
Fig 4.5 true stress-true strain curve of the received Al-alloy 6351
0
50
100
150
200
250
0 0.02 0.04 0.06 0.08 0.1 0.12
tru
e st
ress
, M
Pa
true strain
y = 0.1389x + 5.7028
4.8
4.9
5
5.1
5.2
5.3
5.4
5.5
-7 -6 -5 -4 -3 -2 -1 0
ln (
σ)
ln (ε)
45
Fig 4.6 true stress-true strain curve of solution treatment-water quenched of Al-alloy
6351
0
50
100
150
200
250
300
350
0 0.05 0.1 0.15 0.2 0.25 0.3
tru
e st
ress
, M
Pa
true strain
y = 0.2022x + 5.8984
0
1
2
3
4
5
6
-8 -7 -6 -5 -4 -3 -2 -1 0
ln (
σ)
ln (ε)
46
Fig 4.7 true stress-true strain of artificially aging Al-alloy 6351 at 160oC for 4 h
0
50
100
150
200
250
300
350
0 0.05 0.1 0.15 0.2 0.25
tru
e st
ress
, M
Pa
true strain
y = 0.0771x + 5.6036
y = 0.2274x + 6.1521
5.1
5.2
5.3
5.4
5.5
5.6
5.7
5.8
5.9
-7 -6 -5 -4 -3 -2 -1 0
ln (
σ)
ln (ε)
47
Fig 4.8 true stress-true strain of artificially aging Al-alloy 6351 at 160oC for18 h
0
50
100
150
200
250
300
350
400
0 0.05 0.1 0.15 0.2
tru
e st
ress
, M
Pa
true strain
y = 0.0278x + 5.7261
y = 0.1242x + 6.0954
5.5
5.55
5.6
5.65
5.7
5.75
5.8
5.85
5.9
-8 -7 -6 -5 -4 -3 -2 -1 0
ln (
σ)
ln (ε)
48
Fig 4.9 true stress-true strain of artificially aging Al-alloy 6351 at 160oC for 24 h
0
50
100
150
200
250
300
350
400
0 0.02 0.04 0.06 0.08 0.1 0.12 0.14 0.16 0.18
tru
e st
ress
, M
Pa
true strain
y = 0.0366x + 5.6902
y = 0.1486x + 6.1057
5.4
5.45
5.5
5.55
5.6
5.65
5.7
5.75
5.8
5.85
-7 -6 -5 -4 -3 -2 -1 0
ln (
σ)
ln (ε)
49
Figure 4.10 shows true stress-true strains for as received, as quenched and aging at 160oC for
various times represents under-aging, peak-aging and over aging conditions. It was found that
aging for 18 hours at 160oC show higher level of the true stress-true strain curves in
comparison with the solution treated-water quenched and other aging conditions.
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
Alloy 6351 is a heat treatable Al-Mg-Si alloy where its strength can be improved by
precipitation hardening process. Precipitation hardening is a temperature and time dependent
process; it is achieved by aging Al-alloys artificially in a temperature range 160 – 260oC after
solution treatment for 45 min at 540oC in order to dissolve nearly the whole second phase
particles followed by water-quenching to room temperature in order to have high
concentration of vacancies and a supersaturated solid solution of α- solid solution with mainly
dissolved Mg and Si atoms at room temperature. This treatment results in slight increase in
either yield stress or ultimate tensile strength and slight decrease in total elongation in
comparison with the result obtained from solution treated water-quenched of material. The
high increase in vacancies is necessary in order to enhance the diffusion of magnesium and
silicon atoms through artificial aging in order to form a sequence metastable Mg/Si phases
until the equilibrium phase β (Mg2Si) forms.
0
50
100
150
200
250
300
350
400
0 0.05 0.1 0.15 0.2 0.25 0.3
tru
e st
ress
, M
Pa
true strain
As Quenched
As Received
Under-aging
Peak-aging
Over-aging
51
Precipitation hardening manifests itself within Al-alloy 6351 by peaks either in hardness,
yield stress or ultimate tensile strength in its relation with artificial aging times as shown in
figures 4.1, 4.2 and 4.3 and a drop in elongation as shown in figure 4.4 related generally with
intergranular fracture.
The age hardening response of the alloy is quite significant and the time, temperature
variations principally have an effect on the tensile properties. The peak aging time decreases
with increasing the temperature of aging as shown in figures 4.1, 4.2 and 4.3. At a given
artificial aging temperature the strength increase with increasing artificial time of aging
(under aging condition) until reaching a maximum value of strength (peak-aged condition).
Further increase in the artificially aging time above the peak time reduces the strength of the
alloy (over-aged condition).
The early stage of the artificial aged Al-alloy 6351 that were investigated recently showed
that the initial stages of precipitation involve the separate cluster of Mg and Si atoms after that
Si/Mg co-clustering. After this stage small zones are formed, referred as Guninier-Preston
Zones (GP) [42, 43]. The initial increase in the strength as shown in figures 4.1, 4.2 and 4.3
may also be as a result of vacancies type SSSS and therefore the formation of high volume
fraction of (GP) Zones that disturb the regularity in the lattice [38]. Once this stage is formed,
a needle like formed β” precipitate (≈ 07 nm long) is that the predominant and characterizes
the under-aging condition [42]. This aging treatment significantly increases the strength,
representing that the effectiveness of β” phase in strengthening is greater than that of small
equiaxed GP-zones [42].
Increasing the artificial time of aging, larger needles of β" precipitates are formed which
represents (peak-aging condition), referred to the max age hardening response. It was found
that the β" precipitate grows in its length direction which is coherent with the matrix resulting
in an increase in the irregularity in the lattices and causes an increase in the strength of this
alloy [42].
β” precipitates transformed to β‟ precipitates in the aging sequence by Increasing the artificial
time of aging or aging temperature, which has a rod shaped and less incoherency than the
needle formed precipitates because its cross section is incoherent with the matrix [42]. The
formation of this phase results in a drop in strength and represents the starting of the over-
aging stage. Further heat treatment at higher temperatures and time reduces hardness, yield
stress and ultimate tensile strength of the alloy and results in over-aging of the alloy. During
this case stable β (Mg2Si) phase forms as platelets shape. This precipitate contributes little to
strength of Al-Mg-Si alloys because it is completely incoherent with the matrix [42, 44]. It
was found that the precipitation of phase β (Mg2Si) and the presence of intermetallic phases
that formed during solidification of the alloy itself highly affect mechanical properties. These
formed precipitates have an important effect on lowering the volume fraction of the hardening
phase β” (Mg2Si) [61]. Microstructure of all Al-Mg-Si alloys contained a mixture of Al3Fe,
spherical of α-AlFeMnSi and plate like β-AlFeSi intermetallic phases distributed at grain
boundaries, accompanied sometimes with coarse Mg2Si[62]. The presence of a brittle and
monoclinic hard phase β-AlFeSi causes a poor surface finishing and reduces workability.
51
Generally, the increase in strength may be obtained by inhibit dislocation movement. The
main concept of the precipitation hardening process for increasing strength is the formation of
second phase particles that impede the dislocation motion. The dislocation-particles
interaction can be produced by one of two mechanisms, first cut through the precipitate
particles and second by pass the obstacles forming dislocation loops. Larger particles work as
void nucleation sites so damage are obtained [62].
It was noticed that in case of the under-aged up to peak-aged conditions as the size of
precipitates increases the strength increase as shown in figure 2.7 [9] and the first mechanism
is supposed to occur in the primary stages of aging (under-aged and peak-aged conditions)
where fine coherent precipitates or zones existed and the strengthening depend on the nature
and the size of precipitates.
The formed precipitates increase in size due to clustering of the smaller precipitate particles
into larger particles so the number of particles decreases and therefore the distance between
them increases, which will cause fewer obstacles to the movement of dislocation and it is easy
for moving dislocation by-passing precipitates and then tensile strength decrease [38]. The
second mechanism is characteristic of the over-aged alloy that has coarse dispersed particles
and it is impossible for dislocation to cut through incoherent precipitates and therefore the
strengthening depends on the size of the precipitates, as shown in figure 2.7, wherever
strength decreases as the particle size increases.
The shifting the peak strength to shorter artificial time of aging with higher aging temperature
during this alloy, as shown in figures 4.1, 4.2 and 4.3, this may be due to the high rate of
aging process which is diffusion dependent process i.e. the higher rate of diffusion of Si and
Mg atoms to form the fine coherent metastable β” phase which is responsible for the peak
hardening [42]. So aging at high temperatures produced earlier peak strength in comparison
with aging at lower temperature.
It is noticed that higher peaks of ultimate tensile strength and hardness as shown in figures 4.1
and 4.3 values are observed at lower aging temperature than higher aging temperature this
may be due to lower the artificial aging temperature may increase the nucleation sites,
therefore more number of precipitates are formed which imped dislocation movements [63],
also as a result of straining the matrix during precipitation of intermetallic phases with a
number of intermediate metastable stages. As the aging temperature decreases, tensile
strength and hardness of the material increases with a drop in toughness and ductility as
shown in figures 4.1, 4.3 and 4.4.
The yield stress of this alloy is more sensitive to the artificial aging temperature and also the
artificial time of aging as shown in figures 4.2 than the ultimate tensile strength within the
under-aged condition; this may be due to the higher stresses are required in the early stage of
deformation to cut through precipitates. The higher peak yield stress at 260oC than 160
oC (see
figure 4.2) is also because of the presence of copper that plays an important role. In Al-Si-Cu
alloy, it was found that aging at temperature more than 230oC Copper precipitates rapidly and
may change the precipitation sequence of the investigated alloy, this leads to higher volume
fraction and coarser size of precipitates at 260oC than 160
oC resulting in higher yield stress
[43, 54, 64]. Additional study required for this point
52
The total elongation of Al-alloy 6351 after solution treated-water quenched to room
temperature followed by means of artificially aging showed similar trend in its relationship
with temperature and time as shown in figure 4.4. The drop in total elongation with increasing
time of aging or increasing aging temperature depends on the precipitation behavior of the
alloy. This drop in elongation associated with the occurrence of intergranular fracture in
under-aging condition nearly (≈ 50% of the total fracture) and the peak-aging condition (≈
70% of the total fracture) while it show a completely dimple fracture in the over-aging
condition.
The drop in total elongation associated with intergranular fracture (fracture along grain
boundary) during the course of precipitation hardening may be explained as follows, the
intergranular fracture showed in the fracture surfaces in under-aged condition and peak-aged
condition may be due to the combined effect of precipitation of large Si particles (100 nm)
along the grain boundaries, the presence of intermetallic phases such as Al3Fe, β-AlFeSi and
α-AlFeMnSi a long grain boundaries and the occurrence of planer slip as observed by Zehn
[51] particularly in low excess Si-content Al-Mg-Si alloy. The planar slip occurs in under and
peak aged condition, and it may be more with the increase in the volume fraction of the β"
precipitates and the increase in its critical size [40, 51], where large dislocation pile ups occur
at the precipitate-matrix interface and therefore the dislocation are concentrated in narrow
bands on the slip system.
The effect of planar slip on the occurrence of intergranular fracture can be discussed as
follow, Precipitates free zone (PFZ) form adjacent to the grain boundaries. This zone is softer
than the interior of the grain, thus high stress is accumulating during this region, and the voids
formed by high stress activate fracturing. The interface of grain boundary precipitates and
PFZ is comparatively week. Stress concentration caused by shearing mechanism at the tip of
slip bands nucleates voids at the grain boundary precipitates, resulting in intergranular failure
[40, 51, 57], thus; the main drop in elongation of artificially aged Al-alloy 6351 in (the under-
aged and peak-aged conditions) may be due to the formation of slip bands which nucleate
voids, the precipitation of Si-rich phase particle along the grain boundary and the precipitation
of intermetallic phases formed during solidification.
4.2 Factors Affecting the Artificial Aging
4.2.1 Natural Aging
Room temperature natural aging of Al-alloy 6351 solution treated at 540oC and the water
quenched was studied to investigate the effect of natural aging time on tensile properties and
on artificial aging.
4.2.1.1 The Influence of Natural Aging Duration on Mechanical Properties
Variation in natural aging time is illustrated in figures 4.11, 4.12, 4.14 and 4.14 versus tensile
properties. Figures 4.10 and 4.11 showed that by increasing the natural aging time both yield
stress and ultimate tensile strength increase slightly, for samples naturally aged in the range of
53
24 to 1000 hours. It was observed that a large increase in both yield and tensile strength took
place in the first 20 hours then mild increase from 20 to 1000 hours.
It is interesting to mention that after 1000 hours aging at room temperature the yield strength
and the ultimate tensile strength reached the values of 142.3 and 290.4 MPa respectively
which are much lower than those 253.459 and 304.195MPa that were obtained by artificial
aging for aging time 18 hours at 160oC only (peak aged condition).
Figure 4.13 illustrates the effect of natural aging time on hardness HV; the results show that
by increasing the aging time, hardness increase slightly. The value of hardness 90.6 HV after
1000 hours aging at room temperature much higher than that (64.8 HV) for solution treated–
water quenched condition and much lower than that (108 HV) obtained by artificial aging at
160oC for aging 18 hours (peak aged condition). Figure 4.13 shows that the effect of natural
aging time on total elongation the results show that by increasing that aging time total
elongation decrease slightly. The value of total elongation (19.2449%) after 1000 hours aging
at room temperature is little lower that for solution treated – water quenched condition and
much higher than that (11.5385%) obtained by artificial aging for aging 18 hours at 160oC
(peak aged condition).
Figures 4.15 and 4.16 show true stress-true strain curves for naturally aged conditions at room
temperature for 170 hours and 1000 hours respectively with the calculation of strain
hardening exponent for each condition. All true stress- true strain curve show parabolic
hardening after yielding. Natural aging at room temperature cause slight increase in true
stress-true strain curves to level above that obtained from solution treated-water quenched
condition and lower than that obtained from artificially aged condition (peak-aged condition).
Raising the natural aging time from 170 hours to 1000 hours the level of the true stress-true
strain curve slightly increases.
v
54
Fig. 4.11 Change in 0.2%yield strength, MPa of Al-alloy 6351 due to the effect of natural
aging for various times
Fig. 4.12 Change in ultimate tensile strength, MPa of Al-alloy 6351 due to the effect of
natural aging for various times
50
100
150
1 10 100 1000 10000
0.2
% o
ffse
t yie
ld s
tres
s, M
Pa
Natural aging time, h
50
100
150
200
250
300
350
1 10 100 1000 10000
Ult
imat
Tn
sile
Str
ength
, M
Pa
Natual aging time, h
55
Fig. 4.13 Change in hardness, HV of Al-alloy 6351 due to the effect of natural aging for
various times
Fig. 4.14 Change in total elongation %, of Al-alloy 6351 due to the effect of natural aging
for various times
0
20
40
60
80
100
1 10 100 1000 10000
Ha
rdn
ess,
HV
Natural aging time, h
0
5
10
15
20
25
30
35
40
1 10 100 1000 10000
Tota
l E
lon
gati
on
%
Natural aging time, h
56
Fig 4.15 true stress-true strain of natural aging of Al-alloy 6351 at room temperature for
170 h
0
50
100
150
200
250
300
350
0 0.05 0.1 0.15 0.2 0.25
tru
e st
ress
, M
Pa
true strain
y = 0.0635x + 5.5021
y = 0.2527x + 6.1899
5
5.1
5.2
5.3
5.4
5.5
5.6
5.7
5.8
5.9
-7 -6 -5 -4 -3 -2 -1 0
ln (
σ)
ln (ε)
57
Fig. 4.16 true stress-true strain of natural aging of Al-alloy 6351 at room temperature
for 1000 h
0
50
100
150
200
250
300
350
0 0.02 0.04 0.06 0.08 0.1 0.12 0.14 0.16 0.18
tru
e st
ress
, M
Pa
true strain
y = 0.0568x + 5.4328
y = 0.2307x + 6.1179
5
5.1
5.2
5.3
5.4
5.5
5.6
5.7
5.8
-7 -6 -5 -4 -3 -2 -1 0
ln (
σ)
ln (ε)
58
Fig. 4.17 shows true stress-true strain curves of Al-alloy 6351 for naturally aged condition at
room temperature for 170 and 1000 hours in comparison with solution treated-water quenched
and peak-aging conditions. All true stress-true strain curves show parabolic hardening after
yielding. It was noticed that naturally aging at room temperature for 170 and 1000 hours
raises true stress-true strain to level above that of the solution treated-water quenched only
and much lower than that of the artificially aged condition (160oC for 18 hours). Increasing
the natural aging time the level of true stress-true strain curve will increase slightly.
Fig.2.17 true stress-true strain curves of Al-alloy 6351 for naturally aged condition in
comparison with solution treated-water quenched and peak-aging conditions
Al-Mg-Si alloy possesses a negative strength response at room temperature storage after
solution treatment-water quenched for short time. Although it is already known that this is
due to clustering during room temperature storage or natural aging [65].The increase in in
both tensile properties and hardness during natural aging is due to the high shearing
opposition to dislocations developed by the formed clusters compared to that of the super
saturation matrix [66].The higher diffusion rate and lower solubility of silicon in aluminum
during the initial stage of natural aging in 6xxx series alloys is a dominant factor in the
formation of clusters. After that magnesium takes part in the formation of clusters.
The formation of Si-clusters and Si-Mg co-clusters led to increase in tensile properties when
the investigated alloy spent at room temperature (NA) for long time after solution treatment-
water quenched as shown in figure 4.11 and figure 4.12.
0
50
100
150
200
250
300
350
400
0 0.05 0.1 0.15 0.2 0.25 0.3
tru
e st
ress
, M
Pa
true strain
As Quenched
NA 170 h
NA 1000 h
Peak-aging
59
Natural aging gives the alloy a mild increase in strength with a slight decrease in total
elongation as shown in figures 4.11, 4.13 and 4.14 and this may be due to Si clusters and Si–
Mg co-clusters are too small and week to contribute the alloy with a considerable strength.
The increase in strength depends upon the time spent at room temperature. It was found that
in Al-Mg-Si alloys clusters of Si atoms formed rapidly after quenching in addition to the
gradual formation of Si-Mg co-clusters atoms during natural aging with the aid of vacancies
formed after quenching which consider the main reasons for strengthening of the alloy at
room temperature [67].
The influence of RT storage was described by slower kinetics of precipitation of the β” phase;
the nucleation of β” phase is restricted by the presence of low concentration of vacancies
formed after quenching and the formed clusters and GP zones in the RT-stored specimens,
have a delaying consequence upon the nucleation of β” phase [68].Increasing in hardness
through natural aging may be due to the increasing of a high density of clusters that have a
higher shearing resistance to dislocations than the supersaturated matrix.
4.2.1.2 Effect of natural aging time on artificial aging
In this part of study the effect of natural aging time (NA) on artificial aging (AA) for various
artificial aging times, the effect of intermediate 170 and 1000 hours natural aging times
demonstrated in terms of tensile properties is shown in figures 4.18 – 4.25. It was found that
in general natural aging for 170 hours ( one week) or for 1000 hours followed by artificial
aging at 160oC for several times leads to much higher values of yield for yield stress up to the
over-aged condition obtained by (NA + AA) and little higher values for ultimate tensile
strength up to the peak aged condition obtained by (NA + AA), in comparison with obtained
only by artificial aging at 160oC as shown in figures 4.18 – 4.21.
The 170 h natural aging followed by artificial aging at 160oC shifts the peak for both ultimate
tensile strength and 0.2% offset yield stress to the artificial aging time of 8 h instead of 18
hours in case of artificial aging at 160oC only figures 4.18 and 4.19, while the 1000 h natural
aging followed by artificial aging at 160oC shifts the peak for both ultimate tensile strength
and 0.2% offset yield stress to the artificial time of 4 hours instead of 18 hours figures 4.20
and 4.21.
Figures 4.22 and 4.23 demonstrate the hardness, HV for 170 and 1000 hours natural aging
followed by artificial aging at 160oC for various times and it showed that in case of 170 hours
natural aging, the peak shifts to 8 hours in place of 18 hours obtained in case of artificial
aging and in case of 1000 hours natural aging, the peak shifts to 4 hours instead of 18 hours.
Natural aging for 170 hours or for 1000 hours followed by artificial aging at 160oC for
various time varied from 1 to 8 hours reduce total elongation to values lower than those
obtained by artificial aging at 160oC at all investigated artificial aging times as shown in
figures 4.24 and 4.25.
61
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
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
150
200
250
300
350
0 5 10 15 20 25 30
Ult
imate
Ten
sile
Str
ength
, M
Pa
Artificial Aging Time, h
NA(170 h) + AA
AA
50
100
150
200
250
300
350
0 5 10 15 20 25 30
0.2
% o
ffse
t yie
ld s
tren
gth
, M
Pa
Artificial aging time, h
AA
NA(170 h) + AA
61
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
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
150
200
250
300
350
0 5 10 15 20 25 30 35
Ult
imate
ten
sile
str
ength
, M
Pa
artificial aging time, h
NA(1000)+AA
AA
50
100
150
200
250
300
0 5 10 15 20 25 30
0.2
% o
ffse
t yie
ld s
tren
gth
, M
Pa
Artificial aging time, h
NA(1000h)+AA
AA
62
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
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
20
40
60
80
100
120
0 5 10 15 20 25 30
Hard
nes
s, H
V
Artificial aging time, h
AA
NA(170 h) + AA
20
40
60
80
100
120
140
0 5 10 15 20 25 30
Hard
nes
s, H
V
Artificial aging time, h
NA(1000h)+AA
AA
63
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
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
5
10
15
20
25
30
0 4 8 12 16 20 24 28
Tota
l el
on
gati
on
, %
artificial aging time, h
AA
NA(170 h) + AA
0
5
10
15
20
25
30
0 5 10 15 20
Tota
l el
on
gati
on
, %
Artificial aging, time
AA
NA(1000h) + AA
64
Figure 4.26 and figure 4.27 show true stress-true strains of two conditions which were
naturally aged at room temperature for 170 hours and 1000 hours followed by artificially
aging at 160oC for 8 hours and 4 hours respectively with the calculation of strain hardening
exponent for each condition. It was noticed that natural aging at room temperature followed
by artificial aging for 8 hours at 160oC showed higher level of true stress-true strain curves
above than obtained from peak-aged condition. Increasing the natural aging time for 1000
hours followed by artificially aging for 4 hours at 160oC showed a higher level of true stress-
true strain than that of peak-aged condition but lower than natural aging at room temperature
followed by artificial aging for 8 hours at 160oC. On the other hand, increasing the artificial
aging time for 18 hours for both cases showed lower level of true stress-true strain curve than
that obtained from peak-aging condition.
The effect of natural aging on the subsequent artificial aging at 160oC shown in figures 4.18 –
4.21, was found to depend upon the time spent to room temperature after solution treatment-
water quenching and followed by artificial aging at 160oC (peak-aged condition). This type of
treatment is common and regarded as two step aging.
Natural aging before artificial aging of the alloy 160oC may enhance the precipitation and
shifts the peak strength at 160oC to shorter time of artificial aging for both ultimate tensile
strength and yield stress than the artificial aging only at 160oC for 18 hours. The longer the
time of aging at room temperature the shorter time of artificial aging for peak strength but the
lower the peak in case of yield strength. These could be explained as follow, atomic clusters
such as Si- clusters and Si–Mg co-clusters are formed during natural aging in Al-Mg-Si
alloys; many of these atomic clusters during the earlier time of artificial aging treatment have
been dissolved, but some of the coarse Si-Mg co-clusters and much fine Si-clusters are likely
to survive [22, 43, 69]. The survived coarse Si-Mg co-clusters act as a nucleation sites for the
precipitation of β" phase and may result into a higher strength than that due to artificial aging
only.
On the other hand the quantity of vacancies formed after quenching and solute atoms
decreased as a result of aging at room temperature. The low initial numbers of vacancies and
solute atoms limits the formation of pre- β” and β” phase particles. The transportation of the
Si and Mg atoms from clusters and solid solution to pre-β” nuclei during artificial aging of
naturally aged specimens may be sluggish, which restrict pre-β”phase formation [24]. Also
the formation of Si-clusters during natural aging which are stable even at higher temperature
cannot act as nucleation sites for the precipitation of β” phase because of its very fine size, so
it contributes a negative effect to the subsequent artificial aging [22, 51, 69].The addition of
small amount of copper may reduce the negative effect of natural aging on subsequent
artificial aging; this may be due to copper segregate on Si-clusters during artificial aging and
reduce its bad effect on strength after artificial aging. The end result is that the negative
influence of natural aging is reduced with an overall increase in the strength level after natural
aging followed by artificial aging [64].
After one week of natural aging the clusters formed began to grow consuming more solute
atoms and excess vacancies [67] that may affect the precipitation of β” phase during the
subsequent artificial aging, resulting in the decrease in strength by increasing natural aging
time and behind one week of natural aging.
65
Fig 4.26true 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
0
50
100
150
200
250
300
350
400
0 0.01 0.02 0.03 0.04 0.05 0.06 0.07 0.08
tru
e st
ress
, M
pa
true strain
y = 0.0203x + 5.8172
y = 0.0638x + 5.981
5.68
5.7
5.72
5.74
5.76
5.78
5.8
5.82
5.84
-6 -5 -4 -3 -2 -1 0
ln (
σ)
ln (ε)
66
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
0
50
100
150
200
250
300
350
400
0 0.02 0.04 0.06 0.08 0.1
tru
e st
ress
, M
Pa
true strain
y = 0.0231x + 5.754
y = 0.0839x + 6.0193
5.55
5.6
5.65
5.7
5.75
5.8
5.85
-8 -7 -6 -5 -4 -3 -2 -1 0
ln (
σ)
ln (ε)
67
4.2.2 Pre-aging
Pre-aging at 100oC after solution treatment followed by water quenched and before artificial
aging of the Al-Mg-Si alloy 6351 for 18 hours at 160oC (peak-aged condition) was studied to
investigate the effect of pre-aging treatment on artificial aging behavior.
4.2.2.1 Effect of pre-aging time on artificial peak aging condition
The effect of changing the pre-aging time at 100oC on the artificial aged condition (peak aged
condition) at 160oC for 18 h, was studied in terms of change in tensile properties, hardness
HV and elongation as follow;
It was found that pre-aging at 100oC for various times 5-1000 min after solution treatment
followed by water quenching and before artificial peak aged condition leads to slight increase
in the ultimate tensile strength and a significant increase in the yield stress as shown in figure
4.28 in terms of strength increment (∆σ) which is the difference between tensile stress due to
pre-aging plus the artificial peak aging and that due to the artificial peak aging only.
It was observed that the largest stress increment in both ultimate tensile strength and 0.2%
offset yield stress was obtained after 10 min pre-aging time, the stress increment after 5 min
pre-aging time showed small increase in tensile properties. The stress increment decrease
gradually by increasing pre-aging time until there is a decrement in ultimate tensile strength
after pre-aging time of 1000 min as shown in figure 4.28.
Pre-aging at 100oC after solution treatment and before peak aging leads to slight increase in
total elongation at all investigated pre-aging times as shown in figure 4.29 in terms of ∆E%
which represents the difference between the total elongation obtained from pre-aging at 100oC
plus artificial aging for 18 hours at 160oC and that obtained from artificial peak aging only.
Pre-aging at 100oC after solution treatment followed by water quenching and before peak
aging leads to slight increase in hardness at all investigated pre-aging times as shown in figure
4.30 in terms of hardness, ∆HV which represents the difference between the hardness
obtained from pre-aging at 100oC plus artificial aging for 18 hours at 160
oC and that obtained
from artificial peak aging only.
68
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)
Fig.4.29Change in total 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)
-10
0
10
20
30
1 10 100 1000 10000
∆σ
, Mp
a
Pre-aging time, min
UTS
0.2% YS
0
1
2
3
4
5
1 10 100 1000 10000
∆E
, %
Pre-aging time, min
TE
69
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)
Figure 4.31 illustrates the true stress-true strain of the pre-aged condition at 100oC for 10 min
followed by artificial aging for 18 hours at 160oC with the calculation of strain hardening
exponent. it was found that pre-aging at temperature 100oC (higher than room temperature)
for 10 min before artificial aging at 160oC for 18 hours will raise the level of true stress-true
strain above that of artificial aging only (peak-aged condition).
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.
Heat treatment
Strain hardening
exponent
Strengthening
coefficient, MPa
Solution treated water-quenched 0.2022 364.5
Natural aging 170 h 0.1659 396.5
Natural aging 1000 h 0.158 376.2
Peak-aged condition 0.0801 395.1
NA(170)+160oCfor 8 h 0.0394 367
NA(1000)+160oCfor 4 h 0.0567 378
PA(100)+ 160oCfor 18 h 0.0576 394.1
0
5
10
15
1 10 100 1000 10000
∆ H
ard
nes
s, H
V
Pre-aging time, min
71
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
0
50
100
150
200
250
300
350
400
0 0.02 0.04 0.06 0.08 0.1 0.12
tru
e st
ress
, M
Pa
true strain
y = 0.0225x + 5.7983
y = 0.0886x + 6.0706
5.6
5.65
5.7
5.75
5.8
5.85
5.9
-8 -7 -6 -5 -4 -3 -2 -1 0
ln (
σ)
ln (ε)
71
Pre-aging of Al-alloy at 100oC after solution treated-water quenched followed by artificial
aging for 18 hours at 160oC (peak aged condition) resulted in higher yield stress than that
obtained from the artificial aging only in the early stage of pre-aging i.e. 10 minutes as shown
in figure 4.28 and figure 4.29. The longer the pre-aging time the smaller is the increase in
yield stress. The higher yield stress obtained from pre-aging plus artificial aging for 18 hours
at 160oC than that obtained by artificial peak aging at 160
oC only may be as a result of the
fine GP zones formed during pre-aging at a higher temperature as 100oC which act as the
nuclei for the β” precipitates during artificial aging as found in Al-Mg-Si alloys [42, 56, 69].
The GP zones size that formed during pre-aging temperature 100oC is larger than critical size
which is necessary for nucleation of β" phase at artificial aging temperature, so only smaller
GP zones are resumed at artificial aging temperature and others may aid as nuclei for the β"
precipitates [43, 48].
A negative effect can be realized when pre-aging temperature is 20oC (natural aging), while a
positive effect appears when pre-aging temperature more than 80oC. The size of needle-like
β” precipitate in subsequent artificial aged alloy is much coarser when pre-ageing temperature
is 20oC, which causes a decrease in peak-hardness. The positive effect occurs again when
natural aging time is longer than 3 weeks [70].
It was found that pre-aging at 100oC result in a beneficial effect on the sub-sequent artificial
age-hardening response as found by Sato [47] who stated that the pre-aging at temperature
more than 70oC the Al-Mg-Si alloys tends to precipitate GP zones which act as nucleation
sites for β" phase while the pre-aging at temperature lower than 70oC the Al-Mg-Si alloys
tends to precipitate of small Si-clusters and Si-Mg co-clusters, which consume the quenched
excess vacancies and the solute atoms and do not act as a nucleation sites for β" phase but
may revert at the temperature for artificial aging resulting in a negative effect on the
mechanical properties of the material.
The increase in the pre-aging time result in a decreasing in strength and this may be due to
the zones formed began to grow consuming solute atoms and excess vacancies [67] that may
affect the precipitation of β" phase during the subsequent artificial aging.
Figure 2.32 illustrate the true stress-true strain curves of solution treated-water quenched and
the artificially peak aged condition and three conditions which were two of them were
naturally aged at room temperature for 170 and 1000 hours followed by artificially aging at
160oC for 8 and 4 hours respectively. The third condition was pre-aging at 100
oC for 10
minutes followed by artificially aging for 18 hours at 160oC. It was found that naturally aging
and pre-aging before final artificial aging raises the level of true stress-true strain above that
of artificially aged condition only as shown.
Natural aging for 170, 1000 hours preceded the artificial aging at 160oC for 8 and 4 hours
respectively reduces the true uniform strain sharply from 17.2% for the artificial peak-aged
condition to about 7.3% and 9% respectively. While pre-aging at 100oC for 10 min before
artificial aging for 15 hours at 160oC has lower true uniform strain than the artificial peak
aging about 11 %.
72
Fig. 2.32 true stress-true strain curves illustrate the effect of natural aging and pre-aging
on artificial peak aging
4.3 Microstructure and XRD Examination
Precipitation sequences rely on super saturation solution treatment, the temperature of aging
and time of aging. In case of the alloy used in the present investigation (6351 Al-Mg-Si
alloy), the transition exists from super saturation namely GP zones → Si rich phase → β''
phase → β' phase → Si precipitates→ β (Mg2Si) stable phase. The transformation sequence
mechanism involves uniformly distributed fine precipitates nucleate on the sites developed
earlier. The hardness measurements can be used as indicator of the development of
precipitation [58].
Figures 4.33 and 4.34 showed the microstructure of the as received specimen that contain α-
Al matrix related to some intermetallic phases such as Al3Fe,Al9Si,α-Al8Fe2Si,β-Al5FeSi and
Mg2Si observed by using XRD analysis have hardness of 74.6 HV and the microstructure of
the as quenching specimen which have hardness higher than the as received condition 82.8
HV. According to mechanical properties, hardness and microstructure examination in an
extruded condition (as received) referred to a temper T54, after solution treatment at 540oC
for 45 min followed by water quenching the microstructure nearly one phase α-matrix and
small intermetallic phase such as β-Al5FeSi insoluble during solution treatment observed by
using XRD technique as shown in figure.
0
50
100
150
200
250
300
350
400
0 0.05 0.1 0.15 0.2 0.25 0.3
tru
e st
ress
, M
Pa
true strain
As Quenched
NA 170h + AA
NA 1000h + AA
Peak-aging
Pre-aging 100 + AA
73
Fig 4.33 Microstructure of the as received specimen at magnification
Figures 4.35, 4.36 and 4.37 showed the microstructure of the 6351 Al-Mg-Si alloy that is
artificially aged at 160oC for 4, 18 and 24 hours respectively after solution treatment followed
by water quenching. Figure 4.32 illustrate the under aging condition where in the early stage
of precipitation involve the separate clustering of silicon and magnesium atoms after that co-
clustering of Si-Mg. After this stage a fine scale zones are formed, referred as GP zones after
this stage a needle like shaped β” precipitate is the predominant and characterizes the under-
aging condition. Slightly larger needles of β” precipitate are formed by increasing the aging
time which is coherent with the matrix which increase in its irregularity in the lattices as
shown in figure 4.33 β‟ precipitates from β” in the aging sequence which it is a rod shaped
and less coherency than the needle shaped precipitate. The formation of this phase causes a
drop in strength and represents the beginning of the over-ageing stage as illustrated in figure
4.34.
74
Fig 4.34 Microstructure of the as quenched specimen (540
oC for 45 min) at
magnification
Fig 4.35 Microstructure of 6351 Al-Mg-Si alloy treated at 540
oC for 45 min then
artificially aged at 160oC for 4 h (under-aging condition)
75
Fig 4.36 Microstructure of 6351 Al-Mg-Si alloy treated at 540
oC for 45 min then
artificially aged at 160oC for 18 h (peak-aging condition)
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)
76
Generally the acceptance sequence of precipitation for Al-Mg-Si alloys is 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) where SSSS represent
supersaturation solid solution, GP zones are Guinir Preston zones. The most effective
hardening precipitate for these alloys is β”, β‟ is metastable hardening precipitate while β is
stable phase [59].
There are three points of view for studying the formation of β” phase. One of them considers
that solid solution is the direct source of β”. On the other hand the second assumption is that
GP zones serve as nuclei for β” phase during aging process. The third one assumes that β”
phase formed from atomic clusters [60].
Precipitate phases of 6351 Al-Mg-Si alloy was analyzed by using XRD technique. Due to a
large specimen of the investigated alloy, large angle X-rays must be used to detect all phases
within the alloy. Results of this analysis are displayed on a graph between beam intensity
(counts) versus angle of incidence of x-rays beam (2θ). Results from XRD analysis of
selected specimen are shown in Figure 4.38 – Figure 4.42 as follow:
Fig. 4.38 showed XRD image of the as received specimen before any treatment. As it is
shown in figure 4.38, Al matrix related to some intermetallic phases such as,Al78Mn22, Al9Si,
α-AlFeMnSi,β-Al5FeSi and Mg2Si. Figure 4.39 demonstrates XRD analysis of the alloy after
solution treatment water-quenched and it shows highly peaks of Al-matrix and some insoluble
phases, the presence of some dissolved phase may be due to the time to spend in solution
treatment not enough and need more time to dissolve it.
Figure 4.40 showed XRD image of 6351 Al-Mg-Si alloy that is artificially aged for 4 hours at
160oC represents (under-aging condition) where XRD investigate possible phase exist in this
case. It showed highly peaks of Al-matrix and some precipitates such as Mg/Si co-clusters,
Al78Mn22, Al9Si, β-Al5FeSi and excess Si phase which cause mild increase in strength of the
alloy. Figure 4.41 illustrate XRD image analysis of 6351 Al-Mg-Si alloy that is artificially
aged for 18 hours at 160oCrepresents (peak-aged condition) and it investigate all possible
phases exist such as highly peaks of Al-matrix, Mg5Si6, Mg5Al2Si4(β” phase), β-Al5FeMnSi
and Al78Mn22. Finally, figure 4.42 demonstrated XRD image analysis of 6351 Al-Mg-Si alloy
that is artificially aged at 160oC for 24 hours (over-aged condition) and it shows highly peaks
of Al-matrix, β-Al5FeSi, Al4.01MnSi0.7 and peaks of Mg2Si (β metastable phase) which have
fcc structural.
For all specimens as shown in all figure 4.38 - figure 4.42, highly aluminum peaks are so
pronounced, because of Aluminum is the main element of the alloy which represent
97.35wt% of the composition of the alloy and also the presence of intermetallic phase β-
Al5FeSi which formed during solidification and it is insoluble during solution treatment. This
is the reason why aluminum and β-Al5FeSi substrate was visible to the incident x-rays beam,
so detector detected high intensity due to aluminum peaks and β-Al5FeSi peaks. Also may be
because of cracks induced while cutting specimens to fit machine‟s stage.
77
Fig. 4.38 XRD analysis of the as received condition
Fig. 4.39 XRD analysis of solution treated water-quenched condition
78
Fig. 4.40 XRD analysis of under-aged condition
Fig. 4.41 XRD analysis of peak-aged condition
79
Fig. 4.42 XRD analysis of over-aged condition
4.4 Scanning Electron Microscope (SEM) with Energy Dispersive
X-rays Analysis (EDAX)
Energy dispersive X-ray Analysis or energy dispersive x-ray spectroscopy (EDAX, EDS, and
EDX) is used for determining surface elemental analysis using the principle that each element
has its own characteristic X-ray emission spectrum. Results of this analysis are displayed on
the screen as a graph between beam intensity (counts) versus energy of detected
spectrum. This spectrum is analyzed using computer software and elements detected are
displayed on the screen. As shown in figure 4.43 the results of SEM observation connected
with EDAX analysis of 6351 Al-alloy after solution treatment water-quenched and it shows
highly peaks of Al-matrix and some insoluble phases, the presence of some dissolved phase
may be due to the time to spend in solution treatment not enough and need more time to
dissolve it. Figures 4.44, 4.45 and 4.46 represent SEM observation connected with EDAX
analysis matches with the chemical composition of the artificially aged specimens after
solution treatment for 4 h (under aged condition), 18 h (peak aged condition) and 24 h (over
aged condition), respectively EDAX analysis performed on particles showed in figures
indicate that those generally contained apart from Al, Si, Mn and significant of Fe, these
particles were identified as Al9Mn3Si, β-Al5FeSi and α-(Al8Fe2Si) also identified by formula
α-Al15(FeMn)3Si Fe rich phase. EDAX analysis also records the precipitation of Mg&Si
metastable phases with different composition till reach the equilibrium phase Mg2Si along
grain boundary.
81
Fig. 4.43 SEM microstructure with EDAX of solution heat treated specimen
81
Fig. 4.44 SEM microstructure with EDAX of under-aged condition
82
Fig. 4.45 SEM microstructure with EDAX of peak-aged condition
83
Fig. 4.46 SEM microstructure with EDAX of over-aged condition
84
4.5 Fracture behavior
Scanning Electron Microscope examines the fracture surface of the investigated alloy 6351 Al
alloy at different conditions. After tensile test the specimens were immersed into alcohol with
magnetic or ultra-sonic stirring in order to remove any contamination conducted after tensile
test. Scanning Electron Microscopy imaging was carried out using electron beam under
accelerating potential difference of 20 KV. Analysis of SEM images was done using image
analysis program called Imagej.
The fracture surface of the solution treated – water quenched condition (540oC for 45 min)
and artificially aged (160oC for various times) is shown in figures 4.47 - 4.50. It was noticed
that the solution treated-water quenched condition showed completely dimple fracture (more
ductile) as shown in figure 4.47. Artificial aging for lower aging time (under-aged condition)
after solution treated-water quenched showed about 50% intergranular fracture with 50%
dimple fracture as shown in figure 4.48. In case of the peak-aged condition the amount of
intergranular fracture increase to about nearly 80% as in figure 4.49. On the other hand, the
over-aged condition as shown in figure 4.50 showed nearly completely ductile (dimpled and
shear fracture).
Fig. 4.47 Fracture surface of solution treated water-quenched condition
85
Fig. 4.48 Fracture surface of under-aged condition
Fig. 4.49 Fracture surface of peak-aged condition
86
Fig. 4.50 Fracture surface of over-aged condition
The existence of intermetallic insoluble phases of AlFeSi and Mg2Si phase (reach to 1μm in
size and more) which are present in the material after solution treatment may be the reason for
the dimple fracture which observed in the solution treatment and all aged conditions. The
dislocation pile up at these particles resulting in the creation of voids at particle-matrix
interface (de-cohesion). The formed voids grow and coalesce leading to transgranular dimple
fracture [40].
Subsequently these particles are present in small fraction in the material the elongation of the
solution treated water-quenched condition is quiet high compared to the aged conditions.
Besides the small effect of unsolvable particles on the total elongation of the over-aged
condition there is another reason for the much drop in elongation and also the cause of dimple
fracture. The reason of this is the formation of voids at particle-matrix interface where
precipitates grow during over aging, and become coarser and dispersed in the matrix [40, 51].
Fracture surface observation by using SEM technique in the samples established that fracture
initiates within voids clusters owing to a sequence of void nucleation at particle-matrix
interface, void grow and finally voids coalescence. The presences of different particles are
visible inside the dimples. XRD measurements have established that these particles are
Mg2Si, β-Al5FeSi and α-(Al8Fe2Si). Large dimples around hard intermetallic α-(Al8Fe2Si) and
β-(Al5FeSi) precipitates, and also slightly finer ones among dispersive hardeningβ-Mg2Si and
α-AlFeMnSi precipitates were formed [62].
87
CHAPTER 5: Conclusions
In the present work, a study of the effect of aging temperature, time, natural aging and pre-
aging on artificial aging behavior in terms of mechanical properties (ultimate tensile strength
UTS, yield stress YS and total elongation), hardness HV and fractography examination of Al-
Mg-Si alloy 6351 was carried out. It was found that Al-Mg-Si alloy have a positive response
for age hardening treatment. There is improvement in hardness, tensile properties and of the
alloy, if the treatment is carried out efficiently. The results can be summarized as:
1. Artificial aging of Al-Mg-Si alloy 6351 in the temperature range 160-260oC after solution
treatment-water quenched showed higher values of ultimate tensile strength UTS, yield stress
YS and hardness HV and decrease in total elongation than that obtained from solution
treatment-water quenched condition only.
2. Higher aging temperature, the peaks of UTS, YS and hardness HV shifted to shorter aging
time i.e. aging at 160oC the peak strength attends after 18 h aging time in comparison to aging
at 260oC the peak strength attends after shorter time 1 h. Raising the aging temperature lower
is the strength of the peak aged specimen with increased toughness compared to aging at
lower temperature. So the best mechanical properties can be observed at lower aging
temperature.
3. It was noticed that the peak value YS increase with increasing the artificial aging
temperature, which reached to the maximum value at 260oC for 1h aging time (294 MPa).
4. Natural aging of Al-Mg-Si alloy 6351 at room temperature (25oC ±5) resulted in a mild
increase in UTS, YS and hardness HV and slight drop in total elongation in compared to
artificial aging. Natural aging for 170 hours and 1000 hours after solution-water quenched
followed by artificial aging at 160oC for various time resulted in higher UTS and YS than that
obtained from artificial aging only (peak- aged condition), as the natural aging time increase
followed by artificial aging, the peak aging of UTS and YS decrease.
5. The time to reach peak aging in case of ( NA+AA) shifted to shorter aging times with
respect to artificial aging only i.e. natural aging for 170 h, 1000 h before artificial aging at
160oC require 8 hours, 4hours respectively to reach the peak strength in compared to that
obtained from artificial aging only (160oC for 18 h).
88
6. Pre-aging of Al-Mg-Si alloy 6351 at temperature higher than room temperature i.e. 100oC
for various time followed by artificial aging at (160oC for 18 h) peak-aged condition result in
higher UTS, YS and HV than artificial peak aging only.
7. As the pre-aging time at 100oC before artificial aging (160
oC for 18 h) increases, the value
of YS, UTS and hardness HV decrease tell reach values smaller than that obtained from
artificial aging only. Pre-aging of Al-Mg-Si alloy 6351 for 10 min followed by artificial aging
at 160oC for 18 h (peak-aged condition) increase the value of UTS, YS and hardness HV and
reasonable elongation than that obtained from artificial aging only.
8. Pre-aging at 100oC for 10 min followed by aging for 18 hours at 160
oC (peak-aged
condition) considered the best condition for modification of 6351 Al-Mg-Si alloy compared to
the artificial aging at 160oC only and to natural aging for 170 hours and 1000 hours followed
by artificial aging for 8 hours, 4 hours at 160oC respectively.
89
REFERRENCES
1. S. Venetski. Silver from clay. Metallurgist,1969,13: p.451–453, Quoted in Fredrik
Aleksander Martinsen, M. Sc. Thesis, Norwegian University of Science and Technology
Department of Physics, 2011, p.5-8.
2. Donald R. Askeland and Pradeep P. Phul´e., 2006 “The Science and Enginering of
Materials. Thompson, 5th
edition”, 2006, p.188-195.
3. J. Banhart, C.S.T. Chang, Z. Liang, N. Wanderka, M.D.H. Lay, and A.J. Hill, 2010,
“Natural aging in Al-Mg-Si alloys - a process of unexpected complexity. Advanced
Engineering Materials, 2010. 12(7): p. 559-571.
4. P.B Hirsch. Electron Microscopy of thin crystals. Butterworths, 1th
edition, 1965, Quoted in
Fredrik Aleksander Martinsen, M. Sc. Thesis, Norwegian University of Science and
Technology Department of Physics, 2011, p.8-12.
5. Ruslan P. Kurta, Volodymyr N. Bugaev, and Alejandro Diaz Ortiz., 2010 “Long-
Wavelength Elastic Interactions in Complex Crystals”. Physical Review Letter, 104,
2010.
6. European Aluminum Association (EAA) and MATTER. http://aluminium.matter.org.uk.
May 2010.
7. A. Kelly, R. B> Nicholson, 1971, "strengthening Methods in Crystals", Elsevier
publishing Ltd. UK, 1971, p.12-14, 38.
8. G. Dieter, 1988 “Mechanical Metallurgy, SI Metric Edition, McGraw–Hill, London,
UK”, 1988.
9. I. J. polmear, 2006“Light alloys from traditional alloys to Nano crystals. Fourth edition,
2006, Butterworth-Heinermann. P.131-160.
10. H. j. McQueen, J. P. Ballon, Idichson, J. J. Jonas and M. G. Akben, 1985, "Strength of
metals and alloys", (icsma7)” 1985, p.1754-1769. Quoted in: M. H. Jacobs, “Precipitation
Hardening”, article issued by EAA-Eurpean Aluminum Association, UK, 1999.
11. A Wilm, Metallurgie, vol 8, p.225, 1911. Quoted in: M. H. Jacobs, “Precipitation
Hardening”, article issued by EAA-European Aluminum Association, UK, 1999.
12. P D Mercia, R G Waltenberg and R Scott, Trans AIMME, Vol. 64, p.41, 1920. Quoted
in: M. H. Jacobs, “Precipitation Hardening”, article issued by EAA-European Aluminum
Association, UK, 1999.
13. P. D. Mercia, Trans AIMME, Vol. 99, p.13, 1932. Quoted in: M. H. Jacobs,
“Precipitation Hardening”, article issued by EAA-European Aluminum Association, UK,
1999.
14. A. Guinier, Nature, vol 142, p.13, 1938. Quoted in: M. H. Jacobs, “Precipitation
Hardening”, article issued by EAA-European Aluminum Association, UK, 1999.
15. G. D. Preston, Proc Roy Soc, vol A167. p.526, 1938. Quoted in: M. H. Jacobs,
“Precipitation Hardening”, article issued by EAA-European Aluminum Association, UK,
1999.
16. J. P. Lynch, L. M. Brown and M H Jacobs, 1999, Acta metall., vol. 30, pp1389 - 1395,
1982. Quoted in: M. H. Jacobs, “Precipitation Hardening”, article issued by EAA-
European Aluminum Association, UK, 1999.
91
17. "properties and selection nonferrous alloys and special-purpose material", ASM hand
book, formerly tenth edition, 1991, Vol. 2, p. 3-15, 32, 37-40, 44-57.
18. M. H. Jacobs and Phil Mag, vol.26, p1, 1972. Quoted in: M. H. Jacobs, “Precipitation
Hardening”, article issued by EAA-European Aluminum Association, UK, 1999.
19. B. Chalmers, “Physical Metallurgy. Cambridge University Press”, 1961.
20. D. A. Porter, K. E. Easterling, and M. Y. Sherif. Phase Transformations in Metals and
Alloys. CRC Press, 3th
edition, 1961.
21. C.D Marioara, S.J Andersen, J. Jansen, and H.W Zandbergen. The influence of
temperature and storage time at RT on nucleation of the β" phase in a 6082 Al-Mg-Si
alloy. Acta Materialia, 51:789–796, 2003.
22. M. Murayama, K Hono, M Saga, M Kikuchi, 1998 “Materials Science and Engineering
A250”, 1998, p.127-132.
23. S. B. Kang, L. Zhen, H. W. Kim and S. T. Lee, 1996 “Material Science Forum”, Vol. 217-
222, 1996, p.827-832.
24. C. D. Marioara, S. J. Andersen, J. Jansen and H. W. Azndbergen, 2003, Acta material.
Vol. 51, 2003, p. 789-796.
25. M. Murayama and K. Hono., 1999 “Pre-Precipitate Clusters and Precipitation Process in
Al-Mg-Si Alloys” Acta Materiala, 47:1537–1548, 1999.
26. C. D. Marioara, S. J. Andersen, J. Jansen, and H. W. Zandbergen, 2001, “Atomic model
for GP zones in a 6082 Al-Mg-Si system”,ActaMaterialia,49/2:321–328, 2001.
27. M. Takeda, F. Ohkubo, T. Shirai, and K. Fukui, 1998, “Stability of metastable phases and
microstructures in the ageing process of Al-Mg-Si ternary alloys” Journal of materials
science, 33:2385–2390, 1998.
28. S.J. Andersen, H.W. Zandbergen, J. Jansen, C. Træholt, U. Tundal, and O. Reiso, 1998,
“The crystal structure of the β" phase in Al-Mg-Si alloys”ActaMaterialia,46:3283–3298,
1998.
29. C. D. Marioara, H. Nordmark, S. J. Andersen, and R. Holmestad, 2006, “Post- β” phases
and their influence on microstructure and hardness in 6xxx Al-Mg-Si alloys” Journal of
materials science, 41:471–478, 2006.
30. G.A. Edwards, K. Stiller, G.L Dunlop, and M.J Couper, 1996, “The Composition of Fine-
Scale precipitates in Al-Mg-Si Alloys” Materials Science Forum, 217-222:713–718, 1996.
31. H˚akon S. Hasting, Anders G. Frøseth, Sigmund J. Andersen, John C. Walmsley Rene
Vissers and, Calin D. Marioara, Fr´ed´eric Danoix, Williams Lefebvre, and Randi
Holmestad, 2009, “Composition of β" precipitates in Al-Mg-Si alloys by atom probe
tomography and first principles calculations”. Journal of Applied Physics, 106, 2009.
32. R. Vissers, J. Jansen M.A. van Huis, H.W. Zandbergen, C.D. Marioara, and S.J.
Andersen, 2007, “The crystal structure of the β' phase in Al-Mg-Si alloys”.
ActaMaterialia,55:3815–3823, 2007.
33. S.J. Andersen, C D. Marioara, A. Frøseth, R. Vissers, and H.W. Zandbergen,
2005,“Crystal structure of the orthorhombic U2-Al4Mg4Si4precipitate in the Al-Mg-Si
alloy system and its relation to the β' and β" phases. Materials Science and Engineering
A”, 390:127–138, 2005.
34. M. H. Jacobs. The structure of the metastable precipitates formed during ageing of an Al-
Mg-Si alloy. Philosophical Magazine, 26:1–13, 1972.
91
35. L. F. Mondolfo, "Aluminum alloys: structures and properties", Butterworths, 1976, p.787-
797.
36. R. C. Dorward and C. Bouvier, 1998,“Materials Science and Engineering A254”, 1998,
p.33-44.
37. G. Albetini, G, Caglioti, F. Fiori and R. Pastorelli, Phisica B276-278, 2000, p.921-922.
38. S. Esmaeili, W. J. poole and D. J. lioyd, 2000, “Material Science Forum” Vol. 331-337,
2000, p.995-1000.
39. L. Zhuang, J. Bottema, P. Kaasenbrood, W. S. Miller and P. De Smet, 2004, “material
Science forum”, Vol. 217-222, 1996, p.487-492.Quoted in O. El-Sayed, M. Sc. Thesis,
Cairo University, 2004.
40. R. A. Siddiqui, H. A. Abdullah, K.R. Al-Belushi, 2000,“Journa of Material Processing
Technology, Vol. 102, 2000, p.234-240.
41. D. M Jiang, B. D. Hong, T.C. lie, D. A. Downham and G. W. Lorimer,2004, “Material
science and Technology”, Vol. 7, 1991, p.1010-1014.Quoted in O. El-Sayed, M. Sc.
Thesis, Cairo University, 2004, p.66-69.
42. G. A Edwards, K. Stiller, G.L. Dunlop and M. J. Couper, 1998, “Acta mater”, Vol. 46,
No. 11, 1998, p.3893-3904.
43. S. P. Ringer and K. Hono, 2000, “Materials Characterizations”, Vol. 44, 2000, p.101-131.
44. L. B. Ber, 2000, “Materials Science and Engineering A280, 2000, p.91-96.
45. M. H. Jacobs, Ph.D. Thesis, "The Nucleation and Growth of Precipitates in Aluminum
alloys", University of Warwick, 1969, Quoted in: M.H. Jacobs, "Precipitation Hardening",
article issued by EAA-European Aluminum Association, UK, 1999.
46. D. W. Pashley, M. H. Jacobs and J. T. Vietz, Phil Mag, Vol.16, 1967, p.51, quoted in: M.
H. Jacobs, "Precipitation Hardening", article issued by EAA-European Aluminum
Association, UK, 1999.
47. Tatsuo sato, 2000,“Material Science Forum” Vol.331-337, 2000, p.85-96.
48. M. Murayama and K. hono, 1999, Acta mater, Vol. 47, No. 5, 1999, p.1537-1548.
49. A. K. Gupta, D. J. Lioyd and S.A court, 2001,“Materials Science and Engineering A316”,
2001, p.11-17.
50. G.A Edwards, K. Stiller, G.L. Dunlop and M. J. Couper, 1996, “Materials Science
Forum” Vol. 217-222, 1996, p.713-718.
51. L. Zhen and S.B Kang, 1997, “Metallurgical and Material Transactions A”, Vol. 28A,
1997, p.1489.
52. H. Tanihata, K. Mastuda and S. Ikeno, 1996, “Material Science Forum”, Vol. 217-222,
1996, p.809-814.
53. J. Dutkiewicz and L. Litynska, 2002,“Materials Science and Engineering A324”, 2002,
p.239-243.
54. A. Perovic, D. D. Provic, G.C. Weatherly and D. J. LIoyd, 2000, “ScriptaMaterialia” Vol.
41 Nov. 7, 1999, p.703-708.
55. W.F. MIAO and D.E. LAUGHLIN, 2000, “Effects of Cu Content and Preaging on
Precipitation Characteristics in Aluminum Alloy 6022, 2000, p.1-10.
56. M. Takeda, F. Ohkubo and T. Shirai, 1998, “Journal of material science”, Vol. 33, 1998 p.
2385-2390.
92
57. AizaJaafar, Azmi Rahmat, Ismail Zainol and Zuhailawati Hussain, 2012, „Effects of
composition on the mechanical properties and microstructural development of dilute 6000
series alloy‟, 2012, p.775-779.
58. I. M. Masoud, T. Mansour and J. A. Jarrah, 2012, „Effect of Heat Treatment on
Microstructure and Hardening Properties of 6061 Al-alloy‟, J. Applied Science Research,
2012, p.5106-5113.
59. M. Nowotink, 2010, „Influence of Chemical Composition and Heat Treatment on
Mechanical Properties and Microstructure of 6xxxAl-alloy‟, 2010, p.98-107.
60. T. V. Christy, N. Nurgan and S. Kumar, 2010, “A Comparative Study on The
Microstructure and Mechanical Properties of 6061 Al-alloy and MMC Al 6061/TiB2/12P,
J. Min. and Mat. Charac, 2010, p.9, 57-65.
61. G. Mrówka- Nowotnik, 2010, “Influence of chemical composition variation and heat
treatment on microstructure and mechanical properties of 6xxx alloys”, 2010, Vol.46, p.
98-112.
62. G. Mrówka-Nowotnik, 2008, “The effect of intermetallics on the fracture mechanism in
AlSi1MgMn alloy”, 2008, Vol. 30, p. 1-8.
63. GOWRISHANKAR M. C., SHRAVAN, RAKESH, RAHUL, ACHUTHA KINI, S. S.
SHARMA, 2014, “Effect of Artificial Aging on Strength and Wear Behavior Solutionized
Aluminum 6061 Alloy, 3rd
World Conference on Applied Sciences, Engineering &
Technology”, 2014, p.1-6.
64. J. W. Marti, 2004, “Micromechanisms in particle-hardened alloys”, Cambridge University
press, 1980, p.37, Quoted in O. El-Sayed, M. Sc. Thesis, Cairo University, 2004, p.66-69.
65. C. S. T. Chang, Z. Q. Liangand J. Banhart, 2010, „Natural Ageing of Al-Mg-Si Alloys‟,
2010, p.1-4.
66. John Banhart, Cynthia Sin Ting Chang, Zeqin Liang, Nelia Wanderka, Matthew D.H.
Lay, Anita J. Hill, 2010, Natural ageing in Al-Mg-Si alloys a process of unexpected
complexity, 2010, p.1-37.
67. L. Zhen and S.B. Kang, 1998, materials letters Vol. 37, 1998, p.349-353.
68. A. Hayoune, 2012, „Thermal Analysis of the Impact of RT Storage Time on the
Strengthening of an Al-Mg-Si Alloy‟, 2012, p.1-7.
69. K. Yamada, T. Sato and A. Kamio, 2000, material science Forum Vol. 331-337, 2000,
p.669-674
70. Lizi He, Haitao Zhang and Jianzhong Cui, 0202, „Effects of Pre-Ageing Treatment on
Subsequent Artificial Ageing Characteristics of an Al-1.01Mg-0.68Si-1.78Cu Alloy‟,
2010, p.1-5.
أ
ةملخص الرسال
( استخدامات وتطبيقات فى مجاالت متعددة مثال ٦٪٫٨لسبائك األلومنيوم ماغنيسيوم سيميكون ) سبيكة ومركبات الفضاء و اليندسة وىياكل الطائرات لذلك تستخدم فى اليندسة االنشائية صناعة السيارات
المدنية وغيرىا من االستخدامات و يرجع ذلك لقوتيا مقارنة بوزنيا الخفيف و مقاومة ىذه السبائك لمتأكل )فى اليواء و الماء والزيوت ومقاومتيا لبعض المواد الكيميائية( و قابميتيا العالية لمتشكيل وقابميتيا لمحام
. المنخفض يضا باالضافو الى سعرىاأوصالبة عالية
الحالى الى تحسين الخواص الميكانيكية لمسبيكة وتحديد أنسب ظروف التعتيق لتحقيق ييدف البحث و دقيقة ٪٩لمدة مئويةدرجو ٩٥٪عند درجة حرارة بالسبيكةابة لجميع االطوار الموجوده ذو تم عمل ا٬ىذا
ثم أجريت المعالجات االتية: تبريدا سريعا تم بعد ذلك تم تبريدىا فى الماء فى أزمنو مختمفة. مئوية٧٫٥الى ٦٫٥تعتيق صناعى فى درجات حرارة من .أ
تعتيق طبيعى عند أزمنو مختمفة. .ب
ألزمنو مختمفة. مئويةدرجة ٦٫٥ساعة ثم تعتيق صناعى عند ٦٥٥٥و ٦٬٥تعتيق طبيعى لمدة .ج
٭٦لمدة مئوية درجة ٦٫٥درجة ألزمنو مختمفة ثم تعتيق صناعى عند ٦٥٥تعتيق مسبق عند .د ساعة.
ودرست العالقة بين ظروف التعتيق المختمفة و الخواص الميكانيكية و بنية الكسر ووجد األتى:
جيادات الشد والخضوع وذلك حتى يصل إوجد أنو بزيادة زمن التعتيق الصناعى يزداد كال من .أ
القيمة يبدأ هلو وتسمى ىذه المرحمة بقمة التعتيق الصناعى و بزيادة زمن التعتيق بعد ىذالى أقصى قيمة جياد إما بعد التعتيق. يصاحب الزيادة فى فيما يعرف بمرحمةجياد الشد والخضوع فى النقصان إكال من
لحالة المعالجو الشد والخضوع الناتجة من التعتيق الصناعى نقص شديد فى الممطولية وذلك بالمقارنة وجد أيضا أن أقصى قيمة ألجيادات الشد والخضوع يمكن أن تصل الييا السبيكة فى درجات .األذابية
الحرارة المنخفضة. لذلك أفضل الخواص الميكانيكية يمكن مالحظتيا فى درجة حرارة التعتيق المنخفضة.
ساعة ووجد أنو بزيادة زمن التعتيق ٦٥٥٥تم عمل تعتيق طبيعى فى أزمنة مختمفو وصمت الى .ب جياد الشد والخضوع يصاحب ىذه الزيادة نقصان طفيف فى إالطبيعى تحدث زيادة بسيطة فى كال من
ذابية. وجد أن التعتيق الطبيعى قبل التعتيق الصناعى عند الممطولية وذلك بالمقارنة لحالة المعالجو اإل
ب
باألضافة الى أنو يقمل الزمن الالزم لموصول الى قمة βر "قد يحسن عممية ترسيب الطو ٦٫٥درجة حرارة .التعتيق مقارنة بالتعتيق الصناعى
٦٫٥دقائق ثم تعتيق صناعى عند ٦٥لمدة ٦٥٥وجد أن بعمل تعتيق مسبق عند درجة حرارة .ج جيادات الشد والخضوع وممطولية جيدة مقارنة بقيم إساعة يحدث زيادة كبيرة فى كال من ٭٦درجة لمدة
التعتيق المسبق أفضل حالة لتحسين الخواص الميكانيكية لسبيكة يعتبر العينات المعتقو صناعيا فقط. مقارنة بالتعتيق الصناعى فقط والتعتيق الطبيعى ثم التعتيق ٦٪٫٨الومنيوم ماغنيسيوم سيميكون
الصناعى.
أحمد يحيى أحمد عبد الرحمن :دسـمهن ٮ٭ٮ٦\٨\٩ تاريخ الميالد:
مصرى الجنسية: ٧٥٦٦\٦٥\٦ تاريخ التسجيل:
..........\....\.... تاريخ المنح: ىندسة الفمزات القسم: العموم ماجستير الدرجة:
براهيمإمحمد ممدوح .د. أ المشرفون: السيد محمود البناد. أ. طاىر أحمد البيطارد. أ.
)الممتحن الخارجي( محمد عبد الوىاب والى أ.د الممتحنون: مركز بحوث و تطوير الفمزات
)الممتحن الداخمي( عبد الحميد أحمد حسين أ.د )المشرف الرئيسي( براهيمإمحمد ممدوح أ.د )عضو( السيد محمود البنا أ.د.
)عضو( طاىر أحمد البيطارد. أ. مركز بحوث و تطوير الفمزات
عنوان الرسالة: بالتعتيق عن طريق المعالجه الحراريةسيليكون -وميماغنيس -الومنيوم ١٥٣٦تحسين الخواص الميكانيكية لسبيكه
الكممات الدالة: الماسح الميكروسكوب اإللكترونى ٬ حيود أشعة إكسالتعتيق الصناعى٬ التعتيق الطبيعى٬ التعتيق المسبق٬
:رسالةممخـص ال
( استخدامات وتطبيقات فى مجاالت متعددة و يرجع ذلك لقوتيا مقارنة ٦٪٫٨لسبائك األلومنيوم ماغنيسيوم سيميكون ) سبيكة أيضا باالضافو الى وصالبة عالية السبائك لمتأكل و قابميتيا العالية لمتشكيل وقابميتيا لمحامبوزنيا الخفيف و مقاومة ىذه
و ٬الحالى الى تحسين الخواص الميكانيكية لمسبيكة وتحديد أنسب ظروف التعتيق لتحقيق ىذاييدف البحث و المنخفض. سعرىادقيقة. بزيادة زمن التعتيق ٪٩درجو مئوية لمدة ٩٥٪ة تم عمل اذابة لجميع االطوار الموجوده بالسبيكة عند درجة حرار
الصناعى يزداد كال من اجيادات الشد والخضوع يصاحبو نقص فى الممطولية ووجد أن أفضل الخواص الميكانيكية يمكن سيط مالحظتيا عند درجات الحراره المنخفضة. وبعمل تعتيق طبيعى يحث زيادة طفيفة فى اجيادات الشد و الخضوع ونقص ب
باألضافة الى أنو يقمل الزمن الالزم βفى الممطوليو. التعتيق الطبيعى قبل التعتيق الصناعى قد يحسن عممية ترسيب الطور "دقائق قبل التعتيق الصناعى يحدث ٦٥وجد أنو بعمل تعتيق مسبق لمدة لموصول الى قمة التعتيق مقارنة بالتعتيق الصناعى.
و يعتبر أفضل جيادات الشد والخضوع وممطولية جيدة مقارنة بقيم العينات المعتقو صناعيا فقط إزيادة كبيرة فى كال من الظروف لتحسين خواص ىذه السبيكة.
ضع صورتك هنا
سيليكون عن طريق -وميماغنيس -الومنيوم ١٥٣٦تحسين الخواص الميكانيكية لسبيكه
بالتعتيق المعالجه الحرارية
عداد إ
أحمد يحيى أحمد عبد الرحمن
القاهرة جامعة – الهندسة كلية إلى مقدمة رسالة
العلوم ماجستير درجة على الحصول متطلبات من كجزء
في
هندسة الفلزات
:يعتمد من لجنة الممتحنين
براهيم المشرف الرئيسىإاألستاذ الدكتور: محمد ممدوح
محمود البنا عضواألستاذ الدكتور: السيد
األستاذ الدكتور: طاهر أحمد البيطار عضو
أستاذ بمركز البحوث وتطوير الفلزات
األستاذ الدكتور: عبد الحميد أحمد حسين الممتحن الداخلي
األستاذ الدكتور: محمد عبد الوهاب والى الممتحن الخارجي
أستاذ بمركز البحوث وتطوير الفلزات
القاهــرة جامعــة - الهندســة كليــة
مصـرالعربيــة جمهوريـة -الجيـزة
٥١٦٣
سيليكون عن طريق -وميماغنيس -الومنيوم ١٥٣٦تحسين الخواص الميكانيكية لسبيكه
بالتعتيق المعالجه الحرارية
عداد إ
أحمد يحيى أحمد عبد الرحمن
القاهرة جامعة – الهندسة كلية إلى مقدمة رسالة
العلومماجستير درجة على الحصول متطلبات من كجزء
في
هندسة الفلزات
شراف إتحت
:األستاذ الدكتور
السيد محمود البنا
:األستاذ الدكتور
محمد ممدوح إبراهيم
:األستاذ الدكتور
طاهر أحمد البيطار
وتطوير الفلزاتأستاذ بمركز البحوث
القاهــرة جامعــة -الهندســة كليــة
مصـرالعربيــة جمهوريـة -الجيـزة
٥١٦٣
سيليكون عن طريق -وميماغنيس -الومنيوم ١٥٣٦تحسين الخواص الميكانيكية لسبيكه
بالتعتيق المعالجه الحرارية
عداد إ
أحمد يحيى أحمد عبد الرحمن
لقاهرةا جامعة – الهندسة كلية إلى مقدمة رسالة
العلوم ماجستير درجة على الحصول متطلبات من كجزء
في
هندسة الفلزات
القاهــرة جامعــة -الهندســة كليــة
مصـرالعربيــة جمهوريـة -الجيـزة
٥١٦٣