METALLURGICAL AND MECHANICAL PROPERTIES OF
Ni-BASED SUPERALLOY FRICTION WELDS
by
Sujith Sathian
Department of Metdugy and Materials Science University of Toronto
Toronto, Canada
A thesis submitted in confonnity with the requirements for the degree of Master of Applied Science
Graduate Department of Metaiiurgy and MateriaIs Science University of Toronto
Toronto
Copyright by S. Sathian, 1999
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METALLURGICAL AND MECHANICAL PROPERTIES OF
NCBASED SUPERALLOY FRICTION W L D S
by
Sujith Sathian
Master of Applied Science
Graduate Department of Metallurgy and Materials Science
University of Toronto
1999
ABSTRACT
Friction welding is a solid state joining process and can be used to successfully join
similar and dissimila. Ni-based superalloy base materials. The objective of this project
involves the optimization of welciing parameters and post weld heat treatment procedures
dunng sirnilar Ni-base wrought alloy and dissimilar (Ni-base wrought alloy/Ni-base cast
alloy) friction welding. Extensive microstructural analysis was carried using a combination
of optical, scanning electron microscopy, transmission eIectron microscopy and atomic force
rnicroscopy. The application of a Ire-solution + stabilization + precipitation] post weld
treatment procedure produced the optimum balance of joint t ende strength and ductility
properties in Ni-base wrought alloy fiction joints. TEM analysis provided an insight
concerning the various mechanisms that occur in friction welded joints, particularly with
regard to changes in precipitate chemistry and in particle distributions.
ACKNOWLEDGMENTS
1 would like to express my sincere gratitude to my supervisor, Prof Tom North for
his persistent encouragement and inspiration during the course of this study.
I a h like to acknowledge Prof. G. Benâzsak and Prof' Z, Wang and Prof G.
Weatherly of McMaster University for their participation in useM discussions.
I would like to thank Pratt & Whitney Canada, Longueuil for fiinding this project.
My sincere thanks to Mt. Dave Thomas, Ms. Isabelle Bacon, Mr. Andy Weaver, Dr. Simon
Durham and Mr. Alain Bouthillier of Pratt & Whitney Canada for their usefùl suggestions
and technical discussions.
1 am grateful for the award of a University of Toronto Open Fellowship and a Pratt
& Whitney Canada Graduate Scholarship.
The technical support of Mr. Fred Neub and Mr. Sa1 Boccia and the administrative
efforts of the office st&s are deeply appreciated. I would like to thank the members of our
welding research group for their valuable discussions.
Finally, 1 would like to thank my wife, S h i , and my parents for their support and
encouragement.
TABLE OF CONTENTS
ABSTRACT
ACKNOWLEDGEMENTS
LIST OF FIGURES
LIST OF TABLES
CHAPTER 1: INTRODUCTION
CHAPTER 2: LITERATURE REVIEW
Metallurgy
Friction Welding
Problems Associated with Similar
Friction Welds
Problems Associated with Dissimilar
Friction Welds
CHAPTER 3: EXPERLMENTAL PROCEDURE
3.1 Ni-base Wrought AlloyMi-base
Wrought Alloy Friction Welds
xvii
Preliminary Tests to Generate the 34
Baseline Data
Opthkation of Friction Welding 36
Parameters
ûptimization of the Post Weld Heat 37
Treatment Procedure
3.2 Ni-base Cast Aiioy/Ni-base Wrought AUoy 38
Dissimilar Friction Welds
3.3 Experimental Testing 39
3.4 Transmission Electron Microscopy 41
3 -5 In-Situ Transmission Electron Microscopy 43
3.6 Atomic Force Microscopy 45
CaAPTER 4: RESULTS AND DISCUSSION OF 46 - 108
Ni-BASE WROUGHT ALLOY FRICTION W L D S
Generating the Baseline Data
4.1.1 Base Metal Microstructure
4.1.2 As- Welded Joint Microstructure
4.1.3 Combineci Effects of Friction
Pressure and Time
4.1.4 Rotational Speed and Weid quality
4.1.5 Forging Pressure and Weld Quality
4.1.6 Muence of Base Metai Condition
Optïmization of Welding Parameters
4.2.1 Weld Profiles
4.2.2 Hardness Profiles
4.2.3 Tensile S trength Properties
Optirnization of Post Weld Heat
Treatrnent Procedure
4.3.1 Microstructural Aspects
4.3.2 Hardness Profiles
4.3.3 Tensile Strength Properties
Tensile Strength of Ni-base Wrought Alloy
Base Material
CHAPTER 5: RESULTS AND DISCUSSION OF 109 - 121
Ni-BASE CAST ALLOY/Ni-BASE WROUGHT
ALLOY FRICTION WELDS
5.1 Ni-base Cast M o y Base Metal Microstructure 1 09
5.2 Welding Parameters and Post Weld 11 1
Heat Treatment
5.3 Hardness Profiles 117
CHAPTER 6: CONCLUSIONS 122- 123
FUTURE WORK 123
REFERENCES 124 - 128
LIST OF FIGUIWS
Figure 1
Figure 2
Figure 3
Figure 4
Figure 5
Figure 6
Figure 7
Figure 8
Figure 9
Figure 10
Figure t 1
Figure 12
Figure 13
Figure 14
Figure 15
Rupture strength behavior of three superalloy classes
(Fe-Ni, Ni- and Co-based superalloys)
Evolution of Ni-based superalloy microstructures
Microstructures of Ni-base cast ailoy and Ni-base
wrought alloy base materials
Schematic diagram indicating the relationship between the
percent volume hction of y' and the content of
hardeners (Al+TI) in superalloy base material
Schematic diagram showing the relationship between the
solution treatment temperature and the (AhTi)
content of the Ni-based superalloy
Pseudo binary diagram for Ni-based superalloys
Tirne-Temperature-Transformation diagram for Ni-base
wrought alloy base material
Primary and secondary y' particles in Ni-base cast
alloy superalloy base material
Basic steps during fiction welding of Ni-based superalloys
Direct - drive fiction welding machine
Different stages during direct dr ive fnction welding
Influence of fiction pressure and rotational speed on the
bum-off rate during welding
The relation between fnction pressure and forging
pressure on the notched tensile strengths of dissimilar
Al based MMC/ATSI 304 SS fiction welds
Reference chart indicating the weldability of different
material combinations
DBerent microstructural regions in friction welds:
Zpl is the Mly plasticised region,
vii
Figure 16
Figure 17
Figure 18
Figure 19
Figure 20
Figure 2 1
Figure 22
Figure 23
Figure 24
Figure 25
Figure 26
Figure 27
Figure 28
Figure 29
Figure 30
Zpd is the partly deformed region,
Zud is the undeformed region
Experimentally measured and calculated fluïd flow
region in aluminum aiioy ~ c t i o n welds
The relationship between the width of the flow region
and the fnction pressure applied during welding
operation ( m w s indicate the flow regions)
Predicted solute distribution within the HAZ in Al-Mg-Si
alloy fiction welds
Schematic representation of the HAZ hardness distribution
following fnction welding of Al-Mg-Si alloy base material
Spiral defect formation in aluminum alioy fiction welds
Particle agglomeration in yttcia containhg ODS alloy
fiîction welds
Intennetallic formed at the interface of a dissimilar
Al-based MMC/staidess steel fiction weId
Martensite formation in the stainless steel substrate in
a dissimilar Al-based MMC/stainless steel fiction weld
Materid transfer in a dissimilar Al based MMC/stainless
steel fiction weld
Microcrack formation in a dissimilar Al based
MMC/stainless steel fnction weld
Eutectic Formation at the bondline of a fiction joint
between Nickel and Magnesium
Sectionhg of Ni-base wrought dloy fiction welds prior
to metallography
Schematic showing the measurement of axial shortening
during fiction welding
Location of micro-hardness tests in Ni-base wrought
alloy fiction joints
Tensile sample design (ail dimensions are in impenal Units)
Figure 3 1
Figure 32
Figure 33
Figure 34
Figure 35
Figure 36
Figure 37
Figure 38
Figure 39
Figure 40
Figure 41
Figure 42
Figure 43
Figure 44
Figure 45
Figure 46
Figure 47
Figure 48
Extraction of TEM test sample following friction welding
Extraction of discs h m the foi1 using the EDM Process
Location of off-centered discs for TEM examination
Jet thinning of the discs
Heating chamber used during in-situ TEM microscopy
Working p ~ c i p l e in AFM Microscopy
SEM microstructure of solution-treated Ni-base
wrought d o y base material
Hïgher Magnification Mew of solution-treated Ni-base
wrought ailoy base material
Microstructure of hardened Ni-base wrought alloy
base material
TEM microstructure of solution-treated Ni-base
wrought alloy base material
XEDS pattern of the solution-treated Ni-base
wrought alloy y matrix
XEDS pattern of y' particles in the solution-treated
Ni-base wrought alloy base material
TEM microstructure of hardened Ni-base wrought
alloy base materid
XEDS Pattern fiom a grain boundary carbide in hardened
Ni-base wrought alloy base material
XEDS analysis of y matrix in hardened Ni-base
wrought alloy base material
XEDS pattern fiom a y' precipitate of the hardened
Ni-base wrought alloy base material
TEM micrograph fiom the dynamically recrystallised zone
(solution-treated Ni-base wrought alloy fiction weld
in the as-welded condition)
TEM micrograph showing the hcture of particles
in the partiy deformed region in a solution -treated
Figure 49
Figure 50
Figure 5 1
Figure 52
Figure 53
Figure 54
Figure 55
Figure 56
Figure 57
Ni-base wrought aiioy fnction weld
Particle segregation in the undeformed region of a
solution-treated Ni-base wrought ailoy fnction weld.
As-weldeà condition.
TEM rnicrograph showing the dynamically
recrystaïfised region in hardened Ni-base wrought
alloy fiction weld. As-welded condition.
Segregation of y' in undeformed region of hardened
Ni-base wrought alloy fnction weids.
As-welded condition.
TEM micrograph of a carbide contained in the y matrix of 56
hardeneà Ni-base wrought ailoy fiction welds.
AFM Image of Ni-base wrought alloy base materiai 59
(the arrow shows the grain boundaq)
AFM Image of the HAZ region in a fiction weld produced 60
using the conditions; Friction Pressure: 350 MPa
(tirne: 1 Os), Forging Pressure: 350 MPa (the: 2s),
Rotational Speed: 1000 rpm
Welded joint produced using a low fiction pressure
and fiction time (fnction pressure = 250 MPa,
fnction time = 2s). The other parameters were; forging
pressure = 250 MPa, forging t h e = 1 S.
Grain size variation in a traverse fkom the edge of
the partially recrystallised region into the as-received
Ni-base wrought alloy base matenal. These welds were
made using the conditions:
A. Friction Pressure: 250 MPa (time: 2s) + Forging Pressure:
250 MPa (time: ls), B. Friction Pressure: 300MPa (time: 6s)
t Forging Pressure: 300 MPa (time: ls), C. Fnction Pressure:
400 MPa (time: 14s) + Forging Pressure: 400 MPa (time: 2s).
Muence of rotational speed ( b m 1000 rpm to 64
2000 rpm) on grain size variation measured at the
component centerline. As-welded Ni-base wrought alloy
fiction joint.
Figure 58 Grain size changes on either side of the bondline
(in the stationary and rotating) components of the
fiction weld. As-welded Ni-base wrought alloy fiction joint.
Figure 59 Microhardness profiles in as-welded Ni-base 66
wrought aUoy/Ni-base wrought alloy joints in a traverse fkom
the bondline into the base matenal
(al1 measurernents at the centerline of the component).
The welding conditions comprised:
A. Forging Pressure: 200 MPa (tirne: 2s)
B. Forging Pressure: 350 MPa (time: 2s)
C. Forging Pressure: 500 MPa (time: 2s)
(Friction Pressure: 350 MPa, Rotational Speed: 1000 rpm)
Figure 60 Microhardness values before and &er post weld 68
heat - treatment in a traverse from the bondline into the
solution-treated Ni-base wrought alloy base metal.
The welding conditions compnsed:
Friction Pressure: 350 MPa (Friction Time: 10s)
Forging Pressure: 350 MPa (Forging Tirne: 2s)
Rotational Speed: 1000 rpm
Figure 61 Micro-hardness results before and &er post weld
heat- treatment in a traverse h m the bondline into the
hardened Ni-base wrought alloy base metal.
The welding conditions comprised:
Friction Pressure: 350 MPa (Friction Time: 10s)
Forging Pressure: 350 MPa (Forging Time: 2s)
Rotatiod Speed: 1000 rpm
Figure 62 Ni-base wrought alloymi-base wrought alloy welded
joint produced using a fiction pressure = 350 MPa and a
fiction tirne = 2s.
The other welding parameters were forging pressure
= 350 MPa, forging time = 2s and rotationai speed: 1000 rpm
Figure 63 Influence of fiction pressure on the amount of 73
axial shorteaing during fkiction welding of Ni-base
wrought atloy base material
Figure 64 Flash regions produced using different fiction pressures 73
(275,325 and 375 MPa) during Ni-base wrought
alloy/Ni-base wrought alloy fnction welding.
The other parameters were fnction time = 1 Os, forging
pressure = 350 MPa, forging tirne = 2s and
rotational speed = 1000 rpm.
Figure 65 Weld profiles produced using different fiction pressures 75
(275 and 325 MPa) during Ni-base wrought
aNoy/Ni-base wrought alloy fiction welding.
The other parameters were Friction Time = los, Forging
Pressure = 350 MPa, Forging Time = 2s and
Rotational Speed = 1000 rpm.
Figure 66 (a-c)Hardness profiles at the bondine in Ni-base
Figure 67
Figure 68
wrought alloyMi-base wrought alloy welds produced
using different fiction pressures
(275,300,325 and 375 M'Pa). The other welding parameters
were fiction tirne: 10s. forging pressure: 350 MPa,
forging tirne: 2s, rotational speed: 1000 rpm.
Typical stress -strain cuve produced during testing of 78
Ni-base wrought ailoy fiction joints.
The welding conditions comprised:
Friction Pressure: 375 MPa (Friction Time: los),
Forging Pressure: 350 MPa (Forging Time: 1 Os),
Rotational Speed: 1000 rpm
Tensile strength properties (yield strength and uitimate 79
tensile strength) of Ni-base wrougbt ailoy friction welds
Figure 69 Tende Ductility (total elongation) during tende testing of 79
Ni-base wrought alloy fiction joints
Figure 70a-b Fracture sufaces of broken tende sample fiom 81
Ni-base wrought alloy weld
Figures 7 1 a-b Magnined view of the above sample (see Figure 70a-b) 81
Figure 72
Figure 73
Figure 74
Figure 75
Figure 76
Figure 77
Figure 78
Figure 79
TEM micrograph showing massive precipitation in the 84
dynamicaily recrystailised region of a Ni-base
wrought alloy fiiction weld foilowing p s t weld
heat treatment ( P m
C-curve showing different precipitation mechanisrns in
Ni-based superalioys
XEDS pattern Erom the y matrix in post weld heat treated
PWHT] Ni-base wrought alloy friction weld
XEDS pattern fkom a y' precipitate (using
conventional TEM)
XEDS pattern fkom a y' particle (using STEM mode)
TEM microstructure following [RS + PWHT] treatment of a solution-treated Ni-base wrought
d o y friction weld
Variation in re-solution temperature with [Al+Ti] content 88
in the superalloy
Tirne-Temperature-Transition diagram for Ni-base 88
wrought alloy base material
Figures 80a-d Optical micrograph of the bondline region when an ST 90
heat treatment procedure is followed by PWHT
Figure 8 1 TEM micrograph f?om the dynarnically recrystallised 92
region (solution-treated Ni-base wrought alloy
base material prior to Wction welding)
Figure 82 TEM micrograph f?om the dynamically 92
recrystallised region (solution-treated Ni-base
Figure 83
Figure 84
Figure 85a-c
Figure 86
Figure 87
Figures 88
Figure 89
Figure 90
Figure 9 1
Figure 92
Figure 93a-f
wrought alloy base material pnor to fiction welding)
following in-situ heat treatment
Dynamically recrystallised region (hardened Ni-base 93
wrought d o y base material prior to Ection welding)
TEM micrograph fiom the dynamically recrystallised region 93
(hardened Ni-base wrought alioy base material prior to
fiction welding) foliowing in-situ heat-treatment
Hardness profiles produced using various heat treatment cycles 96
(a) After P WHT Treatment
(b) After RS + P m Treatment
(c) M e r ST + PWHT treatment
Schematic diagram showing the variation in hardness wi th 96
aging tirne and temperature
Schematic illustrating different precipitation effects 98
Typical stress-strain curve for w + P WHT] 98
Ni-base wrought alloy fiction welds
(solution-treated Ni-base wrought alloy base material
pnor to fiction welding)
Tensile strength properties (yield strength and ultimate 99
tensile strength) of solution-treated Ni-base wrought alloy
fiction welds given different heat treatments
Tensile Ductility (total elongation) of solution-treated 99
Ni-base wrought aLloy niction welds given different
heat treatments
Tensile strength properties (yield strength and ultimate
tensile strength) of hardened Ni-base wrought
alloy niction welds given different heat treatments
Ductiiity (total elongation) of hardened Ni-base
w u g h t dloy niction welds given different
heat treatments
Fracture surfaces of broken tensile test samples
extracteci h m heat treated Ni-base wmught alloyMi-base
Figure 94
Figure 95
Figure 96
Figure 97
Figure 98
Figure 99
Figure 100
Figure 101
Figure 102
wrought alloy welds
(a) Prior base metal condition: Solution treated Ni-base
wrought ailoy, Heat treatment: PWHT, (b) Prior base
metal condition: Hardened Ni-base wrought alloy,
Heat treatment: PWHT, (c) Prior base metal condition:
solution-treated Ni-base wrought doy ,
Heat treatment: RS + PWHT, (d) Prior base metal
condition: Hardened Ni-base wrought alloy,
Heat treatment: RS + PWHT, (e) Prior base metal condition:
solution-treated Ni-base wrought alloy,
Heat treatment: ST+PWHT, (f) Prior base metal
condition: Hardened Ni-base wrought alloy,
Heat - treatment: ST + PWHT Tensile strength properties weld strength and ultimate
tensile strength) of as-received Ni-base wrought
alloy base material given different heat treatments
Tensile strength properties weld strength and
ultimate tensile strength) of hardeneci Ni-base wrought
alloy base material given different heat treatments
SEM micrograph of Ni-base cast alloy base material
Magnified view of Figure 96
TEM micrograph of Ni-base cast alloy base material
XEDS pattern kom a cubical y' particle in Ni-base
cast ailoy base material
XEDS pattern h m a spherical y' particle in Ni-base
cast alloy base material
Joint interface in Ni-base cast alloy/Ni-base wrought
alloy fnction joint
As-Welded microstruchire of Ni-base cast
alloy/Ni-base wrought alloy dissimilar fnction weld
Figure 1 03 Crack formation foiiowing pst weld heat treatment 116
of a Ni-base cast ailoy /Ni-base wrought aiioy fiction weld
Figure 104 TEM micrograph of the bondhe region in a Ni-base IL8
cast aüoy/Ni-base wrought ailoy fiction weld
following RS (re-solution) + PWHT (stabilization
plus precipitation) heat treatment
Figure 105 XEDS spectra fkom a sphencal y' particle (hm
Ni-base cast alloy/Ni-base wrought alloy Ection welds)
Figure 106 XEDS pattern h m the y ma& (Ni-base cast
ailoy/Ni-base wrought d o y fiiction weld)
Figure 107 Hardness profiles in Ni-base wrought alloy / Ni-base
cast ailoy fiction welds
A. As- welded condition B. After Re-solution + PWHT
Table 1
Table 2
Table 3
Table 4
Table 5
Table 6
Table 7
Table 8
Ni-base superalloys and their composition
AUoying elements and their effects on superalloys
Typical heat treatment routes for Ni-base wrought aIloy
and Ni-base cast alloy bar materiai
Different types of carbides in Ni-based superalloys
Experhental testing rnatrix during preliminary
fiction welding trials
Experimental test matrix when examining the combined
effects of fiction pressure and time during Ni-base
wrought dloyMi-base wmught alloy friction welding
Experimentai test matrix when examining the
inauence of rotational speed on the Ni-base wrought
dioy Enction welds
Test matrix when exiunining the influence of
forging pressure during fiction welding of 19 mm
diameter Ni-base wrought alloy bars
CHAPTER 1
INTRODUCTION
Ni-based superalloys are dlo ys of Ni - Co - Cr containhg controlled additions of 1 0
or more trace elements. These materials are extensively used during the manufacture of
aircraft aero-engines. However, they are highly susceptible to cracking when conventional
fusion weldùig is used to fabricate them. In contrast, fiction welding is a solid state joining
process and can be used to successfiilly weld a range of Ni-based superalloy base materials.
The key joining parameters during fiction welding operation comprise friction
pressure and time, forging pressure and t h e and rotational speed. It bas been established
[North, 19991 that the fiction pressure is the key variable during the fiction joining
operation. The fiction time mut be sufncient to allow adequate fictional heaîing but not
so long that excess flash is produced during the welding operation. The influence of forging
pressure on nnal weld quality is unclear. However, some investigators have suggested that
longer forging time has a beneficial effect in reducing the residual stress in certain fiction
weld geometry's Pacon, 19991. For this reason, the influence of forging pressure and
joining time on weld quality constituted an important aspect of this program.
Rotational speed is determined by the type of direct drive fiction welding machine
used during fabrication. The torque increases with reduction in rotational speed. The torque
produced during the direct drive fiction joining process is also controlled by fiction
pressure since the range of rotational çpeeds that can be selected is relatively narrow- For
this reason, the present study focused on examining the effects of varying the rotational
speed in the first instance and then, having established an acceptable rotational speed value,
the influence of fiction pressure, fiction the, forging pressure and forging time were
inves tigated.
The fiction joining operation creates a My-plasticised region at the contact region
and heat affect4 zone regions on either side of the bondine. In duminum ailoy welds, for
example, the width of the hliy plasticised region is about 2 mm and the width of the heat
affected zone is about 4 mm on each side of the bondline (depending on the joining
parameters applied). The formation of HAZ regions is important when Ni-based superalloys
are fnction welded, since post weld heat treatment is mandatory. The required heat
treatment temperatures for Ni-base wrought alloy and Ni-base cast alloy base materials are
quite different. The corresponding heat - treatment temperatures are approx. 100-1 SOC
higher for Ni-base cast d o y compared to Ni-base wrought alîoys. When the solution-
treated Ni-base wrought alloy base material is Ection welded the joints are heat-treated
using a re-solution + stabihtion + precipitation treatment. Also in dissimila. Ni-base cast
alloy (available in the hardeneci condition)/Ni-base wrought alloy joints, the welds are heat-
treated using re-solution + PWHT involving a stabilization plus precipitation thermal cycle
corresponding with that for Ni-base wrought alloy base material.
The objective of this thesis is to examine the fnction welding characteristics of
similar Ni-base wrought ailoy/Ni-base wrought alloy joints and ushg this information
develop satisfactory parameter settings for dissimilar Ni-base wrought alloy/Ni-base cast
alloy fiction welding. For this reason the work effort was carrieci out in different stages,
narnel y,
O Investigating the propenies of Ni-base wrought alloymi-bare wrought alloy joints
0 Investigating the properties of Ni-base c m alloyNi-base wrought olloy joints
The factors influencing the mechanical and metallurgical properties of the cornpleted
welds were examineci in detail using a combination of optical, scanning electron
microscopy, fractography, microhardness and mechanical (tensile) testing. Transmission
Electron Microscopy was used to examine the changes in precipitate chemistry, particle
morphology and distribution resulting h m the fnction welding operation and fiorn
subsequent post weld heat-treatments.
CHAPTER 2
Ni-, Fe-,Co- based superaibys
Ni-, Fe- and Co-base s u p d l o y base materials are generally used at temperatures
above 550°C because of their excellent crap strength properties. Although some superalloy
material compositions can be forged and rolled into sheet, the more highly alioyed
chernistries are generaliy processeci as castings. Fabricated components are aiso made using
welding or brazing. However, many highly alloyed superalloy compositi011~ contain large
contents of hardening phases and are difncult to weld since they are highly susceptible to re-
heat cracking followiag post weld heat treatment [Duvall, 1966 and 19691. The mechanical
properties of s u p d o y base materials are controlled by composition and thermal
processing. For example Figure 1 compares the rupture strength properties of three alloy
classes (Fe-Ni, Ni- and Co-base superalloys) [MM, 19721.
Ni-based SuperaUoys
Ni-based superalloys are alloys of Ni, Co and Cr containhg controfled additions of
10 or more trace elements. The evolution of Nickel based superalloy is shown in Figure 2
[ASM Handbook, 19821 and a representative list of Nickel based superalloys and
compositions is presented in Table. 1.
Superalloy base materials derive their strengtb h m a combination of solid solution
hardenen and precipitating phases. Carbides may provide limited stmgthenuig directly or,
more commoniy, indirectly. In addition to those elements that produce solid solution
Figure 1 : Rupture strength behavior of three super alIoy classes (Fe-Ni, Ni- and Co-based superalloys) [MM, 19721.
Figure 2: Evolution of Ni-based superailoy microstructures [ASM Handbook, 19821 (Rm is the creep stress when samples are tested at 870C for 1 0 0 hours)
strengthening and promote carbide and gamma prime formation, other elements (B, Zr, Ce,
Hf) are added to enhance mechanical or chernical properties. Tabk 2 provides a List of
alloying elements and their effects in superailoys. Some elements produce readily
discemible changes in microstructure; while other elanents produce more subtle
microstructurai effects. The most obvious microstructurai effects involve precipitation of
geometrically close - packed phases such as y' and carbides.
Table 1 : Ni-base superdoys and their composition
Others
Elementd additions go into solid solution and provide one or more of the following
effects; improved strength (MoyT* W, Re) pathd, 19851, oxidation resistance (Cr,Al),
phase stability (in case of Ni) or increased volume fiactions (Vf) of favourable secondary
precipitates (Co) Ir\rathal, 19821. Other elements are added which fonn hardenuig
precipitates such as y' (Ni3(Al,Ti) and y" prime (Ni3Cb). y' phase is the key microstructural
feature responsible for the extraordinarily usefbl high temperature strength properties of Ni
and Fe-Ni base superalloys. Minor elernents (C&) are added to fonn carbides and bondes;
these elements plus others (CeJMg) are added for purposes of tramp element control. Some
elements (B, Zr, Hf) are also added to promote grain boundary effects Baldan, 19891.
Many elements (Co,Mo,W,Cr etc.), although added for their favourable alloying
capabilities, c m result, in some circumstances, in undesirable phase formation (sigma(a),
mu@), laves phases).
Table 2 : Alioying elements and their effects on superalloys.
- - E t r r a s n a m cn- Aluniinum, , , . , , , , , , , , , , . , . formr y* NWAî. TU; n t m r d m formrtion ork-gmmml .r
NtTi Titmnium .................. Fonnm r* NidAl. Ti) and MC urbidem Niobium. Tmntœlum, , , - , . , . . Formm body-œntœrœd tmtrrgonrl v" r d MC -bid- C m r b a n , . .................. Formm MC. Me. MaCa rnd &C arbidu; rtrbilitar fcc
matrcx Phorphorus. ............... Promotrr --ml pdpit i t ion of crrbidu Sitragen .................. Forma M t C m wbonitridar Chromium. ................ Oridition rriirtrrrœ; molid molutr'an mtrunehanine MolyW-num; Tua- - . . -. Solid .oIritioir atrœaqChrriis~ forum M& crrbid- Nickel . . -, . - , _ . _ . , _ , _ , _ _ - S t r b i l h C;cc mitrix; rom v*. i d i b i t inforinrtion o fdr leb
rio- p h n m Bor-=: Zi-nium, , . , . . , , , . . Impiovm erœœp propmrliœm; mtad formrtion of m i n bound-
Microstructure
The microstructure of superalloy base materiai is cornplex. It consists of an
austenitic face centred cubic m a t . (y) containhg a varïety of secondary phases. These
secondary phases are gamma prime ((Ni,Co),Al,Ti - intermetallic) and various carbides,
namely, MC &C, M G , and so on. Figure 3 shows optical microstructures of Ni-
base wrought ailoy and Ni-base cast ailoy base materials subjected to hardening treatments.
Cast superalloys generally have coarser grain &es, more aiioy segregation and
ïmproved creep and rupture strength characteristics Wmught superalloys are more uniform
and usuaily have 6aer grain sizes and improved tensie and fatigue strength properties.
Heat treatment routes
The optimum mimstnicture and mechanical properties are achieved using suitable
kat-treatment procedures. Table 3 details the typical heat treatment routes for Ni-base
wrought alloy and Ni-base cast aiioy base maîeriaIs peaver, 19981.
Table 3: Typical heat treatment routes for Ni-base wrought d o y and Ni-base cast alloy base material
Hardeners content
'Hardeners' are the total wt.% of (Ti+Al) in the superalloy. Figure 4 prager, 19681
shows the relationship between the percent volume fiaction of y' phase and the content of
hardeners. A linear relationship is observed. Other elements which are reported to be
present in y' are Nb, V, Hf, Ta, W and Mo.
Heat - Treatmcnt Solution - Tmtment
Stabilhtioh
Recipitation
Solution Treatment Temperatare
Solution treatment temperatures Vary from one alloy to another depending on the
AhTi contents (Le. the content of y') in the base material. Higher solutionizing
temperatures are required for ailoys containing increased Ai+Ti contents. The typical Ai+Ti
content in Ni-base wrought alioy is 4.5% and a solutionizing temperature of 1090°C for 4
hours completely dissolves the y' particles in the base material mi~l~~structure (see Figure 5)
peaver, 1 9981.
Ni-bars wrought .Uoy 109O0C for 4 ho- + Air Cool W C for 2 hours + Ai. Cool 76û°C for 16 hours + Air Cod
Ni-base cast alloy 1 190°C for 2 h o m + Air Cool
1
109o0C for 4 hours + Air Cool 870°C for 20 hours + 1
Air Cool
Figures 3 : Microstructure of Ni-base cast ailoy and Ni-base wrought ailoy base materials
Figure 4: Schematic diagram showing the relationship between the percent volume fiaction of the y' and the content of hardeners (Ai+Ti) in superalloy base material prager, 19681
Phase Diagram
A pseudo binary phase diagram for Ni-based superalloy is s h o w in Figure 6
[Schubert, 19801. The fouowing points should be noted:
The solution temperature varies k m one superalloy to another. For example,
alloy # 3 requires a higher solutionizing temperature thm alloy # 1. n ie
temperature difference is entirely dependent on the amount of y' prime in
superallo y.
Following the solution treatment, the alloys are heat treated to promote further
hardening. Ailoy #1 requires only one heat treatment cycle to precipitate al1 the
particles while alloys #2 and #3 require hvo different heat treatment cycles for
this to occur (see the figure 6). These subsequent heat treatments comprise a
stabilization treatment at 840°C for 2h and a precipitation treatment at 760°C for
16 h.
TTT Diagram
A typical TTT ( t h e - temperature - transformation) diagram for Ni-based superalloy
base material is shown in Figure 7 [Schubert, 19801 and is helpfid in understanding the
infiuence of t h e and temperature on the final microstructure. However, it must be noted
that for extended aging &es, diajgrams are only available for a few superalloy materials.
Stage 1 in this figure is the solution treatment. Stage II involves stabilization treatment.
Stage III is the precipitation treatment. The following effects occming during these stages:
Stage 1 : Complete dissolution of precipitates in the ma&.
Hot forming and working is carried out at this temperature and
high temperature carbides (MC) and gamma prime precipitate out during cooling to
room temperature.
S- 11: Coarsening of existing y' particles occurs [Baldan, 1989 and 19921.
Precipitation of low temperature carbides (e.g. S C & occurs dong the grain
bowidaries.
Figure 5 : Relation between the soiuuon ueatment temperature and the (-Ti) content of Ni-base wrought alloy peaver, 19981
o d u t i o n trmmorrtum for 7' :
- -------.- t---- 1140- t170°C
Seo -1010 OC
gure 6: Pseudo binary
sm Nucleation of secondary y' precipitation occurs within the existing y'
precipitates.
Precipitate growth is limited by low temperature khetics and the primary and secondary y'
particles in Ni-base cast ailoy (see Figure 8) [Antolovich, l982].
Carbides are important coastituents in superalloys. Carbides provide particle /
matrix strengthening- Different types of carbides are present Ui superalioys depending on
their composition and some of the important types are MC, M&, M,C, T C , (where M
stands for one or more metal atoms). Among the most cornmon are M& and MC
carbides. Table 4 lists various carbides found in different Ni-based superalloy base
materials.
Table 4: Different types of carbides in Ni-based superalloy
Allq cm-
ticooel718 MAR M-200
MAR M446 Nimmle 75 Rmb 41
Udimet 700-
The characteristics feahires of the different carbides comprise:
a) Carbide type - MC Tt is stable at elevated temperatures
Precipitation of MC is associated with the effective removal C from the matrix
and thus Little C is available for the forniaton of other carbides (e.g. K 3 C &
Time (s) Figure 7: Tirne - Temperature - TU-L~~U~UVU ~ ~ ~ q a m for Ni-base wrought alloy base material [Schubert, 1980)
Gamma Prime
>amma Prime
Figure 8: Primary and secondary Y' particles in Ni-base cast alloy base materid [Antolovich, 19821
Precipitation of MC carbides improves high temperature tende strength
properties.
Large amouuts of carbide precipitation result into poor weldability.
In Nb-containing superalloys (e.g. N i o n i c 80), formation of NbC during heat
treatment is sluggish compared to Tic. Addition of Hf in ailoy IN 792 replaces Ta in the
MC carbide particle paldan, 1989, Nathal, 19891. Further, Hf additions alter the
solidification sequence of the carbides in superalloys [Lecompte, 19881 during casting and
alter the carbide morphology and distribution during casring operatiom.
b) Carbide type - Mac
M6C (W, Mo&) has a cubic cubic lattice structure and forms, for example, in
Rene77 base material. M,C carbides are more stable than S C 6 carbides. However, at
elevated temperatures M,C carbides degeneration has been noted in Rene 41 and Mar-M-
252 base materials.
C) Carbide tjpe - Mg. M,C, forms at temperatures just below 1 050°C in Nimonic 80 base material. At
lower temperatures, M,C, carbides decompose to &C6 with a reverse transformation being
possible at temperatures exceeding 1060°C [Garosshen, 19851. Formation of M,C, has a
favourable infîuence on superalloy properties because this reduces the amount of &,C6
particles at grain boundaries and improves the creep Me of the material.
4 Carbide type - MJ3C6
%C6 has a cubic structure and forms preferentiaiiy at grain boundary regions.
Discrete carbides are the most favourable since they pin and limit grah boundary
movement. However, blocky and continuous carbides are highly deletenous [ASM
Handbook, 19721. Continuous carbide formation occurs in the temperature range of 870°C
to 980°C with the particles having an aspect ratio of 30 to 50 which favours crack nucleation
at the interface between the carbide and the matrix.
In H,- containing superalloys (e.g.DS 2Oûû He base material), Hf ties up C and
promotes the formation of a continuous network of carbides. The factors that determine
continuous carbide foxmation are the precipitation kinetics, the extent of prior precipitation,
the MC distribution and the yo' chemistry.
2.2. Friction Welding
Basic Mechanism
Friction welding is a solid-state joining process that produces coalescence due to the
heat developed between two surfaces when mechanicaiiy induced nibbuig motion occurs.
Basic steps in the fiction welding pmess are indicated in Figure 9. The welding process
can be characteriseci as a number of distinct stages. During Stage I, the substrates are
brought into contact at a low applied load (see Figure 9a) and the defornation process is
dominated by fictional Wear. In stage II, the applied load is substantially increased and
considerable fictional heating occurs at the bondline (see Figure 9b) and a balance of strain
hardening and softening processes is attained. During Stage III, fictional heat generation is
terminated, and the applied stress is substantiaüy increased to forge the heated material on
either side of the bondlule (see Figures 9c and 9d). An image of U of Toronto's direct-drive
fiiction welding machine is s h o w in Figure 10. The characteristics feahires of continuous
drive fiction welding are shown in Figure 11 [WRC Bulletin # 2041.
Weiding Variables
A number of parameters are important during the ltiction joining and comprise:
Friction pressure
Friction Time
Rotational speed
Forging pressure
Forging Time
Effects of Welding Parameters
Rotational speed produces the necessary relative velocity at the fayhg surfaces.
Lower rotation speeds produce higher torque values and necessarily involve senous work-
piece clamping problems and sometimes have been associatecl with non-uniform upsetting.
The equilibrium torque and the bum-off rates depend on the work-piece diameter, the
applied pressure and the speed of rotation. Equilibrium torque and bum-off rate are linearly
related with fiiction pressure and are inversely related with rotational speed (see Figure 12).
-1- CL.. .
Figure 9: Basic steps during fiction welding of Ni-based superalloys BulIetin # 204, 19821
Figure 10: Direct - drive fkiction welding machine
Figure 1 1 : Different stages during directdrive fiction weldhg Bulletin # 204, 19821
0.20 Sg.clman d m 0.75 In. db.
Pressure, psi
Figure 12: Influence of fiction pressure and rotational speed as the burn-off rate during welding pllis, 19921
[Ellis, 1992 and Hollander, 19631. A 1inear correlation exists between the equilibrium
torque and bum off rate over the range of Wction pressures comrnonly used, the slope being
dependent on the rotation speed selected.
The effects of fiction pressure and forging pressure on the notched tensile strengths
of dissimilar Al based MMCIAISI 304 Stainless Steel iiiction welds are illustrateci in Figure
13 Worth, 19951. It is apparent that the notched tensile strength increased when the fnction
pressure increased. However, it is worth noting that forging pressure had not a neghgible
effect on weld tensile strength,
Materials Welded
In p ~ c i p l e , fiction welding can be used to join alrnost any material combination. The
following matenal classes can be successfully readily joined;
Al1 types of steels
Aluminium and its alloys
Copper and its alloys
Titanium and its alloys
Nickel, cobalt and its aIloys
Dispersion strenghend alioys
However, base materials that contain weakening phases (e.g. graphite in cast irons
and manganese sulphide in steels) produce joints, which have poor strengths and
watisfactory microstnictures. A number of dissimilar metal combinations exhibit poor
weldability, e.g. aluminum alioydstainless steel and aluminum ailoylsteel, titanium
alloys/steel joints. hadequate joint strength properties result nom the formation of
intermetallic compounds at the bondlùie region. A reference weldability chart (indicates
various dissimilar metal combinations) is s h o w in Figure 14 W C Bulletin # 2041.
Figure 13 : The relation between fiiction pressure and forging pressure on the notched tensile strength of dissimilar Ai based MMC/AISI 304 SS Ection welds Worth, 19951
Figure 14: Reference chart indicating the weldabiiity of different matenal combinations [WRC Bulletin # 204, 19821
Microstructurai regions
The fiction welding opemtion markedly alters the microstructure of materiai on
either side of the bondine, i.e., in the heat-affecteci zone. The HAZ region can be divided
into a number of distinct regions, see Figure 15 wddling, 1994- 11. The three main regions
of specific interest are:
The fiilly plasticised region (Zpl) where the material is abIe to accommodate the
plastic strain by dynamic rezovery (or recrystallisation) of the microstructure.
The partly deformeci region (Zpd) where the plastic deformation is
accommodated by an increase in the dislocation density in the matrix grains. In
this region the temperature is sufnciently high to facilitate the dissolution of the
base metal hardening precipitates.
The undeformed region (Zud) that is characterized by partial reversion of the
base metai precipitates.
Modeling of the Welding Process
Bendzsak and North pendzsak, 19961 modeled the width of the flow region during
fiction welding operation by assuming that the materials in the contact region could be
treated as a highly viscous fluid. Assuming steady state conditions (see Figure 16) the width
of the flow regions is determind by the relation:
where p is the viscosity, r is the radius, a is the rotational velocity, p is the coefficient of
fiction, % is the outer radius of the test sample and Pa is the fiction pressure. Satisfactory
agreement was observed between the calculateci and the measured flow widths in Aluminum
alloy 606 1 friction joints. The width of the fully plasticised region in Al based MMC /
MMC friction welds was inversely pmportional to the fiction pressure applied during
welding (see Figures 17) Worth, 1998-11 -
Figure 15: Different microstructural regions in fnction welds wddling, 1994-11 Zpl is the M y plasticised region Zpd is the partly defomed region Zud is the undefonned region
-6CK) O 500 Axial Distance (pm)
Figure 16: Experimentally measured and calculated fluid flow regions in aluminum alloy fiction welds pendzsak, 19961
Figure 17: The relation between the width of the flow region and the Kction pressure applied during welding operation (arrows indicate the flow regious) Forth, 1998- 1 1. The welding conditions for welds A and B were as foiiows:
A. Friction Pressure : 60 MPa (tirne: 2s), Forging Pressure : 60 MPa (time: 2s) B. Friction Pressure : 180 MPa (time: 2s), Forging Pressure : 60 MPa (tirne: 2s) Rotational Speed : 2000 rpm
2.3. Problems Associated with Shilar Meta1 Friction Welds
a) Alterations in Precipitate Distniution and Morphology : Friction welding alters the
particle distribution wasiderably when age-strengthened base materials are fabricated.
Middlùig and Grong Wddhg, 1994 -11 developed a generic mode1 for the coupled
reversion for rod-shaped MhSi particla in Al-Mg-Si aluminium d o y s and applied this
during friction welding. Based on this approach, the volume hction (t) of precipitates
decreases fiom its initial value (fo), according to the relation:
where dt is the tirne increment at a temperature (T), t,' is the time required for complete
dissolution at the temperature concenieci, and a is the tirne exponent. The variation of t,'
with temperature is given by the relation :
where t,' is the time required for complete dissolution at the reference temperature Tr,., QS is
the metastable solws boundary enthalpy and Q, is the activation energy for diflkion of
particles in the undeformed matrix.
Using the above equations it is possible to calculate the solute distribution within the
of fiiction welds. The predicted particle dissolution cuves are shown in Figure 18.
b) Hardness Variations in the HAZ : Aiterations in the particle distribution, morphology and
microstnicture in the HAZ have a considerable influence on the mechanical properties of
completed welds (on the hardness profile, tende strength properties). The HA2 can be
divided into three different regions of interest according to the particle distribution and
general microstructure:
The fûlly plasticised region is where dynamic recovery of the microstructure
occurs. In this region the particles are cornpletely dissolveci.
The partiy deformed region whexe the plastic deformation is accommodateci by
an increase in dislocation density in the matrix grains and the temperature is
su££ïcient to facilitate dissolution of particles.
The undefomed region is characterized by the partial reversion of the base metai
precipitates.
Middiing and Gmng modeled the above regions when examining HAZ hardness
variations in Al-Mg-Si fiction welds wddling, 1994-23.
The net precipitation strength inmement, Amp , was caiculated using the relation :
Anp = k2f7f0 where (W) is the particle volume k t i o n and k, is the kinetic constant. This equation c m
be rewritten as
a = (~-o-/a--omin) = (Aq, /A%& = f7fo where cmin denotes the intrinsic matrix strength following complete particle dissolution and
a- is the original base metal strength.
The above equation provides a basis for obtaining information occurring reaction kinetics;
for example:
(i) Reversion Model: during particle dissolution the volume fiaction fell eom its initial
value, f,. The following equation delineated the dimensionless strength parameter within the
partly reverted region of the HAZ:
(ii) Naturd Aging Model: in fiction welding it is important to predict the final strength of
the HAZ following n a d aging. The precipitation strengthening increment was calculated
fiom the relation:
(iii) Work Hardening: dislocations are generated in the matruc matenal adjacent to bondline
as a result of straining. The strength contribution due to this factor a, was given by the
relation:
where y is a constant, which is a characteristic feature of the material selected. It proved
possible to calculate the HAZ strength distribution d e r friction welding by coupling the
above quations. Figure 19 shows schematic representaîion of the superimposed hardness
profiles produced using the above models. Partial dissolution of particles promp ted
softening in friction welds [MSu, 19911.
c) Spiral Defect Formation: A typical spiral defect obsenred on the fkacture surface of an Al
based MMC/A1 based MMC fiction joint is shown in Figure 20 Worth, 1998- 11. Spiral
defects are formed as a result of the formation of discontinuities in the flow of plasticised
materiai in the contact region. Spiral defects are fluid flow - induced defects formed when
material, magnesium nch segregates and oxide inclusion transfer to embedded regions close
to the stationary boundary of the joint. They remain trapped there for the reminder of the
welding cycle. The magnesium content of spiral defects ranged h m 15 wt % to 55 wt % in
different regions of spiral defects. Further, localized regions of a low melting point eutectic
material were observed and contained high magnesium contents. The spiral defects acted as
the sites for preferential failure sites during both fatigue testing and notch tensile testing of
MMC joints.
d) Particle Agglomeration and Particle Fragmentation: Particle agglomeration was observed
in MA 9561MA 956 fnction welds and was associated with the agglorneration of srna11
diameter (20 to 30 nm) yttria particles (see Figure 21). Oxidation early stage of the joining
operation also facilitated the formation of agglomerated particles. As expected, the creep
rupture properties of the fnction welds were detrimentaily affkcted by the particle
agglomeration process. Similar effects were also observed in FeA140 grade ODS
(containing yttria) alloy fiction welds [Inksoo, 19981. These particle agglomeration results
are quite different fkom those produced when Al based MMC material containhg much
larger (3 Fm to 45 diameter) reinforcing A$03 particles was fiction welded. In
MMC/MMC joints, the welding operation hgmented, not agglomenited, large 40,
particles Forth, 1996- 1,1997 and 1998-41.
1 2 3
Axiil distance . Z [mm]
Figure 18 Predicted solute distribution within the HAZ of Al-Mg-Si alloy friction welds Wddling, 1994- 11
Figure 19: Schematic representation of the HAZ hardness distribution welding of Al-Mg-Si alloy base matenal wyhr, 19911
Figure 20: Spiral defect formation in a l h u m alloy fiction welds FJorth, 1998-11
Figure 2 1 : Particle agglomeration in ytûia-containhg ODS alloy fiction welded joints NO*, 1996-11
2-4, Problems Associated with Dissimilar Metal Friction Welds
During the dissimilar metal welding process the microstructural changes are quite
different h m those in similar metal jcints. This is due to the diffient thermo-physical
properties (the melting points, thermal diftbivities, conductivity and diffusion coefficients)
and mechanical properties (elastic moâulus, yield strength) of the contacting substrates
[Sassani, 19981. Inadequate weld mechanical properties have been reported when dissimilar
substrates are fiction welded due to the formation of brittle intermetah phases at the
bondline- Further, brittle phases cm be formed in the HAZ, e.g. strain induced martensite in
dissimila. stainless steel /Ai based MMC friction welds. Since dissirnilar welding is an
important aspect in this thesis these topic will be discussed fùrther:
a) Intermetallic formation at the joint interface: Many metallic alloy combinations do not
produce satisfactory weld strength properties. This occurs since the joint strength is
detrimentaily affected by the foxmation of hard, brittle intermetallic phase at the bondline
region. During tensile loading, the intermetallic layer acts as a site for crack initiation and
propagation. For example, it has been reported that fnction joints between C u M contained
a thick intermetallic Iayer while Cu-70%WfAl welds contained thin intermetallic layers the
dissimilar joint interface. Similarly, FeAI and Fe2Al,, intermetaiiics have been observed in
fiiction welds between stainless steel and aluminium alloy base materials. Also, Fe,+, and
Fe$,, have been observed in fiction joints between Al and mild steel. The detrimental
influence of intermetallic layers on joint mechanical properties becomes apparent when a
critical intermetallic layer thickness is exceeded. A critical intermetailic layer thickness of 1
to 2 microns has been reported in AYC steel, Al / AIS1 3 16 stainless steel and in TUAISI 304
stainless steel fiction welds (see Figure 22) Forth, 1996-2 3.
b) Martemite Formation: When some dissimilar combinations are welded, local deformation
of the harder substrate occurs. For example, when welding Al based MMC to stainiess steel,
the differences in the flow strrngths of the harder and the softer substrates are considerable
so that almost al1 the deformation and plastic flow occurs in the softer substrate. However,
the harder substrate is still subjected to local deformation and TEM microscopy has
co-ed the formation of strain induced martemite in the stainless steel component of
MMC/AISI fiction welds (see Figure 23) [North, 1998-2 and 1998-31.
c) Matexial Transfer h m One Substrate to Another: The initial stage in the fiction welding
operation is characteriscd by a large number of adhesion/seizure/failure events. These
localiseci events transfer materiai h m one substrate to the other and vice-versa This
explains the high torque values generated at the start of the fiction welding operation. The
torque increases when in the flow stresses of the dissimilar combination increase. It has
been observed that generation of very high torque values can promotes hgmentation and
transfer of stainless steel material at the contact zone of stainless steeVMUC fnction welds
(sec Figure 24) Worth, 19951
d) Microcrack Formation at the Bondline: Microcrack fonnation commonly occurs in
dissimilar Al based MMC/AISI 304 stainless steel fiction welds when the fiction pressure,
fiction t h e and rotational speed are set low values (see Figure 25) Worth, 19951.
e) Trapping of Oxide Films: Trapping of oxides in the contact area is an inberent feature of
fiction welding. #en this occurs it can have a detrimental effect on the strength and
ductility of the cornpleted joints. This problem is severe when the oxide films are highl y
stable and very adherent at the contact surface. This problem is rnost apparent when
aluminium alloy substrates are fiiction welded. In Ti substrates however, the highly adherent
oxides dissolve in the tit;inium substrate at high temperature and their effect on weld tende
strength is negligible.
f ) Eutectic Formation at the Bondline: Eutectics are formed at the interface at low
temperatures (507OC for Ni and Mg) (see Figure 26) wazlett, 19621. Other alioy
combinations that are difficult to weld involve Zr to mild steel, Zr to stainless steel and Al
alloys to Cu.
Figure 22: Intermetaiiïc formed at the interface of a dissimilar Al- fiction weld Worth, 1996-21.
-bas4 MMC/stainless steel
Figure 23 : Martemite formation in the stades steel substrate of MMC/stainless steel friction weld Worth, 1998-21
a dissimilar Al-based
Figure 24 : Material transfer in a dissimilar Al based MMC/stahfess steel fiction weld Worth, 19951
Al based MMC/stainiess steel fiction weld
Figure 26 : Eutectic Formation at the bondhe of a joint between Nickel and Magnesium pazleît, 19621
Objective of the Thesis
Ni-base wrought aiioy and Ni-base cast alloy base materials were selected during
this study of the Mction welding characte&ics of Ni-based superdoys. The general
microstructure of these base materials is very cornplex, with the y matrix being hardened by
means of spherical a d o r cubical y' intermetallic particles and carbides. Ni-based
superalioys are highly susceptiile to cracking (iiquation cracking, re-heat cracking and
ductility dip cracking) during h i o n welding operatioris. Therefore, solid state joining
processes (fiiction welding) are recormnended for their fabrication. However, high
fictional Ioads and torques are rquired in order to produce defect fiee welds. For example,
it will be shown later that fiction welding of 19 mm diameter Ni-base wrought d o y bar
requires a fiction load of 12 T applied for a fictional time of 10s. This compares with
aluminum ailoy fiction welding where the fictional load is 2 T and the friction time is 0.5
S. During fiction welding of AIS1 304 stainless steel base material, the fiction load is 4 T
and the fiction time is 3 S. Moreover, fiiction welding operation generates considerable
microstructural changes at the bondline and in the HAZ regions on either side of the
bondline, and particularly with regard to the particle distribution and chemistry. These
microstructural changes have a significant effect on the tensile strength properties of
completed welds.
The initial work in this thesis focused on Ni-base wought alloy/Ni-base wrought
alloy friction welding. The influence of welding parameter selection and different post weld
heat treatment procedures on the final joint microstructure and tende strength propert-ies
were studied in detail. Transmission electron microscopy, in-situ transmission electron
microscopy and atomic force microscopy techniques were used to characterise the
microstructure. Microstructure - mechanical property correlations were established for Ni-
base wrought alIoy/Ni-base wrought alloy fiction welds. Finally, the results fkom Ni-base
wrought alloy/Ni-base wrought d o y fiction welds were used as the basis for the welding
parameter settings employed during dissimilar Ni-base wrought alloy/Ni-base cast alloy
friction welding operatiom.
CHAPTER 3
EXPERIMENTAL PROCEUDRE
3.1. Ni-BASE WROUGHT ALLOY FRICTION WELDS
Precipitation hardenable Ni-based superalloys retain theu strength at elevated
temperatures. For this reason it would be expected that very high fictionai loads (e.g.; 12
Tons) acting on 19 mm diameter Ni-base wrought alloy bar would be necessary in order to
obtain sound welds fiee of unbondeci regions. However, M t e d information is available in
the literature conceming fiction welding of Ni-based superalloy and for this reason the
initial welding trials were planned with the objective of geaerating baseline information.
The results fkom these preliminary trials are described in Section 4.1 of Chapter 4 in the
present thesis.
Following prelimuiary welding trials, the experimental work effort examined the
process envelope and the optimum welding parameter settings required during Ni-base
wrought alloy niction welding. The most important welding parameters (niction pressure,
Wction t h e , forging pressure and rotational speed) were varied systematically and their
effect on weld quality was investigated using a combination of optical, SEM and TEM
metallography, micro-hardness testhg and teasile tating. The results are discussed in
section 4.2 of Chapter 4.
Following friction welding the joints were pst weld heat- treated to restore their
mechanical properties and also to relieve the residual stresses generated during the fiction
welding operation. As expected, precipitate morphology and distribution were markedly
affecteci by the post weld heat treatment procedure selected. These fiction welds were heat-
treated using different post weld heat treatmcnt procedures and theù eff- on joint
microstructure and tende strength properties are described in Section 4.3 of Chapter 4.
3.1.1. Preliminary Tests To Genente the Bweliire Data
During these tests, the as-received Ni-base wrought dloy base materid was Ui the
solution-treated condition. The initial test joints were produceci using 25.4 mm diameter Ni-
base wrought alioy bar. It was quickiy leamed that joining of this base materiai required
high fiction pressures and long fiction times in order to produce satisfactory joints.
Testing of 25.4 mm diameter bars resulted in the University of Toronto's direct drive
fiction welding machine operating at, or above, its maximum capability. Consequently, it
was necessary to reduce the sample diameter in order to be able to d e l y operate within the
machine's capabiiities. For this reason, 19 mm diameter bar matenal was used during al1
subsequent trials (see Table 7).
During direct drive fiction welding, the weld is made between stationary and
Table 5 : Experimental testing ma& during preliminary fiction welding trials
rotating components. The rotating section is much shorter than the stationary section. In the
tests at U. of Toronto, the length of the stationary wmponent is 300 mm; it is 100 mm in the
case of the rotating component. It should be noted that these dimensions are the minimum
lengths for adequately safely holding the test sections during the fiction welding operation.
BAR A
BAR B
BAR C
' Ni-base wrought alloy bars A and B had different grain sizes ' ~ h e as-received base material was in solution-treated condition
Bar A' - Diameter
Average Grain Size
Tests Conducted
Bar B' - Diameter
Average Grain Size
Tests Conducted
Bar C - Diameter
Average Grain Size
Tests Conducted
25.4 mm
40-60 microns
Combined effects of fiction pressure and tirne
19 mm
120 microns
Effect of rotational speed (rpm)
19 mm
40-60 microns
Effect of forging pressure
Effect of base metal condition2
Methodology
The following welding parameters were examined:
a) Combined Effects Resulting Friction Pressure and Fricton Time Variations
The initial trials (Test # 1 and 2) were cSmed out usbg fiction pressures ranging
fiom 250 MPa (for a fiction thne of 2s) to 400 MPa (for a Eriction time of 14s). Tests # 3
and 4 were carried out using intexmediate fiction pressure and fiction time values (see
Table 8 below):
Table 6: Experimental test matrix when examinuig the combined effects of fiction pressure
TEST #4
400
14
400
2
and time during Ni-base wrought alloy fiction welding
b) Effects of Rotational Speed
in this test series, the rotational speed was varied £iom 500 rpm to 2000 rpm (ail
other joinuig parameters being held constant). The experimentai parameters are shown in
Table 9.
WELDING CONDITION AND
MEASURED VARLABLES
Friction Pressure (MPa)
Friction Tirne (s)
Forging Pressure @Pa)
Forging Time (s)
Table 7: Experimental test matrix when examining the influence of rotational speed on the Ni-base wought alloy fiction welds (using 25.4 mm diameter bar) TEST # 1 ROTATXONAL SPEED (rpm) 1 WELDING PARAMETERS
Rotational speed was kept constant at 2000 rpm Diameter of the bar used: 25.4 mm
TEST # 1
250
2
250
1
TEST # 2
300
6
300
1
Test # 5
Test # 6
Test#7
Test # 8
TEST #3
400
IO
400
2
2000
1500
1000
500
Friction pressure: 350 MPa
Friction thne: 10s
Forging pressure: 350 MPa
Forging the : 2s
c) Effects of Forging Pressure
The forghg pressure was v e e d h m 200 MPa to 500 MPa, all other joining
parameters being held constant (see Table 10).
Table 8: Test matrix when examining the influence of forging pressure during fiction - - - - welding of 19 mm diameter ~i-base-mught ailoy bars 1 TEST # s. 1 Forging Ressure (MPa) 1 Welding Parameters held constant I
d) ïnfluence of Base Metal Condition
Test # 9
Test # 10
Test # I l
In addition, the influence of initial Ni-base wrought alloy base material condition - whether it was in the hardened or solution-treated condition - on friction welding
performance was investigated. In these tests, the welding parameters compnsed:
Friction Pressure : 350 MPa Friction Time : 10s Forgiag Pressure : 350 MPa Forging T h e : 2s Rotational Speed : 1OOO rpm Bar Diameter : 19mm
200
350
500
3.1.2. Optimization of Welding Parameters
Friction Pressure: 3 50 MPa
Friction The: 10 s
Rotational Speed: 1 0 0 rpm
Forging Tirne: 2 s
The work carried out early in this study when generating the baseline data provided a
preliminary estimate of the fiction welding parameters needed to produced sound fiiction
Ni-base wrought alloy fiction welds using 19 mm diameter bar matexid. These parameters
comprised fiction pressure: 350 MPa, fiction time: los, forging t h e : 350 MPa, forging
time: 2s and rotational speed: 1000 rpm. These preliminary welding parameter settings
served as the basis for fiuther work aimed at determining the optimized welding parameters
during M e r Ni-base wrought alloy friction joining operations
The innuence of fiction pressure variation was examined in detaü, using the
conditions indicated below:
1 Solution Treated Ni-base wrought allorbars I v
Friction Welding Friction Pressure Variation: 275,300.325.350 and 375 MPa The aram met ers k e ~ t constant were:
Friction Time: 10s Foraina Pressure: 325 MPa Fornina Tirne: 2s Rotational S~eed: 1000 mm
Following the fiction welding operation al1 welded joints were subjected to a heat
treatment comprising re-solution + PWHT (stabilization plus precipitation). This heat-
treatment cycle comprised:
Re-solution 9 8 0 " ~ for 2 hours + Air Cool Stabilization 8 4 0 " ~ for 4 hours + Air Cool Precipitation 7 6 0 " ~ for 16 hours + Air Cool PWHT' Stabilization + Precipitation
3.1.3. Optimization of the Post Weld Heat Treatment Procedure
As mentioned earlier, as-received Ni-base wrought alloy base material was in
solution-treated condition. Hardened Ni-base wrought ailoy base materiai was produced by
heat treating the as-received base material using a [stabilization + precipitation] treatment
and this base material was employed d u d g a nurnber of welding trials. Tests were carried
out to examine the welding characteristics of solution-treated and hardened Ni-base wrought
ailoy base materials. The Ni-base wrought alloy bar materials were fiction welded using
the following fiction welding parameters:
- .. -- .
1 note that in the remainder of this thesis PWHT means a (stabilization + precipitation) heat treatment. It is not
a acronym for post weld heat treatment
Friction Pressure : 325 MPa Friction T h e : 10s Forging neJsure : 325 MPa Forging Time : 2s Rotational Speed : 1ûûû rpm Bar Diameter : 19 mm
M e r fiction welding the test joints were subjected to three heat treatments; (PWHT,
Solution-trcatad N i - b Friction 1. PWHT wmught d o y base materiai '
J Weldiag 2. RS+PWHT
* b 3. ST+PWHT
The heat-treatment temperatures and holding tirnes were:
Solution-treatment (ST) : 1090°C for 4 hours + Air Cool Re-solution (RS) : 980°C for 2 hours + Air Cool Stabilization : 840°C for 4 hours + Air Cool Precipitation : 760°C for 16 hours + Air Cool PWHT = [StabiIization + Precipitation]
1 PWHT 2 RS+PWH 3 ST+PWHT
H a r d e d Ni-base wrought alioy base materimaterial
3.2. Ni-BASE CAST ALLOYl Ni-BASE WROUGHT ALLOY
DISSIMLAR FRICTION WELDS
The as-received Ni-base wrought dloy base material was in the solution-treated
condition and Ni-base cast alloy was in the fully hardened condition (hardened Ni-base cast
alloy base material was produced by heat-treating the base material using a [stabilization + precipitation] treatment). The initial test welds were produced using a range of different
fiction pressures; 275,300,325,350 and 375 MPa The other welding parameter settings
were fiction tirne: 10 s, forging pressure: 350 MPa, forging tirne: 2s, rotational speed: 1 O00
Friction Welding
+
rpm. 19 mm diameter bars were used throughout. Foliowing fiction welding ali welded
joints were subjected to a heat treatment comprising re-solution + PWHT (stabiiization plus
precipitation), i.e. ;
Re-solution (RS) : 980°C for 2 hours + Air Cool S tabilization : 840°C for 4 hours + Air Cod Precipitatioa : 760°C for 16 hours + Air Cool PWHT = [Stabiiization + Precipitation]
3.3. EXPERIMENTAL TESTING OF SIMILAR AND DISSIMILAR Ni-
BASE ALLOY FRICTION WELDS
a) Metallography: Al1 test welds were cut and sectioned using a water-cooled &O3 cemnic
wheel (see Figure 27). The sectioned Ni-base wrought alioy welds were then rnolded and
polished to a 1 micron finish using diamond spray. The polished samples were then etched
using Marble's reagent (10 grams of CuSO4 + 10 ml HN03 in 100 ml of Water) with fkesh
etchant was being prepared in each case.
Cutting operaîion Slicing of the welded joint
Figure 27: Sectioning of Ni-base wrought alloy fiction welds prior to metallography
b) Measurement of Axial Deformation: The length of the rotating and stationary components
was rneasured pnor to the fiction welding operation. Following fiction welding the axial
length difference was recorded as the amount of axial shortening (see Figure 28 below).
Figure 28: Schematic showing the measurement of axial shortening during fiction welduig (Amount of axial shortenhg : X-Y)
C ) Micro-Hardness Profiles: A 200 g load was used during micro-hardness testing with
indentations being made at 200 microns intervals h m the joint centerline into the stationary
and rotating Ni-base wrought aUoy components (see Figure 29). AU hardness values were
made at the component centerline up to a distance of 10 mm h m the joint interface.
Figure 29: Location of micro-hardness tests in Ni-base wrought ailoy/Ni-base wrought aîloy joints
d) Tende Testing and Fractography: The tensile test specimen geometry is s h o w in Figure
30. Holding the externometer was particularly difncult due to the test specimen gauge
length. A strain rate of 0.05 in./iiJmin was employed during mechanical testing and al1
broken tensile samples examined using SEM hctography.
Figure 30: Tensile sample design (dl dimensions are in imperid units)
3.4. TRANSMISSION ELECTRON MICROSCOPY
TEM specimens were extracteci h m Ni-base wrought aUoy similar and Ni-base cast
alloyMi-base wrought aiioy dissimilm fiction welds which had been made using the
following parameters:
Friction Pressure 350 MPa Friction Tirne 10s Forging Pressure : 350 MPa Forging Time 2s Rotational Speed : 1 0 rpm Bar Diameter 19 mm
The Ni-base wrought aüoy fiction test joints were then given the following heat treatments:
Ni-base cast alloy/Ni-base wrought alloy fiction welds were given a re-solution + PWHT
(stabilization plus precipitation) heat treatment:
1. As-welded 2. PWHT 3. RS+PWHT
*
Solution-trcated Ni-base wrought alloy base ninttrial
i #
1. -4s-wtlded 2. P m 3. RS+PWHT
Hardened Ni-base wrought aüoy base materid
Friction -D Welding
wrought alloy base mterial
r a
- -
Friction Welding
1
1 NiOb- cast aUoy bue 1
*
-, Friction Weldhg b 1. kewtldcd
2. RS+PWHT
The niction welds were sectioned using a water-cooled ceramic wheel as shown in
Figure 3 1. The extracted components were mechanicaily thinned to a thickness of 75 - 150
microns. This was carrieci out m a n d y using Sic emery paper. A foil thickness of 75-100
microns produced the high quality images during TEM microscopy.
Figure 3 1 : Extraction of TEM test sample following fiction weld
An EDM process was used to extract the discs. This process is extremely slow and
time consuming. Discs were extracted at the bondline and also at off-centered locations in
the fiction weld (see Figure 32). The offlcentered locations correspondeci with the region
containing partly re-oriented grains in completed welds (see Figure 33).
Figure 32 Extraction of discs b m the foil using the EDM process
dise extracted .hm the dynamicall y recrystallised region
Figure 33: Location of offkentered discs for TEM examination discs
The jet polishing (Taiupol) technique was u s d to thin the TEM discs. Tub ,., conditions were apptied:
Solution used : 20 % Perchloric Acid + 80 % Ethanol Temperature o f the solution : - 40°C Potential between electrodes : 16 V Average current : 0.05 A
During the thinoing process, the temperature was maintaineci at - 40°C. Extreme
care was taken to maintain the temperature at -OC suice any rise in temperature during
thinning was extremely dangerous (see Figure 34). Adjacent regions (close to the central
hole produced during thinning) were transparent to transmitted electrons and were viewed in
TEM column.
Figure 34: Jet thinning of the discs
3.5 IN-SITU TRANSMISSION ELECTRON MICROSCOPY
The test specimen was placed on a specially designed fitniace containing a M o
resistance-heating coil. The maximum temperature attainable b ide the fiunace was 850°C.
When the specimen was heated it expauded markedly and this made imaging dinicult during
the initiai stages of the heating procas. AU images were coliected using a Video camera
attached to the TEM wlumn. A schematic representation of the equipment is shown in
Figure 35.
Figure 35: Heating chamber used during in-situ TEM microscopy
The as-welded joints were heat -treated inside the TEM column, using the procedure
hdicated below:
Hardened Ni-based wmught He&-treated inside the TEM ailoy in the as-welded condition column at 850°C for 10 minutes
--
Solution - tteatad N i - b d wrought alloy in the as-welded condition r Heat-treated inside the TEM
J
column at 850°C for 20 minutes
3.6. ATOMIC FORCE MICnOSCOPY
The working principle in AFM microscopy is simple and is illustrated in Figure 36.
The specimen holder vibrates due to a piezo-electric mechanism and the fiequency can be
easily varied. The cantilever is aàjusted such that it just touches the specimen and the laser
beam falls on the Sw cantilever and is then reflected back to a spi& diode. This generates a
potential in the split diode and an extemal potential amplifies the signal. The A-B potential
signai is nonnally kept at -2.2 V and the A+B signal is maintaineci at 7 . W . When the
specimen vibrates depending on the d a c e topography, the intensity of the reflected laser
beam on the split diode varies. During AFM examination, the electrical pulses fiom the split
diode are rnonitored continuously and the changes in (A-B)/(A+B) potential are plotted
versus the scanning distance.
Figure 36: Working principle in AFM Microscopy
The thickness of the foil is not very critical; however, great care must be taken not to
bend the foil during poüshing and etching. In the present study the foil was polished
manually using a polishhg wheel to 0.25-micron finish and was then etched using Marble's
(10 grams of CuS04 + 10 ml HN03 in 100 mi of Water) reagent.
RESULTS AND DISCUSSION
CHAPTER 4.1
Ni-BASE WROUGHT ALLOY FFUCTION WELDS
4.1. GENERATING THE BASE LINE DATA
Initial welding trials were carried out to generate the basehe data during Ni-base
wrought alfoy fiction welding. The welding parameters that a81ect the joint quality
comprise fiction pressure, fiction time, forging pressure, forging time, rotational speed and
the prior base material condition. In the present study, Ni-base wrought alloy base material
was weldeà in the solution-treated and hardened conditions. The welding parameters were
varied systematically and their effet on weld quality was investigated using a combination
of metallography and micro-hardtless testing. U of Toronto's Giction welding equipment
has a limitation regarding welding parameter settings. The maximum fictional load that can
be applied to a test piece is restricted to a maximum of 15 T. Further, the machine only
performs well using rotational speeds between 1000 to 2000 rotations per minute. The
initial welding tests provide a rough estimate of the operating envelope required during Ni-
base wrought d o y fnction welding.
4.1.1 Base Metal Microstructure
Ni-based superalloy material microstnictures are complex and comprise an austenitic
fcc matrix (y) contaking a variety of secondary phases. These secondary phases are the
gamma prime ((Ni,Co)&Ti intermetaliic phase and various carbides, namely, MC, MsC,
M7C3, Mac6. Figures 37 and 38 show SEM micrographs of Ni-base wrought alloy base
material in the solution-treated condition. An optical micrograph of hardened Ni-base
wrought alloy is shown in Figure 39. The grain boundary regions are etched preferentially
in this micrograph.
Figure 37: SEM microstructure of solution-treaîed Ni-base wrought alioy base materid
Figure 38: Higher magnification view of solution-treated Ni-base wrought alloy base material
Solution-treated Ni-base wrought alloy is an alloy supersaturatecl containhg different
elements and precipitation occurs when the ailoy is cooled to room temperature foilowing
solution treatment The most cornmon precipitates are y' particles and high temperature
carbides of the MC type. Figure 40 shows a TEM micrograph of solution treated Ni-base
wrought ailoy base material. The maîrix structure comprises unifody distributeci gamma
prime particles contained in a y ma&.
X-ray Energy Dispersive Spectrometry (XEDS) was used to examine the chemistry
of precipitates in the Ni-base wrought d o y base material. High resolution transmission
electron microscopy (HRTEM) was used instead a conventional TEM since the XEDS
pattern produced by the HRTEM microscope was more accurate and diable- Figure 41
shows an XEDS pattern produced by the marrix. The matrix compnsed an ailoy of Ni, Co
and Cr. However, trace amounts of Ti, Al and Mo were also indicated. The XEDS pattem
fiom a coherent precipitate is shown in Figure 42. This particle was low in Cr and Co and
contained mainly Ti, Al and Ni; its composition corresponded with the formdation
Ni3(Cr,Co)AiTi.
The microstructure of hardened base material was quite different fiom that of the
solution-treated Ni-base wrought alloy substrate. Hardened Ni-base superalloy contained
large volume fiaction of gamma prime particles and carbides compared to the alloy in the
solution-treated condition. A h y the hardening treatment promoted extensive carbide
precipitation dong grain boundaries, see Figure 43. An XEDS pattern nom a grain
boundary carbide is shown in Figure 44. No traceable contents of Al or Ti were present in
the carbide. The carbide was rich in Cr and cuntained smaller amounts of Co and Ni and
confirmed with the M d 6 type where M corresponds to Cr and Ni. Figure 45 shows an
XEDS pattern fiom the matrix and Figure 46 shows the XEDS pattem fiom the gamma
prime particles. The maîrix compnsed an alloy of Ni, Co and Cr. However, small amounts
of Ti, Al and Mo were indicated.
Figure 39: Microstructure of hardened Ni-base wrought alloy base material
XEDS of the mat&
-XEDS of the precipitate
Figure 40: TEM microstructure of solution-treated Ni-base wrought alloy base material
Client r Sugi Spthirn Job : Job numbor 387
CP* SpsetNrn 2(Sf12JQe f1:lT)
Energy (kew
Figure 4 1 : XEDS pattem of the solution-treated Ni-base wrought aiioy y matrix
Client : Suai Sathian ~ o b : Job numbor 387 Spectrum 1 < 5 1 l ~ l l : t l >
cps ?
Figure 42: XEDS pattern of y' particles in solution-treated Ni-base wrou&t alloy base material
XEDS spot
- XEDS analysis
Figure 43: TEM microstructure of the hardened Ni-base wrought alloy base material
Figure 44: XEDS Pattern h m a grain boundary carbide in hardened Ni-base wrought ailoy base material
Client : Sugi S m t h l r n Job : Job numbor 367
cps <5/12JB8 lP:OS>
Figure 45 : XEDS analysis of the y matrix in hardened Ni-base wrought aiioy base material
Figure 46: XEDS pattern fiom a y' precipitate of the hardened Ni-base wrought alloy base material
4.1.2. As-Welded Joint Microstructure
TEM discs were extracted h m the dynamically recrystaliïsed region and also nom
the partly defomed region in completed joints (see Figure 33). Figure 47 shows a TEM
micrograph of the dynamïcally recrystaliised grains in the as-welded joint. The grain size
was approx. 2 - 3 microns in diameter. This location corresponds with the M y plasticised
region (Zpl), where the material is able to accommodate plastic strain via dynamic recovery
(or recrystallisation) of the microstructure. Dissolution of y' is facilitated since these
particles are coherent with the rnaûix. Therefore, fiction welding creates a region of highly
alloy super-saturation at the bondline of as-welded joints.
Figures 48 shows the particle segregation and Figure 49 shows the particle fracture. The
TEM specimen was extracteci h m the off-centted location, Le. h m the partly deformed
region (Zpd), where the degree of plastic deformation is accommodated by an increase in the
dislocation density in the matrix gains. The base material grains were orientateci dong the
flow Iines due to fiction welding operation caused particle segregation while movement of
dislocation resulted the particle fhcturing.
Similar results were observed when hardened Ni-base wrought alloy base material
was fkiction welded. A TEM micrograph of the dynamically recrystallised grains in an as-
welded joint is show in Figure 50. No precipitates were observed in the dynamically
recrystallised region. Figure 5 1 shows a micrograph of the partly deformed region (Zpd).
Particle segregation was observed in this region. Figure 52 shows MC carbides, which are
easily differentiated fiom the gamma prime particles. The y' particles were coherent in
nature and were somewhat opaque to the transmitted electrons. However, the carbides
scatter the electrons and therefore appear dark.
TEM microscopy did not produce a complete p i c m of the particle
distribution over an area of 10 microns2. For this reason, Atomic Force Microscopy was
used to image the particle distribution over an area 15 micron x 15 micron. The scanned
area in an AFM cm be varied easily: 0.1 micron x O. 1-micron to 15 micron x 15 micron
(max.) regions can be investigated.
dynamicaily recrystallised grains
Figure 47: TEM mimgraph h m the dynamicaliy recrystallised zone (solution-treated Ni- base wrought alloy fiction weld in the as-welded condition)
Figure 48: TEM micrograph showhg the k t u r e of particles in the partly deformed region in a solution -treated Ni-base wrought alloy fkction weld
particle segregation
Figure 49: Particle segregation in the undefomied region of a solution-treated Ni-base wrought aiioy fkiction weId, As-welded condition
dynamicaiiy recrystalIised grains
Figure 50: TEM micmgraph showing the dynamically recrystallised region in hardeneci Ni- base wrought dloy fiction welds. As-welded condition.
particle segregation
Figure 5 1 : Segregation of y' in undeformed region of hardened Ni-base wrought alloy friction welds. As-welded condition
Figure 52: TEM micrograph of a carbide containeci in the y ma& of hardened Ni-base wrought alloy fiction welds
AFM is generally used for imaging d5u:e topographical features, e.g., the
topographical changes occurring when Si semiconductor material is doped with different
impurity elements. The limitation ofthe AFM technique is that the test specimen
preparation is very difncuit and it is difficuit to identifL the exact location that is being
imaged by the microscope.
Figure 53 shows an AFM image of Ni-base wrought alloy base metal. A grain
boundary region is clearly evident. Also, a topographical image of the dynamically
recrystallised region in the welded joint is shown in Figure 54. As rnentioned above it was
difficult to be precise conceming the area that was being scannai during AFM micmscopy.
4.1.3. Combined Effets of Friction Pressure and Friction T h e
The initial weiding trials were cmied out ushg fiction pressures ranging h m 250
MPa (for a fnction time of 2s) to 400 MPa (for a fiction time of 14s). Joints # 2 and 3
were made using intermediate fiction pressures (300 - 400 MPa) and fiction times (6 - 10s)
(see Table 8).
Unbondeci regions were fo& at the periphery of welded joints produced using low
fiction pressures (c 400 MPa) and friction times (40s). This is apparent in Figure 55,
which shows a micrograph of the welded joint produced using a fiction pressure of 250
MPa and a fiction time of 2s.
Figure 56 shows the grain sue variations in a traverse h m the bondiine into the as-
received base material. The bondline microstructure wmprised partially recrystallised
grains having diameters e 10 microns. Further h m the bondline, the microstructure
contained equiaxed dynamically recrystailised grains (having grain sizes in the range 10 - 20 microns). With increasing distance h m the bondline, the microstructure comprised te-
orienteci base metal grains. In efféct then was a transition h m a fine grain microstructure
at the bondihe to larger re-orieated base metai grains and fbally into the as-receivexi Ni-
base wrought alloy base material (where the grain sue was 60 microns). It would be
expected that the presence of very fine equiaxed grains at the joint centerline rnight have a
detrimental effect on elevated temperature properties (since grain boundary sliding
preferentially occurs at grain boundary regions that are aligned perpendicular to the
direction of tensile stress [Thamburaj, 19951).
The welding parameters applied during welding tests # 1 (friction pressure 250 MPa,
fnction time 2s) and # 2 (fnction pressure 300 MPa, fiction t h e 6s)) resulted in the
formation of unbonded regions near the joint periphery. The formation of these defects can
be explained as follows. When low fiiction pressure and iow fnction times are applied, the
filly plasticised region develops rapidly at the component centerline. However, formation
of the fully plasticised region at the periphery depends on transfer of plasticised material
fiom the component centerline towards the periphery. When the fnction welding time is
v e v short or the fiction pressure is low, this inhibits transfer of M y plasticised material
fiom the component centerline to the joint periphery and creates a region of weahess.
Satisfactory bonding across the whole joint interface and satisfactory weld profiles
were produced when higher fiction pressures (> 350 MPa) and longer fnction times (> 10 s)
were applied. However, since the fiction welding machine at University of Toronto had a
maximum rated capacity of 15 tons, a fiction pressure value around 400 MPa (on 25.4 mm
diameter bars) represented its lirnit of machine performance. For this reason, the peak
fiiction pressure used during subsequent fiction welding trials was necessarily limited to
350 MPa.
Based on the above commentary, it is concluded that the application of higher fiction
pressures (>350 M'a) and long Wction times (>los) are prime requirements during Ni-base
wrought alloy fiction welding and that this welding parameter combination produces welds
fiee of unbonded region located at the joint periphery.
Figure 53: AFM Image of Ni-base wmught aUoy base material (the arrow shows the grain boundary)
Figure 54: AFM Image of the HAZ region in a fiction weld produced using the conditions; Friction Pressure: 350 MPa (tirne: 1 Os), Forging Pressure: 3 50 MPa (the: 2s), Rotational Speed: 1000 rpm
Base me!
- xysîaiiised
iented base
ion
1 , - =
Figure 55: Welded joint produced using a low fiction pressure and fiction time (fiction pressure = 250 MPa, fiction time = 2s). The other parametm were; forging pressure = 250 MPa, forging time = 1 s, Diameter of the bar = 25.4 mm.
Distancm h m tho dge of the partklly nctyst8lli.d mgion (mm)
Figure 56: Grain size variation in a traverse h m the edge of the partially recrystallised region into the as-received Ni-base wrought alloy base material.
The welds were made using the conditions:
A. Friction Pressure: 250 MPa (time: 2s) + Forging Pressure: 250 MPa (time: 1s) B. Friction Pressure: 300MPa (time: 6s) + Forging Pressure: 300 MPa (tirne: 1 s) C. Friction Pressure: 400 MPa (the: 14s) + Forging Pressure: 400 MPa (tirne: 2s)
Rotational speed: 2000 rpm Bar Diameter: 25.4 mm Base metal grain size: 40 - 60 microns
4-1.4, Rotationaï Speed and Weld Qurlity
In this series of tests, the rotationai speed was varied h m 500 rpm to 2000 rpm (dl
other joining parameters k i n g held constant (f?ïction pressure : 350 MPa, fiction time :
los, forging pressure : 30 MPa, forging tirne : 2 s). Details of the experimental parameters
are shown in Table 9. It is weil known that the torque increases significantly when the
rotational speed decreases h m 2000 rpm to 500 rpm. As a result heat generation at the
contact interface markedy increases. Increased heat generation had a signincant infiuence
on the Gnal joint microstructure; e.g., when the rotational speed increased f?om 1000 to 2000
rpm the grain size increased h m 5 to 10 microns. This compares with an as-received Ni-
base wrought alloy base metal grain size of 120 microns (see Figure 57).
The torque produceci by the rubbing suffies substantially increased when the
rotational speed decreases. Consequently, although low rotational speeds inhibited the
formation of unbonded regions at joint peripheries when 25.4 mm diameter Ni-base wrought
alloy bars were welded, this methodology could not be generally applied because it
compromised the safe operation of the fiction welding machine (the safe operating range
being fiom 1000 to 2000 rpm).
Finally, there was a significant difference in the grain sizes measured on either side of
the bondline, i.e., in the stationary and rotating components (see Figure 58). Note that the
diameter of the Ni-base wmught alioy bar used prior to welding was 25.4 mm. This
diameter necessitates the application of high fiction ioads (loads>l ST). That means, the
welding machine was run at its maximum rated capacity and therefore, the results obtained
may not be the true representation. However, it can be presented for the cornparison
purposes. The grain size differences in the stationary and the rotating components resulted
fiom the different rates of heat absoption produceci by the clamping jaws in the case of the
shorter rotating section.
O 0.5 1 1.5 Dirbnce m m the cornponant t.ntreline(mm)
Figure 57: Influence of rotational speed (hm 1000 rpm to 2000 rpm) on grain size variation measured at the component centerline. As-welded Ni-base wrought alloy friction weld.
The measurements were carrieci out f3om the component centerhe
Welding conditions applied:
A. Rotational speed: 2000 rpm B. Rotationai speed: 1500 rpm C. Rotational speed : 1 O00 rpm
(Friction Pressure: 3 50 MPa (tirne: 1 0s) + Forging Pressure: 3 50 MPa (time: 2s) Grain size: 120 microns Bar Diameter: 25.4 mm
Distance from the wmponent centeriine (mm)
Figure 58: Grain size changes on either side of the bondline (in the stationary and rotating) components of the fiction weld As-welded Ni-base wrought alloy fiction joint.
The measurements were carried h m the joint centerline
Welding parameters applied:
Friction Pressure: 350 MPa (tirne: 10s) + Forging Pressure: 350 MPa (time: 2s) Rotational Speed: 1000 rpm Diameter of the bar: 25.4 mm Grain size: 120 microns
4.1.5. Forging Pressure and Weld Qurllty
In these t a c the forging pressure was varied fiom 200 MPa to 500 MPa, aii other
joining parameters being held constant (see Table 10). Forging pressure had no effect on the
grain size of recrystallised grains at the joint centerline. Figure 59 shows micro-hardness
profiles from the bondline into the as-received base metal (all measurments were carried
out at the centerline of the component). It is apparent that increased forging pressure had no
significant effect on hardness values (bearing in mind the typical scatter in resdts produced
during micro-hardness teshg).
4-1.6 Influence o f Base Metaï Condition
In general practice, solution-treated Ni-based superalloy material is fiction welded
and is then post weld heat-treated to produce the required joint strength and ductility
properties for any given application. In the case of Ni-base wrought alloy base material, this
involves the application of a two-stage thermal post treatment comprising stabilization at
8 4 0 ' ~ for 4 hours followed by a precipitation hardening treatment at 760'~ for 16 hours.
The initial stabilization heat-treatment promotes carbide precipitation at grain boundaries
and also coarsens the primary intennetallic phases. The low temperature precipitation
hardening treatment promotes precipitation of srnail secondary intermetallics, which
improve material strength. In the present study, fiction welding of both solution-treated and
hardened Ni-base wrought alloy base materials was investigated. Al1 fiction welded joints
produced using solution-treated and hardened Ni-base wrought alloy base materials were
post weld heat-treated using a stabilization + prezipitation thermal cycle.
Figures 60 and 61 show the micro-hardness results in a traverse h m the bondline into
the as-received Ni-base wrought alloy base metal. The indentation load was 200 g and 10
readuigs were taken on each test sample over a distance of 2 mm. When solution plus heat-
treated Ni-base wrought ailoy base material was fiction welded, this produced a sofiened
zone with a hardness of about 25 Hv lower than the adjacent base material. The hardness
trough was much deeper (about 150 Hv) in fiction welded joints produced ushg hardened
Distance from aie bondline (mm)
Figure 59: Microhardness profiles in as-welded Ni-base wrought alloy friction joints in a traverse fiom the bondline into the base material (al1 measureements at the centerline of the component).
The welding conditions comprised:
A. Forging Pressure: 200 MPa (the: 2s) B. Forging Pressure: 350 MPa (time: 2s) C. Forging Pressure: 500 MPa (the: 2s)
(Friction Pressure: 350 MPa, Rotational Speed: IO00 rpm) Bar Diameter: 19 mm
A - As-welded condition
Distance from the bondline (mm)
Figure 60: Microhardness values before and following post weld heat - treatment in a traverse fiom the bondline into the solution-treated Ni-base wrought alloy base matenal.
The measurernents were taken h m the bondline region
The welding conditions comprised:
Friction Pressure: 350 MPa (time: 1 Os) Forging Pressure: 350 MPa (time: 2s) Rotationl Speed: 1000 rpm Bar Diameter: 19 mm Base metal condition: Solution treated Ni-base wrought alloy
A. As- welded condition B. After PWHT
Distance from the bondline (mm)
Figure 61: Micro-hardness results before and following p s t weld heat treatment in a traverse fiom the bonche into the hardened Ni-base wrought alloy base material.
The measurements were carried out h m the bondline region
Welding conditions compnsed:
Friction Pressure: 350 MPa (time: 10s) + Forging Pressure: 350 MPa (time: 2s), Rotational Speed: 1000 rpm Bar Diameter: 19 mm Base metal condition: Hardened Ni-base wrought alloy
Ni-base wrought aiioy base material. The formation of softened zones on either side
of the bondline resulted h m solution and coarsening (over-aging) of intennetallic phases
and carbide pdc l e s during the thermal cycle in fiction welding. These softened regions
were completely removed foliowing p s t weld heat treatment. Similar hardness troughs
have been reported in the in the HAZ regions of 6Ml- T6 a1uminu.m fiction welds Malin,
19951.
To conclude, similar microstructural features were produced in post weld heat-treated
(stabiiization -t precipitation) joints when the starting Ni-base wrought ailoy base materid
was in the solution treated and hardened conditions. However, this may not mean that the
mechanical properties of the completed joints will necessady be similar. When using a
rotational speed of 1OûO rpm, microstnicturaUy sound, defeît fiee fiction welded joints
having excellent profiles were produced ushg the following welding parameters: a friction
pressure of approx. 350 MPa, a fiction t h e of 10 s and a forging pressure of 35OM.a (see
Figure 62).
Figure 62: Ni-bas fiction pressure =
M k i and a fiction time = 2s. The 0th- welding parameters were forging pressure = 350 MPa, forging time = 2s and rotational spccd: 1ûûû rpm
CHAPTER 4.2
RESULTS & DISCUSSION
4.2. OPTIMIZATION OF WELDING PARAMETERS
Initiai welding trids provideci a preliminary esthate of fiction welding parameter
settings needed to produce sound Ni-base wmught alloyMi-base wrought alloy friction
welds. These welduig parametm comprised a Wction pressure of 350 MPa, a fiiction tirne
of los, a forging pressure of 350 MPa, a forging time of 2s and a rotational speed of 1000
rpm. These parameter senings served as the basis for f.urther studia aimed at optiminng
welding parameter settings. In this work, the fiction pressure was varied systematically,
Le., 275,300,325,350 and 375 MPa The rotational speed was maintaineci constant at 1 O00
rpm for the foliowing reasons. When using rotational speeds exceedhg 1000 rpm, fiction
pressures >>350 MPa were required to produce sound Ni-base wrought alloy friction joints.
However, the application of such high fiction pressures pushed the U. of Toronto fiction
welding machine beyond its operating capability. Similady, the use of rotational speeds
c 1000 rpm had a detrimental influence on the performance of the fiction welding machine
(since the torque increased considerably when rotational speed was reduced). Ail test welds
were subjected to re-solution + PWHT (stabilization + precipitation) heat treatments and
were analyzed using a combination of optical, scanning microscopy, micro-hardness and
tende testing.
4.2.1. Weld Profiles
Figure 63 shows the effects of fiction parameters on the amount of axial shortening
produced during welding. The amount of axial shortening increased when the friction
pressure increased. For example, a fiction pressure of 275 MPa produced an axial
shortening value of 3 mm while a fiction piessure of 375 MPa produced an axial shortening
value of 8 mm.
300 360 Friction Prrwm (MP.)
Figure 63: Influence of fiction pressure on the amount of axial shortenhg during Wction welding of Ni-base wrought alloy base materiai
Figure 64: Flash regions produced using different friction pressures (275,325 and 375 MPa) during Ni-base wrought alloy fiiction welding. The other welding parameters were: fiction time = los, forging pressure = 350 MPa, forging time = 2s and rotational speed = 1000 rpm.
Figure 64 shows CCD images of the flash regions produced during welding using
different fiiction pressures. It is apparent that the amount of flash increased wheu the
niction pressure increased. For example, nibstantiai amounts of flash were produced when
a fiction pressure of 375 MPa was appl id When high niction pressure (>275 MPa) and
the friction times was los, the fbily-plasticised region developed rapidly at the component
centerline and transfer of plasticised material h m the component centeriine towards the
periphery occurred. When the fiction welding t h e was very short or the fiction pressure
was too low, this inhibitecl transfer of My-plasticised material fkom the component
centerline towards the joint periphery. Thus a region of weakness was created at the joint
penphery. Therefore, higher friction pressures (2275 MPa) are recommended based on a
flash formation perspective of Ni-base wrought alloy Ection welding.
The flow region in completed welds was deheated by the presence of dynamically
recrystallised grains. The width of this region decreased when the Wction pressure
increased. (see Figure 1 7). Bendzsak and North modeled pendzsak, 19971 the flow regions
produced during fiction welding and showed that the width of the flow region (h<,))
depended markedly on the welduig parameters selected, e.g.,
where, p is the viscosity, r is any radius of the sample, w is the rotational speed, q is the
coefficient of fiction, & is the outer radius of the sample and Pa is the applied pressure.
This relation indicates an inverse proportionality between fiction pressure and the width of
the flow region produced during the fiction welding operation. Ifit is assumed that the heat
generated at the contact region conducts equdy into the adjoining substrates, it would be
expected that the profiles of the hcat-affected-zone region on either side of the bondiine
would mirror that of the flow region This may explain the weld profiles shown in Figure
65.
4.2.2 Hardness Profiles
Figures 66a to c show hardness profiles in welds produced using different fiction
pressures. As noteâ earlier, a sofiened zone with a hardness of about 25 - 50 Hv lower than
the adjacent base material is produced when solution treated Ni-base wrought ailoy base
material is fiction welded, (see Section 4.1). The softened zones were completely removed
Friction Pressure: 325 MPa Friction Pressure: 275MPa Figure 65: Weld profiles produced using different fiction pressures (275 and 325 MPa) during Ni-base wrought aUoy friction welding. The other parameten were Enction time = los, forging pressure = 350 MPa, forging tirne = 2s and rotational speed = 1000 rpm
following the re-solution + PWHT (stabilization plus precipitation heat treatment).
However, when a fiction pressure of 275 MPa was applied sottened zones were still
apparent in Ni-base wrought alloy base materid immediately adjacent to the bondline.
Kowever, when a fiction pressure of 325 MPa or higher was applied the hardness profiles
were much more uniform (see Figure 66 c - d).
4.2.3. Tensile Strength Properties
Figures 67 shows a typical stress-strain curve produced when testing a Ni-base
wrought alloy/Ni-base wrought alloy weld. In ail the test samples the uitimate tensile
strength was the fiacture strength of the test samples. The ductility was evaluated as the
total elongation measued h m the stress- strain curve. In this thesis, the yield strength was
the stress corresponding to a strain of 0.02%.
Figure 68 compares the yield and ultimate tensile strengths while figure 69 shows the
ductility results. Ni-base wrought alloy fiction joints made using a fiiction pressure of 275
MPa had the poorest strength pmperiies (YS: 840 MPa, UTS: 1270 MPa). The highest
tensile strengths (YS : 925 - 1070 MPa, UTS : 1360 - 1475 MPa) were produced when
intemediate fiction pressures (300MPa to 350 MPa) were applied. In all the cases the
tensile sample failure occurred in base metal away fiom the weld region and the ductility
values exceeded 15% (Figure 69). It should be noted that the elongation measurements were
solely based on the extensometer rendings of broken tensile test samples and therefore
should be considered as approximate estimates only.
The ultimate tensile strength of as-received solution-treated Ni-base wrought alloy
base material was 1100 MPa and the ductilïty was 22%. Higher ultimate tensile strengths
(UTS values > 1275 MPa) and appreciable ductility levels (215%) were obtained when Ni-
base wrought alloy base material was welded using fiction pressures ranging fiom 325 to
375 MPa. Also, as indicated earlier, the tensile test specimen failure occmed in the base
metal away h m the bondline.
(a) Friction Pressure: 275 MPa
4 -2 ,,,je O 2 4
Distance frorn the bondllne (mm)
-4 2 O 2 4
Distance from the bondlino (mm)
(b) Friction Pressure: 325 MPa
-1 O -8 4 4 -2 #' O 2 4 6 8 10
Distance from the bondline (mm)
(c) Friction Pressure: 375 MPa
Figure 66 (a to c): Hardness profiles at the bondiine in Ni-base wrought aiioy friction welds produced using different fkiction pressures (275,325 and 375 MPa). The other welding parameters were Wction tirne: los, forging pressure: 350 MPa, forging tirne: 2s, rotational speed: 1000 rpm.
Figure 67: Typical Stress Strain c w e produceci during testing of Ni-base wrought ailoy fiction joints (Friction Pressure: 375 MPa Friction Time: 10s. Forging Pressure: 350 MPa, Forging Time: 1 Os, Rotational Speed: 1 Oûû rpm)
Yield Strength mtimate Tcosile Strength
Figure 68: Tensile strengths properties (yield strength and ultimate temile strength) of Ni- base wrought alioyMi-base wrought aiioy fiction welds
Figure 69: Tensile Ductility (total elongation) during tensile testing of Ni-base wrought alloy fiction joints
The fkcture surfaces of the broken teusile samples were examineci using scanning
electron mimscopy. Typical k t u r e d a c e s are shown in Figures 70 a and b and
rnagnified views of broken t d l e test specirnens are shown in Figures 71 a and b. Sample
failure occurred at grain boundary regions in an inter-granular fkture mode. In Ni- based
superalloys it is well documenteci thaî the grain ôoundaries are weak compared to the grain
interiors (since precipitation strengthens the grain interiors) [MM, 19821. Therefore, failure
is therefore more likely at grain bomdary regions, Le., the intergrandu failure mode is
more likely. Increasing the fiction pressure oniy had a minor influence, Le. secondary
cracks were observeci in fkactureâ tende test samples produced using high fiction
pressures.
Figure 70 a and b: Fracture SUff'es of broken tensle sample h m Ni-base wrought ailov/Ni-base wrounht allov weld
Figures 71(a-b): Magnified Mew of the above sample (see Figure 70 a-b)
CHAPTER 4.3
RESULTS & DISCUSSION
4.3. OPTIMIZATION OF POST WELD HEAT TREATMENT
PROCEDURE
Post weld heat treatment procedures Vary h m one application to another. For
example, it has been suggested that friction welding of hardened Ni-base wrought alloy bar
produces superior fatigue strength properties compareci to soiution-treated base material.
Following fiction welding the joints were heat-heated using dinerent procedures, namely:
O P m which involves stabilization + precipitation treatments:
S tabihtion : 860°C for 2 hours + Air Cool Precipitation : 740°C for 16 hours + Air Cool
RS (re-solution) + PWHT
Re-Solution : 980C for 4 hours + Air Cool S tabilization : 860°C for 2 hours + Air Cool Precipitation : 740°C for 16 hours + Air Cool
Solution - Treaîment : 1090°C for 4 hours + Air Cool Stabilization : 860°C for 2 hours + Air Cool Precipitation : 740°C for 16 hours + Air Cool
Precipitate morphology and distribution were markedly affectecl depending on the post weld
heat treatment procedure that is applied. In the present study fiction welds produced using
previously found welding parameter settings (fiction pressure: 350 MPa, fiction tirne: los,
forging pressure: 350 MPa, forging tirne: 2s and rotational speed: 1000 rpm) were subjected
to the above heat treatments. The effects of these heat - treatments were investigated using
a combination of transmission electron microscopy, micro-hardness and tende testing.
4.3.1. Microstructurai Aspects
a) Microstructure Following P m (Stabikation + Prenpiation) Treatment
Figure 72 shows massive (0.02-0.03 pm diameter) y' precipitation occurring in the
dynamically recrystalïised region of Ni-base wrought alloy friction welds given a PWHT
(stabilization + precipitation) heat treatment. The precipitation of particles follows a C-
curve relation, i.e., at temperatures exceeding T2 in figure 73, the kinetics of the
precipitaîion are controlled by diaision processes while at lower temperatures (temperatures
a,), they are controlled by nucleation phenornena. At very low temperatures, long tunes
are required for complete precipitation because the diaision rate is very slow. The rate of
precipitation is also very slow at temperatures just below the solvus line (see point 1, Figure
73). In this case nucleation is slow and precipitation is determined by the rate at which
nucleation can occur. Although high diffiision rates exist at temperatures just below the
solidus line, few nuclei are present. At intermediate temperatures, between these two
extremes, the precipitation rate increases to a maximum value so that the time required for
complete precipitation decreases. In this manner, the combination of moderate diffiision and
nucleation rates promote rapid precipitation. This explains why rapid precipitation occurs
during PWHT. In effect, nucleation during the PWHT (stabibtion + precipitation) heat
treatment will be favored thermodynamically rather than via d i h i o n and this will enhance
the precipitation of small particles within the ma&.
The superalloy matrix and the particles were analyzed using X-ray Energy
Dispersive Spectrometer (XEDS). This work was carried out using a conventional TEM
microscope and therefore, the values are not as accurate as in the case of high-resolution
transmission electron microscopy (HRTEM). However, the output is still usefid for
cornparison purposes. The XEDS result fiom the matrix is shown in Figure 74. X-ray
analysis of precipitates was evaluated using two différent techniques; (i) using a
conventional XEDS analyzer, (ii) using the scanning transmission electron microscope
(STEM) mode. In the normal TEM mode, condensing the bearn down to an appropriate size
massive precipitation of 7' particles
- XEDS of particle
XEDS of matrix
Figure 72: TEM micrograph showing massive precipitation in the dynarnically recrystallised grains of a Ni-base wrought alloy fiction welds following pst weld heat treatrnent (P WHT)
kinetics of precipitation: difkion controlled
Figure 73 : C<urve showing different precipitation mechanism in Ni-based superalloys
for micro-analysis can misalign the iiiumination system. Hence, the STEM mode is
preferred for XEDS anaiysis. The pattern produced using the nrst technique is shown in
figure 75 while the chart produced using the STEM technique is shown in figure 76. Both
approaches produced similar d t s , bearing in mind that only comparative elmental output
not absolute values are considemi. The XEDS results indicated that the matrix contained
approx. 3 wt% Ti and approx. 64 wt.% Ni. nie particles comprised 75% Ni and 10% Ti
and contained smaller amounts of Co and Cr. Similar results were produced when the
precipitates were examineci using the STEM mode. In this case, the values comprised 73
W.% Ni and 6.5 wt.% Ti and trace amounts of Cr and Co.
b) Microshrrciure Following RS me-Solution) + PWHT Treatment
Coagulation of y' particles occurred during the re-solution heat treatrnent. Also, fine
secondary precipitation occurred during the stabilization plus precipitation heat treatment
(see Figure 77). The coagulation priocess limited massive precipitation of the fine gamma
prime precipitates that were observe- in samples given the PWHT treatment (see Figure 72).
The re-solution temperature (980°C for 4 hours) exceeded T2 in Figure 73 and were
determined by diffûsion mechanisms and this enhanceci coagulation controlled the
precipitation kinetics.
Re-solution treatment temperatures vary fiom one superaiioy to another, depenciing
on the [Al+Ti] contents (on the y' content) in the Ni-based superailoy. Hïgher re-solution
temperatures are required for alloys containing high [Ai+Ti] contents. The heat-treatment at
980°C for 2 hours partiaiiy dissolves precipitates in the microstructure (see Figure 78) in the
Ni-base wrought alloy.
Figure 79 shows a typical 'IïT (the-temperature-transfomation) diagram that is
helpfbl in understanding the influence of time and temperature on the final microstructure
following heat treatrnent. Stage 1 in Figure 79 is the re-solution treatment while Stage II is
the stabilization treatment. Stage III is the precipitation heat-treatment. The following
processes occur during heat treatment:
X - R f i Y L t : 100s R t : 119s
P r s 1
Rem:
Figure 74: XEDS pattern h m the y matrix in p s t weld heat-treated (PWHT) Ni-base wrought dloy fiction welds.
Conventionai XEDS .:-RAY ~ t , : 100s P r s t : 100s Rem: 8s R t : 1215 17%Dt
Figure 75: XEDS pattem h m a y' precipitate (using conventional TEM)
X - R A Y L t : 1005 P r s t : R t :
100s 1225 18%Dt 1
STEM MODE
Rem: 0s
Figure 76: XEDS pattern fkom a y' particle (using STEM mode)
Figure 77: TEM miciostructure followkg RS + PWHT treatm base wrought alloy fiction weld
.ent of a solution-treated Ni-
Figure 78: Changes in re-solution tempcrature with [AhTi] content in the superalloy [weaver, 19981
Figure 79: Time-Temperature-Transition diagram for Ni-base wrought alloy base material [Schubert, 19801
a) Stage L-
0 There is partial dissolution of precipitates in the matrix, and high temperature
carbides (MC) and gamma prime precipitate during cwling to rmm temperature.
b) Stage lr:
0 Coarsening of existing y' phases occurs and also precipitation of low temperature
carbides (e-g-, MuCs) dong grain boundaries.
c) Stage iJlk
0 Nucleation of secondary fine y' precipitation occurs within the existing y'
precipitates. However, their growth is limited because of the low temperature
kinetics.
These different stages readily explain the coagulation of y' particles observed in welds that
were given a re-solution heat treatment (see figure 79).
c) Microstructure Following ST (Solution - Treatment) + PEUT Treatment
Figures 80 a to d show optical rnicrographs of the dynamicalfy recrystallised region
fomed at the bondline in completed welds. The right-hand micrograph in each set shows
the selected location at higher magnification. Considerable grain growth was observed
when the Ni-base wrought alloy fiction welds were heat-treated using a [ST + PWHT]
thermal cycle. Further, solution-treated Ni-base wmught alloy base material was more
susceptible to grain growth than hardened Ni-base wrought ailoy base material. The
hardened Ni-base wrought alloy base material contains high temperature carbides that are
very effective in pinning the grain boundaries and in reçtncting growth during post weld
heat treatment [Thamburaj, 19851.
d) Microstructure Following In-Situ Heat - Treatment in the TEM Microscope
The samples for in-situ microscopy were extracted hrn the dynamically
recrystallised region in as-welded joints. As pointed out eatlier this region is highly
supersaturateci with ailoying elements. When the test sample was heated in the TEM
column rapid precipitation occumd and the output was imaged. The test section was heated
inside the TEM column using a Mo resistance heating coi1 and the microstructural changes
Base Metal Condition: Hardened Ni-base wrought aiioy; Heat Treatment: PWHT
Base Metal Condition: Solution - Treated Ni-base wrought alloy; Heat Treatment: ST + PWHT
Figures 80 (a to d): Optical micmgraphs of the bondline region when ST heat treatment procedure is followed by PWHT
were imaged using a Video camem. The temperature ioside the fumace was graduaily
raised to 870°C (this was the maximum hanace temperature that could be used safély). This
temperature was marginaüy higher than the stabilisation temperature normaliy employed for
Ni-base wrought alloy base matenal. Figure 81 shows the r d t s produced when examining
as-welded joints made using solution-treated Ni-base wrought alloy base material. During
testing, the sample was heated for approx. 15 - 20 minutes and the ha1 microstructure was
photographed (see Figure 82). Massive nucleation of gamma prime occureci at 870°C and
since this temperature was higher than the stabilisation temperature for Ni-base wrought
alloy, base materid coagulation of the y' phase was apparent late in the heating cycle.
A similar approach was employed when examining as-welded fiction joints
produced using hardened Ni-base wrought alloy base material. The as-welded joint
microstructure is shown in Figure 83. The test specimen was held at 875°C for approx. 5
minutes. Since this holding temperature is higher than the stabilization temperature (860C)
normally used for Ni-base m u g h t alloy, considerable precipitation of y' was observed (see
Figure 84). Also it is worth noting that the jet-thinning process creates minute holes in the
specimen disc. When the test sample is heated these tiny holes expand considerably and
appear in the h a 1 microstnicture as small voids (see Figure 84).
4.3.3. Hardness Profiles
Figures 85 a to c show the hardness profiles in welds produced using different
welding procedures. The circular symbols refer to the use of solution treated Ni-base
wrought alloy base matenai prior to fiction welding; the square symbol indicate the use of
hardened Ni-base wrought alloy base materiai. The influence of using solution treated and
hardened Ni-base wrought alloy base materials prior to ection welding were examined:
a) Solution- Treated Ni-base wrought aihy
When a PWHT (stabiLization + precipitation) treatment was applied following fiction
welding, a hardened zone was fomed in Ni-base wrought ailoy base material adjacent to the
bondline. [Re-solution + P m ] andk [ST + PWHT] resulted in more uniforni hardness
Figure 8 1 : TEM micrograph h m the dynamically recrystailised region (solution-treated Ni- base wrought alloy base materiai prior to fiction welding)
y' precipitates
Figure 82: TEM micrograph h m the dynamicdy recrystallised region (solution-treated Ni- base wrought alioy base material prior to fiction welding) following in-situ heat treatment
Figure 83: Dynamicw tecrystallised region (hardened Ni-base wrought alloy base materials prior to fiction welding)
voids
y' particles
Figure 84: TEM micrograph fbn the dynamicaliy recrystallised region (hardened Ni-base wrought alloy base material prior to fiction welding) following in-situ heat treatment
profiles (see Figures 85 b and 85 c). The base metal hardness values were higher (approx.
4SOHv) when a PWHT treatment was applied. [PS + PWHT] treatment produced lower
hardness values (approx. 42SHv) in Ni-base wrought alloy base material. The most unifom
hardness profiles were obtained when an CRS + PWHT] thermal cycle was applied.
b) Hardened Ni-base woughf alloy
Hardened regions @eak batdlless: approx. 5OOHv) were formed in material adjacent
to the bondline when the fiction welded joints were heat-treated using a PWHT procedure.
When the fiction weids were heat-treated using ST or RS treatments prior to PWHT, the
hardness profile across the weld zone was uniform.
The typical aging (hardness) curve is a hc t ion of temperature (see Figure 86). The
significant feature is that extending the holding time and aging the test specimen too long at
a given temperature will decrease the sample hardness. It c a . be seen h m Figure 86 that
holding the superalloy at temperame Tl may attain a saturation value. However, holding at
temperature T2 wili produce a parabolic hardness relation. In-situ TEM micrographs
showed that the gamma prime particles had an extreme tendency for coarsening (over-
aging). Similar effects have been observed in creep samples of Ni-base cast alloy [Ahmet,
19941.
Temperature Tl corresponds with PWHT heat-treatment while T2 corresponds with
an RS post weld heaî treatmemt. Numerous tiny particles of gamma prime were precipitated
in the dynamically recrystallised region of Ni-base wrought alloy fiction welds following
the PWHT treatment. These particles were uniformly distributed and had an average
diameter of 0.02 microns. Similar precipitation behavior is likely in the Ni-base wrought
alloy base materiai. This would explain the peak hardness of 550 Hv observed in the
dynamically recrystallised zone and also the high hardness values in the base material
(approx. 435HV) foiiowing PWHT heat treatment.
Generally, an RS heat - treatment prior to PWHT (stabihtion + precipitation) is
quite uncornmon for Ni-base wrought d o y base matenal. However, the fiction welding
operation generates residual stress and non-uniform microstructures in the HAZ regions on
(a) After PWHT Treatment
4 -2 p O 2 4
Distance from the bondline (mm)
4 -2 Aje O 2
Distance from the bondline
(b) Mer RS + PWHT Treatment
4 -2 Ap O 2 4
Dîstrnca fmm the bondlino (mm)
(c) After ST + PWIfT treatment
Figure 85 a to c: Hardness profiles produced using various heat treatment cycles.
Figure 86: Schematic diagram showing variation in hardness with aging t h e and temperature.
either side of the bondline. Hence, the RS heat-treatment prior to PWHT is recommended to
relieve residuai stresses and aiso homogenise the weld microstructure. As seen in the TEM
"mgrapghs (see Figure 77), the particles in the dynamically recrystallised region were
0.05 pm in diameter following + PWHT] heat-treatment; this compares with a particle
diameter of 0.02 microns diameter foilowing the PWHT heat-treatment. Consequently, the
RS heat-treatment promotes over-aging and this reduces the haTdlless of the Ni-base
wrought alioy base material. This may explain why an overall reduction in hardness was
observeci in Ni-base wrought aiioy base material following the [PS+PWlWJ heat-treaîment.
However, the ST heat treatment was carried out at high temperature (1090C) for a long
duration (4 hours). This heat-treatment dissolves almost al1 the particles in the base
material. PWHT following the ST heat-treatment has the same effect as the PWHT
treatment with regard to the hardness profile. However, a ST heat-treatment pnor to PWHT
completely eliminates the residual stresses in the welded joint and HAZ regions and even
out hardness variations. In addition, extreme grain growth was observed following the ST
treatment (see Figure 80) and this will have a considerable infiuence on joint mechanical
properties. The different precipitation effects are shown schematicaily in Figure 87.
4.3.4. Tensile Strength Properties
Figure 88 shows typical stress-strain c w e s for Ni-base wrought dloy fiction welds
produced using different post weld heat treatment procedures. In al1 cases, the ultirnate
tensile strength was the fkacture strength of the test sample. The yield strength was the
stress corresponding to a strain of 0.02% and the ductility was the total elongation measured
using the stress- strain cuwe.
Figure 89 compares the yield and ultirnate tende strengths of joints made using
solution treated Ni-base wrought alloy base material while Figure 90 shows the
corresponding ductility values. Figure 9 1 compares the strengths of welded hardened Ni-
base wrought aiioy base material and Figure 92 shows the ductility values.
Foilowing PWHT Following RS+PWHT
Figure 87: Schematic illustrating different precipitation effects.
Figures 88: Typical Stress-Strain c m e s for w+PWHT] Ni-base wrought welds (solution -treated base materiai prior to fiction welding).
alloy fiiction
PWHT R S + M ST +PWHT
Varying heat treatment routes
ml- Yield Strength (MPa)
a Ultimate Tensile Strength (MPa)
Figure 89: Tensile strength properties (yield strength and ultimate tensile strength) of solution-treated Ni-base wrought ailoy fiction welds given different heat treatments.
br i s Mebl : M e c e i v e d Ni48rs wrought dloy
30
PWHT RS+PWHT ST+PWHT
Varying hart treaûnent mutes
Figure 90: Tensile Ductility (total elongation) of solution-treated Ni-base wrought alloy fiction welds following given different heat treatments.
Varying heat treatment routes
Yield Stieagth @Ga)
a t c Tensile Strength (MPa)
Figure 91 : Tensile strength pmperties (yield strength and ultimate tende strength) of hardened Ni-base wrought alloy fiction welds given different heat treatments.
Ba= Metal : Hardoned N i b r - wrought rlloy
PWHT RS + PWHT ST + PWHT
Figure 92: Ductility (total elongation) of hardened Ni-base wrought ailoy fiction welds given different heat treatments.
PWHT treatment produced the poorest weld strength properties (YS: 930 MPa, UTS:
1300 MPa) when solution- treated Ni-base wrought alloy base material was Ection welded.
The highest tensile strengths (YS: 925 MPa, UTS: 1380 MPa) were produced when a re-
solution treatment was applied prior to PWHT. In al1 the cases tende sample failure
occurred in Ni-base wrought dey base metal away h m the weld zone at appreciable
ductility values (elongations X5%) in ali joints. Similar tensile strengths and ductility
results were produced when hardeneci Ni-base wrought ailoy base materiai was employed
prior to welding. Re-solution prior to the stabilization plus precipitation treatment produced
the highest tensile strength properties (YS: 1080 MPa, WTS: 1428 MPa). However, Iower
strengths (YS: 870 MPa, UTS: 12ûûMPa) were found when the fiction welds were heat-
treated using an [ST + PWHT] thermal cycle.
The variation in tende strength properties can be explained via particle - dislocation
interactions. Dislocations are generated during tensile testing and an increase in tensile
strength is synonymous with incfea~ed difficulties in dislocation movement. A moving
dislocation must cut thorough the particles in its path or it must rnove around them.
Dislocations c m move around large diameter particles (Ormwan looping). However,
uniformly distributed srnall particles are cut as the dislocation moves. PWHT heat-treatment
prornoted the precipitation of small particles while the RS + PWHT treatment favored the
formation of both coarse and small particles. Further, RS + PWHT heat-treatment produced
large volume fraction of y' precipitates and therefore produced a matrix haviog high tensile
strength.
In general, solution treated Ni-base wrought alloy base material produced the highest
ductility values while the hardened Ni-base wrought alloy produced highest joint swengths.
In al1 the cases, sample failure occurred away h m the bondline in the adjoining base metal.
This is an indication that the variations in strength and ductility values were largely
determined by the Ni-base wrought alloy base matenal properties. For example, hardened
Ni-base wrought alloy base material containecl large volume fiactions of y' precipitates and
therefore had the highest strength values. High precipitate volume h t i o n in the matrix
decreases ductility.
In conclusion, application of a [RS+ PWHT] treatment produced the optimum
balance of strength and ductiiity properties and in all cases, tensile test specimen failure
occurred in base metal away h m the weld zone.
Typical fiacture surfaces are shown in Figures 93 a to f. The hcture mode was
intergranular since grain boundary regions in Ni- based superalloys regions are weak
compared to the grain interiors (since precipitation strengthens the grain interiors).
Consequently, failure is therefore more likely at grain boundary regions and therefore the
intergranular failure mode is more likely. It is worth noting that hardened Ni-base wrought
alloy base material contained large number of y' particles and carbides. Also, the hardening
treatment promoted extensive carbide precipitation dong grain bomdaries, see Figure 43.
An XEDS pattern h m a grain boundary carbide confimed that these carbides were rich in
Cr. This means that a depleted of Cr layer was created adjacent to the grain boundary
carbides. This depleted zone will produce a thin weak layer in which cracks can easily
nucleate and pmpagate during temile loading- This may explain why secondary cracks were
often observed in the failed hardened Ni-base wrought alloy tensile test samples (see Figures
93 a to 0.
4.4. Tensile Streugth Properties of Base Material
Friction welded joints comprise different microstructural regions, e.g., the region
containhg dynamically recrystallised grains, the heat afTected zone and the as-received base
material. During tensile loading for a given strain rate, these microstructural regions
respond differently. Weaker matenal preferentially necks and tensile failure occurs. In Ni-
base wrought alloy fnction welds the sample failure occurred wholly in the base matenal
indicating that it had the lowest tensile strength. In effect, the tensile strength of Ni-base
base material has a significant infiuence on mechanical properties of the Ni-base wrought
alloy fiction welds.
Ni-base wrought alloy base metai was subjected to different heat - treatment routes;
P m , RS+PWHT and ST + P m . Figures 94 and 95 compare the yield and ultimate
(a) Prior Base Metal Condition: Solution -Treated Ni-base wrought aiioy; Heat Treatment: PWHT
@) Prior Base Metal Condition: Hardened Ni-base wrought aiioy; Heat Treatment: PWHT
103
(c) Base Metal Condition: Solution -Treaîed Ni-1 base wrought doy; Heat Treatment: RS + PWHT
(d) Base Metal Condition: Hardened Ni-base wrc ~ught aiioy; Heat Treatment: RS + PWHT
Figure 93 a-E Fracture d a c e s of broken temile test samples extracted h m heat-treated Ni-base wought alloyMi-base mught dey welds
tensile strengths of solution treated and hardened Ni-base wrought alloy base materials. The
tende strengths of solution-treated Ni-base wrought alloy base material were YS: 590 MPa
and UTS: 970 MPa; these wrnpared with those in hardeneci Ni-base wrought alioy base
material (YS: 740 MPa and UTS: 1140 MPa). The highest tende strength values (YS: 762
MPa, UTS: 1136 MPa) were produced when PWHT hûit-treatment was carrieci out using
solution treated Ni-base wrought alloy base material. The tensile strength values following
the w+PWHT] heat-treatrnent were simila. to the tensile strengths of Ni-base wrought
alloy base material subjected to a PWHT heat treatment cycle: the values were YS: 700 MPa
and UTS: 1090 MPa for the solution treated Ni-base wrought alloy base material and YS:
760 MPa and UTS: 1 153 MPa for the hardend Ni-base wrought alloy base material.
However, the [ST+PWH'l'] heat-treatment produced the lowest tende strength properties; in
solution -treated Ni-base wrought alloy base material the vaiues were YS: 560 MPa, UTS:
864 MPa; in hardened Ni-base wrought d o y base matenal the values were YS: 600, UTS:
93 2MPa receptively .
PWHT heat-treatment pmmoted the precipitation of srnail (0.02 micron diameter) y'
particles while the CRS + PWHT] treatment favored the formation of both coarse (0.05
micron) and small diameter particles (0.01 micron) in the Ni-base wrought alloy base
material. Both these heat- treatments increased the tensile strength of the base material. An
ST heat-treatment pnor to PWHT dissolved ail the y' precipitates in the matrix and the
PWHT treatment promoted fiutber precipitation. However, it is possible that low volume
hctions of y' following [ST+PWHT] heat-treatment led to the lowest tensile strength
values: YS: S60MPa and UTS: 864 MPa for solution - treated Ni-base wrought ailoy base
materiai and YS: 600 MPa and UTS: 932 MPa for the hardened Ni-base wrought alloy base
matenal.
In general, solution treated Ni-base wrought alloy base matenal produced the highest
ductility values (>20%) while the hardened Ni-base wrought alloy produced highest
strengths (11 100MPa). In al1 the cases, specimen failure occurred close to the center of the
gauge length. This is au indication that the strain rate dong the gauge length was essentially
uniform. Hardened Ni-base wrought alloy base matenal had lowex ductility in temis of total
elongation values (approx. 8-100/0 lower) compareci to solution-treated Ni-base wmught
alloy base material.
ST + PWHT heat- tmatment improved the ductility @y 20-27%) in compatison the
to P m and = + P m heat-treatments. Extreme grain growth was observed when the
ST heat- treatment was applied. Moreova, carbides were completely dissolveci during the
heat-treatment cycle resulting in a matrix containing a lower volume h t i o n of the y' and
carbide precipitates. This explains high ductility and the low teasile strengths following the
ST+PWHT heat-treatrnent.
Yield Strength (MPa) m-+ t e Tensile Strength m a )
Figure 94: Tensile strength properties (yield strengh and ultimate tensile strength) of as- received Ni-base wrought alloy base material given different heat treatments.
Yield Strength (MPa) m-+ Ultimate Tensile Strength @Pa)
Figure 95: Tensile strength properties (yield strength and ultimate tensile strength) of hardened Ni-base wrought alloy base material given different heat treatments
CHAPTER 5
RESULTS AND DISCUSSION
Ni-BASE CAST ALLOYINI-BASE WROUGHT ALLOY
FRICTION W L D S
The results found during Ni-base wrought alloy/Ni-base wrought alloy friction
welding served as the basis for the selection of the welding parameters applied during
dissimilar Ni-base cast aiioyNi-base wrought alloy welding. Ni-base cast alloy material
was fiction welded to Ni-base wrought aUoy in the hardend condition and foilowing
welding operation the dissimilar joints were post weld heat - treated. It was previously
found that the optimum mechanical properties (ductility >15% and UTSX 300 MPa) were
produced when a re-solution + PWHT (stabilization plus precipitation) heat treatment was
applied to Ni-base wrought ailoy fiction welds. For this reason, it was recommended that a
RS + PWHT treatment should be applied in the case of dissimilar Ni-base cast alioy/Ni-base
wrought alloy fiction welds. Heat-treated welds were investigated using transmission
electron microscopy.
5.1 Ni-base Cast Ailoy Base Metal Microstructure
Ni-base cast alloy base material is normally available in the hardened condition, i.e.,
the alloy is subjected to a stabilization + precipitation heat treatment. Figure 96 shows the
typicai cast base metal microstructure and coÏnprises of incoherent and coherent y'particles
contained in a y matrix together with MC and M& type carbides. The mismatch between y'
precipitates and the y matriv increases when the volume fiaction of y' in the base material
increases. The shape of the y' particles changes from spherical to cubical when the volume
fiaction of y' in the y matrix exceeds 35 vol.%. In Ni-base wrought alloy base material the
y' particles were spherical; however, they were cuboidal shaped in Ni-base cast alloy base
1 Cast Sb
Figure 96: SEM micrograph of Ni-base cast d o y base mai
Figure 97: Magnified view of Figure 96
110
material (see Figure 97). Further, spherical y' particles have the lowest mismatch while the
cubical y' particles have the highest intemal strah (+ 0.0200% to - 0.0100%)). In spite of
this mismatch, base material containhg cubical y' particles exhibits excellent ductility, a
unique behavior. The inherent ductility of the y' particles prevents severe embrittlement,
uniike the strengthening that is developed by phases which have much higher hafdIless, cg.
alumina in an Al matrix.
Figure 98 shows a TEM mimgraph of Ni-base cast alloy base material, The
structure comprised u n i f o d y distributecl cubical y' particles dong with coherent sphencal
y' particles in a y matrix. The average size of cubic particles was approx. 0.5 microns; the
spherical particles were appox.0.05 microns in diameter. The chemistry of y' precipitates
was examined using X-ray Energy Dispersive Spectrometer (XEDS). High resolution
transmission electron microscopy (HRTEM) was used since the output was more accurate
and reliable. Figure 99 shows an XEDS pattern produced by the cubical particles. The
cubical particles were essentialiy an allay of Ni containhg Al, Ti and W. However, trace
contents of Cr and Co were also indicated. The XEDS pattern h m a coherent sphencal y'
particie is shown in Figure 100. These particles had the same chemistry as previously, but
were high in W and comprised maitily Ti, Al and Ni. These particles corresponded with the
formulation Ni3(Cr,Co)AlTiW.
5.2. Welding Parameters and Post Weld Heat Treatment
The initial friction welding tests were carried out by varying the fiction pressure
nom 275 MPa to 375MPa when using a fiction time of 105, a forging pressure of 350 MPa,
a forging t h e of 2s and a rotational speed of 1000 rpm. The Ni-base wrought alloy
substrate was severely deformed and Ni-base cast d o y base material was largely unaffected
(see Figure 101). However, the final weld interface profile was not perpendicular to the axis
of rotation of the test samples. The high temperature flow stress of Ni-base wrought alloy
base material is lower that that of Ni-base cast alioy base material and this is why the Ni-
base wrought alioy substrate deformed prefmtially and constitutes the principal
component in the flash. During friction welding the contact surface stresses are the highest
Figure 98: TEM micrograph of Ni-base cast d o y base material
Cliant : Sugl Smt t i l rn Job : Job numbœr 387 Spectrum 8 (8/5199 10:32>
Figure 99: XEDS pattern fiom a cubical 7' particle in Ni-base cast ailoy base material
Figure 100: XEDS pattern h m a sphericd y' particle in Ni-base cast alloy base material
Figure 101 : Joint interface in Ni-base cast ailoy/Ni-base wrought alloy fiction joint
at the weld centerline Worth, 19921 and the accumulation of these stretched elements
makes a major contriiution to the flash volume during welding. If deformation of the Ni-
base wrought alloy substrate was the only faetor controlling the shape of the Ni-base
wrought alloy/Ni-base cast aUoy joint interface; it would be curved pnor to the forging stage
in the fiction welding operation. However, since higher temperatures are produceci at the
outer periphery of the joint and this promotes preferential plastic flow in this location so that
the final joint interfce profile is that shown in Figure 101.
Figure 102 shows the as-welded microstructure at the joint centerline. An intennixed
zone was observed adjacent to the bondline. This is the region where elemental diaision
and mechanical mixing occurs as a resuit of the thermal and sWst ra in rate cycle in
friction welding. The width of this intermixeci zone depended on the welding parameters
selected, particularly on the fiction pressure selected. As expected, welds produced using
high friction pressures (375 MPa)-produced wider inter-mixed zones.
Al1 dissimilar Ni-base wrought alloy/Ni-base cast alloy joints were heat-treated using
a re-solution + PWHT involving a stabilization plus precipitation thermal cycle
correspondhg with that for Ni-base wrought alloy base material. Following heat treatment
cracks were observed at the bondline of welds produced using low fiction pressures (( 325
MPa). Figure 1 O3 shows micrographs of weided joints produced using a low fkiction
pressure (Wction pressure = 300 MPa, Wction tirne = 10s) and an intermediate fiction
pressure of 325 MPa (fiction tirne: 10s). Crack fiee dissimilar welds were produced using
friction pressure exceeding 325 MPa.
TEM discs were extracted h m the weld centerline. It is worth noting that Ni-base
cast alloy and Ni-base wrought alloy materials behaved differently while electrolytic
thinnllig process. The Ni-base wrought alloy substrate was thinned preferentially while Ni-
base cast alloy substrate stayed as it was. Thus, the transmission micrograph fiom a
dissimilar weld was of barely acceptable qualïty. For this reasoq a modified TEM
procedure will be required in friture* which will effectvely characterize dissimilar Ni-base
cast alloy/Ni-base wrought alloy welds. A TEM micrograph of the weld centerline
Figure 102: As-Welded microstructure of Ni-base cast aiioyMi-base wrought aiioy dissimilar fiction weld
Figure 103: Crack formation following post weld heat treatment of a Ni-base cast alloyMi- base wrought aiioy fiction weld
following re-solution + P m (stabiiization plus precipitation) is shown in Figure 104. The
location of this sample is close to Ni-base cast alloy substrate. As noted earlier, in Ni-base
cast alloy contains cohererit 0.05 pm diameter spherical particles and 0.5 pm diameter non-
coherent cuboidal y' particles in a y matrix. FoUowing fiction welding considerable
amounts of elemental a i o n take place across the weld interface (into the Ni-base cast
alloy componeni). It is worth noting that Ni-base wrought alloy base material is less rich in
Al, Ti and W than the Ni-base cast alloy alloy. Thus, diaision fiom Ni-base wrought alloy
material into the Ni-base cast alloys substrate w i U dilue the chemistry of the Ni-base cast
alloy alloy. As a result high temperature diffusion during the fiction welding operation
might result in partial dissolution of cuboidal y' particles and dissolve the smaller 0.3 pm
diameter spherical particles in the Ni-base cast dloy base matenal. Since the spherical
gamma prime particles are coherent with the gamma matrix they may be easily dissolved.
The XEDS pattern fkom a spherical type y' particle is shown in Figure 105. The
particle composition corresponds to Ni3 (Cr, Co) AlTiW. This 0.3 pm diameter sphencal y'
particle has approximately the same chemistry as that of the cuboidal particles (see figure
101). Figure 1 O6 shows an XEDS pattern h m the y rnatrix. As expected, the ma&
comprised an alloy of Ni, Cr and Co. However, smaller amounts of Ti, Al and W were aiso
indicated. Also, the matrix contains higher amounts of Cr than the y' particles.
It can be concluded that the application of a higher fiction pressure (>325 MPa) is a
prime requirement during Ni-base cast ailoymi-base wrought alloy fiction joining and
produces crack-fiee welds.
5.3. Hardness Profiles
Figure 107 shows the micro-hardness results produced in a traverse h m the bondline
into the as-received Ni-base wrought alloy base metal. The indentation load was 500 g and
25 readings were taken on each test weld. When solution-treated Ni-base wrought alloy
base material was fiction welded, a softened zone was produced in the Ni-base wrought
alloy substrate and had a hardness about 25 -50 H v lower than the surromding base
Figure 104: TEM micrograph of the bondhe region in a Ni-base cast alloy/Ni-base wrought ailoy friction weld following RS (re-sdution) + PWHT (stabiiization plus precipitation) heat treatment
Clknt : S u g I S m t h l r n Job : Job numhr 367 spmctrurn tS CWS/99 1 1 :46>
Figure 105: XEDS spectra h m a spherical y' particle (fiom Ni-base cast alloy/Ni-base wrought alloy fiction welds)
Figure 106: XEDS pattern ftom the y matrix (Ni-base cast alloyMi-base wrought d o y weld)
Nilase Cast Ailoy & { Ni-base Wrought Alloy
A 1 1
: O t
: 0. O C 0
O 0
E x
1
4 -3 -2 -1 ,#' O 1 2 3 4 5
Distance ftom the bondline (mm)
Figure 107: Hardness profile in Ni-base cast alloy/Ni-base wrought alloy fiction welds
A- As- welded condition B. Mer Re-solution + PWHT
4
r
The measurements were carrieci out at the center (see arrow)
Welding Condition :
Friction Pressure : 350 MPa Friction Time : 10s Forging Pressure : 350 MPa Forging Time : 2s Bar Diameter : 19-
material. This soAened zone r d t e d hm the solution and coarsening (over-aging) of
intemetallic phases during the thermal cycle in fiction welding. These softened zone
regions were completely removed following re-solution + PWHT (involving a stabilization
plus precipitation treatment).
CHAPTER 6
CONCLUSIONS 6.1. Ni-base Wrought AIIoy Friction Wel&
The optimum weldhg parameters were: Friction Pressure: 350 MPa, Friction Time:
los, Forging Pressure: 350 MPa, Forgïng Time: 2s, Rotationai Speed: 1000 rpm.
Highest ultimate tende strengths (UTS values: 1475 MPa) with appreciable ductility
levels (>15%) were obtained when Ni-base wrought alloy base material was welded
using an intermediate friction pressures (325 to 375 MPa). In al1 the cases the tensile
test specimens failure occwed in base metal away fkom the weld zone. In welds
produced using at higher fiction pressures, the hardness profile across the weld zone
was essentially unifonn.
Hardened regions were produced in matenal adjacent to the bondline when the
welded samples are directly heat-treated using a PWHT (Stabilization + Precipitation) procedure. When Wction welds are heat-treated using ST or RS
treatments prior to PWHT, the hardness profile across the weld zone is uniform.
Application of a [RS+ PWHT] heat-treatment produced the optimum balance of
strength and ductility properties. In d l the cases tensile test failure occurred in base
metal away fiom the weld zone.
Solution - treated Ni-base wrought alloy base matenal produced the highest ductility
vales (>20%) wwhi the hardened Ni-base wrought alloy produced the highest
strengths (>Il00 MPa). Hardened Ni-base wrought alloy base material had lower
ductility in texms of total elongation values (approximately 8-10% lower) compared
to solution - treated Ni-base wrought allay base material. An ST+PWHT heat - treatment improved the ductility @y 20-27%) in cornparison to alternative PWHT
and RS+P WHT heat-treatments.
5. Hardened Ni-base wrought alloy base material exhibiteci extensive Cr carbide
precipitation dong the grain boundaries. Coarse and elongated carbides having
aspect ratios > 15 were highly detrimental vis-à-vis mechanical properties since they
can acted as a stress raisers and promoted secondary crack propagation.
6. The base materials tensile strengths were @? WHT: 1 137 MPa; RS+PWHT: 1 O87
MPa; ST+PWHT: 862MPaI and compared with fiction welded joints PWHT: 1302
MPa, RS+PWHT: 1380 MPa, ST+PWHT: 1323 MPa]. The Ni-base wrought
alloy/Ni-base wrought ailoy joint strengths were 20-25% higher than the as-received
base matenal. Ail sarnple failures occurred in the adjohhg base material
6.2. Ni-base Cast AlloyNi-buse Wrought AlZoy Friction Wei&
1. The application of higher fiction pressures (>325MPa) is the prime requirements
during Ni-base cast alloy/Ni-base wrought alloy dissimilar fiction joining and that
this welding parameter combination produces crack-fkee welds following post weld
heat-treatment.
2. The fiction welding operation altered the original particle shape and morphology at
the bondline. Close to the bondline region in the Ni-base cast d o y substrate side,
the cubical gamma prime was altered to spherical in shape.
FUTURE WORK
It is recommended that future work should emphasise modelling of the
microstructural changes produced in Ection welded joints, particularly with regard to
changes in the particle distribution when différent p s t weld heat treatment procedures are
applied. A general mode1 correlating and predicting microstructural variations and their
effects on tensile strength properties is required for different dissimilar superalloy material
combinations.
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[ASM Haudbook, 19721
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