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Characterization of Alurninum Alloy 2618 and Its Composites
Containing Numina Particles
A Thesis Submitted to the ColIege of Graduate Studies and Research
in Partial Fulfillment of the Requirements
for the Degree of Doctor of Philosophy
in the Department of Mechanical Engineering
University of Saskatchewan
Saskatoon
BY Ikechukwuka N. A. Oguocha
Spring 1999
O Copyright Ikechukwuka N. A. Oguocha, 1998. Ail rights reserved.
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UNIVERSITY OF SASKATCHEWAN
Coilege of Graduate Studies and Research
SUMMARY OF DISSERTATION
Submitted in Partial Fulfdlment
Of the Requirements for
DEGREE OF DOCTOR OF PEEOSOPHY
by
Ikechukwuka Oguocha
Department of Mechanical Engineering
University of Saskatchewan
Summer, 1998
D r . J.A. Gillies
Dr. S. Yannacopoulos
Dr. C. Sargent
Dr. B. Hertz
Dr. T, Rezansoff
Externai Examiner:
Dr. A. K. Gupta
AIcm International Ltd.
Dean's Designate, Chair
College of Graduate Studies and Research
Department of Mechanical Engineering
Department of Mechanical Engineering
Department of Mechanical Engineering
Department of Civil Engineering
Kingston Research & Development Centre
Kingston, Ontario K7L 5L9
PERMISSION TQ USE
Whereas this thesis is submitted in partial fuifihent of the requirements for the degree
of Doctor of Philosophy fiom the University of Saskatchewan, the author has agreed that
the Libraries of this University may make it fieeiy available for inspection. Further, the
author has agreed that permission for copying of this thesis for scholarly purposes may
be granted by the professor who supervised the thesis work reported herein or, in his
absence, by the Head of the Department or the Dean of the College in which the thesis
work was carried out. It is understood that any copying or publication or use of this
thesis or parts thereof for financial gain shall not be allowed without the author's written
permission. Furthemore, it is understood that due recognition shaU be given to the
author and to the University of Saskatchewan in any scholarly use which may be made
of any materid in this thesis.
Requests for permission to copy or make other use of material in this thesis in whole or
part should be addressed to:
Head of the Department of Mechanical Engineering
57 Campus Drive
University of Saskatchewan
Saskatoon, Canada S7N 5A9
Metal matrix composites (MMCs) combine a stiff but brittle phase, typically a cerarnic,
with a more ductile metal ma&. The correct fractional combination of materials c m
resuit in a material with irnproved stiffness, creep resistonce, yield stress, and Wear
resistance relative to the monolithic matrix. The use of MMCs in recent years has
become more widespread due to a growing understanding of the dependence of
composite properties on a number of factors (e.g., interface properties, metailurgy of the
matrix, and stress partitioning between the constituent phases) and appreciation of the
problems that can occur in their usage. The purpose of this work was to investigate
rnicrostnictural evolution in ingot metallurgy AA2618 due to the addition of 10 and 15
vol. % angular dumina (A1203) particles.
The primary investigative techniques employed were microhardness measurements,
differential scanning calorimetry (DSC), scanning electron rnicroscopy (SEM), electron
probe microanalysis (EPMA), and transmission electron rnicroscopy (TEM). In addition,
other metallographic and data analysis techniques were used.
The results of this study showed that the addition of N103 particles did not alter the
aging sequence of AA2618, but it altered certain aspects of the precipitation reaction. It
caused the suppression of Guinier-Preston-Bagaryatskii (GPB) zone nucleation,
acceleration of the artificial aging response, lowering of peak hardness value, and non-
unifonn distribution of precipitate and dispersoid phases. However, it did not affect the
growth mechanisms for S' and 9' formation. The gowth parameters obtained for the
unreinforced ailoy and its composites were not significantly different.
Magnesium accumulation around the reinforcing &O3 particles was very prouounced.
Mg-rich intermetdlic particles (suggested to be MgA1204 spinel) were observed existing
in isolation and embedded in Al,O, particles. The presence of these particles was
considered to be responsible for the low peak hardness obtained for the composites.
Also, other intermetallic particles (such as aluminosilicates and Fe-nch particles) were
observed.
Aluminide (Al,FeNi) particles, which usually occur in AA2618, were detennined to
possess a variety of chernical fomulae. Also, the A1,FeNi phase was determined to be
more consistentiy indexed on the basis a Ctentered monoclinic crystal structure with a
= 0.867 nrn; b = 0.900 nm; c = 0.859 nm; and P = 83.50" rather than the primitive
monoclinic structure reported in the Literature.
1 would like to thank my supervisor, Dr. S. Yannacopoulos, for his direction and
assistance in this research. 1 would like to thank Duralcan Aluniinurn Company, San
Diego (USA) for the supply of the test materials. Also, I am very grateful for the
criticisms received from members of my supervisory cornmittee. Technical assistance
received from Mr. Phi1 Siminoff, Mr. Tom BonIi (Geology Department), Dr. Asern
Hedayat, and Dr. M. C . Chaturvedi and his research group at the MechanicaI and
Industrial Engineering Department, University of Manitoba, Winnipeg, is highly
appreciated. I am grateful to Dr. S. O. Kasap (EE Department, U of S) for his assistance
with thermal analysis and Dr. Yan Jin (Department of Physics, Carnegie Mellon
University, PA, USA.) for his assistance with crystallographic analysis.
1 would like to thank my colleagues, Mojdeh Radjabi, Nathan Gennan, Ehab Shaheen,
Ali Abedian, and Ray Taheri for many useful discussions. Also, my thanks go to Mr.
Chike Odigboh, Mr. Kenechukwu Ezeike, Mr. Lfeanyi Odigboh, Dr. Cornelius
Muojekwu, Dr. Cosmas Oguejifor, Dr. Jude Uzonna, and other fnends 1 made in
Saskatoon for their friendship, cooperation, and continued encouragement. 1 am
especially gratefuI to Dr. and Mrs. Davidson Oguocha, Dr. and Mrs. Raphael Idem, Dr.
and Mrs. 'Diran Fasina, Dr. and Mrs. Emeka Oguejiofor, Dr. and Mrs. Edwin Annze,
Dr. and Mrs. Adebayo Adams, and Mr. and Mrs. Joseph Jobi for their continued support
and advice.
This work was made possible by the financiai support from the NSERC gants to my
supervisor and the funding 1 received from (i) the Canadian Commonwealth Scholarship
and Fellowship Plan, (ii) Department of Mechanical Engineering (Graduate Teaching
Fellowship) and (iii) School of Graduate Studies (Graduate Service Fellowship). 1 am
very grateful.
May God bless you dl.
My Paren&
and
Priincss Nonye/um Oguocha
TABLE OF CONTENTS
PERMISSION TO USE ....................................................................................................... i
... ABSTRACT ....................................... ..... ...........-........................................................... ii
ACKNOWLEDGMENTS .................................................................................................. iv
.................................................................................................................... DEDICATION v
................................................................................................... TABLE OF CONTENTS vi
LIST OF TABLES ............................................................................................................. ix
............................................................................................................ LIST OF FIGURES xi
NOMENCLATURE ........................ ... .............................................................................. xvi
Abbreviations ................................................................................................................ xvi . .
Greek Symbols ............................................................................................................. xvri
. 1 INTRODUCTION ........................................................................................................ 1
........................................................................................... 1.1 Alurninum Alloy 26 18 2
...................................................... 1.2 Particle-Reinforced Metai Matrix Composites 3
1.3 Objectives .............................................................................................................. 4
...................................*.............................................. ........ . 2 LITERATURE REVIEW ... 6
..................................... 2.1 Review of Precipitation Hardening in Alurrîinum Alloys 6
2.1.1 Microstnictural Changes in Al-Cu-Mg (k) AUoys ....................... ... ..... 9
......................................................... 2.2 Fabrication of Particle-Reinforced MMCs 11
........................................................................................ 2.2.1 Powder Metallurgy 11
2.2.2 Compocasting ................................................................................................ 12
2.2.3 SprayFomiing .............................................................................................. 13
............................................................ 2.2.4 XDTM Process (Reactive Processing) 14
........................................ 2.3 Engineering Properties of Particle-Reinforced MMCs 16
......................................................................................................... 2.3.1 Stiffness 17
2.3.2 Elongation ..................................................................................................... 18
2.3.3 S trength ......................................................................................................... 2 1
............................................................................................ 2.3.4 Wear Resistance 23
2.4 The Effect of Reinforcement Particles on Precipitation in
Aluminum Alloys ...................... ..,, ...................................................................... 23
........................ 2.5 Methods Used for Kinetic Analysis of Precipitation Reactions .. 25
................................. .................... 2.5.1 Kinetics of Isothermal Transformations ... 28
2.5.2 Non-Isothermal Analysis .............................................................................. 31
........................................... 3 . MATERIALS AND EXPEFXMENTAL PROCEDURE 35
3.1 Materials ..........................................................~................................................... 35
3.2 Expenmental Techniques .................................................................................... 36
3.2.1 Hardness Measurements .....................~........................................................ 4 0
3 .2.2 Differential Scanning Calorimetry ................................................................ 40
3 .2.3 Transmission Electron Microscop y .......................... .... ................................ 4 1
3.2.4 Scanning Electron Microscopy and Electron Probe Microanalysis .............. 43
3.2.4.1 X-ray Mapping .......................................................................................... 4 5
3.2.4.2 Detennination of Reaction Products ......................... .. ............................... 45
4 . RESULTS AND DISCUSSION ......................................................~........................ 46
4.1 Microhardness ...................................................................................................... 46
4.2 SEM and EPMA Results .................................... .... ............................................. 50 4.2.1 The Nature of Alurninide Particles .............................................................. 50
4.2.2 Depletion of Magnesium in the Composite Matrix ...................................... 60
4.2.3 Other htermetallic Phases ............................................................................ 64
4.3 TEM Results ........................................................................................................ 70
. ......-.---....-............* ......................... 4.3.1 Insoluble Particles ...... ........ 70
.......................................................................................... 4.3.2 Precipitate Phases 76
........................................................... 4.3.3 Crystai S tmcture of Alurninide Phase 88
....................................................................................................... 4.4 DSC Results '02
.................................................................................... 4.4.1 General Description 102
.......... .............. 4.4.2 Determination of Kinetic Parameters for Precipitation ... 108
4.4.2.1 GPB Zone Formation (Peak A) ............................................................... 108
4.4.2.2 GPB Zone Dissolution (Trough B) ...................................................... 120
vii
............ ......*-.-......................................--................. 4.4.2.3 S' and 8' Formation .. 130
5 . CONCLUSIONS AND RECOMMENDATIONS .................................................. 149
5.1 Conclusions .................... ..... .............................................................................. 149
5.1.1 Precipitation Kinetics .................................................................................. 149
5.1.2 Microstructure ............................................................................................. 151
5.2 Recommendations ................................................................. .... .. .. . . . . 153
REFERENCES ............................................................................................................... 155
APPENDIX ..................... .. .... ... ........................................................................... 170
A Alurninum Alloy Designation Systems ............................................................... 170
Al Wrought Nurninurn Alloy Designation System ....................... .... .......... 170
A2 Cast AIuminum Alloy S ystems ...................................................................... 171
B Experimental Apparatus ........................................................................................ 173
B 1 Differential Scanning Calonmetry (DSC) ................................................ 173
B2 . Optical and Transmission Electron Microscop y .......................................... 174
B3 Scanning Electron Microscopy ............. .. .................................................... 175
B4 Energy Dispersive X-ray Spectrometry ............................................................. 176
C Theones of Particle Stren,hening ........................................................................ 177
Cl Yield without Particle Shear ........................................................................... 177
C2 Orowan Theory for Ceramic Reinforcement .......... ........ ......................... 178
viii
LIST OF TABLES
.......................... Table 2.1. Typical properties of some unreinforced alloys (1 3) .... .... 18
Table 2.2. Typical properties of some cornmercially available MMCs (1 3) .................... 19
Table 3.1. Composition of experimental materials ........................................................... 35
. Table 4 . L EPMA point analysis from duminide particles ............................................ 59
Table 4.2. Variation of magnesium content (wt . %) with aging
in 2618i10 MMC ............................................................................................. 64
Taole 4.3. EPMA point andysis of Si-rich particles ......................................................... 66
Table 4.4. Reciprocal lattice parameters of 2-dimensional unit ce11
of the X phase .................... ... ..................................................................... 85
Table 4.5. Measured values of the SADPs shown in Figures 4.27(a)-(0. ........................ 92
Table 4.6. Measured values of the SADPs shown in Fiames 4.28(a)-(0. .................... .... 92
Table 4.7. Calculated and measured values of crystallographic parameters
of Al, FeNi using a = 0.6213 nm; b = 0.6290 nm; c = 0.8557 nm;
............................................................................................... P = 94.76" (133) 92
Table 4.8. Calculated and measured values of angles between the zone axis
of the SADPs shown in Figures 4.27 and 4.28 using a = 0.6213 nm;
b = 0.6290 nm; c = 0.8557nm. P = 94.76O (133) ........................................... 93
Table 4.9. CalcuIated and measured values of crystallographic parameters
of Alx FeNi phase based on C-centered structure using a = 0.867 nm;
................................. ................... b = 0.900nm. c = 0.859nm. P = 83.50° .. 100
Table 4.10. Cdculated and measured values of angles between the zone axis
of the SADPs shown in Figures 4.27 and 4.28 based on C-centered
rnonoclinic structure with a = 0.867 nm; b = 0.900 nm; c = 0.859 nm;
................................................................................................... p = 83.509 100
Table 4.1 1 . Reflection conditions of Al, FeNi phase and possible point
........................................................................................ and space groups 101
................................. Table 4.12. Variation of DSC peak temperature with heating rate 107
........................... Table 4.13. Total heat effects of DSC peaks at different heating rates 107
............................ Table 4.14. Kinetic parametes for GPB zone precipitation reactions 1 18
Table 4.15. Variation of inflection point temperature (Ti) with heating
................................. rate for GPB zone formation ..................................... ., 118
Table 4.16. Variation of inflection point temperature (Ti) with heating
rate for GPB zone dissolution .................................................................. 129
Table 4.17. Kinetic parameters for S ' and 8' fonnation in 26 18 and
2618+15 ....................................................................................................... 135
. ............................................. Table 4.18. Diffusion data for aluminum alioys (Ref 17) 148
................................................ . Table A 1 Wrought alurninum alloy designation system 170
....................................................... . Table A2 Cast aluminum dloy designation system 172
LIST OF FIGURES
Figure 2.1. The Aluminum-rich end of the Al-Cu phase diagram (28,29). ........................ 7
Figure 2.2. Schematic illustration of a spray forrning process for
manufacturing MMCs (60). .... .. ................ . .... .... . ..... . . . . . . . . . . . . . .. 15
Figure 3.1. Schematic of dispesoid particle distribution in the MMCs. ...................-.--... 44
Figure 4.1. Variation of microhardness with aging time at 190 OC. ................................. 47
Figure 4.2. Microstructure of as-received sarnples oE (a) unreinforced 26 18
and (b) 26 18+10 composite. ......... .... ..... . ....... ............ . . . . . . . . . . . 5 1
Figure 4.3. Microstructure of aged samples of: (a) unreinforced 26 18 and
03) 26 l8+lO composite. ...................................... .. ..................................... 52
Figure 4.4. SEM micrograph of alumina particles. ......................................................... 53
Figure 4.5. EDX spectra from (a) an Al,FeNi Particle; and (b) the
surrounding Al matrix. ............................................................................... 54
Figure 4.6. X-ray maps showing (a) iron, @) nickel, (c) copper, and
(d) magnesium in overaged A M 6 18. ......... .... ..............,. ...... ................ .. ........ 56
Figure 4.7. X-ray maps showing (a) iron, (b) nickel, (c) copper, and
(d) magnesium in overaged 26 18+ 10 composite. .. ........... .... .... .. . ........ .. . . ... . .. 57
Figure 4.8. (a) Microstructure of a typicai magnesium-nch particle
(b) the EDX spectra. ..................... .. ........................................ ............ ..... 6 1
Figure 4.9. SEM micrograph showing a typical magnesium-nch particle
embedded in an alumina particle ................................................................. 62
Figure 4.10. (a) SEM micrograph of aged sample of 26 18+10 composite
showing a silicon-rich particle. (b) EDX spectra of a . .
srlicon-rich particle ................................................................................... 65
Figure 4.11. (a) A silicon-nch particle attached to an aiumina particle.
(b) the EDX spectra ................................................................................... 67
Figure 4.12. (a) Microstructure of an intermetallic particle rich in silicon,
magnesium, and iron. (II) the EDX. ...... . . .. . . ..... .... . .... . . . . ... ... ... . . . . . . . . . . . . . . . . . 68
Figure 4.13. (a) Microstructure of an intermetallic particle rich in iron.
(b) the EDX. ................... ......... ..........,, 69
Figure 4.14. (a) TEM micrograph of duminide particles. (b) bright field
TEM image of an aluminide particle at the matrix grain boundary. ............ 7 1
Figure 4.15. TEM micrograph of aged 26 18 showing two merging
Al,FeNi particles. ....... . . .... ...... ..... .... ............. . . . . . . .. . . . . . . . . . . 72
Figure 4.16. TEM micrograph of aged 26 18 showing (a) an aluminosilicate
particle and (b) an aluminosilicate particle lying adjacent to
an Al FeNi particle.. . .... . ..... ...... ...................... .. . . . . . . . . . . . . . . . 73 X
Figure 4.17. EDX spectra from an aluminosilicate particle. ............................................. 74
Figure 4.18. TEM bright field image of aged 26 18 showing S ' distribution:
(a) in the ma&; and (b) at the grain boundary, al1 in the
11 12Jmmx direction. ....... .... .. .. ..-.......... .... .. . .... . - . . ............. . . . . . . . 77
Figure 4.19. Bright field image of artificially aged 26 18+ 15 composite
showing S ' distribution in (a) [ 1 12Imhx and (b) [O0 11,- directions.
(c) and (d) are the corresponding SADPs, respectively. .............................. 78
Figure 4.20. (a) Bright field image of aged 26 I8+lO composite showing S '
distribution in the [1 12]m~x direction; @) corresponding SADP ..... ... .... .. . . . 79
Figure 4.2 1. Bright field image of aged samples showing bulky Cu-rich precipitates
in (a) 2618; (b) 2618+15 composite; (c) and (d) corresponding SADPs
in the 100 1 Imh, direction, respectively. .. . .. . -.. ..-. - -. .. . . . . . ... . . . .. . . . . . .. . . . . .. . . . 80
Figure 4.22. Dark field image of the 8 phase in overaged sarnple of 26 l8+ 15
composite; (b) corresponding SADP in the [112],~, direction ................. .. 82
Fi-gure 4-23. (a) Bnght field image of 8" precipitates in naturdy aged sample
of 26 18. (b) corresponding SADP in the [ 1 12JdX direction. ................... ... 83
Figure 4.24. (a) Bnght field image of X precipitates in naturally aged sample
of 2618+15 composite, (b) corresponding SADP in the [112],,
direction, and (c) indexed pattern of (b) ................................................ 84
Figure 4.25. A sketch of two-dimensional lattice cell of the X phase. ............................. 85
Figure 4.26. (a) Dark field image of the X phase in overaged sample of 26 18+ 15
composite; (b) corresponding SADP in the f 1 12lmaix direction ................... 86
xii
Figure 4.27. SADPs of the Ai, FeNi phase obtained by tilting between [IO01
....................................................................................... and 100 11 zone axes 89
Figure 4.28. SADPs of the Al, FeNi phase obtained by tilting between [3 101
and [O101 zones axes ..................................................................................... 90
Figure 4.29. Computer simulated electron diffraction patterns corresponding
to (a) [100Ii (b) [O0 11, and (c) [O101 zone axes ............................................. 94
Figure 4.30. Crystal structure of AlxFeNi phase based on C-centered mode1 .................. 95
Figure 4.3 1 . Indexed patterns of f 1001, [O0 13 and [O 101 zone axes .................................. 95
Figure 4.32. Computer generated stereographic projections for (a) [100],
(b) [O0 11, and (c) [O 1 O] zone axes ................. ..... ...................................... 96
Figure 4.33. First-order Laue zone (FOU) rings of the AI, FeNi phase for
[110],,, and [130],., zone axes ............................................................ 99
Figure 4.34. DSC thermograms of as-quenched 26 18 for various heating rates ............ 103
Figure 4.35. DSC thermograms of as-quenched 26 18+15 composite for various
heating rates ............................................................................................ 104
Figure 4.36. y vs temperature curves for GPB zone formation at different heating
.................................................................................................. rates (26 18) 109
Figure 4.37. (dy/dT)<P vs temperature curves for GPB zone formation at different
..................................................................................... heating rates (26 18) 110
Figure 4.38. y vs temperature curves for GPB zone formation in 26 18+ 15
composite at different heating rates ............................................................ 111
Figure 4.39. (dy/dT)@ vs temperature curves for GPB zone formation in 26 18+15
............................................................ composite at different heating rates 112
Figure 4.40. Arrhenius plots for detemiination of the activation energy for GPB
...................................... zone formation in 26 18 (based on equation 2.25). 113
Figure 4.41. Arrhenius plots for detemiination of the activation energy for GPB
.............. zone formation in 26 18+15 composite (based on equation 2.25). 114
Figure 4.42. Plots after equation 2.27 for GPB zone formation in AM6 18 and
..................................................................................... 26 18+ 15 composite 116
Figure 4.43. Arrhenius plots after equations 2.19 and 2.26 for the determination
............... offly) for GPB zone formation in 2618 (Q = 10 and 20 'Chin) 117
Figure 4.44. y vs temperature curves for GPB zone dissolution in 26 18 at
different heating rates ................ ...-....-.-. ..-. -.--. .. -. .. -.. ... . . . . . . . . . . . . . . . . . . . . . .. . . . . . 12 1
Figure 4.45. [(dy/dT)@] vs temperature curves for GPB zone dissolution
in 26 18 at different heating rates ................-....-..-....--.....-..-....... .. ......... 122
Figure 4.46. y vs temperature curves for GPB zone dissolution in 26 18+ 15
composite at different heating rates. .. ........ .. ... . ... .. ..-. ... -. .-.. ...-.-. .. .. . . .. . . 123
Figure 4.47. [(dy/dT)@] vs temperature curves for GPB zone dissolution
in 26 l8+ 15 composite at different heating rates. ...........................-.......... .. 124
Figure 4.48. Arrhenius plots after equation 2.25 for determination of the
activation energy for GPB zone dissolution in 26 18. . .. . .. .. .. . ... ..... . .. . ... .... ... 125
Figure 4.49. Arrhenius plots after equation 2.25 for determination of the
activation energy for GPB zone dissolution in 26 18+ 15 composite. .......... 126
Figure 4.50. Plots after equation 2.27 for determination of the activation
energy for GPB zone dissolution in 26 18 and 26 l8+l5 composite. ........... 127
Figure 4.5 1. Arrhenius plots after equations 2.19 and 2.26 for determination
of the activation energy for GPB zone dissolution (<P = 5 "Urnin). .......... 128
Figure 4.52. y vs temperature curves for S ' and 8' formation in 26 18 at different
heating rates. ................................................... ........................................ . 13 1
Figure 4.53. [(dy/dT)@] vs temperature curves for S ' and 0 ' formation in 26 18
at different heating rates. .............................. . . . . ........ .......................... 132
Figure 4.54. y vs temperature curves for S ' and 9' formation in 26 18+ 15 composite
at different heating rates. .. .. . . ... . . ... . .... . .. .... . -. . . . .. . .. . . . . .. . . . . .. . . . . . . . . . . . . . . . . . . . . . 1 3 3
Figure 4.55. [(dy/dT)@] vs temperature curves for S ' and 8 ' formation in
2618+15 composite at different heating rates. ..................... ... ............... 134
Figure 4.56. Schematic illustration of Af detennination. .. .... .. . .. . .. . .. . ... . . . . . . . . . .. . . . . . . 135
Figwe 4.57. y vs temperature curves for S ' formation in 26 18 at different
heating rates. .......................................................................................... 136
Figure 4.58. [(dy/dT)<P] vs temperature curves for S ' formation in 26 18 at
different heating rates. ............ . ...... ..... ... ..... .. . . . ........ . . . . . . .. . 137
Figure 4.59. y vs temperature curves for S ' formation in 26 18+ 15 composite
at different heating rates. ............. ...... .......... .... . . ... ......... .. ..... .... .... . .......... ... . 138
xiv
Figure 4.60. [(dy/dT)@] vs temperature caves for S ' formation in 26 18+ 15
composite at àifferent heaticg rates .......................................................... 139
Figure 4.6 1 . y vs temperature curves for 6' formation in 26 18 at different
heating rates ................................................................................................ 140
Figure 4.62. [(dy/dT)@] vs temperature curves for 8' formation in 36 18
at different heating rates ............................................................................ 141
Figure 4.63. y vs temperature curves for 8' formation in 26 l8+l5 composite
at different heating rates ................................................... .... ............. 142
Figure 4.64. [(dy/dT)@] vs temperature curves for 8 ' foimation in 26 18+15
composite at different heating rates ...................................... ... ................... 143
Figure 4.65. Arrhenius plots after equations 2.19 and 2.26 for determination
of the activation energy for S' phase formation (Qi = 20 "C/min) .............. 144
Figure 4.66. Arrhenius plots after equations 2.19 and 2.26 for determination
of the activation energy for 8' formation (a = 20 "Chin) ........................ 145
Figure B . 1 . Schematic diagran of a typical differential scanning calorimetry .............. 173
Figure B.2. Cornparison of a light and electron microscopes ....................................... .. 174
Figure B.3. Schematic diagram of a typical scanning electron microscope ................... 175
Figure B.4. Schematic diagram of: (a) a typical EDS system @) a typicd
S i c i ) EDS detector ...................................................................................... 176
NOMENCLATURE
a, b, c
A
AJM
b
C
CTE
d
D
Do
DSC
e
E
Ec
E m
EP EPMA
G
GPB
h
(hW
k
ko
Lattice parameters
Area under a reaction peak
Avrarni-Johnson-Mehl
Burger's vector
Speed of Iight
Coefficient of thermal expansion
Interplanar spacing
Interparticle distance
Diffusion coeff~cient
Material constant
Differential scanning calorimetry
Electron charge
Elastic modulus
Activation energy
Electron energy
Energy in keV
Elastic modulus of composite
Elastic modulus of matrix
Elastic modulus of particle
Electron probe microanalysis
S hear modulus
Guinier-Preston-Bagaryatskii
Planck's constant
Miller indices of a plane
Rate constant
Pre-exponential factor
xvi
SEM
SSS
TEP
TEM
UTS
Cuvwl
v m
VP
Y
YS
Cdibration constant
Particle length
Sample mass
Metal matrix composite
Numerical exponent
Integer
Universal gas constant
Dislocation core radius
Particle aspect ratio
A metastable phase in 2ax alloys having an orthorhombic lattice structure
with a = 4.04, b = 9.25, and c = 7.18 A. A stable phase in 2wc alloys having a face-centered orthorhombic lattice
structure with a = 4-00 b = 9.23 A and c = 7.14 A. Its chernical formula is
AJzCuMg.
Scanning electron microscopy
Supersaturated solid solution
Absolute temperature
Line tension
Thermoelectric power
Transmission electron microscopy
Ultimate tensile strength
Direction indices in a crystal
Volume fiaction of matrix
Volume fraction of particle
Mole fraction of the excess solute transformed at time t
Y ield strength
xvii
Greek Symbols
Angle
State variable
Physical property measured during the course of a phase transformation
Enthdpy change
Stress intensity factor range
Wavelength
Radius of curvature
Composite streno&
Matrix strength
Shear stress
A coherent transition pha se in îmx alloys having a tetragon attice with
= b = 4.04 A and c = 7.68 A. Generally, 8" precipitates at Iow tempering
temperatures (T 2 150 OC)
A semicoherent transition phase formed at relatively higher aging
temperatures (or after longer aging times) than the 9'' phase. It has a
tetragonal lattice with a = b = 4.04 A and c = 5.8 A. 8 An incoherent equilibrium phase developed from the 0' phase. It has a - - - -
- - - - - - - - - - - - - - - - - - - - - - - - - - - - -
tetragonal lattice with a = b = 6.07 and c = 4.87.
Angle
Heating rate
Poisson's ratio
xviii
INTRODUCTION
Pure aiuminum obtained from the electrolytic reduction of alumina (A1203) is a
relatively weak material. Therefore, for applications requinng greater mechanical
strength, it is alloyed with metals such as copper, zinc, magnesium and manganese,
usually in combinations of two or more of these elements together with iron and silicon.
Wrought aluminum ailoys are divided into seven major classes according to their major
alloying elements. In the internationdly agreed four-digit systern, the first of the four
digits in the designation indicates the principal alloying element of the alloys within the
group (see Appendix A for more details).
Aluminum ailoys cm be divided into two categories: heat treatable and non-heat
treatable alloys. Heat treatable alloys are those in which strength is developed by
precipitation hardening. Heat treatable alloys are usually found in the Zwx (aluminum-
copper), 6xxx (aiuminum-magnesium-silicon), and 7xrr (alurninurn-zinc-magnesium)
series, although a few such dloys occur in the 4 x 3 ~ (copper-silicon) and 5 x r x
(duminum-magnesium) series. lh non-heat treatable alloys, strength is developed mainly
by solid solution and by strain hardening from coldwork. The non-heat treatable alloys
are in the lxxx (aluminum), 3xw: (aluminum-manganese), 4 m , and S m aluminum
series, although a few such alloys occur in the 7 x x ~ and 8- (alurninum-other elements)
series.
The use of commercial aluminum alloys in both structural and nori-structural
applications has witnessed a significant expansion since the beginning of the century.
Next to iron and steel, aiuminum alloys are the most widely used metallic materials.
High-strength aluminum alloys of the W, 6 m or 7 m series are used in many
applications for their particular combinations of strength and corrosion resistance (1). In
aerospace and automobile industries, the high strength-to-weight ratio of aluminum
alloys (which is particularly important in the design of stmctural components) makes
these alloys a very attractive ciass of materials.
While the moderate strength aluminum alloys such as 6061 are readily weldable and
corrosion resistant, they do not develop sufficient strength for most plate applications.
The higher-strength heat-treatable aiuminum alloys such as 2024, 7075, and 7050 c m be
produced with excellent combinations of strength and corrosion resistance but they are
not weldable by conventional techniques. During solidification of the weld material
grain boundary phases form. The presence of these phases can lead to brittleness and
cracking when the welding process creates stresses. AUoy 2024 is widely used in aircrafi
structures, rivet hardware, truck wheels, screw machine products, and other
miscellaneous structural applications (1). On the other hand, alloys such as 2219 and
7039 with narrow fieezing ranges develop better combinations of strength and
weldability. Alloy 7039 has emerged as one of the most widely used high strength
weldable aluminum alloys, but has very low stress corrosion cracking resistance. Alloy
2219 perhaps possesses the best combinations of strength, weldability, and corrosion
resistance. Typical applications for this dloy are in the manufacture of supersonic
aircraft skin and structural components.
1.1 Aluminum Alloy 2618
Aluminum alloy 2618 (AA2618) is a heat treatable Al-Cu-Mg-Fe-Ni forging alloy
developed for high temperature applications (2,3), especially in the manufacture of
aircraft engine components (3). This alloy has good elevated temperature strength up to
204 OC (2). The addition of small amounts of Fe and Ni produces rnicrostmctural
stability under thermal exposure (2). This alloy derives its strength fiom a combination
of precipitation and dispersion hardening. The main precipitates are coherent Guinier-
Preston-Bagaryatskii (GPB) (Cu, Mg) zones (4) which fonn rapidly on aging at
temperatures up to at least 200 OC (S) , and a semi-coherent S' (AlzCuMg) phase. The S'
phase forms as rods or laths on the (210) ma& planes and nucleates preferentially on
dislocation lines. The precipitation of S' in the matrix is known to be facilitated by
silicon due to its effect on increasing the available concentration of vacancies a d o r the
stability of the pre-existing GPB (Cu, Mg) zones (6). The presence of stable
intennetallic particles (such as aluminide particles of the phase A19FeNi) helps to control
grain size and impede dislocation rnovement (7). Here, balanced arnounts of iron and
nickel are needed, othenvise, these elements combine with some of the copper to f o m
stable compounds which reduce the response of the alloy to age hardening (2,7,8).
Recent studies have shown that AA2618 prepared from ingots has a relatively low
quench sensitivity while that prepared fkom rapidly solidified powders is more quench
sensitive (9).
1.2 Particle-Reinforced Metal Matrk Composites
In the past two decades, a strong interest has been shown in the application of metal
matrix composites (MMCs) in the design of many engineering and non-engineering
cornponents (10,ll). Potential uses of these materials are numerous in industries and
they include such areas of application as aerospace (satellite stmts), defense (electronic
instrument racks), automotive (drive shafts and brake discs), sports goods (golf clubs
and mountain bicycle frames), and marine (yacht fittings). When compared with the
unreinforced matrix alloy, MMCs in general have supenor mechanical properties, for
exarnple, high strength, high stiffbess, high Wear resistance, and very good elevated
temperature properties. These properties can be tailored to meet specific requirements.
The early work on MMCs focused mainly on continuous fiber reinforcement. However,
high cost of fibers, complex fabrication techniques, and limited fabncability restricted
their use to those applications where the end could justify the means. This opened the
way for the development of low cost discontinuously reinforced MMCs, such as
particle-reinforced MMCs (1 2,13).
Alumina particles have become one of the popular reinforcing phases for many
aluminum alloy-based metal matrix composites. They are hard but brittle ceramic
particles with high strength, high modulus of elasticity, and high thermal and electrical
resistance. The size of the particles depends on both the manufacturer and the type of
alloy. However, the mean particle dimension normally lies in the range 2-20 pn (14).
Particle-reinforced MMCs are produced via various routes. They have additional
advantages over the continuous fiber-reinforced MMCs especially since they are low
priced and have both high heat treatrnent ability and processing flexibility. Particle-
reinforced MMCs are now being produced comrnercially. There is evidence in the open
iiterature that the presence of a ceramic reinforcement affects the characteristics of age-
hardenable aiurninum alloys. Changes in quench sensitivity, high dislocation density,
and accelerated aging response have been reported in MMCs( 15-27),
1.3 Objectives
The strength of particle-reinforced MMCs is influenced by the matrix properties, for
example, age-hardening. One of the factors limiting the use of MMCs for engineering
components is a lack of property characterization in relation to the unreinforced alloys.
The lack of data extends from processing parameters to final mechanical properties.
Understanding the factors that influence the physicai and mechanical properties of these
materials is very important in the sense that these properties are sensitive to the type of
reinforcement, the mode of fabrication and the details of post-fabrication processing
(13). Based on its mechanical properties, AA2618 is an attractive matrix material for
MMCs.
To date, the characteristics of 2618 aiuminum alloy reinforced with alumina particles
have not been extensively studied. Recent microhardness measurement and differential
scanning calorimetry @SC) studies have shown that the aging response of alumina
particle-reinforced AM618 composites is accelerated when aged at 200 OC (27). This
was attributed to the presence of large taogles of dislocations at the matrix-particle
interfaces. However, the study was not accompanied by vigorous systematic kinetic
analysis of the precipitation reactions involved in the aging process or followed up with
detailed electron rnicroscopy of the microstructures. Kinetic analysis is capable of
showing how each stage of the precipitation process is affected by the presence of
alumina particles. Detailed electron microscopy investigations are needed to fully
understand any microstructural changes due to the presence of the reinforcing particles.
In view of the foregoing, the main objectives of the present work were:
1. To carry out a systematic investigation of the microstructure of cast AA2618 and
its composites reinforced with alumina particies using metallography and electron
microscopy.
2. To veriv precipitation in AA2618 and aiso investigate intermediate phases.
3. To investigate the influence of dumina-particle reinforcement on the precipitation
behavior of a cast AA26 18 composite.
2.1 Review of Precipitation Hardening in Aluminum Aiioys
Aluminum alloys of the ;?ixr, 6- or 7xxx series are used in many applications because
of their high mechanical strength, acceptable weldability characteristics, and superior
corrosion resistance. Strength is developed via the process of precipitation hardening in
which the alloy is subjected to solution heat treatment, quenching and aging in order to
obtain the optimum combination of mechanical properties.
Solution heat treatment results in the dissolution of soluble phases. As such, the
temperature and time for this treatment (which vary from alloy to alloy) are carefully
chosen. The solubility-temperature relationships can be illustrated by using the &-Cu
system in Figure 2.1 (28,29). At temperatures beIow the cr s o h s , the equilibrium state
consists of two solid phases; namely, a solid solution and an intermetallic compound, 6
(A12W.
The solid solubility of copper in a increases with temperature. When the alloy is heated
above its solvus temperature and held in the a range for a sufficient tirne, the solubility
of copper in aluminum will increase, thus giving the single phase, a. This is the required
solution heat treatment. However, at temperatures above the solidus (the incipient
melting temperature), the solubility of copper in aluminum decreases with increasing
temperature because of the formation of a liquid phase which can have a higher copper
content than the solid. Therefore, the solution temperature is usually chosen to lie
between the solvus and the solidus. The single-phase a structure can be retained at
Atomic Percent Cu
1 2 3 C 5 Weight Pet cent Cu
Figure 2.1. The Aluminum-rich end of the AI-Cu phase diagram (28,29).
arnbient temperatures by cooling rapidly (for example, water quenching) from the a
range to prevent the second phase (0) from forming. The resulting structure is
supersahmted solid solution (SSS) (with respect to the solute, i-e., copper), and hence is
unstable. After quenching from the a region, precipitation is achieved by holding the
alloy below the a solvus at a suitable temperature for a given period of time. This is
caiIed aging. During this time, the 8 precipitates nucleate at locdized high-energy
regions.
In Al-Cu alloys, the precipitation process is not as simple as that described above.
Instead of only 8 precipitating, a succession of metastable precipitates is developed upon
aging SSS. These precipitates develop sequentially during aging at elevated
temperatures. This aging sequence is given in equation (2.1) (29,30):
SSS + G. P. zones + 8" + 8' + B (CUAL) (2- 1)
where 0" and 8' are fine metastable precipitates and 9 is a coarser equilibrium
precipitate.
Strength due to precipitation hardening is produced by the finely dispersed precipitates
that forrn during aging. The mechanism of strengthening involves the formation of
coherent precipitates which causes a great deal of strain because of the mismatch in size
between the precipitates and matrk. The strain energy thus generated can be reduced by
dislocations. When dislocations are anchored or trapped by coherent particles, their
rnovement is obstructed. The Orowan bypassing of the particles strengthens and hardens
the alloy (3 1) (see Appendix C for more details). The characteristic that detemiines
whether a precipitate phase is coherent, semi-coherent, or is the closeness of match or
degree of registry between the atomic spacings on the lattice of the matrix and those on
the precipitate. Most of the heat treatable alurninum alloy systems exhibit multistage
precipitation and undergo accompanying strength changes sirnilar to those of the Al-Cu
system. The structurai changes from the solid solution to the high-strength tempers
involve precipitates that are too srnail to be resolved by the light microscope.
2.1.1 Microstnictural Changes in AI-Cu-Mg (2rrr) Alloys
The ûp series of aluminurn alloys are the most comrnonly used age-hardenable
aluminum alloys. In these alloys, the nature of the precipitating phase that develops
during aging depends on the Cu-Mg ratio. The general sequence of precipitation in Al-
Cu-Mg alloys with CulMg weight ratio 2 2.2: 1 is as follows (4, 32-34):
SSS + GPB zones + SN'+ S '+ S (&CuMg) (2-2)
where SSS and GPB zones are as defined previously and S" and S' are the transition
precursors to the equilibrium S-A12CuMg phase.
By means of hardness measurements and X-ray methods, Hardy (32) and Silcock (4)
have shown that GPB zones form as cylinders in Al-Cu-Mg alloys with Cu to Mg weight
ratio 2.2 to 1. These cylinders are about 10A in diameter and 40 A in length in the eariy
stage of aging, coherent with the matrix on (100) planes, and are followed by the
nucleation of lath-like platelets of the S' phase. Other possible shapes and sizes have
dso been reported for the GPB zones (35,36).
The crystal structure of GPB zones appears not to be fully understood. Alekseyev et al.
(37) suggested that the structure is the same as the equilibnurn S phase. Silcock (4)
reported that GPB zones have face-centered tetragonal lattice (with a = 0.55 nm and c =
0.404 nm). On the b a i s of Silcock's work, strains would be generated in the parent
lattice only by growth normal to the axis of the cylinder since the c dimension of the unit
cell exactly matches that of the aluminurn lattice. On the other hand, Gerold and
Haberkorn (38) have suggested that GPB zones have a tetragonal CuAuI lattice. This is
in contrast with the report by Bagaryatskii (39) which States that GPB zones are clusters
of atoms of 1 Cu, 1 Mg, and severai Al atoms. Another type of zones have been found
and designated GPB (II) zones (4). They are formed oniy at elevated aging temperature
above 240 OC. These zones are larger in size and sharper in shape than GPB zones. Aiso,
they give X-ray diffractions that are superimposed on the broad GPB zones. More recent
X-ray diffraction (40) and high resolution transmission electron rnicroscopy (HRTEM)
studies (34) of these phases have not added significantly to the existing knowledge of
these zones.
The existence of the S" phase is still controversial. The S" phase has been reported to
possess a monoclinic structure (35,41) whereas Bagaryatskii (39) suggested that S" is a
slightly distorted S phase. Alekseev et al. (42) reported that the S" precipitates consist
of bundles of a number of ultra fine rods of 1-2 nm in diarneter each. However, other
authors (4,43) did not observe the S" phase in the alloys they studied. The metastable S '
phase is the predorninant precipitate in the Al-Cu-Mg alloy (44). Wilson and CO-workers
(6,7,43) showed that the S' phase precipitates as rods or laths on the (210) planes in the
cO0b directions of the matrix and nucleates heterogeneously on dislocation lines.
At the present time, the crystal structure of the S' phase is a subject of controversy
(43,4547). It has been shown to have the same crystai structure as the equilibnum S
phase (43). In contrast, Gupta et al (33) reported that it is only a slightly strained version
of the S (M2CuMg) phase. As such, the S' phase need not be distinguished from the
equilibrium S phase and that the former grows as a rod parallel to 400> rnatrix
directions. Nevertheless, although the authors in references (45-47) are in agreement
with other workers (33,43) that the S' phase is orthorhombic, they have reported lattice
parameters that indicate that S0 and S phases have different cell dimensions. In addition,
Yan et al. (47) gave the space group to be Pmm2 while that proposed by Mondolfo (45)
is Cmcm. The S phase is reported to have an orthorhombic crystal structure with lattice
parameters a = 4.00, b = 9.23, and c = 7.14 A (48). The space group is given as Cmcm.
Sen and West (49) have carried out a detailed study regarding the coarsening rate of the
S phase in the temperature range of 260-300 OC.
2.2 Fabrication of Particle-Reinforced MMCs
The physical and mechanical properties of MMCs are sensitive not only to the type of
reinforcernent but to the mode of fabrication and the details of any fabrication
processing of the composite after initial manufacture. Different fabrication methods are
now available for making particle-reinforced MMCs (13,50,5 1). However, the details of
these fabncation routes are proprietary. Basically, there are two generic methods:
namely, solid-state and liquid-state methods. The powder metallurgy route is a typical
example of solid-state processing while spray forming and compocasting are iiquid-state
fabrication routes. The choice of fabrication route is dictated by particle size, particle
shape, particle strength, particle-matrix reactivity, cost of particle and commercial
viability of the chosen route. Some of the fabrication routes are described below.
2.2.1 Powder MetaIlurgy
Early attempts to manufacture particle-reinforced MMCs by incorporating ceramic
particles into metallic melts had lirnited success because most metals do not wet cerarnic
particles. The powder metallurgy route was developed to overcome this difficulty.
According to reference (13), the main notable features of this method are that: (i) high
volume fractions of reinforcement cm be used, thereby maximizing the modulus and
minimizing the coefficient of thermal expansion; (ii) practically any metal alloy c m be
used as the matrix; and (iii) any type of reinforcement is allowed since reaction between
the matrix and reinforcement can be reduced by using solid-state processing.
In the powder metallurgy process, the matrix alloy powder is blended with the
reinforcing particles to achieve a homogeneous mixture. The blending can be carried out
dry or in a liquid suspension. To achieve hornogeneity, the sizes of the metal powders
and the ceramic particles need to be carefully chosen so that agglomerates are not left
after blending. The appropriate size ratio will depend on the blending process used.
Lewandowski et al. (52) have reported that a SiClAl particle size ratio of 0.7:l gives a
more uniform reinforcement distribution than a ratio of 0.3: 1. Typically, the atornized
metd powder is in the size range of 20-40 p and the reinforcing particle sizes are 3-20
pn with aspect ratios < 5: 1. Cold isostatic compaction, canning, degassing, and a high
temperature consoIidation operation (e-g. hot isostatic pressing (HP)) usually fo1low the
blending operation. The fuial wrought product is obtained by extrusion, with an
extrusion ratio of about 20: 1 or higher (13). A high extrusion ratio is required to disrupt
the oxide film between matrix powder particles, ailowing rnetal-to-metd contact and the
development of a good bond between the matrix particles. A high extrusion ratio also
improves the distribution of reinforcement because the plastic flow associated with
extrusion tends to disperse any clusters of reinforcing particles.
The powder metallurgy technique has some shortcomings. One of the main
disadvantages of this approach is that the material handling procedures are cumbersorne
and the fabrication route is relatively cornplex. Hence, the product is expensive relative
to wrought products prepared via conventional casting routes. The brittie cerarnic
particles are susceptible to particle fracture during powder metallurgy processing and
there is also a possibility that the powder route may lead to less interfacial contact
between the reinforcing particles and the matrix (50). Detailed discussions conceming
the various procedures involved in the manufacture of powder metdlurgy MMCs are
given elsewhere (53,54).
2.2.2 Compocasting
The casting of a mixture prepared by stir mixing a liquid metal with solid ceramic
particles is referred to as compocasting. It is the simplest and most economically
attractive method of manufacturing MMCs. In this method, liquid alloy at a temperature
of 30 to 50 OC above the liquidus is agitated vigorously and allowed to cool slowly to the
semisolid range. As the agitation continues, the reinforcing ceramic particles are added
to the siurry. In principle, this can be done using fairly conventional processing
equipment and can be carried out on a continuous or semi-continuous basis. The
composite mixture (with a relatively low viscosity) c m be cast directly into a simple
billet. This is termed a rheocast composite, and the process is known as rheocasting
(55). Alternatively, if the semisolid composite slurry is reheated just above the liquidus
and die-cast into net-shape components, the process is cdled cornpocasting. The
reheating is needed to reduce the viscosity so as to allow the composite sluny to flow
into complex die mdds. Compocasting is now in commercial use for producing Al-Sicp
composites (56). One notable disadvantage of this method is that stir casting invoIves
prolonged liquidkerarnic contact and this can cause excessive interfacial reaction
(56,57). Another difficdty encountered with compocasting and rheocasting is
microstructural inhomogeneity that could be caused by either particle agglomeration and
sedimentation in the melt, gas bubble entrapment, porosity from inadequate liquid
feeding during casting or particle segregation.
2.2.3 Spray Forming
The spray fonning process was originally developed by Singer (58) for unreinforced
alloys and was put into commercial use by Osprey Metals (59). It is now being used to
manufacture MMCs. In the basic form of spray fomiing, a molten strearn of metal is
disintegrated by impingement with a high-pressure inert gas jet to form an atomized
spray of droplets, and the droplet spray is then propelled away from the atomization zone
by fast fiowing gas to deposit on a collector plate interspersed in the spray plume.
Synchronous scanning and withdrawal of the collector plate relative to the molten metal
Stream can be used to manufacture cylindrical, tubular, and slab-shaped billets of spray
formed material.
The potential for adapting the spray forming procedure for manufacturing particle-
reinforced MMCs was recognized at an early stage and has been developed by a number
of metal producers (60,61). This is achieved by hjecting ceramic particles into the spray
(a variant of the basic process known as the spray CO-deposition process). Figure 2.2
(60) is a schematic illustration of a typical apparatus designed for the production of
cylindrical ingots. The deposition rate is in the range 6- 10 Kgmin (23). The ceramic
particles are usually injected into the gas Stream rather than the melt Stream, and do not
n o m d y become appreciably heated in flight. In principle, the ceramic content can be
controlled by relative feed rates of melt and ceramic. However, in practice the overspray
loss rate for the ceramic gets higher as the feed ratio is increased. Therefore, 20-25 vol.
% of particles is the upper limit for successful reinforcement incorporation (50).
The notable advantages of the spray forming methods include: (i) the absence of organic
binders, (ii) the short contact time between the reinforcement and the molten metal, (iii)
the use of fewer processing steps than the powder metallurgy routes, (iv) the ability to
make near-net shapes, (v) the low oxide contents, and (vi) the wide choice of suitable
rnatrix alloys. The obvious drawbacks include difficulties in obtaining homogeneous
distributions of reinforcement and the presence of significant porosity levels. The cost of
MMCs produced by spray codeposition is reported to be intermediate between powder
metallurgy (PM) MMCs and MMCs made b y casting techniques ( 1 3).
2.2.4 XDM Process (Reactive Processing)
The XIYM (exothermic dispersion) process was developed by Martin Marietta
Corporation (62.63) for fabricating in situ composites. It is a rather different approach to
MMC manufacture than the methods mentioned in the previous sections. In this process,
the rnatrix metal is mixed with compounds with which it reacts exothermally. When this
mixture is heated to a high temperature (usually above the melting point of the matrix or
to a point where a self-propagating reaction takes place), the constituent components
react exothermally to form a dispersion of submicroscopic reinforcing particles in the
rnatrix. Hence the name "XD". Since the particles of the reinforcing phase are formed
exo themdy at high temperatures, they tend to be very stable during subsequent
Furnace n
To air
Figure 2.2. Schematic illustration of a spray forming process for manufacturîng MMCs (60).
processïng and use at elevated temperatures. A wide range of ceramic compounds c m be
formed by the X D T M process (63). However, the two that have received wide attention
are T B 2 and Tic. These can be formed by the following reactions
and
There is little information in the open literature regarding (i) the cost of materials
produced by the XDm process, (ii) the porosity levels in the as-reacted materials, and
(iii) the control of the size and spacing of the reinforcing particles. The production of a
variety of MMCs by this process has been reported. These include matrices of Ai, Ti, Fe,
Cu, Pb, and Ni as well as intermetallics such as TiAl, Ti3AI, and NiAl ( 5 1,64).
2.3 Engineering Properties of Particle-Reinforced MMCs
A bnef discussion of some engineering properties of particle-reinforced MMCs is
presented here to show how these properties are influenced by the addition of ceramic
particles to duminum alloys. The enhancement of specific stiffness, specific strength,
Wear and creep resistance, and the reduction of density and thermal expansivity are a few
of the most attractive features of MMCs. Stiffness is a critical design parameter for
many engineering components because the avoidance of excessive elastic deflection in
service is the principal ovemding consideration. A typicai potential application of
improved creep resistance is in the development of high-temperature components, such
as turbine engine parts where the aim is to replace some heavy components with
cornponents made of much lighter substitute matenals.
2.3.1 Stiffness
Elastic modulus is one mechanical property that is always significantly increased by the
addition of ceramic particles into a metallic ailoy. The enhancement of stiffness
achieved by the addition of the reinforcement is retained at high temperatures and this is
of great benefit in the design of rotating parts, support members, and stnictural
bodywork. Examples of applications that depend primarily on stiffness include drive
shafts, electronic instrument racks, bicycle frames, and inertial guidance spheres for
missiles. Tables 2.1 and 2.2 (13) list the properties of a few comrnercially available
unreinforced ailoys and some particle-reinforced MMCs, respectively. It c m be seen
from Table 2.2 that the elastic modulus of a composite increases with the volume
fraction of the reinforcing phase and c m be calculated from the rule of mixtures (ROM)
expression. It should be noted that the ROM is appropriate for estimating the Young's
modulus of continuous reinforcement, but it overestimates that of discontinuous
reinforcement. Therefore, that is why this has been modified in the Halpin-Tsai equation
(65) :
where q = (EP 1 42 - 1)
( E , / E, + 2s)
Ec, E,,,, Ep are the elastic moduli of the composite, matrix, and particle, respectively, s is
the particle aspect ratio, and V, the volume fraction of the particle. The elastic modulus
can also be calculated using the Eshelby equivalent inclusion method (66), and this
approach is also known to be in good agreement with experimental data.
Table 2.1. Typical properties of some unreinforced alloys ( 13).
A-110~ YS* @Pa) UTS (MPa) Elongation (%) E (GPa)
6061 (T6) 275 310 20 69
2014 (T6) 476 524 13 73
2 124 (T6) 325 470 12 72
2618 (T6) 370 470 9 74
7075 (T6) 505 570 10 72
8090 (T6j 415 485 7 80
A356 (T6) 205 280 6 76
A380 (IF) 160 320 3 -5 72
AZ9 1 168 311 2 1 49
AZ6 1 157 198 3 .O 38
* 0.2% offset yield strength.
2.3.2 Elongation
It can be seen from a cornparison of Tables 2.1 and 2.2 that a major limitation in the
engineering properties of particle-reinforced MMCs is the rather low ductility (as
quantified by percent elongation). The tensile elongation decreases with increasing
particle content. Sirnilarly, tensile elongation decreases with increased aging time in heat
treatable alloys (3). The opposite changes in stifhess and ductility with increasing
particle content reflect the interactions between particles and the intervening rnatrix
within MMCs.
Previous work has demonstrated that composite failure is associated with particle
cracking and void formation in the ma& within clusters of particles (13, 67-69). Lloyd
(13) has suggested that particle fracture is more prevalent in coarser particles than in the
finer ones due to the higher probability of finding crack-initiating defects in the former
Composite* Wrought 606 1/A.l20~1 Op (T6) 606 1/A1203/ 15, (T6) 606 1/Aln3/20p (T6) 606 1/AI2O3/2Op (T6) 606 l/SiC/15, (T6) 606 l/SiC/15, (T4) 606 1/SiC/2OP (T4) 606 l/SiC/25, (T4) 20 14/A1203/1 Op (T6) 2014/A1203/15p (T6) 20 14/Al2O3/2OP (T6) 2014/SiC/15, (T6) 2124/SiC/l7.8, (T4) 2 124/SiC/2OP (T4) 2 1 241s iCI25, (T4) 2 168/SiC/12, (T6) 7075/SiC/15, (T6) 8090/SiC/L3, (T4) 809O/SiC/13,(T6) Cas t 356/SiC/10, (Tt5 1) 356/SiC/15, (T6 1) AZ9 l/SiC/9.4, AZ91/SiC/15.lP 380/SiC/1Op 0
Table 2.2. Typical propeaies of some commercially available MMCs (13).
Elongation Supplier
Duralcan, Alcan Duralcan, Aican Durdcan, Alcan Comral85, Comalco Cospray, Alcan DWA$ DWA$ DWAZ Duralcan, Alcan Duralcan, Aican Duralcan, Alcan Cospray, Alcan BP" DWAZ BP" Cospray, Alcan Cospray, Alcan Cospray, Alcan Cospray, Alcan
Duralcan, Aican Duralcan, Aican Dow Dow Durdcan, Aican Duralcan, Alcan
* Composite designation: rnatrix/reinforcement/volume fraction of particles; t 0.2% offset yield stress; $ Composite Specialties Inc., Chatsworth, CA; ' British Petroleum; Young's modulus.
than in the latter particles. The failuse associated with particle clusters is attributed to the
higher stress triaxiality generated in such regions. It has been reported that matrix
deformation between closely spaced elastic particles would be highly constrciined
resulting in local stress levels which are many t k e s the rnatrix flow stress (70). This
behavior has been confmed by continuum modeling (7 l,72). A~so, the larger the
particles are, the more they will be loaded by conventional fiber loading and end loading
mechanisms. For Al-Sicp composites, it has been observed that particle cracking is an
important failure mechanism for composites containing 5 20 p size particles (67).
The geometry of the reinforcement in MMCs has been shown to rnarkedly affect rnatrix
deformation behavior (67,69,73,74). This is largely due to the fact that the matrix stress
and strain fields developed in response to external loads Vary appreciably with the
geometry of the reinforcing phase (72,73,75). This has, in turn, been shown to alter
fracture behavior particularly near the rnatrix/reinforcement interface. Song et al. (76)
recently studied the effects of particle shape on the fracture and ductility of a sphencal
and an angular particle-reinforced 6061 Al composite using scanning electron
microscopy (SEM) and transmission electron microscopy (TEM). It was found that
although the spherical particulate composite showed a slightly lower yield strength and
work hardening rate, the ductility was significantly higher than the anp1a.r counterpart.
The SEM fractographic examination showed that dunng tende loading, the sphencal
composite failed via void nucleation and linking in the matrix near the reinforcement-
matrix interface whereas the angular composite failed through particle fracture and
matrix ligament rupture.
Experimental evidence in the literature shows that voids nucleate preferentially at the
sharp corners of the reinforcements (68,72,77,78). Fisher and Gurland (79) have
discussed the factors that tend to favor the formation of voids. Voids often cause
premature failure of the composite. Recent FEM modeling has predicted that composites
with spherical reinforcements have a higher ductility due to the lower matrix triaxiality
(76,80). Therefore, a feasible way to improve composite ductiiity is to use spherical
reinforcements to reduce stress concentrations and thereby bring about changes in the
stress distribution throughout the composite.
The stress distributions created around and within hard particles in a deforming matrix
have been studied (81,82). The way a particle gatherç stress to itself depends on the
elastic misfit between the two phases. The stress concentrations at sharp corners of the
reinforcements give rise to intense localized plastic 80w (68,78,83). The onset of local
plastic deformation leads at fmt to plastic relaxation, but with further deformation,
localized strain hardening once again leads to high stresses next to the particles (81).
Due to the complexity of the stress fields, dislocation glide, void nucleation, and growth
in the matrix during plastic defomation proceed differently from those comrnonly found
in unreinforced ailoys (84). The particles are known to cany much higher stresses than
the mauix. This ansf fer of stress to the particles, and the associated near-particle
perturbations affect the failure modes, stiffening, and strength observed in MMCs.
2.3.3 Strength
A cornparison of Tables 2.1 and 2.2 shows that the enhancement of 0.2% yield strength
due to the addition of ceramic particles can be quite substantial. McDanels (85) carried
out the first extensive study of the strength of several discontinuous MMCs reinforced
with S i c whisker and particle and reported up to 60% increase in yield and ultimate
tensile strengths. The exact value depended on the volume fraction of reinforcement, the
type of alioy and its temper, and processing of the composite. Although subsequent work
by different authors essentially confirmed these findings, the reported experimental data
show a large degree of scatter due to differences in the material quality, processing
routes, and testing parameters. The fatigue resistance at low AK (stress intensity factor
range) c m be enhanced depending on the test mode employed, although the fracture
toughness and ductility are usually reduced.
Numerous strengthening mechanisms that may operate in particle-reinforced MMCs
have been discussed in the literature and the behavior has been extensively rnodeled
mathematically (72,86-88). The strengthening process in the composite has been
modeled based on two different approaches, namely, the continuum approach and the
micromechanics approach. The continuum shear lag model, originally developed by Cox
(89) and later modified by many workers (90.9 l), gives the composire strength (0,) for a
particulate composite as (9 1):
where O, is the matrix yield stress, V, and V, are the volume fractions of the matrix and
particle respecdvely, and s is the particle aspect ratio. The aspect ratios typically used for
particulate MMCs are in the 1-5:l range. The major difficulty with the continuum
approach lies with its inability to account for the influence of the particles on the
micromechanics of deformation. These include the very high work hardening at Low
strains as well as modification in microstructure such as grain size and dislocation
density.
In the micromechanics approach the microstnicniral effects arising from the presence of
the particles are considered. The possible strengthening mechanisms (the details are
given elsewhere (31,50,86-88)) are: (i) dislocation strengthening due to difference in
coefficient of thermal expansion between the matrix and the reinforcing particles; (ii)
Orowan and dispersion strengthening caused by the resistance of closely spaced hard
particles to the passing of dislocations; (iii) strengthening from grain size refinement;
and (iv) workhardening due to the smin rnisfit between the elastic reinforcing particles
and the plastic matrix. The extent to which the different mechanisms operate will
depend on the microstmcture and processing of the particular composite.
In general, there are relatively few applications where the main attraction of using the
MMCs stems from the greater strength offered, especially at room temperature. While
the presence of ceramic particles improves the modulus at higher temperanires, they do
not add significantly to the high temperature strength. Only a small improvement in
strength over the monolithic alloy is retained at higher temperatures. The reason for this
is that the strengthening mechanisms operating in MMCs at low temperatures are
relaxed at high temperatures. Thus the composite strength is primady controlled by the
high temperature strength of the matnx.
2.3.4 Wear ResÏstance
Although different Wear applications require different reinforcement types tc achieve
optimal Wear rate reduction, there are many situations where Wear rates are reduced by
factors of up to ten by the introduction of the reinforcement. This makes MMCs very
attractive for bearings, bushings, cylinder liners, and break rotors. In some cases, it is
advantageous to control the distribution of reinforcement so as to provide matenal of
high Wear resistance in selected surface areas while other regions are suitably tough,
strong, or thermally conducting. This can be done by selective reinforcement of critical
areas through spray deposition or some other route. In general, it is important for Wear
resistance to be combined with other properties such as high thermal conductivity (to
dissipate frictional heat) and high stifhess ( to avoid Wear frorn excessive deflections).
2.4 The Effect of Reinforcement Particles on Precipitation in Alurninum AIIoys
At present, most duminum ailoys employed as matrices in metal matrix composites are
age-hardenable alloys such as 2124, 2618, 6061 and 7475. Strength increment due to
aging is necessary in these ailoys because it helps to develop acceptable mechanical
properties. When heat-treating age-hardenable MMCs, it is often assumed that the
ma& heat-treats in a manner identical to the unreinforced alloy. Therefore, little
consideration is given to the effect of the reinforcement on the structure and properties
of the heat-treated matrix. However, the current literature shows that the reinforcement
can affect the aging kinetics and, hence, the mechanical properties of MMCs (1 3,15-27).
In MMCs fabncated by molten-metal methods, a ceramic reinforcement c m alter the
aging characteristics of the matrix of a MMC by depleting the precipitate-forming
elements during the fabrication stage. The depletion is caused by chernical reactions
between the matrix alloy and the reinforcement, a phenomenon that has been reported in
many MMCs (13,92-95). For instance, aluminum oxide, A1203, is known to be stable in
pure aluminurn, but reacts with magnesium in Mg-containing Al alloys to form Mg0
and MgA1204 (spinel) as shown in equations (2.7) and (2.8).
and
3Mg + 4iU203 t, 3Mg&04 +2Al (AG O = -2 15.1 W at 1000K) . . (2.8)
Mg0 may form at high magnesium levels and lower temperatures whereas the spinel
will form even at very low magnesium levels (13). Therefore, it is not surprising that
&O3 is not thermodynamically stable in most aluminum dloys.
Unlike SIC, which is stable below the solidus, alumina remains unstable in the
soiid state. Therefore, reactions may still take place dunng a normal solution heat
treatment operation or solid-state processing. Reinforcement reactions have been
proposed to explain the presence of intermetallics such as MgAlIo, spinel in the
matrices of MMCs (92,93). The intermetallic particles are usually high-melting point
phases that do not dissolve during normal solution heat treatrnent. Dislocation
generation is another important effect associated with ceramic phase additions to metal
alloys. When a metal matrix containing a ceramic reinforcement is cooled from the
fabrication or solution heat treatment temperature, large tangles of dislocations are
generated in the matrix alloy due to the difference between the coefficients of thermal
expansion (CTEs) of the ceramic reinforcement and the matrix (96,97). The difference
in C'Es gives rise to stresses usually large enough to deform the matnx plastically (98).
Dislocations are high-energy sites that facilitate the nucleation of strengthening phases
in age-hardenable alloys (19,SO). Transmission electron microscopy (TEM) studies have
shown non-uniform distribution of dislocations in many MMCs (96).
Currently, three major hypotheses have been proposed to explain the small increase in
hardness that occurs when discontinuously reinforced MMCs are aged artificially. First,
it has been attributed to the relatively srnaIl amount of solute atoms available for
precipitation reaction in a reactive matrix (Le., a matrix in which one or more of the
solute elements reacts with the reinforcement phase). In such instances, the amount of
available solutes is found to decrease with increasing volume fraction of the reinforcing
phase (99). Usually, test samples are not cut in such a way as to compensate for
unavailable solutes in the composite materïals. Secondly, reduced hardening has been
associated with acceleration of non-coherent precipitation reactions on dislocations and
other high-energy sites (19,99). Finally, reinforcement-induced segregation and selective
precipitation in the vicinity of the matrix-reinforcement interfaces have been reported in
Al-Si-Mg/SiC, composites (99). Interfacial reaction and elementai segregation reduce
the amount of precipitate-fomiing elements avaiiable for hardening. Therefore, it is cIear
that the knowledge of how the reinforcing phase affects the matrix microstmcture is
important in tenns of designing the composition and processing route of new and
improved MMCs.
2.5 Methods Used for Kinetic Analysis of Precipitation Reactions
The mechanical properties of age-hardenable aluminum alloys are determined largely by
chernical composition, rnanufacturing rnethod, and heat treatment. The effects of alloy
composition and rnanufacturing method on the mechanical properties of these alloys
have been documented in the litmature (2,100). The improvement of mechanical
properties by heat treatment is Iargely amibuted to the formation of metastable phases
during aging. The theones that deal with the nucleation and growth of phases in metallic
doys are covered in the literature (101-103). The pnmary objective of kinetic analysis is
to identiQ the kinetic equations controlling the precipitation process.
To date no attempt has been made to determine the kinetic parameters for the
precipitation processes in AA2618 and its composites reinforced with alumina particles.
There are a number of isothermal and non-isothermal techniques which can be used to
study the kinetics of precipitation processes in metals and alloys. The techniques include
electrical resistivity, differentiai scanriing calorimetry (DSC), hardness measurements,
dilatometry, and thermoelectric power (TEP) measurement. In general, TEP, resistivity,
and hardness measurements are used to study the kinetics under isothermai conditions
while DSC and dilatometry techniques are used for non-isothermal conditions. Youdelis
et al. (104-106) have studied precipitation and dissolution reactions in AL-base alloys
using resistivity, hardness measurements, and DSC. Gupta et al. (107) and Jena et al.
(108) used the DSC technique to determine kinetic parameters in Al-Cu-Mg alloys while
Northwood et. al. (109,110) used TEP to study precipitation processes in AA319 and
AA2024.
The overall precipitation and dissolution process can be described by an Avrami-
Johnson-Mehl (AM)-type equation (1 1 1,i 12). Papazian (1 13) used DSC to study the
kinetics of precipitation in wrought aluminum alloys 2219 and 7075 and determined the
kinetic parameters by trial and error, so as to obtain the best possible fit between the
theoretical and expenmental curves. Donoso (1 14) detennined kinetic parameters for the
dissolution of GP zones in Al-Zn-Mg alloy using the DSC technique. DeIasi and Adler
(1 15) determined the total heat effecrs of 7000 series aluminum alloys and calculated the
activation energies by assuming that dissolution reactions obey first order reaction
kinetics. Mittemeijer and coworkers (1 16- 1 1 a), Meisel and Cote (1 I9), and Starink et al.
(120,121) have used DSC, dilatometry, or both to study precipitation reactions in
different metallic alloys and ernployed different versions of the Kissinger (122) analysis
to calculate the kinetic parameters.
Kinetic parameters have been determined frorn hardness measurements. Arrhenius
analysis of aging kinetics in aluminum-based MMCs has been carried out by Thomas
and King (17), Nieh and Karlak (20), and Song and Baker (2 1). The validity of this
rnethod is based on the fact that Arrhenius andysis uses only the time to peak hardness
and temperature of aging rather than the absolute hardness values. In this technique, it is
assumed that: (i) peak hardness corresponds to the same level of transformation in each
case and (ü) aging temperatures are Iow enough for the transformation to be diffusion
controlled. Hence, the activation energy for diffusion can be calculated from the time to
peak hardness. The Arrhenius equation for diffùsion is given as (17):
where D = difision coefficient, Do = material constant, E = activation energy for
diffusion (J/mol), R = universal gas constant (8.34 J/rnol K) and T = absolute
temperature (K). The Iogarithmic form of equation (2.9) is given as:
When D is replaced with Ut (where t is the time to peak hardness) a plot of
b(f]against UT for a range of temperatures gives a straight line with dope = -UR,
dowing the activation energy, E, to be calculated.
There is yet to be a consensus on the most appropriate mathematical recipe for the
extraction of kinetic data from solid-state transformation experiments. There are
objections in the literature against the use of the Kissinger analysis, which was originally
derived for homogeneous reactions (e-g., gas phase chernical reactions), for the study of
the generally heterogeneous solid-state transformations in metallic d o y s (123). The
three main objections are: (i) thermal gradients are inherent in non-isothermal methods.
Therefore, it is claimed that significant inaccuracy will result fiom the appkation of the
Kissinger analysis which does not aUow for the presence of thermal gradients; (ii) the
order of reaction equation often assumed in the Kissinger method is appropriate for
homogeneous transformations but not valid for the heterogeneous transformations which
take place in solid-state reactions; (iii) the reaction rate theory which is considered
appropnate o d y for isothermal experiments is normally assumed in the Kissinger
rnethod. However, its application to heterogeneous reactions has recently been justified
in references (1 16-1 19). On the other hand, the use of the AJM equation or the simple
order rate theory to describe heterogeneous reactions (solid-state transformations) either
for the isothermal case or the non-isothermal case is questionable because there is yet no
valid theoretical justification for its applicability (1 18).
In d l , the continued use of the activation energy-based analyses for describing
nucleation, growth, and overall transformation kinetics in metals has been strongIy
criticized in the Iiterature (124,125). This is largely due to their tendency to produce
meaningless and misleading results. It has been shown that when nucleation and growth
occur simultaneously, the activation energy obtained from an Arrhenius-type plots for
the rate data is the sum of the positive and negative energy values for nucleation and
diffusional growth mechanisms (124). Nevertheless, due to what seems to be the Iack of
a better quantitative method, experimental determinations of activation energies have
continued to be used for describing the kinetics of phase transformations in metals.
In the present work, DSC was used to study the kinetics of precipitation and dissolution
reactions in the monolithic alloy (2618) and the results were compared with those
obtained for 26 l8+l5 composite. The results were compared with previously published
data. A general overview of kinetic treaûnents pertinent to the present work is presented
in the following two sections.
2.5.1 Kinetics of hothermal Transformations
The theoretical basis for unders tanding the kinetics of a phase transformation
(characterized by constant rates of nucleation and radial growth) was developed by
Avrami (1 11) and Johnson and Mehl (1 12) for diffusion-independent transformation
(e.g., recrystallization). Since then, efforts have been made to derive a more general
kinetic equation with wider applicability (104,107,116- 1 19,126- 128) than the classical
AvrÛmi-Johnson-Mehl (AJM) equation. The modified Awami-Johnson-Mehl (AJM)
equation (104,126,127) used for describing precipitation kinetics in metals is given as
where y is the mole fraction of the excess solute transformed at time t , and n is a
numencd exponent whose value can vary from 0.5 to 2.5 for diffusion-lirnited growth
(101). The exponent (growth parameter) n rnay be an integer or a fraction and its value
depends on the precipitate growth mode (i.e., spheres, rods, or discs) (129) and on the
diffusion rnechanism (e-g., dislocation-assisted diffusion) (130). Provided there is no
change in the nucleation rnechanism, n is independent of temperature. k is the rate
constant and depends on the nucleation and growth rates. As such, it is very sensitive to
temperature. k is defined by
where E is the effective activation energy describing the overail process, while ko, T, and
R denote the pre-exponential factor, the absolute temperature, and the gas constant,
respectively.
The logarithmic form of equation (2.11) is:
which is a straight line function from which n (slope) and k (intercept) are obtained.
From the definition of the rate constant k in equation (2.12),
Hence, the slope of the plot of tnk vs (lm, (-ER), may be used to calculate E. For the
transformation rate,
r r
(dy/dt), it follows fiom equation (2.11) that
In the anaiysis of isothermal transformations using the Kissinger rnethod, the basic
assurnption is that the fraction transformed, y, is fully determined by a state variable, P, (1 16-1 M), that is:
For iso thermal transformation,
where k(T) is a constant. For non-isothermal transformation where T, and thus k(T),
depends on t,
By employing the formalism of equations (2.15), (2.16a) and (2.16b), Mittemeijer et al.
(1 16- 1 18) derived an expression from which the activation energy term of the JMA
equation c m be obtained without recourse to any specific rnodel. This is given as
Thc activation energy can be calculated from the aging times tl and tt which correspond
to two fixed stages of transformation, fi andfi, measured at a number of temperatures.
From the definition of P, k(r, - t , ) = Pf2 - P,, = constant. Therefore, the activation
energy E c m be determined from the slope of the straight line obtained by plotting
tn(tf2 - t r i ) vs I/T. The vdue of c m only be determined if P,, and P f 2 are known,
and this implies adopting a specific kinetic model.
2.5.2 Non-Isothermal Analysis
Experiments are frequently perfonned in a non-isothermal manner. For example, mosr
constant heating rate experiments as obtained in the DSC and dilatornetry are
non-isothermal. In such cases, the application of the AJM expressior. given in equation
(2.1 1) is not straightfonvard. The task of denving an exact solution is formidable.
Brown and Phillpotts (131) have reviewed the usual approach to non-isothermal kinetics
in thermal analysis. In most of these kinetic analyses, the rate of transformation is given
by
where f(y) is a function of y ody and the other terms retain their usud definitions. A
cornparison of equations (2.15) and (2.18) shows that:
Most ofien, it is not possible to obtain the transformed fraction as a function of tirne, but
one c m measure a physical property, 5, (e-,o., hardness, specifk volume/len,&, enthalpy,
electrical resistivity, and rnagnetization) of the metal under study as a function of time
and temperature. Then the degree of transformation, y, can be defined as:
where 5 is the physical property measured during the course of transformation, ci and cf are the values of 6 at the beginning and end of the transformation, respectively. In the
case of isothermal analysis, ci and 4 are constants whereas for that of non-isothermd
analysis, they are not normally considered as constants (1 17,118). In the DSC technique,
the reaction peaks in the thennogram are related to the different solid-state
transformations over the temperature range. The fraction transformed at a given
temperature, y(T), is normally defined as the ratio of the area of a reaction peak up to the
temperature under consideration, A(T), to the total area of the peak, Af. Thus.
and y varies between O and 1. This expression is derived on the assumption that the heat
effect due to the formation of a mole of precipitate is constant. The heating rate term,
dT/dt, is introduced through the relationship
In non-isothermal experiments, the heating rate (a) is usually kept constant, so that
cornbining equations (2.12) and (2.22) yields
Most modem thermal analysis equipment corne with utility programs with which to
calculate y(T) and, sometimes, dy/dT. M e r re-arranging tems, the logarithmic fom of
equation (2.23) is given by
Thus, if at a given heating rate Oj, the mole fraction of precipitates obtained is Yi at
temperature 5, equation (2.24) becomes
where (s) is the rate at the mole fraction y,. The plot of ln Yi
should yield a straight line of dope (-UR) from which the value of the activation energy
(E) of the process can be calculated. This is the recipe for the varying-heating-rate
method. The value of ko can only be determined if f(y) is known, and this irnplies trying
out severai expressions for the f(y) function. This fonns the basis for the single-heating-
rate method (104,107,108) for determining the activation energy in non-isothermal
experiments. From equations (2.18), (2.22), and (2.23), one obtains
The activation energy is then obtained by adjusting the f(y) function using an n value
(see equation (2.19) or the solid-state rate expressions summarized in references
(131,132)) that gives a straight line plot for equation (2.26). Hence, the activation
energy, E, is calculated from the slope (-UR). In this technique, the activation energy for
overall process, Q, is given as Q = nE.
The Kissinger-type expressions can be used to extract activation energies from data
obtained fÏom non-isothermal experiments (1 16-121). This method is best suited for
cases where heat effects are caused by a single precipitation process. For non-isothermal
annealing with a constant heating rate, @, (e-g., DSC experiments) Mittemeijer et al.
(1 16- 1 18) have shown that in good approximation, equation (2.27) showing an
expression between the temperature for a fixed stage of transformation, T and the Y/
heating rate, @, holds:
where p is the state variable fully determinhg that fixed state of transformation, yf , Y f
and the other terms retain their usual meanings. They fbrther showed that, to a very good
approximation, the temperature, Ti, where the reaction rate is maximum (Le., the
temperature corresponding to a point of inflection on the y versus T curve), c m be
substituted for T.,,f. The activation energy is determined from the dope of the straight
line obtained by plotting f n ( ~ ; ' / @ ) vs (I/lri).
3. n/LATERIALS AND EXPERXMENTAL PROCEDURE
The materials and the experimental techniques employed in this study are described in
the following sections.
3.1 Materials
Three materiais were used in the present study. These were: two AA2618 (Al-Cu-Mg-
Fe-Ni) dloys containing 10 and 15-vol. % of durnina (&O3) particles, respectiveIy and
an unreinforced AM618 which was used as a control matenal. The geometry of the
dumina particles was angular, with the longest dimension ranging from 2-20 p. Table
3.1 shows the composition of the materials. Both the metal matrix composites (MMCs)
and the unreinforced alloy were manufactured through a proprietary casting route and
followed by extrusion. Duralcan Alumuium Inc., San Diego, USA, supplied al1 these
rnaterials.
Table 3.1. Composition of expenmental materiais.
Material Element (wt.%)
* Composition in vol. %. Balance = Al. 2618 = AA2618; 2618+10 = 10 vol. % dumina/AA26 18 composite; 261 8+lS = 15 vol. 8 alumindAA26 18 composite.
3.2 Experirnental Techniques
Five principal experimental techniques were used in this study. These are hardness
measurements, differential scanning caiorimetry (DSC), transmission electron
microscopy (TEM), scanning electron microscopy (SEM) and electron probe
microandysis (EPMA).
Hardness measurements generaily give an idea of the aging behavior of metallic dloys
and metd matnx composites at a particular temperature. In this technique a smail
indenter is forced into the surface of a materiai to be tested, under controlled conditions
of load and rate of appiication. The depth or size of the resulting indentation is
measured, which in turn is related to a hardness number, The softer the material, the
larger and deeper the indentation, and the lower the hardness index number.
DSC has become a rapid, inexpensive, and quantitative tool for microstmctural
characterization in aluminum alloys (1 13,115). In its simplest form, and under
appropnate conditions, DSC can show which precipitates are present and their volume
fractions. Also, additional information pertaining to reaction kinetics can be derived
frorn the thermograms (104-110). This kinetic information can be used in many
disciplines, for example, polyrner science and chernical processing, as well as in studies
of solid state reactions. In the traditional DSC technique (see Appendix BI), the sample
and reference are maintained at the same temperature while the furnace temperature is
changed at a constant rate. This requires input of electrical energy to heat either the
sample or the reference, depending on whether the sample undergoes an endothermic
(heat absorbing) or exothennic (heat-emitting) reaction. The DSC curve is thus a plot of
heat flow as a function of temperature. The basic equation of the DSC is given as:
where AH is the enthalpy change, rn is the mass of the sample, A is the area under the
reaction peak, and K is a calibration constant which is proportional to the thermal
conductivity and involves sample geometry. The area under the DSC cume is
proportional to the change in enthalpy associated with the reaction.
TEM is one of the most powerful techniques for obtaining information on the
microchemistry, crystal structure on a microscale, and defect structure of crystalline
materiais. It requires slow and specialized sample preparation. In the TEM technique
(see Appendix B2), a parallel beam of electrons is used to illuminate the area of the
sample to be imaged. The transmitted electrons form an image on a fluorescent screen.
For this purpose, either the directly transmitted electrons (bright-field image) or the
scattered electrons (dark-field image) can be used. The selection is made by using an
aperture in conjunction with tilting the incident eiectron beam to direct the selected
emerging beam unto the axis of the microscope.
Like the TEM, SEM is one of the major charactenzation techniques used routinely in
materials science. In the SEM (see Appendix B3) a focused beam of electrons is rastered
across the sarnple. The backscattered electrons, secondary electrons, or x-rays are used
to form images on a cathode ray tube (CRT). The x-rays are used to provide chemical
analysis. The high-energy backscattered electrons are detected by an annular detector
concentric with the beam. The secondary electrons spiral upward around the axial
magnetic field of the objective lens and are drawn into the secondary electron detector
by an accelerating electric field. An x-ray detector is typically placed perpendicular to
the axis of the microscope at the level of the specimen. When this detector is used the
sarnple is tilted 30" or more.
The identification of chernical elements is of great importance for many aspects of
materials science. Materials emit x-rays when bombarded by high-energy electrons and
the wavelength of the x-radiation depends on the matenal. Electron probe x-ray
microanalysis (EPMA) is an elemental analysis technique based upon bombarding a
sarnple with a focused beam of energetic electrons (beam energy 5-30 keV) to induce
emission of charactenstic x-rays. Spatial distributions of constituents can be visualized
qualitatively by x-ray area scans (dot maps) and quantitatively by digital compositional
niaps.
In addition to gathenng images of materials and rnicrostmctures, the other principai
analyticai capability of the SEM, EPMA or TEM is to gather both qualitative and
quantitative chernical compositional data about a sarnple. The principal way this is done
is by analyzing the x-rays generated when an electron beam interacts with the sample. X-
rays are generated when the pnmary beam ejects an inner shell electron thus exciting the
atom. As an electron fiom the outer sheIl drops in to fil1 the vacancy and de-excite the
atorn it must give off energy. This energy is specific to each individual element in the
periodic table and is also specific to what particular electron dropped in to fil1 the
vacancy. X-rays as photons of electromagnetic radiation have an associated wavelength
related:
where h = Planck's constant, c = speed of Iight, e = electron charge, E = energy in keV,
and A. = wavelength in nm. Based on this relationship, two distinct types of x-ray
detector systems are used. These two types of x-ray detector systems are called Energy-
Dispersive x-ray Spectrometry (EDS) and Wavelength-Dispersive x-ray Spectrometry
(WDS).
EDS spectrometers are most frequently attached to electron column instruments such as
SEM, EPMA, and TEM. As the name implies, EDS is a method of x-ray spectroscopy
by which x-rays emitted from a sarnple are sorted out and analyzed based on the
difference in their energy level. An EDS system (see Appendix B4) consists of a source
of high-energy radiation (often electrons); a sarnple, a solid-state detector (usually from
Lithium-drifted silicon (Si(Li)); and a signal processing electronics. When the sample
atoms are ionized by a highenergy radiation, they emit characteristic x-rays. X-rays that
enter the S i a i ) detector are converted into signals (charge pulses) that can be processed
by the electronics h t o an x-ray energy histogram. This x-ray spectrum comprises a series
of peaks representative of the type and relative amount of each element in the sample.
The number of counts in each peak can be further converted into elemental weight
concentration either by cornparison with standards or standardless calculations. In
general, three principal types of data can be generated using an EDS detector: (i) x-ray
dot maps or images of the sample using elemental distribution as a contras mechanism.
(ii) line scan data or elemental concentration variation across a given region, and (iii)
overall chernicd composition, both quaiitatively and quantitatively.
As the name irnplies, WDS is a detection system by which x-rays emitted from the
sample are sorted out and analyzed based on differing wavelengths (A) in the WDS, or
crystal spectrometer. As in EDS or imaging modes, the beam rasters the sarnple
generating x-rays of which a small portion enters the spectrometer. As the fluorescent x-
rays strike the analyzing crystal, they will either pass through the crystal, be absorbed, be
scattered. or be diffracted. Those which satisw Bragg's Law
(where n = an integer, d = the interplanar spacing of the crystal, 0 = the angle of
incidence, and A. = x-ray wavelength) will be diffracted and detected by a proportional
counter. The signal from this detector is arnplified, converted to standard pulse size in a
single channel andyzer, and counted with a scaler or displayed as rate vs time on a rate
meter. By varying the position of the analyzing crystal, one changes the wavelength that
will satisw Bragg's Law. ïherefore, one can sequentially analyze different elements. By
automating crystai movements one c m dramatically speed up the analysis time.
Typically, WDS analysis is used to gain the sarne type of information that the EDS is
used for: that is, qualitative and quantitative cornpositional information. fine scans, and
dot maps for elemental distribution.
3.2.1 Hardness Measurements
In this experiment, slices measuring about 5 mm in thickness were cut from al1
materials. Ail samples were solution heat treated at 53W5 OC for 2 h and then water
quenched. After quenching, they were aged naturally for 30 days and subsequently aged
artificially at 190 OC for up to 100 h using a constant temperature air-fumace. Dunng
aging, the hardness values of the specimens foilowed the decomposition kinetics of the
matrix alloy. Hardness measurements were can-ied out on polished samples using a
Vickers hardness tester (Buehler Microhardness Tester - Micromet E) with a direct load
of lOOg applied for 15 seconds. The small load was chosen to produce indentations
small enough to occur only in the mauix without touching the alumina particles in the
plane of measurement. To obtain a hardness vaiue, an indentation was made in the
aluminum matrix. The two base diagonals of the ensuing pyramid were measured using
a focal scale on the microscope of the testing machine. Cdculation of the rnicrohardness
value was based on the average length of the diagonals. Diagonals that differed by more
than 5 pm were discarded. Each hardness value was the average of at Ieast ten
measurements. The presence of any subsurface particles and voids was identified by
excessively high or low hardness values, respectively, and these values were also
discarded.
3.2.2 Differential Scanning Calorimetry
Srnall slices were cut from the extrudates from which discs (approximately 5 mm
diameter, 1-1.2 mm thick) were prepared. The discs were solution heat treated at 530 OC
for 2h and water quenched to laboratory water temperature. DSC tests were conducted
on each material in the as-quenched condition using a Mettler TA 4000 thermal analyzer
(TA) equipped with a DSC 20 cell. Heating rates ranging £Yom 5 to 30 OC/min were used
for the DSC experiments. The DSC scans were initiated at 30 "C and completed at 520
OC. The output was in mifiwatts (mW). The net heat flow to the reference pan (Le.,
relative to the sample) was obtained by subtracting the baseline data from the heat data
of the test samples and recorded as a function of temperature. The baseline data were
obtained by scanning an empty reference pan through the same temperature range as the
test samples. The sample and reference pans were constnicted with high punty annealed
aluminum. The utility program supplied with the thermal malyzer was used to calculate
the specific heat capacity and enthalpy of reaction data and the volume fractions of
phases transformed. The specific heat data were normalized for unit mass of the rnatrix
material. The average densities of the matrix dloy and the alumina particles were
determined to be 2.679 and 3.746 gkc, respectively. These values were used to calculate
the mass fraction of the rnatrix available for precipitation reactions dunng DSC scans in
the test rnaterials. At least two samples of each material were scanned at each heating
rate to ensure reproducibility. The results obtained were found to be reproducible. Also,
to avoid room temperature aging, the samples were kept at -10 OC when not in use.
3.2.3 Transmission Electron Microscopy
Transmission electron rnicroscopy is a well established technique used to study the
microstructure of aluminum alloys as well as that of metal matrix composites based on
these alloys. TEM studies have provided information concerning: (i) the nucleation and
distribution of precipitate and dispersoid phases (7,34,43,133,134); (ii) the crystal
structure of phases and the orientational relationships between these phases
(33,135,136); (iii) the influence of manufacturing parameters on the resulting MMC
microstructure (95,134); (iv) the development of MMC microstructure during post-
fabrication thermornechanical treatments (1 8,86); (v) the relationship between
mechanical properties and microstructure (137,138); and (vi) the aging characteristics of
MMCs as cornpared to the unreinforced aUoys (139). There is only limited TEM study
of IM A M 6 18 reinforced with almina particles published in the open literature.
It was pointed out in the introduction (Chapter 1) that AM618 contains insoluble
aluniinide (Fe and Ni-rich) dispersoid particles as a result of the simultaneous additions
of small amounts of Fe and Ni. The aluminide phase was identified as Al9FeNi phase
(1,3,4). Recently, the crystal structure of the AlgFeNi phase in spray-formed A M 6 18 has
been reported to be a primitive rnonoclinic structure with lattice parameters of a =
0.6213 nm, b = 0.6290 nm, c = 0.8557 nm, and P = 94-76" (3). In the present work, it
was observed that the selected area diffraction patterns (SADPs) of alurninide particles
could be indexed consistently using a different rnodel, thus indicating that the particles
may be possessing different crystal structures. n ie aim of the TEM investigation was to
study the microstructure of the MMCs and the unreinforced AA26 18 in order to obtain a
bener understanding of both their aging behavior and nature of the identifiable phases.
Samples of the as-supplied materials were solution heat treated at 530 OC for 2 h, water
quenched, and aged naturally for 40 h. Artificial aging was carried out at 200 OC for
various lengths of time. The microstructures of the aged samples were exarnined in the
microscope. Samples were cut from the aged materials by means of a very thin diamond-
tipped circular cutter and were subsequentiy ground using 600-@t emery paper to a
thickness of approximately 200 p. Disks having 3 mm diameter were cut frorn the 200
pm thick materials and a TEM disk grinder was used to further reduce the thickness to
about 100 p. The disk samples were then prepared by both the ion-rnilling technique
and the conventional electro-poiishing method. The ion-milled samples were thinned in
two stages. In the first stage, the disks were dimpled to a thickness of about 25 pm in a
GATAN 656 dimpler grinder using 2-4 pm diamond paste. while in the second stage the
dimpled disks were thinned to perforation using a GATAN 600 dual ion-rniller equipped
with a cold stage and cooled continuously with liquid nitrogen. An accelerating voltage
of 6 V, a total gun curent of 1 rnA (0.5 mA/gun) and an incident ion-beam angle of 13 O
were used.
Electropolishing was dcne in a solution that contained one part (volume) of 90 70
cornmercially pure nitric acid (HNO3) and three parts of methanol (CH30H). The
solution was maintained in the temperature range -50 to -15 OC at a voltage of 20 V. All
the thin foils were examined in a JEOL JEM 2000FX transmission electron microscope
at an accelerating voltage of 150-160 kV, using a combination of bright field image
(BFl) and dark field image (DFI) techniques as well as selected area diffraction pattern
(SADP) technique. Chemical analysis of phases was determined by the X-ray EDS
method using standardless metallurgical thin film (SMTF) program. Tilting of the
phases fiom one orientation to another was carried out in the selected area diffraction
mode by systematically tilting the phase (Le., by tilting the rnounted TEM sampie) about
two rnutually perpendicular axes (say X and Y) using a double-tilt holder. The operation
mostly started from a high-symmetry SADP, then tilted dong a given reciprocal axis
until another high-syrnmetry SADP was reached. The SADPs were analyzed by
program developed by Jin (140) and Sutliït (14 1).
3.2.4 Scanning Electron Microscopy and Electron Probe Microanalysis
Samples in various heat treated conditions were investigated using SEM and EPMA or a
combination of both. Sample preparations for both SEM and EPMA investigations were
in most cases identical. The samples were polished to a high smoothness, mounted on
the specimen stage of the testing apparatus, and coated lightly with carbon. The carbon
coating was necessary to provide surface electrical conductivity for the alumina
particles. Electrical conductivity is needed to maintain beam stability.
In a particular investigation, the distribution and composition of the various phases
present in the samples were studied using a JEOL JXA-8600 elecuon probe
microanalyzer. The samples were solution heat treated at 530 OC for 2 h, water
quenched, and aged naturdly for 40 h d e r which they were aged artificiaily at 200 OC
for up to 1990 h. They were tested in the overaged condition in order to enable the
secondary phases attain a size comparable to the probe diameter of about 5 p. By using
the elernental information in Table 3.1, the following activities were canied out: (i)
calibration of standard elements and setting-up of a standards table; (ii) preparation of an
element table, and (iii) preparation of a point table of several spots. However, there was
no standard element for rnanganese (Mn). Spot analysis of each recognizable phase was
canied using both energy dispersive spectrometry (EDS) and wavelength dispersive
spectrometry (WDS). A large number of secondary particles (dispersoids) were observed
in the materials as shown schematically in Figure 3.1. The size and distribution of the
dispersoids were not unifom. The particles were generally of two types; narnely, those
that appeared very bnght and those that were gray-white (medium bright) under
backscattered electron imaging (BEI). A preliminary EDS analysis of these particles
showed that they consisted of Al, Fe, Ni, Mg and Cu. The particles varied between 0.05
and 2.5 p in length, but were more homogeneously distributed in the unreinforced
alloy. A quantitative WDS scheme was used to determine the elemental concentration in
the particles.
O Dispersoid particle
Figure 3.1. Schematic of dispesoid particle distribution in the MMCs.
A close examination of the alumina particles showed that each had a gray ring around a
dark inner domain as show in Figure 3.1. This ring was considered to be the interface
between the matrix and the reinforcement or a reaction layer around the alumina particle.
Spot analysis was also carried out on selected rings and the dark domains.
3.2.4.1 X-ray Mapping
The distributions of elemental Al, Mg, Fe, O, Si, Cu and Ni in the test materials were
determined by x-ray mapping technique using a JEOL KA-840 scanning microanalyzer
equipped with a Si(Li) detector and 3 prn thick norvar window. The sarnples were in the
same aged condition as those used for compositional analysis. The P P m program
developed by Tracor Northern was used for the image analysis.
3.2.4.2 Determination of Reaction Products
The need to separate the reinforcements from the matrix arose in the course of the
present investigation. The particles were exarnined for reaction products and the
presence of other insoluble/impurity particles. Small pieces of the MMC materials were
dissolved in Aqua Regia (AR) inside a fume charnber and in a water bath. The AR used
was consisted of 1 part (volume) concentrated HN03, 3 parts concentrated HC1 and 1
part water. Afier dissolving the samples, the solution was filtered and the filtrate was
rinsed with water before drying. The dry particles were examined in the EPMA.
4. IU?,SULTS AND DISCUSSION
The purpose of the microhardness measurement experiment was to document the effect
of alumina particle reinforcement on the aging response of a ingot metallurgy (IM) 26 18
composite. Changes in microhardness of the metal matrix composite (MMC) matrix
have been monitored as functions of aging time and reinforcement volume fraction. The
variation of microhardness with aging time at 190 OC for the control alloy (26 18) and the
two composites (2618+10 and 2618+15) is shown in Figure 4.1. The figure shows that
the samples followed the decomposition kinetics of the matrix alioy. For each sample,
the hardness increased initially to a maximum with increasing aging time after which it
decreased with further increase in aging. In the naturally aged condition (the O aging
time in Fig. 4. l), the composite matrices are observed to be harder than the unreinforced
alloy. This is an indication of an increased dislocation density in the composite matrices
due to a coefficient of thermal expansion (CTE) mismatch between the alumina particles
and the unreinforced alloy .
A close examination of the cuves shows that while an appreciable difference in
hardness exists between the unreinforced alloy and the 26 18+ 10 composite (both in the
naturally and artificia.y aged conditions), the difference in hardness between the
2618i10 and 2618-t-15 MMCs is s m d . The results of this study are consistent with
those reported by Dutta et al. (23,139) for AA6061-N203 and AA2014-&O3
composite systems. It is suggested that the dislocation density does not increase with
O 20 40 60 80 100 120
Aging T h e (h)
Figure 4.1. Variation of rnicrohardness with aging t h e at 190 OC.
additional N203 particles. In the case of the study involving AA6061-Al& composites,
increasing the alumina content from O to 10-vol. % increased the mean dislocation
density by about two orders of magnitude (from 3.1 x 10'* to 4.5 x 10" m2). -4 further
Increase in alumina content fiom 10 to 15 vol. % led to a relatively little additional
increase (from 4.5 x 10" to 7.3 x 1012 m-*).
Figure 4.1 shows that the aging kinetics of the two MMCs is accelerated at 190 OC (Le.,
shorter times to peak hardness) as cornpared to that for the unreinforced alloy. The time
to reach peak hardness in the two MMCs is about 14-16 h, whereas it is 19-22 h for the
unreinforced alloy. Also, it can be deduced from the figure that relatively little
acceleration is observed on increasing the alumina volume fraction from 10 to 15 %.
Accelerated aging has been reported in many discontinuously reinforced MMCs ( 15,17-
27,139,142- 144) and is attributed to the increased ease with which precipitate nucleation
and growth take place due to increased dislocation density in the MMC matrix (17-
27,139,142-144). However, it should be noted that this could be attributed to the
fabrication route of the MMC. Papazian (15) investigated the influence of S i c whiskers
and particles on the aging kinetics of alurninurn matrix composites fabncated by both
powder metdlurgy (PM) and ingot casting techniques. It was observed that the aging
kinetics of MMCs fabricated via the PM route were more enhanced than those of IM
MMCs. The PM process results in a fmer grains and incorporation of oxide particles. It
is important to note that although precipitation in IM alloys is slower than in PM
processed alloys, they have also been shown to exhibit accelerated aging (19,27). In
addition, Chawla et al. (19) have reported that the aging kinetics of composites are close
to those of the unreinforced alloys at low aging temperatures but greatly accelerated at
high aging temperatures. The authors compared the aging characteristics of a Sic
particle-reinforced AA2014 composite with its unreinforced counterpart. The results
show that, at 180 and 195 OC, the level of peak hardness achieved in the composite was
lower than that in the unreinforced alloy. On the other hand, the results obtained at low
aging temperatures (120 and 165 OC) were opposite.
Although aging is accelerated in the 2618+10 and 2618+15 composites, the degree of
hardening (as indicated by the level of peak hardness achieved during aging at 190 "C) is
less in the composites than in the unreinforced alloy. Hence, the overall gain in the
mittrix strength obtained through heat treatment decreases with reinforcement addition.
This is consistent with the results reported in the literature ( 19,139,144,145). The
addition of Galumina whiskers to AA6061 resulted in a decrease in the peak hardness of
the composite (145). This was attributed to the absorption of quenched-in vacancies
needed for solute migration at the reinforcement-matriu interfaces. Also, the addition of
&O3 particles to AM014 led to a reduction in the peak matrix hardness (139). The
peak hardness values obtained for sarnples aged at 185 OC were approximately 1480
MPa, 1250 MPa, and 12 15 MPa for the monolithic dloy and the two composites (10 and
15 vol. % A1203), respectively. The decrease was attributed to the depletion of solute
elements (which, in mm, resulted in a reduction in the arnount of h' precipitates formed
in the composites).
The low level of hardening obtained in the present composites contrasts with most of the
results obtained to date in the literature (20,2 l,23,24,72,142, W ) , which report that peak
hardness is higher in the composites than in the unreinforced. It is important to note that
apart from the authors in reference (23) who studied AA6061-A1203 composites, the
other investigators in references (20,2 1,24,72,142,143) studied durninum alloys
reinforced with S i c reinforcements. A1203, is known to react with magnesiuni in Mg-
containing Al alloys to form Mg0 and MgA1204 spinel (13,92-95). Also, Lloyd has
reported that S i c is stable below the solidus whereas N203 remains unstable in the solid
state (13). Therefore, in addition to reactions that may take place during fabrication,
interface reactions c m still take place during a normal solution heat treatment operation
or solid-state processing.
4.2 SEM and EPMA Results
4.2.1 The Nature of Aluniinide Particles
The backscattered eiectron @SE) micrographs in Figures 4.2 (a) and (b) show the
microstructure of typical as-received samples of unreinforced 26 L 8 and the 26 18+ 10
composite, respectively. Figures 4.3 (a) and (b) are the BSE micrographs of solution heat
treated and artificially aged samples of the monolithic alloy and the composite,
respectively. The bright (white) particles in both figures are mostly insoluble (FeNi)-
containing dispersoids. The duIl bright rings around the alumina particles (see Figure
4.3(b)) are suggested to represent the matrix-reinforcement interface or a reaction layer.
The alumina particles are angular, irregularly shaped, and of different sizes as shown in
Figure 4.4. Figures 4.5 (a) and (b) show the energy dispersive spectrornetry (EDS)
spectra of a typical aluminide particle and the surrounding matrix, respectively. In Figure
4.5 (a) the prominent Fe and Ni peaks show that the bright particles indeed contain these
elements. Typically, aluminide particles form during solidification of the alloy but they
are known to transform or othenvise be modified during subsequent solution heat
treatment ancUor fabrication.
A cornparison of Figures 4.2 with Figure 4.3 shows that solution heat treatment and
subsequent aging have a substantid effect on the morphology of the aluminide particles.
It can also be discemed from the figures that the distribution of alurninide particles is
fairly uniform in the unreinforced dloy, while a somewhat preferential distribution can
be observed in the composites as these particles appear to congregate around the
reinforcing alumina particles.
Figure 4.2. Microstructure of as-received samples of: (a) unreinforced 2618 and (b)
26 18+lO composite.
Figure 4.3. Microstructure of aged samples of: (a) unreinforced 26 18 and @) 26 18+ 10
composite.
Figure 4.4. SEM rnicrograph of alurnina particles.
5 3
Figure 4.5. EDX spectra from (a) an Al,FeNi Particle; and @) the surrounding Al
matrix.
One of the important factors that determine the level of hardening obtained in age-
hardenïng alloys is the availability of the solute elements to form precipitate phases. The
distribution of the main alloying elements was examined using a E O L m - 8 4 0 SEM.
Figures 4.6 (a)-(d) are X-ray maps showing the distribution of iron, nickel, copper and
magnesium in overaged 2618, respectively. The distribution of iron is similar to that of
nickel. This corroborates the EDS results, which show that the dispersoid particles are
rich in iron and nickel.
Figures 4.7 (a)-(d) show the distribution of iron, nickel, copper and magnesium in an
overaged sample of 2618+10 composite. The distribution of magnesium and copper is
practically uniform in the mauix, although the areas occupied by the aluminide particles
are deficient in magnesium (except on the fringes). On the other hand, the presence of
copper in these areas indicates that the durninide particles contain a small amount of
copper. However, the amounts of copper were considered not sufficient to anticipate a
new quaternary Al-Cu-Fe-Ni phase. In Figure 4.7 (d), a copious amount of magnesiurn
can be seen around the alurnina particles. Accumulation of magnesium around the
dumina particles (whether due to segregation, interfacial reaction, or stress-assisted
diffision) results in the net amount of free magnesium being smaller in the composite
matrix (the reinforcement-free area where hardness readings are taken) than in the
unreinforced alloy. Although the details of the mechanisms through which solute
depletion occurs in the composite ma& are not fully known, the lower level of
hardening observed in the composites (see Figure 4.1) can be attributed to magnesiurn
depletion from the matrix. This observation is consistent with the results reponed by
Ribes et al. (138) and Dutta et al. (139).
Figure 4.6. X-ray maps showing (a) iron, (b) nickel, (c) copper, and (d) magnesium in
overaged AA26 18.
Figure 4.7. X-ray maps showing (a) iron, @) nickel. (c) copper, and (d) magnesium in
overaged 26 18+10 composite.
It has been reported (7,133) that aluminide particles have a chernical formula of
MgFeNi. Recently, Underhi11 et al- (8) reported the existence of A.i7Cu4Ni particles in
spray-formed AA26 18. Zhang and Cantor (133) further identified the MgFeNi particles
as a monoclinic phase with lattice parameters of a = 0.6213 nm, b = 0.629 nm, c =
0.8557 nm, and P = 94.76". However, a quantitative wavelength dispersive spectrometry
O S ) analysis of representative particles (see Table 4.1) did not conform to the
MgFeNi formula. The fabrication route and post-fabrication processes used to produce
the test materials could be responsibie for the differences in the structural formula (8).
Since the structure formulas obtained for the alurninide phase in the present study are
different from those reported in the literature (7, 133), it has been decided to use the
general phase formula of A1,FeNi in the present study.
Although the x value varies as shown in Table 4.1, the atomic ratio of iron to nickel
remains practically consistent at unity wherever they occur simultaneously. Thus, the
aluminum content seems to dictate the stmcture formula of these particles and this may
alter the crystallography. Further, although the elements found in the aluminide phase
were similar to those reported by other investigators (7,133), some of the selected area
diffraction patterns (SADPs) obtained in this study could not be indexed consistently on
the basis of the lattice parameters suggested previously (133). That is, the Al,FeNi phase
being a primitive monoclinic unit ce11 with lattice parameters a = 0.6213 nm, b = 0.629
nm, c = 0.8557 nm, and P = 94.76O. Rather, analysis o f SADPs obtained in this study
shows that they can be indexed more consistently on the basis of the structure of the
A1,FeNi phase being C-centered monoclinic with a = 0.8673 nm; b = 0.9000 nm; c =
0.8591 nm; and P = 83.504".
Table 4.1 EPMA point andysis from duminide particles.
Particle Element -wt O/o Likely Phase?
# Si Mn Fe Cu Mg Ni Al Total
Al 1 1 FeNi
A1 FeNi
AI1 IFeNi
Al 1 FeNi
Ali FeNi
AI 12FeNi
AI ?FeNi
Al12FeNi
Al 1 zFeNi
A1 1 2FeNi
Ai I3FeNi
AI 3FeNi
A1 13FeNi
Al 13FeNi
AI14FeNi
Al IZeNi
Al IsFeNi
AlzlFeNi
Al l&u2Fe
4.2.2 Depletion of Magnesium in the Composite Matrix
It has been mentioned previously (Section 4.1) diat the level of hardening achieved in
the composite matrix during artificial aging at 190 OC is less than that obtained for the
unreinforced alloy. it is reasonable to attribute this to the smaller amount of precipitate-
fonning elements in the composite matrix as compared to the monolithic ailoy. In the
present study, no attempt was made to have equal rnatrix content in the test materials,
and the microhardness data were not normalized for the unit mass of the matrix material.
In the industry, hardness samples are rarely cut to have equai matrix mass and the
hardness data are seldom norrnaiized. Although the unequal mass of matrix postulation
is tenable to some extent, the preferential segregation of magnesium to the matrix-
reinforcement interface shown in Figure 4.7(d) demands that a new explanation be found
for the lower level of hardening obtained in composites. Segregation of rnagnesium to
matrix-reinforcement interface has been reported (92,93,95). Le Petitcorps et al. (92)
and Lloyd et al. (95) have reported that MgA1204 particles are present at the matrix-
reinforcement interfaces due to reaction between elemental magnesium and the alumina
reinforcements.
The presence of reaction products and other insoluble phases was investigated by
exarnining in the EPMA the dry residues obtained from digesting sarnples of the
composite materials in Aqua Regia (AR). Isolated magnesium-rich particles (suspected
to be MgA1204 spinel or magnesium oxide (MgO)) were observed. Figure 4.8(a) shows
the BSE image of a typical particle while Figure 4.8@) shows the energy dispersive X-
ray (EDX) spectra. There are instances where the magnesium-rich particles are
embedded in alumina particles (see Figure 4.9).
According to Lloyd (13) and Eiibes et al. (93), Mg0 is formed in samples with high
magnesium content whereas MgA1204 spinel c m form at very low magnesium levels.
This suggests that Mg0 would not be the likely phase to forrn from the relatively small
Figure 4.8. (a) Microstructure of a typical magnesium-rich particle @) the EDX spectra.
Figure 4.9. SEM micrograph showing a typical magnesium-nch particle embedded in an
alurnina particle.
amounts of magnesium contained in the samples studied here (ir, Table 3.1, W.% Mg c
2 for al1 matenals). In absence of other supporting evidence (e-g., selected area
diffraction patterns (SADPs) of the Mg-rich crystals), it is therefore concluded that the
magnesium-rich particles formed on the surface of or embedded in the alurnina particles
are MgA1204 spinel.
Also, EDX spectra of supposedly clean alumina particles show strong magnesium peaks.
Only scarcely did the EDX spot analysis show alumina particles with insignificant or no
magnesium peaks. Elementai magnesium is readily soluble in Aqua Regia. Hence, the
presence of isolated and embedded MgA1204 spinel and the appearance of strong
magnesium peaks in the EDX spectra of clean alumina particles indicate that certain
level of chemical reaction has taken place during the fabrication of the composites. The
formation of MgA1204 spinel depletes magnesiurn in the matrix which, in mm, can
influence the aging response and mechanical properties of the composites.
The strong magnesium peaks associated with the EDX spectra of clean alumina particles
could have e s e n from the diffused species. EPMA investigation was can-ied out to
snidy the variation of magnesium content with aging in the free matrix and the ring
region (the matrix-reinforcement interface region as shown in Figure 3.1). Table 4.2
shows the results obtained for 2618+10 composite. At Ieast two readings were taken for
each test condition to check reproducibility. The data in Table 4.2 show that there are
changes in the magnesium content within the composite matrix and at the interface. The
increase in magnesium content at the interface can be attributed to two sources. One is
magnesium fluorescence emanating from precipitate phases such as equilibriurn S
(A12CuMg) phase (which nucleated preferentially around the alumina particles) andlor
Mg-rich intermetallic particles such as MgA.1204 spinel. The other is stress-assisted
migration of free magnesium atoms due to high dislocation density generated around the
alumina particles as a result of CTE rnismatch between the reinforcing alumina particles
and the ma& alloy.
It is not clear which of the rnechanisms predominates. However, it is clear that either
wiII result in causing magnesium atoms to accumulate around the aiurnina particle and
that magnesium segregation at the interface leads to magnesium depIetion from the
composite matrix, a phenornenon which has been reported to impair age-hardening
process in aluminurn rnatrix composites ( 13 8,13 9).
Table 4.2 Variation of magnesium content (wt. %) with aging in 26 18+ 10 MMC.
Thermal Treatment
Source As-Cast Naturally S h at 2000 h at
(No Solution Aged 200 OC 200 O C
Treatmen t)
Matrix 1.65 1.48 1.58 0.96
Ring 1.34 2.1 1 2.83 3.12
42.3 Other Interxnetallic Phases
A common feature of aluminurn dloy systems is the wide range of intermetailic phases,
which occur as a result of the highly electronegative and trivalent properties of
duminum. Some intennetaüics are present in both the unreinforced dloy and composite.
Figure 4.10 (a) is a SEM micrograph of solution heat treated and aged sarnple of
2618+10 composite showing a bnght silicon-rich particle adjacent to an aluminide
particle. The EDX spectra are shown in Figure 4.10@). Table 4.3 shows the EPMA
point analysis of representative silicon-nch particles present. From Table 4.3 and Figure
4.10(b), it is inferred that these particles are aluminosilicate particles (oxygen makes up
the difference in the total). They could also be particles of the quatemary phase h based
on the formulas AlFu2Mg8Si6 and &CuMgsSb (2).
Figure 4.10. (a) SEM rnicrograph of aged sarnple of 2618+10 composite showing a
silicon-rich particle. (b) EDX spectra of a silicon-rich particle.
Si-rich particles were also observed in the residues obtained from the AR treatrnent.
Figure 4.1 1(a) shows the BSE image of a Si-nch particle attached to an alumina particle
(the alumina particle is Iabeled X) and Figure 4.1 l(b) shows the EDX spectra. The EDX
spectra do not show the presence of copper as expected from Table 4.3 or found in
Figure 4.10@). These may be duminosilicate particles, Mg2Si precipitates. or temary
ALMgSi phases. Figures 4.12(a) and @) show respectively the BSE image and the EDX
spectra of another type of intemetallic particle observed in the residues. This seems to
show the presence of the quatemary A18Si6Mg3Fe phase, which is reported to occur in
duminum-nch aluminum alloys (2). Also, Figures 4.1 3(a) and (b) show, respectively,
the BSE image and the EDX spectra of a representative aluminu~n and iron-rich particle
observed in the residues, thus indicating a possible presence of the binary m e or Ai6Fe
phase.
Table 4.3 EPMA point analysis of Si-rich particles.
Particle Element -wt % Likely Phase
# Si Zn Fe Cu Mg Ni Al Total
1 13.76 0.22 0.37 9.62 2.70 0.13 53.84 80.66 ALSiOl
2 15.59 0.12 0.65 6.46 5.68 0.29 51.93 80.72 AI3SiO2
Figure 4.11. (a) A silicon-rich particle attached to an alumina paaicle. (b) the EDX
spectra.
Figure 4.12. (a) Microstructure of an intermetallic particle rich in silicon, magnesium,
and iron. (b) the EDX.
Figure 4.13. (a) Microstructure of an intermetailic particle rich in iron. @) the EDX.
69
4.3 TEM Results
Transmission electron microscopy (TEM) investigations were undertaken to obtain
detailed information about the microstnicture of the test materials with a view to
understanding the phases present, to characterize the sequence of precipitate evolution in
the matrix ailoy, and to observe the effect of dumina particles on the overdl
precipitation process in 26 18.
4.3.1 Insoluble Particles
hsoluble particles were observed to be distributed either within the matrix or at matrix
grain boundaries, matrix/alumina interfaces or sometimes adjacent to one another in the
samples exarnined. The insoluble particles that featured prominently in this study are
duminide and aluminosilicate particles. Figure 4.14(a) shows a bright field (BF)
micrograph of aluminide particles (the large particles) while Figure 4.14@) shows an
aluminide particle situated at the matrix grain boundary. It is seen in Figure 4.14(a) that
the duminide particles are not uniformly shaped. Usuaily, they have an oblong shape
with curved ends. The shapes they assume are aec ted by grain boundaries as shown in
Figure 4.14@), and also by the presence of dumina or other obstructive particles. Figure
4.15 shows that two aiuminide particles are occasionally merged with a small neck
separating them at the joint. These particles also varied in size (about 0.05 - 2.5 pm
wide) and were found to act as obstacles to or sources of dislocations.
Figure 4.16(a) shows a typical aluminosilicate panicle while Figure 4.16@) shows an
aluminosilicate particle located adjacent to an aluminide particle. Figure 4.17 shows the
EDX spectra of a typical aluminosilicate particle. These particles occur randomly in the
sarnples examined and are generally smailer in size than the aluminide particles (see the
smaller and more circular particles in Figure 4.14(a)).
Figure 4.14. (a) TEM micrograph of durninide particles. (b) bright field TEM image of
an duminide particle at the rnatrix grain b o u n d q .
Figure 4.15. TEM rnicrograph of aged 26 18 showing two merging A1,FeNi particles.
72
Figure 4.16. TEM micrograph of aged 26 18 showing (a) an aluminosilicate particle and
@) an aluminosilicate particle lying adjacent to an A1,FeNi particle.
Figure 4.17. EDX spectra from an aluminosilicate particle.
74
The source of the aluminosilicate (AixSi,Oa particles is not completely known. The only
known source is the matrix alloy which contains - 0.18 wt. % silicon, but it is unlikely
that this can account for the high Si content found in the aluminosilicates shown in
Figure 4.16. A possible source is impurity silica particles (sand) that were inadvertently
carried over into the melt from the materiai cleaning stage. In accordance with
thennodynamic data from references (146. 147), SiO, may also be a product of chemical
reaction between free silicon (fiom the alloy) and oxygen. Under favorable conditions,
duminum can combine with oxygen to form alumina according to equation 4.1 while
free silicon can react with oxygen to form SiO, according to equation 4.2. -
Other energetically f-avored reactions that might account for the presence of embedded
and/or isolated alurninosilicates are given in equations (4.3) - (4.5).
Equations (4.2) and (4.3) can account for the presence of isolated aluminosiiicate
particles shown in Figures 4.10(a) and 4.16(a) and (b), but they may not account for
those embedded in dumina particles (see Figure 4.1 l(a)). Equations (4.4) and (4.5) may
explain the attachent of Si-nch particles to alumina particles. However, only equation
(4.5) descnbes the formation of aluminosilicate particIes as a result of a direct reaction
(a solid state reaction) between two ceramics, dumina and silica.
4.3.2 Precipitate Phases
TEM investigation was camied out on sarnples that were solution heat treated at 53W5
OC for 2 h, water quenched, aged naturally for 40 h, and subsequently aged artificidly at
200 OC for various lengths of tirnes. After aging at 200 OC for up to 8 h, S' precipitates
were detected in samples of the monolithic 2618 and the composites. The [OOl],,,,
and [112],,, beam directions were used interchangeably and the confirmation of the
presence of the transitional S' and equilibrium S phases were based on the work of
Gupta et al. (33). Figures 4.18 (a) and (b), show the bright field image of S' precipitates
observed in the matrix and at the grain boundary of AA26 18 samples aged for 8 h at 200
OC, respectively.
Generally, S' precipitates were distributed in the matrix, at grain boundaries, and in the
vicinity of the duminide particles. The BF images of S' precipitates observed in the
2618+15 composite sarnple aged for 8 h at 200 OC are shown in Figures 4.19 (a) and (b)
for the [112], and [OOl], directions, respectively. The corresponding selected area
diffraction patterns (SADPs) in the [112],,, and [OOl],,,, directions are shown in
Figures 4.19(c) and (d), respectively. Figure 4.20(a) is the BF image of S' precipitates
near an alurnina particle in the 2618+10 composite aged for 8 h at 200 OC while Figure
4.200>) shows the corresponding SADP in the [112],,,, direction. The Sr precipitates
in the samples of the composite material aged for 8 h at 200 OC generally gave stronger
diffraction than those observed in the samples of the monolithic alloy given the same
level of heat treatment.
Further, bulky Cu-rich phases were randomly observed CO-existing with the S'
precipitates in samples aged for 8 h at 200 OC. Figures 4.21 (a) and (b) show respectively
the bright field TEM micrographs of this phase in aged monolithic 26 18 and 26 1 8 4 5
composite while Figure 4.21 (c) and (d) show the corresponding SADPs in the [OO1]m~,
direction. Interfacial dislocations can be seen clearly in Figure 4.2 lm).
Figure 4.18. TEM bright field image of aged 2618 showing S' distribution: (a) in the
matrix; and (b) at the grain boundary, al1 in the [112],,,, direction.
Figure 4.19. Bright field image of aaificiaily aged 2618+15 composite showing S'
distribution in (a) [l 12],,~x and (b) [OOl],,~, directions. (c) and (d) are the
corresponding SADPs, respectively.
Figure 4.20. (a) Bright field image of aged 2618+10 composite showing S' distribution
in the [ 1 12],,h, direction; (b) corresponding SADP.
Figure 4.21. Bright field image of aged samples showing bulky Cu-rich precipitates in
(a) 2618; @) 2618+15 composite; (c) and (d) corresponding SADPs in the [O0 l],,h,
direction, respectively.
The diffraction patterns were indexed with the lattice parameters proposed by Perlitz and
Westgren (48) for the S-N2CuMg phase (a = 0.400 MI, b = 0.924, c = 0.714 nm).
Although not al1 the diffraction spots could be indexed as a result of doubIe diffraction,
the results showed that the buIky phases were identical to S precipitates. Zhang and
Cantor (133) have observed buIky S phases in monolithic AM6 18 aged for 24 h at 190
OC and in 2618/SiCp MMC aged for 20 h at 200 OC. However, Gupta et al. (33) have
reported that the S' phase is only a slightly strained version of the equilibrium S-
(Al?CuMg) phase and, as such, both phases need not be distinguished from one another.
Interfacial dislocations c m be seen on these precipitates indicating strain absence.
Therefore, these buky phases could be S' precipitates in their early stage of formation or
precursors to S' precipitates which happen to produce similar diffraction effects as the
transition S' precipitates.
Precipitate phases based on the 8-AI2Cu formula were observed in the overaged samples
of the composites, probably due to depletion of Mg (see section 4.2.2). Figures 4.22 (a)
and @) show respectively the dark field (DF) image and the corresponding SADP (in the
[112],~, direction) of the 0 phase in 2618+15 composite aged for 16 h at 200 OC.
EDX analysis of a typical f3 phase showed that it was ncher in Cu than the parent matrix.
Needle-shaped 8" phases were observed in the samples aged naturally for 45 days.
Figure 4.23 (a) and (b) show the BF image and the corresponding SkDP of the O" phase,
respec tively.
Further, the plate-like X precipitates reported recently by Jin et al. (135) were observed
in the composite samples aged at 200 OC and those aged naturally. Figure 4.24(a) shows
a BF image of a typical X phase in 2618+15 composite aged naturally for 45 days while
Figure 4.24(b) shows the corresponding SADP in the [112Im~, direction. The
parameters of the two-dimensional unit ceII corresponding to Figure 4.24(b) (as
illustrated in Figure 4.25) were measured and are listed in Table 4.4. In the table, RI and
R2 represent the lengths of the shortest and second shortest reciprocal lattice vectors
respectively, a is the angle between Ri and Rz and dl is the plane spacing corresponding
Figure 4.22. Dark field image of the 8 phase in overaged sarnple of 26 18+ 15 composite;
(b) corresponding SADP in the [112],,,, direction.
Figure 4.23. (a) Bright field image of 8" precipitates in nanirally aged sample of 2618.
@) corresponding SADP in the [112],,, direction.
Figure 4.24. (a) Bright field image of X precipitates in naturally aged sample of 26 18+ 15
composite, @) corresponding SADP in the [l 12],0, direction, and (c) indexed pattern
of (W.
Figure 4.25. A sketch of two-dimensional lattice ce11 of the X phase.
Table 4.4. Reciprocai lattice parameters of 2-dimensional unit cell of the X phase.
I Measured 1 Calculated 1 Indices
to RI. The data in Table 4.4 are consistent with those published in the literature (135).
Figures 4.26 (a) and @) show respectively the DF image and the SADP of X phase in a
sarnple aged for 19 h at 200 OC.
The EDX analysis of the X phase shows that it contains Al, Cu and Mg, although Jin et
aL(135) have, in addition to these elements, observed small amounts of Fe and Mn
peaks in the AA2124 they studied. The X phases found in the naturally aged samples
were thinner than those found in the artificially aged samples. This suggests that during
artificid aging, the init idy thin plate-like X phases increase in size and, Iater, with
continued increase in the aging time, attain a buky size. Nso, the amount of X phase is
much smaller than that of the needle-like precipitates. Jin et al. (135) have reported that
although the S and X phases both have an orthorhombic crystal structure, their lattice
Figure 4.26. (a) Dark field image of the X phase in overaged sample of 2618+15
composite; @) corresponding SADP in the [112],,~, direction.
pararneters and atomic compositions are different. The lattice parameters of the S phase
are: a = 0.400 nm, b = 0.923 nm. and c = 0.714 nrn with Cmcm space group (48) while
the X phase has lattice pararneters of a = 0.492 nm, b = 0.852 nm, and c = 0.70 1 nm
with Cmmm space group (135). Further, their orientation relationships with the rnatrix
are different. At this time, it is not clear what influence the X phase has on the overall
precipitation kinetics of the alloy. However, it is believed that its presence is due to
changes in rnatrix composition in the composites.
Phases based on the AllCu formula are known to occur in Al-Cu(Mg) alloys with zero or
low magnesium content (137). Both S' and €la precipitate in ailoys with higher Cu:Mg
weight ratios (4). A simple calcuiation (rvt. % i atomic mass for each element) shows
that the Cu:Mg atornic ratio in the matrix alloy and the composites is - 1 2 . Thus, there
are more magnesiurn atoms than there are copper atoms (per unit volume of
supersaturated solution). lt was expected that precipitation of the S' and S-(A1,CuMg)
phases would strongly prevail in A M 6 18 during aging. However. other precipitate
phases (e-g., X and A12Cu-based phases) occurred, thus making the aging sequence quite
complex and different from that mostly reported in the literanire for the 2w( series Al-
Cu-Mg alloys. (see Equation 2.2). In the composite. depletion of Mg will change this
ratio in the rnatrix. The Cu:Mg atomic or weight ratio would increase and Al,Cu phases
will precipitate. This rnight also explain the presence of the X phase in the composites.
The Mg,Si phase is another phase expected to occur in this alloy. Si combines with Mg
readily to form Mg,Si precipitates which assist precipitation hardening. However, they
were not observed in the TEM.
Based on the results of the present study, the aging sequence in AA2618 c m be
interpreted in t e m s of three main precipitation reactions:
SSS + GP zones + 8" + 8' + 8 . . . (4.7)
However, it is not yet known if the X phase has any transition phases.
4-3.3 Crystal Structure of Aluminide Phase
As mentioned previously, the structural formula of the alurninide particles observed in
the present investigation was different from the MgFeNi formula reponed in the
literature. Therefore, a detailed study of the crystal structure of A1,FeNi phase was
carried out using TEM. Each selected area diffraction pattern (SADP) observed in the
TEM represents a two-dimensional reciprocal lattice plane. Thus, the 3-dimensional
reciprocal lattice of a crystal (phase) c m be obtained by tilting the crystal systematically
dong different crystallographic directions. Usually, the operation starts with the
specimen oriented in such a way that the zone axis of a high symrnetry SADP is parallel
to the incident electron beam. The specimen is then tilted about a chosen reciprocal
lattice axis in a given direction until the specimen orientation with respect to the electron
beam reaches another high-symmetry SADP zone a i s . Similar tilting operations are
then carried out for other directions. This procedure has been used by many investigators
( 135,136,148- 150) for crystal structure determination using the TEM technique. From
the results, the three high-syrnmetry reciprocal lattice planes (which may be
perpendicular to each other) are then selected to form a 3-dimensionai unit ce11 of the
phase.
Figures 4.27 (a)-(f) are the SADPs of the A1,FeNi phase obtained by tilting the crystal in
one direction (example, M to MM direction) while Figures 4.28 (a)-@) are the SADPs
obtained by tilting in another direction (example, N to NN direction). The measured
values of the parameters of the SADPs for these two directions are shown in Tables 4.5
and 4.6, respectively. In the tables, RI and R2 are respectively the shortest and second
shortest vectors of the 2-D reciprocal section, d l is the interplanar spacing corresponding
Figure 4.27. SADPs of the A1,FeNi phase obtained by tilting between [100] and [O011
zone axes.
Figure 4.28. SADPs of the A.i,FeNi phase obtained by tilting between [3 101 and [O 1 O]
zones axes.
90
to RI, and 0 is the angle between R1 and Rr. The SADPs in Figures 4.27 (a)-(f) and
Figures 4.28 (a)-(f) were indexed using the lattice parameters reported in reference (133)
(i-e., a = 0.6213 nm. b = 0.629 nm, c = 0.8557 nm, and P = 94.76O) and the results are
shown in Table 4.7. In this table (HiKILl) and (HZK2L2) are respectiveIy the planes
corresponding to the 2-D reciprocal lattice vecton RI and R2, while [UVW] is the
direction of the reciprocd lattice plane (i.e., the zone axis of the SADP).
The calculated and expenmental values of the angle between the zone axes of the
SADPs in Figures 4.27 (a)-(f) and Figures 4.28 (a)-(f) are shown in Table 4.8. One way
to veriQ the consistency and accuracy of the lattice parameters obtained in previous
studies (133) for indexing the Ai,FeNi particle is to compare the experimental and
calculated values of the paraneters R2/R1. dl, and 0. As shown in Table 4.7, the
measured and calculated angles differ appreciably in rnost cases. The calculated
interplanar angles (angles between the reciprocal planes) obtained for the SADPs in
Figures 4.27 (a)-(f) and Figures 4.28(a)-(f) are cornpared with the corresponding
experimental values in Table 4.8. From the upper section of Table 4.8, it can be seen that
the experimental values are in fairly good agreement with the calculated values of the
angles between the SADPs in Figures 4.27 and 4.28. However, the calculated angle
between Figures 4.27(a) and (f) (Le., between [1-101 and [100]) is not in good agreement
with the measured value. A simi1a.r result was obtained for Figures 4.28(a) and (f) and
this is shown in the lower section of Table 4.8.
Since the SADPs in Figures 4.27 and 4.28 could not be completely indexed with the
lattice parameters suggested by Zhmg and Cantor (133), it became necessary to establish
a more consistent mode1 for the A1,FeNi phase. Using the measured values in Tabie 4.7
and a cornputer-based SADP analysis program (140), it was determined that the A1,FeNi
phase has a monoclinic structure with lanice parameters a = 0.8673 nrn, b = 0.900 nm, c
= 0.8591 nrn, and P = 83.504". The three reciprocal lattice planes shown in Figures
4.27(a), 4.27(f), and 4.28(f) were chosen to fonn a three-dimensional unit ceIl of the
A1,FeNi phase.
Table 4.5. Measured values of t le SADPs shown in Figures 4.27(a)-(0.
Table 4.6. Measured values of the SADPs shown in Figures 4.28(a)-(0.
-
Fig. No.+ 4.26a 4.26b 4.26~ 4.26d 4.26e 4.26f
R21R~ 3 .O3 4.24 1.36 4.24 3 .O3 1.97
dl& 8.59 8.59 8.59 8.59 8.59 8.59
0 (O) 91 91.6 93 95 95 96
Table 4.7. Cdculated and measured values of crystailographic parameters of A1,FeNi
using a = 0.6213 nm; b = 0.6290 nm; c = 0.8557nrn; P = 94.76O (133).
1 Fig. 1 Measured Values Calculated V:
R f l , 1 d , ( h
lues Indices
Table 4.8. Calculated and measured values of angles between the zone axis of the
SADPs shown in Figures 4.27 and 4.28 using a = 0.6213 nm; b = 0.6290 nm; c =
0.8557 nm; P = 94-76" (133).
The three-dimensional reconstruction of these SADPs (i.e., those corresponding to the
S M o 1 2 3
[100], [OOl], and [O101 zone axes) suggests that the AI,FeNi phase is C-centered
monoclinic. Computer simulated electron difiaction patterns corresponding to the three
Orientation -110" -331 -331 " -221 -221 "-111
zone axes are shown in Figures 4.29(a)-(c). The real lattice crystal structure of the
Al,FeNi phase which corresponds to the latter structure is also C-centered monoclinic
CalcuIated 8 (O)
17.54 7 -62 17.28
with lattice parameters as mentioned above. This is shown in Figure 4.30. The indexed
Measured 8 (O)
17.10 7 -76 16.74
patterns corresponding to the [100], [O0 11, and [O101 zone axes are shown schematically
in Figures 4.3 1 (a)-(c) while Figures 4.32 (a)-(c) show the stereographic projections for
these zones.
Table 4.9 shows the results of the calculations. Again, (HIKILI) and (H2K2L2) are the
planes corresponding to R, and &. These are respectively the shortest and second
shortest vectors of the 2-dimensional reciprocal lattice plane. 0 is the angle between Ri
and R2 whereas is the zone axis of the SADP. It can be seen that the calculated
values are fairly consistent with the expenmental values. Again, if the above lattice
parameters are correct, the calculated angles between the zone axis of the SADPs in
Figure 4.29. Computer simulated reciprocal lattice corresponding to (a) [ 1001, (b) [001], and (c) [OIO] zone axes.
Figure 4.30. Crystal structure of A1,FeNi phase based on C-centered model.
[IO01 Zone [O011 Zone [O101 Zone
Figure 4.3 1. Indexed patterns of [100], [O0 11 and [O 101 zone axes.
Figure 4.32. Cornputer generated stereographic projections for (a) [100], (b) [00 11, and
(c) [O IO] zone axes.
Figure 4.32 (continued).
Figures 4.27 and 4.28 should be consistent with the corresponding experimental values.
Table 4.10 shows the cornparison. It can be seen fiom Tables 4.9 and 4.10 that the
results are fairly satisfactory.
The lattice parameters of the aluminide phase were also examined by the convergent
beam electron diffraction (CBED) technique. Figures 4.33 (a) and @) show the fust-
order Laue zone (FOLZ) rings of the phase in the [110],~, and [130],,, zones,
respectively. In addition, the standard seven reflection conditions of the SADPs shown
in Figs. 4.27 and 4.28 were determined. These are given in Table 4.1 1. It is important to
note that double reflection conditions seem to have occurred in some SADPs at the (001)
positions (see Figures. 4.27(a) and 4.28(f) and Figures 4.3 1 (a) and (c)). Baring possible
errors in the observation of double reflection, the possible point groups and space groups
of the Al,FeNi phase were then determined from the relationship between the seven
reflection conditions and the crystai symrnetry as obtained fiom the International Tables
for Crystallography (15 1). The results are also shown in Table 4.11.
Figure 4.33. First-order Laue zone (FOLZ) rings of the Ai,FeNi phase for [1 10],,ui, and
[130Imauix zone axes.
Table 4.9. Calculated and rneasured values of crystallographic parameters of A1,FeNi
phase based on C-centered structure using a = 0.867 nm; b = 0.9CO nm; c = 0.859 nm;
Indices
Table 4.10. Calculated and measured values of angles between the zone axis of the
4a 4b 4c
SADPs shown in Figures 4.27 and 4.28 based on C-centered monoclinic structure with
a = 0.867 nrn; b = 0.900 nm; c = 0.859 nm; P = 83.509
1 S/No 1 Orientation 1 Calculated 8 (O) 1 Measured 0 (O) 1
(HlK,L,)
00 1 020 020
R2/R1 1.91 1.79 1.27
R$Rl 1.90 1.79 1.33
dl (A) 8.59 4.50 4.50
(H?K2k) 020 11-3 11-2
dl (A) 8.54 4.50 4.50
0 (O)
93 70 65
[UVW
100 30 1 20 1
8 (O)
90.00 73.79 63.85
Table 4.1 1. Reflection conditions of Al,FeNi phase and possible point and space groups.
Reflection Conditions Point Group Space Groups
( h m )
h = 2 n 2 m 2/rn C2 Cm C2/m
(OkO)
k = 2 n
(001)
none
(hkO)
h+k= 2n
(OkO
k = 2n
(ho0
h = 2n
(hW
none
4.4 DSC Results
In this section, the results of a DSC investigation undertaken to monitor the effect of
alumina particles on precipitation reactions in AA26 18 and its composites are presented.
The DSC peaks analyzed were assumed to be due to the three reactions given in
equations (4.6)-(4.8). However, in general, the analyses were treated as if the peaks were
due to GPB zone, S-(Ai-CuMg), and Al,Cu precipitation reactions only.
4.4.1 General Description
The DSC thennograms of unreinforced 2618 and 2618+15 composite obtained at
different heating rates are shown in Figures 4.34 and 4.35, respectively. The curves in
both figures consist of sirnilar heat effects and the principal features of interest in this
study are the two exothermic peaks, A and C, md the two endothermic effects, B and D.
Peak A is attributed to GPB zone formation ( 15,17,108) and it is a characteristic thermal
effect in the 2xxx series of alurninum ailoys. Trough B has been ascribed to the
dissolution of GPB zones. The second exotkierrnic reaction zone, C. reveals a doublet
exothermic peaks at heating rates above 5 "C/min. Therefore, at lower heating rates. the
two exothermic peaks seem to be superimposed on each other and, as such, are
indistinguishable. Peak C is different from the single exothemiic peak attributed by
various workers (15,17,108) to precipitation of the S' (&CuMg) phase.
In 26 18 the Cu to Mg atom ratio is less than 1. This implies that the S ' phase may not be
the only phase precipitating in the temperature range of C. As shown in section 4-32 .
TEM examination of the microstructure of artificially aged samples of 2618 and its
composites revealed the presence of phases based on the Ai2Cu-structure as well. Gupta
et al. (33) have suggested that the S' phase is only a slightly strained version of the S
phase. Thus, the thermal effects associated with S' -t S transformation may not be
sufficient to exhibit a separate peak during the DSC scan. Hence, it is proposed diat the
Figure 4.34. DSC thermograms of as-quenched 2618 for various heating rates.
O 100 200 300 400 500 600
Temperature (OC)
Figure 4.35. DSC thermograrns of as-quenched 2618+15 composite for various heating
rates.
doublet exotherms are due to the CO-precipitation of the S' and 0' phases.
The literature is not definitive regarding the precise temperatures at which S' and 8'
formation reactions are maximum during DSC scan because the peak reaction
temperatures of these phases Vary with heating rate, previous thermal history, and the
material. Therefore, two postulations cm be suggested regarding the doublet exothermic
peak (C), namely: (i) the first sub-peak is due to the simultaneous precipitation of S' and
0' whereas the second sub-peak is due to their s ~ u l t a n e o u s growth (1 15); (ii) the first
sub-peak is due to the formation of S' phase whereas the second sub-peak is due to the
formation of 8 ' phase.
The results obtained for IM aluminum ailoys scanned at 10 " C h i n in the as-quenched
condition shows that 8' formation reaction peaked at 279 OC in AA 2219 whereas S'
formation reaction peaked at 268 OC in AA2124 (15). Reaction peak temperatures of 285
OC, 278.7 OC, and 277.5 OC were reported for 8' formation in as-quenched AA2014 and
its composites containing 10 and 15 vol. % A1203 particles, respectively (139). For
N I 2 4 aged naturally for 30 min before DSC scan at 10 "Chin , the S' phase
formation reaction peaked at 264.7 OC while the data obtained for AA2219-T3 1, aged
naturally for six months and scanned at the same heating rate, shows that 0' formation
reaction peaked at 250 OC (17). The difficulty caused by the unresolved peaks
notwithstanding, postulation (ii) seems to descnbe the occurrence of the doublet
exothermic peaks better than posrulation (i) on the basis of the quoted literature data.
The final dissolution trough D is attributed to the dissolution of the equilibnum phases,
S (A12CuMg) and 8 (AlzCu).
In nsing-temperature techniques such as the DSC, the extent of transformation at any
particular temperature cm be controlled by thermodynarnic equilibrium or kinetic
limitations (152). In thermaily activated processes the heat effects shift to higher
temperatures with increasing heating rate. Figures 4.34 and 4.35 show that the formation
and dissolution of GPB zones and the formation and dissolution of the S ' and 8 ' phases
are dominated by their reaction kinetics. The average values of the reaction peak
temperatures and the heat effects associated with these transformations are shown in
Tables 4.12 and 4.13, respectively. It is not possible to quanti& the reaction enthalpies
for the dissolution reaction under trough D because of the difficulty in establishing the
onset and endset temperatures. The peak temperatures correspond to the points of
maximum enthalpy of formation or dissolution. The area under the DSC peaks gives the
total enthalpy of formation or dissolution and is directly related to the molar heat of
reaction and the volume fraction of the precipitate formed or dissolved (15).
The data in Figures 4.34 and 4.35 as well as Tables 4.12 and 4.13 show that:
the volume fraction of phases formed in the unreinforced 26 18 is larger than that in
the composite matenal. This has been reported by other workers (23,25,32) and has
been explained in terms of the free vacancy concentration. Vacancies are required
for the nucleation of GPB zones. The relativeiy high dislocation density found in
the MMCs (97) on cooling from the fabrication/solution heat treatment temperature
gives nse to vacancy annihilation. Hence, with the vacancy concentration being
lower in the MMCs than the unreinforced dloys, fewer stable GPB zones are
formed in the MMC during DSC scans. However, this hypothesis cannot explain
the observed higher GPB zone dissolution enthaipy usuaily obtained in the MMCs.
Therefore, it is suggested that the smaller amount of GPB zones formed in the
MMC is due to a synergistic action of two competing phenornena, namely: (i) the
prior precipitation of stable GPB zones dunng quenching and, (ii) the decreased-
vacancy-content phenomenon.
the peak temperature for GPB zone formation is less in the monolithic dloy than in
the composite mater;.al. This suggests that GPB zone formation requires a lower
driving force in the unreinforced alloy compared-to the MMC
the peak temperatures for GPB zone dissolution and 8' + S' formation are less in
the composite compared to the unreinforced alloy.
Table 4.12. Variation of DSC peak temperature with heating rate.
Peak Temperature (OC)
Heating Peak A Trough B Peak C
Rate GPB Zone GPB Zone S '+8 ' Formation
(O C h i n ) Formation Dissolution
Peak 1 Peak 2
Table 4.13. Total heat effects of DSC peaks at different heating rates.
("Chin) Total Heat Effect (J/g)
Peak A Trough B Peak C
GPB Zone GPB Zone S ' 4 '
Formation Dissolution Formation
2618 2618+15 2618 2618+15 2618 2618+15
4.4.2 Determination of Kinetic Parameters for Precipitation
4-4-2.1 GPB Zone Formation (Peak A)
The DSC peak A in both Figures 4-34 and 4.35 is due to GPB zone formation. The mole
fraction, y, of precipitates formed and the rate of precipitation, (dy/dT), which were
calcdated from the DSC peak A as functions of temperature are shown in Figures 4.36
and 4.37 for unreinforced 26 18. Similar figures are plotted for the 26 18+15 composite in
Figures 4.38 and 4.39. The y vs T curves (Figures 4.36 and 4.38) have sigmoid shape.
AIso, the y-T curves and the maxima of the (dy/dT)-T curves shifi to higher temperatures
with increasing heating rate. These results imply that GPB zone formation reaction is
kinetically controlled. Arrhenius plots of the data extracted from Figures 4.36-4.39 are
made using equation (2.25) for various values of yi (0.15-0.9). In Figures 4.40 and 4.41,
representative data are shown for the unreinforced alloy and 2618+15 composite. The
average of the least square slopes of the straight lines obtained fiorn the [ (dy /d~)y , O,]
vs. (11~) plots gives the activation energy, E, for GPB zone formation in each material.
The kinetic parameters obtained for the two materids are summarized in Table 4.14.
Kinetic parameters for GPB zone formation were also calculated using equation (2.27),
with the last term neglected. PJthough some error would be incurred, it h a been shown
(1 16- 1 18) that for non-isothermal kinetics PT G 1 and the error made by discarding this
last term is srnail, especially when the imposed heating rate is not more than 40 "C/min.
Higher heating rates can cause serious temperature gradients in the samples on heating.
The average value of the Section-point temperature (Ti) of the y vs T curve was
determined (for the GPB zone formation peak) for different heating rates. Table 4.15
shows the Ti data for the test rnatenals. B y plotting fn(q2 /a) against (11~) as shown in
Figure 4.42 the effective activation energies for GPB zone formation in both materials
were obtained (see Table 4.14). These values agree fairly well with the results obtained
40 60 80 100 1 20 140 160
Temperature ( O C )
i,z Figure 4.36. y vs temperature curves for GPB zone formation at different heating rates
(26 18).
100 120
Temperature (OC)
Figure 4.37. (dy/dT)@ vs temperature curves for GPB zone formation at different
heating rates (26 18).
60 80 100 120 160 180
Temperature (OC)
Figure 4.38. y vs temperature curves for GPB zone formation in 26 l8+l5 composite at
different heating rates.
60 80 100 120 140 160 180
Temperature (OC)
Figure 4.39. (dy1dT)O vs temperature curves for GPB zone formation in 2618+15
composite at different heating rates.
Figure 4.40. Arrhenius plots for determination of the activation energy for GPB zone
formation in 26 18 (based on equatim 2.25).
Figure 4.41. Arrhenius plots for detemiination of the activation energy for GPB zone
formation in 2618+15 composite (based on equation 2.25).
from the A M (Avrarni-Johnson Mehl) analysis. Baring extrapolation errors, the pre-
exponential factors, ko, were also calculated from the intercepts of the straight lines
shown in Figure 4.42.
The single-heating rate method (equation (2.26)) was also used to determine E. The
exponent n which yielded a straight Iine fit to the expenmental data (see Figure 4.43)
was determined by an iterative process using equations (2.19) and (2.26). The activation
energy, E, was calculated for each heating rate from the least square slope (-E/R)of the
resulting straight line. The values obtained for the unreinforced alloy are: E = 68.38d.4
Wmol; ko = 1.2 x 10l0 s-'; and the growth parameter n = 0.99I0.02. Given that
nucleation of GPB zones occurs during the early stage of DSC scan and further
thickening occurs by a diffusion-controlled rnechanism, the value of n obtained is
consistent with low temperature precipitation of GPB zones. It is within the range of
values reported by Luo er al. (104) for GPB zone formation in AA8090, and is also in
close agreement with data published in reference (100). The constant ko is in close
agreement with that obtained from the Kissinger method.
The migration energy of vacancies in Al-Cu alloys has been reported to be in the range
41-66.9 kJ/mol (153,154). The values of activation energy for GPB zone formation
obtained in the present study for AA2618 using the varying-heating rate method are
outside this range, and are higher than those reported in the literature for Al-Cu-Mg
alloys. Jena et al. (108) obtaùied E = 55.6 I d h o 1 for Al-1.53 wt. % Cu-0.79 wt. % Mg
alloy using the DSC technique. Luo et al. (104) reported E = 6 1 - 0 e . 2 kJ/mol and E =
53.9t 1.6 I d h o 1 for AA8090 using the DSC and resistivity techniques, respectively;
while Horiuchi and Minonishi (153) obtained E = 64.0 kVmol for the A l 4 2 wt. % Cu-
1.4 wt. % Mg alloy based on resistivity measurements. Unlike the aforernentioned
alloys, AA2618 is known to contain an appreciable amount of insoluble intermetallics
(such as alurninide particles) which create a large amount of interfaces. Like the ceramic
particles in MMC, the intermetaliic particles also have different coefficient of thermal
expansion from the Al-nch matrix. As such, on cooling from the solution heat treatment
Figure 4.42. Plots after equation 2.27 for GPB zone formation in 26 18 and 26 1 8 4 5
composite.
Figure 4.43. Arrhenius plots after equations 2.19 and 2.26 for the determination offly)
for GPB zone formation in 26 18 (@ = 10 and 20 "Umin).
Table 4.14. Kinetic parameters for GPB zone precipitation reactions.
Parame ter GPB Zone Formation GPB Zone Dissot ut ion 2618 2618-1-15 2618 2618t15
AJM Kissinger AJM Kissinger AJM Kissinger AJM Kissinger
E (kl/mol) 76.8î0.78 7 1.67 82.32 1.4 74.66 157.917.0 166.1 143.617.3 127.3
Table 4.15. Variation of inflection point teinperature (Ti) with heating rate for GPB zone formation.
GPB Zone Formation Heating Rate ("Chin)
(Peak A) 5 10 15 20 30
temperature they give rise to large arnount of misfit dislocations. The dislocations and
the large interfaces can act as vacancy sinks that annihilate excess vacancies needed for
GPB zone nucleation.
The above calculations also show that GPB zone precipitation in the 26 l 8 t 15 composite
requires a higher dnving force than in the unreinforced alloy. The suppression of GPB
zone formation and its elevation to higher temperatures have been reported in some
MMCs (15,50,155). This has been attributed to the annihilation of quenched-in
vacancies by misfit dislocations. The amount of misfit dislocations generated in the
composite material is larger than in the unreinforced ailoy due to the large difference in
the coefficients of thermal expansion (CTEs) of the matrix and the alumina particles
(15,97). The presence of alumina particles also creates large arnounts of interfaces that
c m act as vacancy sinks. The low vacancy concentration in the composite matrix
reduces the vacancy migration contribution to the GPB zone nucleation while increasing
the contribution of substitutionai atom (Cu. Mg) migration. The latter process requires
the creation of additional vacancies which consequently results in an increase in the
calculated activation energy, E, (23).
4.4.2.2 GPB Zone Dissolution (Trough B)
The endothermic trough B in Figures 3.34 and 4.35 is associated wiîh GPB zone
dissolution. Figures 4.44 and 4.45 show respectively the variation of y and (dy/dT) with
temperature at various heating rates for unreinforced 2618. Sirnilar information as in
Figures 4.44 and 4.45 is shown in Figures 4.46 and 4.47 for the composite material. The
y vs T and (dy/dT) vs. T curves shift to higher temperatures with increasing heating rate.
Therefore, the GPB zone dissolution reaction is kinetically controlled in both
unreinforced 2618 and the composite. The dissolution of GPB zones in some aluminum
alloys has been attributed to a diffusion-limited mechanisrn ( 1 13,114). The calculation
of the kinetic parameters was carried out using equations (2.25) and (2.27) and the
results are shown in Table 4.14. The results obtained using these equations are shown in
Figures 4.48-50 whereas Table 4.16 shows the inflection point temperature data used
with equation (2.27). The single-heating rate equation (2.26) was also used to calculate
the kinetic parameters for GPB zone dissolution in both materials.
The radial dimensions of the cylindricd particles of GPB zones change when GPB zone
dissolution reaction takes place under diffusion-limited conditions (1 14). The results
obtained by iterating for the growth parameter (n) between equations (2.19) and (2.26)
are shown in Figure 4.51 for the two materiais (a = 5 OC/min). For AA2618, n =
0.4dl.02, and Q = 167.818.4 kl/mol whereas n = 0.5a.02 and Q = 139.6k4.9 Wmol
for the composite material. It has been reported in reference (101) that n = 0.5 for the
thickening of large plates and n - 213 for a dislocation-assisted process. The present
results suggest that GPB zone dissolution in AA2618 and its 15-vol. % composite
involves precipitate thickening andor dislocation-assisted diffusion of solutes.
The value of activation energy obtained for 2618 using the M M analysis (E = 157.9k
7.0 kJ/mol) is greater than the value of 123.9 kUmol and 12833.5 Id/mol reported in
the Literature (104, log), respectively. The concentration of quenched-in vacancies would
be lower in A M 6 18 than in the alloys studied by these workers (104,108), as previously
120 160 200 240 280
Temperature (OC)
Figure 4.44. y vs temperature curves for GPB zone dissolution in 2618 at different
heating rates.
Temperature (OC)
Figure 4.45. [(dy/dT)@j vs temperature curves for GPB zone dissolution in 26 18 at
different heating rates.
Temperature ( O C )
Figure 4.46. y vs temperature curves for GPB zone dissolution in 26 18+15 composite at
different heating rates.
Figure 4.47. [(dy/dT>@] vs ternperature curves for GPB zone dissolution in 26 l8+l5
composite at different heating rates.
Figure 4.48. Arrhenius plots d e r equation 2.25 for determination of the activation
energy for GPB zone dissolution in 26 18.
-4 ' i r I 1 1 I I I I I I 1 l I I
Figure 4.49. Arrhenius plots after equation 2.25 for determination of the activation
energy for GPB zone dissolution in 26 l8+lS composite.
Figure 4.50. Plots after equation 2.27 for determination of the activation energy for
GPB zone dissolution in 26 18 and 26 18+ 15 composite.
Figure 4.51. Arrhenius plots after equations 2.19 and 2.26 for determination of the
activation energy for GPB zone dissolution (<P = 5 "Urnin).
Table 4.16. Variation of inflection point temperature (Ti) with heating rate for GPB zone dissolution.
GPB Zone
Dissolution
Heating Rate ("Chnin)
(Trough B) 5 10 15 20 30
explained in section 4.4.2. The activation energy for solute difision during the GPB
zone dissolution reaction includes the activation energy for formation of excess
vacancies and that for solute transport. This implies a higher effective activation energy
for GPB zone dissolution in a vacancy-deficient environment than that in a vacancy-rich
environment. The consequence is that the effective activation energy for GPB
dissolution in the composite matenal wili be larger than that in the unreinforced alloy.
However, in Table 4.14 the value of the effective activation energy obtained for the
composite matenal (E = 143.6k7.3 Wmol) is smaller than that obtained for the
unreinforced alloy (E = 157.9k7.0 kT/mol). Therefore, the vacancy-content theory is not
capable of completely explaining the results obtained for GPB zone dissolution reaction
in the present study.
Starink and Mourik (120) have reported that GPB zone dissolution reaction is essentially
a two-step process. After the formation of GPB zone during a DSC scan, at least two
reactions c m occur on continued heating: (i) GPB zone dissolution (endothermic) and
(ii) GPB zone coarsening (exothermic). In the present study, the temperature range of
GPB zone dissolution during the DSC scan is conducive for solute diffusion. In the
composite matenal, the local stress fields induced in the matrix by the presence of the
reinforcements can give rise to enhanced matrix diffbsivity (and hence accelerated
dissolution kinetics) and a high dislocation density. Dislocations have been shown by
both theoretical and experimental analyses (156) to serve as short-circuit paths for solute
difision. Therefore, this could explain why GPB zone dissolution kinetics are enhanced
in the composite material as compared to the unreinforced alloy.
4.4.2.3 S ' and 8' Formation
The DSC doublet peak C in Figures 4.34 and 4.35 is due to S' and 8' precipitation.
Figures 4.52 and 4.53 show respectively the y vs. T and (dyldT) vs. T plots obtained at
various heating rates for unreinforced 2618 while Fiopes 4.54 and 4.55 show the
respective plots for the composite material. These figures show that the curves shift to
higher temperatures with increasing heating rate. However, the unresolved peaks
introduce some complications to the expected continuous and smooth sigmoid behavior.
The 8' and S' formation reactions occur in overlapping temperature intervals and there is
no error-free method to deconvolute their reaction peaks. Therefore, they are difficult
candidates for accurate kinetic analysis. In the present study, the total area (Af) under the
S' peak was calculated by doubling the area between the onset temperature and the peak
reaction temperature. Similady, the total area under the 8' peak was calculated by
doubling the area under the endset temperature and the peak reaction temperature (see
Figure 4.56). The modified y vs. T and [(dy/dl)J vs T plots of the S' phase are shown in
Figures 4.57 and 4.58, respectively, for unreinforced 2618 while Figures 4.59 and 4.60
show the respective plots for the 2618+15 composite. Sirnilar plots obtained for the 0'
phase are shown in Figures 4-61-64 It can be seen from these figures that the
precipitation of the S' and 8' phases are kinetically controlled in the unreinforced alloy
and its composites. The kinetic parameters were determined using equations (2.19) and
(2.26) and the results are shown in Figures 4.65 and 4.66 and Table 4.17 for a heating
rate Q> = 20 "C/min The overall activation energy was cdculated from Q = nE.
240 260 280 300 320 340 360 380
Temperature (OC)
Figure 4.52. y vs temperature curves for S ' and 8' formation in 26 18 at different heating
rates.
Figure 4.53. [(dy/dT)<P] vs temperature curves for S' and 0' formation in 2618 at
different heating rates.
240 260 280 300 320 340 360 380
Temperature (OC)
Figure 4.54. y vs temperature curves for S' and 8' formation in 26 18+15 composite at
different heating rates.
Figure 4.55. [(dy/dT)Q] vs temperature curves for S' and 8' formation in 26 18+ 15
composite at different heating rates.
Energy
Figure 4.56. Schematic illustration of Al determination.
Table 4.17. Kinetic parameters for S ' and 0 ' formation in 26 18 and 26 18+ 15.
Parameter* S ' Formation 0' Formation
* Heating rate = 20 "C/min.
260 270 280 290 300 310 320 330
Temperature (OC)
Figure 4.57. y vs temperature curves for S * formation in 26 18 at different heating rates.
260 270 280 290 300 3 10 320 330
Temperature (OC)
Figure 4.58. [(dy/dT)@] vs temperature curves for S' formation in 2618 at different
heating rates.
Figure 4.59. y vs temperature curves for S ' formation in 26 18+15 composite at different
heating rates.
Figure 4.60. [(dy/dT)Q] vs temperature curves for S ' formation in 26 l8+15 composite
at different heating rates.
Figure 4.61. y vs temperature curves for 0' formation in 261 8 at different heating rates.
Figure 4.62. [(dy/dT)@] vs temperature curves for 0' formation in 26 18 at different
heating rates.
Figure 4.63. y vs temperature cuves for 8' formation in 2618+15 composite at different
heating rates.
Figure 4.64. [(dy/dT)Q] vs temperature curves for 8' formation in 2618+15 composite
at different heating rates.
Figure 4.65. Arrhenius plots after equations 2.19 and 2.26 for determination of the
activation energy for S' phase formation (a = 20 "Chin).
Figure 4.66. Arrhenius plots after equations 2.19 and 2.26 for determination of the
activation energy for 8' formation (a = 20 "C/rnui).
Table 4.17 shows that, for S' precipitation (the first exothermic reaction peak under peak
C), n = 1.54 for both unreinforced 2618 and its composite. This is close to the value (n =
1.56) obtained for S' formation in Al-2.5 wt.%Cu-1 .Zwt.%Mg alloy (157). This result
indicates that the growth mechanism for the S' phase is not affected by the addition of
dumina particles. The activation energy for S' formation (obtained for each material
from the dopes of the Arrhenius plots in Figure 4.65) is outside the range of values
reported by previous workers. Jena et al. (108) obtained E = 129.9 kJ/mol for Al-
1.53wt.%Cu-0.79wt.%Mg alloy, Youdelis and Fang (157) reported E = 109.6 I 12.4
kJ/moI, while Luo et al. (104) docur~ented E = 132.6 2 3.4 W/rnol for S' formation in
Al-Li-Cu-Mg alloy. The discrepancy between the results obtained in the present work
and the quoted data is attributed to differences in material composition, mass transfpr
resistance, and environment (matrix environment). Nevertheless, the E values obtained
for S' precipitation in this study are within the range of values reported in reference (17)
(E = 147 kJ/rnol) for PM 2214 aluminum alloy and by Martinod et al. (158) (E = 158
kJ/mol),
For the 8' phase precipitation (the second exothermic reaction under peak C in Figures
4.34 and 4-35)? n = 1.63 for unreinforced 2618 and this compares fairly well with n =
1.52 obtained for the composite. This is an indication that the growth mechanism for the
8' precipitates is unaffected by the addition of -0, particles. The values of the growth
parameter n are within the range of that previously reported (n = 1.68) (157). Papazian
(1 13), using DSC to study 8' formation, reported that the overall process could be
descnbed by the Avrarni equation, with n = 1-1.2 and that the reaction is diffusion
controlled. Chen and Doherty (159) have also proposed that 8' thickening cm also be
described as a volume difision process.
On the b a i s of the growth models proposed by Ham (130), the value of n = 1.52, 1.63 is
indicative of a disc-like growth which is intermediate between constant plate thickness
2 (n = 2) and constant eccentricity (n = - oblate spheroid). The value of the activation
2 ,
energy obtained for unreinforced AM6 18 (Q =133.08 kJ/moI) is outside the activation
energy ranges reported by Aaronson and Laird (160) for 8' thickening (E = 76.6-1 15.5
kJ/mol) and lengthening (E = 79.5-98.6 kVmo1) in Al-4 wt.% Cu. However, it is within
the range of the activation energy for lattice diffusion of Cu in Al (E = 130.616.3
kJ/rnol) reported by Murphy (161). Thomas and King (17) have sumrnarized the data for
elemental diffusion in pure aluminum and simple binary aluminum ailoys (see Table 4.
18).
The activation energy obtained for 8' formation in 26 l8+l5 composite (Q = 123.95
kJ/mol) is somewhere between the values for bdk difision of Mg in Al and grain
boundary difision of Cu in Al published in Table 4.18. The value of E for the MMC is
relatively lower than that for the unreinforced alloy, suggesting that diffusion of 8' is
easier in the MMC than in the unreinforced alloy. There are few published data in this
area with which to compare the current results. Dutta et al. (L39), working on 2014
matrix composites, found the same trend as in the present work. The values of activation
energy they obtained were 82.031, 73.455, and 72.589 kJ/mol for 8' precipitation in
unreinforced AA.20 14, 10 and 15 vol. % alurnina/AA20 14 composites, respective1 y.
Nieh and Karlak (20), working on 6 m matnx MMCs, also found that the activation
energy for diffusion was approximately 37% Iower in the MMC than in the unreinforced
alloy. Nieh and Karlak attributed the reduction in activation energy for diffusion in the
MMC to enhanced diffusion of solute along dislocations to growing transition
precipitates. This could be the case in the current MMC, with additional diffusion along
the Alto3,-matrix interfaces also contributing. As previously mentioned, the composite
has a larger grain boundary area than the unreinforced alloy due to the smaller sub-grain
size. This is likely to contribute to enhanced solute diffusion.
Table 4.18. Diffusion data for aluminum alIoys (Ref. 17)
Type of Diffision E (kJ/mol)
Bulk diffusion - AI in Al 123-126
Bulk diffusion - Cu in Al 126-136
Bulk diffusion - Mg in Al 115-130
Bulk diffusion - Cu in AVCu 120
Bulk diffusion - AI + Cu in Al-Cu 122-130
Grain boundary difision - Al in Al 60-87
Grain b o u n d q difision - Cu in Al 105
Grain boundary diffusion - Cu in Ai-Cu 100
Grain boundary diffusion - Al -t Cu in AI-Cu 84-98
5. CONCLUSIONS AND RECOMMENDATIBNS
In the present work, different experirnental techniques have been used to characterize the
microstructure and behavior of AA26 18 and its composites containing different volume
fractions of Al,O, particles. The metal matrix composites studied contained 10 and 15
vol. % &O3 particles. The evaluation of the microstructure and behavior of the test
materials were carried out using microhardness measurements, differential scanning
calorimetry (DSC), scanning electron microscopy (SEM), electron probe microanalysis
(EPMA), and transmission electron microscopy (TEM). Frorn the results presented and
discussed in the preceding chapter, the following conclusions can be made.
5.1.1 Precipitation Kinetics
1. The addition of Al,Q particles to A M 6 18 does not alter the aging sequence, but it
alters certain aspects of the precipitation reaction. Although aging is accelerated in
the composites, the presence of A1,03 particles in AA26 18 reduces the peak matrix
hardness. This is attributed to the depletion of solute elements such as magnesium.
2. The kinetics of GPB zone formation in the MMC is retarded when compared with
those of the unreinforced alloy. On the other hand, the kinetics of GPB zone
dissolution are more enhanced in the MMC than in the unreinforced alloy. The
addition of N203 particles affects not only the precipitation het ics , but also the
relative arnounts of the various phases present. It reduced the arnount of GPB zones
formed in the MMC relative to the unreinforced ailoy. It is believed that this is due
to the fact that some GPB zones have formed and stabilized prior to the DSC scan.
That is, during quenching from the solution heat treatrnent temperature the A1,0,
particles cool more slowly than the matrix (since the particles have a lower themial
conductivity). This causes the rnatrix around the particles to be warmer than the
bulk matrix. The high dislocation density and the high solubility centers (the warm
parkle-rnatrix interfaces) in the composite are favorable conditions for GPB zone
formation. With a sipfïcant amount of GPB zones formed (in the composite)
during quenching, the subsequent DSC scan yields a srnailer amount of GPB zones
in the composite than in the unreinforced alloy. This mechanism is a possible
explanation for the higher volume fraction of precipitates obtained in the composite
during the endothemiic GPB zone dissolution process than dunng the exothermic
GPB zone formation.
3. The precipitatian and dissolution of GPB zones and the metastable S' and 8'
phases are kinetically controlled in both unreinforced AA26 18 and the composites.
The growth pararneters for the formation of the S' and 8' phases are practicaily the
same for both the monolithic alloy and the composite. Hence, the addition of A1103
particles does not seem to alter the growth mechanisms for the formation of these
metastable phases.
4. The kinetic parameters determined for the formation and dissolution of GPB zones
using the Avrami-Johnson-Mehl equation are, within the lirnits of experimental
errors, in good agreement with those obtained by the Kissinger-like expression.
5. Nucleation of precipitates in a cast AM618 matrix composite takes place on two
different sites, narnely: (1) quenched-in vacancy loops and (2) matrix dislocations
generated as a result of the CTE mismatch between the matrix and the
reinforcement. Increasing reinforcement content increases nucleation on
dislocations but, at the same time, it decreases vacancy concentration. Therefore,
the contribution of quenched-in vacancies to the overall nucleation process is
reduced.
6. A precipitation sequence for AA2618, based on the current DSC investigation, is
proposed as: (i) formation and dissolution of GPB zones; and (ii) concurrent
formation and dissolution of metastable phases (S', 8', and X'). That is:
1. Mg-rich intermetallics, considered to be MgA1204 spinel or magnesia (MgO), are
present in the AA2618-alumina composite systems studied. The are suggested as
one of the main causes of magnesium depletion in the composite matrix. X-ray
maps have revealed Mg accumulation around the reinforcing &O3 particles.
Therefore, it c m be deduced that increasing the reinforcement-rnatrix interfacial
area (i.e., increasing reinforcement content or decreasing reinforcement size) in
Al,O,-AA2618 composites gives rise to increased Mg depletion from the matrix.
This causes a reduction in the overall peak hardness achievable in the composites.
To achieve the same Ievel of hardening in the alumina-AM618 composites as in
the monolithic alloy, the following suggestions are made:
(a) increase the arnount of solutes (especially magnesium) added to the composite
matrix d o y pnor to casting;
(b) coat the N203 particles with a matenal which will neither react with the
solute elements nor alter the matrix-particle interface chemistry (e-g., by
degrading the interface bonding); and
(c) employ short interaction times between N2O3 particles and the reactive
elements in conjunction with Iow fabrication temperatures.
2. Aiuminide particles possess a variety of chernical forrnulae. The formula assumed
by each particle is strongly dependent on the rnatrix environment, especially
aluminurn content. Also the morphology is dependent on the heat treatment history
and thermomechanical processing.
3. The aluminide phase has been detemiined to have a C-centered rnonoclinic
structure with lattice parameters a = 0.867 nrn, b = 0.900 nm, c = 0.859 nrn, and P = 83.504". This structure gives a more consistent indexing than the primitive
monoclinic structure reported in the literatue.
The current study has yielded some important information about the microstructurc and
aging behavior of AA26 18 and its Al1O, particle-reinforced composites. In the course of
the investigation, new areas of research have been identified. These areas need further
investigations.
1. The migration of solute atoms, especially magnesium, during aging needs to be
rnonitored much more closely than has been done in the present study. It is
important to measure the level of solute migration from solution treatrnent to the
overaged condition of the test matenals. Such an investigation will yield
information which is needed to clariQ the mechanisms of solute depletion in the
composite matrix. It will also help in developing a mathematical mode1 for solute
difision in metd matrix composites derived from age hardenable alloys.
2. The occurrence of reaction products in the MMC has been inferred from Aqua
Regia (AR) anaiysis. It is acknowiedged that AR c m digest some species of
insolubIe particies thereby making them unavailable during the subsequent SEM
examination. The matrix-particle interface could not be resolved by the
SEMIEPMA used in the current study. It was also difficult to get a detailed
inspection using the TEM because of the limitation imposed by the small size of
areas inspected. It is recommended that the matrix-particle interface be inspected
using high-resolution scanning electron microscopy (HRSEM) or high-resolution
transmission electron rnicroscopy (HRTEM) to reveal the presence of a reaction
zone and c o n f m the existence of intermetallics such as MgM204 spinel.
3. The role of the X phase in the aging behavior of AA26 18 is not certain. Its relative
distribution in cornparison with other well established age hardening precipitates
could not be deterrnined in the present investigation. Also, there could still be other
precipitate phases in 2618 which are yet to be identified. Therefore, further work is
required in this area.
4. The doublet e x o t h e d c peak produced during the DSC investigation has been
attnbuted to CO-precipitation of the S' and 8' phases. In the current study, a
cornparison of peak reaction temperatures obtained from different DSC studies was
used to identim which peak is due to the formation of S' phase or the 0' phase. A
more accurate identification can be obtained by in situ DSC experiment (especidy
constant-temperature scanning as opposed to varying-temperature scanning
commonly used) or, where possible, electrical resistivity rneasurements during
aging coupled with TEM examination.
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A Aluminum Alloy Designation Systems
The aim of this section is to explain in greater detail the designation systems used for
duminum alloys. The materials here have been sourced from the handbooks ( 1,162) and
special books on afuminum (2,100).
A l Wrought Aluminum Alloy Designation System
A four-digit numencal system is used to identiQ wought and alurninum alloys systems.
As shown in Table Al , the first of the four digits in the designation indicates the major
alloying element of aUoys within the group.
Table A 1. Wrought aluminum alloy designation system.
Designation Main Alloying Element
Alirminum; 2 99.00%
Copper
Manganese
Silicon
Magnesium
Magnesium + silicon
Zinc
Other elements
Unused series
In the lm group, the series lOxs is used to codify unalloyed compositions that have
natural impurity lirnits. The last two of the four digits in the designation indicate the
minimum aluminum percentage. Designations having second digits other than zero (Le.,
integers 1 - 9 assigned consecutively as needed) indicate special control of one or more
individual impurities .
In the 2xxx - 8- alIoy groups, the second digit in the designation indicates alloy
modification. If the second digit is zero, it indicates the original alloy. Integers 1-9,
assigned consecutively, indicate modifications of the original ailoy. The 1 s t two of the
four digits in the 2Kxx - 8:m groups have no special significance. They serve only to
identiw the different aluminurn alloys in the group. For the 2ar - 7xxx senes, the dloy
group is determined by the alloying element present in the greatest mean percentage. The
6 m series is an exception where the proportions of Mg and Si available to form Mg2Si
(magnesium silicide) are predorninant. If the greatest mean percentage is the sarne for
more than one element, the choice of group is in order of group sequence (Le., copper,
manganese, silicon, magnesium, magnesium silicide, zinc, or others).
A2 Cast Aluminum Alloy Systems
Like in the wrought aluminurn alloy systems, a system of four digits numerical
designations incorporating a decimal point is used to identify alurninum and alurninum
alloys in the form of castings and foundry ingot (see Table A2). The first digit indicates
the alloy group. The second digits identim specific aiuminum alloy or. for the alurninum
(1xx.x) senes, indicate purity. The last digit, which is separated from others by a decimal
point, indicates a product form, Le., whether casting or ingot: xxr.0 indicates castings
and xcx. 1 indicates ingot.
Table A2. Cast aluminum alloy designation system.
Main Alloying Element
Aiuminurn; 2 99.00%
Copper
Silicon, with added copper
Silicon
Magnesium
Zinc
Tin
Other elements
Unused senes
a d o r magnesium
B Experirnental Apparatus
The aim of this section is to provide eiiciting sketches of some of the experimental
apparatuses mentioned in Chapter 3. They are presented mainly without m e r
descnp tions .
B1 Differential Scanning Calorimetry @SC)
Differential scanning caiorirnetry (see Figure B.1) is a thermal technique in which
differences in heat flow into a substance and a reference are measured as a function of
sample temperature while the two are subjected to a controlled temperature program.
The fundamental difference between DSC and WerentiaI themal analysis (DTA) is
that the former is a cdorimetric method in which differences in energy are measured. In
contrat, in DTA. differences in temperature are recorded. Nevertheless, the temperature
programs for the two methods are similar.
Readour r
Sample cell
Hea t stiield
Programming 4 dwce 1 Figure B. 1. Schematic diagram of a typical differential scanning calorirnem
173
B2. Opticai and Transmission EIecîron Microscopy
Condenser lem
(photographic plate)
I I I Electron source
%'
Magnetic objective Rp Magnctic projector
Fiuorescent screen (photographic plate)
(a)
Figure B.2. Cornparison of a light and electron microscopes.
B3 Scanning Electron Microscopy
ELECTR,ON GUN 1 FI LAMENT I GUN SUPPLY
SHIELD I f- 1 ANODE
SCANN INGi COILS
Fl RST
n
f FINAL ~ E N S O LENS, 1
TOR
LENS SUPPLY H
1 DISPLAY AND 1
ISIGNALCO,LLE?R IIMAGNIFICATION UNIT
VACUUM SYSTEM
RECORD UNITA '(17 ' SIGNAL
AMPLIFIER
Figure B.3. Schematic diagram of a typical scanning electron microscope.
Energy Dispersive X-ray Spectrometry
(a)
Pulse Energy-to- Preamp processor digital converter
O
4
Sample C 7 9
Vide0 4 - 4 Multi-
Channel analyter
Keyboara C
1 1 . 1 A
Mini- compuier
idow 1
. L f t h i u m D r f f t e d I n t r i n s f c Region .
\.lu i n a c t i v e p S i r eg fpn
Figure B.4. Schematic diagram of: (a) a typical EDS system (b) a typical Si(Li) EDS detector.
6: Theuries of Particle Strengthening
The effect of a dispersed phase on the mechanical properties of an alloy may be
considered in terms of changes both in the yield stress (163-166) and also in the work-
hardening behavior (167). The yield stress is that stress which must be applied to a
crystd to rnove a single dislocation over a distance that is large compared to the
precipitate spacing. In the stress-strain curve, it is defined as the point of departure from
the linearity of the elastic part of the c w e . When a material is permanently (plastically)
deformed, it is found that its yield stress and hardness is increased while its ductility is
duninished. This phenomenon is known as work hardening. The work-hardening
behavior of dispersion-hardened alloys is conveniently approached by dividing the alloy
into two distinct groups. Group 1 alloys are those whose tensile curves of single crystals
have the same shape and charactenstics as the curves of crystals of pure metals and solid
solution alloys. The particles in these alloys deform when the alloy is deformed. Group 2
alloys have single crystal stress-strain curves quite different from those of pure single
crystais. Most alloys containing incoherent particles beiong to this group. In this section,
the interest is in the yield stress-based theories, especially the Orowan theory.
C l Yield without Particle Shear
The smallest radius (p) to which a dislocation of Burger's vector b and line tension T
can be bent by an applied shear stress s i s given by (168):
If the separation of particles in the glide plane is d, then a dislocation would have to be
bent to a radius of the order of 6/2 in order to bow between the particles rather than
shear them. The shear stress to cause this is given by
The preceding paragraph outlines what is known as the Orowan theory of yield of
dispersion-hardened crystals. The increase in the yield stress due to the presence of the
particle is given by equation (C.2), that is, T = - 2T . If T is approximated to be equd to bd
1 - ~b' , equation ((2.2) become 2
so that the yie1d stress of a dispersion-hardened alloy is approximately given by (168)
provided the particles are not sheared at yield. G is the matrix shear modulus and .ri,,
,, is yield stress due to long-range interactioris between dislocations and precipitates
(see Ref. (168) for more details).
C2 Orowan Theory for Ceramic Reinforcement
It has been shown in Section C l that Orowan bypassing of particles by dislocations can
increase the strength of a matenal. Consequently, it has been proposed that reinforcing
ceramic whiskers and particles may result in an Orowan-type strengthening effect (169).
The bypassing stress z for particles with an appreciable aspect ratio can be calculated
fkom the modified Orowan relation (170)
where G is the ma& shear modulus, b is the Burger's vector, v is Poisson's ratio, t is
the particle length, t is the particle width (thickness), ro is the dislocation core radius,
and L is the inter-particle spacing dcfined by
where V, is the particle volume fraction.
