modification of mechanical properties of 6351 al-mg-si … 2021. 1. 2. · fig. 2.14 the phase...

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 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].

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2.4.1 Precipitation Hardening in Al-Mg-Si alloys

2.4.1.1 Pseudo-binary Al-Mg2Si

Al-Mg-Si alloy is a ternary system. Engineering Al-Mg-Si alloys are based on the pseudo-

binary composition Al-Mg2Si %. Fig.2.16. the equilibrium precipitate in the Al-Mg-Si is

Mg2Si which known as a balanced compositions contain magnesium and silicon in same

atomic ratio of 2:1 as the equilibrium precipitate. In terms of Wt%, this translates to the

ratio1.73:1.

Fig. 2.16 pseudo-binary diagram of Al-Mg2Si

2.4.1.2 Phase co-exist and precipitation sequence

For a balanced alloy, the precipitation sequence is specifically as follow:

Embryo clusters →needle-shaped GP zones β” →intermediate β‟→ β (Mg2Si)

The expression of “embryo cluster” is introduced into this sequence. The recent work in this

field by Murayama et al [22] who studied the pre-precipitation stages of Al-0.70Mg-0.33Si

and Al-0.65Mg-0.70Si alloys by using Atom Probe Field Ion Microscopy (APFIM) and High

Transmission Electron Microscopy (HTEM) claim to have detected the separation of Mg and

Si clusters atoms. They were incapable of detect either separate clusters or co-clusters in a

High Resolution Transmission Electron Microscope. The smallest clusters that can be

detected within the TEM are needle like-shaped zones that grow in length and rather more

slowly in diameter, with increasing aging time. They proposed the following precipitation

sequence:

Separate Mg and Si clusters →co-clusters of Mg and Si →small equiaxed precipitation → β''

precipitates → β' precipitates → β (Mg2Si)

The effect of aging treatment on mechanical properties and precipitation behavior in Al-Mg-

Si alloy (0.95%Mg, 1.55%Si and 0.1%Zr) were studied by Kang et al [23]. The results

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 direction

of the aluminum lattice with size nearly ∼ (4x4x50 nm) [28]. The β" precipitate has monoclinic crystal structure with a = 1.516 nm, b =0.405 nm, c = 0.674 and β = 105.3

o as

shown in figure 2.17, and it is ordered relative to the host aluminum lattice in such a way that

(001)Al|| (010)Β", [310]Al||[001] Β" and [230]Al||[100]β". the angle between the β" a-vector and

[010]Al is 33.69oand therefore the angle between the β" c-vector and [100]Al is 18.43

o [28].

d) The β' precipitate

Increasing the aging time or aging temperature, β" phases will start to dissolve or transform

and a new phase will create known as β' [29]. Which is bigger than β" phases and have

dimensions nearly∼ (10x10x500 nm) in compared to∼4x4x50 nm for β'' precipitate. It has a hexagonal unit cell with a = 0.705 nm and c = 0.405 nm, and the latter coinciding with the

4.05 ˚A lattice parameter of fcc aluminum making it fully coherency with the Al. Fig.

2.18 show the hexagonal unit cell of the β' precipitate. The unit cell of β' doesn‟t have a

required orientation in the aluminum (001) plane and may be observed with many different

orientations unlike β" [32].

Fig.2.17 Pictures of the β" precipitate taken with conventional TEM. (a) shows the

original picture, while (b) shows a filtered version. The precipitate eyes can be

seen as small rings, and denote the unit cell centers [28].

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Fig.2.18 Picture of the β' precipitate taken with conventional TEM. The unit cell can be

observed to be hexagonal with lattice parameters a = b = 7.05o A [32].

e) The B', U1 and U2 precipitates

The B‟, U1 and U2 precipitates or also known A, B and C which are coexist with β‟. U1 are

Si-rich and belongs to space group P3m1 which have a hexagonal rod-shaped, semi-coherent

phase which is often found on dislocations, while U2 have orthorhombic with space group

Pnma [33]. Table 2.3 gives more information about their crystal structure and Fig.2.19 shows a

conventional TEM-picture of the B‟-phase.

Fig.2.19 Picture of the B’ precipitate taken with conventional TEM. The precipitate eyes

can be seen as hexagonal rings, and denote the unit cell centers. The unit

cell can be observed to be hexagonal with lattice parameters a= b= 10.4 ˚A

[32].

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 t

modification of mechanical properties of 6351 al-mg-si … 2021. 1. 2. · fig. 2.14 the phase...

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