THE EFFECT OF INTERLAYERS ON DlSSlMlLAR FRICTION WELD PROPERTIES
Cuauhtemoc Maldonado-Zepeda
A thesis subrnitted in conformity with the requirements for the degree of Doctor in Applied Science Graduate Department of Metallurgy & Materials Science University of Toronto
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The author retains ownership of the L'auteur conserve la propriété du copyright in this thesis. Neither the droit d'auteur qui protège cette thèse. thesis nor substantial extracts f?om it Ni la thèse ni des extraits substantiels may be printed or otheTWise de celle-ci ne doivent être imprimés reproduced without the author's ou autrement reproduits sans son permission. autorisation. THE EFFECT OF INTERLAYERS ON DlSSlMlLAR FRICTION WELD
PROPERTIES
Doctor in Applied Science (2001)
Cuauhtemoc Maldonado-Zepeda
Graduate Department of Metallurgy & Materials Science
University of Toronto
ABSTUACT
The influence of silver interlayers on the metallurgical and mechanical properties of dissimilar aluminium alloy/stainless steel friction welds are investigated.
An elastic contact model is proposed that explains the conditions at and close to the contact surface, which produce A1D3 particle fracture in dissimilar MMCIAlSl 304 stainless steel friction welds.
lntermixed (IM) and particle dispersed (PD) regions are forrned in Ag-containing dissimilar friction welds. These regions forrn very early in the joining operation and both contain AgdI. Therefsre, an interlayer (Ag) introduced with the specific aim of preventing Fe,AI, compound formation in MMCIAISI 304 stainless steel friction welds promotes the formation of anofher intermetallic phase at the bondline. Since IM and
PD regions are progressively removed as the friction welding operation proceeds thinner interrnetallic layers are pmduced when long friction welding times are applied.
ii This type of behavior is quite different from that observed in silver-free dissimilar
MMWAISI 304 stainless steel welds.
Nanoparticles of silver are formed in dissimilar MMC/Ag/AISI 304 stainless steel
welds produced using low friction pressures. Nanoparticle formation in dissimilar friction
welds has never been previously observed or investigated.
The introduction of silver interlayers decreases heat generation during welding,
produces narrower softened zone regions and improved notch tensile strength
properties. Ali research to-date has assurned per se that joint mechanical properties
wholly depend on the mechanical properties and width of the intemetallic layer fomed
at the dissirnilar joint interface. However, it is shown in this thesis that the mechanical
properties of MMC/AISI 304 stainless steel joints are determined by the combined
effects of interrnetallic formation at the bondline and softened zone formation in MMC
base maten'al immediately adjacent to the joint interface. A methodology for calculating the notch tensile strength properties of dissimilar friction welds is presented and is
based on a combination of FEM with a ductile failure criterion. There is excellent correspondence behveen actual and calculated joint strength results.
iii ACKNOWLEDGEMENTS
Doy las gracias mas profundas por el invaluable apoyo recibido del pueblo de
México a través dei Consejo Nacional de Ciencia y Tecnologia (CONACYT) y la
Universidad Michoacana de San Nicolas de Hidalgo (UMSNH). ABSTRACT ...... II
ACKNOWLEDGEMENTS ...... IV
TABLE OF CONTENTS ...... v
LIST OF TABLES ...... xi
LIST OF FIGURES...... XII
CHAPTER 1. THESIS PROPOSAL...... 1
1. 1. INTRODUCTION...... 1
1.2. THE INTERLAYER APPROACH ...... 2
1.3. DETAILED THESIS OBJECTIVES...... 6
1.4. THE PROCEDURES ...... 7
1.5. THESIS ORGANIZATION ...... 7
CHAPTER 2. THE FRICTION WELDING PROCESS ...... 9
2.1. INTRODUCTION...... 9
2-1.1. lnertia friction welding ...... 1 O
2.1.2. Direct-drive fmion welding ...... 12
2.1 .2. i. Heating period ...... 12 v 2.1 .2.2. Forging stage ...... 16
2.1.2.3. Direct-drive friction welding parameten ...... ~...... 17
2.2. INTERMETALLIC COMPOUNDS IN DISSlMlLAR JOtNlNG ...... 18
2.2.1 . Fe-AI intermetallic compounds ...... 18
2.2.2. Ni-AI intemetallic compounds ...... 20
2.2.3 .Ag-AI intermetallic compounds ...... 21
2.2.5. Influence of composition on the intermetallic layer-...... 23
2.2.6. lntermetallic compounds and mechanical properties ...... 25
2.3. FRICTION AND WEAR ...... 27
2.3.7. SoR coatings and the coetficient of fn'cfion ...... 27
2.3.2. Wear and friction welding ...... 30
2.4. THERMAL ASPECTS OF FRICTION WELDING ...... 31
2.4.7. Calculafion of heat input in friction welding ...... 31
2.4.2. Temperature distribution during friction welding...... 33
2.5. WELDING METALLURGY OF ALUMINUM ALLOYS ...... 34
2.5.7. AI-Mg-Si Alloys ...... 34
2.5.2 . Effect of the thermal welding cycle on dissolution and reprecipitation...... 34
2.5.3. The softened zone ...... ,,...... 36
2.6. MECHANICS OF SLlDlNG CONTACT ...... 38
2.6.1. Line loading of a semi-infinite half-space ...... 39
2.6.2. Torsional loading ...... ,...... 40
2.6.3. Normal pressure and torsional loading...... 41
2.6.4. Contact of rigid-ideally-plastic materials ...... 44
2.6.4.1. Combined effect of shear and pressure on a plastic surface ...... 45 vi 2.6.4.2. The wave model ...... 46
2.6.4.3. Wave removal model...... 48
CHAPTER 3. EXPERIMENTAL PROCEDURES ...... 51
3.1. MATERIALS...... 51
3.2. FRICTION JOINING ...... 54
3.3. METALLOGRAPHIC EXAMINATION ...... 55
3.4. MECHANICAL TESTING ...... 58
CHAPTER 4 . INTERLAYERS AND PARTICLE FRACTURE IN DlSSlMlLAR FRICTION JOINTS ...... 61
4.1. INTRODUCTION...... ,...... 61
4.2. RESULTS ...... 62
4.2.1. influence of friction welding on particle fracture ...... 62
4.2.2. Influence of friction pressure ...... 64
4.2.3. Influence of friction time ...... ,...... 64
4.2 .4. lnterlayers and parficle fracture ...... 68
4.2.5. Influence of silver intelayer on particle accumulation ...... 69
4.3. DISCUSSION...... 71
4.3.1. Friction welding and reinforcement particle morphology ...... 71
4.3.2. Effect of friction pressure ...... 72
4.3.3. Influence of fn'cion time ...... 72
4.3.4. Modelling of pmtruding parficle fracture ...... 74
4.3.5. Retention of fractured particles at the bondline ...... 78 vii 4.3.6. Efiect of the interiayer on the stress distribution...... 80
4.3.7. Summation ...... 83
CHAPTER 5 . MICROSTRUCTURE OF A DISSIMILAR FRICTION JOINT CONTAlNlNG A SlLVER INTERLAYER ...... 84
5.1. INTRODUCTION...... 84
5.2. RESULTS ...... 87
5.2.1. lnterlayer microstructure ...... 87
5.2.2. As-welded joint rnicmst~cfure...... 88
5.2.3. Friction pressure and TEM micmstmcture ...... 102
5.3. DISCUSSION ...... 121
5.3.7. Influence of friction welding on the silver intedayer ...... 121
5.3.2. Similanfies befween sliding Wear and Stage 1 of friction welding...... 122
5.3.3. Formation and removal of PD and IM mgions ...... 125
5.3.4. Transition layer in dissimilar MMC/AISI 304 stainless steel welds ...... 129
5.3.5. Intennetallic compounds and joint strength ...... 129
5.3.6. Mode1 for interlayer rernoval ...... 131
5.3.7. Surnmation ...... 134
CHAPTER 6 .THERMAL ASPECTS AND SOFTENED ZONE FORMATION...... 136
6.1. INTRODUCTION...... 136
6.2. RESULTS ...... 140
6.2.1 . Effect of friction pressure ...... 140
viii 6.2. 2. Influence of the silver interlayeri...... 141
6.3. DISCUSSION...... 143
Heat entering the stainless steel substrate...... 143
Heat paitition during dissimilar friction welding ...... 145
Friction pressure and peak temperafure ...... 146
Influence of the silver interlayer...... 148
Summation ...... 149
CHAPTER 7. INFLUENCE OF THE INTERLAYER ON THE MECHANICAL PROPERTIES OF DISSIMILAR FRICTION JOINTS ...... 155
7.1. INTRODUCTION...... 155
7.2. RESULTS ...... 156
7.2.1. Joint mechanical properfrés ...... 156
7.3. CALCULATING THE TENSILE STRENGTH OF DlSSlMllAR FRICTION
JOINTS ...... 162
7.3.1. Introduction ...... 162
7.3.2. The notched fensile specimen ...... 164
7.3.3. Finite element mode1 of the notched tensile specimen ...... 168
7.3.4. Stress and strain distribution in notched tensile specimens ...... 171
7.3.5. Calculating the notch tensile strength of dissimilar fncton welds ...... 175
7.4. DISCUSSION ...... 179
7.4.1. lntermetallic compound formation and joint tensile strength ...... 179
7.4.2. Soffened zone and joint tensile strength ...... 190
7.4.3. Failure mode in dissimilar fricion joints ...... 192
ix 7.4.4. FEM analysis of notched tensile test specimens ...... 194
7.4.5. The notch tensile strength of the dissmilar friction joints ...... 195
7.4.7. Summation ...... 196
CHAPTER 8. CONCLUSIONS ...... 197
REFERENCES ...... I...... 201 LIST OF TABLES
CHAPTER 3. EXPERIMENTAL PROCEDURES ...... c...... 51
Table 3.1. Chernical composition of materials (wt%) ...... 53
Table 3.2. Geometric characteristics of reinforcing particles...... 53
Table 3.3. Mechanical properties of materials ...... 59
CHAPTER 6 .THERMAL ASPECTS AND SOFTENED ZONE FORMATION...... 136
Table 6.1 . Thermal properties of base materials ...... 138
Table 6.2. Relationship between specific heat input entering the stainless steel
substrate (q&) and friction pressure ...... 145 LIST OF FIGURES
CHAPTER 1. THESIS PROPOSAL ...... 1
Figure 1-1. The effects of an interlayer on dissirnilar friction welds...... 8
CHAPTER 2. THE FRICTION WELDING PROCESS ...... 9
Figure 2.1 . lnertia friction welding process (Wang and Rasmussen 1972)...... 10
Figure 2.2. Directdrive friction welding process...... 13
Figure 2.3. Idealised traces of the variations with time of speed, torque and axial
shortening in the direct-drive friction welding process during the frictioning stage. 1-
Peak torque. 2-Equilibrium torque. 3-Terminal torque (Duffin and Bahrani l976).. 14
Figure 2.4. The heating period based on the work of DuMn and Bahrani (Duffin and
Bahrani 1976)...... 1S
Figure 2.5. FeAl binary diagram (Kattner et al. 1987)...... 19
Figure 2.6. AI-Ni binary diagram (Hultgren et al. 1973)...... 20
Figure 2.7. AgAl binary diagram (McAlister 1987)...... 22
Figure 2.8. Effect of holding time at 823 and 873 K on the tensile strength of Ti/(6xlo4
wt% Si) Al and Til(0.12 wt% Si) Al joints O 873 K, Til(0.12 wt% Si) Al; O 873 K, TV(6x1 o4 wt% Si) Al; 0 823 K, Til(6xl o4 wt% Si); * fractured at the interface (Fuji et al. 1995~)...... 24 xi i Figure 2.9. Thickness of the intemetallic layer (Ti3AI) vs. heating time in a TiAI
diffusion joint (Suzuki et al. 1994)...... ~~..~...... ~~~...... ~..~~...... ~...~~~...... 26
Figure 2.10. Joint strength vs. heating time in a TVA1 diffusion joint (Suzuki et a/. 1994).
..,...... ~...... ~....~~...... 27
Figure 2.1 1. The dependence of the strength of the weld between stainless steel and
1100 aluminium alloy on the total thickness of the intermetallic layer (Calderon et al.
1985)...... --...... + ...... 28
Figure 2.12. Examples of dependence of the coefficient of friction of p for a soft coating
on harder substrate materials: (A) pLfKB>l-3; (B) 1.32~L/p B21 ; (C) p L/p B >1; (D)
pL/pB4(T = thickness) (Heilmann and Rigney 1981)...... 30
Figure 2.13. Schematic arrangement of friction welding of a solid rod (Grong 1994)... 33
Figure 2.14. Schematic diagram showing the hardness distribution following P"(Mg2Si)
dissolution in the HAZ of 6082-T6 aluminium welds (Grong 1994)...... 35
Figure 2.15. Precipitation of P' (Mg2Si) dispersoids during the weld cooling cycle (Myhr
and Grong 1991a)...... -...... -.---.-..-.36
Figure 2.16. Schematic representation of the HAZ hardness distribution after welding
and subsequent natural ageing. (A) Short duration thermal cycle. (B) Long duration
thermal cycle (Midling and Grong 1994b)...... -...... ~...... 38
Figure 2.17. Line loading of a semi-infinite space...... ~.~..~...... 39
Figure 2.18. Hemispherical normal pressure distribution (Heteny and McDonald 1954).
...... ~...... ~...... ~...... ~~...... 41
Figure 2.1 9. Displacement components at the surface (Heteny and McDonald 1954). 42
Figure 2.20. Contact area growth of a plastic wedge under the action of a constant
normal load P and an increasing tangential load Q (Johnson 1985)...... 45
xiii Figure 2.21. Wave formation mode1 (Challen and Oxley 1979)...... 47
Figure 2.22. Theoretical and experimental results for the coefficient of friction
(Kopalinsky and Oxley 1995)...... 48
Figure 2.23. Wave removal slip-line model (Challen and Oxley 1979)...... 49
Figure 2.24. Sirnplified wave model used in calculating strains (Kopalinsky and Oxley
1995) ...... 50
CHAPTER 3. EXPERIMENTAL PROCEDURES ...... 51
Figure 3.1. An optical micrograph showing the microstructure the reinforced 6061 AI-10
vol% A1203-...... 52
Figure 3.2. A backscattered SEM micrograph showing the silver interlayer at the
bondline of dissimilar MMCfAgfAISI 304 stainless steel friction joints ...... 53
Figure 3.3. Some examples illustrating the relationship behiveen (a) regular particles,
(b) irregu lar particles, and (c) misaligned particles, and the cylindrical particles
which were used to represent thern assurning the bar axis is from left to right (Lewis
and Withers 1995)...... 57
Figure 3.4. U-Notch tensile testing specimen configuration...... 58
Figure 3.5. Round test tension specimen ...... 59
Figure 3.6. Stress-strain diagram for reinforced 6061-T6110 vol% AI2O3...... 60
Figure 3.7. Stress-strain diagram for AlSI 304 stainless steel ...... 60
CHAPTER 4 . INTERLAYERS AND PARTICLE FRACTURE IN DlSSlMlLAR FRICTION JOINTS ...... 61
xiv Figure 4.1 Modes of particle failure observed dunng mechanical testing of a particulate reinforced metal matnx composite materiai...... 62
Figure 4.2. Particle fracture in dissimilar MMWAISI 304 stainless steel friction welds. 63
Figure 4.3. Effect of friction joining on the aspect ratio distribution in a 6061-1O vol%
AI2O3MMC/AISI 304 friction joint. Friction pressure, 120 MPa; forging pressure,
120 MPa; friction time, 4.5 s; forging time 1.5 S. (A) As-weld condition; (B) As-
received condition...... 65
Figure 4.4. Effect of friction joining on the number of same-size particles in 6061-10
vol% AI2O3 MMC/AISI 304 friction joints. Friction pressure, 120 MPa; forging
pressure, 120 MPa; friction time, 4.5 s; forging time, 1.5 S. (A) As-weld condition;
(B) As-received condition...... 66
Figure 4.5. Effect of friction joining on the particle volume in a 6061-10 vol% AI2O3
MMCIAISI 304 friction joint. Friction pressure, 120 MPa; forging pressure, 120 MPa;
friction tirne, 4.5 s; forging time, 1.5 S. (A) As-weld condition; (6)As-received
condition...... 67
Figure 4.6. Effect of friction pressure and forging pressure on the average particle
radius in the region adjacent to the bondline in 6061-10 vol% AI203/1020 mild steel
friction joint. for a friction tirne of 4.5 s, and a forging time of 1.5 S. All
measurernents were made at the half radius location...... 68
Figure 4.7. Effect of friction tirne on the percentage of fractured particles in MMCIAISI
304 stainless steel and MMCIAglAISI 304 stainless steel friction joints. Friction
pressure, 120 MPa; forging pressure, 30 MPa; friction time from 0.2-4.5s; forging
time, 1.5 S. Al1 measurements were made at the half radius location...... 69
Figure 4.8. Effect of a silver interlayer on the average particle radius in the region
adjacent to the bondline in an MMCIAISI 304 stainless steel friction joint. Friction XV pressure, 120 MPa; forging pressure, 120 MPa; friction time, 4.5 s; forging time, 1.5
S. All measurements were made at the haif radius location...... 70
Figure 4.9. Accumulated particle volume in dissimilar 6061-10vol% AI2O3MMCIAISI
304 stainless steel and MMC/Ag/AISI 304 friction joints. Friction time, 4 s; forging
pressure, 240 MPa; forging time, 1 S. Ail measurements were made at the half
radius location...... 71
Figure 4.10. Regular wavy surface model...... 75
Figure 4.1 1. Effect of the combined elastic modulus (E*)on the contact pressure @,).
Assuming a nominal pressure, p. of 0.5 MPa, an interparticle distance (A) of 104.8
pm, and protuberance (A) of 1 .O pm...... 78
Figure 4.12. Effect of oxide layen on the notch tensile strength of MMCIAISI 304 . . -- stainless steel friction joints...... 79
Figure 4.13. Distribution of normal stress (ax)produced by the combination of normal
and tangential line loads along the contact surface (y = O unit of length). P = 1 unit
of force per unit of length...... 81
Figure 4.14. Distribution of normal stress (fi)produced by the combination of normal
and tangential line loads along the line at y = 0.1 unit of length. P = 1 unit of force
per unit of length...... 82
CHAPTER 5. MICROSTRUCTURE OF A DISSIMILAR FRICTION JOINT
CONTAlNlNG A SILVER INTERLAYER ...... 84
Figure 5.1. TEM micrograph of silver interlayer: (A) bright field image; (B)
corresponding selected area diffraction pattern, and (C) key diagram showing Ag.
xvi Figure 5.2. Optical micrograph of the bondline of a dissirnilar MMC/Ag/AISI 304 friction
joint, showing the wavy interface rnorphology produced by the friction welding
operation...... ~...... ~...... ~...... 89
Figure 5.3. Backscattered rnicrograph of a dissimilar MMC/Ag/AISI 304 stainless steel
friction joint showing the intermixed region (IM) containing alurnina particles. The
nickel layer is not visible in this micrograph...... ~...... 90
Figure 5.4. EDX analysis of the intemetallic compound of the IM region. Friction
pressure, 240 MPa; forging pressure, 240 MPa; friction time, 1 S. Location at 1 mm
from the joint periphery. See Figure 5.3...... 91
Figure 5.5. TEM micrograph of the IM region: (A) bright field image, the dark phase is
AgdI; (B) corresponding selected area diffraction pattern, and (C) key diagram
confirming Ag3AI formation in the IM region. Friction Pressure, 90 MPa; forging
pressure, 15 MPa; friction time, 1.2 S. Half-radius location ...... 92
Figure 5.6. X-ray diffraction pattern at the bondline of a dissirnilar MMC/Ag/AISI 304
stainless steel joint. Friction pressure, 90 MPa; forging pressure, 15 MPa; friction
time, 1.2 S...... 93
Figure 5.7. PD region in dissimilar MMC/Ag/AISI 304 comprising particles of AgAI, Ag,
and Ai as identified by EDX. Friction pressure, 90 MPa; forging pressure, 15 MPa;
friction time, 0.2 S. Half-radius location...... ~...... ~...... 94
Figure 5.8. Optical micrograph showing the presence of a PD region in a
6061/MMC/AISI 304 stainless steel friction joint. The arrow shows the joint
centreline...... 95
Figure 5.9. Effect of friction tirne on the microstructure of a dissirnilar MMC/Ag/AISI 304
stainless steel friction joint: (A) friction time, 0.2 s, and (B) friction time, 1.2 S.
xvii Friction pressure, 90 MPa; forging pressure, 15 MPa; forging tirne, 1.0 S. and
rotational speed, 1500 rpm. Backscattered micrograph. Centreline location ...... 96
Figure 5.10. Effect of friction time on the thickness of the PD and IM regions in
dissimilar MMC/AISI 304 stainless steel friction joints. Friction pressure, 90 MPa;
forging pressure, 15 MPa; forging time, 1.0 s, and rotational speed, 1500 rprn. Half-
radius location ...... 97
Figure 5.1 1. Distribution of IM and PD regions along the bondline of dissimilar
MMC/Ag/AISI 304 stainless steel friction joints as function of friction tirne. A)
Friction time, 0.2 s; 5) friction time, 1-2 S...... 98
Figure 5.12. Effect of friction pressure on the microstructure of dissirnilar MMC/Ag/AISI
304 stainless steel friction joints: A) friction pressure, 30 MPa, and B) friction
pressure, 240 MPa. Friction tirne, 4s; forging time, 1 s, and rotational speed, 1500
rpm. Backscattered micrographs. Centreline location ...... 99
Figure 5.1 3. Effect of friction pressure on the thickness of the PD and IM regions in
dissimilar MMC/Ag/AISI 304 stainless steel friction joints. Friction time, 4s; forging
pressure, 240 MPa; forging tirne, s, and rotational speed, 1500 rpm. Half-radius
location...... 100
Figure 5.14. Distribution of IM and PD regions along the bondline of dissirnilar
MMC/Ag/AISI 304 stainless steel friction joints as function of friction pressure. (A)
friction pressure, 30 MPa, and (6)friction pressure, 240 MPa with a small amount
of IM region (1 prn thick) at the centreline location...... 101
Figure 5.1 5. TEM micrographs of the iM region from the bondiine of a dissimilar
MMC/Ag/AISI 304 stainless steel friction joint: (A)-(B) bright field images; (C)
corresponding ring diffraction pattern, and (D) key diagrarn showing the presence of silver. Friction pressure, 30 MPa; forging pressure, 240 MPa; friction time, 4 s;
rotational speed, 1500 rpm. Centreline location...... 104
Figure 5.1 6. TEM micrograph of the IM region: (A) bright field image; (9) corresponding
selected area diffraction pattern, and (C) key diagram showing Ag. Friction
pressure, 120 MPa; forging pressure, 240 MPa; friction tirne, 4 s; rotational speed,
1500 rpm. Half-radius location...... 106
Figure 5.1 7. TEM micrographs of the IM region: (A)-(6) bright field images; (C)
corresponding selected area diffraction pattern, and (D) key diagram showing Ag.
Friction pressure, 240 MPa; forging pressure, 240 MPa; friction time, 4 s; rotational
speed, 1500 rpm. Centreline location...... 108
Figure 5.18. TEM micrographs of the IM region: (A)-(B) bright field images. Friction
pressure, 240 MPa; forging pressure, 240 MPa; friction time, 4 s; rotational speed,
1500 rpm. Centreline location...... 109
Figure 5.1 9. Microstructural features detected in a dissimilar MMCIAgfAISI 304
stainless steel friction joint showing the boundary between the nickel layer and the
stainless steel substrate. Friction pressure, 240 MPa; forging pressure, 240 MPa;
friction tirne, 4 s; forging time, 1 S. The arrow indicates the particle from which the
EDX pattern was obtained. Centreline location...... 111
Figure 5.20. Micrograph of a dissirnilar MMC/AISI 304 stainless steel friction joint
showing the transition layer. Friction pressure, 30 MPa; friction time, 4 s; forging
pressure, 30 MPa; forging time, 1 S...... 112
Figure 5.21. Cracking at the half-radius location in a dissimilar MMClAlSi 304 stainless
steel joint. Friction pressure, 30 MPa; friction time, 1 s; forging pressure, 30 MPa;
forging tirne, 1 s ...... 113
xix Figure 5.22. EDX analysis of the transition layer. Friction pressure, 30 MPa; friction
time, 4 s; forging pressure, 30 MPa; forging time, 1 S...... 114
Figure 5.23. Transition layer distribution along the interface...... 115
Figure 5.24. TEM micrographs at the bondline of a MMCIAISI 304 stainless steel
friction weld: (A)-(B) bright field images; (C) corresponding selected area diffraction
pattern, and (D) key diagram confirming FeAl formation. Interrnetallic layer width,
1.9 Fm. Friction pressure, 30 MPa; forging pressure, 240 MPa; friction time, 4 s;
rotational speed, 1500 rpm. Centreline location...... 116
Figure 5.25. TEM micrographs at the bondline of a MMCIAISI 304 stainless steel
friction weld: (A)-(8) bright field images; (C) corresponding selected area diffraction
pattern, and (C) key diagram confirming Fe2AI5formation. Intermetallic layer width ,
0.5 p.Friction pressure, 240 MPa; forging pressure, 240 MPa; friction time, 4 s;
rotational speed, 1500 rpm. Centreline location...... 117
Figure 5.26. TEM micrographs at the bondline of a MMClA1SI 304 stainless steel
friction weld: (A) bright field images; (C) corresponding selected area diffraction
pattern, and (C) key diagrarn confiming Fe4AIl3 formation. Friction pressure, 30
MPa; forging pressure, 240 MPa; friction tirne, 4 s; rotational speed, 1500 rpm. Half
radius location...... 11 8
Figure 5.27. Comparison of intermetallic layer widths at the centreline and at the half
radius location in the dissimilar joint produced using friction pressure of 30 MPa and
friction time of 4 s; (A) at the half radius location, and (B) at the centreline...... 119
Figure 5.28. Effect of friction pressure (Pl) on the interrnetallic layer thickness in a
dissimilar MMCIAISI 304 stainless steel friction joint. Friction time, 4 s; forging tirne,
1 s, and rotational speed, 1500 rpm...... 120 Figure 5.29. Schematic representation of the wave model when f = 0.90...... 133
Figure 5.30. Schematic representation of the wave model when f = 0.1 0...... 134
CHAPTER 6. THERMAL ASPECTS AND SOFTENED ZONE FORMATION...... 136
Figure 6.1. Effect of friction pressure on the temperature at 0.2 mm from the bondline
and 4.5 mm from the periphery in dissimilar MMC/Ag/AISI 304 stainless steel
friction welds. Friction time, 3 s; forging pressure, 240 MPa; forging time, 1 S..... 139
Figure 6.2. Effect of friction pressure on the temperature at 0.2 mm from the bondline
and 4.5 mm from the periphery in dissimilar MMCIAISI 304 stainless steel friction
welds. Friction time, 2 s; forging pressure, 240 MPa; forging time, 1 S...... 140
Figure 6.3. Effect of friction pressure on the temperature at 0.2 mm from the bondline
and 4.5 mm from the periphery in dissimilar MMC/Ag/AISI 304 and MMC/AISI 304
stainless steel friction welds considering a friction time of 1 s ...... 141
Figure 6.4. Effect of friction pressure on the temperature-time curves at 0.2 mm from
the joint interface and 4.5 mm from the joint periphery in dissimilar MMC/Ag/AISI
304 stainless steel friction joints. Friction time, 3 s; forging pressure, 240 MPa;
forging time, 1 s...... 142
Figure 6.5. Effect of friction pressure on the temperature-time curve at 0.2 mm from the
joint interface and 4.5 mm from the joint periphery in dissimilar MMC/AISI 304
stainless steel friction joints. Friction time, 2 s; forging pressure, 240 MPa; forging
time, 1 S...... 143
Figure 6.6. Relationship between axial shortening and friction pressure in dissimilar
MMC/AISI 304 and MMC/Ag/AISI 304 stainless steel friction joints. Friction time, 4
s; forging time, 1 s; forging pressure, 240 MPa ...... 144
xxi Figure 6.7.Vickers hardness distribution across the joint interface. Friction pressure, 30
MPa; friction time, 4 s; forging pressure, 240 MPa; forging time, 1 S...... 150
Figure 6.8. Vickers hardness distribution across the joint interface. Friction pressure, 60
MPa; friction time, 4 s; forging pressure, 240 MPa; forging time, 1 S...... ,...... 151
Figure 6.9. Vickers hardness distribution across the joint interface. Friction pressure,
120 MPa; friction tirne, 4 s; forging pressure, 240 MPa; forging tirne, 1 S...... 152
Figure 6.1 0.Vickers hardness distribution across the joint interface. Friction pressure,
180 MPa; friction time, 4 s; forging pressure, 240 MPa; forging time, 1 S...... 153
Figure 6.1 1. Vickers hardness distribution across the joint interface. Friction pressure,
240 MPa; friction time, 4 s; forging pressure, 240 MPa; forging time, 1 S...... 154
CHAPTER 7. INFLUENCE OF THE INTERLAVER ON THE MECHANICAL
PROPERTIES OF DlSSlMlLAR FRICTION JOINTS ...... 155
Figure 7.1. Influence of friction pressure on the notch tensile strength of MMCIAISI 304
and MMC/Ag/AISI 304 stainless steel friction joints. Friction tirne. 4s; forging
pressure, 240 MPa; forging time, 1s ...... 156
Figure 7.2. Relationship between the effective plastic strain and notch tensile strength
of MMCIAISI 304 and MMC/Ag/AISI 304 stainless steel friction joints. Friction time,
4 s; forging pressure, 240 MPa; forging tirne, 1 s ...... 157
Figure 7.3. Fracture surface morphology in MMC/Ag/AISI 304 stainless steel friction
joints: A), 6) friction pressure, 30 MPa; C, D) friction pressure, 240 MPa. Friction
tirne, 4 s; forging pressure, 240 MPa; forging time, 1 s; rotational speed, 1500 rpm.
xxii Figure 7.4. Fracture surface morphology in MMC/AISI 304 stainless steel friction joints:
A), 6)friction pressure, 30 MPa; C, D) friction pressure, 240 MPa. Friction tirne, 4 s;
forging pressure, 240 MPa; forging time, 1 s; rotational speed, 1500 rpm...... 160
Figure 7.5. lnfluence of friction pressure on the percentage fraction of ductile failure.
Friction time, 4 s; forging pressure, 240 MPa; forging tirne, 1 s; rotational speed,
1500 rpm ...... 161
Figure 7.6. Fracture surface morphology in a MMCfAgfAISI 304 stainless steel friction
joint. Friction pressure, 240 MPa; friction time, 2 s; forging pressure, 240 MPa;
forging time, 1 s; rotational speed, 1500 rpm...... 162
Figure 7.7. Influence of friction pressure on the softened zone width in MMC/AISI 304
and MMC/Ag/AISI 304 stainless steel friction joints. Friction time, 4 s; forging
pressure, 240 MPa; forging time, 1 s; rotational speed, 1500 rpm ...... 163
Figure 7.8. Relation between softened zone width and the notch tensile strength of
dissimilar MMCfAISI 304 and MMC/Ag/AISI 304 stainless steel friction joints.
Friction time, 4 s; forging pressure, 240 MPa; forging time, 1 s; rotational speed,
1500 rpm...... 164
Figure 7.9. lnfluence of friction pressure on the hardness of the adjoining MMC
substrate in dissimilar MMC/AISI 304 stainless steel and MMC/Ag/AISI 304
stainless steel friction joints. Measured at the location 0.125 mm from the bondline.
Friction time, 4 s; forging pressure, 240 MPa; forging time, 1 s; rotational speed,
1500 rpm...... 165
Figure 7.10. Relationship between the hardness of the adjoining MMC substrate and
the notch tensile strength in dissimilar MMCfAISI 304 stainless steel and
MMC/Ag/AISI 304 stainless steel friction joints. Measured at the location 0.125 mm
xxiii from the bondline. Friction time, 4 s; forging pressure, 240 MPa; forging time, 1 s;
rotational speed, 1500 rpm...... 166
Figure 7.1 1. The necked region of a hornogeneous notched tensile specimen showing
the directions of the normal (Q), radial (O,), and tangential stress (oe)...... 167
Figure 7.12. Half of the notch tensile specimen...... 169
Figure 7.1 3. Finite element idealisation of the notch tensile specimen...... 170
Figure 7.14. Equivalent stress distribution along the Iine located at 1 mm from the
bondline for two MMCIAISI 304 welds stainless steel friction joints. For an applied
stress = 160 MPa...... 172
Figure 7.1 5. Total equivalent strain distribution along the line located at 1 mm from the
bondline for two MMC/AISI 304 stainless steel friction joints. For an applied stress =
160 MPa ...... 173
Figure 7.16. Triaxiality factor distribution along a Iine located at 1 mm from the bondline
for two MMC/AISI 304 stainless steel friction joints. For an applied stress = 160
MPa ...... ~~~...... ~...... 174
Figure 7.17. Equivalent stress distribution along Iines parallel to the bondline in a
dissimilar MMCIAISI 304 stainless steel friction joint. Pl= 240 MPa. Sofîened zone
width = 4.80 mm. For an applied stress = 200 MPa...... 175
Figure 7.18. Total equivalent strain distribution along lines parallel to the bondline in a
dissimilar MMC/AISI 304 stainless steel friction joint. Pl= 240 MPa and softened
zone width = 4.80 mm. For an applied stress = 200 MPa...... 176
Figure 7.19. Triaxiality factor distribution along lines parallel to the bondline in a
dissimilar MMC/AISI 304 stainless steel friction joint. Pl= 240 MPa and softened
zone width = 4.80 mm. For an applied pressure = 200 MPa...... 177
xxiv Figure 7.20. Stress and triaxiality factor distributions along a line located at 0.1 25 mm
from the bondline. Applied stress = 204 MPa. Notch tensile strength = 293.8 MPa.
...... 181
Figure 7.21. Stress and triaxiality factor distributions along a line located at 0.125 mm
from the bondline. Applied stress = 232 MPa. Notch tensile strength = 334.08 MPa.
...... 182
Figure 7.22. Stress and triaxiality factor distributions along a line located at 0.125 mm
from the bondline. Applied stress = 240 MPa. Notch tensile strength = 345.6 MPa.
...... 183
Figure 7.23. Stress and triaxiality factor distributions along a line located at 0.125 mm
from the bondline. Applied stress = 152 MPa. Notch tensile strength = 21 8.9 MPa.
...... 184
Figure 7.24. Stress and triaxiality factor distributions along a line located at 0.1 25 mm
from the bondline. Applied stress = 226 MPa. Calculated notch tensile strength =
325.44 MPa...... 185
Figure 7.25. Stress and triaxiality factor distributions along a line located at 0.125 mm
from the bondline. Applied stress = 260 MPa. Calculated notch tensile strength =
374.4 MPa...... 186
Figure 7.26.Calculated and measured notch tensile strength values in dissimilar
MMC/Ag/AISI 304 stainless steel friction joints. Friction time, 4 s; forging pressure,
240 MPa; forging time, 1 s; rotational speed, 1500 rpm...... 187
Figure 7.27. Calculated and measured notch tensile strength values in dissimilar
MMCfAISI 304 stainless steel friction joints. Friction time, 4 s; forging pressure, 240
MPa; forging tirne, 1 s; rotational speed, 1500 rpm...... 188 CHAPTER 1. f HESlS PROPOSAL
1.l . INTRODUCTION Dissimilar joining involves the combination of base materials having different thermophysical, mechanical, and chernical properties. If a traditional fabrication rnethodology such as fusion welding is applied to join aluminium alloy and steel base materials, for example, the mechanical properties of cornpleted welds will be extremely poor. Firstly, aluminium alloys and steel have quite different melting temperatures which makes the fusion welding process difficult since there will be a strong tendency to preferential melting of the low melting point base material. Secondly, the fusion welding operation will produce very thick brittle AI-Fe intermetallic layers at the bondline because the diffusion rates in the liquid are considerable. Also, an aluminium alloy base material has a higher coefficient of thermal expansion and during cooling following the welding operation, thermal stresses will be generated that are high enough to promote joint failure.
Solid-state joining provides an alternative to fusion joining. In this case the bonding occurs wholly in the solid state at temperatures lower than the solidus temperature of the base material. Typical solid-state bonding processes comprise diffusion bonding, friction welding, and ultrasonic bonding. Even when these welding processes are applied, the formation of brittle intermetallic is still of concern. Diffusion bonding and friction welding can be applied directly or with the introduction of an interlayer material at the contact interface (Elliot and Wallach 1981a, b). Noble or 1 2 reactive interiayer materials have been applied during joining of dissimilar materials, with much of the research effort being associated with diffusion bonding. It has been suggested that improved weld properties can be achieved by introducing an interlayer
between the contacting substrates prior to the friction welding operation (Elliot and
Wallach 1981b). t-iowever, this particular approach has not been researched in detail.
Only limited research has been carried out on friction welding using interlayer materials.
7.2. THE INTERLAYER APPROACH lnterlayer materials have been applied during dissimilar friction welding of metallmetal combinations such aluminium and steel base materials (Hartwig and
Kouptsidis 1977, Dunkerton 1982, 1983, and Sassani and Neelam 1988). However, almost al1 this work has concentrated on evaluating the effects of interlayer selection on joint mechanical properties. For this reason, almost al1 of the detailed information that has been published conceming the metallurgical and mechanical effects produced when interlayers are introdüced during dissimilar joining applies to the diffusion bonding process.
lnterlayer rnaterials have been used as buffers that bridge the wide differences in properties (the coefficient of thermal expansion, thermal conductivity, yield strength or elastic modulus, and so on) and in mutual solid solubility of the contacting components.
It has been suggested that brittle intemetallic layer formation at the bondline can be prevented through judicious selection of an interlayer material (Elliot and Wallach
198ib). lnterlayer materials have also been applied with the objective of taking advantage of the speed and reliability of the friction welding technique.
Much of the published information is incomplete or contradictory concerning the detailed effects produced when an interlayer is introduced during dissimilar friction 3 welding. For example, it has been reported that the introduction of a 25 pm-thick silver interlayer produced a substantial irnprovement in tensile strength properties durhg dissimilar aluminium/stainless steel friction welding (Hartwig and Kouptsidis 1977). No evidence of intermetallic formation was reported and no expianation for the strength improvement produced by introduction of the silver interlayer was presented. In contrast, Dunkerton found only a modest increase in the joint strength when they introduced a silver interlayer during aluminium/steel friction welding. This researcher reported the formation of an intermetallic phase at the dissimilar joint interface but did not identify it or characterise its effect on joint mechanical properties (Dunkerton l982,i983).
The effects of interlayer materials on joint tensile strength properties has been the main focus during dissimilar friction welding and this has resulted in a dearth of information concerning the influence of interlayers on intemetallic phase formation at the dissimilar joint interface. For example, it is not possible based on the published literature to detemine whether a particular interlayer will prevent intermetallic layer formation at the bondline, how it will do SO, under what conditions it is best applied, and what its overall effect will be vis-à-vis joint mechanical properties. It is even possible that use of an interlayer material could produce per se harmful intermetallic phases during the welding operation, i.e. by combining chemically with alloying elements contained in one or both of the adjoining metallic substrates. The scantiness of information concerning the influence of interlayer materials on intermetallic layer formation is al1 the more surprising since it has been stressed that the avoidance of harmful intermetallic layer formation is the prime incentive for introducing interlayers during dissirnilar welding operations (Elliot and Wallach 1981a, b). 4 Based on the above discussion, it is apparent that there is need for proper identification of the composition, morphology, dimensions, location, and mechanical properties of the intermetallic layers forrned when an interlayer material is applied during dissimilar friction welding. The effects of the intemetallic layer on weld mechanical properties also require much clarification. This is highlighted using the following example. It has been suggested that the interrnetallic layer width has a critical influence on joint mechanical properties. A critical intermetallic width ranging from 0.2 to 1 pn has been proposed during dissimiiar aluminium/stainless steel friction welding (Jessop et al. 1978, Elliot and Wallach 1981b). The tensile strength of completed welds markedly decreases when this critical dimension is exceeded. In effect, it has been suggested that a geometric feature (the intemetallic layer width) is more important than the composition and mechanical properties (yield strength and ductility) of the intemetallic phase. Not only is a geometric feature emphasised but also the implication is that intermetallic layer formation and its subsequent growth depends on interdiffusion. Therefore, the width of the intermetallic layer fomed at the bondline will increase when the peak temperature attained at the bondline increases and when the friction tirne increases. There is support for this proposal. For example, the highest mechanical properties in dissimilar Ti/aluminiurn alloy A5083 welds were produced using a combination of short friction times and high friction pressures (Fuji et al. 1995a, b).
However, intermetallic Iayers can be formed as a result of mechanical mixing during the friction welding operation. When this occurs, the interrnetallic layer width in completed joints may depend, not on peak temperature or friction time, but on whether the intermetallic cornpound is forrned early in the joining operation and is then retained or removed as the welding operation progresses. When intermetallic layer formation 5 during dissimilar welding depends on a combination of interdiffusion and mechanical
mixing early in the welding process, the situation will be even more complicated.
The proposal that a critical intermetallic layer width exists during dissimilar friction
welding may not be applicable when age-hardened aluminium alloys are joined to steel.
The thermal cycle produced by the friction welding operation will create a softened
region in age-strengthened base material irnmediately adjacent to the bondline. With
this in mind, it would be expected that the influence of welding parameters changes on
final joint strength properties would depend on the interplay of intermetallic layer
formation and softened zone in aluminium alloy base material adjacent to the bondline.
It is well documented that contaminants such as lubricating oils, oxide films and
the surface roughness prior to friction welding have a marked influence on joint
mechanical properties (Jessop et al. 1978, Fuji et al. 1992). Lubricating oils and oxide
films alter the coefficient of friction between the contacting surfaces, the amount of heat
generated during the welding operation and the final joint strength (Jessop et al. 1978).
It has been reported that the tensile strengths of dissimilar TifAlSI 304 stainless steel friction welds increase when the surface roughness decreases from 2.2 pm to 0.05 Pm.
This improvement in joint mechanical properties has been associated with decreased
magnitude of the residual strains and with the removal of surface oxide films and contaminants from the contacting substrates (Fuji et al. 1992).
Introduction of an interlayer material prior to friction welding will alter the surface conditions (roughness, oxide film composition, and adherence) and the frictional coefficient when the adjoining substrates corne into contact, see figure 1.l. Also, introduction of an interlayer material will affect joint mechanical properties because it will modify the type, chernical composition, dimensions, location, and mechanical properties of the intermetallic layer formed at the bondline. Finally, since the 6 introduction of an interlayer affects heat generation dunng the welding process, this will rnodiw the dimensions and mechanical properties of the softened zone formed immediately adjacent to the dissimilar joint interface.
1.3. DETAILED THESIS OBJECTIVES In the present thesis, friction welding of dissirnilar aluminium alloy 6061(T6)
MMC/Ag/AISI 304 stainless steel is investigated. The behaviour of a silver interlayer and the nature, chemical composition, and location of the different intermetallic phases fonned when using a range of welding parameter settings is examined. The presence of reinforcing A1201 particles in MMC base material provides information conceming the events that occur in the aluminium-based substrate very early in the friction joining process. The specific objectives of this thesis comprise:
Examining the fracture of reinforcing AI2O3 parficles during dissimilar MMCfAlSl
304 stainless steel a~dMMCfAg/AISI 304 stainless steel friction welding to
elucidate the effects of silver interlayers on the mechanical processes that occur
in mâterial close to the bondline. Reinforcing particle fracture behaviour is also
used to monitor changes as function of tirne in the region close to the bondline.
Developing a deeper understanding concerning the role that silver interlayers
perform when the friction welding parameters are varied, in particular, how
intermetallic layers are formed and removed at the bondline. It is worth noting that
no published information exists concerning how interlayer rnaterial performance is
affected by welding parameter selection.
Examining the factors determining the mechanical properties of dissirnilar
MMC/Ag/AISI 304 stainless steel and MMCIAISI 304 stainless steel friction welds
that contain both intermetallic layers and softened zones in MMC base matenal
irnmediately adjacent to the bondline. The detailed effects of interrnetallic layer 7 and softened zone formation on the tensile strength of dissimilar aluminium alloy
MMCIAISI 304 stainless steel friction welds are evaluated using finite element
modelling (FEM).
1.4. THE PROCEDURES Dissirnilar aluminium alloy 606 1(T6)-base MMCIAISI 304 stainless steel friction welds with and without silver interlayer are produced using a range of welding
parameters (friction pressure and friction tirne). The joint strength is rneasured using typical notch tensile testing specimens. In particular, the influence of silver interlayers on the mechanical properties of weids containing both intemetallic layes and a sofîened zone in MMC material adjacent to the bondline is exarnined and rnodelled
using FEM. This work effort is complemented by extensive microstructural characterisation (optical, SEM, and TEM microscopy) of rnaterial at and close to the dissimilar weld interface.
1.5. THESlS ORGANIZATION Chapter 1 provides a summary of the thesis objectives. Chapter 2 examines background information concerning the principles of the friction welding process, friction welding phenornena, dissimilar joining, intermetallic layer formation, and softened zone. Chapter 3 describes the experimental procedures that were applied.
Chapter 4 discusses how the silver interlayer affects particle fracture during dissimilar
MMC/AISI 304 stainless steel friction welding operations. Chapter 5 examines the microstructural changes produced in dissimilar MMC/AISI 304 stainless steel and
MMC/Ag/AISI 304 stainless steel friction wefds. Chapter 6 examines softenea zone formation. Chapter 7 examines the interplay of intemetallic layer and softened zone and their effects on dissimilar weld mechanical properties. Finally, the conclusions are presented in Chapter 8. DISSIMI'AR FRICTION MATERIALS 4 O PARAMETERS
-.
DEFORMATION II' II' 1 t 1
HEAT GENERATION -1 AND TRANSFER I
F - I lNTERMETALLIC ' 1 FRACTURE 1 II' COMPOUNDS 1 I J I
Z 1 FRICTION JOINT 1 PROPERTIES 1
I-Iœ-I-iœœ-I
Figure 1.1. The effects of an interlayer on dissimilar friction welds. CHAPTER 2. THE FRICTION WELDING PROCESS
2.1. INTRODUCTION Friction welding is classified by the American Welding Society (AWS) as a solid- state joining process in which bonding is produced at temperatures lower than the meiting point of the base matenals (Johnson et al. 1991). However, the underlying mechanisms that determine both similar and dissimilar friction welding are not completely understood and the results of recent research do not support the above view. For example, it has been shown that the solidus temperature is attained at the contact zone of rotary AIIAI friction welds and at the rotating toolladjoining material interface in friction stir welded aluminium base material (Bendzsak et ai. 2000 and
North et ai. 2000). The high heating rate during friction welding operation can facilitate local melting, e.g. undissolved q' particles segregated at the grain boundary regions in
Al 7030 T6 base material promote local melting when the eutectic temperature (475 OC) is attained (Grong 1994). As a result the above definition may be considered to be
Iimited and controversial.
In an earfy book on the subject (Vill 1962), it is mentioned that in 1956, the
Russian machinist, A.I. Chudikov successfully demonstrated the possibility of achieving high-quality butt welds between steel rods. It is in this period that study of friction welding started in order to understand its fundamentals. Friction welding (FRW) produces a joint under compressive force when the workpieces rotate or move relative 10 to each another producing heat and displacing matenal frorn the faying surfaces.
Finally, the weld is consolidated applying a forging pressure (Johnson et al. 1991,
Elmer and Kautz 1993). Currently two types of friction welding process are used in industry, the inertia friction welding process and the continuous-drive or direct-drive friction welding process, see figures 2.1-2.2.
2.7.7. lnerfia fnction welding In inertia friction welding (see figure 2.1) a rotating spindle with a replaceable flywheel (or combination of flywheels) is brought up to a speed of N (rpm). The welding cycle starts when the non-rotating piece is pushed against the rotating cornponent. The
stage ~"";"'t-""'~-~ III f
Start Time - Welding Welding Complete
Figure 2.1. lnertia friction welding process (Wang and Rasmussen 1972). 11 drive power is cut off simultaneously and produces a rapid drop of rotating speed so that it stops completely in a short time period. The inertia friction welding process blends the heating and forging phases. On the other hand the direct-drive process has a distinct heating period before forging takes place (Wang 1975).
One important difference with the direct-drive process is the number of welding parameters to control. During inertia friction welding, there are three key parameters: the friction pressure, the initial rotating speed, and the total moment of inertia of the rotating mass. Compared with the direct-drive fiction welding, the inertia welding process is a much simpler process (Wang and Rasmussen 1972).
The inertia welding process produces a torque curve containing two peaks, one in the initial stage of friction welding and other near the end of the welding cycle. In the initial stage, the torque reaches the first peak and drops to an equilibrium level during the steady stage of the friction welding process. It should be pointed out that in elastic contact the torque produced between a rotating body and a surface depends directly on the coefficient of friction, the applied load, and the apparent area of contact (Johnson
1985). Since the applied friction pressure remains constant in Stage 1, see figure 2.1, the increasing torque values can be ascribed to an increase in the area of contact. The second torque peak occurs when the rotating speed is near to the end of the process.
This particular behaviour results of the cornbined effect of the material viscosity and strain rate. During deceleration, the material viscosity increases while the strain rate drops to zero.
Appreciable shortening (burn-off) of the workpieces is produced during welding since material close to the interface is heated up and extruded producing the typical bulge or flash of the friction welds. 12 2.1.2. Direct-drive friction welding
2.1.2.1. Heating period The heating penod of the direct-drive friction welding process can be divided into
three distinct phases, see figures 2.3-2.4.
Stage 1.
During this stage the components are brought into contact under an applied
constant rotational speed (N) and compressive friction pressure (Pl). The torque rises
until it reaches a peak value and the axial shortening is almost equal to zero, see figures 2.2 and 2.3. The strong adhesion between asperities forms junctions that can
be stronger than the adjoining substrates. The tangential movement between surfaces
breaks these junctions, resulting in mass transfer and heat generation. Since the large transferred fragments ride over each other, they push the two components apart and therefore al[ rubbing is confined to an annular area. The size of transferred fragments grows until they form a continuous plasticised layer, see figure 2.4.
It is convenient to point out that the plasticised layer term has been used to describe the layer formed close to the bondline during the friction welding process. This terni cannot be applied during the whole process because during a period of tirne this
layer behaves as a viscous material and is affected by the temperature and strain rates close to the contacting surfaces. Friction Time (tl) Deceleration Time (t2) Delay Time (t3) - 1 Forging Rotational Speed (N) Time (t4) 1I Friction pressure (Pl)
Welding Time - Welding Starts Ends
Figure 2.2. Direct-drïve friction welding process.
Stage II
The flow strength of the plasticised layer decreases and a steady-state condition
is established. The plasticised layer then behaves as a viscous material. The torque decreases to a minimum value, metal is heated by a viscous dissipation mechanism.
Material on both sides of the joint interface is heated up and extruded producing the collar observed in friction welds.
Stage III.
In this deceleration phase the viscous layer is displaced radially outwards from the component centreiine. The torque rises reaching a terminai peak value again and then decreases to zero, see figure 2.3. I 0: : 1 1I a! Time ---)
1 l I 0: Time + I I 11 1 t I 1I I : l 1 I
1 I I I 1 l 1 1 l I I l I I
0; i I1 I Time k-b I OI II 1 l II I 1 11 Stage II ?=y1 I T Stage ïïï i i I Stage 1 I l l Frictioning wlI I Stage I Figure 2.3. ldealised traces of the variations with time of speed, torque and axial shortening in the direct-drive friction welding process during the frictioning stage. 1-Peak torque. 2-Equilibrium torque. 3-Terminal torque (Duffin and Bahrani 1976). Stage 1. Dry fiction occurs in an annular area at the interface- A plasticised region starts developing in the surface and subsurface of the annula area The size of the contact area depends on the applied fkïction pressure and increases with 6nction time along with'the resisting torque. Axiai shortening starts at the end of this stage.
Stage II. The material of the plasticised region starts behaving as a viscous material. The annular contact area expands and the resisting torque decreases unt il it reac hes the equih'brium value. Axiat shortening continues. The tipicai collar of fiction joints starts fonning.
The viscous Iayer region increases in size and decreases in thickness. The radial displacement is inwards.
Stage III. Deceleration occurs, the rotating piece is stopped- The viscous Iayer covers most of the contact dace. The radial displacement is outwards. The resisting torque reaches a second peak value.
Figure 2.4. The heating period based on the work of Dufh and Bahrani (DufFm and Bahrani 1976). 16 When low friction pressures or short friction times are applied, there is a tendency
to lack of bonding at the periphery of the joint. Good quality friction welds require
spreading of the fully viswus layer across the contacting surfaces- Duffin and Bahrani
consider that Stage III of the friction welding process plays an important role by
spreading this region across the contact region, see figure 2.4, and therefore affecting
the weld mechanical properties (Duffh and Bahrani 1976). However, it is the opinion of
this author, that most of the spreading of the plasticised layer must occur previous to
this stage. This situation is attained applying higher friction pressures and longer
friction tirnes.
2.1 -2.2. Forging stage It has been suggested that a key objective in the forging stage is to expel the
plasticised layer into the flash and with it, the trapped oxides and contaminants. Rich
and Roberts developed an interfacial dispersion parameter (q) for similar welding, which they daim it can be used to estimate the removal of oxides, voids and contamination from the surface during forging (Rich and Roberts 1971a). The interfacial dispersion parameter is the ratio of the original interfacial cross sectional area, which becomes expelled into the flash during forging. It ranges from a value of zero where no dispersion takes place to one where al1 of the original interfacial area has been dispersed into the flash, see equation 2.1.
where q is the dispersion factor, U, is the axial shortening (burn-off) of each component 17 during the forging stage, and ho is the thickness of the plasticised zone at the onset of forging (Rich and Roberts 1971). From equation 2.1, it is noted that friction pressure has a great influence on the dispersion parameter. Friction pressure (Pl) detemines the plasticised layer width (ho) prior to forging because this parameter is inversely related to the friction pressure (Bendzsak et al. 1997). As result, the higher the friction pressure, the higher will be the 7 value.
However, the prior model proposed is of Iimited application. Rich and Roberts assumed that flash formation is produced completely during the forging stage, but flash formation starts in the frictioning stage of the welding process. At the time that forging pressure is applied, a large axial bum-off has already been produced. Their suggestion that forging plays an important role on expelling trapped oxides, voids, and contamination into the flash is not well supported. In successful friction welds alf impurities are rernoved previous to the beginning of the friction welding process, a dirty surface is associated with a low coefficient of friction, and low heat input in Stage I of friction welding. These welds show low joint strength, see Chapter 4.
2.1.2.3. Direct-drive friction welding parameters The direct-drive friction welding machine has more parameters that can be set up compared to the inertia welding machine. This situation makes welding parameter optimisation more difficult. The basic welding parameters in the direct drive welding variables comprise: rotational speed (N), friction pressure (Pl), friction time (tl), upsetting or forging pressure (P2), and forging time (t4), see figure 2.2. Some authors have investigated as many as seven parameters to determine their respective effects on weld properties, e.g. deceleration time (t2) and delay tirne (t3) have been included
(Wang 1975). 18 Murîi and Sundaresan applied a statistical approach when optimising dissimilar friction weld mechanical properties. Parameter optimisation was carried out using a 2" factonal experimental design. They examined the following combinations: low alloy steel/stainless steel, high-speed steelfmediurn carbon steel, and alurninium/stainless steel. They studied the effect of friction pressure, friction time, and forging pressure on the notch tensile strength. They found that in dissimilar aluminium/stainless steel, high values of friction pressure, friction time, and forging pressure produced the optimum joint mechanical properties (Murti and Sundaresan 1983). However, a number of
investigators do not support the idea that long friction times improve the joint strength
(Jessop et al. 1978, Fuji et a/. 1995 a, b, and Zhou et al. 1995).
2.2. INTERMETALLIC COMPOUNDS IN DlSSlMlLAR JOlNlNG The formation of brittle intemetallic layers is of major concern in dissimilar friction welds. Therefore, in the following sections the metallurgical and mechanical properties of relevant intemetallic compounds are reviewed, and reiated to the problem of dissirnilar friction welds.
2.2.1, Fe-AI intemetaliic compounds lntermetallic compounds have a well-defined crystal structure and exist within narrow composition ranges, see figure 2.5. The iron-aluminium binary equilibrium diagram indicates five stable intemetallic compounds (Fe3AI, FeAI, FeAI2, Fe2AI5,and
FeAI3). At 49.0 wt% aluminium, the FeAI2 compound appears. The intemetallic compound Fe2AI5, is produced by a congruent melting reaction at 1169 OC and contains 55 wt% Al. At 1160 OC, a peritectic reaction between the melt and Fe2Ai5 produces FeAI3. The FeAI3 composition ranges from 58.5 to 61 wt% AI. Whereas between 49 and 53 wt%, Fe2AI5and FeAI2 coexist. In this connection, the presence of
Fe2AI5, Fe3AI and FeAl has been detected in dissimilar 1050 aluminium alloy/AISI 304 Atomic Percent Alumulurn
--- Fe Weight Percent Aluminum Al
Figure 2.5. Fe-AI binary diagram (Kattner et al. 1987).
stainless steel friction welds produced using a friction time of 1.5 s (Fukumoto et al.
1997).
The mechanical properties of the iron aluminides are sensitive to many factors, inciuding aluminium content, heat treatment, test temperatures, alloying additions, environment, microstructure, and defects. In general, the room temperature tensile yield strength of iron aluminides tends to rise with increasing aluminium addition, reaches a peak near the Fe3AI stoichiometric composition and then decreases again.
At room temperature, binary compositions with aluminium contents 4 8 at% Al fail in a ductile manner by void nucleation and coalescence and exhibit tensile ductilities of over
20%. However, Fe3AI and FeAl characteristically exhibit limited tensile ductility. Typical tensile elongation values are less than 10% for Fe3Al-base binary compositions, and less than 4% for compositions above 35 at% Al (Stoloff et al. 1994). WEiGHT PERCENT NICKEL
Figure 2.6. AI-Ni binary diagram (Hultgren et al. 1973).
2.2. 2. Ni-AI internetallic compounds The binary phase diagram indicates the presence of four intermediate phases:
Ni& (E), Ni2AI2 (ô), NiAl (P'). and Nidl (a').see figure 2.6. There is no previous work on intermetallic NiAl compounds in dissimilar friction welding but extensive research has been carried out on the physical and mechanical metallurgy of Ni3AI (a')and NiAl
(P') compounds. The yield strength of NiAl (a') is in the range from 70 to 100 MPa for aluminium contents from 23 to 25 at% and increases substantially with aluminium 21 content above 25 at% reaching a room temperature yield strength of 220 MPa with
aluminium contents of 27 at%. With the exception of chromium al1 substitutional elements can increase the yield strength. The mechanical behaviour of NiAI (fi') is similar to the behaviour of Nidl (a'),it has a low-temperature brittleness. Also, its yield strength depends on the aluminium content, Le. when the aluminium content increases from 45 to 50 at%, there is a drop in yield strength from 660 to 260 MPa. However, the
use of small additions of iron, rnolybdenum, and gallium can improve the ductility from
î to 6 %. As is the case with most intemetallics, low-temperature ductility is inadequate
(Stoloff et al. 1994)
2-2.3. Ag-AI intennetallic compounds The phase diagram of Ag-AI, shows the presence of three intermediate phases:
Ag3N (p at high temperature), Ag3AI (p at low temperature), and Ag2AI (5). There is no published information about the mechanical properties of Ag-AI intermetallic compounds but the effects of these intermetallic compounds in dissimilar aluminium alloys/Ag/stainless steel diffusion welds are discussed in section 2.2.6.
2.2.4. Growth of intennetallics Intemetallic formation has been examined during contact between rnolten aluminium and solid iron. Two phenornena occur simultaneously at the contact interface, i.e. dissolution of iron into Iiquid aluminium and the formation and growth of an alloy layer comprising intermetallic compounds produced due to the migration of molten aluminium into iron. The intermetallic layer was mainly Fe2At5 immediately adjacent to the iron substrate and FeAI3 was immediately adjacent to liquid aluminium
(Vaillant and Petitet 1995). WEIGHT PERCENT ALUMINUM
Figure 2.7. Ag-AI binary diagram (McAlister 1987).
The growth rate of the intemetallic layer was limited by the diffusion of reacting
species through the intemetallic layer and not by interface chernical reaction (Vaillant
and Petitet 1995). As result, the intermetallic growth rate (&dl) was inversely
proportional to the interlayer thickness (x), see equation 2.2.
where k, is the parabolic growth rate constant (m2s") and r the contact tirne (s). The intermetallic layer thickness increased in accordance with equation 2.3. From equation 2.3 it is apparent that the intermetallic thickness (x) depends on holding time (t). In dissirnilar diffusion bonding, holding time plays an important role, the width of the intermetallic layer increases with the holding time (Calderon ef al. 1985).
However, there is an important difference between friction welding and diffusion bonding, interdiffusion is not the only mechanism involved in intermetallic formation in friction welding, mechanical mixing is also important, and its role on intemetallic layer formation is not completeiy understood.
In friction welding, friction tirne is an important parameter and it has been suggested that decreased intermetallic layer thickness is promoted when short friction times are applied (Fuji et al. 1995a, b). However, high friction pressures produces thin intermetailic layers. Fuji et al. suggested that a high friction pressure squeezes the plasticised region and the intermetallic layer from the bondline. This statement is Iimited because the plasticised region close to the bondline in dissimilar welds can be considered as stationary. Therefore, the complete removal of the intermetallic layer is difficult. The only possible effect of friction pressure is to decrease the width of the plasticised layer and consequently the growth of the intermetallic compounds.
2.2.5. Inmence of composition on fhe intemetallic layer There is no information about the effects of chemical composition on the intemetallic layer in dissimilar aluminium/steel friction welds. However, during the contact of molten aluminium and steel it has been observed that alloying additions such as silicon, nickel, and copper, reduce the intennetalfic layer thickness (Vaillant and Heating tirne (h)
Heating time (ks)
Figure 2.8. Effect of holding time at 823 and 873 K on the tensile strength of ~il(6xl0~wt?h Si) Al and Til(0.12 wt% Si) Al joints O 873 K, Til(O.12 wt?! Si) Al; O 873 K, ~il(6xl0~wt?! Si) Al; 823 K, ~il(ôx10~ wt% Si); * fractured at the interface (Fuji et al. 1995~).
Petitet 1995). Though, the case is not the same because the diffusivity of an element in
Iiquid metal is much higher than in a solid material. It can be expected that the presence of some elements may decrease the width of the intemetallic layer at the bondline of a friction weld.
Although there is no much information on the effect of alloying elements on the intermetallic layer formed in dissimilar aluminium alloys/steel friction welds, there is plenty of information on the effect of alloying additions on the strength and ductility of monolithic iron aluminides. The results of past studies show that only chromium improves the tensile ductility of FeAl-based iron aluminides. Additions of 2-6% of chrornium have resulted in doubling the room temperature ductility. The high temperature tensile and creep strengths in ail iron-alurninide compositions have been 25 improved with additions of elements such as silicon, thallium, cerium, titaniurn, molybdenum, zirconium, hafnium, or niobium, but usually at the expense of room temperature ductility (Stotoff et al. 1994).
It is convenient to examine the influence of silicon on the intermetallic layer formed during dissirnilar friction welding of titanium and aluminium base materials (Fuji et al. 1995~).Fuji et al. examined the influence of two residual silicon contents (6x1o4 and 0.12 wt%) on the mechanical properties of dissimilar TVA1 friction welds following post-weld heat-treatment. The joint tensile strength and bend test properties were quite different in the base materials containing different silicon contents, see figure 2.8. The tensile strength and bend test properties of dissimilar TVA1 welds containing the lowest silicon content (6x10~wt% Si) were markedly decreased by heat treatment compared to Ti/AI joints made using an aluminium alloy containing higher silicon content (0.12 wt% Si). Weld failure was associated with AI3Ti formation at the dissimilar joint interface and occurred when the width of the intermetallic layer formed at the bondline exceeded
10 Pm. The rate of growth of the AIJTi intermetallic layer was faster in post-weld heat- treated TVA1 containing low residual silicon contents. More silicon segregated to the titanium side of the dissirnilar joint suggesting that silicon segregation retarded AI3Ti formation by acting as a barrier to titanium and aluminium diffusion (Fuji et al. 1995~).
2.2.6. lntermetallic compounds and mechanical propetfies The formation of brittle intemetallic layers is of major concern in dissirnilar aluminium alloys/steel friction welding. In a recent study, it is suggested that the poor weld strength found during tensile testing of aluminium-based MMWAISI 304 stainless steel friction welds resulted from the retention of an FeAI3 intermetallic layer and an oxide Fe(AI,Cr)204 or FeO(AI.Cr)203 film at the dissimilar joint interface (Pan et al. Figure 2.9. Thickness of the intermetallic layer (Ti3AI) vs. hea.ing tirne in a TilAl diffusion joint (Suzuki et al. 1994).
1 996).
The deleterious effect of interrnetaliic layers on the mechanical properties of dissirnilar diffusion welded joints have also ibeen investigated. In pure aluminium/pure titaniurn diffusion welds, the width of the intenetallic layer was 100 pm following a holding tirne of 400 Ks (approximately 100 hrs) at 600 OC. The joint strength increased and attained a peak value when the intermef!iataliic thickness was 35 pm, see figures 2.9-
2.1 0. However, joint strength decreased when this intermetallic thickness was exceeded. Although these test results are consistent with the suggestion that a critical intennetallic thickness exists, which produces optimum weld mechanical properties, this result is not always the case. For example, in dissimilar aluminiurn/silver/stainless steel diffusion joints. the joint tensile strengtti decreased linearly frorn 240 MPa to 100
MPa when the intermetallic layer thickness imcreased frorn O to 16 pm, see figure 2.1 1
(Calderon et al. 1985). Heating time I ks
Figure 2.10. Joint strength vs. heating time in a TilAl diffusion joint (Suzuki et al. 1994).
The proposal that improved joint strength properties are produced when the
intermetallic layer width is less than a critical value in dissimilar friction welds has also been suggested (Fuji et al. 1992). For example, Fuji et al. found that the mechanical properties of dissimilar Almi friction welds decreased markedly when the intermetallic layer width exceeded 0.5 Pm. In addition, they suggested that the critical width depended on the mechanical properties of the adjoining substrates. This particular issue is discussed in detail in Chapter 7 of the present thesis.
2.3. FRICTION AND WEAR
2.3.1. Soft coatings and the coefficientof friction Friction is the resistance encountered when one body rnoves over another with which is in contact. This resistance to sliding is described using the coefficient of friction. This parameter is particularly important during friction welding. Heat generation, This Work : Unaged O Aged at 473 "K A Aged at SI3 "K Other Results : x Ref (10) - --- Best Fit, Ref (10) \ x
Figure 2.11. The dependence of the strength of the weld between stainless steel and 1100 aluminium alloy on the total thickness of the intermetallic layer (Calderon et al. 1985).
the onset of plastic yielding, delamination, and other friction related phenomena depend on the coefficient of friction (Johnson 1985).
The coefficient of friction depends on many factors, among them the contacting materials, the environmental conditions, and so on. The coefficient of friction can t;e rnodified using a soft coating or interlayer material. In a system containing an interlayer, the interlayer thickness, the mechanical properties of the interlayer and adjoining substrate materials play important roles. For exampie, if a steel substrate is covered with a silver interlayer, the coefficient of friction decreases (Iliuc 1980).
In a sliding system containing an interlayer, the energy-mode1 proposed by
Heilmann and Rigney provides some insight concerning the effect produced by 29 interlayers on the coefficient of friction value (Heilmann and Rigney 1981). According to this model, the coefficient of friction in a coated substrate depends on the strength and thickness of the selected rnaterials. For a soft interlayer on a hard base material, e-g. a silver film on a steel substrate, the coefficient of friction has values between the coefficient of friction of the interlayer material Of) and that of the base material OB), see figure 2.12. The expressions obtained by Heilmann and Rigney are cumbersome.
They indicate that the coefficient of friction p is a function of the interlayer thickness (T), the applied shear stresses in the interlayer and the base material and #),and their respective shear strengths (kL. and kB). Figure 2.12 shows the coefficient of friction as function of the interlayer thickness (T). This figure indicates that the coefficient of friction can reach a minimum value when the interiayer is in the range from 0.1 to 10
Fm. This prediction seems to be in agreement with results obtained by Jahanmir et al.
These researchers examined the influence of the 0.1-pm thick cadmium plating on annealed AlSI 1018 steel. However, the influence of the thin cadmium layer depended on the testing atmosphere; it was not effective in an oxidising atmosphere. The coefficient of friction (p) of the 0-1-pm thick cadmium plated steel was 0.4 in air and 0.2 in argon. In unplated steel p = 0.65 in air and 0.8 in argon (Jahanmir et al. 1975). Thus, the introduction of a thin soft interlayer material can decrease the frictional coefficient. Figure 2.12. Examples of dependence of the coefficient of friction of p for a soft coating on harder substrate materials: (A) pL 1p BM.3; (B) 1.3~~L1p 8SI ; (C) pLlpB>l;(O) pL 1p B4 (T = thickness) (Heilmann and Rigney 1981).
2.3.2. Wear and fricion welding In Wear, avoiding seizure between contacting substrates has been a matter of much concern for tribologists. However, in friction welding one is interested in the conditions that promote seizure (since this leads to joint formation). Seizure has been associated with severe Wear between contacting surfaces. The conditions that promote seizure have been investigated for aluminium alloysfsteel pairs. Since this research is of direct relevance to friction welding, Wear in such systems will be examined.
Wear mechanisms when testing aluminium alloys and aluminium metal matrix composites have been examined in detail (Ames and Alpas 1995, and Zhang and
Alpas 1996). Ames and Alpas studied the Wear properties of a combination comprising an hybrid A356 composite containing 20 vol% Sic and steel. Zhang and Alpas 31 examined the sliding Wear behaviour of an aluminium alioy 6061-1 0 vol% A120&teel combination. They found that Wear properties were related to a transition temperature, severe Wear and seizure occurred when the temperature in the surface was higher than the transition temperature of the test material. The transition temperature is an inherent feature of the aluminium alloy base material. For example, A356 aluminium alloy MMC base material containing 20 vol% Sic had a transition temperature of 162 2 OC.
However, Wilson and Alpas indicated quite different values for the transition temperature of the same composite material (440450 OC). ln unreinforced 6061 aluminium alloy the transition temperature ranged from 110 to 175 OC. Aluminium alloy
6061-MMC containing 20 vol% AI2O3 particles had transition temperatures ranging from
310 to 350 OC (Wilson and Alpas 1996). These results are of interest in friction welding of aluminium alloys, because they indicate a possible minimum temperature above which welding may happen. Therefore, in friction welding it would be necessary to reach temperatures higher than the corresponding material transition temperature. For example, in similar friction welding of aluminium alloy substrates, material close to the bondline is subjected to peak temperatures around 550 OC (Midling and Grong 1994a).
Such temperatures are much higher than the transition temperature values found in
Wear testing of aluminium alloy base materials.
2.4. THERMAL ASPECTS OF FRICTION WELDlNG
2.4.1. Calculation of heat input in friction welding The temperature distribution in friction welds is deterrnined by factors such as the power input, the thermophysical properties of the adjoining base materials and by flash formation. A key problem during modelling of the friction welding process is in obtaining an accurate description of the heat generated at the bondline. In sliding friction, heat generation has been described using equation 2.4: where Q (W) is the heat generation rate during frÏctional sliding, p is the coefficient of
friction, P (N) is the applied load, and ~(ms-')is the velocity (Johnson 1985).
In friction welding, the problem is complicated by the effect of the radius on factors
such as the rubbing speed, the pressure distribution, and the coefficient of friction. The
heat generation rate (g)is expressed by equation 2.5:
where n is the rotational speed (rpm), r is the distance from the centre, p is the
coefficient of friction, and P(r) is the pressure distribution. The coefficient of friction (p)
and pressure distribution P(r) are affected by the rubbing speed, the surface temperature, the material hardness, and the surface condition (Wang 1975).
When a torque T is applied on the contact surface, see figure 2.13, almost al1
plastic work produced by shear loading is transfomed into heat; the average heat input
is therefore:
where qd.2 (~rnrn-~)is the net power, P (MPa) is the friction pressure, and V,, (ms-') is the maximum speed at the periphery of the rotating cornponent. Equation 2.6 serves as 33 the basis for calculating heat generation at the bondline (Grong 1994). However, it is worth to keep in mind that equation 2.6 is valid only for Stage I of friction welding. which involves dry sliding friction or relative velocity between contacting surfaces. Once, the relative velocity between contacting surfaces is zero, heat is generated mainly by a viscous dissipation rnechanisrn.
2.4.2 Temperatum distribution during friction welding The factors deterrnining the temperature distribution in friction welds have been studied by a number of researchers (Cheng 1963, Rich and Roberts 1971b, Sluzalec and Sluzalec 1993, Midling and Grong 1994a. The rnodel originally proposed by Rikalin considers a continuous plane heat source in a long rod (Midling and Grong 1994a). In this mode1 the temperature of the contact section at the end of the heating periodi is Th the temperature distribution is given by equation 2.7:
Figure 2.13. Schernatic arrangement of friction welding of a solid rod (Grong 1994). where Th = temperature at the end of the heating period, OC,
T, = ambient temperature, OC,
t = tirne, s,
t,: = duration of the heating period, s,
x = distance to the contact surface, mm,
2 1 a = thermal diffusivity, mm s- .
Equation 2.7 describes the temperature at different distances from the contact surface during the heating period (Midling and Grong 1994a). This equation is applied in Chapter 6.
2.5. WELDING METALLURGY OF ALUMINUM ALLOYS
2.5=1. AbMg-Si Alloys Age-hardenable AI-Mg-Si alloys are widely used as structural materials and offer tensile strength values exceeding 350 MPa when the base material is in the artificially- aged T6 condition. Strengthening results from the presence of very fine, needle-shaped precipitates P" (Mg2Si) fomed along the crystallographic direction clOO> within the aluminium matrix. The problem with this type of aluminium alloy is that softened zones are formed due to the thermal cycle in friction welding. Softened zones are produced by the reversion (dissolution) of the p" (MgzSi) precipitates during the weld thermal cycle. This reversion produces a reduction of solute content that softens the matrix.
The presence of softened zones impairs the mechanical properties of completed welds
(Grong 1994).
2-52 Effect of the thermal welding cycle on dissolution and reprecipitatîon The microstructural changes in the softened zone are summarised as follows:
1. reversion (dissolution) of the Pn(Mg2Si)phase; a-. Peak Tem perature
9 reverted
1 HAZ b Distance from fusion line -+ Figure 2.14. Schematic diagram showing the hardness distribution following pS'(Mg2Si)dissolution in the HAZ of 6082-T6 aluminium welds (Grong 1994).
2. precipitation of P' (Mg2Si) during the cooling period following the welding
cycle; and,
3. natural ageing.
The effects of these microstructural changes on the hardness profile are shown in figure 2.14. The partially reverted region attains temperatures in the range from 250 OC to 500 OC. The peak temperature is 250 OC at the location where the heat affected zone teminates. It follows from figure 2.14 that reversion of P' (MgrSi) precipitates will occur to an increasing extent for peak temperatures ranging from 250 to 500°C. This reversion produces a continüous decrease in the softened zone hardness until the particle dissolution process is completed. During cooIing following friction weiding, some solute recombines to form coane metastable P' (Mg2Si) precipitates, see figure
2.1 5; these precipitates do not contribute to strengthening. However, in the region close to the bondline where al1 the precipitates have been dissolved, the material will be Figure 2.15. Precipitation of (3' (MgzSi) dispersoids during the weld cooling cycle (Myhr and Grong 1991a)-
supersaturated in alloying elements. Extensive hardening (natural ageing) of this
region will occur during a period of 5 to 7 days at room temperature. However,
enhanced strength recovery can be achieved through the use of artificial ageing in the
temperature range from 150 to 180°C (Grong 1994).
2-53,The sotlened zone The strength distribution in the softened zone following the welding operation and
the subsequent natural ageing can be calculated. Figure 2.16 shows sketches of the
superimposed hardness profiles, as evaluated from the equations developed by Myhr
and Grong (Myhr and Grong 1991a, b). The resulting strength level in the partly
reverted region depends on the interplay between two competing processes, dissolution and re-precipitation.
It can be observed from figure 2.16 that particle dissolution is the major softening
mechanism during friction welding of age-strengthened aluminium alloys and 37 aluminium-based MMC materials. Also, a substantial strength recovery will occur as result of plastic deformation in combination with intrinsic precipitation of hardening P"
(Mg2Si) precipitates following prolonged room temperature ageing. Depending on the operational conditions, this recovery may produce different shaped hardness distributions following friction welding. From figure 2.16(A) it is apparent that short friction times produce the narrowest HAZ regions. In contrast, high heat input during welding increases the width of the HAZ region. A similar effect is obtained when longer friction tirnes are applied, see Figure 2.16(B). Mowever, it should be pointed out that the model of Midling and Grong (MIDLING 1994) is not clear in one aspect, it does not compare different heat inputs for the sarne friction tirne. Also, this model does not consider the effect of friction pressure on the softened zone width. As result, it is not of general application. It will be shown later in Chapter 6 that an increase of the heat input produced by higher friction pressure results in narrower softened zones 4 , (*) I I Reversion mode1
- - Work hddening model I t 1- '- Nat 7'al ageing model r Axial distance -b
Reversion mode1
Unaffected base material
- Work har+ning mode1
i % - ,Naturd ageing mode1 Ï--YI Axial distance -b Figure 2.16. Schematic representation of the HM hardness distribution after welding and subsequent natural ageing. (A) Short duration thermal cycle. (B) Long duration thermal cycle (Midling and Grong 1994b).
2.6. MECHANICS OF SLIDING CONTACT Friction welding is a process that involves the application of normal and tangential loads when the substrates to be weld are in contact. The problem of sliding loading in a serni-infinite space has already been examined in contact mechanics. Although, there are some differences cornpared to friction welding, it is convenient to examine the effect of this type of loading on the nature and distribution of the stress and strain at and close to the contacting surface in a semi-infinite space. Figure 2.17. Line loading of a semii'nfinite space.
2.6.1. Line loading of a semi-infinite haif-space During Stage I of the friction welding process, loads are transferred frorn one body to the other via surface asperities. A qualitative picture of the stress distribution in the region close to the contact interface in a semi-infinite space can be derived assuming lineal loading contact and plain strain conditions. Plane strain conditions are justified because the dimensions of the asperity contacts are small compared to the surface area of the contacting substrates. The semi-infinite space is subjected to normal and tangential forces (see figure 2.17) with the relation between the normal load (P)and tangential load (Q)being given by equation 2.8, where p is the coefficient of friction.
The normal and shear stresses in a semi-infinite space can be calculated using equations 2.9 and 2.10 (Johnson 1985). These equations are applied to calculate the stress distribution at and close to the contact surface, see Chapter 4. 2.6.2. Torsional loading Although in friction welding contact occurs between the flat surfaces of two rods, the tangential load acting on a circular area in a semi-infinite surface is relevant to the friction welding process. The tangential load on a circular surface is applied perpendicular to the circle radius and results in surface rigid rotation. Considering the displacements produced by this type of loading, the only non-zero displacement is in the circumferential direction (us). Thus, torsional loading does not result in radial displacement (u,) or axial displacement (u=). Moreover, the torsional load does not affect the distribution of pressure on the contact region (Johnson 1985). Figure 2.18. Hemispherical normal pressure distribution (Heteny and McDonald 1954).
2- 6-3. Normal pressure and torsional loading When a circular punch applies normal and torsional loading to a semi-infinite
space, the surfaces are subjected to the combined action of normal pressure and
torque, see figure 2.18 (Heteny and McDonald 1954). In this problem, Heteny and
McDonald calculated the stresses and displacements at and near the contact surface.
It was assumed that at any point on the surface, the tangential load (Q) produced by
the applied torque was proportional to the normal load (P),with the factor of
proportionality being the coefficient of friction (p),see equation 2.8. in the semi-infinite
solid problem loaded on its plane-surface with a normal pressure distribution P(r), the
distributed shear or tangential loading Q(r) equals pP(r). Considering the cylindrical
system of coordinates (r, 8, and z), the displacements in the positive directions are denoted by u,, ua and u=,respectively. Figure 2.19. Displacernent components at the surface (Heteny and McDonald 1954).
In a combined loading problem, the normal pressure distribution does not produce a circumferential displacement (us}, which is the only displacement produced by the torque (see section 2.6.2). The effects of torque and normal pressure are independent of each other and consequently can be considered separately.
Heteny and McDonald obtained a corresponding solution for the displacements for elastic and plastic conditions. They assumed there was slippage on the contacting surface. This slippage produced a hemispherical normal pressure distribution, see figure 2.18.
At z = 0, and da 51, assuming a Poisson's ratio of 0.3, which is typicai of rnetals, figure 2.19 indicates the calculated surface displacements. The radial displacement (u,) is negative suggesting that rnaterial at the surface moves inwards reaching a peak value near the periphery. Extending this result to friction welding, it is suggested that 43 any trapped oxide or contaminant at the surface will be pushed towards the centre of the weld. As result, cleaning the interface from any oxide layer in the subsequent stages of the friction welding process is more difficult. On the other hand, the
circumferential displacernent (ue) reaches a peak value at the half radius location, while the axial displacement (u=)attains a maximum at the centre of the contact surface.
Since surface displacements are not unifonn, strains are not uniform in the surface and the region close to the contact interface. As result, straindependant phenornena such as heat generation, and particle fracture during friction welding of an MMC materiai will be affected.
At the surface, z = 0, the radial displacement u, is given as:
where, where p, = hemispherical distribution of Hertzian pressure, Pa,
P = applied normal load, N,
A = Lame's constant, Pa,
v= Poisson's ratio,
G = modulus of elasticity in shear, Pa,
E = Young's modulus, Pa,
a = radius of the punch, m.
When Poisson's ratio is equal to 0.5, the material is in the plastic condition and
Lame's constant A= CO, see equation 2.13. In this case the radial displacement becomes u = 0. see equation 2.11 (Heteny and McDonald 1954). Once again, this result suggests that when the surface is in the plastic condition. it is difficult to expel the impurities. Although, this result stands for material in the plastic condition at the surface, it suggests that in friction welding the case might be the same. This result confirms the importance of a clean surface at the beginning of the friction welding process.
2.6.4. Contact of rigid-ideally-plastic materials In the present study a silver interlayer is applied during friction welding of dissirnilar substrates. The interlayer is deforrned and worn away during the welding process. Deformation and Wear in the bondline are analysed using methods such as slip-line field theory and upper bound lirnit analysis. Upper bound lirnit procedures are applied in the present thesis, and their basics are found elsewhere (Dieter 1986,
Hosford 1993). Figure 2.20. Contact area growth of a plastic wedge under the action of a constant normal load P and an increasing tangential load Q (Johnson 1985).
2.6.4.1. Combined effect of shear and pressure on a plastic surface The application of a tangential force Q produces an increase in contact area even
though the nomal load P remains constant. This process has been called junction
growth, and it can be examined using a sliding wedge and a flat. The relation between
the coefficient of friction, junction growth, and interface strength is shown in figure 2.20.
The interface strength (the interface shear q/k parameter) is plotted versus the
coefficient of friction or traction coefficient (Q'P), and area growth (ALA,). In the event of
contamination or lubrication the interface shear (q/k) will be less than one and the junction growth process will be interrupted prematurely as result of sliding at the
contact interface. For values of q close to k (q/k-1) the coefficient of friction approaches
to one, a value obsenred in a chemically clean ductile material tested in vacuum
(Rabinowicz 1995). If contamination decreases the interface shear (qh) by one tenth, 46 the coefficient of friction (Q/P) decreases by one haIf. This situation explains the effects of surface contamination during friction welding; the wntaminants will decrease the coefficient of friction, avoid junction growth and surface plastic deformation during the friction welding process. As result, the generated heat input is low, and the temperature at the bondline will not be high enough to gromote formation of a plasticised iayer, a necessary condition for the weld to proceed.
2.6.4.2. The wave model Wear of the silver interlayer has been observed in dissimilar MMCfAgIAISI 304 stainless steel welds (see Chapter 5). Also, the silver interlayer affects the coefficient of friction during Stage I of the friction welding process (see Chapter 6). It is believed that slip-line field models such as the wave model (see figure 2.21) provide some insight concerning Wear and frictional behaviour of the soft interiayer during the friction welding operation. This model is applied in the present thesis in Chapter 5.
The wave model was developed to explain Wear and friction during the interaction between rigid and soft rnaterials. This model can explain different types of friction and
Wear regimes that are present when two surfaces contact. It should be pointed out that this type of slip-line field rnodel has also been proposed by other researchers for the analysis of problems such as junction formation, rolling contact, and wire drawing
(Johnson 1985 and Suh 1986).
It is assumed that the soft material is a rigid-ideally plastic material behaving under plain strain conditions. During contact between soft and rigid materials a wave forms as shown in figure 2.21, and the area of contact grows along ED until the system is in equilibrium. This process is aiso known as junction growth. 2Sofi layer Figure 2.21. Wave formation model (Challen and Oxley 1979).
In the slip line model the independent variables are the rigid asperity dope (a),
and the nonnalised strength V), see equation 2.15, of the interfacial film along ED:
where ris the interfacial film strength and k is the soft substrate shear strength.
The values of the nonnalised strength are within the range O values off correspond with a well-lubricated surface, or surfaces containing weak oxide layers or other type of soft organic films. A low value off results in limited junction growth, narrow contact area, low surface shear strain, and consequently low temperature in the contact region. In friction welding the presence of lubricant films or surface contamination markedly decreases the joint strength. For exarnple, it has been reported low strength properties of dissimilar pure aluminium/stainless steel friction welds when the contacting surfaces were initially contaminated using oil or grease (Jessop et al. 1978). Figure 2.22. Theoretical and experimental results for the coefficient of friction (Kopalinsky and Oxley 1995). The values of the coefficient of friction increases with a and5 see figure 2.22. For example, for f = 0.9 and a = IO0, the predicted coefficient of friction is 0.7; for f = 0.14 and a = IO0, the coefficient of friction is 0.15. These predicted values are in good agreement with experimentally measured results for 5086-H32-aluminum alloy base material (Kopalinsky and Oxley 1995). Detailed procedures on how to apply the slip- line model of figure 2.21 to calculate the effects of the interfacial film and surface roughness on the coefficient of friction are detailed in the work of Challen and Oxley. 2.6.4.3. Wave removal rnodel The repeated deformation of the contact surface will eventually remove the wave formed at the contact surface to produce a particle. Theoretically, the wave is removed when the angle ( is ( O (see figure 2.23) and this condition is given by equation 2.16 Figure 2.23. Wave removal slip-line model (Challen and Oxley 1979). (Challen and Oxley 1979). The value of # depends on a and f. High values of a and f increase the chances of particle formation. These conditions correspond to the presence of a rough or clean oxide-free surfaces. An important aspect is the amount of strain that the surface can sustain. The strain can be calculated using the upper bound rnethod. The slip line field of figure 2.21 is modified as shown in figure 2.24 and the upper bound method is applied to calculate the shear strain (y) at the surface, see equation 2.17. 50 At' Pde Figure 2.24. Simplified wave model used in calculating strains (~o~alinskyand Oxley 1995). Where v, , V, , and v;, are the velocity changes parallel to the discontinuity lines AB, BE, and CD, see figure 2.24. v>, vL,and v&-, are the normal velocities to the corresponding discontinuity line. The values of the previous velocity parameters depend on a and f: Kopalinsky and Oxley rneasured experimentally the shear strains on the surface during wave formation and wave rernoval. They found that the waves did not exhibit cracks for shear strains 40;above this value cracks were apparent and the wave was removed producing a particle. This model is applied in Chapter 5 to explain particle formation and Wear of the silver interlayer. CHAPTER 3. EXPERIMENTAL PROCEDURES 3.1. MATERIALS Al! dissimilar friction welds were made using 19-mm bars of 6061-T6 base material containing '!O vol% of reinforcing AI2O3particles. This material was supplied by ALCAN LTD of Kingston, Ont. The MMC material displayed some evidence of particle clustering and had a banded morphology as a result of particle alignment and agglomeration during processing. This morphology produced a slightly inhomogeneous and anisotropic composite base material. Figure 3.1 shows the typical base rnaterial microstructure. The chernical compositions and metallographic characteristics of the MMC and AlSI 304 stainless steel base materials are given in tables 3.1 and 3.2. One objective of this research work is to study the properties and behaviour of dissimilar friction welds with and without a silver interlayer. The required silver interlayer is electrodeposited onto the stainless steel substrate, which has been previously coated with a nickel strike layer. The nickel strike serves as a base for subsequent coatings and it has been used in dissimilar friction welding and dissimilar diffusion bonding; in these cases a 5-pm thick nickel strike was applied when electroplating stainless steel with chromium, zinc, and silver (Dunkerton 1982 and Dini et al. 1983). Figure 3.1. An optical micrograph showing the microstructure the reinforced 6061 AI-10 vol% Al2&. Table 3.1. Chernical composition of materials (wt %)' MMC Al Mg Si C Cu Fe Zn 606 1 97.76 1.15 0.535 0.099 0.225 0.121 0.022 STEEL C Si Mn Ni Cr Mo V AlSl 304 0.040 0.006 1.15 9-5 17.9 0.540 0-08 *As deterrnined by a commercial laboratory Table 3.2. Geometric characteristics of reinforcing particles Reinforced 6061 Al Average radius Average particle Aspect ratio Pm area Blocky AI2O3 5.82 156.81 1 .O21 (Mean value) 1 Standard deviation 1 5.30 1 289.3 1 0.0434 Figure 3.2. A backscattered SEM micrograph showing the silver interlayer at the bondline of dissimilar YMClAglAlSl 304 stainless steel friction joints. 54 The electroplating procedure cornprised surface degreasing and cleaning in! 10 vol% NaOH for 2 min, followed by deposition of an 8 pm thick nickel barr-ier layer in a nickel chloride bath for 5 minutes. The current density was 538 Ah2; the temperature bath was 25 OC. Silver electroplating was carried out for 20 min using a silver potassium cyanide plating solution (with a current density of 50 ~rn~and a temperature bath of 25 OC). The average thickness of the silver interlayer was 20 Pm. The plating procedure applied to the stainless steel substrate is based on the procedu re outlined for dissimilar 6061 aluminium alloylAlS1 304 stainless diffusion joints by Dini et al. The electroplating was carried out by a commercial plating Company in Toronto, Ont. Figure 3.2 shows a completed weld containing the silver interlayer. 3.2. FRICTION JOlNlNG The contacting surfaces of stainless steel and MMC substrates were machined perpendicular to the axis of the as-received bar. The test sarnple deviation from perpendicular axis was 4 degree. Al1 stainless steel test sampies were subsequently polished using 1 pm diarnond particles prior to friction joining, the MMC substraaes were polished using 1200 grade emery paper and the electroplated stainless steel and the MMC specimens were cleaned using acetone. Adhesion between the nickel barrier layer and the stainless steel substrate substrates was improved via vacuum heat treatment of the electroplated stainless steel samples at 650 OC for 1 hour and at 800 OC for 15 minutes (Hartwig and Kouptsidis 1977). All welds were produced using a continuous-drive machine rated at 15 kW transmission power with a maximum axial thrust of 110 KN. The friction welding machine was provided with a self-centring hydraulic clamp and a butt recognition system with a resolution of 76.2 Pm. 55 The friction pressure (Pl) was varied from 30 to 240 MPa with the rotational speed held constant at 1500 rprn and the forging pressure equal to 240 MPa. However, in some trials friction pressures of 30 and 120 MPa were applied. A friction time of 4.0 s and a forging tirne of 1.0 s were applied in almost al1 cases. During short- term welding of dissimilar MMC/Ag/AISI 304 stainless steel welds, these joints were produced using friction times ranging from 0.2 to 1.2 S. During these tests the friction pressure was 90 MPa, the forging pressure was 15 MPa, and the rotational speed was 1500 rpm. This range of friction times encompasses the initial heating period (Stage 1) and part of the steady-state period (Stage II) in friction joining, see figure 2.2-2.3. The microstructure of material in the region immediately adjacent to the bondline of MMCIAISI 304 was investigated using specially-designed test joints where an MMC capping was applied to particle-free alloy 6061-T6 base material. A 500 pm thick disk of alloy 606A/A1203composite material was joined to a 19-mm diarneter section of alloy 6061-T6 base material using TLP (Transient iiquid Phase)-bonding. TLP-bonding was carried out at 580 OC using a 10 pm-thick copper foi1 for a holding time of 4 hours. Following TLP-bonding the aluminium alloy test sections were heat-treated at 530 OC for 2 hours and then aged at 160 OC for 16 hours. 3.3. METALLOGRAPHIC EXAMINATION All joints were sectioned perpendicular to the joint interface and polished. The dimensions of the silver and nickel interlayers and the reinforcing particle characteristics (average radius, particle area, and particle ratio) at the half-radius locations in the test joints were evaluated using a Global Lab SP0550 image analyser and SEM microscopy. During image analysis the magnification was x250 and the 56 measurernent involved examining 0.158 mm2 fields at 0.318 mm distances up to 6 mm from the dissimilar joint interface. It is important to stress that the particle size distributions and the aspect ratio in the as-received base material are quite different from those characteristics in material irnmediately adjacent to the bondline- The as-received base material had a wide range of particie dimensions and aspect ratios. It will be shown later in this thesis that a much narrower particle size distribution and aspect ratio exists in material close to the bondline because of particle fragmentation during the joining operation. The friction welding process produced clearly measurable differences in the dimensions of particles at the bondline compared to those dimensions in material away frorn the joint interface. Scanning electron microscopy was used to examine the frequency of fractured reinforcing particles and the aspect ratio of reinforcing particles at and close to the bondline. These parameters were examined in 0.232 mm2 areas at the joint interface (using scanning eiectron micrographs taken at x400 and x500). The aspect ratios of the reinforcing particles were evaluated using the procedure outlined by Lewis and Withers- In this procedure, the particles are modelled as cylinders. Figure 3.3 shows how the dimensions of particles were assigned. In this case the particle diameter, dl is related to the measured width, w, by a factor 4/~see equation 3.1 (Lewis and Withers 1995). Once the diameter d, is obtained the volume of each particle can be calculated and therefore the accumulated particle volume in a particular location. Figure 3.3. Some examples illustrating the relationship between (a) regular particles, (b) irregular particles, and (c) misaf igned particles, and the cylindrical particles wh ich were used to represent them assuming the bar axis is from left to right (Lewis and Withers 1995). The average radius of each particle was taken as half the square of the mean of the particle cross-sectional area observed using the Global Lab SP0550 image analyser. When the short-term friction welding tests were examined, al1 measurements were made at the half-radius locations in these joints. During TEM microscopy the sample preparation technique was as follows. A section of the dissimilar joint was cut into slices and polished to a thickness of 0.120 mm. 3-mm dis- were then punched out from each of the slices. Each disc was then thinned to a thickness of 15 pm using a dimple grinder. A perforation was made in the centre of the test specimen using an ion beam-milling machine with the test specimen rotated during bombardrnent at incident angies of 15'. 58 3.4. MECHANICAL TESTING Notch-tensile testing was applied throughout this study since the notch stimulates bondline failure and altows a quantitative estimate of final weld quality- Since notch- tensile testing promotes bondline failure, it well-illustrates the effects produced when welding conditions are varied. Weld - d=3.01 R 6 R 0.8 133.4 -- - C Units: mm Figure 3.4. U-notch tensile testing specimen configuration. The design of the notch tensile test specimen is shown in figure 3.4. The tensile strength properties of the MMC and AlSI 304 stainless steel base materials were evaluated using standard round tension test specimens. The aluminium alloys were tested according to the ASTM B 557M standard, while the steel was tested using the standard ASTM A 370 test. The dimensions of the test specimens are shown in figure 3.5. The mechanical properties of materials are given in table 3.3, and the corresponding stress-strain diagrams in figures 3.6 and 3.7. ALUMINIUM ALLOYS STEEL G-Gage length 45.00 2 0.09 mm 35.0 2 0.1 0 mm D-Diarneter 9.00 -+ 0-10 mm 8.75 2 0.18 mm - 1 R-Radius of fillet, min 1 8 mm 6 mm A-Length of reduced section. min Figure 3.5. Round tension test specimen. Table 3.3. Mechanical properües of materials MATERIAL ULTIMATE YIELD ELONGATION E (GPa) POISSON'S- TENSILE STRENGTH (%) RATIO (v) STRENGTH (MPa) (MW MMC 338(303)'11 296(255)'" 7.6"' 81 -4"' (10 vol%) A1203 MMC 322"' 290'~' 7.gtL' - 1. Supplied by ALCAN? The number in parentheses is the minimum value reported by the material supplier. 2. Measured by the author 3. Page 932. (Mochida et ai. 1991) 4. Page 8. (Hertzberg 1989) 5. Page 34. Metals Handbook, Vol. 3., Ninth Edition 6. Page 384. (Kim et al. 1995) Figure 3.6. Stress-strain diagram for reinforced 6061-T6110 vol% A1203. Figure 3.7. Stress-strain diagram for AlSI 304 stainless steel. CHAPTER 4. INTERLAYERS AND PARTICLE FRACTURE IN DlSSlMlLAR FRICTION JOINTS 4.1. INTRODUCTION During Stage I of friction welding, the physical situation is similar to that in siidhg Wear testing. In sliding Wear testing of aluminium-based alloy A356120 vol% Sic base rnaterial, strains as high as 30 are produced near the contact interface and pairticle fracture occurs when the applied load exceeds the fracture strength of the reinfodng material (Alpas and Zhang 2992). It is possible to modify the particle fracture frequiency by modifying the loading conditions at the contact surface. In the present thesis the effect of a silver interlayer on particle failure is exarnined and compared with that observed in dissirnilar MMCfAISI 304 stainless steel welds produced without interlayer. The presence of a soft interlayer material can affect the coefficient of friction, plastic deformation in the substrate materials, and particle failure. The processes that occur during Stage I of the friction welding operation are also investigated by examining changes in size and distribution of reinforcing particles in MMC base rnaterial immediately adjacent to the joint interface. The likelihood of particfe fracture decreases during tensile testing ai temperatures above 200 OC because the matrix flow strength decreases and the local stresses are not high enough to break the reinforcing particles (Zhao et al. 1994). Consequently, there is a transition from particle fracture to particle void nucleation at the ends of particles and in the regions between particle clusters when the testing temperature 61 Cracking Shattering Debonding Figure 4.1. Modes of particle failure observed during mechanical testing of a particulate reinforced metal matrix corn posite material. rises. In friction welding the temperature rises with friction tirne, therefore it is convenient to observe particle fracture for short friction times. 4.2. RESULTS 4.2.1. influence of friction weiding on particle fracture Three particle failure modes are commonly observed during mechanical testing of wrought MMC base material, namely: cracking, shattering, and debonding (see figure 4.1) (Mochida et ai. 1991). However, in friction welded joints only debonding and particle cracking were observed in material close to the bondline (see figure 4.2). It is possible to have particle shattering, but it is difficult to observe because large plastic strains are produced close to the friction weId interface and the broken particles rnight have moved apart and separated; in such an event, evidence of this particle failure mode may have been obliterated. Figure 4.2. Particle fracture in dissimilar MMCIAISI 304 stainless steel friction welds. Figure 4.3 shows the changes in the number of reinforcing particles and the particle aspect ratio in rnaterial close to the bondline (in an MMC/AISI 304 stainless steel weld and in as-received MMC base material). The measurements were made at the half-radius location in the joint (see Chapter 3). The average aspect ratio of reinforcing particles at the dissirnilar joint Interface was 0.55 while that in the as- received MMC base material was 1.02. The aspect ratio was measured parallel to the longitudinal axis of the rod. The friction welding process altered the particle orientation and re-aligned particles parallel to the joint interface. Figure 4.4 shows the influence of the friction welding operation on the average particle radius at the bondline of a dissimilar MMCIAISI 304 stainless steel friction joint; this situation is compared with that found in as-received MMC base material. Larger numbers of small-diameter reinforcing particles were observed in rnaterial close to the 64 joint interface. Also, the total volume of reinforcing particles (529 369 in an area (0.232 mm2) close to the bondline was up to 2.8~higher than in the as-received MMC base material (187 922 see figure 4.5. All measurements were made at the half- radius location in the weld. The particle volume was evaluated by taking into account the volume of individual particles in the area concemed assuming that al1 particles had cylindrical shapes, section 3.3 of Chapter 3 (Lewis and Withers 1995). 4.2.2. Influence of friction pressure Figure 4.6 shows the influence of friction pressure on the average particle radius in material close to the bondline in dissimilar MMCfAISI 7020 mild steel welds. The average particle radius decreased when higher friction and forging pressures were applied. However, friction pressures in the range from 90 to 120 MPa did not exert a noticeable effect on the average particle radius. 4.2.3. Influence of friction tirne Figure 4.7 shows the changes in the percentage of fractured particles as a function of fiction time. In welds produced using a short friction tirne (0.2 s) and a friction pressure of 120 MPa. The percentage of fractured particles in MMClAlSl 304 stainless steel friction welds decreased markedly with friction time. The percentage of fractured particles decreased from 11-1% to about 3.8 % when friction tirne increased from 0.2 to 4 S. In this connection, the percentage of fractured particles was measured along a line perpendicular to the dissimilar interface at the half-radius location of the dissimilar weld. . , Mdaxidition 100- -an fit Mean: 0.55358 80- Sd: 0.27940 w- 40- i .I,. O. O O. 5 1.0 1-5 20 25 ASPE- RAT10 ASPECT RAliO Figure 4.3. Effect of friction joining on the aspect ratio distribution in a 6061-10 vol% AI203 MMCIAISI 304 friction joint. Friction pressure, 120 MPa; forging pressure, 420 MPa; friction time, 4.5 s; forging time 1.5 S. (A) As-weld condition; (B) As-received condition. AVERAGE P.4RTICi-E RADIUS, pm Figure 4.4. Effect of friction joining on the number of same- size particles in 6061-10 vol% AI2O3 MMCIAISI 304 friction joints. Friction pressure, 120 MPa; forging pressure, 120 MPa; friction time, 4.5 s; forging time, 1.5 S. (A) As-weld condition; (B) As-received condition. CI) F8 fY d 8 fY W m Z4 Figure 4.5. Effect of friction joining on the parade volume in a 6061-10 vol% A1203 MMCIAISI 304 friction joint Friction pressure, 120 MPa; forging pressure, 120 MPa; friction time, 4.5 s; forging time, 1.5 S. (A) As-weld condition; (B) As- received condition. Pl= 30MPa PI=60MPa A Pl =A20 MPa Figure 4.6. Effect of friction pressure and forging pressure on the average particle radius in the region adjacent to the bondllne in 6061-10 vol% A1203/1020 mild steel friction joint For a friction time of 4.5 s, and a forging time of 1.5 S. Ail measurements were made at the half radius location. 4.2.4. lnterlayers and particle fracture Introduction of a silver interlayer between the contacting MMC and AISI 304 stainless steel substrates affected the particle fracturing process. The percentage of broken reinforcing particles ranged from 6.1 to 2.5% in MMCfAgJAISI 304 stainless steel friction joints this result compared to dissimilar MMCIAISI 304 stainless steel friction welds where particle fracture ranged from 11.1 to 3.8%, see figure 4.7. It is apparent from figure 4.7 that the percentage of fractured particles is affected by the introduction of a silver interlayer during dissimilar MMCfAISI 304 stainless steel friction TlME (SEC) Figure 4.7. Effect of friction time on the percentage of fractured particles in MMClAlSl 304 stainless steel and MMCfAgfAISI 304 stainless steel friction joints. Friction pressure, 120 MPa; forging pressure, 30 MPa; friction time frorn 0.24.5 s; forging time, 1.5 S. All measurements were made at the half radius location. welding. lntroducing a silver interlayer during MMCIAISI 304 stainless steel joining also changed the average particle radius at the bondline from 4.6 to 3.9 Pm, see figure 4.8. 4.2.5. Influence of sihrer interlayer on particle accumulation When a silver interlayer is introduced, the volume fraction of particles decreased in an area adjacent to the joint interface at the half-radius location of the joint (see figure 4.9). However, friction pressure did not markedly affect the volume fraction of particles in MMC/Ag/AISI 304 stainless steel welds. The particle volume fraction at the bondline was 8.2 % in a joint produced using a friction pressure of 30 MPa, while the particle volume fraction was 7.0 % in a joint produced using a friction pressure of 240 Figure 4.8. Effect of a silver interlayer on the average particle radius in the region adjacent to the bondline in an MMCIAISI 304 stainless steel friction joint Friction pressure, 120 MPa; forging pressure, 120 MPa; friction time, 4.5 s; forging tirne, 1.5 S. All measurements were made at the half radius location. MPa. In effect, friction pressure had negligible influence on particle volume fraction at the bondline. The above results apply to particle fracture in MMClAglAISI 304 stainless steel friction welds. In welds produced without silver interlayers and a friction pressure of 30 MPa, the particle volume fraction was 15.7 %. When the friction pressure increased to 240 MPa, the particle volume fraction was 17.8 %, see figure 4.9. In effect, the particle volume fractions were higher in dissimilar MMCIAISI 304 stainless steel friction welds free of silver interlayers. 30 MPa 240 MPa FRICTION PRESSURE, MPa Figure 4.9. Accumulated particle volume in dissimilar 6061-10 vol% AI2O3MMClAlSl 304 stainless steel and MMClAglAlSl 304 friction joints. Friction time, 4 s; forging pressure, 240 MPa; forging tirne, 1 S. All measurements were made at the half radius location. 4.3. DlSCUSSlON 4.3.1. Friction welding and reinforcemen t particle morphology Examination of the friction welds provided evidence of particle realignment, with the longer axis of the reinforcing particles being aligned parallel to the bondline (figure 4.3). This feature combined with an increased particle frequency (figure 4.4), markedly altered the particle aspect ratio measured in the direction perpendicular to the joint interface. The aspect ratio changed from 1.02 to 0.55 after friction welding. Also, the total volume of reinforcing particles at the bondline (at the half-radius location in the joint) was up to 2.8~greater than in the as-received MMC base materia! (see figure 4.5). The influence of friction welding on the fragmentation of reinforcing particles and on particle realignment at the bondline is similar to the changes that occur in the grain structure of complete friction welds. Plastic deforrnation resulting from the combined action of normal and torsional loads during friction welding promotes grain refinement and grain alignment in matrix material close to the interface (Midling and Grong 1994b), while the reinforcement particles close to the interface are realigned and reduced in size. 4.3.2. Effect of friction pressure. Figure 4.6 shows that the average particle radius decreased when high friction pressures were applied. However, friction pressures higher than 120 MPa did not further decrease the average particle radius. This observation may be due to the relative mechanical properties of the reinforcing particle and the matrix, which has Iimited shear strength. It is worth noting that the matrix shear strength limits the load that can be transferred to the reinforcing particles (Clyne and Withers 1993). 4.3.3. Influence of friction time The similarities between friction welding and Wear are restricted to the initial phase of the friction welding process (Stage 1). This stage involves short friction times, probably lower than 1 S. As pointed out previously, particle fracture has been observed in both sliding Wear (Alpas and Zhang 1992) and friction welds (Zhou et al. 1995). During sliding Wear a highly strained region close to the contact surface of the aluminium alloy substrate facilitates delamination Wear and particle fracture. Strains as high as 30 are produced immediately adjacent to the contact surface when an aluminium alloy A 356 20 vol% Sic substrate contacts a high carbon steel slider (Alpas and Zhang 1992). The highly strained region close to the interface exhibits extensive particle fracture in a similar manner to that observed in friction welded MMC base material. As friction joining proceeds, dry friction rapidly increases the temperature of the contact zone and the base material deforms plastically. As result, the load transferred to the reinforcing particle decreases and the Iikelihood of void formation at the matrix/particIe interface will increase. The dramatic decrease in the percentage of fractured particle observed at the bondline of welds produced using friction times exceeding 1 s (see figure 4.7) can be explained due to an increase in the matrix temperature (this promotes void formation rather than particle fracture). Figure 4.7 shows the change in percentage of fractured particles with friction time. The percentage of fractured particies decreased with friction time, attaining a stable value of 3.8 % in MMCIAISI 304 stainless steel friction welds and 2.4% in MMC/Ag/AISI 304 friction welds. These results can be explained as follows. Stage II initiates when a fully-plasticised region foms across the whole component diameter, see figure 2.4. In MMCfAISI 304 stainless steel welds, this condition may occur when a friction time between 0.8 s and 1.2 s is applied. Stage II in friction welding is characterised by low torque and high peak temperature values. The peak temperature that can be attained during similar friction welding of AI-Mg-SIC base material is 550 OC,and the measured peak temperature at 2.5 mm from the interface for a friction time of 0.9 s is 440 OC (Midling and Grong 1994a). The Iikelihood of particle fracture is very low for temperatures ranging from 440 to 550 OC; in this temperature range the MMC flow strength is around 25 MPa (Singh and Alpas 1996). This situation readily explains the negligible change in the percentage of fractured particles in joints produced using friction times 5 1 s (see figure 4.7). The percentage of fractured particles in material within 100 pm of the joint interface was as high as 11.1% in joints produced using a friction time of 0.2 s (see 74 figure 4.7). It is worth noting that the percentage of fractured particles in MMC friction joints is much higher than that observed in aluminium-based composite material which has been tensiie tested (Lloyd 1990). For example, during tensile test samples of 6061- 15 vol% AI2O3 base material, Lloyd found that less than one percent of the alumina particles fractured. However during sliding Wear testing of A356-20 vol% SIC base material, 20% of the reinforcing Sic particles were fractured in material 100 prn from the contact surface (Alpas and Zhang 1992). It must be borne in mind that during friction welding the particles are subjected to different types of loading (compression plus torsion). 4.3.4. Modelling of protruding particle fracture The following rnodel considers fracture of reinforcing particles at the contact surface behveen stainless steel and MMC base material with a normal contact without sliding. When the two substrates corne together, contact occurs between the stainless steel and reinforcing particles (assurning that ail the particles at the surface protrude uniformly from the surface of the MMC substrate). The magnitude of the contact pressure will depend of the combined elastic modulus of the contacting substrates and on the surface topography. Since the MMC substrate has a wavy surface profile (see figure 4-10), the initial substrate contact is considered as Hertzian. In this wavy model, the distance between the contact points is equal to the interparticle distance (A). The interparticle distance is calculated using equation 4.1 (Zhang et al. 1995a), where d is the particle diarneter and f,the particle volume fraction. MMC Figure 4.10. Regular wavy surface rnodel. The load acting on any protuberance is given by equation 4.2 (Johnson 1985): P = (4-2) where p = nominal pressure, Pa , A = interparticle distance, ml P = normal load acting on the asperity, N. From equation 4.1, for an MMC cuntaining 10 vol% of reinforcing particles having an average particle diameter of 11.64 Pm, the interparticle distance (A) 2)s 104.8 Pm. Considering an example with a nominal pressure (p) of 0.5 MPa, using equation 4.2 the normal load acting on every asperity is 5.49~10~~N. The contact pressure (p,) can be calcuiated using equation 4.3. The equivalent elastic modulus (E*)is given by the relation: where VI, vz , El, and E2 are the Poisson's ratio and elastic modulus of substrate 7 and substrate 2. respectively (both substrates are in contact). The equivalent ratio of cutvature (R) is obtained from equation 4.5: where A is the asperity protuberance, see figure 4.1 0. The elastic rnoduli and Poisson3 ratios of MMC (E = 81.4 GPa and v = 0.322) and AIS1 304 stainless steel (E = 196 GPa and v = 0.3) can be combined to produce the equivalent elastic modulus (E* = 63.88 GPa). For a protuberance (A) of 1 Pm, the equivalent radius of cutvature (R) is 2.784~10~m, see equation 4.5. and the contact pressure @,) acting on the protuberance can be obtained using equation 4.3. Substituting these values, the contact pressure (p,) is 384.4 MPa. For a protuberance (A) of 5 Fm, from equation 4.5, the equivalent ratio of curvature (R)is 5564x10" m. and 77 from equation 4.3 the contact pressure (p,) is 1118.8 MPa. The fracture strength of Al& is about 2620 MPa in compression and 262 MPa in tension (Dowling 1993). The alumina strength is in the range of the calculated contact pressure (384.4-1118.8 MPa). Thus, the present results indicate that fracture of reinforcing particles may occur at low friction pressure. It is worth mentioning that this calculated value is in the range of the experimentally measured and calculated normal pressure value of 1 -06 MPa found during sliding Wear testing of an aluminium-based composite containing 10 vol% A1203 (Zhang et al. 1995b). This value is the critical nominal pressure necessary to produce extensive Wear and involves the detachment and fracture of reinforcing particles at and close to the surface of the MMC substrate. Since the friction pressure applied during friction ranges from 30 to 120 MPa and higher, the fracture of reinforcing particles must occur early in the friction welding process, immediately following contact between the adjoining substrates. Figure 4.1 1 shows the variation in the contact pressure, p,, for different base materials (MMCIMMC, 606?/MMC, Ag/MMC, and MMC/AISI 304 stainless steel). This contact pressure was calculated assurning a nominal pressure, p, of 0.5 MPa during welding. Based on a bi-dimensional wavy surface model assumption, the contact pressure will increase when the equivalent elastic modulus (E) increases. However, the particle fracture tendency during friction joining cannot be wholly ascribed to changes in the contact pressure resulting from the combined elastic modulus value at the contact interface. This mode1 can only be applied to reinforcing particles on the contacting substrates and therefore to friction welds produced using short friction times (4 s). However, this model cannot explain the particle fracture observed at a distance Figure 4.11. Effect of the combined elastic modulus (E*)on the contact pressure (po). Assuming a nominal pressure, p, of 0.5 MPa, an interparticle distance (A)of 104.8 Fm, and protuberance (A) of 1.O Pm. of 100 Orn from the contact surface. In this case, other factors such as plastic deformation rnust be considered. Although the contact model has limitations, it indicates that particle fracture at the surface can be produced at low appiied pressures (0.5 MPa). The calculated value in the present thesis therefore indicate that there is no possibility of avoiding particle fracture based on a consideration of the friction pressure values typically ernployed during fabrication. For example, 30 MPa was the lowest friction pressure ernployed during friction welding in this study. 4.3.5. Retention of fractured particles et the bondline MMC/AISI 304 stainless steel friction welds show evidence of an increase in particle volume at the bondline, see figures 4.4 and 4.5. This effect was not observed in joints produced using a silver interlayer, see figure 4.9. Figure 4.12. Effect of oxide Iayers on the notch tensile strength of MMCIAlSl304 stainless steel friction joints. In similar MMC/MMC welds it has been confirmed that material is retained at the bondline during the whole of the steady-state period (Stage II). It has been suggested that particle retention results from fluid flow during the friction joining operation (Bendzsak et al. 1997). In similar MMClMMC joints, there is preferential transfer of reinforcing particles from the stationary to the rotating boundary of the joint. In dissimilar MMC/AISI 304 stainless steel joints, a quiescent layer forms immediately adjacent to the stainless steel boundary and it has been suggested that the absence of 17owat the joint interface promotes retention of alumina particles at the bondline (North et al. 1997). However, retention of fractured particles at the joint interface rnay be explained using another approach. Heteny and McDonald examined the behaviour of a serni- infinite surface under the simultaneous action of normal and torsionai ioads. When this semi-infinite surface becomes fully plastic (Poisson's ratio equal to 0.5), the radial 80 surface displacement (u,) (see equation 2.11) will be zero and the only non-vanishing displacement components will be the circumferential (u~)and axial (w)components of displacement. In the absence of radial displacement, it follows that particles are retained at the bondline. Particle retention can be explained using either Heteny's plasticised surface model or Bendzsak's fluid flow model, or combining both models. The critical point is that the expulsion of surface oxide debris or trapped reinforcing particles from the bondline becomes difficult as the friction welding operation proceeds. The above mechanisms can afso be used to explain the welt-documented influence of surface finish prior to joining on the joint strength of dissimilar friction welds, see figure 4.12. Entrapment of oxide layers at the bondline of completed joints will decrease the notch tensile strength and promote preferential failure at the joint interface. 4.3.6. EHect of the interlayer on the stress distribution Analytical methods (Johnson 1985) can be used to provide a simplified picture of the stresses produced at the contact region (see equations 2.9 and 2.1 0).Consider that the surface is subjected to a line load (see figure 2.17). This Iine load is cornposed by a normal load (P)and a tangential load (Q) related by the coefficient of friction (p). In this particular analysis two coefficients of friction are considered (p=0.35 and 0.50). It is shown later in this thesis (Chapter 6) that electroplating silver onto the stainless steel substrate decreases the coefficient of friction during dissirnilar friction welding reaching a calculated value of 0.35. Figure 4.13 shows the stress (o,)distribution produced along the contact surface (y = O) when a semi-infinite substrate is Iine loaded. Both tensile and compressive stresses are generated at the contact surface (at y = O) in the region close to the a -25 -20 -1.5 -1-0 5 0.0 0.5 1.0 1.5 20 25 DlSTANCE (x), UNIT OF LENGTH Figure 4.13. Distribution of normal stress (a,) produced by the combination of normal and tangential line loads along the contact surface (y = O unit of length). P = 1 unit of force per unit of length. loading point. With a coefficient of friction of 0.35, at the contact surface and distances x = + 0.15 frorn the loading point the stresses can be as high as 1.49~(tensile) and - 1.49~(compression) the magnitude of the applied normal pressure (P). At the same location but with the coefficient of friction of 0.5, the stress values are + 2.12~the magnitude of applied pressure. At y = 0.10 the calculated tensile stresses range from 0.18~to 0.35~for coefficients of friction from 0.35 to 0.5, respectively, see figure 4.14. These calculated values show that a lower coefficient of friction (p=0.35) decreases the value of stresses developed at the contact sutface and in the subsurface region. Consequently, in dissimilar MMC/Ag/AISI 304 stainless steel, reinforcement particles close to the contact surface will be subjected to lower stresses. This condition might explain the DlSTANCE (x), UNiT OF LENGlH Figure 4.14. Distribution of normal stress (G,J produced by the combination of normal and tangential line loads along the line at y = 0.1 unit of length. P = 1 unit of force per unit of length. lower particle fracture tendency observed in MMC/Ag/AISI 304 srtainless steel welds (see figure 4.7). It must be borne in mind that a sirnplified approach has been employed in the present thesis and more detailed methods considering other loading situations have been elucidated elsewhere (Johnson 1985 and Suh 1986). Finally, based on the previous results, the contact surface wlll be subjected to a combination of tensile and compressive stresses during friction joining whose magnitude will decrease with depth below the contact surface. The nature of the applied stresses (tensile and compressive) will influence the surface damage mechanisrn. 83 4.3.7. Summation The percentage of fractured reinforcing particles markedly decreases when silver interlayers are introduced during dissimilar MMC/AISI 304 stainfess steel friction welding. An elastic contact mode1 is proposed, which explains the conditions at and close to the contact surface that produce A1203 particle fracture and explains the beneficial effects observed when silver interlayers are introduced. This mode! is the first time that this particular aspect has been explained during dissimilar friction welding of MMC substrates CHAPTER 5. MICROSTRUCTURE OF A DlSSlMILAR FRICTION JOINT CONTAlNlNG A SILVER INTERLAYER 5.1. INTRODUCTION In past studies the influence of interlayer materials on the friction joining process and on interrnetallic phase formation at the bondline region has received limited emphasis. With this need in mind, the present chapter examines the effects produced when a silver interlayer is introduced during friction joining of aluminium-based MMC and AIS1 304 stainless steel base materials. Prior research on dissimilar friction joining has emphasised the influence of interlayer materials on joint mechanical properties. For example, Sassani and Neelam investigated the influence of rnetallic interlayers on the mechanical properties of dissimilar brass/copper, bronzekteel, and titaniumlnickel friction. The introduction of a pure copper interlayer during bronzelsteel joining increased the room tensile joint strength by 40% joints (Sassani and Neelam 1988). In dissimilar aluminium/steel friction joints the introduction of a silver interlayer produced welds with tensile strengths exceeding the weaker (aluminium) substrate (Hartwig and Kouptsidis 1977). In this connection, Dunkerton examined dissimilar friction joints produced using AISI 321 and AlSI 304 stainless steels, and 5154 and 5083 aluminium alloys base mâterials A nurnber of interlayers (zinc, chromium, and silver) were investigated. 60th 5 pm thick zinc and chromium electroplated layers produced poor quality welds between 5154 aluminium alloy and 321 stainless steel. The zinc interlayer 85 did not bond readily to the aluminium alloy and the chromium electroplate demonstrated poor adherence with the steel. When 13 pm thick silver electroplate was applied to join 5083 aluminium alloy to 304 stainless steel there was a significant improvement corn pared to interlayer-free joints (Dunkerton 1982, 1 983). The introduction of a soft (low yield strength) interlayer material such as silver at the bondline during aluminium and stainless steel friction joining can produce unique mechanical properties in this region. Prior work has shown that the introduction of a silver interlayer at the bondline between aluminium and stainless steel substrates may not result in poor strength properties compared to the monolithic interfayer material (Henshall et al. 1990). Also, if the silver interlayer inhibits brittle FeAl intennetallic layer formation at the bondline the final joint strength can be greater than in joints completed without an interlayer. A key feature of any interlayer material is that it must promote excellent adhesion between itself and the contacting substrates (Sassani and Neelam 1988). However, It has been pointed out in his compatibility chart that there is poor compatibility between silver and steel (Rabinowicz 1995). The compatibility is considered as the tendency between two metals to form solid solutions. Since substrates such as silver and iron have low compatibility, poor mutual adhesion is expected. However, the combination of aluminium and silver substrates has good compatibility and as result, there is good adhesion between them. Another aspect to consider when using an interlayer is the coefficient of friction. The introduction of an interlayer modifies the coefficient of friction between contacting substrates (see Chapter 6). If the interlayer decreases the coefficient of friction, less heat will be generated during the friction welding process. With this behaviour in rnind, it would be expected that the joining parameters required during dissimilar friction 86 welding using an interlayer would be quite different from those parameters needed when an interlayer was not present. lnterlayers have been applied with the specific airn of preventing interrnetallic phase formation and weld cracking during aluminium/stainless steel friction joining. The presence of FeAI3, FeAI. and Fe2AI5intennetallics has also been detected in dissimilar aluminium/stainless steel joints (Elliot and Wallach I98la. b). Also, in a recent study. it is suggested that the poor joint mechanical properties of aluminium-based MMCIAISI 304 stainless steel friction welds resulted from the retention of a mixture of intermetallic (FeA$) and oxide Fe(AI, Cr)*04 or F~O(AI,CT)~O~films at the bondline (Pan et al. 1996). Although intermetallic phases do not form in Ag-Fe binary alloys, the Ag-AI binary equilibrium phase diagram does indicate the formation of intermediate phases, Ag3AI and Ag2AI phase formation has been observed in 1 100 aluminium/Aglstai~lesssteel and 6061 aluminium/stainless steel diffusion welds. When Calderon et al. heat-treated their 6061 aluminium/stainless steel diffusion welds at 473 K for 100 hours the width of the intermetallic layer increased from 0.5 to 4 pm and the joint mechanical properties decreased markedly (Calderon et al. 1985). It is important to point out that remarkably different conditions exist at the bondline of diffusion and friction welded joints. During diffusion weiding, dissimilar aluminium and steel substrates are bonded at a specified temperature and interrnetallic layer formation results from interdiffusion during the holding period at the bonding temperature. During friction weiding of aluminium alloy substrates, the contact zone is at high temperature (around 550°C), the strain rate is extremely high (up to lo4 s-') (Midling and Grong 1994b) mechanical mixing occurs and flow of fully-plasticised material redistributes material from the stationary to the rotating boundary of the 87 welded joint (Bendzsak et al. 1997). In dissimilar MMClstainless steel joints it has been suggested that the steady-state penod in friction joining is characterised by the absence of material Row due to the formation of dynamically quiescent layers immediately adjacent to the stainless steel boundary (North et al. 1997). It has been suggested that the presence of the quiescent layers may affect the size and formation of the intermetallic layers at the dissimilar interface. This suggestion may not be the case in diffusion bonding because the presence of quiescent layer has not been detected in this specific process. The present chapter investigates the influence of a silver interlayer on intermetallic formation durÎng MMC/stainless steel friction welding. MMC/AISI 304 stainless steel friction welds produced without silver interlayers were examined for cornparison purposes. 5.2. RESULTS 5-2- 1. lnterlayer microstmcture The microstructure of the silver interlayer prior to the friction welding operation was examined using transmission electron microscopy. The silver interlayer exhibited spot diffraction patterns, see figure 5.1. The spot diffraction pattern suggests that a single grain is diffracting the electron beam. The micrograph was obtained at relatively high magnification (x100 000) and no grain boundaries were observed. It can be concluded that the silver interlayer is polycrystalline prior to the friction welding operation. Figure 5.1. TEM micrograph of the silver interlayer: (A) bright field image; (6)corresponding selected area diffraction pattern, and (C) key diagram showing Ag. 5.2.2. As-welded joint microstnrcture A wavy interface profile formed at the bondline of MMCIAgIAISI 304 stainless steel friction joints (see figure 5.2). The different regions in figures 5.2 and 5.3 comprise AlSI 304 stainless steel, the nickel barrier layer separating the AlSI 304 and the silver interlayer, an intermixed region (IM) containing a mixture of silver, an intemetallic phase, and entrained Al& particles. The veined microstructure in the IM region comprised 4 prn thick silver ribbons that separated regions of the intermetallic phase (see figure 5.3). EDX analysis of materiai in the IM region indicated that the 89 intemetallic phase had an approximate chemical composition comprising 19.1 at% Al - 80.9 at% Ag (see figure 5.4). Ag3AI formation was confirmed during subsequent transmission electron microscopy and X-ray diffraction analysis, see figures 5.5 and 5.6. The composition of the intemetallic phase corresponds to the p phase in the binary equilibrium phase diagram (see figure 2.7). In this connection, Dunkerton also indicated the formation of an intermetallic layer in alum inium/Ag/stainless steel friction joints (Dunkerton 1982, 1983). However, this investigator did not identify the intermetallic compound. A wide particle-dispersed region (PD) region was observed immediately adjacent to the IM region (see figure 5.7). EDX analysis confirmed that the PD region cornprised a mixture of granular silver, Ag3AI, aluminium plus entrained A1203 particles. A PD Figure 5.2. Optical micrograph of the bondline of a dissimilar MMClAglAlSl 304 friction joint, showing the wavy interface morphology produced by the friction welding operation. Figure 5.3. Backscattered micrograph of a dissimilar MMCIAgIAISI 304 stainless steel friction joint showing the intemixed region (IM) containing alumina particles. The nickel layer is not visible in this micrograph. region was also observed in MMC/AISI 304 stainless steel joints where the alloy 6061- T6 substrate was capped with a 500 prn thick MMC layer prior to the joining operation (see figure 5.8). The TLP bonding procedure used in capping the 6061-T6 is detailed in Chapter 3. section 3.2. In this case. the PD region was 500 pm wide at the component centreline. The characteristic rotational flow patterns in figure 5.8 are the result of mechanical rnixing. Figures 5.9-5.1 1 show the development of the MMC/Ag/AISI 304 stainless steel joint interface when the friction time increased from 0.2 to 1.2 S. The width of IM and PD regions varied along the bondline. large IM and PD regions were observed in the 0.2 s friction weld; the widths of the IM and PD regions at the half radius location were X-ER'?: O - 20 KeU Livc: 50s Préset.: 50s Remai ri i ng: O 4 Real: 102s 51% Dead Figure 5.4. EDX analysis of the intermetallic compound of the IM region. Friction pressure, 240 MPa; forging pressure, 240 MPa; friction time, 1 S. Location at 1 mm from the joint periphery. See Figure 5.3. 44 and 120 pm, respectively (see figure 5.10). However, the widths of the IM and PD regions markedly decreased in joints produced using a friction time of 1.2 s, the width of the IM region was 10 pm while the width of the PD region was 39 pm (see figure 5.10). Dissimilar MMC/Ag/AISI 304 stainless steel welds produced using a friction time of 1.2 s showed a pronounced shift in the location of thick PD region from the component centreline towards the weld periphery compared to dissimijar friction welds using a friction time of 0.2 s, see figure 5.1 1. The present results indicate that the widths of the IM and PD regions change as the friction welding operation proceeds. Figure 5.5. TEM micrograph of the IM region: (A) brigth field image, the dark phase is Ag3AI; (8) corresponding selected area diffraction pattern, and (C) key diagam confirming AgJAl formation in the IM region. Friction Pressure, 90 MPa; forging pressure, 15 MPa; friction time, 1.2 S. Half-radius location. Figure 5.7. PD region in dissimilar MMCIAglAISI 304 comprising parücles of Ag3A1, Ag, and Al as identified by EDX. Friction pressure, 90 MPa; forging pressure, 15 MPa; friction time, 0.2 S. Half-radius location. 'igure 5.8. Optical micrograph showing the mesence of a PD region in a 60611MMCIAlSl 304 rtainless steel friction joint The arrow shows the oint centreline. Figure 5.9. Effect of friction time on the microstructure of a dissimilar MMC/Ag/AISI 304 stainless steel friction joint: (A) friction time, 0.2 s, and (B) friction tirne, 1.2 S. Friction pressure, 90 MPa; forging pressure, 15 MPa; forging time, 1.0 s, and rotational speed, 1500 rpm. Backscattered micrograph. Centreline location. FRICTION TIME, s Figure 5.10. Effect of friction time on the thickness of the PD and IM regions in dissimilar MMCIAISI 304 stainless steel friction joints. Friction pressure, 90 MPa; forging pressure, 15 MPa; forging time, 1.0 s, and rotational speed, 1500 rpm. Half-radius location. SPECIMEN RADIUS, mm "j- 12 rn 4 SPECIMEN RADIUS, mm Figure 5.11. Distribution of IM and PD regions along the bondline of dissimilar MMCIAglAISI 304 stainless steel friction joints as function of friction tirne. A) Friction time, 0.2 s; B) friction time, 1.2 S. Figure 5.12. Effect of friction pressure on the microstructure of dissimilar MMCIAglAlSl 304 stainless steel friction joints: A) friction pressure, 30 MPa, and B) friction pressure, 240 MPa. Friction time, 4s; forging time, 1 s, and rotational speed, 1500 rprn. Backscattered micrographs. Centreline location. 120 MPa FRICTION PRESSURE, MPa Figure 5.13. Effect of friction pressure on the thickness of the PD and IM regions in dissimilar MMCIAglAISI 304 stainless steel friction joints. Friction time, 4s; forging pressure, 240 MPa; forging time, 1 s, and rotational speed, 1500 rpm. Half-radius location. SPECMEN RADIUS. mn P- s m- 8 'O- 0 -r -10 SPMMEN RAMUS, mn Figure 6.14. Distribution of IM and PD regions along the bondline of dissimilar MMCIAglAlSl 304 stainless steel friction joints as function of friction pressure. (A) friction pressure, 30 MPa, and (B) friction pressure, 240 MPa with a small amount of IM region (1 pm thick) at the centreline location. 102 Figures 5.12-5.14 show the changes in dimensions of the IM and PD regions at different radial locations along the joint interface when friction pressures ranging from 30 to 240 MPa were applied. In the joint produced using a friction pressure of 30 MPa. the width of the PD region was greatest at the half-radius location, see figure 5.14(A). When the friction pressure was 240 MPa, the IM region was only observed in locations near the cornponent centreline and close to the joint periphery. In this particular joint, the width of the IM region was 1 pm at the component centreline and was 12 pm in regions close to the joint periphery (see figure 5.14(8)). In summation, when the friction pressure increased from 30 to 240 MPa the width of the PD and IM regions decreased resulting in the removal of the Ag interlayer. Although, there is no experimental evidence, it is believed that the silver interlayer is totally or partially expelled during the flash formation. 5.2.3. Friction pressure and TEM microstrucfure The influence of friction pressure on the microstructure in dissirnilar MMC/Ag/AISi 304 stainless steel and MMC/AISI 304 stainless steel friction joints was examined in detail. The friction welding process deforms and removes the silver interlayer, promoting microstructural changes at and close to the dissirnilar joint interface. Using a friction pressure of 30 MPa, figure 5.15 shows that the interlayer at the bondline consists of small silver particles. Examination of the bondline region at high magnification (x100 000 and x200 000) together with the specific rnorphology of this region indicated the formation of silver nanoparticles, having sires in the range from 10-20 nm. When the friction pressure increased to 120 MPa, TEM examination of region at half radius location in the weld did not indicate the presence of silver nanoparticles, figure 5.16 shows the presence of silver grains. When the friction 1O3 pressure increased to 240 MPa, the microstructure at the centreline comprised recrystallized silver and Ag3AI grains (see figures 5.17-5.1 8). The interface between the nickel barrier film and the stainless steel substrate was also examined using STEM microscopy and EDX analysis. This interface confinned the presence of a complex Fe-AI-Ag-Ni-Cr compound (see figure 5-19) indicating that silver and aluminium diffused into the stainless steel and formed a complex intermetallic phase. Because of equipment limitations it was not possible to determine quantitatively the chemical composition of the observed intemetallic compound. A detailed examination of dissimilar MMC1'AISI 304 stainless steel friction welds confirmed the presence of a discontinuous transition layer along the length of the joint interface (see figures 5.20-5.23). EDX chemical analysis of the transition layer indicated peaks of aluminium and iron, see figure 5.22. The presence of FeAl and Fe4Al13 was confinned using TEM rnicroscopy at the centreline and at the half radius location in dissimilar MMCIAISI 304 stainless steel friction welds produced using friction pressures of 30 and 240 MPa, see figures 5.24-5.26. Figures 5.24 and 5.25 show the influence of the friction pressure on the dimensions of the intermetallic layer formed at the joint centreline, when the friction pressure increased from 30 to 240 MPa the thickness of the intermetallic layer decreased from 1.9 to 0.5 pn at the centreline. Analysis of diffraction patterns confirmed the formation of FeAI. These particular joints were produced using a friction time of 4 S. Figure 5.26 shows the presence of Fe4AIl3 af the dissimilar MMCIAISI 304 stainless steel friction joint interface produced with a friction pressure of 30 MPa and friction time of 4 s, the intemetallic compound was observed at the half radius location. Figure 5.27 shows the intermetallic layer at the half-radius and centreline locations in dissimilar MMC/AISI 304 welds produced using a friction pressure of 30 1O4 MPa and a friction time of 4 S. it is observed that the intermetallic layer is narrower at the centreline location compared to the half-radius location. Figure 5.28 shows the influence of the friction pressure on the width of the intermetallic layer formed ai the half-radius location of completed welds. When the friction pressure increased from 30 to 240 MPa, the thickness of the intermetallic layer decreased from 4.6 to 2 pm at the half-radius location. Figure 5.15. TEM micrographs of the IM region from the bondline of a dissimilar MMCIAgIAISI 304 stainless steel friction joint: (A)-@) bright field images; (C) corresponding ring diffraction pattern, and (D) key diagram showing the presence of silver. Friction pressure, 30 MPa; forging pressure, 240 MPa; friction time, 4 s; rotational speed, 1500 rpm. Centreline location. Figure 5.15 (Cont) TEM micrographs of the IM region from the bondline of a dissimilar MMClAglAlSl 304 stainless steel friction joint: (E) bright field images; (F) corresponding ring diffraction pattern, and (G) key diagram showing the presence of silver. Friction pressure, 30 MPa; forging pressure, 240 MPa; friction time, 4 s; rotational speed, 1500 rpm. Centreline location. Figure 5.16. TEM rnicrograph of the IM region: (A) bright field image; (B) corresponding selected area diffraction pattern, and (C) key diagram showing Ag. Friction pressure, 120 MPa; forging pressure. 240 MPa; friction time, 4 s; rotational speed, 1500 rpm. Half-radius location. Figure 5.16. (Cont) TEM micrographs of the IM region: (D) bright field image; (E) corresponding selected area diffraction pattern, and (F) key diagram showing Ag in the IM region. Friction pressure, 120 MPa; forging pressure, 240 MPa; friction time, 4 s; rotational speed, 1500 rpm. Half- radius location. Figure 5.17. TEM micrographs of the IM region: (A)-(B) bright field images; (C) corresponding selected area diffraction pattern, and (D) key diagram showing Ag. Friction pressure, 240 MPa; forging pressure, 240 MPa; friction time, 4 s; rotational speed, 1500 rpm. Centreline location. Figure 5.18. TEM micrographs of the 11111 region: (A)-(B) bright field images. Friction pressure, 240 MPa; forging pressure, 240 MPa; friction time, 4 s; rotational speed, 1500 rpm. Centreline location. Figure 5.18. (Cont) TEM micrographs of the IM region: (C) bright field images; (D) corresponding selected area diffraction pattern, and (E) key diagram showing Ag& Friction pressure, 240 MPa; forging pressure, 240 MPa; friction time, 4 s; rotational speed, 1500 rpm. Centreline location. III Figure 5.19. Microstructural features detected in a dissimilar MMClAglAlSl 304 stainless steel friction joint showing the boundary between the nickel layer and the stainless steel substrate. Friction pressure, 240 MPa; forging pressure, 240 MPa; friction time, 4 s; forging time, 1 S. The arrow indicates the particle from which the EDX pattern was obtained. Centreline location. Figure 5.20. Micrograph of a dissimilar MYCIAISI 304 stainless steel friction joint showing the transition layer. Friction pressure, 30 MPa; friction time, 4 s; forging pressure, 30 MPa; forging time, 1 S. Figure 5.21. Cracking at the half-radius location in a dissirnilar MMClAlSl 304 stainless steel joint. Friction pressure, 30 MPa; friction time, 1 s; forging pressure, 30 MPa; forging time, 1 S. Figure 5.22. EDX analysis of the transition layer. Friction pressure, 30 MPa; friction tirne, 4 s; forging pressure, 30 MPa; forging time, 1 S. DISTANCE ALONG BONDUE, mm Figure 5.23. Transition layer distribution along the interface. Figure 5.24. TEM micrographs at the bondline of a MMClAlSl 304 stainless steel friction weld: (A)-(B) bright field images; (C) corresponding selected area diffraction pattern, and (D) key diagram confirming FeAl formation. lntermetallic layer width, 1.9 Pm. Friction pressure, 30 MPa; forging pressure, 240 MPa; friction time, 4 s; rotaüonal speed, 1500 rpm. Centreline location. Figure 5.25. TEY mictographs at the bondline of a MMCIAISI 304 stainless steel friction weld: (A)-(B) bright field images; (C) corresponding selected area diffraction pattern, and (C) key diagram. Intennetallic layer width, 0.5 Fm. Friction pressure, 240 MPa; forging pressure, 240 MPa; friction üme, 4 s; rotational speed, 1500 rpm. Centreline location. Figure 5.26. TEM micrographs at the bondline of a MMCI.AIS steel friction weld: (A) bright field images; (C) corresponding selected area diffraction pattern, and (C) key diagram confinning Fedl13 formation. Friction pressure, 30 MPa; forging pressure, 240 MPa; friction tirne, 4 s; rotational speed, 1500 rpm. Half radius location. Figure 5.27. Comparison of intemetallic layer widths at the centreline and at the half radius location in the dissimilar joint produced using friction pressure of 30 MPa and friction time of 4 s; (A) at the half radius location, and (8)at the centreline. 30 MPa 120 MPa 240 MPa FRICTION PRESSURE, MPa Figure 5.28. Effect of friction pressure (Pl) on the intemetallic layer thickness in a dissimilar MMClAlSl 304 stainless steel friction joint Friction time, 4 s; forging time, 1 s, and rotational speed, 1500 rpm. 5.3. DISCUSSION 5.3.7. Influence of friction welding on the silver inferiayer Friction welding involves rapid heating caused by a combination of dry friction and numerous adhesion/seizure/failure events (Stage 1) prior to the establishment of a steady-state period when the torque, temperature distribution, and rate of axial shortening become essentially uniform (Stage II).The welding process is completed by the forging operation. During the steady-state period, the welding process depends on deformation at high temperature. At first sight it might appear that the steady-state period in friction joining would be the determining phase of the friction joining operation. However, the results in the present thesis suggest that the metallurgical and mechanical properties of dissimilar friction joints are determined by processes that occur very eariy during the first 0.5 s in the welding process. This observation is supported by the results exarnining particle fracture in friction welded joints (that show that almost al1 of the particte fracture process occurs early in Stage f of the welding process). The joint interface regions in dissirnilar friction welds had wavy surface profiles. This wavy profile was produced early in the process (it was observed in welds produced using a friction tirne of 0.2 s). This interface profile is in many respects similar to the surface layer observed during sliding Wear testing. During sliding Wear testing, the surface layer results from the transference of material from the substrate to the counterface. For example, in a study examining sliding Wear between a steel pin and a Cu-15 wt%Ni disk, nickel transfers preferentially to the 3-mm diameter steel pin and a widespread system of patches is formed on the surface (Rigney et al. 1984). These patches form after a relatively few (15) cycles of the 6-mm rotating disk (for a sliding 122 speed of 50 mm/s, a sliding distance of 56 mm, and a nominal load of 0.78 N). In friction welds produced using a friction time of 0.2 s and rotating speed of 1500 rpm, the sliding distance at the half-radius location is 113 mm. Thus, it would be expected that the wavy surface is fomed early in the friction joining process. 5.3-2. Similarities between sliding Wear and Stage 1 of friction welding Since the physical situation early in friction joining is similar to that during sliding Wear testing (Ames and Alpas 1995, Zhang and Alpas 1996) the results produced during Wear testing using aluminium-based MMC and steel substrates are of direct interest. A detailed microstructural examination of samples produced during sliding Wear testing using aluminium alloy and steel substrates has confirmed that fully- plasticised material forrns when the temperature at the contact interface exceeds 162 * 10 OC (Ames and Alpas 1995). This critical temperature marks the change from mild to severe Wear. Severe Wear is characterised by adhesion of aluminium material to the steel counterface and thermal softening which produces dynamically recrystallised grains in the aluminium alloy substrate. With this model in mind, it is suggested that the microstructures in the IM and PD regions (in the 0.2 s friction weld) are formed when the average temperature of material close to the joint interface exceeds a transition temperature which depends on the contacting substrates. One important aspect is the formation of nanoparticles in dissimilar MMC/Ag/AISI 304 stainless steel joints produced using a low friction pressure (30 MPa) and a friction tirne of 4 S. This nanocrystalline microstructure was observed in the region close to the component centreline and had particle dimensions ranging from 10 to 20 nm. Figure 5.15 shows bright field TEM micrographs of the interface region in a dissimilar MMC/Ag/AISI 304 stainless steel friction weld produced using a friction pressure of 30 MPa. Figures 5.17-5.1 8 show bright field TEM micrographs of the centreline location in 123 a friction weld produced using a friction pressure of 240 MPa. In this weld there was no evidence of nanoparticle formation. It is suggested that the nanoparticle formed in the silver interlayer (in joints produced using low friction pressures) may result from the formation of a transfer layer early in the welding process resulting from Wear or mechanical alloying. A nanocrystalline microstructure was observed in the transfer layer formed during sliding Wear testing involving OFHC copper and M2 tool steel counterface. The sliding speed was 0.1 mis, the load was 1.33 N, the testing time was 1 hr in argon atmosphere and the relative humidity was 0.21%. The Wear surface temperature was 56 OC (Rigney ef al. 1984). The size of the grains in the transfer layer depended on the sliding speed; Rigney et al. observed that the dimensions of the crystallites were larger when faster speeds were applied. These investigators associated the increased size with higher temperature being produced during the Wear process. Rigney et al. also indicated that the microstructure of the transfer layer was similar to that produced during mechanical alloying of OFHC copper and M2 steel balls. During mechanical alloying, the rnean processing temperature was 60 OC. Based on this similarity, they suggested that the transfer layer formation involved processes sirnilar to those process in mechanical alloying, i.e. plastic defornation and fracture of powder particles. The temperature at the contact interface had an important influence on nanoparticle formation and on the size of the crystals formed during the Wear process. The size of nanoparticles increased and the likelihood of nanoparticle formation decreased when higher temperatures were produced at the contact interface (Rigney et al. 1984). Nanocrystalline alloys are produced during mechanical alloying of intermetallic compounds or metallic powders. To produce a nanocrystalline structure it is necessary that the arnorphous phase should be less stable than the nanocrystalline phase since 124 this situation avoids transformation into the amorphous state (Oehnng and Bormann 1991). This requirement has been applied when examining the Ag-AI binary system. Paruchuri et al. produced nanocrystals of the a, p, and < phases when they investigated mechanical alloying of AgAl alloys. They studied four nominal compositions corresponding with the p, and < single phase fields, and with a+p, and pfc; phase fields in the binary phase diagram. The compositions of the alloys were 23, 37, 20, and 26 at% Al, respectively. The apparent particle sizes for a, p, and < phases were approximatety 13, 21, and 18 nm. During mechanical alloying they estimated that the temperature rise was around 100 OC. This increase in temperature occurred since the heat of mixing of the intermediate phases in the Ag-AI system is negative. It is worth mentioning that the t; phase forrned first when the alloy contained 23 at% Al; 8 hours of milling were required to produce the phase p (Paruchuri et al. 7994). In a dissimilar friction weld produced with a friction pressure of 120 MPa, the presence of silver was identified at the half radius location, but no nanoparticles were observed, see figure 5.16. The highest temperature is attained at the half radius location in the joint. Moreover, at this particular location the silver interlayer is worn away preferentially, see figures 5.9-5.14. As a result, the presence of nanoparticles would not be expected at the half radius location, especially in welds produced using increasing friction pressures. At the centreline location, in dissimilar friction welds produced with a friction pressure of 240 MPa, there is no evidence of silver nanoparticles. The TEM bright field micrographs show the presence of silver and AgAl grains, see figures 5.17-5.18. It is worth of noticing that the presence of Ag and AgAI grains is associated to the IM region. 53.3- Fo-n and removaf of- and 1M mgbms Thepresenceofentrained r&nfOrciingm partides,theveined sihrwm plus mer morphology in the IM region and the fad that the IM and PD regions were okewed in W-on joints produ& using a friction tirne of 02 s suggeçt that these mÏ~~OSfi'Ucturaffeatures formed as resutt of mechanid mEng earfy in the friction welding procesç. When the friction time increased to 1.2 s the widoi of the IM region decreased and locations at half radius of wrn~letedwelds showed no evidence of IM region formation (see figures 5.10-5-1 1 and 5.13-5.14). In a similar rnanner very thick PD regions were observed in a weld made using friction time of 0.2 S. In contrast, the joint produced using a friction time of 1.2 s had much thinner PD regions at the bondline. Also, there was a pronounœd shift in the location of the thick PD region from the component centreline towards the joint periphery when the friction time increased from 0.2 s to 1.2 S. When a higher friction pressure was applied (240 MPa) the IM region was only retained in locations near the component centreline and at the joint periphery. Also, there was no evidence of PD region formation along the rernaining length of the dissimilar joint interface (see figures 5.13-5.14). The shift in the location of the PD region from the component centreline towards the joint penphery can be explained as follows. When the two substrates contact each other, the normal pressure has its peak value at the component centreline and the shear stress has its highest value at the hat radius location in the cornponent (Heteny and McDonald 1954). (Johnson 1985). However, as friction welding progresses the normal and shear stress distributions are altered, Le. the normal pressure decreases at the centreline and increases near the periphery of the component. This type of behaviour has been observed when examining rotating spherical indenters (Johnson 1985). Bearing in rnind the changes in location of the normal and shear stress during 126 friction welding, it is suggested that thick PD region forms initially at the component half radius since this region deforms preferentially compared to the component centreline and periphery Another explanation for the shift in the location of the PD region relates to changes in surface displacement. For example, the contact surface is subjected to radial (u,), circumferential (II& and axial surface displacements (uJ. During Stage I of the welding process the radial displacement (u,) has an inward direction with a maximum value at the weld periphery. The circumferential displacement (uO) has a maximum value at the half radius location and decreases as the weld periphery and the centreline are approached. The axial displacement (uJ has a maximum value at the sample centreline (Heteny and McDonald 1954). Consequently, it is suggested that the shift of the IM and PD regions may result frorn the combined action of the radial and tangential displacements. The influence of friction time on the widths of the IM and PD regions can be explained due to the progressive wearing away of these regions as the friction welding process progresses. This suggestion is consistent with the marked influence of increased friction pressure in decreasing the widths of the IM and PD regions (see figures 5.12-5.14). For example, the maximum width of the PD region decreased from ,190 Fm to zero when the friction pressure increased from 30 MPa to 240 MPa. In this connection, Fuji et al. also suggested that the intemetallic layer at the dissimilar joint interface in TVA5083 aluminium alloy joints was acted on and was removed by the action of plasticised A5083 material being continuously squeezed out of the weld (Fuji et al. 1995). This particular process may explain the removal of the PD region in welds produced using high friction pressures. Although it is generally assumed that intemetallic layer formation in dissimilar friction welds results from interdiffusion and/or mechanical mixing during the welding operation, the exact manner in which this occurs is unclear. For example, Fukumoto et al. suggested that FeAI, FeAI and Fe2AI, formation in 1060 aluminium/ AlSI 304 stainless steel friction joint resulted from interdiffusion during the joining operation. Since the aluminium diffused 450 pm into the steel, the joint interface temperature was calculated as 911 K (638 OC). However, Fukumoto et al. presented no evidence confirming how the FeAI, FeAI, and Fe2AI intemetallics actually formed in completed joints and considered the forging time, not the friction time, during friction welding as the period during which interdiffusion occurred (Fukumoto et al. 1997). In effect, these investigators assurned that intemetallic formation occurred following completion of the steady-state period in friction joining when the fabricated component was cooling to room temperature. This assumption is not supported by the results presented in a nurnber of investigations indicating that intemetallic thickness and joint mechanical properties are essentially determined by the friction pressure and friction time during the joining operation (Jessop et al. 1978 and Fuji et al. 1992, 1995a, b). Assuming that al1 the deformation occurs in the dissimilar substrate, which has the lowest flow strength at high temperature, the contacting interface of the other (higher flow strength) substrate will be essentially stationary during the welding operation. When this condition occurs, intemetallic growth at the dissimilar joint interface can be visualised as occurring along a planar boundary and the width of the intenetallic layer will depend on the tirne available for interdiffusion and on the temperature at the bondline. This form of intemetallic growth explains the numerous references indicating that thin intemetallic layers and optimum joint mechanical properties are produced when the friction time is as short as possible and the average temperature at the joint 128 interface is lowered via selection of high friction pressure values (Jessop et al- 1978 and Fuji et al. 1992, 1995a, b). The results in the present study have confirmed that when a silver interlayer is introduced during MMCIAISI 304 stainless steel friction welding a wavy-shaped interface is produced and large mechanically-mixed IM and PD regions containing AgAI are formed adjacent to the joint interface in the MMC base material. In a joint produced using a friction pressure of 30 MPa, the maximum thickness of the IM and PD regions were 1 and 23 Pm, respectiveiy. In joints produced using a friction pressure of 240 MPa, the IM and PD regions were not observed at the half radius location. The influence of friction pressure on the width of IM and PD region can be cornpared with its effect on the dimensions of the intermetallic interlayer formed in dissimilar MMCIAISI 304 stainless steel friction welds. In dissimilar MMCIAISI 304 friction joint produced using a friction pressure of 30 MPa and a friction time of 4 S. the width of the intermetallic layer was 4.6 pm at the half-radius location and 2 pm when the friction pressure was 240 MPa (see figure 5.28). In a joint produced using a friction time of 0.2 s, the maximum thickness of the IM and PD regions were 45 pm and 120 Pm, respectively (see Figure 5.10). In joints produced using a friction time of 1.2 s, the width of the IM and the PD regions were 40 and 10 pm at the half radius location, see figure 5.10 and 5.14(8). The critical feature is that Ag3AI forms early in the dissimilar friction welding operation (it is part of the IM and PD regions). Since the width of the PD and IM regions decreases with friction time (see figures 5.9-5.11) the amount of Ag3AI retained in compieted welds decreases, not increases, when longer friction times are applied. 5.3.4. Transition layer in dissimilar MMC/AISI 304 stainless steel welds During friction welding, material is transferred from the stationary side of the joint to the rotating side during Stage I of the friction welding process. A transition layer (see figure 5.20-5.21) forrns because of the movement of the plasticised region. An annular area of transferred material is preferentially forrned near the half-radius location in the friction joint. This particular area is subjected to higher friction pressure and temperature resulting in the formation of a plasticised region (Duffin and Bahrani 1976 and Bendzsak et al. 1997). Wide interrnetallic layers are fonned at half-radius location in dissimilar MMC/AISI 304 stainless steel produced using a low friction pressure (30 MPa), see figure 5.22. However, in this case, growth of the intermetallic layer is not detenined by diffusion across a cornpletely stagnant region in the MMC base material since it has been confimed that dynamically quiescent regions fonn in MMCIAISI 304 stainless steel welds produced using a low friction pressure. This dynamically quiescent terminology is employed since fluid flow occurs in this region, although it is only of the order of a few micrometers per second (North et ai. 1997). The dynamicalfy quiescent regions are widest at half radius location where the peak temperatures occur during dissimilar friction welding. It is believed that the width of the quiescent region is directly related to the width of the intermetallic layer. 5.3.5. Intennetallic compounds and joint strength It has been suggested that the formation of brittle intermetallics at the joint interface has a markedly detrimental effect on the rnechanical properties of dissimilar friction welds. For exarnple, Fuji et al. indicated that the bend ductility of dissimilar TVAfSI 304 L stainless steel friction weld was detrimentally affected by the formation of intemetallics at the dissimilar joint interface. However, interrnetallic layer formation had iittle influence on tensile strength when it was measured using a standard unnotched 130 tensile testing design (Fuji et al. 1992). The bend ductility properties were improved through post-weld heat treatment (PWHT) at temperatures ranging from 500-600 OC. However, PWHT at temperatures exceeding 700 OC and the use of long holding times at that temperature had a markedly detrimental influence on both tensile strength and bend ductility properties. When the PWHT temperature equalled 900°C, FezTi formation was confirmed and was associated with the poor mechanical properties obtained during mechanical testing. In another study, Suzuki et ai. found a relationship between the tensile strength and the width of the intermetallic layer fomed at the dissimilar joint interface in pure Allpure Ti diffusion joints. Weld tensile strength fell dramatically when the intermetallic layer width reached 200 Fm. However, Suzuki et al. did not indicate the nature of the intermetallic compound formed at the bondline (Suzuki et al. 1994). Based on these results, it has been suggested that the detrimental influence of intemetallic compounds on weld properties depends on the rnechanical properties of the intennetallic phase (its yield strength, tensile strength, and ductility) and on the width of the intermetallic layer formed at the dissimilar joint interface. Although much of the published research has dealt with dissimilar joints that contain well-documented mechanical properties e.g. Fe2Ti, FeAI, and FefiI, it is worth noting that the influence of intermetallic layers on weld rnechanical properties will be affected by the rnechanical properties of the adjoining substrates (Fuji et al. 1992). in the present thesis FeAl and Fe4AIi1 formation were confirmed following TEM examination of dissimilar MMCfAISI 304 stainless steel welds (see figures 5.24-5.26). The FeAl intermetallic phase has an aluminium content in the range from 35 to 50 at%. lron aluminides are brittle at room temperature and the strength of the FeAl is dependent on aluminium content and temperature; for example, the yield strength 431 increases and ductility decreases when the aluminium content increases (Vedula 1994). The iron aluminide with the highest yield strength is the Fe3AI compound and has yield strength of about 780 MPa and an elongation of 8% at room temperature. The FeAl intermetallic has yield strength of 350 MPa and an elongation of 2% at room temperature (McKarney et al. 1991). In summary, the strengths of the intermetallic phases and adjoining substrates (MMC and AlSI 304 stainless steel) are similar. However, their ductilities are markedly different. Stainless steel AlSl 304 has an elongation of 64% while the MMC (10 vol% AI2O3) substrate has an elongation of 7% (at room temperature). Hence, if the total plastic strain at the bondline equals the failure strain of the intermetallic layer, the friction weld will fail along the dissirnilar joint interface. 5.3.6. Model for interlayer removal Silver coatings have been used as solid lubricants to reduce friction and Wear between sliding steel substrates (Iliuc 1989 and Holmberg and Matthews 1994). However, these coatings have limited life and operate in conditions involving the application of low normal sliding loads and substrate materials with low compatibility with silver. In the present thesis, dissimilar friction welding involves contact between a silver interlayer and an aluminium alloy based MMC. Since silver and aluminium form intermetallic compounds and solid solutions, Wear of the silver interlayer should be expected during the friction welding operation. Evidence presented in this thesis supports the contention that the silver interlayer is rernoved from the bondline during the friction welding operation. During Stage 1 of friction welding, a sequence of adhesionlseizure events takes place when the silver interlayer and the aluminium substrate contact each other. This 132 contact process can be explained using the wave mode1 to decribe the different Wear regimes when soft and hard materials corne into contact. In this rnodel, the critical parameters are the aspenty angle of the hard surface (a) and the nomalised film strength,f; see equation 5.1 where .r is the film shear strength and k the shear flow strength of the defoming material (Le. the silver interlayer). Assuming that the silver interiayer is oxide free, the value of the nomalised film strength V) is high. When the surface is well-lubricated or covered with an oxide film, fwill be low. In the schematic representation in figures 5.29 and 5.30, a rigid asperity angle (a) of 10 O, high and low values off are assumed and the shear strain (fi of the soft rnaterial is calculated using the upper bound method. For a high value off (O.SI), the resulting shear strain (j) is 20.7, see figure 5.29. For a low value off (0.1),the calculated shear strain is 1.29, see figure 5.30. These results indicate that clean surfaces undergo increased shear deformation, which may result in particle formation, and an intimate contact between substrates improving the final joint quaiity. Since low values off are associated to the presence of an oxide layer, the mechanical properties of dissimilar joints with oxide layers wiil be impaired as shown in figure 4.12. Figure 5.29. Schematic representation of the wave model when f = 0.90. The angle in the wave model is important in terms of explaining the formation of silver particles found in the PD region of dissimilar MMC/Ag/AISI 304 stainless steel friction welds. When the angle CD is zero, the wave in the soft material is removed resulting in the formation of a silver particle. It is worth noting that the wave model was developed to explain the behaviour of rigid ideally plastic materials. However, particle formation is also infiuenced by strain hardening and low-cycle fatigue properties of the interlayer material. In strain hardening materials, particle detachment may occur at high values of (20') and at low shear strains (y = 10.0) (Kopalinsky and Oxley 1995). In figure 5.29 for f = 0.9, the resulting shear strain is 20.7, as result a particle may form. The wave model can therefore be used to explain the presence of silver particles observed in the PD region of welded joints produced using low friction times. It follows that the conditions which favour minimal retention of AgAI in MMCIAgIAISI 304 stainless steel joints are the employment of high friction pressures Figure 5.30. Schematic representation of the wave model when f = 0.10. and long friction times (> 2 s). In direct contrast, long friction times favour the growth of intermetallic layers at the bondline of MMC/AISI 304 stainless steel welds produced without silver interlayers. When a silver interlayer is introduced, the intermetallic layer forms early in the welding process as result of mechanical rnixing and is progressively removed from the bondline region as the friction welding operation proceeds. In effect, the introduction of a silver interlayer radically alters the rnanner in which intermetallic iayers form during dissirnilar friction welding of MMC and AlSI 304 stainless steel base materials. The remarkable effect of silver depends on its ability to be worn away as the friction welding operation proceeds. The present study is the first detailed examination of the effects produced when a silver interlayer is introduced prior to friction welding. In partieular, the chernical 135 compounds formed at the joint interface have been characterized. lntermixed (IM) and particle dispersed (PD) regions are formed in Ag-containing dissimilar friction welds. The IM region contains a mixture of silver, aluminium and Ag3AI while the PD region contains a granular mixture of silver, aluminium, Ag3AI, and entrained Al& particles. These regions form very early in the joining operation and both contain Ag3AI. It follows that an interlayer (Ag) introduced with the specific aim of preventing Fe,AI, compound formation in MMWAISI 304 stainless steel friction welds promotes the formation of another intemetallic phase at the bondline. Since the IM and PD regions are progressively removed as the friction welding operation proceeds thinner intermetallic layers are produced when welding using long friction times. This type of behavior is quite different from that observed in silver-free dissimilar MMC/AISI 304 stainless steel welds. Thin Fe,AI, intermetallic layers are formed when the friction time decreases in Ag-free welds. Therefore, the present thesis has found quite new results, which provide a rnuch better understanding of the factors that determine intermetallic layer formation and its removal at the bondline of dissirnilar friction welds. Nanoparticles of silver were formed in dissimilar MMC/Ag/AISI 304 stainless steel welds produced using low friction pressures. Nanoparticle formation in dissimilar friction welds has never been previously observed or investigated. CHAPTER 6. THERMAL ASPECTS AND SOFTENED ZONE FORMATION 6.1. INTRODUCTION Friction welding is a joining operation that alters the microstructure and mechanical properties in material close to the joint interface. In age-hardened materials such as aluminium alloys, friction welding produces a softened zone that has a marked influence on the joint strength. Previous studies have established that the width of the softened zone depends on the specific heat input and the friction time (Midling and Grong 1994a. b). In particular, narrow softened zones are produced when high specific heat inputs and short friction times are applied (Grong 1994). Heat generation during friction welding depends on the applied friction pressure, the rotational speed, and the coefficient of friction (Wang 1975 and Grong 1994). The presence of an interlayer such as silver will alter the surface conditions and consequently, the coefficient of friction, heat generation, and heat partition into the adjoining substrates. It is shown in Chapter 7 that changes in the softened zone rnarkedly affect the strength of completed dissimilar welds. The influence of mechanised fusion welding (GMAW) on the dimensions of the softened zone in age-hardenenable aluminium alloy base material was examined and modelled for 6082 aluminium alloy base material (Shercliff and Ashby 1990a, b and Mhyr and Grong 1991a, b). Midling and Grong examined the problems of softened 137 zone formation during similar friction welding of 6082 aluminium alloy and AI-Sic composite base materials (Midling and Grong 1994a, b). The calculated width and hardness of the softened zone correlated well with the measured dimensions and hardness in both friction and GMA welds. However, Midling and Grong investigated friction welding of similar base materials and Iimited their research tu friction pressures ranging from 30 to 50 MPa. The present thesis examines the influence of friction pressures ranging from 30 tu 240 MPa on the softened zone during dissimilar friction welding when a silver interlayer is applied. During sliding friction, heat generation at the contacting surfaces depends on the coefficient of friction, the normal pressure, and the sliding speed. In similar manner, in friction welding heat generation at the bondline depends on the coefficient of friction between contacting substrates, the friction pressure, the rotational speed, and the rod diameter. Calculation of the heat generated at the bondiine during friction welding involves specific problems such as the selection of the coefficient of friction, the pressure distribution at the contacting interface when a rod is employed during the welding operation. In spite of these difficulties, an equation has been proposed indicating the heat generated per area unit during friction welding, see equation 6.1, (Grong 1994). where q, = net power, W, A = cross section, mm2, P = friction pressure, MPa, ,u = coefficient of friction Frm = maximum velocity at the outer periphery of the test sample, m/s. In modelling the softened zone in similar friction welding of aluminium alloys, Grong applied a coefficient of friction of 0.5 and assumed a uniform friction pressure at the contact interface. The calculated heat input served as the basis for predicting the dimensions of the softened zone formed in the aluminium alloy friction welds. However, the problem is more cornplex in dissimilar friction welding since the joints involve aluminium alloy and stainless steel substrates, which have different thermophysical properties, see table 6-1. Table 6.1. Thermal properties of base materials ------ Tm= Melting Temperature Since aluminium has a higher thermal conductivity it is expected that the aluminium substrate will conduct more heat. A simple analysis involving two materials in static contact proposes a heat partition ratio depending on the thermal conductivities of the contacting substrates (Suh 1986). The partition ratio is calculated using equatior: 6.2: mi?-- m m a- I 350- 300- m- m- 150 x I 1 I I I I O 60 la) 180 240 300 360 4ZC FRiCllON PRESSURE, MPa Figure 6.1. Effect of friction pressure on the temperature at 0.2 mm from the bondline and 4.5 mm from the periphery in dissimilar MMCIAglAlSl 304 stainless steel friction wetds. Friction tirne, 3 s; forging pressure, 240 MPa; forging time, 1 S. where kr and k2 are the thermal conductivities of the contacting substrate materials. The thermal conductivity of the aluminium alloy 6061-T6 is 167 W/m-K while stainless steel has a thermal conductivity of 16.2 Wh-K at 20 OC. Hence, using equation 6.2, it would be expected that 91% of the heat generated at the bondline would transfer into the aluminium component and 9% into the AlSI 304 stainless steel substrate. Although the presence of an interlayer rnay alter the heat partition ratio at the bondline, no published information is available indicating exactly how an interlayer will affect this parameter and the softened zone. In the following sections the influence of friction pressure and the silver interlayer on the softened zone are discussed in detail. FRICTION PRESSURE, MPa Figure 6.2. Effect of friction pressure on the temperature at 0.2 mm from the bondline and 4.5 mm from the periphery in dissimilar MMCIAISI 304 stainless steel friction welds. Friction time, 2 s; forging pressure, 240 MPa; forging time, 1 S. 6.2. RESULTS 6.2.1. Effectof friction pressure In the present study the temperature in the stainless steel substrate was measured at 200 pm from the joint interface and at a depth of 4.5 mm from the component periphery. In dissimilar MMCIAglAISI 304 stainless steel friction weld produced using a friction pressure of 30 MPa and friction time of 3 s, figure 6.1 shows that the temperature attains values as high as 435 OC. When the friction pressure increased from 120 to 240 MPa the measured temperature decreased in about 30 OC. In dissimilar MMC/AISI 304 stainless steel welds produced without an interlayer the temperature was 340 OC when using a friction pressure of 30 MPa and friction time FRICTION PRESSURE, MPa Figure 6.3. Effect of friction pressure on the temperature at 0.2 mm from the bondline and 4.5 mm from the periphery in dissimilar MMCIAgIAISI 304 and MMCIAISI 304 stainless steel friction welds considering a friction time of 1 S. of 2 S. The temperature decreased when higher friction pressure were applied, see figure 6.2. The decrease in temperature is in the same range observed in dissimilar MMCIAgIAISI 304 stainless steel welds. In both types of joints, friction pressure had a niarked influence on the temperature values. 6.2.2. Influence of the silver interlayer Considering a friction time of 1 s, the introduction of a silver interlayer decreases the measured temperature in about 50 OC, see figure 6.3. The temperature values in figure 6.3 are obtained from the temperature-time cutves for MMC/Ag/AISI 304 stainless steel welds and MMCIAISI 304 stainless steel welds produced using different friction pressure values (30 MPa, 120 MPa and 240 MPa), see figures 6.4-6.5. Figure 6.4. Effect of friction pressure on the temperature-tirne curves at 0.2 mm frorn the joint interface and 4.5 mm from the joint periphery in dissimilar MMClAglAlSl 304 stainless steel friction joints. Friction time, 3 s; forging pressure, 240 MPa; forging tirne, 1 S. Figure 6.6 shows the relationship between the axial shortening of the MMC substrate and the friction pressure applied during welding. In dissimilar MMC/AglAISI 304 stainless steel joints, the friction welding process produced less axial shortening than welds produced without silver interlayers. Also, dissimilar MMCIAglAISI 304 stainless steel friction joints have narrower softened zones than dissimilar MMCIAISI 304 stainless welds (see figures 6.7-6.1 1). For example, in MMCIAglAISI 304 stainless steel welds the sofiened zone width varied from 10.02 mm (in joints produced using a friction pressure of 30 MPa) to 3.40 mm (in welds produced using a friction pressure of 240 MPa). The softened zone width ranged from 17.25 to 4.80 mm in dissimilar MMC/AISI 304 stainless steel friction welds Figure 6.5. Effect of friction pressure on the temperature-time curve at 0.2 mm from the joint interface and 4.5 mm from the joint periphery in dissimilar MMClAlSl 304 stainless steel friction joints. Friction time, 2 s; forging pressure, 240 MPa; forging time, 1 S. produced using friction pressures ranging from 30 to 240 MPa. The introduction of a silver interlayer, therefore, decreased the width of the softened zone adjacent to the bondline of dissimilar MMCIAISI 304 stainless steel friction welds. 6.3. DISCUSSION 6.3.1. Heat entering the sfainless steel substrate Heat input is an important quantity when examining friction welding because together with friction time detemines the softened zone width and the temperatures in material at and close to the joint interface. The heat input entenng the stainless steel I 30 60 90 120 150 180 210 240 FRICTION PRESSURE, MPa Figure 6.6. Relationship between axial shortening and friction pressure in dissimilar MMClAlSl 304 and MMClAglAlSl304 stainless steel friction joints. Friction time, 4 s; forging time, 1 s; forging pressure, 240 MPa. component during the heating period is obtained using the temperature-time curves in figures 6.4-6.5 and equation 6.3. where % = net power, W, pc = volume heat capacity, ~/rnrn~OC a = thermal diffusivity, mm2/s, z = distance from the contact section, mm, e$c (U) = complementary error function. it is worth noting that equation 6.3 is valid for the heating period of the friction welding process. Table 6.2 shows the calculated heat inputs entering the stainless steel substrate for dissimilar MMCIAlSl 304 and MMC/Ag/AISI 304 stainless steel friction welds produced using friction pressures of 30, 120, and 240 MPa. 6.3.2. Heat partition during dissimilar friction welding An assumption during similar friction welding is that heat generated at the bondline is divided equally between the contacting substrates. Equation 6.2 was developed assuming static conditions when two substrates contact each other and the heat originating at the bondline partitions depending on the thermal conductivities of the adjoining components. Considering the properties of stainless steel and 6061 alloy substrates (see table 6.1) the static partition ratio suggests that 91% of the generated heat input will transfer into the aluminium alloy MMC substrate. Table 6.2. Relationship between specific heat input entering the stainless steel substrate @/A) and friction pressure Friction Pressure MMCfAISI 304 JOINTS MMC/Ag/AISI 304 JOINTS (MW w/mm2 w/mm2 30 MPa 3.25 2.98 120 MPa 3.5 3.34 240 MPa 3.72 3.59 146 Singh and Alpas measured a coefficient of friction of 0.5 during sliding Wear testing of MMC and stainless steel substrates (Singh and Alpas 1996). Midling and Grong applied a frictional coefficient of 0.577 and 0-5when analysing heat generation during simifar MMC friction welding. In the present thesis, a coefficient of friction of 0.54 is applied to dissimilar MMC/AISI 304 stainless steel friction welds. In dissimilar MMC/AISI 304 stainless steel produced with a friction pressure of 30 MPa, the specific heat input is 16.1 w/mm2. Assuming a static heat partition ratio, 91% of the generated heat enters the MMC substrate (14.49 wlmm2). As result, with a heat input of 14.49 w/mm2, the calculated temperature in the MMC would be 809.65 OC,see equation 6.3. However, this temperature is much higher than the temperature of melting of aluminium alloys and considerably higher than the temperature assumed by Midling and Grong when examining heat transfer during similar friction welding of MMC materials (555 OC) (Midling and Grong 1994a). Also, using equation 6.3 and data presented in figure 6.5, the calculated heat input entering the stainless steel substrates is 3.25 w/mmz, and this heat input is greater than the 1.3 w/mm2 value predicted using equation 6.2. It is therefore apparent that the static heat partition ratio (equation 6.2)cannot be applied to dissimilar friction welding. 6.3.3. Friction pressure and temperature The influence of friction pressure on the temperature produced during friction welding is discussed. The Zener-Hollomon constitutive equation for MMC base material is shown below (Davies et al. 1992): where Z = Zener-Hollomon parameter, s= strain rate, s-A Q = activation energy, J/mol, R = universal gas constant, Jlmol K T = absolute temperature, K a=material flow stress, Pa A, a and n are material constants. For 6061 base material containing 10 vol% of particulate material, Q = 216 kJ1mol; a =O.OZ3 MP~";n = 5.24 and A = 9.42~1oi4 s-l. To understand the influence of friction pressure on the temperature values let us consider the situation during forging of a 9.5 mm radius cylinder made of MMC base material. A 5 mm upset is considered here since this value equals the amount of burn-off produced during dissimilar MMCfAISI 304 stainless steel friction welding (see figure 6.7). The forg ing pressure required to upset the cylinder is given by equation 6.5 (Mielnik 1991): where p = applied pressure, Pa m = sticking friction factor, DO= material flow stress, Pa R = cylinder radius, m h = upset height, m. 148 For an applied pressure (p) of 30 MPa (which is typical of friction welding operations), rn = 1, R = 9.5 mm, and h = 5 mm, equation 6.5 indicates that the required material flow stress (O,) to upset the cylinder is 17.3 MPa. With a flow stress of 17.3 MPa and a strain rate of 0.5 s-', using equation 6.4, the temperature of the MMC cylinder should be 579 OC. However, the material temperature is rnuch lower (266 OC) when the MMC cylinder is deformed using a higher pressure (240 MPa). In effect, deformation initiates at lower temperatures when the applied pressure is increased. It is suggested that the temperature attained in dissirnilar MMC/AISI 304 stainless steel friction welds relates to the temperature at which the MMC base material starts deforming plastically. This suggestion rnay explain the lower temperatures measured when higher friction pressures were applied during dissirnilar MMC/AISI 304 stainless steel and MMCfAgfAISI 304 stainless steel friction welding. 6.3.4. Influence of the silver intedayer The introduction of a silver interlayer decreases the softened zone width during dissirnilar friction welding. Softened zones are produced in age-strengthened aluminium base materials because of the thermal cycle in friction welding and wider softened zones are produced when higher heat inputs are applied. In the present thesis, the softened zone is the region extending from the bondline to the location where the temperature was 250 OC in the adjoining MMC base material (Grong 1994). In dissimilar MMC/Ag/AISI 304 stainless steel friction welds, the width of the softened zone was 10.02 mm in a joint made using a friction pressure of 30 MPa and a friction tirne of 4 s (see figure 6.7(A)). The amount of axial shortening in this weld was c0.5 mm, see figure 6.6. When this softened zone width (10.02 mm) and a temperature of 250 OC are substituted into equation 6.3 the calculated bondline temperature is 422 149 OC. This temperature is produced in the MMC substrate when a specific heat input of 7.6 w/rnm2 is applied during a friction time of 4 S. During MMC/Ag/AISI 304 friction joining the calculated heat input in the stainless steel substrate is 2.98 w/rnm2. Therefore, for a dissimilar MMC/Ag/AISI 304 stainless steel weld produced with a friction pressure of 30 MPa the specific heat input is 10.33 wlmm2. With this heat input, and using equation 6.1, the calculated frictional coefficient is 0.35. It follows that the introduction of a silver interlayer during friction welding of MMC and AlSI 304 stainless steel substrates reduces the arnount of generated heat during the dissimilar weld ing operation. In fact, the frictional coefficient decreases to 0.35 by the introduction of a silver interlayer. There is support for this contention in the published literature. For example, silver coatings have been used to decrease the frictional coefficient and the arnount of Wear between sliding rnetals (Iliuc 1980, Holmberg 1994). Since dissimilar MMC/Ag/AISI 304 stainless steel friction welding presents lower heat input entering the MMC side of the joint, lower axial shortening and narrower softened zone are produced in the MMC substrate, see figures 6.6-6.1 1. 6.3.5. Summation Introduction of a silver interlayer decreased the coefficient of friction at the contact region during dissimilar welding and decreased the heat generated during the welding process. This effect explained the decreased particle fracture tendency and the narrower softened zone regions in the MMC base rnaterial and the decreased axial shortening (burn-off) cornpared to dissimilar welds produced without interlayers. Frictiorr Prasoure = 30 MPa HAt = 10.02 mm aprica 290 Min. hardness = 73.7 HV at 1.- mm -10-88-20 2 4 6 81012 DISTANCE FROM INTERFACE, mm MMClAlSl3û4 Friction Pressure, 30 MPa HAZ = 17.25 mm Min. Hardness = 70.8 ai 2-14 mm DISTANCE FROM INERFACE. mm Figure 6.7. Vickers hardness distribution across the joint interface. Friction pressure, 30 MPa; friction time, 4 s; forging pressure, 240 MPa; forging time, 1 S. 246 - (4 MMCIAe(AlSI: 304 220- I FWon Preusure, 60 MPa mm .mm HAZ = 7.24 mm 200 - Irn Min. hudnars = 76.3 MI at 1.73 mm g 180- m- 160- CD W: Z 140- n 3 l2O: 100 - 80- NS1304 - STAINLESS STEEL MMC 60 alvl-~-~-I-I-I~I.~~ -10 -8 ô 4 -2 O 2 4 6 8 10 12 DISTANCE FROM INTERFACE, mm -- - - MMClAlSl304 FMon Pressure. 60 MPa W=11.10 mm Min. fiarcines = 81.8 HV at 254 mm DISTANCE FROM INTERFACE, mm Figure 6.8. Vickers hardness distribution across the joint interface. Friction pressure, 60 MPa; friction time, 4 s; forging pressure, 240 MPa; forging time, 1 S. irmrilCIAs/AlSI 304 Friction Pretssure, 120 MPa HAZ = 5.43 mm Min. Hardness: T72 at 0.94 mm -10844-20 2 4 6 81012 DISTANCE FROM INTERFACE, mm MMCIAISl3û4 Friction Pressure = 120 MPa HAL = 6.34 mm Mn-Hardness = 85.5 HV at 209 mm t.l'l.l.,.~.l.l.l.,.l.l -10 -8 -6 4 -2 O 2 4 6 8 10 12 DISTANCE FROM INTERFACE, mm Figure 6.9. Vickers hardness distribution across the joint interface. Friction pressure, 120 MPa; friction time, 4 s; forging pressure, 240 MPa; forging time, 1 S. 240- (4 MMClAs/AlSl304 I Friction Pragsum = 180 MPa 220 - W\Z = 3.81 mm a Min- hardness: 83.7 MI MO- ===m.. . at 0.89 mm P 1,: cn' *w 160- Z 140- E 2 rm- 100- I AlSI 304 80 - STAJNLESS STEEL MMC iml-l~l~1.1-I.l.I.1~ -10 4 ô 4 -2 O 2 4 6 8 ?O 12 DISTANCE FROM INTERFACE, mm MMCIAiSi 304 Friction Pressum = 180 MPa HAZ = 5.14 mm Mi. Hardm: 867 HV 200 at 241 mm 180 DISTANCE FROM INTERFACE, mm Figure 6.10. Vickers hardness distribution across the joint interface. Friction pressure, 180 MPa; friction time, 4 s; forging pressure, 240 MPa; forging time, 1 S. WAgJAlSI 304 FWon FVesm= 24û MPa HAZ = 3.48 mm Min. hardrmss: 84.2 MI 200 et 0.53 mm 1 1, AlSl3W STAJNLESS STEEL -10 -8 ô 4 -2 O 2 4 6 8 10 12 DISTANCE FROM INTERFACE, mm 304 I WAlSI Friction Ressum = 240 MPa m I HAZ = 4-80 mm Min. Hardness: 91.6 I =. at 1-76 mm DISTANCE FROM INTERFACE, mm Figure 6.1 1. Vickers hardness distribution across the joint interface. Friction pressure, 240 MPa; friction time, 4 s; forging pressure, 240 MPa; forging time, 1 S. CHAPTER 7. INFLUENCE OF THE INTERLAYER ON THE MECHANICAL PROPERTIES OF DlSSlMlLAR FRICTION JOINTS 7.1. INTRODUCTION It has been shown that the introduction of a silver interlayer decreases heat generation during friction welding (by lowering the frictional coefficient at the contact region during Stage 1), the width of the softened zone (Chapter 6), and the percentage of particle fracture in the MMC side of the joint (Chapter 4). However, in both dissimilar MMCIAISI 304 stainless steel and MMC/Ag/AISI 304 stainless steei friction welds, increasing friction pressure produces thinner intermetallic layers at the bondline and narrower softened zones. From the tensile testing results it is not possible to know if the strength improvement in dissimilar joints is produced by thinner intemetallic layers, by inhibiting brittle Fe-AI intermetallic compounds or by changes in the width and hardness of the softened zone. It will be shown in this chapter that the improved tensile strengths of dissimilar friction welds depend, not only on the prevention of intermetallic cornpound formation at the bondline, but mainly on the softened zone formed in completed joints. The influence of the softened zone is evaluated using finite element modelling, and the calculated and actual joint strengths are compared for both MMC/AISI 304 stainless steei and MMC/Ag/AISI 304 stainless steel friction welds. la 1 1 r I I 1 I O 50 100 150 200 2SO FRlCliON PRESSURE, MPa Figure 7.1. Influence of friction pressure on the notch tensile strength of MMClAlSl 304 and MMClAglAlSl 304 stainless steel friction joints. Friction time, 4s; forging pressure, 240 MPa; forging time, 1s. 7.2. RESULTS 7.2.1. Joint mechanical properties Figure 7.1 shows the relation between friction pressure and the notch tensile strength properties of dissimilar MMC/Ag/AISI 304 stainless steel and MMCIAISI 304 stainless steel joints. During these tests al1 other friction welding parameters (friction tirne, rotational speed, forging pressure, and forging time) were held constant. Higher notch tensile strengths were produced in MMC/Ag/AISI 304 stainless steel welds. Also, the effective plastic strain (see equation 7.1) measured on the fractured section of the tensile test specimen corresponding with the MMC substrate was higher in dissimilar joints produced using silver interlayers. 200 - 2bo - 3100 - i50 - 400 NOTCH TENSILE STRENGTH, MPa Figure 7.2. Relationship between the effective plastic strain and notch tensile strength of MMClAlSl 304 and MMCIAglAISI 304 stainless steel friction joints. Friction time, 4 s; forging pressure, 240 MPa; forging time, 1 S. where E, = the effective plastic strain, which is constant across the notch, do = the original diameter, d = the final diameter of the specimen at the neck. The effective plastic strain increased from 0.25 to 5.0% in dissimilar MMC/Ag/AISI 304 stainless steel friction welds, see figure 7.2. These joints have the highest tensile strength. In dissimilar MMCfAISI 304 stainless steel friction welds, the effective plastic 158 strain values ranged from 0.25 to about 0.1 % when the friction pressure was increased from 30 to 240 MPa, see figure 7.2. The failure modes observed in MMCfAgfAISI 304 stainless steel friction welds were quite different from those failures in joints made without silver interlayer. In MMCfAgfAISI 304 stainless steel welds produced using a low friction pressure (30 MPa) failure occurred via a combination of brittle, interfacial, and ductile fracture, Le. brittle failure through regions containing Ag3AI, interfacial failure at the silver/aluminium interface, and ductile fracture through the MMC base material, see figure 7.3 (A-8). However, when the friction pressure was raised to 240 MPa the failure mode was wholly ductile through the MMC base material, see figure 7.3 (GD). In dissimilar MMCfAISI 304 stainless steel friction welds produced using a friction pressure of 30 MPa, failure resulted from a combination of ductile and brittle failure, see figure 7.4 (A-B). When the friction pressure increased from 30 MPa to 240 MPa brittle failure became the dominant mode of fracture (see figure 7.4(C-D)) and the area fraction of ductile failure decreased rnarkedly, see figure 7.5. The effect of friction the on the mechanical properties of dissimilar MMC/Ag/AISI 304 was also investigated. Figure 7.6 shows the fracture surface of a joint produced using a friction time of 2 s and a friction pressure of 240 MPa. Joint failure occurred via a combination of brittle, interfacial and ductile fracture, Le. brittle failure through regions containing Ag3AI and ductile fracture through the MMC base material. The resulting fracture surface was similar to the fracture of MMCIAglAISI 304 produced with a friction pressure of 30 MPa and 4 S. Figure 7.3. Fracture surace morphology in MMClAglAlSl 304 stainless steel friction joints: A), 6) friction pressure, 30 MPa; C, D) friction pressure, 240 MPa. Friction time, 4 s; forging pressure, 240 MPa; forging time, i s; rotational speed, 1500 rprn. Figure 7.4. Fracture surface morphology in MMCIAlSl 304 stainless steel friction joints: A), B) friction pressure, 30 MPa; C, D) friction pressure, 240 MPa. Friction time, 4 s; forging pressure, 240 MPa; forging time, 1 o; rotational speed, 1500 rpm. O ', I 1 I 1 1 O 50 100 150 200 250 FRlCTiON PRESSURE, MPa Figure 7.5. Influence of friction pressure on the percentage fraction of ductile failure. Friction time, 4 s; forging pressure, 240 MPa; forging tirne, 1 s; rotational speed, 1500 rpm. When friction pressure increases from 30 to 240 MPa, the width of the softened zone decreased in both the MMC/Ag/AISI 304 stainless steel friction welds (from 10.02 mm to 3.48 mm) and the MMCIAISI 304 stainless steel friction welds (from 17.25 mm to 4.8 mm), see figure 7.7. The relation between notch tensile strength and the softened zone width is shown in figure 7.8. The mechanical properties of the softened zone were monitored using microhardness testing. In particular, the mechanical properties of MMC base material immediately adjacent to tne bondline have a critical influence on the tensile strength of dissimilar joints. In the present thesis the hardness values at the location 0.125 mm from the bondline were taken as representative of the mechanical properties of material close to the bondline. Figure 7.9 shows the relation behnreen friction pressure and Figure 7.6. Fracture surface rnorphology in a MMCIAgIAISI 304 stainless steel friction joint Friction pressure, 240 MPa; friction tirne, 2 s; forging pressure, 240 MPa; forging time, 1 s; rotational speed, 1500 rpm. microhardness values at 0.125 mm from the bondline. Higher hardness values were produced in MMCIAgIAISI 304 stainless steel and MMCIAISI 304 stainless steel friction welds made using high friction pressures, see figure 7.1 0. Also, the lowest hardness values were observed in dissimilar MMC/Ag/AISI 304 stainless steel friction welds. 7.3. CALCULATING THE TENSILE STRENGTH OF DISSIMILAR FRICTION JOINTS 7.3.1. Introduction Alrnond et al. used finite etement analysis when they examined the tensile strength properties of brazed steellCu1steel joints. The stress and strain distribution in the copper interlayer were calculated using FEM analysis and satisfactorily predicted FRlCTlON PRESSURE, MPa Figure 7.7. influence of friction pressure on the softened zone width in MMCIAISI 304 and MMCIAgiAISI 304 stainless steel friction joints. Friction tirne, 4 s; forging pressure, 240 MPa; forging time, 1 s; rotational speed, 1500 rpm. the stresdstrain relations in welds containing different interlayer thicknesddiameter ratios (t/D) (Almond et al. 1983). Henshall et al. also applied finite element analysis when examining the properties of brazed dissimilar AlSI 304/Ag/AISI 304 stainless steel joints. These investigators suggested that finite element method could be used to calculate the final strength of dissimilar brazed joints provided that a satisfactory failure criterion was applied (Henshall et a/. 1990). However, these investigators did not suggest a specific failure criterion. Since the finite elernent method has been applied to examine the mechanical behaviour of dissimilar brazed joints. FEM is applied in the present thesis when Figure 7.8. Relation between softened zone width and the notch tensile strength of dissimilar MMClAISI 304 and MMCIAglAlSl 304 stainless steel friction joints. Friction time, 4 s; forging pressure, 240 MPa; forging tirne, 1 s; rotational speed, 1500 rpm. analysing the influence of softened zones on the mechanical properties of dissirnilar MMC/AISI 304 stainless steel and MMC/Ag/AISI 304 stainless steel friction joints. 7.3.2. The notched tensile specimen A "U" notched tensile test specirnen design was employed during the present research, see figure 3.4. When this tensile specimen test design is applied, the notch acts as stress concentrator that produces located plastic deformation in material close to the notch root. During the tensile test, an element located at the notch root is acted on by a triaxial state of stress consisting of a radial stress (a,),a tangential stress (a) Figure 7.9. Influence of friction pressure on the hardness of the adjoining MMC substrate in dissimilar MMCIAISI 304 stainless steel and MMCIAgIAISI 304 stainless steel friction joints. Measured at the location 0.125 mm from the bondfine. Friction time, 4 s; forging pressure, 240 MPa; forging time, 1 s; rotational speed, 1500 rpm. and an axial stress (a,),see figure 7.1 1. The triaxial state of stress must be considered during failure of the notch tensile specimen; for this reason a triaxiality factor is introduced, see equation 7.2: where o, is the average stress given by the relation, Figure 7.10. Relationship between the hardness of the adjoining MMC substrate and the notch tensile strength in dissimilar MMCiAlSl 304 stainless steel and MMCIAgIAISI 304 stainless steel friction joints. Measured at the location 0.125 mm from the bondline. Friction time, 4 s; forging pressure, 240 MPa; forging time, 1 s; rotational speed, 1500 rpm. where a,, a2, and are principal stresses. The equivalent stress (m) is calculated by the relation: The triaxiality factor is related to the failure strain. Metals yield when the equivalent stress reaches a critical value, the yield strength (5).However, when voids Figure 7.11. The necked region of a homogeneous notched tensile specimen showing the directions of the normal (cQ), radial (a,), and tangential stress (ae)- nucleate and grow, or when cleavage cracks nucleate and propagate, the volume of the test sample increases and the failure process depends on both OË and o, (Teirlinck et ai. 1988). This situation explains why the triaxiality factor must be taken into account when the failure process is investigated. The effective plastic strain to failure initiation (E,) depends on the triaxiality factor. For example, E, was 0.35 in notched HY 80 steel samples, which had a triaxiality factor of 0.75 (Hancock and McKenzie 1976). However, E, was 0.62 when the triaxiality factor was 0.6. In unnotched HY80 steel test specirnens, E,- was 1.1. In effect, when the triaxiality factor decreased the plastic strain to failure initiation (&,) increased. In unnotched testing specirnens the triaxiality factor attains a value of 113 at the periphery and a maximum value at the centre of the notch (Hancock and McKenzie 168 1976). The maximum value of the triaxiality factor is determined by the relation: where a = d/2 is the radius of the minimum cross-section of the notched specimen, R = the profile radius of the notched specimen as shown in figure 3.4. The triaxiality factor is required to calculate the failure stress during testing of notched and unnotched tensile specirnens. Although, equation 7.5 is applied to calculate the maximum triaxiality factor in homogeneous specirnens, this relation cannot be applied to dissimilar friction joints since the mechanical properties of the MMC and AISI 304 stainless steel substrates friction joints are markedly different. Also, the softened zone in MMC base material adjacent to the bondline makes the situation even more complex. For these reasons the mechanical situation during failure of notched tensile specimens extracted from dissimilar friction welds was investigated using finite element modelling. 7.3.3. Finite element mode1 of the notched tensile specimen Commercial software (ANSYS 5.5m, university version) was used when modelling the mechanical situation in notch tensile specimens extracted from dissimilar welds. All calculations were performed using triangular Snode elements. The geometry and meshing of the finite element rnodel are shown in figures 7.12 and 7.13. Because of the axy-syrnrnetrical shape of the tensile test specirnen, a two dimensional analysis of one half of the specimen was possible. The left edge of the rnodei (r = 0) corresponds with the centre axis of the notch tensile specimen. The mesh contained 2257 elernents and 4694 nodes. Zone Units: mm Figure 7.1 2. Half of the notch tensile specimen. At r = 0, the nodes are allowed to move only in the axial direction and at z = -15 mm, the bottom edge is fixed (the axial displacement, Uz = O). Simulated tensile loads were applied on the upper edge of the mode1 (at z = +15 mm), with the applied stress continuously increasing with time from O to 400 MPa. Elastic behaviour was represented using linear isotropic elasticity theory, while the time-independent plastic deformation was represented using isotropic Von Mises plasticity theory. Both substrates (MMC and AlSI 304 stainless steel) were assumed to behave as bilinear kinematic hardening materials. The yield strength (5)and tensile strength (a~)of the softened zone in the MMC substrate were calculated using equations 7.6 and 7.7; these mathematical expressions are rnodified versions of the regression equations developed by Myhr and Grong (Myhr and Grong 1994a). These equations relate microhardness to yield and BON wc/ns/nrsI Figure 7.13. Finite eiement idealisation of the notch tensile specimen. tensile strength values and agree with measured tensile strength properties of the MMC material, see table 3.3 and figure 3.6. Throughout this thesis, the unnotched failure stress au^) of the MMC material was assumed to be equal to the tensile strength of the metal matrix composite material calculated using equation 7.7: The hardness values across the softened zone were transformed into yield strength values using equation 7.6. For example, a location having a microhardness of 88 HVN has a yield strength of 205.5 MPa and a tensile strength of 262.9 MPa. Since it is assumed that the MMC and the stainless steel behave as bilinear kinematic 171 materials, the value of the tangential modulus in the softened zone (ET) equals the tangential modulus of the MMC base material prior to the friction welding operation. This assumption is based on the knowiedge that the softened zones in aluminium alloys are produced as result of an over-ageing process (Shercliff and Ashby 1990a, b and Grong 1994). Also, Hval et al. examined the properties of simulated softened regions with different yield strength in aluminium alloy 6083 and confinned there was no change in the tangential modulus value (Hval et al. 1998). The tangential rnodulus (ET)was calculated using equation 7.8 and figures 3.6-3.7.The tangential moduli were 375 MPa for the MMC substrate and 2260 MPa for the stainless steel (for a true strain of 0.1). where 8, = equivalent plastic stress, q,= yield strength, E = Young's modulus, El= tangential modulus, Ê = equivalent plastic strain. 7.3.4. Stress and strain distribution in notched tensile specimens The finite element method is used to investigate the influence of softened zones during mechanical testing of dissimilar rnaterial notch tensile specimens. In this work effort, the influence of softened zone width on the equivalent stress (Von Mises stress), the total equivalent strain, and the triaxiality factor are investigated. The distribution of key quantities such as the equivalent stress are calculated along lines parallel to the - - - -i- MMUASl304- Pl= 30 MPa 4- MMUASI 304. Pl=24O MPa \ Figure 7.14. Equivalent stress distribution along the line Iocated at 1 mm from the bondline for two MMCIAlSl 304 welds stainless steel friction joints. For an applied stress = 160 MPa. bondline and are expressed as function of the r/a ratio, where r is the radial distance measured from the centreline of the tensile test specimen to the point under consideration, and o is the specimen radius in the notch region, see figure 3.4. The equivalent stress distribution in dissimilar MMCIAISI 304 welds produced using friction pressures of 30 and 240 MPa is examined along the horizontal line located at 1 mm from the bondline. An applied stress of 160 MPa is considered during FEM modelling. Frorn O to 0.8 da, the equivalent stress is lower in a weld produced using a friction pressure of 240 MPa (softened zone width = 4.8 mm) than in the weld produced using a friction pressure of 30 MPa (softened zone width = 17.25 mm), see figure 7.14. In a dissimilar MMC/AISI 304 stainless steel friction weld produced using a Figure 7.15. Total equivalent strain distribution along the line located at t mm from the bondline for two MMCIAISI 304 stainless steel friction joints. For an applied stress = 160 MPa. friction pressure of 30 MPa, the equivalent stress is 191.32 MPa (at 0.85 da); this calculated value exceeds the yield strength at this particular location (180.6 MPa, calculated using equation 7.6). As result. the softened zone is plastically deforrned at this location, see figure 7.15. In dissimilar MMCIAISI 304 stainless steel welds produced using a friction pressure of 240 MPa, applied stresses exceeding 160 MPa are required to attain the material yield strength at this particular location. The triaxiality factor reaches a maximum value at the joint centreline in dissirnilar MMC/AISI 304 stainless steel welds produced using a friction pressure of 240 MPa. 1 n contrast, the peak value of the triaxiality factor is reached at the 0.81 r/a location in a weld produced using a friction pressure of 30 MPa, see figure 7.16. As result, the failure mode in joints produced using low and high friction pressures will be quite 1 I I I '1.1 O~O ' 0.2 0.4 O. 6 O-8 1.0 1.; Figure 7.16. Triaxiality factor distribution along a line located at 1 mm from the bondline for two MMCIAISI 304 stainless steel friction joints. For an applied stress = 160 MPa. different, with joints made using low friction pressures exhibiting ductile failure between the centreiine and the periphery of tensile test specimen. This feature is illustrated in figures 7.4 (A) and 7.4 (C). In a dissimilar MMClAlSl 304 stainless steel friction weld produced using a friction pressure of 240 MPa, finite element modelling indicates that the equivaient stress acting on the line located at 0.5 mm from the bondline shows a peak value close to the specimen periphery. However, when the equivalent stress is calculated at 1.O and 1.5 mm from the bondline, the peak stress moves towards the centreline of the tensile test specimen and away from the bondline, and the equivalent stress bemmes uniform, see figure 7.17. The calculated total equivalent strain follows the same tendency, see figure 7.18. The triaxiality factor has higher values in regions close to the bondline, Figure 7.17. Equivalent stress distribution along Iines parallel fo the bondline in a dissimilar MMClAlSl 304 stainless steel friction joint. Pl= 240 MPa. Softened zone width = 4.80 mm. For an applied stress = 200 MPa. decreasing its magnitude in regions away frorn the bondline, see figure 7.19. Simiiar changes in location have been previously observed in dissimilar maraging steel/Ag/maraging steel joints (Kassner et al. 1992); the peak stresses moved from the location at 0.8 r/a to the specimen centreline. 7.3.5. Calculating the notch tensile strength of dissimi'lar friction welds Although FEM modelling allows determination of stress and strain distribution in notch tensile specirnens, it does not indicate at which applied stress the specimen wili fail. To calculate the notch tensile strength it is necessary to introduce a failure criterion. This failure criterion must consider both the material properties and the test specimen geometry. Figure 7.18. Total equivalent strain distribution along lines parallel to the bondline in a dissimilar MMCIAISI 304 stainless steel friction joint. Pl= 240 MPa and softened zone width = 4.80 mm. For an applied stress = 200 MPa. The model proposed by Teirlinck et al. is applied to calculate the strength of notched tensile specimens (Teirlinck et al. 1988). In this procedure, it is assumed that failure occurs as result of void formation and void coalescence in the MMC substrate. One important feature of Teirlinck's rnodel is that is stress-based. Teirlinck et al. proposed this model to predict the stress at which ductile fracture occurred in unnotched and notched tensile test samples (Teirlinck et al. 1988). Applying this model to dissirnilar friction welding involves the following steps: 1. Calculating the volume change parameter D. 2. Calculating the faiiure stress. Figure 7.19. Triaxiality factor distribution along lines parallel to the bondline in a dissimilar MMCIAlSl 304 stainless steel friction joint. Pl= 240 MPa and softened zone width = 4.80 mm. For an applied pressure = 200 MPa. In unnotched and notched samples the volume change parameter deterrnined by the relation: 3 a", D = 0.56sinh - -- 2 O-" In equation 7.9 the triaxiality factor (O,&) can be obtained either analytically (using equation 7.5) or numerically. Since MMC/Ag/AISI 304 stainless steel and MMCIAlSl 304 stainless steel welded welds involve dissimilar materials and softened zones the use of an analytical solution (see equation 7.5) greatly oversimplifies the problem. The limitations of the analytical approach were overcorne by using finite 178 element analysis to calculate the tnaxiality factor. The failure stress was calculated using equation 7.1 0: where ONF = failure stress in notched tensile specimens, MPa, cm = failure stress in unnotched tensile specimens, MPa, m = strain hardening index. For clarity, the procedure used to calculate the notch tensile strength of dissimilar MMCfAgIAISI 304 stainless steel friction welds produced using a friction pressure of 240 MPa will be explained in detail. In this particular joint, the softened zone (3.48 mm) was divided into eight rectangular areas, which extended from the centre to the periphery of the tensile specimen, see figure 7.12. These rectangular areas had a height of 0.5 mm. The hardness values and the calculated yield strengths of the softened zone are indicated in the following table: - - 8 1 Area 1 1 2 3 I 4 5 6 7 1 - HVN 94.0 91.8 86.2 88.9 100.0 107.5 112.4 117.5 GY 223.5 217.1 200.0 208.4 241.4 263.9 278.6 294.0 These yield strength values are used as material properties in the finite element program. The finite elernent model is subjected to a tensile load applied on the upper edge of the model. The applied load increases from O to 400 MPa in steps of 4 MPa. The effective stress and the tnaxiality factor are calculated along a line parallel to the bondline located at 0.125 mm from the bondline. This location is selected since ductile 179 failure occurred close to the bondline in MMC base material in dissimilar MMC/Ag/AISI 304 stainless steel and MMC/AISI stainless steel friction welds produced using a low friction pressure (30 MPa). This fracture behaviour suggested that the MMC properties of the region close to the bondline had a major influence on the joint strength properties. Figures 7.20-7.25 show the equivalent stress, hydrostatic stress, and triaxiality factor distributions corresponding to the calculated failure condition during mechanical testing of MMC/AISI 304 stainless steel and MMC/Ag/AISI 304 stainless steel weld specimens made using friction pressures of 30, 120, and 240 MPa. In dissimilar MMClAgIAISl 304 stainless steel weld produced with a friction pressure of 240 MPa, the failure conditions are attained when a load of 260 MPa is applied in the upper edge of the finite element model. From figure 7.25, at r/a =0.7,the triaxiality factor is 0.91, substituting this value into equation 7.9, the volume change parameter, D = 1.043. Using equation 7.7 and a hardness value of 94 HVN, the unnotched failure stress (wF)is 279.7 MPa. Substitution of the above values in equation 7.1 0 (D = 1.043, vup 279.7 MPa, and rn = 0.075) produces a notched failure stress of ahF= 270.01 MPa. This failure stress is reached in the area close to the bondline (94 HVN) when the applied load is 260 MPa producing a notch tensile strength of 374.4 MPa (1.44X the applied load of 260 MPa). The calculated strength values are close to the measured values in both dissimilar MMCIAglAISI 304 stainless steel and MMC/AISI 304 stainless steel friction welds, see figures 7.26-7.27. 7.4. DISCUSSION 7.4.1. lntemetallic compound formation and joint fensile strength It has been suggested that the mechanical properties of dissimilar friction welds 180 markedly decrease when the width of the intermetallic Iayer exceeds a critical value (Elliot and Wallach 1981b, Fuji et al. 1992, Fuji et al. 1995a, b). Also, it has been suggested that the critical intermetallic layer thickness depends on the mechanical properties of the neighbouring substrates (Fuji et al. 1995b). Critical intemetallic layer widths ranging from 0.2 to 1.O pm have been suggested. However, much higher critical thickness values have been found when examining dissimilar diffusion bonded joints. For example, in AüTi joints produced without an interlayer, the notch tensile strength reached a peak value for an intemetallic layer width of 200 prn and then decreased substantially (Suzuki et al. 1994). This result suggests that the critical intermetallic width is 200 pm in TifAl diffusion welds. In contrast, the joint strength of aluminium alloy/steel diffusion welds produced using a silver interlayer decreased frorn 238 to 80 MPa, when the intermetallic layer increased from 0.5 to 16 pm (Calderon et al. 1985). (A, Figure 7.20. Stress and triaxiality factor distributions along a line located at 0.125 mm from the bondline. Applied stress = 204 MPa. Notch tensile strength = 293.8 MPa. Figure 7.21. Stress and triaxiality factor distributions along a line located at 0.125 mm from the bondline. Applied stress = 232 MPa. Notch tensile strength = 334.08 MPa. Figure 7. 22. Stress and triaxiality factor distributions along a line located at 0.125 mm from the bondline. Applied stress = 240 MPa. Notch tensile strength = 345.6 MPa. az vZ E Il) Figure 7.23. Stress and triaxiality factor distributions aiong a Iine Iocated at 0.925 mm frorn the bondiine. Applied stress = 152 MPa. Notch tensile strength = 218.9 MPa. Figure 7.24. Stress and triaxiality factor distributions along a fine located at 0.125 mm from the bondline. Applied stress = 226 MPa. Calculated notch tensile strength = 325.44 MPa. in. Figure 7.25. Stress and triaxiaiity factor distributions aiong a line located at 0.125 mm from the bondline. Applied stress = 260 MPa. Calculated notch tensile strength = 374.4 MPa. Figure 7.26. Calculated and measured notch tensile strength values in dissimilar MMClAglAlSl 304 stainless steel friction joints. Friction time, 4 s; forging pressure, 240 MPa; forging timke, 1 s; rotational speed, 1500 rpm. llI'l'i=l 1 150 180 210 24 270 FiüCilON PRESSURE, MPa Figure 7.27. Calculated and measured notch tensile strength values in dissimilar MMClAlSl 304 stainless steel friction joints. Friction time, 4 s; forging pressure, 240 MPa; forging tirne, 1 s; rotational speed, 1500 rpm. 189 In dissimilar MMCIAglAISI 304 stainless steel welds, the width of the IM and PD regions decreased from >20 pm to 1 pm and from >190 pm to zero when the friction pressure increased frorn 30 to 240 MPa, see figures 5.13-5.14. Since ASSAI is contained in both the IM and PD regions (see Chapter 5) it would be expected that the mechanical properties would be improved in welds produced using high friction pressures. Figures 7.3 (A-B) and 7.6 (8)confinn that failure occurs in friction welds containing thick IM and PD regions. Also, since the widths of the IM and PD regions decrease when frktion time increases, it would be expected that joints produced using low friction times would have low tensile strengths. In dissimilar MMC/Ag/AISI 304 stainless steel friction joints produced using a friction time of 2 s and friction pressures of 120 and 240 MPa had notch tensile strengths of 232 and 292 MPa, respectively. When the friction time increased to 4 s, the notch tensile strengths were 309.6 and 403.8 MPa when using similar friction pressures. Therefore, the general recornmendation indicating that improved dissirnilar joint mechanical properties are produced when the friction time is substantially decreased is not applicable to MMCIAglAISI 304 stainless steel friction welds. FeAl (see figure 5.24) and Fe&13 (see figure 5.26) formation were confirmed at the bondline of dissimilar MMC/AISI 304 stainless steel friction welds. The width of the intemetallic layer markedly decreased when high friction pressures values were applied during dissimilar MMCIAISI 304 stainless steel (see figure 5.28). At first glance, one rnight conclude that the improved dissimilar joint strength was associated with the formation of thin intemetallic iayers at the bondline. Similar arguments have been put foward (Calderon et al. 1985 and Suzuki et al. 1994) to explain the poor tensile strengths of aluminium alloy 6061/Ag/AISI 304 stainless steel and Alm diffusion welds, respectively. This proposal also conforms with the critical intemetallic layer width 190 proposal put forward by a number of investigators (Jessop et al. 1978 and EIliot and Wallach l98la, b). However, friction welding not only produces intermetallic layers but also creates softened zone regions in MMClAISI 304 stainless steel and MMCfAgIAISI 304 stainless steel friction welds. As pointed out earlier, the softened zone width decreases when the friction pressure increases. Therefore, the influence of softened zone formation on joint tensile strength properties cannot be ignored. 7.4.2. Softened zone and joint tensile strength The width of the soffened zone region formed irnrnediately adjacent to the bondline decreases when the friction pressure increases, see figure 7.7. This decrease occurs in both MMCIAglAISI 304 stainless steel and MMCIAISI 304 stainless steel friction welds. A sirnilar relation is found when examining aluminium alloy 606116061 and MMCIMMC friction joints and suggests that the strength improvernent in 606116061 welds produced using high friction pressure resulted frorn the formation of narrow softened zones adjacent to the bondline. For example, joints produced using a friction pressure of 300 MPa had a notch tensile strength similar to that of the as-received base material (Uenishi et al. 1998). It is therefore proposed that if an interlayer X was introduced during MMClstainless steel joining that could absolutely guarantee freedorn from intermetallic layer formation at the bondline, the mechanical properties of MMCNAISI 304 stainless steel friction joints would still be improved when welds were made using high friction pressure since narrower softened zones would be formed. The influence of narrower sofiened zones in promoting higher tensile strength properties in MMClAglAlSi 304 stainless steel and MMCIAISI 304 stainless steel welds (see figure 7.8) is important. Therefore, the mechanical properties of dissimilar welds will depend on the interplay of softened zone hardness (yield strength) (see figure 7.9) and intermetallic layer 191 formation at the bondline. When the friction pressure is increased from 180 to 240 MPa the difference in hardness values in the softened zone region close to the bondline is 4 HVN. This value corresponds to a calculated yield strength difference of 12 MPa (using equation 7.6) while the measured joint strength increment is 11-1 MPa. Consequently, the difference in joint strengths of these welds can be ascribed to an increment in Aardness in the softened zone not to the reduction in softened zone width. The hardness difference between dissimilar MMC/Ag/AISI 304 stainless steel welds produced using friction pressures of 120 and 180 MPa is 3 HVN, which results in a calculated yield strength increment of 12 MPa (using equation 7.6).However, the measured strength difference in actual joints was approximately 60 MPa. The wide difference between actual and calculated joint strength values may be associated with a decrease in softened zone width from 5.43 to 3.81 mm in these welds. Dissimilar MMC/AISI 304 stainless steel produced with high friction pressures had wide softened zones. For example, a dissimilar MMC/AISI 304 stainless steel weld produced using a friction pressure of 30 MPa has a softened zone width of 17.4 mm, this value is greater than the gauge length of the notched tensile specimen (15.9 mm), see figure 3.4. As result, the effect of the softened zone width on the mechanical behaviour of MMCIAISI 304 stainless steel friction welds produced with low friction pressure was not observed. In these welds, the actual joint strength depends on the mechanical properties of the MMC region close to the bondline. In dissimilar MMCfAISI 304 stainless steel welds produced using high friction pressures (120-240 MPa), higher joint strength results from the combined effects of decreasing width and increasing hardness in the softened zone. In dissimilar MMCIAISI 304 stainless steel friction welds produced using high friction pressures, the difference 192 between the calculated and measured joint strength was approximately 10 MPa, see figure 7.27, which shows relatively good agreement. 7.4.3. Failure mode in dissimilar friction joints There were striking differences in the failure modes observed in MMCIAglAISI 304 stainless steel and MMCIAISI 304 stainless steel friction weids produced using high friction pressures. Ductile failure was the dominant mode of fracture in MMCfAgfAISI 304 stainless steel joints produced using friction pressures ranging from 120 to 240 MPa, see figure 7.5. In contrast, the area fraction of ductile failure markedly decreased in MMCIAlSl 304 stainless steel joints produced using high friction pressures and the dominant failure mode became brittle, see figure 7.4 (C-D). Ductile failure is associated with extensive plastic deformation, while brittle failure involves little or no plastic deformation. As pointed out earlier, narrow softened zones inhibit plastic deformation during tensile testing, see figure 7.1 5. Therefore, the different failure modes can be associated with changes in softened zone. Based on this argument, it would be expected that friction welded joints containing wide softened zones would exhibit failure modes involving extensive ductile failure. Also, dissimilar welds containing narrow softened zones would be prone to failure involving little plastic deformation. However, these effects were not observed in dissimilar MMC/AglAISI 304 stainless steel friction welds. For example, joints containing wide softened zones exhibited brittle failure through the silver interlayer and the Ag3AI interrnetallic phase contained in the IM region, see figure 7.3 (A-B). Moreover, friction joints having narrow softened zones exhibited extensive ductiie failure. Also, higher effective plastic strains were measured in welds that contained narrow softened zones and these welds also had high notch tensile strengths, see figure 7.2. The failure modes observed in dissimilar wefds can be explained by taking into 193 consideration the influence of the intermetallic layer formation at the bondline regions. The mechanical properties of the intermetallic phases are determined especially by their aluminium content. Iron-aluminium intemetallics have elongation values in the range of 2-3% and yield strengths from 300 to 400 MPa. When the aluminium content is higher than 50 wt% Al, iron aluminides are considered as unsatisfactory for structural applications. For exarnple, the Fe4AIl3 intermetallic phase has an aluminium content of 61 wt% and as resuit the elongation should be less than 2% (Vedula 1994). Assuming that a brittle Fe,AI, compound formed at the bondline during dissimilar friction welding has a yield strength of 300 MPa and elongation of 2-3%, joint failure may occur when the strain during tensile testing is in the range from 0.02 to 0.03. The plastic strain produced at the bondline will depend on the width of the softened zone and on the local hardness (yield strength) of MMC base material in the softened zone. Figure 7.15 indicates that for the same applied tensile stress, dissimilar welds containing wide softened zones will have larger total equivalent strains. Since this behaviour occurs, dissimilar welds containing intermetallic layers will be prone to failure at this location. In effect, a combination of wide softened zones and intermetallic layer formation will facilitate preferential bondline failure. This result is the situation that occurs in MMC/Ag/AISI 304 stainless steel welds made using a low friction pressure (30 MPa). Also, these dissirnilar welds will fail at low applied toads and the joints will have low notch tensile strength values, see figure 7.1. Narrow softened zones are formed in dissirnilar welds produced using high friction pressures (240 MPa) and therefore higher apptied loads are required during tensile testing to cause joint failure. The intemetallic layer located at the bondline region still plays an important role since failure is also promoted at this location. In effect, the mode of the joint failure during tensile testing wili be markedly affected by intermetallic layer formation at the dissimilar 194 joint interface. Extensive ductile failure was observed in dissimilar MMC/Ag/AISI 304 stainless steel and MMCIAISI 304 stainless steel friction welds produced using high friction pressures, see figures 7.3 (C)-(D) and 7.4 (C)-(D). This behaviour is a quite different situation frorn that observed in tests samples taken from joints produced using low friction pressures, see figures 7.3 (A)-(B) and 7.4 (A)-@). In welds produced using high friction pressure, ductile failure occurred through the MMC of the softened zone close to the tensile specimen centreline in a cup-cone morphology. In dissimilar welds produced with low friction pressure ductile failure was observed close to the bondline and to the specimen periphery. Although no prior research has been published concerning the failure behaviour of dissimilar friction welds, it has been confimed that ductile failure may initiate close to the sample periphery or at the centreline during notched tensile testing of ductile materials. This behaviour depends on the geometry and dimensions of the notch and material properties (Hancock and McKenzie 1976, Sun et al. 1989). 7.4.4. FEM analysis of notched tensile test specimens In dissimilar MMCIAISI 304 stainless steel friction welds, the calculated equivalent stress values using FEM analysis indicated the presence of peaks along lines parallel to the joint interface, see figure 7.17. Similar behaviour was previously observed during FEM analysis of dissimilar maraging steeVAglmaraging steel brazed joints (Henshall et al. 1990). In these joints the peak location and its magnitude depended on the ratio of the silver interlayer thickness (t) and the diameter of the tensile specimen (D). (t/D) ratios equal to or lower than 1110 resulted in a uniform distribution of equivalent stress values. A peak equivalent stress was found at 0.7 r/a from the samplo centreline (where a is the radius of the tensile test specimen). Henshall et al. suggested that these 195 stress peaks were the result of inhomogeneous strain. In dissimilar MMCIAglAISI 304 stainless steel friction welds, the peak formation may result from stress and strain concentrations promoted by the notch and the difference in mechanical properties between the substrate materials. 7.4.5. The notch tensile strength of the dissirnilar friction joints In the present thesis, calculated and experimentally measured notch tensile strengths are shown in figures 7.26 and 7.27. Using a ductile failure criterion, the calculated joint strength along the line located at 0.125 mm from the bondline is close to the actual value. The presence of ductile failure on the fractured surfaces of tensile test specimens from MMCIAglAISI 304 stainless steel friction welds confirms that the mechanical properties of the MMC material irnmediately adjacent to the bondline play an important role on the joint strength. In dissimilar MMCIAglAISI 304 stainless steel friction welds the actual notch tensile strength values were slightly lower than the calculated results, see figure 7.26. In dissimilar MMCIAglAISI 304 stainless steel friction joints produced with high friction pressures (120-240 MPa) the fracture surface morphology was ductile in nature, see figure 7.5. The absence of a thick intermetallic layer at the bondline in MMC/Ag/AISI 304 stainless steel welds promote failure in locations between the dissimilar joint interface and the region having the lowest hardness in the softened zone (see figure 7.3 C-D). The present result suggests once again that the joint strength depends on the softened zone adjacent to the dissimilar bondfine. The calculated notch tensile strengths at the location 0.125 mm from the bondline in MMCIAISI 304 stainless steel friction welds correspond well with the actual values during mechanical testing of dissimilar welds produced using low friction 196 pressures (30-60 MPa), see figure 7.27. The presence of ductile failure confirms the validity of the procedure applied to calculate the joint strength, which is consistent with the ductile fracture observed close to the bondline and through the MMC material, see figure 7.4. The results in both dissimilar MMC/Ag/AISI 304 and MMCIAISI 304 siainless steel friction welds suggest that the joint strength of dissimilar joints produced using low friction pressures (30-120 MPa) depends on the mechanical properties of the material close to the bondline. In dissimilar joints produced with high friction pressure, the joint strength depends on the width of the softened zone and the properties of the region close to the bondline. 7.4.6. Summation When silver interlayers are introduced this arrangement decreases the amount of heat generation during welding, results in narrower softened zone regions and improves notch tensile strength properties. In al1 research to-date, it has been assumed perse that joint mechanical properties wholly depend on the mechanical properties and width of the intemetallic layer formed at the dissimilar joint interface. However, it is shown in the present thesis that the mechanical properties of MMCfAISI 304 stainless steel joints are determined by the combined effects of intermetallic formation at the bondline and softened zone formation in MMC base material immediately adjacent to the joint interface. Finally, for the first time, this thesis provides a methodology for calculating the notch tensile strength properties of dissimilar friction welds. The approach is based on a combination of FEM with a ductile failure criterion and shows an excellent correspondence between actual and calculated joint strength results. CHAPTER 8. CONCLUSIONS The research work in this thesis examined the influence of a silver interlayer on the microstructure and mechanical properties of dissimilar friction welds between aluminium-based MMC and AlSl 304 stainless steel base materials. The principal conclusions are as follows: 1. The effects produced when a silver interlayer was introduced were examined by considering its influence on the particle fracture tendency during dissimilar friction welding of aluminium-based MMC composite and AlSI 304 stainless steel base materials. Introduction of a silver interlayer decreased the particle fracture tendency (the percentage of fractured particle and the average particle radius) in material close to the bondline. The influence of silver was associated with a decrease in the coefficient of friction between the contacting substrates. 2. The calculated normal pressure required to fracture of AI2O3 particles at the contact region was less than 1 MPa. This calculated result is sirnilar to an experirnentally measured value of 1.06 MPa found during sliding Wear testing aluminium-based composite rnaterial. Because the lowest normal pressure applied during friction welding was 30 MPa, this condition means that particle fracture occurs very early in the friction welding operation (imrnediately following contact between the adjoining substrates). 198 3. The introduction of a silver interlayer during MMCIAISI 304 stainless steel friction welding promoted the formation of interrnixed (IM) and particle dispersed (PD) regions containing AgdI. The IM region contained a mixture of silver, AgaIl and entrained Alz03 particles, whereas the PO region contained a granular mixture of silver, AgAI. and aluminium. The widths of the IM and PD regions decrease markedly when the friction pressure increased during dissirnilar welding. 4. The intermetaliic Ag3AI fomed as result of mechanical mixing very early in the dissirnilar friction welding operation. The amount of Ag3AI retained in cumpleted MMC/Ag/AISI 304 stainless steel friction welds decreased when longer friction tirnes were applied. This situation is quite different from that occurring when the intermetallic layer is fonned at the bondline of dissirnilar MMC/AISI 304 stainless steel welds. It is suggested that the beneficial influence of friction pressure and friction time in minimising Ag3AI retention in MMCIAglAISI 304 stainless steel friction joints depends on wearing away of rnaterial at the joint interface. In contrast, in dissimilar MMC/AISI 304 stainless steel welds, growth of the intemetallic layer depends on interdiffusion between adjoining substrates. 5. In dissimilar MMCfAgIAISI 304 stainless steel friction welds produced using a low friction pressure (30 MPa), the presence of silver nanoparlicles was detected. It is suggested that nanocrystal formation and the formation of Ag3AI intemetallics of the PD and IM regions are the result of mechanical mixing during the welding operation. However, in joints produced with high friction pressure (240 MPa) and long friction time (4 s), the presence of nanoparticles was not observed. The absence of nanoparticles seems related to Wear of the silver interlayer. 6. In dissimilar MMC/AISI 304 stainless steel the formation of FeAl and Fe4AIl3 interrnetalfics was observed. In these welds, the application of high friction 199 pressures (240 MPa) promoted the formation of thin intermetallic layers at the dissimilar joint interface. In dissimilar MMC/AISI 304 stainless steel friction welds, this effect cannot be wholly attributed to a decrease in the temperature at the bondline since this temperature was only affected marginally by an increase in friction pressure from 30 to 240 MPa. 7. Introduction of a silver interlayer decreased the coefficient of friction at the contact region during dissimilar welding and decreased the heat generated during the weld ing process. This effect explained the decreased particle fracture tendency and the narrower softened zone regions in the MMC base rnaterial and the decreased axial shortening (burn-off) compared to dissimilar welds produced without interlayers. 8. The notch tensile strengths of both MMC/AglAISI 304 stainless steel and MMCIAISI 304 stainless steel friction welds increased when high friction pressures were applied during the joining operation. However, the highest notch tensile strength properties were obtained in MMC/Ag/AISI 304 stainless steel friction joints. The higher notch tensile strength properties of dissimilar MMCIAgIAISI and MMCIAISI 304 stainless steel friction welds depended on the formation of narrow softened zone dimensions in MMC material immediately adjacent to the bondline. The equivalent stress and total equivalent strain were decreased in dissimilar friction welds containing narrow softened zones. As a result, higher applied loads were required to produce joint failure during mechanical testing. 9. The influence of softened zone width and mechanical properties on the notch tensile strength was analysed using finite element modelling. This numerical method was applied in combination of a ductile faifure criterion to calculate the notch tensile strength of dissimilar joints. This procedure was applied in the region 200 close to the bondline in the softened zone; the intermetallic layer was not wnsidered. The rnodelling results suggest that the substrate properties and the softened zone width play an important role with regard to the weld strength properties. In MMC/AISI 304 stainless steel welds produced using low friction pressure (30-60 MPa), the calculated notch strength is similar to the actual strength test results. In dissimilar MMC/Ag/AISI 304 stainless steel friction welds produœd with friction pressures from 30-180 MPa the calculated notch strength is also in the range of the actual strength values. The calculated results suggest that properties of the substrate close to the bondline determine the joint strength in dissimilar welds with a wide softened zone. 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