MODIFICATION OF MECHANICAL PROPERTIES OF 6351 Al-Mg-Si...

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



By






MASTER OF SCIENCE

In


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)


GIZA, EGYPT

2015



By






MASTER OF SCIENCE

In


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)


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)


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)


Title of Thesis:

MODIFICATION OF MECHANICAL PROPERTIES OF 6351

Al-Mg-Si ALLOY BY AGING HEAT TREATMENT

Key Words:

Artificial Aging; Natural Aging; Pre-aging; XRD; SEM; EDAX

Summary:

The present study is dealing with modification of mechanical properties of Al-Mg-Si alloy

6351 by age hardening. The study investigates the effect of aging temperature, time, natural

aging and pre-aging on artificial aging behavior in terms of mechanical properties and

fractography examination. Artificial aging after solution treatment-water quenched resulted

in a sharp increase in both ultimate tensile strength UTS and yield stress YS, can lead with a

decrease in total elongation. As the time of aging increase the strength increase slightly till

reaches peak strength after that it starts to decrease with increasing time of aging. Better

mechanical properties are observed at lower aging temperature. Natural aging at room

temperature (25 ±3oC) after solution treated-water quenched resulted in a mild increase in

tensile properties with a slight drop in total elongation, natural aging for 170 h and for 1000

h after solution treatment followed by artificial aging of this alloy at 160oC, shifted the time

to reach peak strength to shorter aging time (8- 4 h respectively) in comparison to peak-aged

condition (160oC for 18 h). Pre-aging at 100

oC for various times before artificially aging at

160oC for 18 h was investigated. It was found that the pre-aging for 10 min followed by

artificially peak aging led to slight increase in ultimate tensile strength and yield stress YS

associated with a reasonable total elongation.

ere

mailto:[email protected]

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


for 18 h.

47


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


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


160oC for various times on 0.2% offset yield stress, MPa of Al-

alloy 6351.

60


160oC for various times on ultimate tensile strength, MPa of Al-

alloy 6351.

61


160oC for various times on 0.2% offset yield stress, MPa of Al-

alloy 6351.

61


160oC for various times on hardness, HV of Al-alloy 6351.

62


160oC for various times on hardness, HV of Al- alloy 6351.

62


160oC for various times on total elongation, % of Al-alloy 6351.

63


160oC for various times on total elongation, % of Al-alloy 6351.

63


temperature for 170 h followed by artificial aging for 8 h at 160oC.

65


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


then artificially aged at 160oC for 18 h (peak-aging condition).

75


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


82


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


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


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

, %


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


0.0801

395.1


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


0

5

10

15

20

25

30

35

40

1 10 100 1000 10000

Tota

l E

lon

gati

on

%


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


AA

NA(170 h) + AA

61


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


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


AA

NA(170 h) + AA

20

40

60

80

100

120

140

0 5 10 15 20 25 30

Hard

nes

s, H

V


NA(1000h)+AA

AA

63


for various times on total elongation, % of Al-alloy 6351


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


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

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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



عداد إ


القاهرة جامعة – الهندسة كلية إلى مقدمة رسالة

العلومماجستير درجة على الحصول متطلبات من كجزء

في


شراف إتحت

:األستاذ الدكتور

السيد محمود البنا


محمد ممدوح إبراهيم


طاهر أحمد البيطار

وتطوير الفلزاتأستاذ بمركز البحوث

القاهــرة جامعــة -الهندســة كليــة


٥١٦٣



عداد إ


لقاهرةا جامعة – الهندسة كلية إلى مقدمة رسالة

العلوم ماجستير درجة على الحصول متطلبات من كجزء

في


القاهــرة جامعــة -الهندســة كليــة


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Date post:	03-Jun-2020
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MODIFICATION OF MECHANICAL PROPERTIES OF 6351 Al-Mg-Si...

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