MICROSTRUCTURE ? MECHANICAL PROPERTY RELATIONSHIPS IN
TRANSIENT LIQUID PHASE BONDED NICKEL-BASED
SUPERALLOYS AND IRON-BASED ODS ALLOYS
Except where reference is made to the work of others, the work described in this
dissertation is my own or was done in collaboration with my advisory
committee. This dissertation does not include proprietary, restricted
or classified information.
______________________________
Sreenivasa Charan Rajeev Aluru
Certificate of Approval:
_________________________ _________________________
Jeffery W. Fergus William F. Gale, Chair
Associate Professor Professor
Mechanical Engineering Mechanical Engineering
_________________________ _________________________
Barton C. Prorok Winfred Foster
Assistant Professor Professor
Mechanical Engineering Aerospace Engineering
_________________________ _________________________
Pradeep Lall Stephen L. McFarland
Associate Professor Dean
Aerospace Engineering Graduate School
MICROSTRUCTURE ? MECHANICAL PROPERTY RELATIONSHIPS IN
TRANSIENT LIQUID PHASE BONDED NICKEL-BASED
SUPERALLOYS AND IRON-BASED ODS ALLOYS
Sreenivasa Charan Rajeev Aluru
A Dissertation
Submitted to
the Graduate Faculty of
Auburn University
in Partial Fulfillment of the
Requirements for the
Degree of
Doctor of Philosophy
Auburn, Alabama
May 11, 2006
iii
MICROSTRUCTURE ? MECHANICAL PROPERTY RELATIONSHIPS IN
TRANSIENT LIQUID PHASE BONDED NICKEL-BASED
SUPERALLOYS AND IRON-BASED ODS ALLOYS
Sreenivasa Charan Rajeev Aluru
Permission is granted to Auburn University to make copies of this dissertation at its
discretion, upon the request of individuals or institutions and at their expense.
The author reserves all publication rights.
_________________________
Signature of Author
_________________________
Date of Graduation
iv
VITA
Sreenivasa Charan Rajeev Aluru, son of Venkata Ramana Rao and Jhansi Rani
Aluru, was born on June 17
th
, 1980 in Nellore, India. He graduated from St. Xavier?s
High School, Ongole, India in March 1995 and Intermediate school from Adarsha Junior
College, Ongole in March 1997. He joined Nagarjuna University, Guntur, India in
August 1997 and graduated with a Bachelor of Technology in Mechanical Engineering in
May 2001. He entered Auburn University in January 2002 as a graduate student in
Mechanical Engineering and graduated with a Master of Mechanical Engineering degree
in December 2004. He started working towards the Doctoral program in Materials
Engineering at Auburn University in August 2002.
v
DISSERTATION ABSTRACT
MICROSTRUCTURE ? MECHANICAL PROPERTY RELATIONSHIPS IN
TRANSIENT LIQUID PHASE BONDED NICKEL-BASED
SUPERALLOYS AND IRON-BASED ODS ALLOYS
Sreenivasa Charan Rajeev Aluru
Doctor of Philosophy, May 11th, 2006
(M.M.E, Auburn University, 2004)
(B.Tech, Nagarjuna University, 2001)
150 Typed Pages
Directed by Dr. William F. Gale
The research work presented here discusses the microstructure-mechanical
property relationships in wide gap transient liquid phase (TLP) bonds, between the single
crystal nickel-base superalloy CMSX-4 and two polycrystalline superalloys, IN 738 and
IN 939, using wide-gap style composite interlayers. Fabrication of complicated
geometries and successful repair development of gas turbine engine components made of
superalloys requires a high performance metallurgical joining technique and a complete
understanding of microstructure-mechanical property relationships. A number of joining
vi
processes have been investigated, but all of them have significant disadvantages that
limit their ability to produce sound joints. TLP bonding has proved to be a successful
method and is the most preferred joining method for nickel-based superalloys, with
microstructures and compositions of the joint similar to that of the bulk substrates
resulting in mechanical properties close to that of the parent metal.
The current joining process used two proprietary wide-gap style composite
interlayers, Niflex-110 and Niflex-115, consisting of a nickel-based core with boron-rich
surfaces, and a conventional rapidly solidified metallic glass foil interlayer BNi-3 was
chosen for comparison. When composite interlayers were employed, competition
between wetting of the faying surfaces and formation of the eutectic along the grain
boundaries was observed to lead to non-bonded regions at the faying surfaces, unless a
boron-rich interlayer was employed. Composite interlayers resulted in the suppression of
bondline boride formation. With the exception of this competition, adequate wetting of
the substrates occurred for all interlayers.
Two factors dominated the room temperature mechanical properties of the wide-
gap bonds. The first was the extent of gamma-prime formation at the bondline. Results
from shear testing and fractography of the bonds indicated ductile shear failure at the
bondline. This was due to the formation of insufficient gamma-prime within the joint,
which left a relatively soft bondline region. The second factor was the presence of second
phases in the diffusion zone of the polycrystalline substrate. This led to the formation of
brittle secondary cracks. Overall, it is evident that the room temperature shear strength of
the bonds was more dependent on the extent of formation of ?? on the bondline than on
the secondary phases in the diffusion zone of the polycrystalline substrate.
vii
ACKNOWLEDGEMENTS
The author would like to express his heartfelt gratitude to Dr. William F. Gale for
his continued support, guidance and encouragement throughout this period of
investigation. The author would like to emphasize the extensive knowledge and genuine
concern for students in Dr. William F. Gale, which benefited him scientifically as well as
a person. Thanks are also due to Dr. Jeffrey W. Fergus for his invaluable advice and help
and the committee members for their useful suggestions.
The author is grateful to his parents and brother for their love, prayers, endless
support. Wholehearted thanks to Srilatha Punna for her support and encouragement,
Rajesh Guntupalli, Srinivas Sista, Shakib Morshed for their valuable discussions, Nanda
Ravala, Viswaprakash Nanduri, for their constant motivation and all other friends for
their invaluable friendship. The author would also like to express sincere thanks to all of
his colleagues in AU Physical Metallurgy & Materials Joining group for their assistance
and friendship.
Finally, the author would like to dedicate this dissertation to the lotus feet of his
beloved Lord Venkateswara and his parents (Mrs. Jhansi Rani and Mr. Venkata Ramana
Rao), without whose grace, love and forbearance it would not have been possible to learn
many things in science as well as in general aspects of life in the course of these years
and throughout my life.
viii
Style manual or journal used Metallurgical Transactions A
Computer software used Microsoft Office XP
ix
TABLE OF CONTENTS
LIST OF FIGURES???????..?????????.........................................xii
LIST OF TABLES...???????..?????????......................................xvii
LIST OF ACRONYMS????????????????????????..xviii
1. INTRODUCTION ...................................................................................................... 1
2. LITERATURE REVIEW ........................................................................................... 4
2.1 Origin and Development of Superalloys ...................................................4
2.2 Microstructures and Mechanical properties ..............................................6
2.3 Joining of nickel-based superalloys.........................................................12
2.3.1 Need for Joining Superalloys ...........................................................12
2.3.2 Fusion Welding ................................................................................13
2.3.3 Diffusion Bonding............................................................................15
2.3.4 Brazing .............................................................................................15
2.3.5 Transient Liquid Phase Bonding......................................................16
2.3.5.1 Advantages and Disadvantages of TLP bonding .........................20
2.3.6 Wide-gap Transient Liquid Phase Bonding .....................................21
2.4 Development of Oxide Dispersion Strengthened Superalloys ................22
2.5 Mechanical Alloying ...............................................................................23
2.6 Properties and Applications.....................................................................24
2.7 Joining of ferritic based Superalloys .......................................................25
2.7.1 Conventional Techniques and Limitations.......................................25
2.7.2 Diffusion Bonding............................................................................27
2.7.3 Transient Liquid Phase bonding of ferritic based ODS Alloys .......29
3. RESEARCH OBJECTIVES ..................................................................................... 30
4. MATERIALS AND EXPERIMENTAL PROCEDURE.......................................... 34
4.1 Nickel-based superalloys.........................................................................34
4.1.1 Materials...........................................................................................34
4.1.2 Joining Procedure.............................................................................37
4.1.3 Post Bond Heat Treatment (PBHT) .................................................37
4.1.4 Post Bond Thermal Exposure (PBTE) .............................................38
4.1.5 Metallographic Preparation..............................................................40
4.1.6 Microstructural Characterization .....................................................40
x
4.1.7 Mechanical Testing ..........................................................................40
4.1.7.1 Shear Testing ................................................................................41
4.1.7.2 Hardness Tests..............................................................................44
4.1.8 Wettability Studies ...........................................................................44
4.2 Oxide Dispersion Strengthened Iron Based Superalloys.........................45
4.2.1 Materials...........................................................................................45
4.2.2 Joining Procedure.............................................................................47
4.2.3 Post Bond Heat Treatments..............................................................47
4.2.4 Metallographic Preparation..............................................................47
4.2.5 Microstructural Characterization .....................................................48
4.2.6 Oxidation Studies .............................................................................48
4.2.7 Mechanical Testing ..........................................................................49
5. RESULTS AND DISCUSSION............................................................................... 50
5.1 Nickel-based superalloys.........................................................................50
5.1.1 Microstructural Characterization .....................................................50
5.1.1.1 Porosity at the bondline ................................................................50
5.1.1.2 Bondline Boride Formation..........................................................58
5.1.1.3 Microstructural Bond Development.............................................58
5.1.1.4 Secondary Phases in the diffusion zone .......................................61
5.1.2 Comparison with the wettability studies..........................................70
5.1.2.1 Effect of substrate on wettability..................................................71
5.1.2.2 Effect of Boron Content ...............................................................80
5.1.2.3 Effect of Boride formers...............................................................80
5.1.3 Gamma-prime at the bondline..........................................................81
5.1.4 Structure-Property Relationships of As-bonded TLP bonds............84
5.1.4.1 Shear tests.....................................................................................84
5.1.4.2 Hardness testing............................................................................86
5.1.5 Characterization of Post-Bond Heat Treated TLP bonds.................94
5.1.5.1 Microstructure after PBHT...........................................................94
5.1.5.2 Microstructure after PBTE ...........................................................94
5.1.6 Structure-Property Relationships of TLP Bonds following a
PBHT and PBTE ............................................................................101
5.1.6.1 Shear tests...................................................................................101
5.2 Oxide dispersion strengthened iron based superalloys..........................103
5.2.1 Microstructural Characterization ...................................................103
5.2.1.1 Diffusion bonding of MA 956....................................................103
5.2.1.2 Diffusion bonding of PM2000....................................................104
5.2.2 Structure-property relationships of ODS alloys.............................105
6. CONCLUSIONS..................................................................................................... 118
7. FUTURE WORK.................................................................................................... 121
xi
8. BIBLIOGRAPHY................................................................................................... 124
xii
LIST OF FIGURES
Figure 1: Schematic representation of the typical microstructure of nickel-base
superalloys [43]......................................................................................................... 10
Figure 2-The crystal lattice structures of (a) NiAl and (b) Ni
3
Al [44] ............................. 11
Figure 3 Stages of TLP bonding - Interlayer melting and substrate dissolution [25, 27]. 18
Figure 4 Stages of TLP bonding ....................................................................................... 19
Figure 5: Flow chart of project objectives and investigations for TLP bonding of
dissimilar nickel-based superalloys .......................................................................... 33
Figure 6 Microstructure of as-received single crystal CMSX-4 [97] ............................... 35
Figure 7: Schematic of wide gap composite interlayer..................................................... 36
Figure 8(a): Schematic of sample used for shear testing [114] ........................................ 43
Figure 9 - SEM micrographs in SEI mode, of CMSX-4 -Niflex-110 ? IN 939 joint
showing (a) grain boundary eutectic formation after 0 minutes at 1160
0
C and
(b) porosity on the substrate- interlayer interface in the initial stages after 4
minutes at 1160
0
C. The secondary phases in the diffusion zone of the
polycrystalline substrate are also shown................................................................... 52
Figure 10(a) - SEM micrographs in SEI mode, of CMSX-4 -Niflex-115 ? IN 738 joint
showing (a) grain boundary eutectic formation after 0 minutes at 1160
0
C and (b)
substrate- interlayer interface with no porosity after 60 minutes at 1160
0
C. The
secondary phases in the diffusion zone of the polycrystalline substrate are also
shown ........................................................................................................................ 53
Figure 11- SEM micrographs in SEI mode, of CMSX-4 -Niflex-115 ? IN 738 joint
showing (a) grain boundary eutectic formation after 0 minutes at 1160
0
C and (b)
substrate- interlayer interface with no porosity after 60 minutes at 1160
0
C. The
secondary phases in the diffusion zone of the polycrystalline substrate are also
shown. ....................................................................................................................... 56
Figure 12 ? (a)LM micrograph of CMSX-4 ? BNi-3 ? IN 738 after 0 minutes at
1160
0
C, showing borides at the bondline. No grain boundary eutectic formation
xiii
was observed. (b) SEM micrograph in BEI mode after 0 minutes at 1160
0
C,
showing borides at the bondline. The substrate- interlayer interface here is free
of non-bonded regions. ............................................................................................. 57
Figure 13: Comparison of Vickers microhardness across bondline of CMSX-4 ?
IN 939 bond after 240 minutes at 1160
0
C using Niflex-110 and BNi-3 foil
interlayers. Note the high hardness values for the BNi-3 joint, that might be
due to the borides present. ........................................................................................ 63
Figure 14: SEM micrographs in SEI mode, of CMSX-4 ? IN 738 bonds after 240
minutes at 1160
0
C showing (a) porosity at the bondline in Niflex-110 and
(b) bond free from porosity and free from secondary phases in the diffusion
zone, using Niflex-115 interlayer. ............................................................................ 64
Figure 15: SEM micrographs in SEI mode, of CMSX-4 ? BNi-3 - IN 738 bond
after 240 minutes at 1160
0
C, showing borides at the joint and secondary phases
in the diffusion zone of IN 738................................................................................. 65
Figure 16: SEM micrographs in SEI mode, of CMSX-4 ? IN 939 bond after 240
minutes at 1160
0
C showing (a) secondary phases in diffusion zone of
polycrystalline substrate in Niflex-115 interlayer bond, and (b) borides at the
joint and secondary phases in the diffusion zone of IN 939, for bonds using
BNi-3 interlayer. ....................................................................................................... 66
Figure 17: Composition profile, obtained using SEM-based EDS analysis, across
bondline of CMSX-4 ? Niflex-110 ? IN 738 joint after 4 minutes of bonding time
at 1160
0
C Note the chromium and cobalt peak observed at the diffusion zone of
the polycrystalline substrate that might be carbides, borides, TCP phases formed.. 67
Figure 18: Composition profile obtained using SEM-based EDS analysis, across
bondline of CMSX-4 ? Niflex-110 ? IN 738 joint after 240 minutes of bonding
time at 1160
0
C. Notice that the composition of the joint is fairly uniform after
allowing time for isothermal solidification............................................................... 68
Figure 19: Composition profile obtained using SEM-based EDS analysis, across
bondline of CMSX-4 ? Niflex-110 ? IN 939 joint after 240 minutes of bonding
time at 1160
0
C. Notice that the composition of the joint is fairly uniform after
allowing time for isothermal solidification............................................................... 69
Figure 20: LM micrograph, showing eutectic formation in the grain boundaries of
inner Niflex core for a Niflex-110 interlayer wetted on (a) CMSX-4 and
(b) IN 738. The secondary phases formed in the diffusion zone of the
polycrystalline substrate IN 738 are also visible [121]............................................. 73
xiv
Figure 21: LM micrograph, showing eutectic formation in grain boundaries of
inner Niflex core for Niflex-110 interlayer wetted on (a) CMSX-4 and
(b) IN 939. The secondary phases formed in the diffusion zone of the
polycrystalline substrate IN 939 are visible [121]. ................................................... 74
Figure 22: LM micrograph, of the cross-section of IN 939 wetted with BNi-3 foil
showing the formation of borides at the joint interface [121]. ................................. 75
Figure 23: Wettability of BNi-3 on CMSX-4, IN 738 and IN 939 CMSX-4(R). [121] ... 76
Figure 24: Wettability of Niflex-110 on CMSX-4, IN 738, IN 939 and CMSX-4(R).
Compare this to previous Figure and notice the reduction in time required for
termination of spreading compared to BNi-3 [121].................................................. 77
Figure 25: Wettability of Niflex-115 on CMSX-4, IN 738, IN 939 and CMSX-4(R).
Compare this to previous two Figures and notice the reduction in time required
for termination of spreading compared to BNi-3 [121]............................................ 78
Figure 26: Wettability of BNi-3 on CMSX-4 on two different runs in identical
conditions showing repeatability of data. [121]........................................................ 79
Figure 27: SEM micrographs in SEI mode, of CMSX-4 ? Niflex-110 ? IN 939 bond
showing (a) gamma-prime (??) at the bondline (b) Polycrystalline substrate
IN 939 of the bond with no resolvable (??)............................................................... 82
Figure 28: SEM micrographs in SEI mode, of CMSX-4 ? Niflex-110 ? IN 738 bond
showing (a) gamma-prime (??) at the bondline (b) Polycrystalline substrate
IN 738 of the bond with no resolvable (??)............................................................... 83
Figure 29: SEM micrograph in SEI mode, showing fracture surface of as-received
(a) IN 939 bulk material and (b) IN 738 bulk material. The fracture path was
macroscopically flat and dominated by ductile shear fracture with dimpling.......... 87
Figure 30 : Schematic showing crack path upon application of shear force on the
bonded samples......................................................................................................... 90
Figure 31: SEM micrograph in SEI mode, of shear tested CMSX-4 ? Niflex-110 ?
IN 939 bond fracture surface of (a) IN 939 polycrystalline substrate (b) CMSX-4
single crystal substrate. Notice the secondary cracking on IN 939 substrate due to
the secondary phases precipitated in the diffusion zone of the polycrystalline
substrate. ................................................................................................................... 91
Figure 32: SEM micrograph in SEI mode, of shear tested CMSX-4 ? Niflex-110 ?
IN 738 bond fracture surface of (a) IN 738 polycrystalline substrate
xv
(b) CMSX-4 single crystal substrate. Notice the ductile shear fracture on both
of the substrates and dimpling on IN 738. ................................................................ 92
Figure 33: Vickers microhardness across bondline of CMSX-4 ? IN 738 and
CMSX-4 ? IN 939 bonds after 240 minutes at 1160
0
C using Niflex-110
interlayer. Notice that the bondline is of lower microhardness than that of the
substrates................................................................................................................... 93
Figure 34: SEM micrograph in SEI mode, of As bonded CMSX-4 ? Niflex-110 ?
IN 738 showing gamma-prime (??) at the bondline.................................................. 96
Figure 35: SEM micrograph in SEI mode, of As bonded + PBHT CMSX-4 ?
Niflex-110 ? IN 738 showing gamma-prime (??) at the bondline. Notice the
increase of volume fraction of ?? after PBHT (compare it to previous Figure)........ 97
Figure 36: SEM micrograph in SEI mode, of As bonded CMSX-4 ? Niflex-110 ?
IN 939 bond showing gamma-prime (??) at the bondline......................................... 98
Figure 37: SEM micrograph in SEI mode, of as bonded + PBHT CMSX-4 ?
Niflex-110 ? IN 939 showing gamma-prime (??) at the bondline. Notice the
modification of shape of ?? into circular, and the reformed ?? distribution, which
was dissolved during bonding (compare to previous Figure)................................... 99
Figure 38: SEM micrograph in SEI mode, of asbonded+PBTE CMSX-4 ?
Niflex-110 - IN 738 bond showing coarsened as-bonded gamma-prime (??) and
deleterious phases. .................................................................................................. 100
Figure 39: Average shear stress versus condition of substrates. ASB refers to
Asbonded, PBHT-post bond heat treatment, PBTE- post bond thermal exposure.
Note that the increase in shear strength in CMSX-4 ? IN 738 after PBHT and the
decrease after PBTE. (Error bars denote the 1.96 times standard deviation Error
bars denote 1.96 times standard deviation to represent a 95% confidence limit)... 102
Figure 40 : LM micrograph, showing diffusion bond of MA956 fine grain - fine grain
(longitudinal orientation) at 1250 ?C for 121 s, followed by PBHT [ 1 h at
1300 ?C]. Notice the invisible bondline free of unbonded regions at the
bondline [127,128].................................................................................................. 107
Figure 41: LM micrograph showing diffusion bond of MA956 fine grain - fine grain
(longitudinal direction) at 1250 ?C for 121 s, followed by PBHT [1 h at
1300 ?C]. Note the grain growth across the bondline and also the voids present
at the bondline [127,128]. ..................................................................................... 108
xvi
Figure 42: LM micrograph showing diffusion bond of MA956 coarse grain -
fine grain (longitudinal direction) at 1250 ?C for 170 s. Note the bondline
with no voids present at the bondline [127,128]..................................................... 109
Figure 43: LM micrograph showing diffusion bond of MA956 coarse grain - coarse
grain (longitudinal direction) at 1250 ?C for 174 s. Note the bondline with large
unbonded regions present at the bondline127,128]. ............................................... 110
Figure 44: LM micrograph showing diffusion bond of PM2000 in L-L orientation at
1250 ?C for 310 s, in as-bonded (unetched) condition. Note the bondline free of
unbonded regions [127,128]. .................................................................................. 111
Figure 45: LM micrograph showing diffusion bond of PM2000 in Transverse
orientation at 1250 ?C for 300 s, in as-bonded (unetched) condition. Note the
bondline free of unbonded regions occasionally seen at the bondline 127,128]. ... 112
Figure 46: Average shear stress versus condition of substrates. PBHT-post bond
heat treatment. Note that the shear strength of the bulk material decreased after
recrystallization treatments and that of as bonded followed by PBHT is of the
order of 70% of the strength of the bulk material. (Error bars denote 1.96 times
standard deviation to represent a 95% confidence limit)........................................ 113
Figure 47: SEM micrograph in SEI mode, of MA 956 coarse grained bulk material
room-temperature shear fracture surface. Notice the main shear fracture surface
is planar, but there is extensive secondary cracking labeled. ................................. 114
Figure 48: SEM micrograph in SEI mode, of coarse grained MA 956 bulk material,
showing detail of fracture surface of room-temperature shear test. Notice the
cleavage cracking and shear bands on fracture surface labeled.............................. 115
Figure 49 : SEM micrograph in SEI mode, of fine grained bulk PM 2000, showing
fracture surface of (a) longitudinal sample and (b) transverse sample. Notice
that the main fracture surface is slightly less planar, but there is no secondary
cracking................................................................................................................... 116
Figure 50: SEM micrograph in SEI mode, showing fracture surface of PM2000
transverse-transverse bonds at 1250 ?C for 309 s, followed by PBHT 2 h,
1385 ?C. Notice the secondary cracking associated with planar shear................... 117
xvii
LIST OF TABLES
Table 1: List of some commercially available iron, nickel and cobalt based..................... 5
Table 2: Role of some common alloying elements in nickel-base superalloys [42] ......... 9
Table 3 Nominal compositions of nickel-based superalloys (given in wt %) .................. 39
Table 4: Nominal compositions of Iron-based ODS superalloys (given in wt %) ........... 46
Table 5: Microstructural bond development, for all substrate-interlayer combinations
(CMSX-4 ? IN 738 and CMSX-4 ? IN 939 with Niflex-110, Niflex-115, BNi-3)
{Italics signify the microstructural features that influenced mechanical properties}62
Table 6: Shear testing data of CMSX-4 ? IN 939 bonds using Niflex-110 interlayer at
different conditions. PBHT- Post bond heat treatments as mentioned in
section 4.1.3 and PBTE- post bond thermal exposure as explained in
section 4.1.4 .............................................................................................................. 88
Table 7: Shear testing data of CMSX-4 ? IN 738 bonds using Niflex-110 interlayer
at different conditions. PBHT- Post bond heat treatments as mentioned in
section 4.1.3 and PBTE- post bond thermal exposure as explained in
section 4.1.4. ............................................................................................................. 89
xviii
LIST OF ACRONYMS
APB - anti phase boundary
APBE ? anti phase boundary energy
CBN - cubic boron nitride
CG - coarse grain
EBPVD - electron beam physical vapor deposition
EDM - electric discharge machining
EDS - energy dispersive spectroscopy
FCC - face centered cubic
FG - fine grain
GAR - grain aspect ratio
HAZ - heat affected zone
INEEL - Idaho national engineering and environmental laboratory
L-L - longitudinal - longitudinal
LM - light microscope
MA - mechanical alloying
MPD - melting point depressant
ODS - oxide dispersion strengthened
PBHT ? post bond heat treatment
PBTE ? post bond thermal exposure
PC - poly crystal
PWHT ? post weld heat treatment
SC - single crystal
SEI - secondary electron imaging
SEM - scanning electron microscope
TCP - topologically close-packed phases
TEM - transmission electron microscope
TIG - tungsten inert gas
TLP - transient liquid phase
T-T - transverse- transverse
VIM - vacuum induction melting
1
1 INTRODUCTION
Superalloys are well known for their many high temperature applications,
including their use in critical gas turbine components in aerospace and in land based
power generation. They offer excellent high temperature tensile strength, stress rupture
and creep properties, fatigue strength, oxidation and corrosion resistance, microstructural
stability at elevated temperatures i.e., 600
0
C or above [1-6]. A variety of strengthening
mechanisms, such as precipitation hardening and solid solution strengthening, are
employed for the manufacture of these materials [7]. There are three classes of
superalloys: nickel-based superalloys, cobalt-based superalloys, and iron-based
superalloys [3,5]. Nickel-based superalloys are the most complex and the most widely
used in the high temperature applications. They currently comprise over 60% of the
weight of advanced aircraft engines [8,9].
At high temperatures, superalloys are subjected to severe operating environments
that include creep, thermo mechanical stresses, oxidation and corrosion. This results in
the development of creep damage, thermal fatigue cracks and surface degradation.
Repairing these service damaged components can result in substantial cost savings for the
aerospace and power generation industry as compared with replacing them with new
components. Thus joining of superalloys is a fundamental requirement for extensive post-
service repair of gas turbine engine components [8-14]. Joining is also needed for
primary fabrication, which includes high tolerance components such as jet fuel spray
2
nozzles [15], honeycombs [16] and solid walled structures [17]. Joining of dissimilar
alloys (single crystal to polycrystalline) is also required owing to the difficulty in
manufacturing large single crystal turbine blades. Conventional methods such as fusion
welding [18-21], diffusion bonding [22-24], and brazing [23, 25] suffer from severe
limitations when joining nickel-based superalloys. However, ?transient liquid phase
(TLP) Bonding? or ?diffusion brazing? as the American Welding Society refers to this
process has proven to be a successful joining technique for bonding nickel-based
superalloys [25-28], although successful TLP bonding of superalloys still requires a
better understanding of the microstructure-mechanical property relationships of the joint.
Also, in many industrial applications, wide joint gaps (with a gap width of 100 ?m or
more) are not unusual. Thus, this research investigates the relationship between the
microstructure, wettability, and mechanical properties of TLP bonded dissimilar nickel
base superalloys using wide gap style composite interlayers.
This research investigates the relationship between microstructure and mechanical
properties of TLP bonded dissimilar nickel-based superalloys. The microstructural bond
development, long term microstructural stability of the bonds with time, mechanical
properties of the joint using wide-gap style composite interlayer are examined and the
results compared to those with a conventional interlayer (BNi-3 foil), chosen for
comparison. Finally, a correlation between the microstructure-mechanical property
relationships and wettabilty is drawn.
Precipitation hardened superalloys are very reliable at higher operating
temperatures. However, as the precipitation hardened nickel base superalloys reach
higher temperatures (0.6 times the melting temperature), the second phase precipitates
3
coarsen rapidly [29-31], losing some of their strengthening abilities. This is due to the
transition in the strengthening mechanism from dislocation-particle shear to dislocation
bypassing. At even higher temperatures (about 1100
o
C), the second phase constituents
dissolve into solution, losing their strengthening ability altogether [32]. On the other
hand, the increasing demand for higher efficiencies drives the need for ever higher
temperatures.
This demand has led to the development of a new class of materials, ?oxide
dispersion strengthened (ODS) superalloys? [16]. ODS alloys derive their superior high
temperature strength from finely dispersed oxide particles (dispersoids) with high melting
temperatures and low solubility in the primary phase, which are stable at extremely high
temperatures and do not suffer from coarsening or dissolution effects [17]. There are two
classes of ODS alloys - ferritic (iron - based) and austenitic (nickel - based) ODS alloys.
In addition, ferritic ODS alloys are of interest as a fuel can material in the nuclear
industry due to their better void swelling resistance and irradiation embrittlement
properties compared to austenitic ODS alloys. These fuel can applications require the
ability to metallurgically join end caps, which are also made of ferritic ODS alloys to the
cans. Hence, the current research also focuses on the development of TLP bonds that
retain their creep resistant microstructure for use in fuel can applications of ODS alloys.
Thus, the primary objective of the current research is to study the structure-
property relationships of TLP bonded nickel-base superalloys and iron based ODS alloys.
4
2 LITERATURE REVIEW
2.1 Origin and Development of Superalloys
The origin and development of superalloys can be considered to be synonymous
with the birth and advancement of gas turbines. Although the principles of gas turbines
had been known for many years, industrial application was delayed for more than two
decades due to the lack of adequate materials [34]. The availability of other suitable
materials (chrome cast iron) for high temperature industrial and other strategic
applications also restricted the rate of progress in this field. The development of
superalloys has grown rapidly since World War II [35, 36]. Further contributions to this
progress include many techniques such as vacuum induction melting (VIM) [35] to avoid
oxidation of reactive elements and, advanced coatings to achieve superior environmental
resistance [37, 38]. Other improvements include the use of directionally solidified (DS)
and single crystal (SC) materials to achieve a combination of superior mechanical and
environmental performance [37, 39].
Out of the three classes of superalloys (nickel-based, cobalt based and iron based),
nickel-based superalloys are the principal high temperature material used for hot-zone
components in an aircraft engine. The ongoing research and development of new
superalloy designs and processes development has resulted in today?s availability of
many well-characterized nickel base superalloys with varying compositions, in-service
temperatures, and a wide variety of applications as shown in Table 1.
5
Table 1: List of some commercially available iron, nickel and cobalt based superalloys and their typical
compositions in wt% (Bal- Balance) [40, 41]
SUPERALLOY Alloy Ni Fe Co Cr V Ta Mo W Re Zr Al Ti C Other
A-286 26 Bal 0 15 0.2 - 1.25 - - - 0.2 2.2 0.1 -
N-155 20 Bal - 21 - - 3 2.5 - - - - 0.2 1Nb
IN 706 40 Bal - 16 - - - - - - 0.3 1.6 0.4 3Cb
HK 40 20 Bal - 24 - - - - - - - - 0.5 -
Fe-Base
CG 27 38 Bal - 13 - - 5.5 - - - 1.5 2.5 0.1 0.06Nb
IN 718 Bal 19 - 19 - - 3 - - - 0.6 0.8 0.1 5.2Nb
Mar-M247 Bal - 10 8.2 - 3 0.6 10 - 0.1 5.5 1.4 0 1.5Hf
Udimet-700 Bal - 19 15 - - 5.2 - - - 4.3 3.5 0.1 -
CMSX-2 Bal - 46 8 - 5.8 0.6 7.9 - - 5.6 0.9 0 -
IN 713 C Bal - - 12 - - 4.2 - - 0.1 6.1 0.8 0.1 2Nb
PWA 1480 Bal - 5 10 - 12 - 4 - - 5 1.5 - -
Waspalloy Bal - 13 19 - - 4.3 - - 0.1 1.3 3 0.1 -
N-4 Bal - 75 9.2 - 4 1.6 6 - - 3.8 4.3 0 0.5Nb
Rene-150 Bal - 12 5 3 6 1 5 2.2 0 5.5 - 0.1 1.5Hf
Alloy-D Bal - - - 15 2 2 4 - 0.2 4.5 2.5 0.1 -
Ni-Base
NTaC-13 Bal - 3 4 5.4 8.1 - 3.1 6.3 - 5 - 0.5 -
HS 188 22 3 Bal 22 - - - 14 - - - - 0.1 -
Co-Base
X-40 10 - Bal 25 - - - 7.5 - - - - 0.5 -
5
6
The first nickel-rich alloy developed for high temperature applications was
?Nichrome? (Ni-20 wt% Cr) heater wire. The chromium provides oxidation and
corrosion resistance at elevated temperatures; Chromium reacts with oxygen and results
in the formation of a stable and continuous chromia (Cr
2
O
3
) layer. This serves as a barrier
for the diffusion of oxygen into the bulk metal and provides excellent environmental
resistance for these alloys. Likewise, different alloying elements may be added to
improve the high temperature performance and environmental resistance. Table 2 shows
the different types of alloying elements and their roles in improving nickel-base
superalloy performance.
2.2 Microstructures and Mechanical properties
The microstructure of a typical nickel-based superalloy is shown in Figure 1, and consists
of the following [1]:
1. Gamma matrix (?): The gamma phase is an FCC nickel-based austenitic solid-
solution phase and has a random distribution of different species of atoms. This
comprises a high percentage of solid-solution elements such as cobalt, chromium,
molybdenum, and tungsten.
2. Gamma Prime (??): Aluminum and/or titanium, the essential solutes, are added in
amounts and mutual proportions with a total concentration of typically less than
10 wt % in order to precipitate high volume fractions of primitive cubic ??
[Ni
3
(Al, Ti)] coherent with the ? matrix. The structures of the ? and ?? phases are
shown in Figure 2.
7
3. Carbides: Carbon (0.05 - 0.2 wt %), combines with refractory and reactive
elements such as titanium, and chromium, to form MC carbides, which after heat
treatment and service generate lower carbides such as M
23
C
6
, and M
6
C at grain
boundaries.
4. Grain Boundary ??: Heat treatments and service exposures generate a film of ?'
along the grain boundaries, which helps to improve creep rupture properties.
5. Borides: Carbon and boron are added as solid solutions for grain boundary
strengthening (and thus are only used with polycrystalline alloys). When the solid
solubility of boron is exceeded borides will be formed.
The high temperature strength of nickel-based superalloys arises from a combination of
different strengthening mechanisms, including contributions from solid-solution
elements, particles, and grain boundaries. The major strength provider, however is the
coherent intermetallic compound ?? precipitated in a ? matrix [46].
At high temperatures, the thermal energy makes it easy for dislocations to move
and for this reason the strength of most metals decreases. However, nickel-based
superalloys are resistant to temperature and their strength increases with temperature up
to a certain range. The normal burger?s vector for an fcc ? lattice is (a/2) <110>.
However, for ??, a primitive cubic structure, a <110> is the lattice vector, as shown in
Figure 2. Thus, the motion of an (a/2) <110>
?
dislocation into ?? disrupts the order of the
??, and an anti-phase boundary (APB) is formed. To minimize the energy increase
associated with disordering, penetration of ?? has to occur by pairs of dislocations, known
as super-dislocations. This requirement for pairing of dislocations makes it more difficult
8
for dislocations to penetrate through ??, thereby making the alloy stronger and giving it
excellent creep resistance [47 - 49]. Neither single phase ? nor ?? are particularly creep
resistant. In contrast, two phase ?/?? is highly creep resistant. In both phases, there is
elastic {square of burger?s vector (b
2
)} repulsion between like dislocations. In ? there is
nothing to offset this repulsion. ?However, in ?? the antiphase boundary energy (APBE) of
the antiphase boundary (APB) created by ? <110> dislocations encourages these to move
in pairs (the second dislocation removes the APB). Hence, the equilibrium spacing
between the ? <110> dislocations is governed by the balance of the elastic repulsion and
the attraction due to the APB. Thus, dislocations are impeded in passing through two
phase ?/??, since the dislocations wish to be paired in ?? and separated in ?.
9
EFFECT ALLOYING
ELEMENTS
Solid-solution strengthening of ?
Ti, W, Mo, Cr, Co, Al
Solid-solution strengthening of ??
Cr, Mo, W, Ti, Nb, Ta
Stability of ??
W, Mo, Nb
Lattice mismatch increase Nb, Ti, Ta
Lattice mismatch decrease Cr, Mo, W
Antiphase boundary energy increase Ti, Co, Mo
Antiphase boundary energy decrease Al, Cr
Oxidation resistance improvement Cr, Al
Precipitate formers Ti, Ta, Al
Precipitation modification Co
Grain boundary phases B, C
Surface protection Al, Cr
Grain boundary strengthening Hf
Table 2: Role of some common alloying elements in nickel-base superalloys [42]
10
Figure 1: Schematic representation of the typical microstructure of nickel-base
superalloys [43]
11
Figure 2-The crystal lattice structures of (a) NiAl and (b) Ni
3
Al [44]
Crystal structure of ??
(cubic primitive)
Lattice vector <110> shown
Ni or Al
Al Ni
Crystal structure of ?
(cubic face centered)
Lattice vector a/2 <110>
12
2.3 Joining of nickel-based superalloys
2.3.1 Need for Joining Superalloys
The need for joining superalloys can be classified into three categories:
? Repair joining (pre- and post service): The higher operating temperatures that
improve the efficiency in gas turbines also result in rapid degradation of the
components by enhancing levels of thermomechanical cycling, oxidation and
creep. This damage necessitates the repair of these components in order to extend
their service life and to produce cost savings for the industry. This often requires
the damaged components to be repaired by joining new material to them [8-14].
? Primary fabrication: The complexity of gas turbine engine parts increases with the
improved efficiencies. In addition, the increasing size of land-based turbines
makes the large section components more prone to defects such as freckle
formation. To minimize the problems and allow the manufacture of geometrically
complex shapes, parts can be made in sections and then joined together.
Therefore, a joining technique that bonds several small and relatively simple parts
is often required for the primary fabrication of large complex components [15-17,
50].
? Joining of dissimilar (single crystal to polycrystalline) materials: Joining of
dissimilar superalloys is also needed, for example as a result of the difficulty in
manufacturing large single crystal turbine blades in land-based power generation
applications.
The compositions and microstructures of superalloys, especially the latest generation
alloys, have been designed and developed to meet the demanding requirements of
13
elevated temperature service. Likewise, the joining techniques used for these materials
must also be carefully tailored in order to be compatible with these operating
temperatures, stresses and environments. Fusion welding, diffusion bonding and brazing
are the three main joining techniques used in the industry for joining high temperature
materials [14]. As mentioned previously in the introduction, a number of joining
processes, such as conventional fusion welding [18-21], diffusion bonding [22-24][23,
25, 27] and brazing [23, 25, 27], have been investigated, but all have limitations when
used in joining superalloys. The disadvantages associated with these techniques are
discussed below.
2.3.2 Fusion Welding
The use of conventional fusion welding, to join superalloys often results in hot
cracking, post-weld heat treatment cracking, as well as the reduction in material
properties of the joint [8, 26, 51, 52-60]. Weld metal and heat affected zone (HAZ)
cracking, occurs during welding and during subsequent post weld heat treatments. The
length of these cracks varies from microns to tens of mm and they propagate like brittle
intergranular cracks exhibiting little plastic deformation. These cracks occur in the final
stages of solidification and can be classified as solidification cracking (also known as
liquation cracking [61].
As the solidification of the weld metal occurs, columnar grains grow from the
fusion boundary towards the central region of the weld pool. This results in the rejection
of impurities and alloying elements ahead of the weld pool and getting concentrated in
the last liquid to solidify [56, 60]. This rejection results in the lowering the freezing point
of the liquid, thereby allowing it to srvive for a period of time, even after the grains have
14
begun to meet the center of the weld pool. [54, 58, 60].These liquid films formed reduce
the surface contact of the grains forming regions of low ductility and thereby low strength
in the weld. Also, the residual liquid is often insufficient to fill the voids present and the
contraction stresses due to the cooling of weldment produces intergranular cracks along
the weakened grain boundaries [54, 56, 58, 60]. In low ?? containing alloys, manipulation
of weld parameters might minimize these effects, but in high ?? alloys, preheating the
substrate to reduce the thermal gradients and strains in the weld zone might work. In
addition, hot cracking is also dependant on other contributory factors such as
solidification and microsegregation, viscous flow of liquid metal, crack initiation and
propagation [52]. An extensive study of factors leading to solidification cracking and the
analytical and finite element modelling methods used to simulate these phenomena has
been discussed elsewhere [52].
Post weld heat treatment cracking (PWHT) or strain age cracking, generally,
occurs in precipitation hardened superalloys during post-weld heat treatment or
subsequent high temperature service. This is due to the presence of either residual
stresses developed during welding, or applied stresses arising from service. These PWHT
cracks are characterised by intergranular micro-cracking in either the HAZ or weld bead
and form as a result of precipitation and hardening of the alloy during thermal exposure
and transfer of solidification strains onto the grain boundaries [54, 55]. A common
practice is to attempt to relieve residual stresses arising from the welding procedure by
means of post weld heat treatment. However, often the stress-relieving temperature is
greater than the aging temperature of the alloy and this leads to a transient precipitation
period during post weld heating. This hardens the alloy resulting in excessive strain
15
localization on grain boundaries within the HAZ and weld bead during heating, leading to
cracking. [54].
Other problems associated with fusion welding are the precipitation of brittle
second phases in the weld joint and/or HAZ which might lead to reduction in mechanical
properties such as tensile or yield strengths and ductility [53]. Also, the segregation of
impurities and alloying elements at the joint may lead to a reduction in oxidation
resistance. It has also been reported that fatigue strength of the weld joints was reduced
relative to the base metal [65, 66].
2.3.3 Diffusion Bonding
Superalloys are designed to form a stable oxide layers to enhance their oxidation
resistance. However, diffusion bonding suffers from the limitation of being unable to
readily join systems with stable oxide layers [67]. Since diffusion bonding relies on the
diffusion of interlayer species into the substrates, it thus requires longer bonding times to
attain a perfect bond, with the elimination of non-bonded regions. This also requires high
pressures, longer processing times, and if not conducted carefully may lead to some
deformation in the components during bonding, making the process too expensive for
many applications [25, 26]. When complex geometries are involved, it would be difficult
to apply uniform bonding pressures normal to the mating surfaces and it might require
expensive and sophisticated tooling.
2.3.4 Brazing
Brazing suffers from the limitations of the bond remelting temperature being less
than the operating temperature of the superalloys and the formation of intermetallics
16
along the bondline [25, 26]. This results in bonds with poor mechanical and physical
properties.
2.3.5 Transient Liquid Phase Bonding
In TLP bonding, a liquid forming interlayer (usually rich in the substrate material,
with the addition of a melting point depressant (MPD)) is selected such that the melting
point of the interlayer (initially) is significantly lower than the melting point of the
substrates, and hence has the same initial condition as a braze joint [25, 26]. Once the
bonding temperature is reached, which is greater than the melting point of the interlayer
and below the melting point of the substrates, the interlayer melts forming an eutectic or
alternatively can form a liquid by reaction between the interlayer and substrates. The
further development of the bond can be described using three different, distinct stages.
Figure 3 and Figure 4 show the different stages of TLP bonding.
? Substrate dissolution brings the compositions of the interlayer (usually starting
from an eutectic) and the solid into local equilibrium i.e., the compositions of
liquidus and solidus for the liquid and adjacent solid respectively, by diluting and
widening the liquid interlayer. This stage is very rapid, since it does not rely on
long-range diffusion in the solid phase.
? Isothermal solidification: In this stage, the solute (melting point depressant)
diffuses from the interlayer into the substrates under local equilibrium
concentrations and constant bonding temperatures, resulting in re-solidification of
the joint at the bonding temperature. This stage is governed by long-range
17
diffusion in the solid phase. The time required for isothermal solidification to be
completed depends on the system (substrates and interlayer) employed.
? Solid-state homogenization In this stage, diffusion of remaining MPD at the
bondline into the substrates occurs and reduces the bondline concentration to
below the room temperature solubility of MPD into the substrates. This produces
a uniform solute concentration, homogenization of remaining solute and
eliminates the formation of secondary phases at the bondline on cooling. Hence,
the microstructure and mechanical properties of TLP bonds can approximate
those of the substrate material, unlike brazing.[68, 69]
Transient liquid phase bonding is the preferred joining method for nickel-base
superalloys because of its better joint properties compared to a brazed joint, the lower
pressures required for bonding compared to a diffusion bond, its tolerance of stable oxide
layers on the faying surfaces, the avoidance of both hot cracking in the joint and (as an
isothermal process) the residual stresses that lead to post-weld heat treatment cracking,
plus the elimination of brittle secondary phases along the joint.
18
Figure 3 Stages of TLP bonding - Interlayer melting and substrate dissolution [25, 27]
Terminology:
A = substrate C
A
= Composition of the substrate, C
E
= Initial (eutectic)composition of the liquid T = Temperature
T
B
= Bonding temperature C
s
= Solidus composition C
L
= Liquidus composition
MPD = Melting point depressant C
R
= Room temperature solubility of MPD in A
Melting of interlayer
T
T
B
Substrate
Substrate
C
A
C
E
Liquid
A B
Substrate dissolution
A B
Substrate
Substrate
C
L
T
T
B
Liquid
C
S
18
19
Figure 4 Stages of TLP bonding
Isothermal Solidification and
Solid-State Homogenization [25, 27]
The terminology is the same as Figure 3.
Substrate
Substrate
C<C
R
T
T
B
A B
Solid-state homogenization
20
2.3.5.1 Advantages and Disadvantages of TLP bonding
While some of the advantages of TLP bonding have been mentioned already, the
advantages and disadvantages can be summarized as follows [98]:
Advantages:
? TLP bonding, being an isothermal process unlike fusion welding, is ideal
for joining materials susceptible to hot cracking or post-weld-heat
treatment cracking [21].
? Large and intricate geometries such as honeycomb panels can be easily
joined without expensive and complex tooling.
? The formation of undesirable brittle intermetallics in the bond- line can be
avoided, resulting in a microstructure, and hence mechanical properties, of
the joint that is similar to those of the bulk material.
? TLP bonding is tolerant to stable oxide layers on the faying surfaces and
requires minimal fixturing pressures and surface preparation, unlike
diffusion bonding [99]
These advantages have led to the use of TLP bonding for joining a wide range of
materials, such as nickel-base superalloys [23, 26 ,100-103], oxide dispersion
strengthened alloys [94, 95, 104, 105], metal matrix composites[106-108], and structural
intermetallics [109-111]. TLP bonding also offers the advantage of joining dissimilar
materials (for example, metals to ceramics) [112, 113], via derivatives of conventional
TLP bonding.
21
Inspite of its wide variety of applications, it is also important to note some of the
disadvantages associated with TLP bonding:
Disadvantages:
? Rapid heating rates must be employed in order to avoid the formation
of brittle intermetallics during the heating process. This helps in
minimizing the diffusion of interlayer and /or substrate constituents
during heat-up.
? Long bonding and/or post bond heat treatment (PBHT) process times
are often needed in order to attain complete microstructurally and
chemically homogenized bond- line.
? Selection of interlayer composition and MPD becomes very difficult in
some cases, and the bonding process is strongly dependent on this
choice
2.3.6 Wide-gap Transient Liquid Phase Bonding
In TLP bonding, when thick interlayers are used to fill a wide joint gap (100 ?m
or more) extended bonding times are required in order to complete isothermal
solidification and subsequent solid-state homogenization. Hence, conventional TLP
bonds rarely employ a joint gap greater than 50 ?m, and in some cases these may have an
order of magnitude less than 50 ?m [27]. However, wide joint gaps are often expected in
industrial applications [70], so for TLP bonds using a wide joint gap, composite
interlayers consisting of a liquid forming constituent coated on either side of a non-
melting core can be employed. The liquid forming constituent is typically selected such
that its melting point is slightly below the bonding temperature or to form a liquid by
22
reaction with the substrates and/or interlayer core, while the non-melting core is selected
to have a composition and the solidus temperature similar to the base material [27, 71,
72]. The melting point depressant in the composite interlayers serves two functions. One
is to increase the interfacial area between the liquid and solid phases, so that a more
efficient path for the diffusion of the solute from liquid to solid will be created. These
interlayers also reduce the amount of liquid necessary to fill the gap, so that less solute
needs to be diffused in order to produce isothermal solidification and solid-state
homogenization [27, 73, 74].
In the wide gap TLP process, the most important parameter is the choice of an
appropriate amount of liquid former to be deposited on the non-melting core. If
insufficient liquid former is deposited, premature isothermal solidification occurs and the
liquid ceases to spread before it has had time to penetrate throughout the joint [75, 76].
This leads to excessive porosity at the joint. On the other hand, if a great deal of liquid is
formed, then the amount of the solid-phase dissolved in order to produce a local
equilibrium at the solid- liquid interface (i.e., solid-state dissolution) would be very high
and the advantages of using the composite interlayer would be lost [27,77,78]. The use of
excessive liquid former in a composite interlayer may also result in undesirable second
phase precipitation in the bondline [54].
2.4 Development of Oxide Dispersion Strengthened Superalloys
Precipitation hardened superalloys are very consistent at higher operating
temperatures in terms of their strength. However, as precipitation hardened nickel base
superalloys reach higher temperatures (0.6 times the melting temperature), the second
phase precipitates coarsen rapidly [29-31], losing some of their strengthening abilities.
23
This is due to a transition in the strengthening mechanism from dislocation-particle shear
to dislocation bypassing. At even higher temperatures (around 1100
o
C), the second
phase constituents dissolve into solution, losing their strengthening ability altogether
[32]. Thus, the increasing demand for higher efficiencies has driven a need for ever more
elevated temperatures, which has led to the development of a new class of materials,
?oxide dispersion strengthened (ODS) superalloys? [80].
2.5 Mechanical Alloying
Although ODS alloys are manufactured using different techniques, such as
mechanical mixing, internal oxidation, selective reduction and rapid solidification, these
processes all suffer from drawbacks, namely too irregular interparticle spacing, non-
uniform oxide size, difficulty in reduction of stable oxides, and limited solid-solubility,
respectively [80, 81]. In contrast, mechanical alloying (MA) minimizes the above
mentioned drawbacks and results in the optimum interparticle spacing and uniform oxide
size that are required for high temperature operation. Mechanical alloying minimizes the
interaction between the matrix and dispersoids and the oxide particles are included as an
incoherent phase within the alloy matrix.
In mechanical alloying, the alloying powders are continuously cold welded, work
hardened and fractured using a dry, high-energy ball milling technique. This process,
which requires no liquids or surfactants, produces composite metallic powders with a
controlled and fine microstructure. The mechanically alloyed powders are consolidated at
900
0
C ? 1100
0
C by extrusion or by hot isostatic pressing in vacuum sealed cans [82-
84]. No melting of the material takes place throughout the whole process. Final heat-
24
treatments are used to recrystallize the heavily deformed material, which will result in
coarse columnar type grains [80,83,84,86].
The disadvantage of using the mechanical alloying process is the difficulty in
controlling the final microstructure of the ODS alloys, which leads to inconsistent batch
to batch properties and reproducibility. In addition, the operating temperatures of the
ODS alloys are limited by the melting temperature of the matrix.
2.6 Properties and Applications
The dispersion of particles in a metal matrix produced by mechanical alloying
markedly increases the strength and environmental resistance of these materials at high
temperatures. The finely dispersed stable oxide particles (dispersoids), with high melting
temperatures and low solubility in the primary phase, are stable at extremely high
temperatures and do not suffer from either coarsening or dissolution effects[87].
The outstanding strength and corrosion resistance of ferritic ODS alloys is due to
? Nanosized yttria dispersoids, which are non-shearable and therefore act as
obstacles to the dislocations by pinning them and help to achieve favorable high
temperature creep and stress rupture strength.
? Coarse grain structure with a favorable texture and high grain aspect ratio (GAR),
which improves the high temperature creep properties at intermediate and high
temperatures (above 1000
0
C). The coarse grain structure also helps in preventing
grain boundary sliding and Nabarro-Herring creep at high temperatures [88].
? Adherent surface layer of protective alumina scale, which is formed during
exposure at high temperatures, due to the high aluminium content in the matrix
(5%), and provides excellent oxidation resistance. The presence of yttria (around
25
0.5%) is beneficial in improving the oxide scale adherence during thermal cycling
[89].
There are two groups of ODS alloys.
? Austenitic (Ni-based) ODS alloys strengthened by oxide dispersions and ?? phase
precipitates, which are used as blade materials for advanced gas turbines.
? Ferritic (Fe-based) alloys strengthened by oxide dispersions and have applications
in aero engine and gas turbine chambers, and as cladding materials for fast
breeder reactors and fuel cans in the nuclear industry [44].
The austenitic ODS alloys such as MA 6000 used in advanced gas turbine
engines, combine the strengthening effects from the incoherent oxide particles and the
coherent ?? precipitates in the ? matrix. Thus, even after the solutionizing temperature of
?? (at around 1000
0
C), high temperature strength is provided by the dispersoids present
in the alloy [85]. Ferritic ODS alloys find their potential applications in the nuclear
industry as a fuel can material due to their better resistance to void swelling and
irradiation embrittlement than austenitic superalloys.
2.7 Joining of ferritic based Superalloys
2.7.1 Conventional Techniques and Limitations
The need to be able to join ODS alloys arises from their applications in primary
fabrication, pre and post service repair, and the difficulty in manufacturing large complex
shaped parts. Conventional joining techniques such as fusion welding, friction welding
and brazing have all been investigated by various researchers for use in joining ODS
alloys, and these will be discussed in turn below.
26
As noted in section 2.3 for nickel-based superalloys, fusion welding is limited by
melting of the base metal leading to the formation of dispersoid agglomerates in the joint,
porosity in the joint, undesirable secondary fine recrystallized grains in the joint (boron
present in the weld filler material diffuses into the substrates resulting in lattice strains
and thereby providing additional energy for recrystallization), grain growth in the joint
normal to that of the base metal [90] and formation of brittle second phases in the joint
[91], all of which can lead to premature failure of the joint. In other experiments by
Moilan et al. [92] using tungsten inert gas (TIG) welding, bondlines of lower hardness
values were obtained due to oxide agglomeration and loss of grain orientation was
observed. Yttria agglomeration resulted in a discontinuous chromia scale, as a result of
which the sulphidation resistance of the TIG weld was low. The grain size in the welded
joint was coarser than that of the substrate material. This resulted in bonds with poor
room temperature mechanical properties.
Solid-state joining processes such as explosion welding and friction stir welding do
not result in melting of the base material, so an effective undisturbed yttria dispersion can
be retained in the joints using mechanically alloyed ODS alloys. However, adiabatic
shear bands were observed at the bondline when two ODS alloy substrates were
explosion welded and showed an impact on mechanical properties [91]. In conventional
friction welding, the chemistry, dimensions and shape of the particles were all altered
[93]. Reorientation of some of the grains has also been observed due to the final forging
process in friction stir welding. Due to the high strain rates and oxidation involved, yttria
agglomeration, large titanium-rich and aluminum-rich oxide particles were observed [93].
27
Brazing of ODS alloys also suffers from several disadvantages. The major drawback
in brazed components is that the joint remains as a low-remelting region, hence
restricting the joint to lower operating temperatures. In addition, a brazed joint assembly
results in the formation of brittle secondary phases, leading to poor mechanical properties
[26]. The formation of a protective oxide layer on ODS alloys also inhibits the flow of
the braze filler metal.
2.7.2 Diffusion Bonding
Diffusion bonding is a solid-state joining technique in which the two surfaces to
be joined are brought into contact and held together for extended periods of time under
high pressures at elevated temperatures [115, 116]. The high temperatures allow the
microscopic deformation of the points of contact between the two materials, thereby
greatly increasing the true contact area. Diffusion bonding can be described in terms of
three different stages [117].
Stage I: This stage corresponds to the initial state in which the two surfaces to be
joined are brought into contact. When high pressures and temperatures are applied, the
plastic deformation of asperities on the two faying surfaces results in the formation of a
bond plane with a large number of pores.
Stage II: In this stage, further heating allows the migration of grain boundaries
and the elimination of porosity through diffusion. Growth of metal crystals across the
bond interface takes place, thereby causing the pores to shrink in size. This stage does not
require the application of external pressure.
Stage III: This is the final stage of the bonding process, in which volume
diffusion results in minimization of the final porosity. However, complete elimination of
28
the remaining final pores is extremely difficult at this stage and requires high pressures
irrespective of heat treatments and longer bonding times. A careful selection of bonding
pressures and temperatures in stage I and stage II yields better bonds with less porosity.
Diffusion bonding of ODS alloys using unrecrystallized and recrystallized ODS
materials in the working direction results in fine grain layer formation at the bond-line
[118-120]. This layer often produces deterioration in the properties of the joint. In
addition, the need for a high bonding stress is a disadvantage for this process. For
example, when recrystallized MA956 material was diffusion bonded by Zhang et
al.[120], a thick layer of oxide was observed at the bondline at low bonding stresses.
Higher bonding stresses (56 MPa and 72 MPa) resulted in bonds with tensile strengths
comparable to that of the parent metal. However, the presence of circular voids at the
bondline and a transverse grain boundary extending through out the length of the joint
resulted in a reduction of the joint strength at high temperatures. An increase in bond time
and bonding stress was suggested by the authors, which makes the process uneconomic
for industrial applications. Diffusion bonding of MA6000, a nickel-base ODS alloy was
studied by Moore and Glasgow [118]. Although, they obtained successful macroscopic
joints, low stress rupture strength of the joints was observed.
Although, some success in diffusion bonding ODS alloys has been achieved for
the bonding of unrecrystallized to recrystallized material and recrystallized material to
itself, the very high stresses (200 ? 300 MPa), high temperatures, longer processing times
and extensive specimen preparation required make this process uneconomical.
29
2.7.3 Transient Liquid Phase bonding of ferritic based ODS Alloys
The advantages of TLP bonding, discussed earlier in section 2.3.5, indicate that it
is a promising joining technique for oxide dispersion strengthened alloys when compared
to conventional techniques. This is because it prevents base metal melting although
localized dissolution of the substrates still does occur, eliminates brittle secondary phases
in the joint and results in minimal or negligible microstructural disruption so that the
bond will have the same properties as the bulk material. The vital issues in the TLP
bonding of ODS alloys are the distribution of dispersoids after melting of the interlayer
and the consequences of this for microstructural development and, hence, mechanical
properties. Therefore, the study of microstructure-mechanical property relationships is
crucial in TLP bonded ODS alloys.
Khan et al.[94, 95] successfully employed the TLP bonding technique to join
ferritic ODS superalloys MA956 and PM2000 in the fine grained condition using an iron-
based foil as an interlayer, with boron and silicon as the melting point depressants. The
bonds obtained were free from precipitates, pores, and excessive agglomerations of
dispersoids. However, the recrystallized grain size did not match that of the bulk material
annealed under similar conditions. In addition, the joining process used is only suitable
for narrow joint gaps (of a few ?m), which are difficult to achieve consistently in many
industrial applications.
30
3 RESEARCH OBJECTIVES
This research is aimed at the investigation of strategies for joining dissimilar
superalloys, (single crystal to polycrystalline) for use in land-based gas turbines for
electric power generation. As already mentioned in section 2.3.5 and 2.7.3, TLP bonding
is much more suited for the joining of superalloys than conventional techniques.
However, the successful application of TLP bonding to superalloys requires a complete
understanding of microstructure-mechanical property relationships. In addition, this
better understanding of TLP bonding process can be achieved only through a close
correlation between real-world systems and analytical or numerical models. Also, the
extended bonding times required for thicker (greater than 50 ?m) interlayers and the wide
joint gaps (a gap width of 100 ?m) in real industrial applications requires an in depth
study of wide gap TLP bonding.
Thus, the main objective of this research is to examine the details of
microstructural bond development and their relationship with the mechanical properties
in wide gap style TLP bonded superalloys. The objective also includes a study of the
long-term microstructural stability of these bonds and the production of a joint free of
brittle second phases, with a microstructure and, hence mechanical properties similar to
the unbonded material. This work takes into account the effects of interfacial segregation
and gravitational influences on the liquid phase and solid-state diffusion. Further, this
study will address the wettability issues for wide-gap style composite interlayers on
31
the three substrates used in joining. A flow chart to aid in the understanding of the goals
and investigations conducted is shown in Figure 5.
As mentioned already, the ever increasing demand for higher efficiencies drives
the need for materials that retain their properties at elevated temperatures (higher than the
solution temperature of ??), have led to the development of mechanically alloyed ODS
alloys. However, ODS alloys also need to be joined due to the various reasons mentioned
in section 2.3.1. and as mentioned in section 2.7.3, TLP bonding has shown some
promising results in joining ODS alloys.
Of the two classes of ODS alloys (ferritic and austenitic), ferritic ODS alloys are
of particular interest as a potential fuel can material for use in the nuclear industry due to
their better void swelling resistance and irradiation embrittlement compared to austenitic
ODS alloys [96]. However, these fuel can applications require the ability to
metallurgically join end caps, which are also made of ferritic ODS alloys. Hence, the
current research also aims at developing TLP bonds that retain a creep resistant
microstructure, for use in fuel can applications of ODS alloys.
Thus, the overall objective of the current research is to study the structure-
property relationships of TLP bonded nickel-based and iron based superalloys. The
specific objectives related to nickel-base superalloys and oxide dispersion strengthened
superalloys are listed below.
Nickel-based superalloys:
? Investigate the bond development and long-term microstructural stability
of wide gap style TLP bonded dissimilar nickel-based superalloys
32
? Produce a joint free of brittle second phases, with mechanical properties
similar to that of the unbonded material
Oxide dispersion strengthened superalloys:
? Design and implement a TLP bonding process that provides reasonable
dispersoid continuity across the bondline of ferritic ODS alloys
? Produce a continuous, recrystallized microstructure across the bondline,
resulting in mechanical properties similar to that of the parent material.
33
Figure 5: Flow chart of project objectives and investigations
1
for TLP bonding of dissimilar nickel-based
superalloys
1
Wettability studies are conducted by Subhadra Chitti and microscopy was done by Nofrijon Sofyan
Shear testing
Hardness tests
Recrystallized
Hot stage light
microscopy
(HSLM)
Interrupted
bonding runs
Wettability
1
Mech. Props.
Microstructure
Preoxidized
As recieved
PBHT
As bonded
TLP BONDING
of NICKEL-
BASED
SUPERALLOYS
33
34
4 MATERIALS AND EXPERIMENTAL PROCEDURE
4.1 Nickel-based superalloys
4.1.1 Materials
Single crystal CMSX-4, and polycrystalline IN 939 and IN 738 were the three
nickel-based superalloys used in this research. CMSX-4 was joined to IN 738 and to IN
939. The joining process used proprietary wide-gap style composite interlayers Niflex-
110 and Niflex-115 (the latter has a higher boron content than the former); a conventional
foil interlayer BNi-3 (Ni-4.5 wt%, Si- 3.2 wt%, B) supplied by Metglas Inc. was chosen
for comparison. The composite interlayers and substrate materials were provided by
Siemens-Westinghouse, Orlando, Florida.
The as-received blocks were 52x48x23 mm. The microstructure of as-received
CMSX-4 is shown in Figure 6. A high volume fraction of ?? with no rafting was observed
in the as-received microstructure of the substrates. The substrates were ground to a 1000
grit finish using SiC paper.
The as-received interlayers were 50 ?m in thickness. The composite interlayer
consisted of a non-melting core with a melting point depressant (in this case, boron)
coated on both sides of it. Thus, the liquid formed upon melting of the interlayer was in
direct contact with the substrate surfaces, allowing the melting point depressant to diffuse
readily into the substrate constituents. A schematic of a composite interlayer is shown in
Figure 7. The advantages of using a wide-gap interlayer were discussed in section 2.3.6.
35
Figure 6 Microstructure of as-received single crystal CMSX-4 [97]
1 ?m
36
Figure 7: Schematic of wide gap composite interlayer
BORON
BORON
Niflex non-melting core
37
4.1.2 Joining Procedure
The substrates were ground to a 1000 grit finish using SiC paper and
ultrasonically cleaned in an acetone bath. The interlayer was placed between the prepared
CMSX-4 and IN 738/IN 939 substrates (Refer to
Table 3 for compositions); the whole assembly was wrapped with Ta wire for
fixturing. The whole assembly was inserted in a tube furnace and joined in a vacuum (at
least 13.3 mPa) using the bonding treatment conditions described below.
The conditions of the as-bonded samples were:
? Bonding temperature: 1160?C
? Ramp time: ~140 min
? Bonding time: 4 hrs
? Cooling: Removed from the vacuum tube furnace and fan-cooled at room
temperature after 4 hr of bonding
All the bonding trials were performed with the (001) plane of the CMSX-4 as the faying
surface joined to the polycrystalline substrates IN 738 and IN 939.
To study the microstructural bond development, a series of interrupted bonding
tests (with bonding times consisting of 0 min {i.e., heating to the bonding temperature
and immediately cooling}, 1 min {i.e., heating to the bonding temperature, dwell for 1
min and then cooling} and so on ) were conducted. Interrupted bonding studies at 2 min,
4 min, 10 min, 1 hr, 2 hrs, and 4 hrs (all at 1160
0
C) were also performed.
4.1.3 Bond Heat Treatment (PBHT)
?Post-bond heat treatment? (PBHT) was performed on the as-bonded samples to
allow the diffusion of aluminum from the substrates in order to form more ?? on the
38
bondline and to reform the ?? distribution in the bulk, which was dissolved during
bonding. The assembly was heated (30 min ramp time) to 1093 ?C and held for 1 hour.
This is the temperature at which the ?? goes into the solution.
After 1 hour, the furnace was cooled to 1000 ?C in vacuum and held for 4 hrs in
vacuum. This heat treatment of 4 hrs at 1000 ?C allowed the precipitation and growth of
primary ??.
The furnace was then cooled to 843 ?C and held for 16 hours in a vacuum. This
generated considerable quantities of secondary, fine ??, after which the sample was fan
cooled to room temperature.
In summary, the PBHT consisted of:
? Heat to 1093?C and hold for 1 hour in vacuum;
? Cool to 1000?C and hold for 4 hrs in vacuum;
? Cool to 843?C and hold for 16 hrs in vacuum;
? Fan-cool to room temperature.
4.1.4 Post Bond Thermal Exposure (PBTE)
The ?post bond thermal exposure? (PBTE) refers to studies designed to examine the
microstructural stability of the bonds. In order to determine the microstructure-
mechanical property relationships of the bond with time, post-bond thermal exposures of
up to 1 week (168 hrs) at 1000
0
C, were performed on the as-bonded samples.
39
Table 3 Nominal compositions of nickel-based superalloys (given in wt %)
Element CMSX-4 IN 939 IN 738
Ni 60 48 61
Cr 6.5 22.4 16
Co 10 19 8.5
Mo 0.5 - 1.7
W 6.5 2 2.6
Ta 6.5 1.4 1.7
Nb 5.5 1.9 3.4
Al 5.5 1.9 3.4
Ti 1 3.7 3.4
Others 3 - 1
40
4.1.5 Metallographic Preparation
The as-bonded, PBHT and PBTE samples were sectioned using a wire EDM /
Struers Accutom-5 high speed wafering saw employing a cubic boron nitride (CBN)
blade. Once sectioned, the samples were mounted using a Struers Labpress-3. Samples
were ground to a 1200 FEPA grit finish and then polished using a 6 ?m diamond spray,
followed by a 0.04 ?m alumina suspension solution. The metallographic samples were
etched to show ?? using a mixture of 20 vol% conc. nitric acid and 80 vol% conc.
hydrochloric acid for 5-10 seconds at room temperature.
4.1.6 Microstructural Characterization
The samples, as-bonded, PBHT and PBTE, were primarily characterized using
light microscopy (LM) and scanning electron microscopy (SEM), supplemented with
SEM based energy dispersive X-ray spectroscopy (EDS). These studies
2
were conducted
on a JEOL JSM840 instrument operated at an accelerating voltage of 20 kV. Secondary
electron imaging (SEI) was used to collect all SEM micrographs. The SEM based EDS
analysis used ultra-thin window detectors and Oxford Instruments ISIS analyzers. EDS
analysis was used to determine the composition of phases observed in the bond- line and
compositional profiles across the joint.
4.1.7 Mechanical Testing
The study of bond mechanical properties was performed by means of shear and
hardness testing. All specimens for shear testing were extracted using an EDM and the
mechanical testing was supplemented by fractographic investigations using SEM and
2. As mentioned already, microscopy was conducted by Nofrijon Sofyan.
41
EDS analysis. The details of the mechanical testing procedures employed are discussed
below.
Initial mechanical testing was done on specimens extracted from an as-bonded
and PBHT samples via EDM, and then tested in four-point bending by Tao Zhou [97].
However, these samples deformed to the maximum displacement of the testing rig
without failing, thus rendering the four-point bend test unusable. Double lap shear testing
was conducted for this study, but problems were encountered due to the test grips
slipping off the samples. Tensile testing was also unsatisfactory for these samples due to
the limitation of testing the bulk in addition to the bond.
Thus, the shear testing method developed by Yan and Wallach [114], which is an
easy, effective and inexpensive method of testing the bond was selected for use in this
study as an acceptable alternative means of testing the bond strength.
4.1.7.1 Shear Testing
The shear tests were carried out with a specially designed grip developed by Yan
and Wallach [114]. The schematic in Figure represents the assembled unit, with a loaded
specimen in the geometry of Figure 8(a). The adjustable screw in the line drawing was
adjusted in order to place the bond-line in the plane of shear. When the assembled unit
was then placed in tension, the outer sheath maintained the shear along the bond-line
until the specimen failed. The shear test samples were composed of a circular piece and a
rectangular piece (see Figure 8). The disc shaped substrates were 12-13 mm in diameter
and 3.5 mm thick and the rectangular substrates were 8 x 4 mm for faying surface
dimension and 2 mm thick. All the tests were performed at room temperature. An MTS
42
Q-Test 100 screw-driven machine was employed to conduct these tests. Testing was
performed at a grip separation rate of 0.5 mm/min.
43
Figure 8(a): Schematic of sample used for shear testing [114]
1. Sample 2. Down fixture 3. Up fixture
4. Hardened inserts 5. Screw 6. Shield
Figure 8(b): Schematic of grip used for shear testing [114]
1
2
3
4
5
6
6
12- 13mm
4mm
8mm
8mm
~3.5mm
2 mm
4mm
44
4.1.7.2 Hardness Tests
Vickers microhardness testing was also employed in order to determine the
hardness of the intermetallics present in the bond region and the hardness profile across
the bond- line. Microhardness testing was conducted using a 300 g load for a duration of
10 s.
4.1.8 Wettability Studies
Wettability tests were conducted
3
using the sessile drop technique [121], in order
to investigate the mechanisms involved and the ability of the Niflex interlayers to wet the
superalloy substrates. Substrate samples with a geometry of 4mm X 4mm and 1mm
thickness were cut using a Struers Accutom-5 and ground to 1200 FEPA grit. Disks with
a diameter of 1 mm were mechanically punched from a 50 ?m thick proprietary Niflex
interlayer. The prepared substrates and the interlayer were then ultrasonically cleaned in
an acetone bath before the experiments.
The sessile drop tests were dynamically observed on a real time dynamic hot
stage light microscope, using a Leitz 1750 heating stage mounted on a Leitz DMR light
microscope. The experiments were conducted in a 1.3 MPa vacuum atmosphere at the
bonding temperature. These tests were video recorded, so that the displacement of the
solid- liquid interface was measured
4
as a function of holding time. The wetted samples
were cross-sectioned in order to study the wetting mechanisms and the microstructure.
The displacement measurements were conducted by Robert Love.
3. As mentioned already, wettability tests were conducted by Subhadra Chitti
4. The displacement measurements were conducted by Robert Love
45
4.2 Oxide Dispersion Strengthened Iron Based Superalloys
4.2.1 Materials
The primary materials used were MA956, PM2000 and boron. MA956 and
PM2000, the two ferritic based ODS alloys, were supplied by Idaho National Engineering
and Environmental Laboratory (INEEL). MA956 was received in fine grain (11 mm x 11
mm square billets) and coarse grain (20 mm diameter rods) condition, and PM2000 (25
mm diameter rods) was received in fine grain condition. The compositions of the alloys
are presented in Table 4. After the initial experiments on MA956, further more detailed
investigations were done on PM2000 due to availability of the material.
Samples with dimensions of 11 mm x 10 mm and 2 mm in thickness were cut
using an electric discharge machine (EDM) in longitudinal (along the direction of
extrusion) orientation. The substrates were surface ground to maintain their flatness and
to remove the copper layer formed during EDM cutting. A surface grinder from Gold
International Machinery, Inc. was utilized for this. These substrates were then
ultrasonically cleaned in an acetone bath and then stored in acetone to prevent any oxide
formation prior to bonding.
Electron beam physical vapor deposited (EBPVD) boron on the substrates was
used as the interlayer for TLP bonds. Boron was selected as the MPD due to its high
diffusion coefficient in ferritic materials. Boron films with thicknesses of 250 nm, 500
nm and 1 ?m were used as interlayers.
46
Table 4: Nominal compositions of Iron-based ODS superalloys (given in wt %)
Composition (wt %)
Alloy
Fe
B
Cr Al Ti Y
2
O
3
C
MA956
bal - 20 4.5 0.3 0.5 0.04
PM2000 bal - 20 5.5 0.5 0.5 0.01
47
4.2.2 Joining Procedure
Bonding was performed on the Gleeble 1500, a thermo-mechanical processing
system that can provide high vacuum, thermomechanical cycling, and bonding force
using compressed air or hydraulic loading system.
The two substrates (boron coated and uncoated) to be joined were brought into
contact via carbon blocks on either side. Carbon blocks were used in order to provide
effective thermal conduction between the water cooled copper blocks and the base metal.
To prevent the formation of carbon-iron eutectic, niobium foil was inserted between the
carbon blocks and the substrates. A K-type thermocouple, spot-welded to the substrate
no more than 2 mm from the bondline, was used to monitor the bonding temperature.
Bonding was conducted at 1250 ?C, above the binary Fe ? B eutectic temperature
of 1,174 ?C in a 1.3 MPa vacuum atmosphere. Joining was performed using the base
metal in both the unrecrystallized fine grain (FG) and recrystallized coarse grain (CG)
forms. Compressive stresses of 1-5 MPa were used to extrude the excess liquid formed at
the bondline.
4.2.3 Post Bond Heat Treatments
After the bonding process, post bond heat treatment (PBHT) was conducted in
order to induce controlled recrystallization across the bondline. PBHT was performed in
a radiantly heated Brew furnace, in a vacuum, under a pressure of 1.3 MPa, at 1300 ?C
and 1385 ?C, the recrystallization temperatures of for MA956 and PM2000 respectively.
4.2.4 Metallographic Preparation
After PBHT, the samples were cut using a Struers Accutom-5 precision cutting
machine and were mounted using a Labopress-3. The samples were ground to a 1200
48
FEPA grit finish and then polished using a 6 ?m diamond spray, followed by a 0.04 ?m
alumina suspension solution. Polishing of the samples was done using a TegraSystem
autopolisher. The samples were then water cleaned and air dried, followed by etching at
room temperature in a solution of 40 vol.% hydrochloric acid and 60 vol.% methanol for
2 - 10 seconds and further cleaned with water.
4.2.5 Microstructural Characterization
The samples, both as-bonded and PBHT, were primarily characterized using light
microscopy (LM) and scanning electron microscopy (SEM), supplemented with SEM
based energy dispersive X-ray spectroscopy (EDS). These studies were conducted on a
JEOL JSM840 instrument operated at an accelerating voltage of 20 kV. Secondary
electron imaging (SEI) was used to collect all SEM micrographs. The SEM based EDS
analysis used ultra-thin window detectors and Oxford Instruments ISIS analyzers. EDS
analysis was used to determine the composition of phases observed in the bond-line and
compositional profiles across the joint and detection of porosity or secondary phases.
4.2.6 Oxidation Studies
The excellent high temperature oxidation resistance of ferritic ODS alloys results
from the formation of the stable, adherent alumina scale formed during the high
temperature exposure. The yttria present in the matrix plays a significant role in the
adhesion of the alumina scale [42-46]. In order to study the effect of TLP bonding on the
oxidation resistance of the ODS alloy (PM 2000), oxidation tests were conducted on the
bulk material to provide a baseline for comparison.
Samples with dimensions of 10 mmx9 mm and 2 mm in thickness were cut using
EDM and the Accutom-5. The samples were surface ground to remove the oxide layer.
49
The samples were exposed to air at 1200
0
C for 6, 12, 24, 48, 96 hours. The weight of the
samples was measured before and after oxidation. The samples were then cut using the
Accutom-5 apparatus mounted using Labopress-3, and polished as explained in section
4.2.4. The bonds were then characterized using SEM, EDS, and LM.
4.2.7 Mechanical Testing
Shear testing was perfomed using the rig designed by Yan and Wallach, as
explained previously in section 4.1.7.1.
.
50
5 RESULTS AND DISCUSSION
5.1 Nickel-based superalloys
5.1.1 Microstructural Characterization
5.1.1.1 Porosity at the bondline
The results of bondline porosity for joints incorporating the two different
composite interlayers Niflex-110 and Niflex-115, were compared to those joints obtained
using a conventional BNi-3 interlayer.
5.1.1.1.1 Niflex-110 interlayer:
Microstructural studies showed that all of the TLP bonds between CMSX-4 - IN
738 and CMSX-4 ? IN 939 with a Niflex-110 wide-gap style interlayer formed an
eutectic along the grain boundaries of the Niflex core, as shown in Figures 9(a) and 10(a).
After allowing time for isothermal solidification and solid-state homogenization, non-
bonded regions were seen at the substrate-interlayer interface.
This grain boundary eutectic formation infers that a competition existed between
wetting of the faying surfaces and eutectic formation in the grain boundaries of the Niflex
core, thus leaving non-bonded regions at the interface. Note that, once formed, these
pores did not change with time at the bonding or post-bond heat treatment temperature
51
and so probably are not Kirkendall pores. The pores were visible before etching and so
are not etching artifacts, although etching did change the shape of the pores somewhat by
removing adjacent ?? and hence leaving a facetted appearance. Additionally, the grain
boundary dissolution of the composite interlayer leads to an increase in the solid-liquid
interfacial area, permitting the interlayer elements to readily diffuse from the liquid phase
into the interlayer?s matrix, ultimately reducing the amount of liquid available for
spreading, resulting in porosity at the substrate-interlayer interface.
As mentioned in the literature review, a key parameter in the wide gap bonding
process is the selection of the ratio of liquid former to non-melting phase for the
substrates to be joined. Thus, if insufficient liquid former is employed, the liquid will
begin to isothermally solidify, and hence cease to spread, before it has been able to
penetrate throughout the joint. This would result in a joint with porosity, as shown in
Figures 9(b) and 10(b).
52
(a)
(b)
Figure 9 - SEM micrographs in SEI mode, of CMSX-4 -Niflex-110 ? IN 939 joint
showing (a) grain boundary eutectic formation after 0 minutes at 1160
0
C and (b)
porosity on the substrate- interlayer interface in the initial stages after 4 minutes at
1160
0
C. The secondary phases in the diffusion zone of the polycrystalline substrate
are also shown
CMSX-4
Niflex-110
IN 939
10 ?m
CMSX-4
Niflex-110
IN 939
10 ?m
Eutectic
Pores
Pores
Secondary Phases
Secondary Phases
53
Figure 10(a)
Figure 10(b)
Figure 10 -SEM micrographs in SEI mode, of CMSX-4 -Niflex-110 ? IN 738
joint showing (a) grain boundary eutectic formation after 0 minutes at 1160
0
C and (b) porosity on the substrate- interlayer interface after 4 minutes at
1160
0
C. The secondary phases in the diffusion zone of the polycrystalline
substrate were also shown.
CMSX-4
Niflex-110
IN 738
10 ?m
CMSX-4
Niflex-110
IN 738
10 ?m
Eutectic
Secondary Phases
Pores
Pores
54
5.1.1.1.2 Niflex-115 interlayer:
The microstructural studies of the bonds made using Niflex-115 interlayer also
showed the formation of eutectic along the grain boundaries of Niflex core, as illustrated
in Figure 11(a). However, no porosity was seen in the bonds made using a Niflex-115
interlayer. This can be attributed to the higher boron content in the Niflex-115 interlayer
available for wetting of the faying surfaces as compared to Niflex-110. Although, eutectic
formation was seen at the grain boundaries of Niflex-115, the greater amounts of liquid
present at the interface eliminated the problem of non-bonded regions. The same
behavior was observed both in CMSX-4 ? IN 738 and CMSX-4 ? IN 939 joints.
5.1.1.1.3 BNi-3 interlayer:
Microstructural studies of bonds employing a conventional BNi-3 interlayer
showed that the bonds were free of non-bonded regions as shown in Figure 12(a). This
can be attributed to the nature of the interlayer. BNi-3 is an interlayer with a completely
melting core and boron distributed throughout the foil, unlike the two composite
interlayers which had boron only at the outer surfaces. Thus, the amount of liquid formed
at the bondline is higher than that obtained for a composite interlayer and this helps in
better wetting of the faying surfaces, thereby resulting in the elimination of non-bonded
regions. Although BNi-3 foil, as expected, did not show any porosity or premature
isothermal solidification, formation of borides at the bondline was observed with the
BNi-3 foil interlayer. This will be discussed in detail in Section 5.1.1.2.
Although it is desirable from the standpoint of reducing the amount of MPD that
must be diffused, reducing the boron content of the interlayer results in an insufficient
55
boron content, which leads to insufficient liquid for wetting and to premature isothermal
solidification. On the other hand, a high boron content might result in the formation of
borides at the bondline and in the diffusion zone of the substrates. Therefore, it is
necessary to optimize the amount of boron in the interlayer. BNi-3 contains boron
throughout its thickness, whereas Niflex-115 and Niflex-110 contain boron-rich surfaces
and a boron free core; the net effect is a reduction of the amount of boron that must be
diffused per unit gap width. In summary, higher boron content foil (BNi-3) produced
good joint after wetting, but the comparatively low boron content foil Ni-Flex 110,
showed a significant amount of porosity at the substrate-interlayer interface. This
problem was overcome by using a relatively boron-rich Niflex-115 wide gap style
composite interlayer.
56
(a)
(b)
Figure 11- SEM micrographs in SEI mode, of CMSX-4 -Niflex-115 ? IN 738 joint
showing (a) grain boundary eutectic formation after 0 minutes at 1160
0
C and (b)
substrate- interlayer interface with no porosity after 60 minutes at 1160
0
C. The
secondary phases in the diffusion zone of the polycrystalline substrate are also
shown.
Secondary Phases
CMSX-4
Niflex-115
IN 738
10 ?m
CMSX-4
Niflex-115
IN 738
10 ?m
Eutectic
57
(a)
(b)
Figure 12 ? (a)LM micrograph of CMSX-4 ? BNi-3 ? IN 738 after 0 minutes at 1160
0
C, showing borides at the bondline. No grain boundary eutectic formation was
observed. (b) SEM micrograph in BEI mode after 0 minutes at 1160
0
C, showing
borides at the bondline. The substrate- interlayer interface here is free of non-
bonded regions.
10 ?m
CMSX-4
BNi-3
10 ?m
CMSX-4
BNi-3
IN 738
IN 738
Borides
Borides
58
5.1.1.2 Bondline Boride Formation
Borides are hard, brittle intermetallics formed during the heat ramp up stage of the
bonding process and/or during cooling of the joint. These secondary phases formed in the
joint may have undesirable effects on bond properties, for example decrease in bond
strength.
Microstructural investigations of bonds made using BNi-3 foil as an interlayer
confirmed the formation of borides at the bondline during the initial stages of bond
formation as shown in Figure 12(a) and (b). Even after holding for sufficient time for
isothermal solidification and solid-state homogenization to occur, residual borides were
still visible at the bondline. In contrast, bondline boride formation was suppressed by
using wide-gap style interlayers, as shown in Figures 9, 10, and 11.
The suppression of borides in composite interlayers can be attributed to their
structure, with boron present on either side of the non-melting core, unlike the
conventional BNi-3 foil interlayer, where boron is distributed uniformly throughout the
interlayer with a fully melting core. Also, the amount of boron present in the composite
interlayers is relatively low compared to a BNi-3 foil interlayer. Hardness tests on the
bondline of the joints made using a BNi-3 foil interlayer resulted in high hardness values,
as shown in Figure 13. Microstructural investigations using SEM on the tested samples
revealed that borides were present at the joint.
5.1.1.3 Microstructural Bond Development
As mentioned in section 4.2.2, interrupted bonding runs were conducted in
order to study the bond microstructural development. Bond development in the initial
stages (0 minutes to 10 minutes), intermediate stages (60 min to 120 min), and final
59
stages (120-240 min) at bonding temperature (1160
0
C), are discussed in turn below. The
initial stage represents the eutectic formation and substrate dissolution. The intermediate
stage represents isothermal solidification, where the MPD elements start diffusing into
the substrates. The final stage represents solid stage homogenization that produces a
uniform solute concentration at the bondline. Table 5 presents a summary of the
microstructures observed in these three stages.
In the initial stages, when the wide gap style composite interlayer Niflex-110 was
used to join CMSX-4 to IN 738 and IN 939, a eutectic was formed in the grain
boundaries of the Niflex core as mentioned earlier in Section 5.1.1.1. [Figure 9(a) and
Figure 10(a)]. As a result, a competition existed between liquid wetting the faying
surfaces and the grain boundary eutectic formation which resulted in porosity in the
bonds, that was overcome by using Niflex-115, which has boron content relatively higher
than that of Niflex-110. When a BNi-3 interlayer was used to join CMSX-4 to IN 738
and IN 939, no porosity was visible at the substrate-interlayer interface due to the
uniform distribution of boron and better wetting [Figure 12].
The wide-gap style interlayers Niflex-110 and Niflex-115 also resulted in the
suppression of bondline boride formation. In comparison, the conventional BNi-3
interlayer resulted in substantial boride formation at the bondline in the initial stages
(Figure 12), which was confirmed by a back scattered image of a CMSX-4 ? IN 738 joint
obtained using BNi-3 interlayer. In addition, a variety of second phases which are
assumed to be borides/carbides/TCP phases were formed in the substrate diffusion zone
in the joints made using both composite interlayers and the conventional BNi-3
60
interlayer. This formation of secondary phases was observed in all of the interlayer and
substrate combinations in the initial stages, and is shown in Figures 9 through 11.
In the intermediate stages (60 min-120 min), once time has been allowed for
isothermal solidification to occur, the porosity at the bondline in Niflex-110 interlayer
bonds was not eliminated, unlike that in bonds with the other two interlayers (Niflex-115
and BNi-3). This can be attributed to the competition mentioned earlier and the lower
boron content present in the Niflex-110 interlayer compared to Niflex-115. On the other
hand, in bonds using a BNi-3 interlayer, although some of the boron diffused due to
isothermal solidification, borides were still present in the intermediate stages. In addition,
secondary phases in the diffusion zone were present in the case of the CMSX-4 ? IN 738
joint with BNi-3 foil, unlike the other two interlayers. However, for CMSX-4 ? IN 939
joints, secondary phases in the diffusion zone were present for all of the interlayer
combinations.
In the final stages of the bond formation, use of the Niflex-110 interlayer resulted
in a joint with significant porosity at the substrate-interlayer interface as shown in Figure
14, secondary phases in the diffusion zone of the polycrystalline substrate in the case of
the IN 939 joint, and with no borides at the bondline as shown in Figure 15. The other
composite interlayer, Niflex-115, resulted in a joint with no porosity and no bondline
borides although secondary phase formation was seen in the polycrystalline substrate
diffusion zone when IN 939 was used as the polycrystalline substrate. (Figure 16(a) and
the points of interest highlighted in Table 5). In contrast, BNi-3 foil interlayer bonds had
numerous borides present at the bondline, and secondary phases present in the diffusion
zone, although there was no porosity [Figure 16(b)]. A compositional analysis of the
61
composite interlayer bonds revealed a fairly uniform composition across the CMSX-4 ?
Niflex-110 - IN 738 andCMSX-4 ? Niflex-110 - IN 939 bondline after 240 minutes of
bondtime at 1160
0
C, as shown in Figure 18 and Figure 19.
5.1.1.4 Secondary Phases in the diffusion zone
As mentioned in previous sections, a variety of second phases, which appear to
include borides/carbides and/or topologically close-packed (TCP) phases, were observed
to form in the substrate diffusion zone on the polycrystalline side of both the CMSX-4 ?
IN738 and CMSX-4 ? IN939 bonds, as shown in Figure 15 and Figure 16 respectively. In
the case of bonds involving IN738, the formation of these second phases could be
suppressed by increasing the bonding time to 4 h at 1160
o
C when using either of the two
composite interlayers although these phases could not be avoided in the bonds using BNi-
3 foil as an interlayer, as shown in Figure 15. The compositional analysis done on a
CMSX-4 ? Niflex-110 - IN 738 joint after 4 minutes of bonding time can be seen in
Figure 17. Chromium and cobalt peaks were observed in EDS spectra from the diffusion
zone of the polycrystalline substrate, which can be directly related to the secondary
phases seen in the microstructural studies.
In contrast, when bonding IN939, the second phases remained stable for all
times examined and for all the interlayers, including post-bond thermal exposures of up
to 1 week at 1,000
o
C, which will be discussed in Section 5.1.5.2. These secondary
phases also had an impact on the mechanical properties of the bonds which will be
discussed in section 5.1.4. This is perhaps due to the high chromium content present in IN
939, since chromium can be involved in the formation of both borides and TCP phases.
62
Substrate Interlayer
Eutectic
formation in grain
boundaries (gbs)
of Niflex core
Porosity
Bondline
boride
formation
Secondary
phases in
diffusion zone
of PC
CMSX-4 - IN 738 Yes Yes No Yes
CMSX-4 - IN 939 Yes Yes No Yes
CMSX-4 - IN 738 Yes No No Yes
CMSX-4 - IN 939 Yes No No Yes
CMSX-4 - IN 738 No No Yes Yes
CMSX-4 - IN 939 No No Yes Yes
Substrate Interlayer
Eutectic
formation in gbs
of Niflex core
Porosity
Bondline
boride
formation
Secondary
phases in
diffusion zone
of PC
CMSX-4 - IN 738 No Yes No No
CMSX-4 - IN 939 No Yes No Yes
CMSX-4 - IN 738 No No No No
CMSX-4 - IN 939 No No No Yes
CMSX-4 - IN 738 No No Yes Yes
CMSX-4 - IN 939 No No Yes Yes
Substrate Interlayer
Eutectic
formation in gbs
of Niflex core
Porosity
Bondline
boride
formation
Secondary
phases in
diffusion zone
of PC
CMSX-4 - IN 738 No Yes No No
CMSX-4 - IN 939 No Yes No Yes
CMSX-4 - IN 738 No No No No
CMSX-4 - IN 939 No No No Yes
CMSX-4 - IN 738 No No Yes No
CMSX-4 - IN 939 No No Yes Yes
Initial stages (0-10 minutes)
Intermediate stages (60-120 minutes)
Niflex-110
Niflex-115
Niflex-110
Niflex-115
BNi-3
BNi-3
BNi-3
Final stages (240 minutes)
Niflex-110
Niflex-115
Table 5: Microstructural bond development, for all substrate-interlayer combinations
(CMSX-4 ? IN 738 and CMSX-4 ? IN 939 with Niflex-110, Niflex-115, BNi-3) {Italics
signify the microstructural features that influenced mechanical properties}
63
200
250
300
350
400
450
-600 -400 -200 0 200 400 600
The distance from the center bond line (?m)
Mi
croh
ardn
e
s
s
Niflex-110 BNi-3
Figure 13: Comparison of Vickers microhardness across bondline of CMSX-4 ? IN
939 bond after 240 minutes at 1160
0
C using Niflex-110 and BNi-3 foil interlayers.
Note the high hardness values for the BNi-3 joint, that might be due to the borides
present.
64
(a)
(b)
Figure 14: SEM micrographs in SEI mode, of CMSX-4 ? IN 738 bonds after 240
minutes at 1160
0
C showing (a) porosity at the bondline in Niflex-110 and (b) bond
free from porosity and free from secondary phases in the diffusion zone, using
Niflex-115 interlayer.
CMSX-4
Niflex-110
IN 738
10 ?m
CMSX-4
Niflex-115
IN 738
10 ?m
Pores
65
Figure 15: SEM micrographs in SEI mode, of CMSX-4 ? BNi-3 - IN 738 bond after
240 minutes at 1160
0
C, showing borides at the joint and secondary phases in the
diffusion zone of IN 738.
CMSX-4
BNi-3
IN 738
10 ?m
Secondary Phases
Borides
66
(a)
(b)
Figure 16: SEM micrographs in SEI mode, of CMSX-4 ? IN 939 bond after 240
minutes at 1160
0
C showing (a) secondary phases in diffusion zone of polycrystalline
substrate in Niflex-115 interlayer bond, and (b) borides at the joint and secondary
phases in the diffusion zone of IN 939, for bonds using BNi-3 interlayer.
10 ?m
CMSX-4
Niflex-115
IN 939
10 ?m
CMSX-4
BNi-3
IN 939
Secondary Phases
Borides
Secondary Phases
67
(a)
0
10
20
30
40
50
60
70
80
-60 -40 -20 0 20 40 60
Distance across the bondline (microns)
C
o
mp
o
s
it
io
n
s
(
a
t
%
)
Al
Ti
Cr
Ni
Co
(b)
Figure 17: Composition profile, obtained using SEM-based EDS analysis, across
bondline of CMSX-4 ? Niflex-110 ? IN 738 joint after 4 minutes of bonding time at
1160
0
C Note the chromium and cobalt peak observed at the diffusion zone of the
polycrystalline substrate that might be carbides, borides, TCP phases formed.
CMSX-4
10 ?m
Niflex-110IN 738
IN 738 CMSX-4
68
0
10
20
30
40
50
60
70
80
-60 -40 -20 0 20 40 60
Distance across the bondline (microns)
Co
mp
ositio
ns (at%)
Al
Ti
Cr
Ni
Co
Figure 18: Composition profile obtained using SEM-based EDS analysis, across
bondline of CMSX-4 ? Niflex-110 ? IN 738 joint after 240 minutes of bonding time
at 1160
0
C. Notice that the composition of the joint is fairly uniform after allowing
time for isothermal solidification.
IN 738 CMSX-4
69
0
10
20
30
40
50
60
70
80
-60 -40 -20 0 20 40 60
Distance across the bondline (microns)
Compos
itions (a
t%
)
Al
Ti
Cr
Ni
Co
Figure 19: Composition profile obtained using SEM-based EDS analysis, across
bondline of CMSX-4 ? Niflex-110 ? IN 939 joint after 240 minutes of bonding time
at 1160
0
C. Notice that the composition of the joint is fairly uniform after allowing
time for isothermal solidification.
IN 939 CMSX-4
70
5.1.2 Comparison with the wettability studies
The results of the microstructural investigations performed for this study were
correlated with the results of wettability studies conducted by Chitti et al. [131]. The
eutectic formation in the grain boundaries of the Niflex core, which was in competition
with the wetting of the faying surfaces of the composite interlayers, was also evident in
the cross sections of the wetted samples, as shown in Figure 20 and Figure 21. As
mentioned in Section 5.1.1.1.1, the non-melting core acts as a sink for boron diffusion.
With a Niflex-110 interlayer, due to the competition between the wetting of the faying
surfaces and the grain boundary liquid formation, the available boron content was
insufficient to wet the faying surfaces, leading to premature isothermal solidification. In
addition, the grain boundary dissolution of the composite interlayer led to an increase in
the solid-liquid interfacial area. This resulted in an increase in the rate of boron loss from
the liquid and consequent diffusion into the substrate matrix, ultimately reducing the
amount of liquid available for spreading and resulting in porosity at the substrate-
interlayer interface.
Wettability studies also confirmed the formation of secondary phases in the
diffusion zone of the polycrystalline substrate. This can be seen in Figure 20(b) and
Figure 21(b). Due to the high heating rates employed in the wettability studies, the
formation of borides/carbides during heating ramp up would be minimal, so the
secondary phases such as borides and carbides are likely to have been formed during
cooling of the joint to room temperature.
Wettability studies also confirmed the formation of bondline borides. In general,
high heating rates [high temperature increments per unit time] are employed to avoid the
71
formation of borides during the heating stage. The hot stage light microscope with the
high heating rates employed in the wettability studies, minimized boride formation during
heating to the bonding temperature and thus any borides formed during cooling could be
identified.
5.1.2.1 Effect of substrate on wettability
Wettability studies were conducted on single crystal CMSX-4 in unrecrystallized
and recrystallized condition and on polycrystalline IN 738 and IN 939. All the three
interlayers exhibited a greater extent of spreading on the single crystal alloy CMSX?4
when compared to the polycrystalline alloys IN 738 , IN 939 and recrystallized CMSX-4
[referred to here as CMSX-4 (R)] [Figure 23, Figure 24 and Figure 25]. The reduction in
the extent of spreading of the liquid in CMSX-4 (R) compared to CMSX-4 can be
attributed to two factors. The polycrystalline/recrystallized alloy contains grain
boundaries, which are potential paths for the diffusion of interlayer species (boron),
allowing it to easily diffuse into the substrate. In addition, substrate dissolution also
occurs at the substrate grain boundaries which increases the solid-liquid interfacial area
and hence hastens isothermal solidification. Note that as the homologous temperature
increases, the relative role of preferential grain boundary diffusion as compared to bulk
diffusion decreases. Given that bonding was conducted at approximately 85-90 % of the
absolute melting point of the substrates, preferential grain boundary diffusion is unlikely
to have made a very significant contribution to microstructural development. In contrast,
grain boundary dissolution can have a marked effect.
In addition, the time taken for significant premature isothermal solidification and
hence the termination of spreading was much lower for composite interlayers (Niflex-110
72
and Niflex-115 foils), when compared to the BNi-3 foil. [Refer to the difference in
termination of spreading time in Figure 23, Figure 24 and Figure 25]. For instance, the
termination of spreading for BNi-3 interlayer on CMSX-4 occurred approximately at 190
seconds, in contrast to 50 seconds and 115 seconds approximately for Niflex-110 and
Niflex-115 interlayers respectively. This was because of the smaller amount of liquid
available for wetting, as well as the preferential flow of the eutectic liquid into the
interlayer grain boundaries in the composite interlayers, creating a competition in wetting
between the interlayer non-melting core and the substrate itself. Figure 26 shows the
spreading measurements of BNi-3 foil on CMSX-4 in two different runs under identical
conditions, showing repeatability of spreading measurements.
73
(a)
(b)
Figure 20: LM micrograph, showing eutectic formation in the grain boundaries of
inner Niflex core for a Niflex-110 interlayer wetted on (a) CMSX-4 and (b) IN 738.
The secondary phases formed in the diffusion zone of the polycrystalline substrate
IN 738 are also visible [121].
IN 738
Niflex-110
20 ?m
Niflex-110
20 ?m
CMSX-4
Secondary Phases
Eutectic
Eutectic
74
(a)
(b)
Figure 21: LM micrograph, showing eutectic formation in grain boundaries of
inner Niflex core for Niflex-110 interlayer wetted on (a) CMSX-4 and (b) IN 939.
The secondary phases formed in the diffusion zone of the polycrystalline substrate
IN 939 are visible [121].
Niflex-115
20 ?m
Niflex-115
20 ?m
IN 738
IN 939
Eutectic
Eutectic
Secondary Phases
75
Figure 22: LM micrograph, of the cross-section of IN 939 wetted with BNi-3 foil
showing the formation of borides at the joint interface [121].
IN 939
BNi-3
76
BNi-3
0
0.2
0.4
0.6
0.8
1
1.2
1.4
1.6
1.8
2
0 20 40 60 80 100 120 140 160 180 200
Time (s)
S
p
r
e
ad
i
n
g
D
i
a
m
et
e
r
(
m
m
)
BNi-3 on CMSX-4(R) BNi-3 on IN 738 BNi-3 on IN 939 BNi-3 on CMSX-4
Figure 23: Wettability of BNi-3 on CMSX-4, IN 738 and IN 939 CMSX-4(R). [121]
77
NiFlex-110
0
0.2
0.4
0.6
0.8
1
1.2
1.4
1.6
1.8
0 5 10 15 20 25 30 35 40 45 50 55 60 65 70
Time(s)
Spreadin
g Dia
m
eter(
m
m
)
Niflex-110 on IN 939 Niflex-110 on IN 738
Niflex-110 on CMSX-4 Niflex-110 on CMSX-4(R)
Figure 24: Wettability of Niflex-110 on CMSX-4, IN 738, IN 939 and CMSX-4(R).
Compare this to previous Figure and notice the reduction in time required for
termination of spreading compared to BNi-3 [121].
78
NiFlex-115
0
0.2
0.4
0.6
0.8
1
1.2
1.4
1.6
1.8
2
0 20 40 60 80 100 120
Time(s)
S
p
read
i
n
g
Di
ame
t
er(
mm)
Niflex-115 on CMSX-4 Niflex-115 on IN 738
Niflex-115 on IN 939 Niflex-115 on CMSX-4(R)
Figure 25: Wettability of Niflex-115 on CMSX-4, IN 738, IN 939 and CMSX-4(R).
Compare this to previous two Figures and notice the reduction in time required for
termination of spreading compared to BNi-3 [121].
79
0
0.2
0.4
0.6
0.8
1
1.2
1.4
1.6
0.0 50.0 100.0 150.0 200.0 250.0
Time (s)
S
p
re
ad
i
n
g
D
i
am
e
t
er(
m
m
)
Run 1 Run 2
Figure 26: Wettability of BNi-3 on CMSX-4 on two different runs in identical
conditions showing repeatability of data. [121].
80
5.1.2.2 Effect of Boron Content
Boron was used as the MPD in all three of the interlayer materials chosen.
Although reducing the boron content of the interlayers is desirable from the standpoint of
reducing the amount of MPD that must be diffused, an insufficient boron content also
creates insufficient liquid for wetting, leading to premature isothermal solidification. On
the other hand, a higher boron content is likely to result in increased isothermal
solidification times leading to higher bonding process times, and also the formation of
borides at the bondline and in the substrates. Therefore, it is important to carefully select
the amount of boron in the interlayer. BNi-3 contains boron throughout its thickness,
whereas Ni-Flex 115 and Ni-Flex 110 are composed of boron-rich surfaces on a boron
free core; the net effect reduces the amount of boron that must be diffused per unit gap
width. Spreading diameter vs. time plots revealed that the extent of spreading of the BNi-
3 foil was greater compared to the other two foils. Post wetting SEM/EDS work on cross-
sectioned samples of higher boron content foil (BNi-3) showed good joint with no
porosity, but the comparatively low boron content foil, Ni-Flex 110 showed substantial
porosity at the substrate-interlayer interface [Figure 9]. However, the problem was
overcome by using a relatively boron-rich Niflex-115 interlayer.
5.1.2.3 Effect of Boride Formers
The primary boride formers in superalloys are molybdenum, titanium, chromium
and nickel with molybdenum and titanium forming the most stable borides, followed by
chromium and then nickel. Boride formation tends to cause premature isothermal
solidification during bonding leading to porosity at the faying surfaces. Post wettability
analysis showed the presence of significant amounts of titanium and chromium at the
81
outer rings of the solidified liquid on the IN 939 substrates which corresponded to
borides, visible in cross-sections. This showed that some interdiffusion of the substrate
and interlayer species did occur.
5.1.3 Gamma-prime at the bondline
In all the bonds examined, there was relatively little gamma-prime at the bondline
when compared with the bulk substrates. This might be due to the lack of heat treatments
to reform the ??, in the bonding cycle. For a given wide-gap interlayer, more gamma-
prime was formed at the bondline of CMSX-4 ? IN939 when compared to that of CMSX-
4 ? IN738, as shown in Figure 27 and Figure 28. This can be attributed to the difference
in solution temperatures of the alloys (IN 939 and IN 738) and the cooling rates which
might have affected the formation of ??. The amount of gamma-prime formed at the
bondline, did however have an impact on the mechanical properties of the bond, which
will be discussed in the next section. This is suprising, given that IN 939 is lower in Al
(1.9 wt% versus 3.4 wt%) and higher in chromium (22.4 versus 16 wt%) than IN 738,
and it can only be assumed that the presence of ?? formers other than Al and ? stabilizers
other than Cr must have influenced the overall behavior of the bonds.
82
(a)
(b)
Figure 27: SEM micrographs in SEI mode, of CMSX-4 ? Niflex-110 ? IN 939 bond
showing (a) gamma-prime (??) at the bondline (b) Polycrystalline substrate IN 939 of
the bond with no resolvable (??).
1 ?m
CMSX-4
Niflex-110
1 ?m
IN 939
gamma-prime
83
(a)
(b)
Figure 28: SEM micrographs in SEI mode, of CMSX-4 ? Niflex-110 ? IN 738 bond
showing (a) gamma-prime (??) at the bondline (b) Polycrystalline substrate IN 738 of
the bond with no resolvable (??).
Notice the smaller amount of ?? on IN 738 bondline when compared to IN 939
bondline [Figure 27]
1 ?m
CMSX-4
Niflex-110
1 ?m
IN738
gamma-prime
84
5.1.4 Structure-Property Relationships of As-bonded TLP bonds
5.1.4.1 Shear tests
Before conducting shear tests on as-bonded samples, bulk samples were machined
from as-received substrates and tested using a shear rig. For this study, IN 939 and IN
738 substrates were machined and shear tested, to serve as a baseline for a comparison of
the bond strengths for CMSX-4 ? IN 939 and CMSX-4 ? IN 738 bonds respectively.
Figure 29 shows fractographic images of a shear tested bulk IN 939 and IN 738 samples.
The fracture path was macroscopically flat and was dominated by a ductile shear fracture
with dimpling.
Shear testing conducted on as-bonded CMSX-4 ? IN 939 resulted, on an average
of 92% of the bulk shear strength and 69% of the UTS of the weakest bulk material in
the bond (in this case IN 939) [Table 6]. Shear testing of CMSX-4 ? IN 738 resulted, on
an average of 73 % of the bulk shear strength and 55% of the UTS of the weakest bulk
material in the bond [Table 7].
A schematic showing the crack path upon application of shear force is given in
Figure 30. Figure 32 shows fractographic images of shear tested CMSX-4 ? Niflex-110 ?
IN 939 and CMSX-4 ? Niflex-110 ? IN 738 bonds. A ductile shear fracture was seen on
the bondline in all of the bonds. In the CMSX-4 ? IN 939 bonds, secondary cracking was
seen on the fracture surfaces in addition to the ductile shear, which is also shown in
Figure 31. In addition, no secondary cracking, but rather dimpling, was seen on the
fracture surfaces of the CMSX-4 ? IN 738 bonds as shown in Figure 32.
The shear testing of the as-bonded specimens correlated well with the
microstructural investigations. The ductile shear fracture observed in all of the bonds can
85
be attributed to the insufficient ?? formation at the bondline. As mentioned earlier, the
shear strengths of bonds between CMSX-4 ? IN939 are higher than those of CMSX-4 ?
IN738 bonds. The higher shear strength of CMSX-4 ? IN 939 relative to CMSX-4 - IN
738 using a Niflex-110 interlayer might be due to the greater amount of ?? formed at the
bondline in the former. In addition, the secondary phases in the diffusion zone of the
polycrystalline substrate of the bondline might have resulted in secondary brittle cracking
in addition to the main ductile shear fracture for all the CMSX-4 ? IN 939 bonds. In
contrast, CMSX-4 ? IN 738 bonds free of secondary phases in the diffusion zone showed
no secondary cracking but rather dimpling on the fracture surfaces, as shown in Figure
32. No evidence was found for a role of the porosity observed at the faying surfaces in
the shear strength of the bonds. However, it should be cautioned that this porosity could
have a large effect on fatigue resistance, although this was not investigated here.
From the above observations, on CMSX-4 ? IN 939 and CMSX-4 ? IN 738, the
extent of bondline gamma-prime formation (relatively high in bonds with IN 939, low in
bonds with IN 738) would appear to be more important for the room-temperature
properties than the presence of brittle second phases in the polycrystalline substrate
(which are invariably present in IN 939, but avoidable with a suitable bonding treatment
in IN 738). Also, fatigue tests conducted on aluminide coated superalloys by Veys et al.,
and Grunlig et al., [122, 123] revealed that, brittle secondary phases present in the
transition zone (similar to the secondary phases present in the diffusion zone in the
current observations) were expected to change the final crack path and behavior of
fatigue crack originated at a surface [124].
86
5.1.4.2 Hardness testing
The hardness profile of the bonds is presented in Figure 33, which shows that the
Vickers microhardness of the bond line is similar to that of the IN 939, while the CMSX-
4 alloy has a higher microhardness than the bond line and the IN 939. This infers that the
formation of insufficient gamma-prime within the joint left a relatively soft bondline
region, as is evident from the hardness data for the bond. Bonds made using a BNi-3 foil
interlayer resulted in higher hardness values at the bondline compared to the bonds made
using Niflex-110 and Niflex-115 interlayers. This can be attributed to the formation of
borides at the bondline when the BNi-3 interlayer was used.
87
(a)
(b)
Figure 29: SEM micrograph in SEI mode, showing fracture surface of as-received
(a) IN 939 bulk material and (b) IN 738 bulk material. The fracture path was
macroscopically flat and dominated by ductile shear fracture with dimpling.
10 ?m
20 ?m
IN738
IN 939
88
CMSX-4/IN 939 %Bulk Shear strength (IN939) % of UTS of Bulk IN 939
Sample ID Bulk Shear strength=783.25 UTS Bulk =1050 MPa [2]
1 8.00x3.7 741.82 0.95 0.71
2 8.00x3.7 778.9 0.99 0.74
3 8.12x3.88 743.4 0.95 0.71
4 8.12x4.100 695.1 0.89 0.66
5 8.12x4.100 672.1 0.86 0.64
Mean -------- 726.3 0.93 0.69
Std.Dev. -------- 42.5 0.05 0.04
6(low load fail) 8.00x3.750 189.8 0.24 0.18
CMSX-4/IN 939
1 7.81x3.96 709 0.91 0.68
2 7.79x3.98 673.9 0.86 0.64
3 7.85x3.94 741.5 0.95 0.71
4 7.83x3.87 664.8 0.85 0.63
5 7.85x3.94 806.4 1.03 0.77
6 8.06x3.89 681.1 0.87 0.65
Mean 712.8 0.91 0.68
Std.Dev. 53.7 0.07 0.05
CMSX-4/IN 939
1 8.12x4.00 605.9 0.77 0.58
2 8.22x4.00 571 0.73 0.54
3 7.91x3.88 604.6 0.77 0.58
4 8.33x3.98 582.9 0.74 0.56
5 8.06x3.86 591.7 0.76 0.56
6 7.96x4.00 544.5 0.70 0.52
Mean 583.43 0.74 0.56
Std.Dev 23.20 0.03 0.02
Asbonded + PBTE [CMSX-4 - IN 939]
Asbonded CMSX-4 - IN 939
Asbonded +PBHT [CMSX-4 - IN 939]
Dimensions
(mm)
Peak stress
(MPa)
Table 6: Shear testing data of CMSX-4 ? IN 939 bonds using Niflex-110 interlayer at
different conditions. PBHT- Post bond heat treatments as mentioned in section 4.1.3
and PBTE- post bond thermal exposure as explained in section 4.1.4.
89
CMSX-4/IN 738 %Bulk Shear strength IN 738=785.6 %UTS Bulk IN 738=1035MPa
1 8.24x4.020 617.6 0.79 0.60
2 8.24x3.88 509.8 0.65 0.49
3 8.00x3.870 589.3 0.75 0.57
4 8.3x3.83 618 0.79 0.60
5 8.04x3.94 617.8 0.79 0.60
6 8.23x3.84 496.6 0.63 0.48
Mean 574.85 0.73 0.56
Std.Dev 56.74 0.07 0.05
CMSX-4/IN 738
1 7.93x3.85 727.1 0.93 0.70
2 8.13x3.95 628 0.80 0.61
3 7.93x3.97 819.7 1.04 0.79
4 7.93x3.95 728.4 0.93 0.70
5 7.79x4.03 649.1 0.83 0.63
Mean 710.5 0.90 0.69
Std.Dev 76 0.10 0.07
CMSX-4/IN 738
1 7.82x4.00 492.8 0.63 0.48
2 8.040x4.00 501.8 0.64 0.48
3 8.06x3.88 511.5 0.65 0.49
4 8.06x3.88 571.5 0.73 0.55
5 8.02x4.000 483.9 0.62 0.47
6 7.93x3.94 544.5 0.69 0.53
Mean 517.7 0.66 0.50
Std.Dev 33.7 0.04 0.03
Asbonded + PBTE [CMSX-4 - IN 738]
Asbonded CMSX-4 - IN 738
Asbonded +PBHT [CMSX-4 - IN 738]
Table 7: Shear testing data of CMSX-4 ? IN 738 bonds using Niflex-110 interlayer at
different conditions. PBHT- Post bond heat treatments as mentioned in section 4.1.3
and PBTE- post bond thermal exposure as explained in section 4.1.4.
90
Figure 30 : Schematic showing crack path upon application of shear force on the
bonded samples
CMSX-4
IN 939/ IN 738
Bond
~~~~~~~~~~~~~~~~~~~~~~~~~~~~
crack
Secondary
phases
91
(a)
(b)
Figure 31: SEM micrograph in SEI mode, of shear tested CMSX-4 ? Niflex-110 ? IN
939 bond fracture surface of (a) IN 939 polycrystalline substrate (b) CMSX-4 single
crystal substrate. Notice the secondary cracking on IN 939 substrate due to the
secondary phases precipitated in the diffusion zone of the polycrystalline substrate.
10 ?m
10 ?m
CMSX-4
IN 939
Brittle Cracking
92
(a)
(b)
Figure 32: SEM micrograph in SEI mode, of shear tested CMSX-4 ? Niflex-110 ? IN
738 bond fracture surface of (a) IN 738 polycrystalline substrate (b) CMSX-4 single
crystal substrate. Notice the ductile shear fracture on both of the substrates and
dimpling on IN 738.
10 ?m
10 ?m
CMSX-4
IN 738
Ductile Cracking
93
200
250
300
350
400
450
-600 -400 -200 0 200 400 600
The distance from the center bond line (?m)
M
i
cr
oh
ard
n
e
s
s
CMSX-4 - IN 939 CMSX-4 - IN 738
Figure 33: Vickers microhardness across bondline of CMSX-4 ? IN 738 and CMSX-
4 ? IN 939 bonds after 240 minutes at 1160
0
C using Niflex-110 interlayer. Notice
that the bondline is of lower microhardness than that of the substrates.
CMSX-4
IN 738/ IN 939
94
5.1.5 Characterization of Post-Bond Heat Treated TLP bonds
5.1.5.1 Microstructure after PBHT
Post-bond heat-treatment (PBHT), described in Section 4.1.3, was conducted to
diffuse aluminum from the substrates, resulting in more ?? at the bondline and to reform
the substrate ?? distribution which was dissolved during bonding. CMSX-4 ? IN 738
bonds using a Niflex-110 interlayer, after being subjected to a two stage PBHT, resulted
in an increase of volume fraction of ?? and also reformation of ??, which was already
dissolved during the bonding process. This is clearly visible in Figure 34 and Figure 35.
In addition, coarsening of the ?? phase and modification of the shape of the ?? into
circular particles was observed. This might have been caused due to the reduction in total
interfacial energy, as explained by Sharghi-Moshatghin et al. [29]
On the other hand, in CMSX-4 ? IN 939 bonds, the PBHT process did not change
the volume fraction of ?? to any great extent, although it was already greater than in the
CMSX-4 ? IN 738 bonds. The increase in size of the ?? phase and a shape change of ??
from rectangular to circular were also observed [Figure 36 and Figure 37]. Also, the
secondary phases seen in the diffusion zone of the polycrystalline substrate [See Section
5.1.1.4] in the as bonded samples, were still observed and remained stable even after
PBHT, unlike for the CMSX-4 ? IN 738 combination.
5.1.5.2 Microstructure after PBTE
Post-bond thermal exposure (PBTE), as described in Section 4.1.4, was conducted
to investigate the microstructural stability of the bonds. Coarsening of ?? was seen on the
bonds subjected to PBTE. Also, a significant region of deleterious phases, such as sigma,
and Laves, was observed in the polycrystalline substrate of the joint after the high
95
temperature exposure [Figure 38]. In addition, in bonds involving CMSX-4 ? IN 939, the
secondary phases in the diffusion zone of the polycrystalline substrate remained stable for
all times examined for all interlayers, even after post-bond thermal exposure of up to 1
week at 1,000
o
C.
96
Figure 34: SEM micrograph in SEI mode, of As bonded CMSX-4 ? Niflex-110 ? IN
738 showing gamma-prime (??) at the bondline.
1 ?m
CMSX-4
Niflex-110
97
Figure 35: SEM micrograph in SEI mode, of As bonded + PBHT CMSX-4 ? Niflex-
110 ? IN 738 showing gamma-prime (??) at the bondline. Notice the increase of
volume fraction of ?? after PBHT (compare it to previous Figure)
1 ?m
CMSX-4
Niflex-110
98
Figure 36: SEM micrograph in SEI mode, of As bonded CMSX-4 ? Niflex-110 ? IN
939 bond showing gamma-prime (??) at the bondline.
Compare this with Figure 34 and note the high amount of ?? at the bondline and in
the bulk, which was dissolved during bonding.
1 ?m
CMSX-4
Niflex-110
99
Figure 37: SEM micrograph in SEI mode, of as bonded + PBHT CMSX-4 ? Niflex-
110 ? IN 939 showing gamma-prime (??) at the bondline. Notice the modification of
shape of ?? into circular, and the reformed ?? distribution, which was dissolved
during bonding (compare to previous Figure).
1 ?m
CMSX-4
Niflex-110
100
Figure 38: SEM micrograph in SEI mode, of asbonded+PBTE CMSX-4 ? Niflex-110
- IN 738 bond showing coarsened as-bonded gamma-prime (??) and deleterious
phases.
CMSX-4
Bondline
IN 738
20 ?m
Deleterious Phases
Gamma-prime
101
5.1.6 Structure-Property Relationships of TLP Bonds following a PBHT and
PBTE
5.1.6.1 Shear tests
Shear testing was done on CMSX-4 ? Niflex-110- IN 738 and CMSX-4 ?Niflex-
110 - IN 939 bonds after PBHT. Shear testing conducted on as-bonded +PBHT CMSX-4
? IN 738 joint resulted, on average, in 90% of the bulk shear strength of the weakest bulk
material in the bond (IN 738), an increase of about 15% when compared to the strength
of an as bonded joint [Table 7]. This increase in shear strength of CMSX-4 ? IN 738
bonds can be attributed to the increase of volume fraction of ?? at the bondline and the
reformed ?? distribution, after PBHT [Figure 35]. PBHT might have caused diffusion of
aluminum from the substrates, resulting in more ?? at the bondline, and also helped in
reforming the ?? distribution dissolved during bonding.
In contrast, shear testing of CMSX-4 ? IN 939 resulted, on average, in 92 % of
the bulk shear strength, which showed no increase in shear strength, unlike the
comparable IN 738 combination [Table 6]. This might be due to the lack of change in the
volume fraction of ??. However, fractography of bonds subjected to PBHT, showed
secondary cracking in the polycrystalline substrate corresponding to those observed in
microstructure.
Also, bonds that were thermally exposed to 1,000
o
C for 1 week showed a
decrease in shear strength which can be attributed to the coarsening of ?? and the
formation of secondary phases in the polycrystalline substrate. Bonds prepared using IN
939 as a polycrystalline substrate showed similar results, with decrease in the shear
strength after PBHT.
102
0
100
200
300
400
500
600
700
800
900
1000
BULK(AS RECEIVED) AS BONDED(ASB) ASB+PBHT ASB+THERMALLY
EXPOSED
CONDITION OF THE SUBSTRATES
A
VE.
PEA
K
ST
R
E
S
S
(MPa
)
CMSX-4 - IN 939
CMSX-4 - IN 738
Figure 39: Average shear stress versus condition of substrates. ASB refers to
Asbonded, PBHT-post bond heat treatment, PBTE- post bond thermal exposure.
Note that the increase in shear strength in CMSX-4 ? IN 738 after PBHT and the
decrease after PBTE. (Error bars denote the 1.96 times standard deviation Error
bars denote 1.96 times standard deviation to represent a 95% confidence limit)
n=4 n=3 n=5 n=6 n=6 n=5 n=6 n=6
103
5.2 Oxide dispersion strengthened iron based superalloys
5.2.1 Microstructural Characterization
As mentioned earlier, MA956 and PM2000 are the two substrates studied. The
microstructure of diffusion bonds made with the substrates cut in the working direction
will be discussed in this section.
5.2.1.1 Diffusion bonding of MA 956
Diffusion bonding of MA 956 was conducted in both the fine grain and coarse
grain conditions. In the fine grain condition, satisfactory bonds were achieved in the
Longitudinal-Longitudinal (L-L) bond orientation, where the bondline was parallel to the
direction of extrusion of the ODS bulk material, under low stresses, 1 ? 5 MPa at 1250
?C, for bonding times up to 6-8 minutes. The bonding process was followed by PBHT to
induce recrystallization across the bondline. The microstructure of the bond can be seen
in Figure 40. Grain growth across the bond-line was achieved in the fine grain to fine
grain bonds, as can be seen in Figure 41. However, occasional voids were visible at the
bond-line even after the PBHT. In addition, satisfactory bonds were obtained when a fine
grain condition substrate was joined to a coarse grain condition substrate in the L-L
orientation as shown in Figure 42. However, when two MA956 substrates in the coarse
grain condition were joined, bonds with large unbonded regions were observed, as shown
in Figure 43.
Based on the above observations, acceptable bonds were obtained only when a
fine grain condition material was involved. This can be attributed to the huge amounts of
stored energy possessed by unrecrystallized (fine grained) ODS alloys [129]. In their
studies on diffusion bonding, Bokstein et al. and Orhan et al. [131] explained that a
104
submicron grain size can lead to higher grain boundary diffusion, and during the bonding
process the grain boundaries intersect these voids. Also, these grain boundaries may act
as favorable paths for the atoms to diffuse to the voids during the diffusion bonding
process, thereby resulting in closing of the voids and, hence an increase in the bonding
rate. Similar results showing that coarse grained substrates result in poor bonds have been
reported in diffusion bonding of MA6000 [102], a Ni-base ODS alloy, and MA956
[118,119] in Transverse-Transverse (T-T) orientation, where the bondline was
perpendicular to the direction of extrusion of the ODS bulk material.
5.2.1.2 Diffusion bonding of PM2000
Diffusion bonding of PM2000 was conducted for the fine grain samples only. In
ODS alloys, yttria stringers are aligned in the direction of extrusion, resulting in
elongated grains and the number of yttria stringers cut by the bondline is governed by the
orientation of the substrate. To study the effect of substrate orientation on microstructural
bond development at the bondline, bonding was performed in both the T-T orientation
and L-L orientations. Satisfactory bonds were obtained for both orientations, under low
pressures 1 - 5 MPa, and for bonding times up to 6-8 minutes. The microstructures of the
bonds are shown in Figure 44 and Figure 45. However, a few unbonded regions were
observed at the bondline in the transverse orientation.
Bonding was followed by PBHT at 1385 ?C for 2 hours to induce recrystallization
across the bondline. Recrystallization occurred in the substrate, but no grain growth
across the bondline was found. An increase in porosity was also observed after PBHT
[128]. This can be attributed to the release of the hydrogen gas that was entrapped in the
bulk material during alloying, which resulted in pore formation in the bonds during
105
PBHT. A detailed survey of as-received PM2000 bulk material is presented elsewhere
128]. Substrate preparation techniques such as electropolishing and post bond moving
zone recrystallization heat treatment might lead to minimization of unbonded regions and
to attain recrystallization across the bondline [119].
5.2.2 Structure-property relationships of ODS alloys
Room temperature shear tests were performed to evaluate the mechanical
properties of the bonds and to derive the structure-property relationships of the bonds.
Shear tests were initially conducted on as-received bulk material in both the as-received
unrecrystallized fine grain samples and the recrystallized coarse grain samples and these
served as a baseline for the comparison with the bonded materials. Tests were conducted
on PM2000 diffusion bonds of unrecrystallized fine grain substrates bonded to similar
material followed by PBHT, in both the T-T and L-L orientations. However, it should be
noted that the final bonded area did not exceed 90% of the total substrate faying surface
area. The peak shear stresses of the bonds in the transverse and longitudinal orientations,
along with the peak shear stress of as-received bulk material for comparison, are shown
in Figure 46.
The peak shear stress of the recrystallized bulk material is lower than that of the
as-received unrecrystallized bulk material, which is expected due to the reduction in grain
boundaries and the resulting decrease in grain boundary strengthening in the
recrystallized material. Some of the variation in the peak shear stress with orientation of
the substrate was also observed. The transverse orientation showed higher peak shear
stresses than for the longitudinal orientation in the bulk material. Shear strengths of the
106
post bond heat treated bonds in both the T-T and L-L orientations were on the order of
70% of those in the bulk recrystallized material.
Shear tested specimens were examined using SEM, and fracture was observed to
occur at the bondline. Shear testing on coarse grained MA956 revealed that cleavage
cracking was associated with a planar shear fracture, as shown in Figure 47 and Figure
48. In contrast, shear testing on PM2000 bulk material resulted in a planar shear with no
secondary cracking, which can be attributed to the increase in grain boundaries and thus
the grain boundary strengthening that is found in a fine grained material, as shown in
Figure 49. Fractography on as bonded plus post bond heat treated PM2000 bonds
revealed some secondary cracking, which can be attributed to the recrystallization in the
substrate due to PBHT. This can be seen in Figure 50.
107
Figure 40 : LM micrograph, showing diffusion bond of MA956 fine grain - fine
grain (longitudinal orientation) at 1250 ?C for 121 s, followed by PBHT [ 1 h at 1300
?C]. Notice the invisible bondline free of unbonded regions at the bondline
[127,128].
Bondline
100 ?m
108
Figure 41: LM micrograph showing diffusion bond of MA956 fine grain - fine grain
(longitudinal direction) at 1250 ?C for 121 s, followed by PBHT [1 h at 1300 ?C].
Note the grain growth across the bondline and also the voids present at the bondline
[127,128].
20 ?m
Porosity at the bondline
Recrystallization across the joint
109
Figure 42: LM micrograph showing diffusion bond of MA956 coarse grain - fine
grain (longitudinal direction) at 1250 ?C for 170 s. Note the bondline with no voids
present at the bondline [127,128].
Coarse grain
Fine grain
Bondline
110
Figure 43: LM micrograph showing diffusion bond of MA956 coarse grain - coarse
grain (longitudinal direction) at 1250 ?C for 174 s. Note the bondline with large
unbonded regions present at the bondline127,128].
Coarse grain
Coarse grain
Bondline
Large unbonded regions
111
Figure 44: LM micrograph showing diffusion bond of PM2000 in L-L orientation at
1250 ?C for 310 s, in as-bonded (unetched) condition. Note the bondline free of
unbonded regions [127,128].
Longitudinal
Bondline
Longitudinal
112
Figure 45: LM micrograph showing diffusion bond of PM2000 in Transverse
orientation at 1250 ?C for 300 s, in as-bonded (unetched) condition. Note the
bondline free of unbonded regions occasionally seen at the bondline 127,128].
Transverse
Bondline
Unbonded regions
Transverse
113
0
200
400
600
800
1000
Bulk (As -received Fine
Grain)
Recrystallized Bulk Asbonded + PBHT
Condition of Substrates
She
a
r
Str
e
s
s
(M
P
a
)
Transverse Orientation
Longitudinal Orientation
Figure 46: Average shear stress versus condition of substrates. PBHT-post bond
heat treatment. Note that the shear strength of the bulk material decreased after
recrystallization treatments and that of as bonded followed by PBHT is of the order
of 70% of the strength of the bulk material. (Error bars denote 1.96 times standard
deviation to represent a 95% confidence limit)
n=3 n=8 n=3 n=3 n=3 n=4
114
Figure 47: SEM micrograph in SEI mode, of MA 956 coarse grained bulk material
room-temperature shear fracture surface. Notice the main shear fracture surface is
planar, but there is extensive secondary cracking labeled.
500 ?m
Secondary Cracking
115
Figure 48: SEM micrograph in SEI mode, of coarse grained MA 956 bulk material,
showing detail of fracture surface of room-temperature shear test. Notice the
cleavage cracking and shear bands on fracture surface labeled.
20 ?m
Shear Bands
116
(a)
(b)
Figure 49 : SEM micrograph in SEI mode, of fine grained bulk PM 2000, showing
fracture surface of (a) longitudinal sample and (b) transverse sample. Notice that
the main fracture surface is slightly less planar, but there is no secondary cracking.
500 ?m
500 ?m
117
(a)
(b)
Figure 50: SEM micrograph in SEI mode, showing fracture surface of PM2000
transverse-transverse bonds at 1250 ?C for 309 s, followed by PBHT 2 h, 1385 ?C.
Notice the secondary cracking associated with planar shear.
500 ?m
500 ?m
Secondary Cracking
Planar Shear
Secondary Cracking
118
6 CONCLUSIONS
? Competition between wetting of the faying surfaces and formation of the eutectic
along the grain boundaries of the composite interlayer was observed. This led to
non-bonded regions at the faying surfaces, unless a boron-rich interlayer was
employed.
? Adequate wetting occurred in the case of conventional BNi-3 foil interlayers, but
composite interlayers exhibited premature isothermal solidification and
sometimes led to porosity at the faying surfaces due to inadequate amount of
melting point depressant (MPD) present.
? The single crystal CMSX-4 substrate showed a greater extent of wetting
compared to the recrystallized CMSX-4, due to the absence of grain boundaries.
The results from the wettability studies on recrystallized CMSX-4 were
comparable to those for IN 738 and IN 939.
? Suppression of bondline boride formation was achieved using both the Niflex-115
and Niflex-110 interlayers. However, wettability studies confirmed the formation
of borides in the early stages of bond development using the Niflex-110
interlayer, which were not seen with Niflex-115 interlayer.
? In all the bonds examined, there was relatively little gamma-prime at the bondline
when compared with the bulk substrates. For a given wide-gap interlayer, more
119
gamma-prime was formed at the bondline of CMSX-4 ? IN939 compared to
CMSX-4 ? IN738.
? A variety of second phases, which appeared to consist primarily of
borides/carbides and/or topologically close-packed (TCP) phases, were observed
to form in the substrate diffusion zone on the polycrystalline side of both the
CMSX-4 ? IN738 and CMSX-4 ? IN939 bonds, as shown in the early stages of
bond development. However, in the case of bonds involving IN738, the formation
of these second phases could be suppressed after the completion of solid-state
homogenization, by using either of the two composite interlayers, but could not be
avoided in the bonds using BNi-3 foil as an interlayer. In contrast, when bonding
IN939, the second phases remained stable for all times examined for all
interlayers, including post-bond thermal exposures of up to 1 week at 1,000
o
C.
? Two factors dominated the room temperature mechanical properties of the wide-
gap bonds. The first was the extent of gamma-prime formation at the bondline.
Results from shear testing and fractography of the bonds indicated ductile shear
failure at the bondline, which might have been caused due to the formation of
insufficient gamma-prime within the joint, leaving a relatively soft bondline
region. The second factor was the presence of second phases in the diffusion zone
of the polycrystalline substrate, which led to the formation of brittle secondary
cracks. Overall, it is evident that the room temperature shear strength of the bonds
was more dependent on the extent of formation of ?? on the bondline than on the
secondary phases in the diffusion zone of the polycrystalline substrate.
120
? In view of the microstructure-wettability-mechanical property relationships,
Niflex-115 interlayer might be best suited for joining of single crystal superalloy
CMSX-4 with the polycrystalline superalloys IN 738 and IN 939, followed by
post bond heat treatments.
? In oxide dispersion strengthened alloys, diffusion bonding produced successful
results for MA956 and PM2000 material in the unrecrystallized fine grain
condition under low stresses in the range of 1 ? 5 MPa. Shear strengths of the
bonds attained 70% of those of the bulk material. Diffusion bonding of ODS
alloys with substrate orientation aligned in the direction of extrusion might be a
promising technique to achieve bond strengths comparable to that of the bulk
material.
? In view of the structure-property relationships revealed in this study, longitudinal-
longitudinal orientation in a fine grain condition followed by PBHT is likely to
produce the desired results in joining of ODS alloys. However, it should be noted
that the batch to batch inconsistency still has a great impact on the microstructure
and the mechanical properties of bonds, and this needs to be improved for the
process to become commercially viable.
121
7 Future Work
? TEM investigations: The extent of ?? formation was found to have a greater
influence on bond mechanical properties than the formation of brittle second
phases in the diffusion zone of the polycrystalline superalloy substrates. Hence,
extensive TEM studies need to be conducted including an examination of the
microstructural bond development of TLP bonded nickel-based superalloys using
composite interlayers.
? Oxidation and hot corrosion studies of TLP bonded nickel-based superalloys: In
use, superalloys will be subjected to aggressive environments including oxidation
and hot corrosion. Hence, it is vital to study the effects of the TLP bonding
process on the oxidation and hot corrosion properties of TLP bonds using
composite interlayers.
? High temperature mechanical testing of nickel-based superalloys: Since
superalloys will inevitably be subjected to high temperature environments in their
real world applications, it is vital that high temperature mechanical testing,
including high temperature tensile testing and creep testing of as bonded and
PBHT TLP bonded superalloys, be conducted.
? Porosity in ODS alloys: Inconsistencies in the batch to batch properties of
mechanically alloyed ODS alloys, for example porosity, had a significant impact
122
on the microstructure-mechanical property relationships of the bonds. Hence,
comprehensive investigation of the porosity variations, along with a study of ways
to improve the consistency is urgently needed. Hipping the as-cast components,
which successfully eliminated non-surface connected porosity considerably in the
precipitation hardened superalloy castings [132, 133], seems to be promising and
needs further study.
? Dislocation interactions of ODS alloys: Iron based ODS alloys were intended for
use in temperatures higher than the solution temperature of ??. However, the
interaction of dislocations with the dispersoids at high temperatures is still a
subject for debate. At low temperatures (below 400
0
C), mutual dislocation
interactions and Orowan looping [134] is the major mechanism. In this
mechanism, as the dislocation advances, dislocation loops are left around the
dispersoids. On the other hand, at 400-600
0
C, dislocations were seen attaching at
the departing side of the dispersoids and as the dislocation advances, no
dislocation debris was left behind [135, 136]. At this temperature, it is believed
that transition of mechanism from Orowan looping to bypassing is occurring and
the normalized yield stress drops at that temperature, which is much less than the
yield stresses below that temperature [137-139]. Thus, extensive TEM studies
associated with high temperature mechanical testing of bulk and TLP bonded
ODS alloys with agglomerated dispersoids would yield valuable information.
? Finite element analysis of shear testing: Shear testing is an easy and effective
method of eliminating poor bonds. However, the application of shear forces on
different types of alloys [nickel-based superalloys, ODS alloys, titanium alloys]
123
revealed a difference in their behavior, and consequently an extensive
computational analysis of shear testing using the finite element method is
recommended.
? High temperature testing of TLP bonded ODS alloys: ODS alloys were designed
to withstand temperatures higher than those at which nickel-based superalloys can
operate, and since TLP bonding results in a reasonable dispersoid continuity and
elongated grain structure, high temperature testing, including creep, cyclic
oxidation tests, and high temperature tensile testing of TLP bond should be
performed and the results compared with those for substrates joined using other
techniques.
124
8 BIBLIOGRAPHY
1. E.F. Bradley: ?SOURCE BOOK ON MATERIALS FOR ELEVATED-
TEMPERATURE APPLICATIONS?, American Society for Metals, Metals Park,
Ohio, 1969, pp. 275-298
2. G.P.Sabol, R.Stickler: ?MICROSTRUCTURE OF NICKEL-BASED
SUPERALLOYS: REVIEW ARTICLE? Phys. Stat. Sol. 35(11),1969,
Westinghouse Research Laboratories, Pittsburgh, Pennsylvania.
3. C.T.Sims: J. Met., 1966, October, pp.1119-1130.
4. A.K.Jena, M.C. Chaturvedi: J. Matl. Sci, 1984, (19), pp.3121-3139.
5. E.F. Bradley: ?SUPERALLOYS: A TECHNICAL GUIDE?, Metals Park,OH;
ASM International, 1988.
6. P.S. Kotval: Metallography, 1, 1969, pp.1119-1130.
7. C.T. Sims: ?SUPERALLOYS: GENESIS AND CHARACTER? in ?Superalloys
II?, edited by C.T. Sims, N.S. Stoloff, W.C. Hagel, 1987, John Wiley & Sons,
pp. 3-21.
8. M.Ott, H. Mughrabi: Mater. Sci. & Eng. A, 1999, 272, pp.24-30.
9. Michael F. Ashby and David R.H. Jones: ENGINEERING MATERIALS I: AN
INTRODUCTION TO THEIR PROPERTIES AND APPLICATIONS, Oxford;
Butterworth-Heinemann, 1996.
10. J.W. Lee, J.H. Mc Murray and J.A. Miller: Weld. J., 64(10), 1985, pp.18-21.
11. G.M. Slaughter: Weld. J., 58(10), 1979, pp.17-28
12. W.A. Owczarski: ?PHYSICAL METALLURGY OF METAL JOINING?, edited
by R. Kossowsky and M.E. Glicksman, Metallurgical Society of AIME,
Warrendale, PA, 1980, pp.166-190
13. C.Y. Su, W.C. Lih, C.P. Chou and H.C. Tsai: J. Mater. Proc. Technol., 115, 2001,
pp.326-33
125
14. K. Schneider, B.Jahnke and R. Burgel: Mater. Sci. Technol, 1(8), 1985, pp.613-
619
15. R. Creal: Heat Treatment, 19(10), 1987, pp.30 ? 32
16. T.J. Kelley: Weld. J., 62(10), 1983, pp.57-58
17. J.R. Kelley: Weld. Des. Fabricat., 57(7), 1985, pp.43 ? 45.
18. D. Mckeon: Weld. J., 3, 1971, pp.201s-206s
19. K.C. Wu, R.E. Herfert: Weld. J., 46, 1967, pp.32s-38s
20. W.P. Hughes, T.F. Berry: Weld. J., 46, 1967, pp.361s-367s
21. J.R.Thamburaj, W.Wallace, and J.A.Goldak: Int. Metall. Rev., 28(1), 1983, pp.1-
22.
22. A.E.Dray: Ph.D.Thesis. 1985, University of Cambridge, Cambridge, UK.
23. X. Wu, R. S. Chandel, and H. Li: J.Matl. Sci., 36(6), 2001, pp. 1539-1546.
24. S.J. Lee, S.K. Wu, and R.Y. Lin: Acta. Mater., 46, 1998, pp.1283-1286.
25. I. Tuah-poku, M. Dollar, and T.B. Massalski: Metall. Trans. A, 19A(3), 1988, pp.
675-686.
26. D.S. Duvall, W.A. Owczarski, D.F. Paulonis: Weld. J., 53(4), 1974, pp. 203-214
27. W.F. Gale, D.A. Butts: Sci. Technol. Weld. Join., 2004, 9(4), pp. 283-300.
28. Y. Zhou, W.F. Gale, and T.H. North: Int. Mater. Rev., 40(5), 1995, pp.181-196
29. R. Sharghi-Moshatghin, S. Asgari: J. Alloy. Compd. 368, 2004, pp.144-151.
30. R.A. Stevens, P.E.J. Flewitt: Mater. Sci. Eng., 37, 1979, pp.234?247.
31. P.K. Footner, B.P. Richards: J. Mater. Sci., 17, 1982, pp.2141?2153.
32. G.J. Karen: Ph.D Thesis. 1983, Stanford University, USA.
33. J.S. Benjamin: Met. Trans. A., 1(10), 1970, pp.2943-2951.
126
34. L. Habraken: ?MATERIALS AND GAS TURBINES?, in ?High Temperature
Alloys for Gas Turbines?, edited by D. Coutsouradia, P. Felix, H. Fichmeister, L.
Habraken, and M.O. Speidel, 1978, Applied Science Publishers Ltd., pp.xiii-xiv.
35. W. Boesch: ? INTRODUCTION-SUPERALLOYS?, in ?Superalloys, Super
composites and Superceramics?, edited by J.K. Tien and T. Caulfield, 1989,
Academic Press Inc. Ltd., pp.1-7
36. S. Ahmed: JOM, 42(8), 1990, pp. 24-26
37. M.H. Van de Voorde: ?ADVANCED HIGH TEMPERATURE MATERIALS? in
?Research and Development of High Temperature Materials for Industry?, edited
by E. Bullock, R. Brunetaud, J.F. Conde, S.R. Keown, and S.F. Pugh, 1989,
Elsevier Applied Science, pp.1-12
38. J.H. Wood and E.H. Goldman: ?PROTECTIVE COATINGS? in ?Superalloys II?,
edited by C.T.Sims, N.H. Stoloff, and W.C. Hagel, 1987, John Wiley & Sons,
pp.359-384
39. O.D. Sherby and P.M. Burke: ?MECHANICAL BEHAVIOR OF
CRYSTALLINE SOLIDS AT ELEVATED TEMPERATURE? in ?Progress in
Materials Science?, edited by B. Chalmers and W. Hume-Rothery, 1968, Vol.13,
Pergamon Press, pp.325-390
40. C.T. Sims: ?HIGH TEMPERATURE ALLOYS IN HIGH-TECHNOLOGY
SYSTEMS? in ?High Temperature Alloys for Gas Turbines?, edited by
D.Coutsouradia, P. Felix, H. Fichmeister, L. Habraken, and M.O. Speidel, 1978,
Applied Science Publishers Ltd., pp.13-79.
41. J.R. Stephens: ?RESOURCES-SUPPLY AND AVAILABILITY?, in
?Superalloys, Super Composites and Superceramics?, edited by J.K. Tien and T.
Caulfield, 1989, Academic Press Inc. ltd., pp.9-48.
42. J.R. Davis: ?NICKEL, COBALT AND THEIR ALLOYS?, in ?ASM Specialty
Handbook?, ASM International. 2000.
43. C.R.Brooks: ?HEAT TREATMENT, STRUCTURE AND PROPERTIES OF
NON-FERROUS ALLOYS?. 1982.
44. M. J. Donachie and S. J. Donachie: in ?SUPERALLOYS A TECHNICAL
GUIDE?, 2nd edition, ASM, 2002. pp.26.
45. E.W. Ross, and C.T. Sims: ?NICKEL-BASE ALLOYS? in ?Superalloys II?,
edited by C.T. Sims, N.S. Stoloff, W.C. Hagel, 1987, John Wiley & Sons, pp.101.
127
46. N.S. Stoloff: ?FUNDAMENTALS OF STRENGTHENING? in ?Superalloys II?,
edited by C.T. Sims, N.S. Stoloff, W.C. Hagel, 1987, John Wiley & Sons, pp.61-
96.
47. T.L. Lin: Matl. Sci. Eng. A, 128, 1990, pp. 23-31.
48. M.Yamaguchi: Philos. Mag. A, 45, 1982, pp. 867-871.
49. V. Paidar: Philos. Mag. A, 45, 1982, pp.883-895.
50. C.E. Campbell and W.J. Boettinger: Metall. Mater. Trans. A, 31A, 2000, pp.2835-
2847.
51. P.R. Mobley, G.S. Hoppin III: Weld. J., 40, 1961, p.610s.
52. B.L. Gruzdev: Weld. J., 53, 1974, p.543.
53. W.A. Owczarski, D.S. Duvall, C.P. Sullivan: Weld. J., 45, 1966, 145-s.
54. M. Prager, C.S. Shira: Welding Research Council Bulletin, 128 (1968).
55. M.H. Haafkens, G.H. Matthey: Weld. J., 61 (1982) pp.25-32.
56. T.J. Kelly, in: R.A. Patterson, K.W. Mahin (Eds.), Proceedings of the Conference
on Weldability of Materials, ASM International, USA, 1990, pp.151?157.
57. W.B. Busch, H.D. Kunze, Proceedings of the Conference on Materials
Development in Turbo-Machinery Design, The Institute of Metals, UK, 1988, pp.
81?87.
58. G. Cam, M. Kocak: Int. Mater. Rev. 43 (1998) 1.
59. M.B. Henderson, D. Arrell, M. Larsson, G. Marchant: Sci. Technol. Weld. Join.,
9, 2004, pp.13-18.
60. O. Hunziker, D. Dye, S.M. Roberts and R.C. Reed: Proc. Conf. on Numerical
Analysis of Weldability, Graz-Seggau, Austria, 1999.
61. M.B. Henderson, D. Arrellb, M. Heobel, R. Larsson and G. Marchant: Conf. in
the Microstructure & Performance of Joints in High-Temperature Alloys, London,
UK,Nov. 2002.
62. A. Dhooge and A. Vinckier: Int. J. Pressure Vessels and Piping, 27, (1987), pp.
239-269.
128
63. G. Cam and M. Kocak: Int. Mater.Rev., 43, No. 1, (1998), pp. 1-44.
64. M. Prager, M. Shira: ?WELDING OF PRECIPITATION-HARDENING
NICKEL-BASE ALLOYS? in ?Welding Research Council?, Bull. No.128, New
York, 1968
65. R.P. Newman, T.R. Gurney: Brit. Weld. J., 11, 341, 1964.
66. R.R. Gatts, G.J. Sokol: ?EFFECTS OF CLOSURE WELD CONFIGURATION
ON THE FATIGUE STRENGTH OF OMEGA SEALS? , KAPL-2000-21,
Reactor Metallurgy Report No.24, Metallurgy, 1.1-1.11, 1964
67. W. Choi, S. Kim, C. Lee, J. Jang: ?EFFECT OF HEATING RATE ON
TRANSIENT LIQUID PHASE DIFFUSION BONDING OF NI-BASED
SUPERALLOY GTD-111? in ?Designing, Processing and Properties of
Advanced Engineering Materials: Proceedings on the 3rd International
Symposium on Designing, Processing and Properties of Advanced Engineering
Materials?, Nov 5-8 2003. Jeju Island, South Korea, Trans Tech Publications Ltd,,
Zurich-Ueticon, Switzerland.
68. Y.Zhou, W.F. Gale, T.H. North: Int. Mater. Rev., 40(5), 1995, pp.181-196.
69. Y. Nakao and K. Shinozaki: Matl. Sci. and Technol., 11, 1995, p.304.
70. Y. S. Kwon, J. S. Kim, J. S. Moon, and M. J. Suk: J. Mater. Sci., 35, 2000, pp.
1917-1924
71. W. F. Gale, X. Wen, T. Zhou, and Y. Shen: Mater. Sci. Tech., 17, 2001, pp. 1423-
1433
72. Y. S. Kwon, J. S. Kim, J. S. Moon, and M. J. Suk: J. Mater. Sci., 35, 2000, pp.
1917-1924.
73. E.Lugscheider, T. Schittny and E. Halmony, Weld. J.(Res. Suppl.), 68(1), 1989,
9s-13s.
74. W.D. Zhuang and T.W. Eagar: Weld. J.(Res. Suppl.), 76, 1997, pp. 157s-162s.
75. W.F. Gale, Y. Xu, X. Wen, and Z.A.M. Abdo: Metall. Mater. Trans. A, 30A,
1999, pp.2723-2726.
76. Y.Guan and W.F. Gale: Mater. Sci. Technol., 15, 1999, 207-212
77. W.F.Gale, D.A. Butts, M. Di Ruscio and T.Zhou: Metall. Mater. Trans. A,
33A(10), 2002, 3205-3214.
129
78. W.F. Gale and Y.Guan: Mater. Sci. Technol., 15, 1999, 464-467.
79. M. K. Kitchings, Y. Guan, and W. F. Gale: "MICROSTRUCTURAL
DEVELOPMENT DURING TRANSIENT LIQUID PHASE BONDING OF
STRUCTURAL INTERMETALLIC COMPOUNDS," in ?Analysis of in-Service
Failures and Advances in Microstructrual Characterization?,edited by E.
Abramovici, D. O. Northwood, M. T. Shehata, and J. Wylie, ASM International,
Materials Park, OH, 1999, pp. 425-430.
80. J.S.Benjamin: Met. Trans. A, 1970, 1(10), pp.2943-2951.
81. C. Suryanarayana, E. Ivanov, and V.V. Boldyrev: Matl. Sci. and Eng. A, 304-306,
2001, pp. 151-58.
82. M.Ruhle: in Proc. 1st Intl. Conf. Heat Resistant Matls. Sep 23-26, 1991. Fontana,
Wisconsin: ASM.
83. G.H.Gessinger: ?POWDER METALLURGY OF SUPERALLOYS?, 1984,
Elsevier, pp.213.
84. A.Czyrska-Filemonowicz, B.Dubiel: J. Matl. Proc. Technol, 64, 1997, pp.53-64.
85. C. White: ?ADVANCED HIGH TEMPERATURE MATERIALS? in ?Research
and Development of High Temperature Materials for Industry?, edited by E.
Bullock, R. Brunetaud, J.F. Conde, S.R. Keown, and S.F. Pugh, 1989, Elsevier
Applied Science, pp.71-78.
86. P. Krautwasser, A. Czyrska-Filemonowicz, M. Widera, F. Carsughi: Matl. Sci.
and Eng. A, 177, 1994, pp.198-208.
87. H.R.Seth: Ph.D Thesis, 1991, Northwestern University, USA.
88. R.C.Benn: Frontiers of High Temperature Materials, Inco Alloy International.
1983.
89. A.Czyrska-Filemonowicz: J. Matl. Proc. Technol., 53, 1995, pp.93-100.
90. T.J. Kelly: ?WELDING OF MECHANICALLY ALLOYED ODS
MATERIALS?, in ?Proceedings of Trends in Welding Research?, edited by S.A.
David, ASM, Metals Park, OH, 1982, pp. 471-478.
91. L.E. Shoemaker: ?JOINING TECHNIQUES FOR A FERRITIC OXIDE
DISPERSION STRENGHTENED ALLOY?, in ?Proceedings of the International
130
Conference on Trends in Welding Research?, edited by S. A. David, ASM,
Metals Park, OH, 1986, pp. 371-77.
92. P.A. Molian, Y.M. Yang, P.C. Patnaik: J. Mater. Sci., 27, 1992, pp.2687-2694.
93. C.Y. Kang, T.H. North, D.D. Perovic: Metall. Mater. Trans. A, 27A, 1996, pp.
4019-4029.
94. T.I. Khan, E.R. Wallach: J. Matl. Sci, 30, 1995, pp.5151-5160.
95. T.I. Khan, E.R. Wallach, J. Matl. Sci, 31, 1996, pp.2937-2943.
96. H. Kinoshita, N. Akasaka, H. Takahashi, I. Shibahara, and S. Onose:
?Microstructural Change on Electron Irradiated Oxide Dispersion Strengthened
Ferritic Steels?, J. Nucl. Mater., 1992, Vol. 191-194(2), pp. 874-78.
97. D.A. Butts, T. Zhou, W.F. Gale: Internal Communication on Joining of
Superalloys, Aug.2002.
98. E.H. Bradley: Source Book on Materials for Elevated-Temperature
Applications.Wiley, 1979.
99. S. V. Orel, L. C. Parous, and W. F. Gale: Weld. J. Suppl., 74, 1995, pp. 319s-
324s.
100. D. S. Duvall, W. A. Owczarski, D. F. Paulonis, and W. H. King: Weld. J. Suppl.,
1972, 53, pp. 41s-49s.
101. A. Rabinkin: Weld. J., 68(10), 1989, pp. 39-45.
102. T. J. Moore, and T. K. Glasgow: Weld. J. Suppl., 64(8), 1985, pp. 219s-226s.
103. W. Li, T. Jin, X. F. Sun, Y. Guo, H. R. Guan, and Z. Q. Hu: Scripta Mater., 48,
2003, pp. 1283-1288.
104. T. J. Kelly: Weld. J. Suppl., 61(10), 1982, pp. 317s-319s.
105. A. Ekrami and T. I. Khan: Mater. Sci. Tech., 15, 1999, pp. 946-950.
106. Y. Zhai and T. H. North: J. Mater. Sci., 32, 1997, pp. 5571-5575.
107. C. A. Blue, V. K. Sikka, R. A. Blue, and R. Y. Lin: Metall. Mater. Trans. A, 27A,
1996, pp. 4011-4017.
108. Z. Li, Y. Zhou, and T. H. North: J. Mater. Sci., 30, 1995, pp. 1075-1082.
131
109. W. F. Gale and S. V. Orel: J. Mater. Sci., 31, 1996, pp. 345-349.
110. W. F. Gale, D. A. Butts, M. D. Ruscio, and T. Zhou: Metall. Mater. Trans. A,
33A, 2002, pp. 3205-3214.
111. S. V. Orel, L. C. Parous, and W. F. Gale: Weld. J. Suppl., 74, 1995, pp. 319s-
324s.
112. Y. Guan, and W. F. Gale: Mater. Sci. Tech., 15, 1999, pp. 207-212.
113. W. F. Gale and Y. Guan: Metall. Mater. Trans. A, 27A, 1996, pp. 3621-3629.
114. P.Yan and E. R. Wallach: Intermetal., 1(2), 1993, pp. 83-97.
115. A. Hill and E. R. Wallach: Acta Metall.? 37(9), 1989, pp. 2425-37.
116. M. G. Nicholas: ?JOINING PROCESSES?, Kluwer Academic Publishers, The
Netherlands, 1988.
117. T. W. Eager: ?Processes of Joining Materials?.
http://web.mit.edu/3.37/www/chapter.shtml (06/09/2005).
118. T.J. Moore, and T.K. Glasgow: Weld. J., 64, 1985, pp. 219-226.
119. I.A. Bucklow: ?Diffusion Bonding a Creep Resistant Fe-ODS Alloy?, conf. proc.,
Advances in Joining Newer Structural Materials, Proceedings of the International
Conference, Montreal, Canada: Elmsford, NY: Pergamon. (1990), pp. 299-304.
120. G. Zhang, R. S. Chandel, H. P. Seow and H. H. Hng: Mater. Manuf. Processes,
18, 2003, pp. 599-608.
121. S.V. Chitti: Masters thesis, Auburn University, 2005
122. J.-M. Veys and R. Mevrel: Mater. Sci. Eng., 88, 1987, pp.253-260.
123. H.W. Grunling, K.Schneider and L. Singhesier: Mater. Sci. Eng., 88, 1987,
pp.177-189.
124. T.C. Totemeier, W.F. Gale and J.E. King: Matl. Sci. Eng., A169, 1993, pp.19-26.
125. T. Zhou, Ph.D. thesis, Auburn Uiniversity, 2002.
126. Section 14.95 of the ANSYS Theory Reference
132
127. V.G. Krishnardula, R.Aluru, N.I. Sofyan and W.F. Gale: Proceedings of 7
th
International Conference on Trends in Welding Research VI, Pine Mountain, GA,
in Press, 2005.
128. V.G. Krishnardula: Ph.D. dissertation, Auburn University, 2005.
129. H.K.D.H. Bhadeshia: Mater. Sci. Eng., A223, 1997, pp.64-77
130. B.S. Bokstein, V.S. Gostomelskii, V.A. Ivanov, A.L. Petelin, and S.A. Petelin:
Mater. Lett., 39(1), 1999, pp. 77-79.
131. N. Orhan, M. Aksoy, and M. Eroglu: Mater. Sci. Eng., A271, 1999, pp. 458-68.
132. W.R. Freeman, Jr.:?INVESTMENT CASTINGS? in ?Superalloys II?, edited by
C.T.Sims, N.H. Stoloff, and W.C. Hagel, 1987, John Wiley & Sons, pp.411.
133. S. Reichman and D.S. Chang: ?POWDER METALLURGY? in ?Superalloys II?,
edited by C.T.Sims, N.H. Stoloff, and W.C. Hagel, 1987, John Wiley & Sons.
134. E.Orowan: Symp. on Internal Stresses in metals and Alloys., 1948. London
135. P.Guyot: Acta. Metall., 12, 1964, pp. 941-949.
136. R.Lagneborg: Scripta. Metall., 7, 1973, pp. 605-610.
137. J.Rosler and E.Arzt: Acta. Metall., 36, 1988, pp.1043-1049.
138. D.J.Srolovitz, M.J.Luton, and R.A.Petkovic-Luton: Acta Metall., 31, 1988,
pp.2151-2158
139. D.J.Srolovitz, et al.: Acta Metall., 32, 1984, pp.1079-1083.