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