Solidification Cracking Performance and Metallurgical Analysis of Filler

SOLIDIFICATION CRACKING PERFORMANCE AND

METALLURGICAL ANALYSIS OF FILLER METAL 82

THESIS

Presented in Partial Fulfillment of the Requirements for the Degree Master of Science in the Graduate School of The Ohio State University

Michael Romanoff Orr

Graduate Program in Welding Engineering

The Ohio State University

2016

Master's Examination Committee:

Dr. John Lippold, Graduate Advisor / Faculty Emeritus

Dr. Antonio Ramirez, Graduate School Chair / Professor

Michael Romanoff Orr

2016

ABSTRACT

The main objective of this research is to investigate and eventually optimize the solidification cracking resistance of filler metal ERNiCr-3 (FM 82) in thick-section, highly-restrained, hot wire automatic gas tungsten-arc welding (AGTAW-HW) weld deposits used on reactor vessels in the nuclear energy industry. This investigation was achieved with experimental verification of cracking responses of three FM 82 heats using the Cast Pin Tear Test (CPTT), metallurgical/chemical analysis of each heat, and computational analysis. An intentional nitrogen addition study was conducted to observe the effect of increased weld metal N on solidification cracking susceptibility of FM 82.

Experimentally, results were considered using the CPTT with varying heats of

ERNiCr-3 exhibiting either superior or poor solidification cracking resistance behavior.

A verification study was implemented using the CPTT and its designated procedure with three heats of FM 82 that were previously tested. Excellent correlation results showed accuracy and repeatability of the CPTT process. Quantitatively, zero-pin deviations were found at the lower cracking threshold (LCT) of each heat. The LCT is designated as the highest pin length to exhibit 0% and is the main ranking criterion of solidification cracking susceptibility. Qualitatively, the results proved to verify the same ranking of solidification cracking susceptibility as the previous study—this ranking also mimics what is observed in production mockups. Scanning electron microscope (SEM) imaging of fracture surfaces exhibited classic solidification fracture morphology to ensure failure

mode. Additional work was completed using the CPTT for two dilution studies (10wt% and 25wt %) of a “resistant” heat by a “susceptible” heat. It was observed that 10% dilution of Heat B by Heat A reduced the LCT by one pin-length, whereas 25% dilution reduced it by three pin-lengths.

Furthermore, computational thermodynamic software, Thermo-CalcTM, was utilized to determine solidification temperature ranges and terminal phase formations of the FM 82 heat chemistries. This technique constituted a Scheil simulation which estimate alloying segregation behavior during solidification and phase stability after welding. For ease of comparison, one resistant heat (Heat B) was compared against one susceptible heat (Heat A). Heat B was predicted to have a solidification temperature range (STR) of 216°C and 0.1985 volume fraction NbC upon solidification to 99% solid.

Heat A was predicted to have a lower STR (203°C) and higher volume fraction NbC

(0.2822) upon solidification to 99% solid.

Lastly, intentional nitrogen additions were included to see the effects the interstitial has on the susceptibility to solidification cracking of a FM 82 heat.

Weldability testing included CPTT of a resistant heat welded with additions of 0.2, 0.4, 1, and 5 vol% N2 in the shielding gas mixture. SEM/EDS techniques were utilized to determine which constituents form based on the nitrogen content of the gas mixture. MX- type precipitates rich in Nb and Ti were found both in grain boundaries and in regions throughout the matrix. Although TEM work was not done to confirm this, the literature determines these MX precipitates to be (Nb,Ti)(C,N).

iii

DEDICATION

This work is dedicated to God, my loving family, and my close friends. A special thank you to my parents, Jonathan and Karin Orr, whose encouragement to think and act independently drive me to live and love life every day. I am thankful for all of their wisdom, generosity, humility, and emotional and financial support throughout my life. Also, an immense feeling of gratitude to my brothers, Benjamin and Matthew Orr, for being my best friends and biggest role models.

Michael Romanoff Orr The Ohio State University, 2016

ACKNOWLEDGEMENTS

I would like to thank:

Dr. John Lippold—my graduate advisor, for his constant wisdom and support for my research

Frank Argentine—my contact at BWXT, whose passion for his work, and constant support and communication made for a close and enjoyable collaboration between research and industry

Ed Pfeifer and Ken Copley—the lab supervisor and technician at OSU, for their timely maintenance and much-appreciated advice with the machinery and technology at the Edison Joining Technology Center

My fellow Graduate Research Associate—for teaching me in the labs, discussing metallurgy in the lounge, and making graduate school a fun experience altogether

Connor Guarino—my mentee from Metro High School, for his hard work and persistence in the metallurgy lab at OSU preparing samples for testing

VITA

June 2010……………..…………..…………………………….Woodridge High School

2015…………………………….…..B.S Welding Engineering, Ohio State University

2016………………………………...M.S Welding Engineering, Ohio State University

Field of Study

Major Field: Welding Engineering

TABLE OF CONTENTS

ABSTRACT ...... ii

DEDICATION ...... iv

ACKNOWLEDGEMENTS ...... v

VITA ...... vi

LIST OF TABLES ...... xi

LIST OF FIGURES ...... xiii

CHAPTER 1: INTRODUCTION ...... 1

CHAPTER 2: BACKGROUND ...... 4

2.1 Welding Metallurgy ...... 5

2.1.1 Solidification Subgrain Boundaries (SSGBs)...... 7

2.1.2 Solidification Grain Boundaries (SGBs) ...... 7

2.1.3 Solute Redistribution ...... 8

2.2 Solidification ...... 9

2.2.1 Solidification of Fusion Welds ...... 10

2.2.2 Solidification in Castings ...... 12

vii

2.3 Weldability ...... 16

2.3.1 Solidification Cracking ...... 17

2.3.2 Avoiding Solidification Cracking ...... 29

2.4 Weldability Testing ...... 32

2.4.1 Original Cast Pin Tear Test Method ...... 32

2.4.2 OSU Second Generation Cast Pin Tear Test Method...... 35

2.4.3 OSU Third Generation Cast Pin Tear Test Method ...... 37

2.5 Nitrogen Effects on Ni-based alloys ...... 38

2.5.1 Investigation on the use of Nitrogen in Shielding Gas during Welding of Ni-

based Alloys and Materials with Austenitic Microstructures ...... 41

2.5.2 Effects of Nitrogen on Microstructure and Mechanical Properties ...... 43

2.5.3 Internal Nitridation of Ni-based Alloys ...... 47

2.5.4 Nitrogen Effects on Solidification Cracking in Ni-based Alloys ...... 49

CHAPTER 3: OBJECTIVES ...... 53

3.1 Cast Pin Tear Test ...... 53

3.2 Solidification Cracking Performance and Metallurgical Analysis of FM 82...... 54

CHAPTER 4: MATERIAL AND EXPERIMENTAL PROCEDURE ...... 55

4.1 Material ...... 55

4.2 Weldability Testing Procedure ...... 56

viii

4.2.1 Material Preparation ...... 57

4.2.2 Button-Melting ...... 59

4.2.3 Cast Pin Tear Testing ...... 63

4.2.4 Crack Response Analysis ...... 72

4.2.5 Sample Preparation and Metallurgical Characterization ...... 75

4.2.6 Computational Modeling of Solidification Behavior and Eutectic Formation . 76

CHAPTER 5: RESULTS AND DISCUSSION ...... 77

5.1 CPTT Initial Verification of FM 82 Heats ...... 77

5.1.1 Heat A: High Susceptibility to Solidification Cracking ...... 78

5.1.2 Heat Y: Low Susceptibility to Solidification Cracking ...... 82

5.1.3 Heat B: Low Susceptibility to Solidification Cracking ...... 85

5.1.4 CPTT Dilution Studies ...... 88

5.2 Thermodynamic Simulations ...... 95

5.3 Metallurgical Characterization of Sample Wire Heats ...... 102

5.3.1 Metallography ...... 102

5.3.2 Fractography ...... 108

5.3.3 Compositional Analysis ...... 113

5.4 Effects of Nitrogen Additions ...... 114

5.4.1 Heat YT: 140 ppm N ...... 115

5.4.2 Heat YT: 200 ppm N ...... 120

5.4.3 Heat YT: 600 ppm N ...... 125

5.4.4 Heat YT: 1300 ppm ...... 128

CHAPTER 6: CONCLUSIONS ...... 140

6.1 CPTT ...... 140

6.2 FM 82 Cracking Performance ...... 141

CHAPTER 7: RECOMMNEDATIONS FOR FUTURE WORK ...... 143

7.1 CPTT ...... 143

7.2 Metallurgical Evaluation of FM 82 ...... 146

REFERENCES ...... 147

APPENDIX: Nitrogen Study CPTT Results ...... 153

LIST OF TABLES

Table 1: Second Generation CPTT Parameters ...... 36

Table 2: Chemical Compositions (wt%) of Test Materials ...... 44

Table 3: Welding Conditions for GTAW and pulsed GMAW Processes [20] ...... 51

Table 4: Measured Chemical Compositions of Tested FM 82 Heats ...... 56

Table 5: CPTT Pin Length (in) Per Given Button Mass (g) for FM 82...... 58

Table 6: Nominal Parameters for Button-Melting ...... 62

Table 7: CPTT Behavior from Previous Study [1] ...... 78

Table 8: Initial CPTT Parameters ...... 78

Table 9: Cracking response characteristics of Heat A ...... 79

Table 10: Cracking response characteristics of Heat Y ...... 82

Table 11: Cracking response characteristics of Heat B ...... 85

Table 12: Cracking response characteristics of 10% Dilution of Heat B with Heat A ..... 88

Table 13: Cracking response characteristics of 25% Dilution of Heat B with Heat A ..... 90

Table 14: Cracking response criteria for each heat ...... 94

Table 15: Parameters Used in Thermodynamic Simulations using Thermo-Calc 2015a

Software ...... 96

Table 16: Chemical Compositions used for Computational Modeling ...... 96

Table 17: CS, CL and k (CS / CL) at the End of Solidification (99% solid) for Heats B and A ...... 101

Table 18: Thermo-Calc Scheil Simulation Results for Heats B and A...... 101

Table 19: EDS/Vendor Chemical Composition Comparison ...... 114

Table 20: Heat YT Chemical Composition ...... 115

Table 21: Nitrogen Addition Weld Sample Gas Mixtures and Content ...... 115

Table 22: CPTT Results for Nitrogen Addition Study ...... 138

Table 23: Nitrogen Study Summary Table ...... 139

xii

LIST OF FIGURES

Figure 1: CPTT cracking response curves among 7 heats of FM 82 [1] ...... 2

Figure 2: Schematic of grain boundaries seen in austenite-solidifying weld metal ...... 6

Figure 3: Typical FM 82 microstructure. SGBs and SSGBs in austenitic matrix (Heat A,

1000X magnification) ...... 6

Figure 4: Solidification grain morphologies due to welding [7]...... 11

Figure 5: Change in weld pool geometry from (L) elliptical to (R) teardrop based on welding speed [7] ...... 12

Figure 6: Schematic of (a) competitive growth of crystals from mold wall (b) inward [7]

...... 13

Figure 7: (L) Schematic of cast grain structure exhibiting (R) Photomicrograph taken at

1000X magnification of cast pin grain structure in Heat Y. Zones: (1) an outer chill zone of equiaxed crystals, (2) a columnar zone of elongated grains, and (3) a central equiaxed zone. [7] ...... 14

Figure 8: Development of thermal primary, secondary, and tertiary dendrite arms. (L)

Schematic, (R) Fracture surface of Heat Y [7] ...... 15

Figure 9: Shrinkage-Brittleness theory based on a eutectic phase diagram ...... 19

Figure 10: Strain theory according to Pellini [6] ...... 21

Figure 11: Stages of the Generalized theory of solidification cracking [6] ...... 22

xiii

Figure 12: Modified Generalized theory on interaction of solidification behavior, fractography, and cracking susceptibility. [16]...... 24

Figure 13: Observed stages of Modified Generalized Theory of solidification cracking [6]

...... 25

Figure 14: Three characteristic fractographic regions of a solidification crack [17] ...... 26

Figure 15: Graphical illustration of the hypothesis of the Technological Strength theory

[18] ...... 28

Figure 16: Simplified version of hypothesis of the Technological Strength theory [6] ... 28

Figure 17: Effect of weld D/W ratio and bead size and location on solidification cracking

[6] ...... 31

Figure 18: Schematic of original levitation melter setup [19] ...... 34

Figure 19: Pin mold dimensions [19]...... 34

Figure 20: CPTT ranking of stainless steels [18]...... 35

Figure 21: Second generation CPTT at OSU [13] ...... 36

Figure 22: CPTT average cracking response curve labeled with ranking criteria ...... 38

Figure 23: Nitrogen solubility curve versus alloying element content in at under a nitrogen partial pressure of pN2=1bar [25] ...... 41

Figure 24: Grain refinement in weld metal due to increasing N2 additions in Ar gas. (A)

Pure Ar and (B) Ar-4%N2 for Alloy 263, (C) Pure Ar and (D) Ar-4%N2 for Alloy X .... 45

Figure 25: Internal nitridation in Ni-based superalloys due to localized erosion in the form of surface cracks [36] ...... 48

Figure 26: PVR test results of different base metals susceptible to hot cracking using

GTAW with varying gas composition [20] ...... 51

xiv

Figure 27: Wire-cutter setup at OSU ...... 58

Figure 28: Progression of filler metal sample via steps for CPTT (Top) and profile of

"button" used for CPTT (Bottom) ...... 59

Figure 29: Button-melter setup ...... 60

Figure 30 Button-melting schematic ...... 60

Figure 31: CPTT setup schematic ...... 64

Figure 32: CPTT custom water-cooled copper coil ...... 64

Figure 33: LabVIEW control screen ...... 66

Figure 34: (L) Quartz funnel used in CPTT and (R) engineering drawing of large funnel

...... 66

Figure 35: Engineering isometric drawing (L) and photograph of copper relief disc (R)-- green represents area of acceptable flashing, while red represents area unacceptable of flashing ...... 69

Figure 36: Mold retainer assembly from R-to-L: Spacer, spring, copper relief disc, mold half ...... 69

Figure 37: CPTT purge cycle communication schematic ...... 71

Figure 38: Circumferential crack-counting apparatus and setup ...... 74

Figure 39: Crack-counting device which converts rotation of the pin into circumferential cracking percentage based on 360° ...... 74

Figure 40: CPTT complete curve (A) and average circumferential cracking percentage

(B) versus pin length comparing two studies of Heat A ...... 80

Figure 41: CPTT maximum cracking (A) and minimum cracking percentage (B) versus pin length comparing two studies of Heat A ...... 81

Figure 42: CPTT complete curve (A) and average circumferential cracking percentage

(B) versus pin length comparing two studies of Heat Y ...... 83

Figure 43: CPTT maximum cracking (A) and minimum cracking percentage (B) versus pin length comparing two studies of Heat Y ...... 84

Figure 44: CPTT complete curve (A) and average circumferential cracking percentage

(B) versus pin length comparing two studies of Heat B ...... 86

Figure 45: CPTT maximum cracking (A) and minimum cracking percentage (B) versus pin length comparing two studies of Heat B ...... 87

Figure 46: CPTT complete curve (A) and average circumferential cracking percentage

(B) versus pin length of Heat B diluted with 10wt% Heat A ...... 89

Figure 47: CPTT complete curve (A) and average circumferential cracking percentage

(B) versus pin length of Heat B diluted with 25wt% Heat A ...... 91

Figure 48: CPTT cracking response curves comparing 10wt% and 25wt% dilutions of

Heat B by Heat A ...... 92

Figure 49: Graphical representation of cracking criteria in each heat ...... 94

Figure 50: Scheil STRs of Heats B (A) and A (B) ...... 97

Figure 51: Composition profiles for primary FCC (austenite) solidification for Heats B

(left) and A (right) ...... 98

Figure 52: Composition profiles for secondary FCC (MC-carbide) solidification for Heats

B (left) and A (right) ...... 99

Figure 53: Phase STRs for Heats B and A ...... 102

Figure 54: Evidence of cracking along SGBs in heats: A (A) and B (B); ...... 104

Figure 55: Grain size and secondary phase dispersion in heats: A (A) and B (B); ...... 105

xvi

Figure 56: Amount of eutectic liquid surrounding crack in heats: A (A) and B (B); ..... 106

Figure 57: No crack backfilling in Heat A (A), and complete crack healing in Heat B (B);

...... 107

Figure 58: Liquid coverage of a dendrite in Heat A fracture surface ...... 109

Figure 59: Fracture surfaces of Heats: A (A) and B (B) showing evidence of solidification cracking by “egg-crate” dendritic fracture surface ...... 110

Figure 60: Titanium-rich precipitates in Heats: A (A) and B (B) ...... 111

Figure 61: Precipitate morphologies in Heats: A (A)—rod-like; B (B)--spherical ...... 112

Figure 62: (Nb,Ti)-rich particle observed in sample containing 140 ppm N...... 117

Figure 63: NbC of irregular morphology ...... 118

Figure 64: Irregular NbC ...... 119

Figure 65: Evidence of "micro" cracks in 140 ppm N sample...... 119

Figure 66: Nb-rich particles decorating the matrix of a 200 ppm N sample ...... 121

Figure 67: Evidence of Nb-rich precipitation or continual growth from existing Ti-rich constituent leading into crack surface ...... 122

Figure 68: (Ti, Nb)-rich and Nb-rich particles ...... 123

Figure 69: (Nb,Ti)-rich particle ...... 124

Figure 70: Cuboidal (Nb,Ti)-enriched precipitates in 600 ppm N sample ...... 126

Figure 71: Intergranular Nb-rich precipitate ...... 127

Figure 72: Intragranular and intergranular titanium-rich precipitation ...... 129

Figure 73: EDS of intergranular Nb-rich carbides ...... 130

Figure 74: Large, intragranular, agglomerated titanium-rich particles surrounding a pore

...... 130

xvii

Figure 75: Evidence of agglomeration of large titanium nitrides ...... 131

Figure 76: Nb-rich precipitation and/or continual growth from existing TiN ...... 131

Figure 77: (2) Nb-rich precipitation and/or continual growth from existing TiN ...... 132

Figure 78: (3) Nb-rich precipitation and/or continual growth from existing TiN ...... 133

Figure 79: (4) Nb-rich precipitation and/or continual growth from existing TiN ...... 134

Figure 80: Interface between Ti-rich precipitate and NbC ...... 135

Figure 81: Dendritic morphology of shrinkage porosity ...... 135

Figure 82: CPTT results for the nitrogen addition study ...... 139

Figure 83: CPTT complete curve (top) and average circumferential cracking percentage

(bottom) versus pin length of Heat YT0159 with 19 ppm N [12] ...... 154

Figure 84: Figure 68: CPTT complete curve (top) and average circumferential cracking percentage (bottom) versus pin length of Heat YT0159 with 140 ppm N...... 155

Figure 85: CPTT complete curve (A) and average circumferential cracking percentage

(B) versus pin length of Heat YT0159 with 200 ppm N...... 156

Figure 86: CPTT complete curve (A) and average circumferential cracking percentage

(B) versus pin length of Heat YT0159 with 600 ppm N...... 157

Figure 87: CPTT complete curve (Top) and average circumferential cracking percentage

(Bottom) versus pin length of Heat YT0159 with 1300 ppm N ...... 158

xviii

CHAPTER 1: INTRODUCTION

FM 82 is widely used in the nuclear industry in dissimilar metal applications in transition joints and as overlays to avoid PWHT after field welding. In such applications, it is advantageous to use this alloy for its high temperature strength and good corrosion resistance. It can also be used to prevent creep failures during intermediate temperature exposure. However, it has been determined that thick section, high-restraint overlays may be subject to solidification cracking or ductility-dip cracking (DDC) during welding.

Previous work by BWXT has shown considerable variability in cracking susceptibility among different heats of FM 82. Previous Cast Pin Tear Testing performed at OSU

(Figure 1) demonstrated these differences and correlated well with fabrication experience based on cracking response among different heats of FM 82.

The approach to understanding this variability issue can be divided into several phases. First, CPTT of “resistant” and “susceptible” heats of FM 82 must be completed to ensure valid correlation with fabrication experience. Three heats in particular (AB8573,

YB8908, B8142), whose cracking responses have been tested previously, will be tested again to validate the repeatability and reproducibility of the CPTT. From now on, these heats will be identified as Heat A, Heat Y, and Heat B. Dilution effects using a

“susceptible” heat and a “resistant” heat will also be tested.

Figure 1: CPTT cracking response curves among 7 heats of FM 82 [1]

Next, computational modeling of Scheil solidification behavior will be conducted to better understand the phases that form on solidification and their composition. The solidification characteristics that will be calculated and compared among heats are: liquidus and solidus temperatures, overall solidification temperature range (STR), individual phase STR, phase compositions, elemental partitioning coefficients, and volume fraction of each solid phase.

Metallographic analysis is necessary to observe any differences in grain size and structure, crack initiation, precipitation dispersion, and crack backfilling. Fractographic and EDS analyses will be utilized to observe precipitate morphology and gather compositional data necessary to compare and contrast the bulk matrix and constituents among heats. Lastly, the effects on solidification cracking of N2 in the welding gas mixture will be observed. Four mockups of varying N weld metal contents will be tested using the CPTT and analyzed using SEM and EDS techniques. The four mockups are

made up of a resistant heat, YT0159 (Heat YT), and will be welded with 0.2, 0.4, 1 and 5 vol%N2 in an Argon-Nitrogen gas environment.

CHAPTER 2: BACKGROUND

During initial construction of many nuclear power plants, filler metal 82 was commonly used for dissimilar metal welds in components used in the reactor coolant system [2]. In using Ni-based alloys, this eliminated the need for a postweld heat treatment after field construction. In recent years, it has been observed that Alloy 82/182 welds made in these components may exhibit extensive cracking due to primary water stress corrosion

(PWSCC)—a subject that has received much attention in the nuclear power field [3].

A much less studied subject is the more recently observed solidification cracking in FM 82. “Micro” solidification cracks have been seen in FM 82 butter layers but are undetected by nondestructive evaluation techniques. Moreover, different heats with slightly varying chemical compositions are showing drastically different cracking responses. In order to reduce the presence of cracking and to understand the variability in cracking response in FM 82, it is necessary to have a fundamental understanding on the mechanism which causes cracking, as well as have a reliable and reproducible way to rank cracking susceptibility in an alloy system.

Chapter 2 is dedicated to the background of the welding metallurgy and weldability of Ni-based alloys. The mechanisms of solidification cracking are discussed, as is the evolution of the Cast Pin Tear Test; a reliable weldability test for ranking susceptibility to solidification cracking.

2.1 Welding Metallurgy

ERNiCr-3 is a solid-solution strengthened Ni-based alloy that solidifies as a fully austenitic microstructure upon cooling to room temperature. Microscopic segregation of alloying elements causes local compositional gradients in weld metal on both the solidification grain and subgrain levels. These gradients are present in the bulk matrix but are preferential to the grain boundaries where segregation of some alloying and impurity elements is prominent. Such segregation often results in the formation of secondary phases that may affect the weldability of these alloys. In order to understand the nature of segregation and the formation of secondary phases, it is critical to understand the fundamentals of solidification and the solidification structure. The three distinct boundary types associated with solidification can be observed in Figure 2. They are:

[1] Solidification subgrain boundaries

[2] Solidification grain boundaries

[3] Migrated grain boundaries

Figure 3 is a photomicrograph of the austenitic microstructure of FM 82 (Heat A) at 1000X magnification. Apparent solidification structures in the image include solidification grain boundaries (black) and solidification subgrain boundaries (containing dark spots within grains).

Figure 2: Schematic of grain boundaries seen in austenite-solidifying weld metal [4]

Figure 3: Typical FM 82 microstructure. SGBs and SSGBs in austenitic matrix (Heat A, 1000X magnification)

2.1.1 Solidification Subgrain Boundaries (SSGBs)

Solidification subgrains represent the finest structure that can be resolved under a light optical microscope (LOM). These structures are typically in the form of cells or dendrites, depending on the solidification behavior upon cooling. The boundaries dividing the subgrains from each other are known as solidification subgrain boundaries

(SSGBs). These boundaries are distinguished by their compositional differences from the bulk matrix of the microstructure. SSGBs are considered to be crystallographically low- angle boundaries due to their preferential solidification in “easy” growth directions. Due to this low-angle structural misorientation, SSGBs tend to contain low dislocation densities and therefore, less induced strain [5]. Cracks and voids sometimes form along

SSGBs, but mainly tend to form along solidification grain boundaries.

2.1.2 Solidification Grain Boundaries (SGBs)

Solidification grain boundaries (SGBs) are the distinct intersection between groups, or packets, of solidification subgrains. Grain boundaries can be observed along the trailing end of the weld pool at the intersections of subgrains that preferentially grow to their crystallographic orientation. Since each “packet” of subgrains solidifies in a different orientation, grains grow into each other by what is called competitive growth. Where these packets meet is known as the solidification grain boundary. Contrary to SSGBs,

SGBs are characterized as boundaries with high-angle misorientation because of the competitive growth phenomenon and thus, create large networks of high density dislocations. Cracks tend to form at the SGBs under the right conditions.

Additionally, SGBs contain compositional gradients due to macroscopic (Case 3) solute distribution of alloying and impurity elements. Solute distribution to the grain boundaries may depress the melting temperature of the boundary and lead to an array of weldability concerns in Ni-based alloys such as solidification cracking and liquation cracking [5].

Solidification grain boundaries are made up of both compositional and crystallographic components. The high-angle, crystallographic component of SGBs are called migrated grain boundaries (MGBs). These boundaries tend to “migrate” from the compositional portion of the SGB upon cooling; the driving force for migration is to lower the inherent grain boundary energy. SGBs tend to have high boundary energy due to their tortuosity upon solidification. In order to lower their energies, MGBs pull away from the “parent” grain boundary in order to straighten its orientation. Migration can occur in weld metal upon solidification or upon reheating from multi-pass welds.

However, with proper formation of secondary phase particles, a grain boundary

“pinning” effect can ensue and inhibit any migration of the crystallographic boundary at all [5]. Grain boundary pinning has several weldability effects that will be discussed in this chapter.

2.1.3 Solute Redistribution

During solidification, solute redistribution occurs such that elements partition between the liquid and solid as dictated by the phase diagram of the alloy. In instances where the partitioning coefficient, k, is less than one, increased solute resides in both the solid and liquid, but more so in the liquid phase. During redistribution, mass transport of atoms

from solid-to-liquid and liquid-to-solid must take place to maintain equilibrium of the interface in the system.

This transport can occur via mixing in the liquid or by diffusion in the solid.

These forms of transport are what make up the boundary conditions by which solidification occurs. If mass transport is unrestricted in both phases, redistribution can occur according the equilibrium phase diagram. However, if solute redistribution is restricted in either the solid or liquid, solidification proceeds in a non-equilibrium fashion. In the case where solid diffusion is considered negligible and there is mixing in the liquid, microscopic weld solidification can be approximated. If mass transport is restricted in both phases, the case approximates macroscopic weld solidification where long range diffusion or mixing is not allowed in the system [6]. Macroscopic and microscopic weld solidification characteristics are fundamentally well understood today in the world of welding metallurgy and are necessary to understand alloy microstructure and characteristics.

2.2 Solidification

Many factors affect the solidification of a material which influences its properties and behavior. Material factors include composition, thickness, and physical properties like coefficient of thermal expansion and melting temperature. Process parameter factors include heat input, cooling rate, welding travel speed, and shielding gas used during welding. To avoid solidification cracking, it is important to recognize the fundamentals of fusion weld solidification.

2.2.1 Solidification of Fusion Welds

Certain aspects of weld solidification can lead to cracking and possible catastrophic failure of components especially when exposed to high restraint where solidification control is of upmost importance. One such aspect of solidification is the phenomenon of dilution. Contact between the melting filler metal and the base metal will initially cause melting back of the material and dilution of the filler metal by the base metal is inevitable

[7]. If the composition of the base metal is at all different from that of the filler metal, the amount of dilution is not insignificant and segregation is bound to occur due to compositional gradients in the microstructure. Dilution also affects overlays as it does weld joints. If, during the process of overlaying, varying heats of the same material are interchanged, significant implications may result. Dilution affects the weld metal, in that, the melt composition changes and the surface layer of the base metal is removed and may penetrate the melt.

From a weldability perspective, it is critical to understand the solidification of fusion welds, not only from a “micro” standpoint, concerning dilution and elemental segregation, but also from a “macro” standpoint. In large welding operations such as in the nuclear industry with the fabrication of pressure vessels, there is much inherent restraint with thick section welding of complex components. The joint geometry plays a large role in accounting for the weldability even if the material is known to be highly weldable. Figure 4 is a schematic which represents typical solidification grain morphologies in a fusion weld. Based on solidification behavior, columnar dendritic and equiaxed dendritic solidification modes are achieved. The transition between solidification modes is thought to be related to segregation patterns in these regions upon

the last stages of solidification. These regions are similarly observed in castings but to a coarser degree due to slower cooling rates. In addition to inherent restraint from self- fixturing, parameters such as welding speed have an influence on cracking susceptibility.

The change in travel speed will cause a change in solidification growth of columnar crystals according to the isotherms produced in welding [7]. Thus, the weld pool shape can vary from the acceptable elliptical geometry to a potentially detrimental teardrop geometry which is the cause for many centerline solidification cracks. These geometries are shown in Figure 5. The “micro” and “macro” issues must both be considered when investigating solidification cracking in any alloy. The “macro”, or controllable, aspects of welding are easier to control and are generally the first prospect to consider if cracking is observed. If parameter-based aspects are accurately controlled and it can be determined that they are not the root cause for such cracking response variations, the behavior of impurities and other elements must be closely analyzed. Lastly, the effects of other solidification parameters on the microstructure and therefore properties of the materials in fusion welds must be considered [7].

Figure 4: Solidification grain morphologies due to welding [7]

Figure 5: Change in weld pool geometry from (L) elliptical to (R) teardrop based on welding speed [7]

2.2.2 Solidification in Castings

Regarding this study, solidification in both fusion welds and castings are taken into consideration. The Cast Pin Tear Test relies on the solidification of small castings to produce solidification cracks in the presence of restraint inherent to the geometry of the pin. The solidification of castings exhibits the same general trends as in fusion welds but with different solidification orientations and slower cooling rates. Due to the presence of a mold in CPTT, competitive growth of crystals takes place at heterogeneous nucleation sites along the mold wall and grow inward (Figure 6). Such sites can include residue from previous castings, defects along the mold, or uncleaned particles on the surface. Figure 7 shows the three different zones of solidification morphologies that are seen in castings.

These regions are classified as the (1) outer chill zone of equiaxed crystals, (2) a columnar zone of elongated grains, and (3) a central equiaxed zone [7]. The schematic in

Figure 7 is representative of the grain structure observed in a square bar upon solidification. The CPTT produces round, rod-shaped pins but follows the same solidification grain orientations.

The chill zone is representative of the outermost layer of seemingly spherical grains in Figure 7. This region forms when during the pour when the molten liquid comes

into contact with the much cooler mold wall. The liquid is then cooled rapidly below the liquidus temperature and many solid nuclei form along the wall and grow into the liquid towards the center of the cast object. Pour temperature plays a major influence on the solidification structure of the grains in the cast. With a low pour temperature, the whole of the liquid cools below the liquidus and any crystals swept into the melt may continue to grow and the liquid instantaneously solidifies as an equiaxed ingot structure (known as big bang nucleation) [7].

The columnar zone represents the middle region shown in Figure 7. Soon after pouring, the temperature gradient at the mold walls decreases drastically and crystals from the chill zone preferentially grow along certain crystallographic planes in the direction of the heat flow, or perpendicular to the walls. Just like fusion welding solidification, primary dendrite arms link up and form solid walls where the mushy zone is characterized as the region where the last liquid droplet solidifies between dendrite tips

[7].

Figure 6: Schematic of (a) competitive growth of crystals from mold wall (b) inward [7]

Lastly, the equiaxed zone contains randomly-oriented equiaxed grains in the center of the cast. It is understood that the presence of such grains are due to the melting- off of dendrite side-arms or secondary and tertiary dendrite arms (Figure 8). This phenomenon occurs when the melt is hotter than the formed dendrite—which induces the narrowest part of the dendrite to melt and be carried away by the convection patterns of the melt to reform as a dendrite or serve as a nucleation site for other arms to form in the last to solidify middle region of the cast [7]. An SEM image of a fracture surface of resistant Heat Y demonstrates all three levels of dendrite arms (Figure 8).

Figure 7: (L) Schematic of cast grain structure exhibiting (R) Photomicrograph taken at 1000X magnification of cast pin grain structure in Heat Y. Zones: (1) an outer chill zone of equiaxed crystals, (2) a columnar zone of elongated grains, and (3) a central equiaxed zone. [7]

Figure 8: Development of thermal primary, secondary, and tertiary dendrite arms. (L) Schematic, (R) Fracture surface of Heat Y [7]

In summary, despite the obvious differences between the two, several comparisons can be made between the solidification of fusion welds and castings—more particularly between the solidification of weld metal and castings. In fact, due to their commonalities, it is sometimes convenient to consider the fusion zone of a weld as a minicasting [8]. Most knowledge of weld pool solidification actually comes from the theories associated with the knowledge of freezing castings, ingots and single crystals at lower thermal gradients and slower growth rates. From these derivations, extrapolated theories have been made on rapid solidification of welds solidified at very high cooling rates. Just as in castings, the microstructure development in the FZ depends on the solidification behavior of the weld pool, or the region of metal that melted. In both processes, the principles of solidification dictate segregation behavior, solidification defects, and the size, shape and orientation of the grains. Such parameters can include growth rate (R), temperature gradients (G), undercooling (DT), and alloy composition. It

is when all parts of a weld are considered that weldments become very different from castings. So many characteristics of solidification are related to the moving heat source in welding—which creates simultaneous changes in growth rates and thermal gradients across the weld pool [8]. However, all in all, if the solidification of a weld is isolated to its fusion zone, many comparisons can be drawn between it and a casting. Thus, using the

CPTT and evaluating cast pins as fusion welds is sufficient.

2.3 Weldability

Weldability is defined by The Welding Institute (TWI) as:

“A measure of how easy it is to make a weld in a particular parent material,

without cracks, with adequate mechanical properties for service, and resistance to

service degradation.”

The weldability of an alloy depends on various factors such as weld design, welding technique, and the inherent nature of metals and encompasses all defects, residual stresses, distortion exhibited in its microstructure. Ni-based alloys, such as FM

82, solidify as a Ni-rich, face-centered cubic (FCC) phase known as austenite—which produces strong segregation of alloying elements and impurity elements. These alloys are susceptible to both “hot” and “warm” cracking phenomena—namely, solidification cracking and ductility-dip cracking, respectively.

This section will focus on the background of the mechanism behind solidification cracking, as well as how to avoid it.

2.3.1 Solidification Cracking

Solidification cracking is the featured failure mode studied in this work. Solidification cracking is a form of cracking that occurs during the terminal stage of solidification by a mechanism that is generally well understood. Cracking occurs due to the presence of liquid films at the SGBs, and sometimes SSGBs, at the end of solidification. If the inherent restraint due to welding overcomes the strength and ductility of the material with liquid accumulation at the GBs, shrinkage strains occur at these sites and intergranular failure will ensue. Thus, solidification cracking susceptibility is a function of both metallurgical factors as well as local strain present when solidification has concluded.

Although the mechanism is known to be well understood and certain precautions are known to help avoid cracking, many different theories have been developed on the matter. The most notable theories, listed from oldest to most recent, are [9] :

[1] Shrinkage-Brittleness theory

[2] Strain theory

[3] Generalized theory

[4] Modified Generalized theory

[5] Technological Strength theory

The Shrinkage-Brittleness theory, proposed by Bochvar [10], was developed in the 1940’s as result of a study done on aluminum castings. Data shows that upon solidification of aluminum alloys, some “coherency temperature” is reached where a coherent dendritic structure is formed. At this determined temperature, solid formations begin to interact and create rigid networks of solid-solid bridges. These bridges can accumulate shrinkage strain and begin the onset of what is called the “effective interval”

or, in other words, the brittle temperature range [11]. This interval is positioned between the the coherency temperature and the effective solidus of the alloy. Fracture occurs at the solid-solid bridges where strain accumulations are too high for the material’s inherent strength and ductility to handle. Figure 9 shows these regions based on a eutectic phase diagram compared with the cracking susceptibility of an alloy at varying compositions [6].

The size of the effective interval, which is based on composition and amount/distribution of eutectic liquid in the solid-solid network, dictates an alloy’s susceptibility to solidification cracking. As seen in Figure 9, an alloy of composition A exhibits the worst cracking behavior. At this composition, the least amount of liquid is present and thus, the strain cannot be accommodated. Compositions B and C demonstrate much better cracking behaviors due to the amount of liquid present at the end of solidification. Composition C in particular undergoes a eutectic reaction which leads to

“crack healing” when sufficient amount of liquid is present. Understanding the coherency temperature, brittle temperature range, and amount of liquid at the end of solidification will help better understand the cracking susceptibility of an alloy.

Figure 9: Shrinkage-Brittleness theory based on a eutectic phase diagram

The Strain theory was a theory developed by Pellini in the 1950’s. It was originally developed to describe hot tearing in castings, and later to describe weld solidification cracking, by casting samples with chilled and restrained ends inducing the onset of solidification cracking. This theory introduces the concepts of what Pellini called the “mushy” stage and the “film” stage during the solidification process [12]. During the welding process, the molten weld pool travels along the workpiece underneath the arc and extends closely behind the arc until the “mushy” stage is reached. In this stage, cracking is not possible due to the presence of sufficient liquid and no stable solid network [13]. As the near-solidus-temperature film stage is reach, the onset of liquid films is present on the grain boundaries within the matrix. The combination of the presence of GB liquid and the inherent strain focused on the GBs from solidification

shrinkage is a setup for solidification cracking. At this onset, liquid films are thick and continuous and the strain required to cause GB separation is very little. Furthermore, the film stage may extend below the equilibrium solidus temperature. Such segregate liquid films increase the solidification temperature range (STR) and thus, increase the susceptibility to cracking [14]. With this, the occurrence of fracture correlates with the time a solidifying material exists in the film stage and the strain rate imposed on the liquid films in this stage. This relationship is described in Figure 10 under segregate and non-segregate conditions.

Alloys under equilibrium (non-segregate) conditions cool through the film stage fairly quickly and are more resistant to solidification cracking. Contrarily, under segregate conditions, the effective solidus is depressed allowing for more strain accumulation during solidification through an expanded film stage and thus, increasing susceptibility to solidification cracking. Such intergranular cracking exhibits a smooth dendritic fracture surface morphology.

Figure 10: Strain theory according to Pellini [6]

The Generalized theory, proposed by Borland, is a modification of both the

Shrinkage-Brittleness and the Strain theories. Borland explains how liquid quantity and distribution during solidification impacts susceptibility to cracking [13] [14] [15].

Furthermore, solidification is described in four stages as illustrated on a binary phase diagram in Figure 11. These stages include:

Stage 1: Mushy stage-- Primary dendrite formation

Stage 2: Coherency range--Dendrite interlocking

Stage 3: Critical range--Grain boundary development

Stage 4: Solidification is complete

The Generalized theory, like the Shrinkage-Brittleness theory, introduces the concept of a coherency temperature (Figure 11: a-d) where a coherent solid network begins to form. Next, Borland divided what Pumphrey and Jennings proposed as the effective interval into two distinct regions. Stage 2 depicts a “coherency” range where a

solid dendrite structure is present and dendrite interlocking is observed. Cracks that form in this range may be healed due to the presence of sufficient surrounding liquid. Further solidification brings the onset of stage 3, or the “critical” range, where “packets” of dendrites grow together to form grain boundaries and there is not sufficient surrounding liquid to heal cracks formed in this region. Lastly, stage 4 represents the region of complete solidification where cracks do not form due to the absence of liquid in the system [6].

Figure 11: Stages of the Generalized theory of solidification cracking [6]

The Modified Generalized theory of solidification cracking was proposed by

Matsuda and coworkers from the Joining and Welding Research Institute (Osaka

University) in the 1980’s as an extension of the Generalized and Strain theories. This theory is based on direct observation of weld solidification cracking in carbon steels, stainless steels, and Inconel alloys during welding using an optical microscope [16].

Notable differences introduced in this theory occur in stages 1, 2, and 3. In all stages,

Matsuda et al. proposed these regions are shifted more near the liquidus. Additionally, in the cases of stages 1 and 2, they are said to occur over more narrow temperature ranges.

Contrary to Pellini’s Strain theory, solid-solid networks were dynamically observed to form much more rapidly in the early stages of solidification just below the liquidus than originally proposed [6] [12]. Thus, this region where cracking susceptibility is highest is shifted up closer to the liquidus temperature. This relationship is demonstrated in Figure 12.

Previous to this work, Borland described a critical temperature where cracking susceptibility is highest. In the Modified Generalized theory, this point is located between the liquid mass stage and liquid film stage. As shown from the ductility curve, crack initiation occurs in the liquid film stage and propagates in both stages on either side of it—propagating slower in the liquid mass stage than the liquid droplet stage due to crack

“healing”.

Additionally, stage 3 as proposed by Borland was further divided into two distinct regions. These regions are listed below and illustrated, along with the modified stages 1 and 2, in Figure 13:

Stage 3H: Film stage

Stage 3L: Droplet stage

The liquid film stage (3H) is the region where it was observed that weld solidification crack initiation, as well as propagation, can occur. The liquid droplet stage

(3L) is the region where it was observed that solely crack propagation occurs; crack initiation is not possible due to the extent of the solid-solid network created by this time.

Figure 12: Modified Generalized theory on interaction of solidification behavior, fractography, and cracking susceptibility. [16]

Figure 13: Observed stages of Modified Generalized Theory of solidification cracking [6]

It was further observed that fractographic features play a major role regarding crack behavior in stages 3H and 3L.The authors of this study describe three characteristic regions of a solidification crack as listed below and shown in Figure 14.

Type D: Dendritic fracture surface

Type D-F: Dendritic-to-flat transition fracture surface

Type F: Flat surface

Dendritic-to-flat fracture morphologies help characterize crack initiation and propagation in the 3H stage while flat features are representative of crack propagation in the 3L stage [6].

Figure 14: Three characteristic fractographic regions of a solidification crack [17]

Finally, the Technological Strength theory, proposed by Dr. Prokhorov in 1962, is a theory developed based on the mechanical behavior of materials during welding [18].

This theory took into consideration the fundamentals of the strengths of metals during the recrystallization process by understanding the deformability of the metal throughout the entire welding cycle. Prokhorov stated that the deformability of a material could be determined by its ductility, referring to the brittle temperature range (BTR) of the material where the magnitude of the elastic component in negligible to the plastic component. A reduction in ductility is apparent in metals when transitioning from liquid into the solid state and meanwhile, the deformations increase. Thus, it is expected that failure will occur in the BTR when there is an exhausting of the ductility in an alloy. The width and depth of the BTR are controlled by the solidification temperature range and the nature of the grain boundary liquid films [6]. Prokhorov illustrated this concept in the graph below in Figure 15. The solid lines represent the solidus and liquidus for a given alloy while the dotted lines show the lower and upper boundaries of the BTR [18].

This concept compares the minimum ductility/deformability, IImin, of a material with the accumulated deformation, ∆е, while passing through the BTR. The accumulation

of deformation through the BTR takes into account both the deformation caused by the change in shape of the work, еf, and the free contraction deformation, ew. Altogether the technological strength can be summed up in one equation using these components:

∆е푚 = 퐼퐼푚푖푛 − (∆е푤 + ∆е푓)

Thus, the potential for solidification cracking is said to be a competition between said strain accumulation and the recovery of ductility during weld cooling. It should be noted that this theory does not take into account weld microstructure. For a simplified understanding, another diagram of this theory is shown in Figure 16. Referring to this plot of Ductility or Deformation (%) vs. Decreasing Temperature we can use the dotted lines to describe the different behaviors of the material:

[A-B]: Represents the thermal contraction of the material (in the BTR)

[A-C]: Represents the contraction deformation along with other mechanical forms

of deformation in the system

[A-D]: Represents the critical amount of deformation to cause cracking

*A line traversing through the BTR would represent a situation in which the strain accumulation exhausted the ductility of the material and cracking would occur.

Figure 15: Graphical illustration of the hypothesis of the Technological Strength theory [18]

Figure 16: Simplified version of hypothesis of the Technological Strength theory [6]

2.3.2 Avoiding Solidification Cracking

Fundamentally, three factors contribute to the solidification (and liquation) cracking mechanism: alloy composition, welding parameters, and restraint. Controlling these factors may be difficult in some cases but is necessary in order to avoid solidification cracking. These factors are described below.

Compositionally, if the base metal or filler material cannot be substituted from the procedure specifications, it can be very difficult to control the contents of certain alloying additions, let alone their behavior upon solidification. One principal theory on avoiding solidification cracking relates directly to the solidification temperature range (STR) of the material. If weld metals solidify over a more narrow temperature range, and therefore avoid, or at least suppress, the formation of low-temperature melting eutectic constituents, its susceptibility to solidification cracking is known to be generally low.

Much of the time, low-temperature melting eutectic constituents are known to be comprised of alloying additions such as Nb, Ti, and Si. It is also common practice to avoid impurities such as C, N, P, S and B. STR does not, however, take into account eutectic backfilling and crack healing which increases crack resistance.

Welding parameters are known to largely influence susceptibility to solidification cracking, even in weldable materials. Weld heat input is one parameter known to affect solidification cracking response. With a higher heat input, weld size increases and thus increases the amount of grain boundaries where solidification cracks tend to initiate.

Larger weld sizes also increase solidification shrinkage strains in the material, leaving them susceptible to cracking. Additionally, higher heat input leads to smaller temperature

gradients at the trailing edge of the weld pool and in the base metal HAZ, and thus increase the size of the region known to be susceptible to hot cracking.

Another parameter that affects solidification cracking susceptibility is welding speed. With increased welding speed, the chances of a centerline solidification crack also increases due to the nature of the orientation of the solidifying crystals in a teardrop shaped weld pool. For this case, solidification shrinkage strains are highest perpendicular to the weld centerline and can cause cracking at this location under the right conditions.

The effects of travel speed for an ideal weld and a weld producing a centerline crack are illustrated previously in Figure 5.

As previously explained, certain parameters can affect weld pool shape which may influence the susceptibility to cracking. Such shapes affect the grain structure of the weld and the inherent restraint of the weld. These concepts can be controlled by modifying the depth-to-width (D/W) ratio of the weld. Due to the inherent nature of grains growing perpendicular to the solid/liquid interface, a lower D/W ratio will produce a proper elliptical weld pool resistant to solidification cracking. Higher D/W ratios represent the case previously described with a teardrop weld pool shape. Regarding strains, even the weld surface contour can contribute to cracking potential. When the contour of a weld is concave, tensile strains are introduced to the weld and cause cracking. Tensile strains typically develop due to mechanically-induced residual stresses in the weld. Alternatively, convex weld contours introduce compressive strains that reduce susceptibility to solidification cracking in welds. Figure 17 illustrates the effect on solidification cracking of weld D/W ratio as well as bead size and location resulting in differing weld contours.

Weld D/W Ratio

Bead Size and Location

Too wide Washed up Flat or slightly and concave too high and convex, not concave quite full width

Figure 17: Effect of weld D/W ratio and bead size and location on solidification cracking [6]

In summary, if compositional changes are minimal or non-existent, the appropriate combination of process-controlled variables can help deter solidification cracking in Ni-based alloys. Low heat-input welding processes that produce weld shapes and contours that produce compressive strains will exhibit more resistant behavior to cracking. In fact, depending on the application, multi-pass welding of smaller welds may be the proper solution to cracking issues. A larger number of smaller passes will decrease residual stresses built-up in the joint, as well as decrease the heat input and amount of shrinkage strains in the welds. This will help resist solidification cracking (and liquation cracking) in the material.

2.4 Weldability Testing

This section explains the origin and evolution of the principle testing method involved in this research—the Cast Pint Tear Test. The various iterations of its evolution are:

[1] Original Cast Pin Tear Test Method

[2] Second Generation Cast Pin Tear Test

[3] Third Generation Cast Pin Tear Test

Each iteration since its beginning in the 1950s has further optimized this approach to rank a material’s susceptibility to solidification cracking.

2.4.1 Original Cast Pin Tear Test Method

The CPTT was developed in 1959 by F.C Hull to evaluate the susceptibility of alloys to hot cracking or, more notably, solidification cracking [19]. The method goes as follows:

Samples weighing nominally 19 grams are levitation-melted and cast in molds in the shape of tapered pins. Restraints at either end of the pin impose tensile stresses on the sample that may cause cracking as the mold expands and the metal contracts upon solidification. This test was made to both evaluate a material’s susceptibility to “hot” cracking and to observe the effects of alloying additions and impurities on said cracking.

In short, hot tearing is caused by casting molten metal into a pin mold and thus—Cast Pin

Tear Test [19].

Levitation was chosen as the method of melting for several reasons when compared to crucible melting; most notably, it avoids contamination of the melt. It also simplifies the pouring of the metal by utilizing a magnetic field which allows for adjustable pouring rates. Furthermore, electromagnetic stirring of the charge simplifies

the ability to study the effects of alloying additions on “hot” cracking susceptibility. The original levitation melter setup is shown in Figure 18. Inside the glass chamber of the melter sits a pedestal rack where the pin molds are allowed to rotate underneath the levitation coil. The coil is fed high frequency power through a water cooled coaxial lead in order to induce an electromagnetic field which lifts the charge. When ready, the power is lowered in the coil and the charge is cast in a controllable and repeatable manner.

The mold itself is designed to introduce restraint at both ends which imposes tensile stresses to the solidifying pin. The pin molds blueprint is shown Figure 19 where pin volume is kept at a constant 2.4 cm3. While the dimensions are arbitrary, they serve to rank susceptibility to hot cracking as a qualitative comparison. This ranking is better quantified in the OSU version of the CPTT described later in this chapter.

Circumferential cracks are counted systematically with a stereo-binocular microscope at 30 times magnification which permits angular measurement of crack length. Cracks observed along the axis of the length of the pin are not counted.

Furthermore, the cracking percentage is a summation of all the cracks observed in a pin, which may exceed 100% and thus, calls for the need of what is considered the “cracking index”. The susceptibility to hot tearing in materials is ranked using a plot of the cracking index versus mold number. An example of this plot is shown in Figure 20 where susceptibility curves for two stainless steel alloys are plotted. An important principle of this version of the CPTT is that as mold number increases, the mold length decreases and subsequently so do restraint and cracking percentage. Thus from the figure, Modified 316 is more susceptible to hot tearing than 7299.

Figure 18: Schematic of original levitation melter setup [19]

Figure 19: Pin mold dimensions [19]

Figure 20: CPTT ranking of stainless steels [18]

2.4.2 OSU Second Generation Cast Pin Tear Test Method

Several iterations of this test have been developed since its beginning in order to optimize the casting environment and parameters for repeatable testing of different alloy systems.

The second and third generation tests were developed at the Ohio State University. The second generation test utilized the button-melter setup at OSU where a sample was melted using a GTA welding torch and allowed to flow into a pin-mold below a water- cooled copper hearth. This setup can be seen in Figure 21 and the melting parameters are displayed in Table 1. A schematic of the present button-melter setup is shown in Figure

30. At this time, there were two other major alterations made in the second generation

CPTT regarding the pin molds. In the second generation test, the molds were made of a copper-beryllium alloy and they were not tapered along the length of the gauge. The

geometries of the head and foot impose restraint on the pin, allowing stress to accumulate in the last to solidify liquid. With higher restraint, the occurrence of cracking is more prevalent. Besides dimensional changes in the pin diameters and lengths, the concept in this version is similar to that created by Hull.

Figure 21: Second generation CPTT at OSU [13]

Table 1: Second Generation CPTT Parameters

2.4.3 OSU Third Generation Cast Pin Tear Test Method

The third generation CPTT was developed at OSU by Luskin [14]. This iteration is what is currently used to rank alloy systems on their susceptibility to solidification cracking.

To begin the process, wire or sectioned samples of the desired material are melted into “buttons” which weigh between 10.3-15.0g. The button is placed into a water-cooled copper coil and is levitation melted at an AC current at 434 Hz until a set temperature is reached. Once it is reached, the current in the coil is ramped down and the charge is cast into the molds of constant diameter and varying lengths (depending on the mass of the button). A schematic of the current third generation CPTT setup can be observed in

Figure 31.

The second and third generation CPTT procedures both use the same method for analyzing cracking in pin samples. Cast pins are examined at 10-70X magnification and circumferential cracking is measured where 100% cracking is translated to one full 360° rotation of the pin. For the second generation CPTT, a response curve of Maximum

Circumferential Cracking (MCC) vs. Pin Length was plotted. More recently, the Average

Circumferential Cracking vs. Pin Length is utilized (Figure 22). The maximum pin length with an average of 0% cracking and the minimum pin length with an average of 100% cracking are the quantitative criteria by which alloys are ranked. The former is designated as the Lower Cracking Threshold (LCT), whereas the latter is the Upper

Cracking Threshold (UCT). A more in-depth description of the Third Generation CPTT procedure is described in Chapter 4.

Figure 22: CPTT average cracking response curve labeled with ranking criteria

2.5 Nitrogen Effects on Ni-based alloys

With such slight variations in composition among heats, it is unlikely for such a profound variation in cracking susceptibility to exist. Therefore, it is necessary to examine what role interstitials may be playing if a solution to the variability is not so apparent. For the sake of this work, background probable cause has raised suspicion of the effects of nitrogen on FM 82.

It is well known that the origin of hot cracks, such as solidification cracks, can be traced back to unfavorable combinations of material behavior and composition, design, and manufacturing influences; not to mention process parameters, temperature dependent properties of the material and existence of low- melting temperature eutectics [20]. Ni- based alloys are known to raise weldability concerns in certain circumstances and researchers are continually testing the effects of new alloying addition combinations to resist cracking in these materials. That being said, nitrogen is a low-cost and easily available element but is not commonly used as an alloying element for Ni-based alloys.

In fact, it is almost always present as an impurity in solid-solution austenite matrices, and its presence in weld metal can be very difficult to control. Interstitials have been known to be beneficial in combating some cracking phenomena—such as ductility-dip cracking

(DDC) [21]; solidification cracking, however, is not one of them. Impurities tend to combine with other elements in the matrix to form low-melting temperature constituents, expanding the solidification temperature range and produce GB liquid films which spread and cause cracking.

Small elements such as carbon (C), oxygen (O), phosphorous (P), sulfur (S), and nitrogen (N) are often difficult to detect in microscopy and energy dispersive spectroscopy (EDS) analyses due to their size, sparsity, and miniscule contents relative to other elements in the matrix. However, there is less difficulty capturing the presence of these elements in particles formed in the matrix such as nitrides and carbides—both of which are proven to form in FM 82. In fact, it is reported in the literature that stable, high-temperature forming nitrides (TiN) may act as sites for continual growth of other lower-temperature forming constituents (NbC) [22]. Thus, nitrogen is the interstitial element of interest in this work.

In an austenitic matrix, nitrogen is known to be a strong solid-solution strengthening element and austenite stabilizer, especially in austenitic stainless steels

[23]. In case of Ni-based alloys, however, N is an interstitial and often known as an impurity. If the permeation of nitrogen into the metal A is appreciable, the precipitation of the solute B, in the form of nitrides, may also occur inside the alloy. It is said that the precipitation reactions of these nitrides depend on the activity of the nitrogen in the atmosphere, the activity of solute B in the solid solution matrix A(B), and the diffusivities

of both nitrogen and the solute element in A(B) [24]. One objective of this work is to investigate the interaction between nitrogen and the matrix of FM 82 and to determine if these resultant constituents contribute to cracking susceptibility.

According to Sievert’s Law, nitrogen dissolves in liquid metal based on the following reaction:

1 푁 ↔ 푁 2 2(푔)

Also, the equilibrium constant of nitrogen can be calculated as follows:

푎 푓 ∗ [%푁] 퐾 = 푁 = 푁 √푝푁2 √푝푁2

Where aN, activity of nitrogen; fN, activity coefficient of nitrogen in liquid alloys; [%N], wt% nitrogen dissolved in the alloy; pN2, nitrogen partial pressure in the gas phase.

The knowledge of the solubility of nitrogen in both the liquid and solid state of

Ni-based alloys is pertinent to supplement research on its influence on solidification cracking, and its physical and chemical properties. In 2002, Kowanda and Speidel conducted a study to investigate the solubility of nitrogen in liquid nickel and binary Ni-

Xi (Xi= Cr, Mo, W, Mn, Fe, Co) under elevated pressure [25]. The study found that the highest increase of nitrogen solubility in Ni-based alloys was correlated with increasing

Cr content when compared to the other elements (Figure 23). This is important when discussing FM 82 and Ni-based alloys in general, because Cr is often the second-most abundant element in the matrix besides Ni. It is in the interest of this work to develop a preliminary look into the effects of both isolated and interactive nitrogen in the weld metal of FM 82.

Figure 23: Nitrogen solubility curve versus alloying element content in at under a nitrogen partial pressure of pN2=1bar [25]

The benefits of nitrogen additions on mechanical and chemical properties of Fe- based alloys are well known, but much less are known about the influence of the interstitial on Ni-based alloys. The following is a brief compilation of the work that has been done on nitrogen effects in Ni-based alloys regarding nitrogen in shielding gas during welding, effects of nitrogen on mechanical properties, and internal nitridation of

Ni-based alloys. Additionally, few mentions are made on the few studies conducted relating nitrogen effects on weld solidification cracking.

2.5.1 Investigation on the use of Nitrogen in Shielding Gas during Welding of Ni-based Alloys and Materials with Austenitic Microstructures

Several works have been conducted regarding the influence of nitrogen in shielding gas during the welding of Ni-based alloys, austenitic materials, and materials in general [20]

[23] [26] [27]. Shielding gas is often supplemented with nitrogen additions for certain alloys, however, adding N2 to the gas mixture while maintaining strict control of the amount of N transferring into the weld metal is very difficult. Additional troubles include weldability concerns with additions in austenitic stainless steels—those being electrode erosion, excessive spatter and arc instability [28]. It is, however, commonly seen that nitrogen increases tensile strength and hardness in the weld metal—especially when using the A-TIG process [23] [27]. More detail about mechanical properties is discussed in the next section.

An interesting study by Murphy et al. was done to put forth a computational investigation of the effectiveness of different shielding gas mixtures for GTAW—one of which was nitrogen [26]. Since the heat flux generated through an argon arc is relatively low, different additions were observed in order to get better penetration depth of the weld pool. Several conclusions were made on the effects of the argon-nitrogen gas mixture for welding. First, it was observed that molecular gases such as nitrogen and hydrogen created greater arc constriction due to the pinch effect they exhibit—meaning that less amounts of these need to be added for the same temperature increase in the solid anode if another gas such as helium were being used. Due to the better arc constriction, a greater magnetic pinch force will allow the weld pool to penetrate deeper. Computational modeling of these mixtures might create a push for further gas mixture development without spending money and time using actual mixtures. It was even mentioned that there is the possibility of utilizing ternary gas mixtures to expose the good properties of each gas on and in the weld pool.

Although, like Ni-based alloys, much contradiction in the literature provides for some skepticism to the real effect nitrogen plays on austenitic stainless steels. It has been stated that nitrogen both increases [23] and decreases [27] angular distortion in austenitic stainless steel plates when welded using GTAW—this discrepancy may have to do with differing weld penetration depths. In the work done by Huang, spot varestraint testing showed that hot cracking susceptibility dramatically increased with increasing nitrogen contents [27]. However, Ogawa et al. stated in their work that nitrogen additions, depending on whether or not the (γ+δ) dual phase is stable, can both increase and decrease susceptibility to solidification cracking in austenitic stainless steel [28].

Furthermore, recent research has shown that higher energy density processes decrease cracking susceptibility. The A-TIG process helped to decrease hot cracking susceptibility in austenitic stainless steels along with lower tensile stresses exhibited from lower angular distortion. It can be observed that testing and characterization of nitrogen additions in all alloy systems, including Ni-based alloys, needs to be further conducted.

2.5.2 Effects of Nitrogen on Microstructure and Mechanical Properties

The literature contains several works dealing with the effect of nitrogen (and other alloying additions) on the solidification, microstructure, and mechanical properties of Ni- based alloys [29] [30] and superalloys [31] [32].

Nabavi et al. more recently studied the weld metal microstructures and mechanical properties due to welding nickel alloys, Alloy 263 and Alloy X, with different Ar-N2 shielding gas combinations [29]. The compositions of these materials are shown in Table 2. Autogenous GTA welds were employed with additions of 0-4 vol% N2

in Ar. It was reported that, as suspected, with increasing amounts of N2 in Ar gas, the N level in the weld metals of each materials increased. Additionally, Figure 24 shows that the dendritic structures of each material in the weld metal were refined with increasing N2 in Ar gas, which was identified as the main cause for augmented toughness with increasing N contents in the weld metals of both alloys. Dendritic refinement by N has also been confirmed in other reports [31]. Additionally, the UTS of each material increased with increasing N content in the weld metal. Other conclusions include the increase in MC-type precipitates with increasing N content in the weld metal due to heterogeneous nucleation form the refined dendritic microstructure and increase nitride formation in the weld metal.

Other reports have shown N to influence other microstructural degradations such as microporosity, precipitate morphology, and eutectic fraction in cast structures [31].

This work is described in more detail later on.

Table 2: Chemical Compositions (wt%) of Test Materials Alloy C Si Mn Cr Co Mo Fe 0.05 0.25 0.3 19.58 19.1 5.9 0.48 Alloy Al Ti Cu W S N Ni 263 0.2 2.38 0.1 - 0.0017 0.0062 Bal. C Si Mn Cr Co Mo Fe 0.1 0.15 0.75 22.17 1.3 9.02 18.5 Alloy X Al Ti Cu W S N Ni 0.15 - 0.3 0.6 0.0053 0.0137 Bal.

Figure 24: Grain refinement in weld metal due to increasing N2 additions in Ar gas. (A) Pure Ar and (B) Ar-4%N2 for Alloy 263, (C) Pure Ar and (D) Ar-4%N2 for Alloy X

An experiment on the effect of minor alloying additions on the solidification of single-crystal Ni-based superalloy, CMSX-4, was conducted by Cutler et al [32]. They studied the effects of the combination of N and C additions on segregation, castability and solidification defects in CMSX-4 [32]. Furthermore, they studied the effects of these additions on the primary MC-type carbide morphology and composition in this single crystal Ni-based superalloy. Nitrogen additions were added just before pouring by means of CrN powder wrapped in Ni-foil, while C was added by appropriate amounts of graphite wrapped in Ni-foil. One baseline heat and three modification heats were cast and examined. Solidification defects being examined included high-angle boundaries, low- angle boundaries, freckles and slivers. It was observed, and proved to confirm findings in

other reports [33] [34], that nitrogen in fact increases the amount of microporosity in the microstructure during solidification; meanwhile carbon additions did the opposite. It is also hypothesized that the blocky-type carbides were less effective at reducing fluid flow in the interdendritic regions and this results in an increase in observed defects.

Transformation temperatures of sample containing C+N additions did not change much which is believed to be due to the small amounts of the additions. However, it was observed that carbide precipitation temperatures of the N-containing alloy were lower than that of the C-only alloy. This observation indicates that these carbides form lower in the mushy zone. Carbide morphology was also studied in this experiment. It was observed that the N-addition alloy exhibited blocky carbides/nitrides as opposed to the C- only samples which produced carbides of dendritic-morphology [32].

Huang, Zhang, and Hu conducted an experiment on the effects of small amounts of nitrogen on properties of a cast Ni-based superalloy, MAR-M002—it is important to keep in mind there are large differences in composition between this material and FM 82.

After additions of CrN, each melt was held at temperature for 5 minutes and three castings were produced: (1) virgin with no additions (5 ppm N analyzed), (2) virgin +

CrN (22 ppm N analyzed), and (3) virgin + CrN (30 ppm analyzed). The conclusion from their study was that increasing the nitrogen content in the alloy resulted in a change in carbide morphology from “Chinese script” to blocky form—a similar morphology alteration trend noticed in the work of this research on FM 82. Furthermore, they observed an increase in microporosity throughout the specimens which is affected by the evolution of dissolved nitrogen, whose solubility decreases with the drop in temperature during solidification after casting [31] [35]. The increase in porosity could also

potentially be due to reaching the alloy’s solubility limit of nitrogen. Such porosity was hypothesized to be the main reason for the change in properties observed as more nitrogen was introduced from sample-to-sample. It was found that the ductility, tensile strength, and stress-rupture life of the conventional castings dramatically decreased as nitrogen content increased [31]. All of which can may contribute to crack susceptibility.

It was also found that with increasing nitrogen contents, there was no change in eutectic morphology; however, there was a reduction in the volume fraction of the eutectic microconstituents.

2.5.3 Internal Nitridation of Ni-based Alloys

Much of the studies on general nitride formation in metals are in the form of nitridation behavior noted in the literature in ferrous alloys due to the well-known commercial applications that utilize nitrides for several advantageous property characteristics.

However, several recent works have reported on several approaches to the concept of internal nitridation in Ni-based and Ni-Cr alloys in certain temperature ranges using different environments [24] [36] [37] [38]. Internal nitridation is the formation of nitrides in the matrix attributed to the instability of the protective oxide layer resulting from erosion of the surface, growth stresses in the layer, crack formation due to creep, or fatigue loading, and superimposed thermal cycling [36]. This phenomenon usually takes place as a consequence of spalling and cracking of the protective oxide layer of a metal, which is typically considered impermeable to nitrogen. Figure 25 illustrates the process of internal nitridation in Ni-based superalloys. Since the attack site of nitrogen into the material is extremely localized, penetration into the matrix is rather heterogeneous. This

phenomenon takes place at relatively high temperature in solid solution, and therefore cannot be directly applicable to welding; however the behavior of nitrogen in the presence of other alloying elements is applicable to this work.

When unstable phases are precipitated out of solution in the process of internal precipitation reactions, atomic distribution can be predicted using thermodynamic considerations of the respective system being observed. Thus, the interactions between solute and penetrating atoms, N in this case, in the matrix of the solid solution must be taken into account. Work done on high temperature nitridation of Ni-Cr alloys by

Kodentsov et al. determined N, in fact, diffuses “uphill” its own concentration gradient in the FCC Ni-Cr solid solution matrix due to the interactions produced in the presence of

Cr [36]. Interestingly enough, it was experimentally determined that higher Cr concentrations (~17-20wt %) are necessary for any appreciable nitride formation to occur at all. FM 82 has an approximate Cr content of 18-22wt%.

Figure 25: Internal nitridation in Ni-based superalloys due to localized erosion in the form of surface cracks [36]

Furthermore, in ternary systems of Ni-Cr-Al and Ni-Cr-Ti it was observed that increased nitridation only occurred due to strong dependence on the respective Cr content of the alloy [37]. Several constituents were observed at higher Cr contents than 20wt%.-- those being CrN, Cr2N, in addition to TiN and AlN. Black and blocky nitrides were identified as AlN that resided in a zone of homogenously distributed TiN. AlN were not seen FM 82, but TiN were often analyzed in conjunction with NbC. Although both elements are present in some amount in the solution of the material, TiN are thermodynamically more stable and are more commonly seen in FM 82 in the literature.

Also, the presence of such nitrides was seen to reduce the progress of AlN precipitation

[37]. However, it was unseen in the literature of any form of Cr-nitrides in FM 82, but rather Nb and Ti-carbonitrides. It is likely that comparisons can be made between chromium in Ni-Cr alloys and titanium in FM 82. In Ni-Cr alloys it was hypothesized that grain boundary precipitation of chromium nitrides can be attributed to the likely presence of Cr at the GBs due to segregation while GBs remain preferential precipitation sites [24]. Cr, however, does not segregate to the GBs in FM 82 as in these Ni-Cr alloys.

In fact, Ti and Nb can be found at the GBs while chromium remains in the bulk matrix.

Lastly, work has been done on the thermodynamic characteristics and numerical modeling of internal nitridation of Ni-based alloys and shows promise to predicting secondary phase formation in Ni-based alloys [38].

2.5.4 Nitrogen Effects on Solidification Cracking in Ni-based Alloys

Regarding direct influence of nitrogen on solidification cracking susceptibility, there are contradicting claims in the literature. One particular experiment testing the effects of

nitrogen on welding Ni-based alloys was done by Herold, Zinke, and Hübner [20]. This study was a systematic investigation to the metallurgical effect of nitrogen additions as a gaseous component in weld shielding gas for avoidance of hot cracks of all types in fully austenitic weld metal. As such, externally loaded, 5mm samples of hot-crack sensitive

Ni-based (and Fe-base) materials were tested using the Programmable Deformation

Cracking Test (PVR test). In this test, a horizontally arranged flat tensile specimen is drawn at a linearly increasing tensile speed, VPVR, in the welding direction as welding is occurring. The measure for hot cracking susceptibility is the critical speed, or VCR, when the first crack occurs. Thus, the higher the critical speed, the higher the hot cracking resistance. The Ni-based alloys utilized in the experiment included Alloy 602 CA (BM,

FM), Alloy 601 H (BM), Alloy 690 (BM, FM) and Alloy 617 (BM, FM). The authors concluded of all the alloys tested, a positive effect of nitrogen was observed first and foremost only in the primary carbide containing Alloy 602 CA. This alloy contains approximately 0.19wt%C and varying additions of Yttrium and Zirconium which readily form carbides in welding. The other alloys in the study, however, did not appear to have been positively influenced in a consistent manner by increasing nitrogen content. Thus, further work must be conducted to determine which alloys are affected by nitrogen and in which manner. The welding conditions (Table 3) for the study and PVR results for

“dummy” welds of these alloys using differing gas mixtures (Figure 26) are presented below. Only results from the PVR tests utilizing GTAW are shown.

Table 3: Welding Conditions for GTAW and pulsed GMAW Processes [20]

Figure 26: PVR test results of different base metals susceptible to hot cracking using GTAW with varying gas composition [20]

Although an effort to correlate secondary phase presence to hot cracking susceptibility was not made, several differences in the microstructure in Alloy 602 CA welds were made compared to those of a sample welded in 100% argon. Relative size and shape of the carbides in the presence of nitrogen were shorter and less compact than those of a complete argon environment. Also, carbide morphology of the chromium-rich precipitates was noted to be different. Lastly, and perhaps most importantly, other small globular particles analyzed in this work were observed to contain higher contents of

nitrogen as well as titanium, zirconium and yttrium —the former two not being present

FM 82. Further consideration of their influence on hot cracking susceptibility was not noted [20].

CHAPTER 3: OBJECTIVES

3.1 Cast Pin Tear Test

The CPTT has proven to be a reliable and continuously optimizable tool for ranking susceptibility to solidification cracking in Ni-based alloys [13] [14]. It is in the interest of this work to evaluate the reproducibility of the third generation CPTT at OSU by retesting three heats of FM 82 to compare cracking responses observed by different users.

Meanwhile, further development with the CPTT procedure for optimized casting will be implemented. Developing these characteristics of the test will help produce more accurate evaluations of alloys’ rankings of solidification cracking susceptibility. The second aim is to rank solidification cracking susceptibility of other FM 82 variations to see how different dilutions and alloying additions affect the cracking responses.

1. Verify lower cracking thresholds and overall cracking responses of three heats of FM

82 using CPTT

2. Determine the effect of diluting a “resistant” heat by a “susceptible” heat on weld

solidification cracking using CPTT

3. Evaluate the effect of nitrogen additions on the solidification cracking susceptibility

of a resistant heat of FM 82

4. Further optimize CPTT parameters for Ni-based alloys

3.2 Solidification Cracking Performance and Metallurgical Analysis of FM 82

FM 82 is widely known to have delayed failures pertaining PWSCC in nuclear power applications. However, “micro” solidification cracking is being observed is many heats of this material, with different levels of crack susceptibility within a range of heats.

Determining the root cause of the solidification cracking is being investigated in this work, along with ways to optimize the cracking resistance for this filler metal. To do this, it must first be determined what distinguishes a resistant heat from a susceptible one.

Metallographic and SEM/EDS techniques will be used to compare features among heats as well as compositional data.

1. Conduct thermodynamic simulations of three main heats of FM 82 to compare their

solidification behaviors

2. Characterize three heats of FM 82 (of “typical” compositions), and one resistant heat

with four levels of increased nitrogen content.

3. Compare (by observation) bulk microstructural features such as grain size, crack

initiation sites, secondary phase formation, and eutectic backfilling

4. Compare (by EDS) compositions of bulk matrix and observed constituents in each

heat

5. Conduct fractographic study to observe precipitate morphologies of each heat

CHAPTER 4: MATERIAL AND EXPERIMENTAL PROCEDURE

4.1 Material

The material studied in this work was INCONEL® Filler Metal 82—classified as

ERNiCr-3 (AWS A5.14), or known as EN82H or FM 82. Filler Metal 82 is a Ni-based filler metal used for gas tungsten arc, gas metal arc, and submerged arc welding processes. It is commonly used to weld INCONEL® alloys 600, 601, and 690 or

INCOLOY® alloys 800, 800HT and 330. Weld metal depositions of FM 82 tend to exhibit high strength and good corrosion resistance, along with oxidation resistance and good creep-rupture strength at elevated temperatures. Common dissimilar-welding applications associated with FM 82 include joining stainless steels to nickel alloys and carbon steels, often in the form of “buttering” layers used in pressure vessel construction in the nuclear power industry [39].

Three different heats of FM 82 were tested using the CPTT and compared using

SEM/EDS techniques to better understand the solidification cracking behavior of this alloy. The compositions of each heat varied slightly from each and are shown in Table 4.

Table 4: Measured Chemical Compositions of Tested FM 82 Heats AB8573 YB8908 B8142 Heat (susceptible) (resistant) (resistant) Al 0.026 0.02 0.051 C 0.038 0.049 0.027 Co <.001 0.001 0.002 Cr 19.27 19.52 19.68 Cu 0.005 0.003 0.01 Fe 1.11 1.14 1.05 H 0.0001 0.0001 0.0001 Mg 0.024 0.01 0.001 Mn 2.84 2.82 2.84 Mo <0.001 0.005 0.002 N 0.0257 0.0210 0.0110 Nb 2.37 2.47 2.37 Ni 73.8 73.4 73.5 O 0.0020 0.0020 <.001 P <.005 <.005 <.005 S <0.001 0.002 <0.001 Si 0.17 0.16 0.096 Ti 0.35 0.37 0.36 V 0.006 0.005 <.005 W <.001 0.001 0.004

4.2 Weldability Testing Procedure

The overall testing procedure can be divided into six separate stages. Those stages are:

[1] Material Preparation

[2] Button Melting

[3] Cast Pin Tear Testing

[4] Crack Response Analysis

[5] Metallurgical Characterization

[6] Computational Modeling of Scheil Solidification

4.2.1 Material Preparation

In order to carry out a valid and repeatable procedure, the sample must be prepared the same. The three steps to preparing the filler material are cutting, cleaning, and weighing the wire. After receiving the material on a spool, safely contain it from surrounding environmental contaminants in a sealed plastic bag until further use.

The first step is cutting the wire appropriately so it can fit the hearth of the button- melter. In this case, the wire spools were cut using a custom wire-cutting machine in the

Ohio State welding metallurgy laboratory. This machine utilizes a Miller Electric motor- driven, automatic wire feed system (Figure 27) to continually feed the wire into a rotating blade. This blade utilizes constant rotation to ensure consistent cuts of approximately 1 inch per piece. Since only same-diameter filler metal was used in this study, each piece of wire weighed approximately 0.3 grams. Different weights must be considered if trying to conduct a multi-material, or dilution, study.

After cutting the material, it is transferred to the ultrasonic bath for a minimum of

15 minutes to remove any existing surface contamination. At OSU, the Branson 2510 ultrasonic bath is used. From this point in the process, nitrile gloves were worn to avoid any contamination from dirt or oils on the skin.

After cleaning, the wire is weighed on a scale capable of measuring to the 0.001 grams of the sample. Each sample is weighed to the appropriate amount based on what pin-length is desired. This relationship is outlined in. After each sample is appropriately weighed and sealed in a small plastic bag, they are ready for the button melting process.

Figure 27: Wire-cutter setup at OSU

Table 5: CPTT Pin Length (in) Per Given Button Mass (g) for FM 82 Mass Length (g) (in) 9.5 0.5 10 0.625 10.5 0.75 11 0.875 11.5 1 12 1.125 12.5 1.25 13 1.375 13.5 1.5 14 1.625 14.5 1.75 15 1.875 15.5 2 16 2.125 16.5 2.25 17 2.375 17.5 2.5

4.2.2 Button-Melting

The button-melter apparatus is used to melt the filler wire contained within its chamber into “buttons” for various applications (Figure 28). One application is to do optical microscopy on the sample. The process that takes place in the apparatus represents the high-temperature-melting, complete liquid mixing, and ambient cooling rate of a welded sample. Another use for button-melting is to transfer the sample to the Cast Pin Tear

Tester to perform testing of susceptibility to solidification cracking. Figure 29 shows the button-melting apparatus while Figure 30 represents a schematic of the apparatus.

Figure 28: Progression of filler metal sample via steps for CPTT (Top) and profile of "button" used for CPTT (Bottom)

Figure 29: Button-melter setup

Figure 30 Button-melting schematic

Before using the button-melter, it must be properly cleaned. Using a piece of 800- grit sandpaper to remove any contaminant residue from previous melting and create a visibly clean, copper surface. The surface is then cleaned with ethanol on a Kimtech wipe

until no residue is visible on the wipe. The apparatus is now ready for the material to be placed on the copper hearth. Place the wires in a unidirectional orientation for ease of complete melting and adjust the torch to a position capable of producing an arc. Next, it is critical to maintain certain purging parameters before any melting ensues. The purging procedure is a simple six-step process [13]:

[1] Ensure there is sufficient argon in the tank for melting

[2] Adjust argon flow rate to 20 CFH.

[3] Open the inlet valve to allow argon flow into the chamber.

[4] Close the upper exhaust valve until the chamber pressure exceeds 10 psi.

[5] Open the upper exhaust valve until the chamber pressure decreases to 0-5 psi.

[6] Repeat steps 4 and 5 six times

At this point, the argon concentration is homogenous throughout the chamber and the material is ready to be melted. Adjusting the supplied current to approximately 160-

170 Amps, melting is initiated using a GTA torch. Rotate the torch slowly around the material and observe how the surface tension of the material and the shape of the copper hearth combine to force the material into a “button” shape. For further mixing, continue to “hover” the torch above the molten button for approximately 3-5 seconds after the round and symmetrical “button” shape is achieved. This step may be unnecessary and/or detrimental to future results if noticeable evaporation takes place during the melting procedure. If evaporation is noticeable at this stage, eliminate it altogether. Lastly, let the sample cool for approximately 2-3 minutes. Other studies may call for more caution when melting—such as flipping the button over to melt it a second time for a more thorough mixing. This protocol is necessary for some dilution studies and for cases where

the filler wire being melted contains flux. However, this was not done with these heats of

FM 82. The next step of levitation melting in the Cast Pin Tear Tester also provides for further mixing. The parameters for the button-melting process are listed in Table 6.

Table 6: Nominal Parameters for Button-Melting Button-Melting Parameters Button Mass (g) 9.5-17.5 Argon Shielding Gas (99.998%) Gas Flow Rate (CFH) 20 Purge Pressure (psi) >10 Release Pressure (psi) 0-5 Number of Cycles 6 Power Supply Current 170 (A) Melt Time (sec) Varies on mass "Hover" Time (sec) 3-5 Cool Time (sec) 120-180

4.2.3 Cast Pin Tear Testing

The new generation Cast Pin Tear Test is a weldability test that ranks materials based on their susceptibility to solidification cracking via levitation melting and casting into a pin mold of a certain length. A basic description of the CPTT procedure is described here.

The main components of the machine are pictured in Figure 31. A 10kW induction power supply and work head operating at 224kHz charges current through a custom-designed, rectangular tubed, copper coil (Figure 32) which creates the magnetic field that controls the levitation of the button. Its geometry has been optimized to promote steady levitation and uniform heating of the button. The coil is constantly supplied with current until a casting set temperature is reached. An optical pyrometer attached to the top of the chamber reads the temperature of the button and once this set temperature is reached, the power supply ramps down the current in the coil to induce casting of the button into the pin mold.

Figure 31: CPTT setup schematic

Figure 32: CPTT custom water-cooled copper coil

Most of the purging and casting procedure in the CPTT is automated through communication with a configurable program using LabVIEW software. The software communicates through series of machines and hardware including a number of valves, an optical pyrometer and a proportional-integral-derivative (PID) controller (Figure 31). A screenshot of the LabVIEW control screen is shown in Figure 33 . This interface is where the user inputs given values of parameters depending on the material being used.

Before casting, the CPTT apparatus must be thoroughly cleaned to ensure there is no contamination from other materials. Cleaning involves an ethanol-soaked Kimwipe and complete wiping of the optical pyrometer, quartz funnel, coil, chamber glass and the rest of the inside of the chamber. The quartz funnel is then inserted into the coil to separate the charge from the coil prior to levitation and melting. The size of the funnel used depends on the mass of the charge being melted in the particular test. For charges 13 grams or less, a “small” funnel (23 mm upper diameter, 13 mm lower diameter, 13 mm height) is used. For charges greater than 13 grams, a “large” funnel (23 mm upper diameter, 15 mm lower diameter, 10mm height) is used (Figure 34).

Figure 33: LabVIEW control screen

Figure 34: (L) Quartz funnel used in CPTT and (R) engineering drawing of large funnel

After completing the cleaning steps and preparing the apparatus, the charge can be inserted into the coil, resting on top of the funnel via the top hatch of the chamber.

Next, the mold retainer can be assembled using the four main components that provide for quality cast pins. The first component of the retainer assembly is a combination of aluminum spacers at the bottom of the retainer. These spacers act as shims for the rest of

the assembly. The next component of the retainer assembly is the spring system. The spring plays a crucial role in applying the proper amount of “spring” pressure to the pin mold to inhibit liquid metal flashing. Springs with varying stiffness constants should be available for different mold lengths. The third component of the retainer assembly is the copper relief disc (Figure 35). The copper disc is important for two reasons: (1) if the disc is not present, molten metal will be cast through the retainer assembly without any barriers to stop it. This will ruin the other components of the assembly and possibly the retainer itself once it solidifies. (2) The grooves in the disc slow the flash and avoid spill- through to the rest of the assembly.

Lastly, and most importantly, is the pin-mold. Select the mold of the proper length corresponding to the mass of the charge (Table 5). Before a casting session, it is necessary to scrub the molds being used with a copper and (then) stainless steel wire brush. Once fully scrubbed of any oxide residue, the molds are cleaned using a felt string impregnated with 12 µm diamond paste over the gauged section of the mold in order to remove any surface contamination. This procedure helps further remove any oxide on the surface of the mold. Cleaning the head and foot of the mold is also critical so that these parts of the pin are clean upon casting and solidifying. Next, take the molds to the ultrasonic bath where they will bath in ethyl alcohol for approximately 12 minutes. After cleaning, dry the molds with a heat gun and ensure the pin holes do not contain any liquid—dry if so. This liquid will cause oxidation after casting. Next, insert alignment pins in the mold assembly and roll a rubber O-ring over the shaft towards the top of the mold.

Finally, the retainer assembly can be inserted into the retainer as shown in Figure

36. If the right spring pressure is applied, some force will be needed to push the mold down into the retainer. Screw the mold into the retainer, making sure that the seams between the two halves are perpendicular to the set-screw ensuring sufficient compression to the mold. The entire retainer can then be screwed in the bottom of the

CPTT apparatus, the gas outlet hose should be connected, and the top hatch of the chamber should be closed. Now the purge process may begin.

Figure 35: Engineering isometric drawing (L) and photograph of copper relief disc (R)-- green represents area of acceptable flashing, while red represents area unacceptable of flashing

Figure 36: Mold retainer assembly from R-to-L: Spacer, spring, copper relief disc, mold half

Referencing Figure 33 again, the LabVIEW control screen is where the purge parameters are set. Such parameters include: Initial Purge Time, Pressurize Time, and

Release Time. Once the parameters have been set, the purge process can be initiated by pressing Run and the Start Purge button. A flow chart of this process can be understood using the schematic below in Figure 37. The steps of the process are:

[1] Initial purge: High-flow valve and top release valve open (80 sec)

[2] Pressure build: Top release valve closes (4 sec)

[3] Pressure release: Top release valve opens (4 sec)

[4] Argon inlet: High-flow valve closes; Low flow valve opens and argon flows

into chamber through a 0.5 psi regulator; Meanwhile, variable flow rate valve

set to flow rate of 2.5 CFH and Chamber Pressure of 0.3 psi

[5] Green light illuminated on LabVIEW control screen ensures purge cycle has

completed

After the purge process is completed, the casting process begins by pressing Start

Melt from the control screen. In doing so, the power supply sends a current through the coil which ramps up causing the charge of material to levitate from an electromagnetic field created by the induction process. Once the Set Temperature is reached, the current ramps down from the Casting Current to the Ramp Down Current which allows the molten charge to cast into the mold. The current being supplied to the coil is then shut-off and the material begins to solidify. Once complete, press Stop Record and save the thermal history into a designated folder. The history will save as thermal data in a text file in case of the need for future reference. Remove the mold retainer from the chamber by disconnecting the gas outlet tube and unscrewing the retainer from the chamber.

Figure 37: CPTT purge cycle communication schematic

Loosen the set screws on the side of the mold holder and remove the mold carefully; keeping in mind that it is loaded with spring pressure. Separate the mold halves by prying them apart with a flathead screwdriver—some additional work may be needed to do so. Once removed, the molds and apparatus must be cleaned using the steps previous mentioned. The casting process is now complete.

The following parameter alterations were used and should be used for future work with FM 82 and possibly other Ni-based alloys:

a. Set Temperature was changed from 1430°C to 1450°C to ensure proper

complete mixing without excessive evaporation or “splashing” upon casting.

The exact average Set Temperature throughout 5 heats and 141 successful cast

pin samples was 1451°C. Due to coverage from evaporation of the sample, a

higher Set Temperature may be necessary to achieve this average.

b. Casting Current, especially for large buttons (>13g), was increased from 375

A to 380 A for two reasons: 1) to avoid any potential for the material to spark

the coil upon levitation—typical of heavier samples. 2) To allow the bottom

of a large button to be even with that of a small button to ensure near-equal

casting distances. The exact average Casting Current throughout 5 heats and

141 successful cast pin samples was 380 A.

4.2.4 Crack Response Analysis

After casting is complete, a crack response analysis is required of each pin. To analyze, each pin is secured in the rotating fixture shown in Figure 38 to be observed using a binocular microscope at 10-70X magnification. The circumferential cracks perpendicular to the pin axis are measured in degrees relative to a complete rotation of the pin (360°) using the circumferential crack counting device (Figure 39).

The percentage of cracking is calculated by the equation as follows:

° 퐿푇⁄ %퐶푟푎푐푘푖푛푔 = 360 × 100

Each user is trained to observe cracks alike so data can be accurately compared between different studies if necessary. A light source is used to aid in visual examination of cracks along the pin surface. If a pin exhibits cracking in multiple locations and/or is not a continuous crack, the total circumferential cracking for the pin is the sum of all cracks present (without overlap), where maximum cracking is 100%. If multiple cracks overlap, the overlap region is counted as a single crack. Pins with 0% cracking are examined at the highest magnification to insure they are crack free. Pins with 100% cracking normally exhibit complete separation of the pin head upon removal from the pin mold. A flaw is determined to be a crack when the light source is shined on it from all angles and no light reflects back at all.

During the crack counting process, many other features may be observed in the pins. Being able to decipher between these features and an actual crack is of upmost importance to this process. Such features include:

[1] Shrinkage cavities

[2] Solidification layers

[3] Solidification defects

[4] Casting defects

[5] Oxidation

Figure 38: Circumferential crack-counting apparatus and setup

Figure 39: Crack-counting device which converts rotation of the pin into circumferential cracking percentage based on 360°

4.2.5 Sample Preparation and Metallurgical Characterization

Pin samples were selected and carefully prepared for metallography, fractography and microscopy. Fractography samples were created from pins which exhibited 100% or near

100% circumferential cracking. These samples can be mounted in the SEM for analysis in the as-cast condition. Metallography and microscopy samples, however, were prepared differently. Cast pins were sectioned in both the longitudinal and transverse directions using a Techcut 5 precision sectioning saw. Sections were then mounted in a LECO PR-

36 mounting press with Buehler conductive resin. Once mounted, samples were polished using 240, 400, 600, and 800 grit silicon carbide paper, followed by 6, 3 and 1 µm diamond paste on a LECO Spectrum System 1000. Lastly, samples were polished using a

Buehler Vibromet™ 2 vibratory polisher with 0.5 µm diamond paste for a fine finish.

Between each step, samples were ultrasonically cleaned in water and then ethanol for several minutes to avoid cross-contamination between polishing pads.

Next, samples were electrolytically etched in 10% volume chromic acid (CrO3) solution. Samples were submerged in the solution in a stainless steel (cathode) dish where the current and voltage of the constant voltage DC power supply were preset to 0.25A and 3V respectively. To etch, a tungsten (anode) electrode was placed in direct contact onto the surface of the desired etched region for approximately 2-5 seconds depending on the size of the sample and age of the acid solution. The samples were then ultrasonically cleaned in water and ethanol.

Light optical microscopy and scanning electron microscopy were used to evaluate cast pin samples. LOM was performed using an Olympus GX51 metallograph where basic photomacrographs and micrographs were captured. Three scanning electron

microscopes were used to perform high-resolution microscopy on the sample: (1)

Quanta-200 general purpose SEM, (2) XL-30F ESEM field emission gun SEM, and (3)

FEI Sirion SEM. Backscatter electron (BSE) and secondary electron (SE) modes were utilized in all SEMs for fractographic and microscopic analyses. Energy dispersive spectroscopy was used to identify elemental dispersion and segregation as well as matrix, precipitate, crack-tip and near crack-tip compositions. Working parameters were dependent on the SEM and detector being used. Approximate working distances ranged from 5-12mm, with an accelerating voltage from 20-25 KV, and a spot size from 3-5.

4.2.6 Computational Modeling of Solidification Behavior and Eutectic Formation

The computational modeling of each heat of FM 82 used the Scheil-Gulliver module in

Thermo-Calc™ 2015a software. The TCNI8 database was used to calculate solidification temperature ranges (STR), eutectic solidification start temperatures, volume fraction eutectic, and elemental partitioning coefficients (k). Calculations were carried out with the following assumptions:

[1] Equilibrium of the solid-liquid interface

[2] Diffusion occurs in the solid (C)

[3] Complete mixing of the liquid phase

Each simulation descended from 2000°C to a temperature where 99% solid phase fraction was reached using a cooling rate of 1°C/sec. Present phases included liquid and

FCC crystal structure phases (austenite and NbC).

CHAPTER 5: RESULTS AND DISCUSSION

This chapter contains the results and discussion pertaining to:

[1] CPTT Initial Verification of FM 82 Heats and Dilution Studies

[2] Thermo-Calc™ Scheil Solidification

[3] Metallurgical Characterization of Sample Wire Heats

[4] FM 82 with Nitrogen Additions

5.1 CPTT Initial Verification of FM 82 Heats

In order to examine the reproducibility of the CPTT, a verification study of cracking responses for three heats of FM 82 was conducted to compare with results obtained previously at OSU by E. Przybylowicz [1]. The tested heats exhibited either poor or excellent solidification crack resistance behavior. These heats are identified in Table 7.

This work was done in order to further develop the validation and verification study of the third generation CPTT conducted by E. Przybylowicz [13]. These tests not only refined some of the parameters in the casting process, but also ensured excellent correlation between studies conducted at different times using the same materials.

In order to achieve accuracy and repeatability of the CPTT, the same parameters were used in the current study as the previous study, and are shown in Table 8. It is important to note that, in some cases, parameters may have been changed depending on the casting conditions.

Table 7: CPTT Behavior from Previous Study [1] Cracking Solidification Heat Threshold Pin Cracking ID Length Resistance (in) AB8573 0.75 Bad B8142 1.125 Good YB8908 1.125 Good

Table 8: Initial CPTT Parameters Set Casting Ramp Ramp Flow Pressure Temp Current Down Down Rate (psi) (°C) (A) (A) (Sec) (CFH) 1430 400 275 0.85 2.5 0.3

Although full curves were developed over the entire pin-range for all heats, a more refined focus was maintained on the lower cracking thresholds of each heat.

Evidence of scatter in the data suggests this threshold is the best indicator of susceptibility to solidification cracking. Data scatter is most often seen above the lower cracking threshold and can be better observed in the plots of Maximum/Minimum

Cracking versus Pin Length in any alloy system.

5.1.1 Heat A: High Susceptibility to Solidification Cracking

Heat A is a heat of FM 82 which is known both in-service and in previous testing, to exhibit poor solidification cracking behavior. The comparative study for this heat proved to resemble similar results as the previous study using the CPTT. The lower cracking threshold in the first study (2013) was determined to be 0.75 inches while the newer study (2015) was determined to be the same—although the most recent study exhibited just 1% cracking at the next pin length. Table 9 shows the CPTT data obtained for this

heat in both experiments. It should be noted that less pins were tested in regions of the curve that were of less interest to this study. More pins casted in those areas will provide more refined results. Figure 40 demonstrates the plot of the full range of cracked pins tested (A) and the plot of the average Circumferential Cracking versus Pin Length (B), which determines the cracking susceptibility ranking of the material. Again, it was in the interest of this research to target the lower cracking threshold rather than the entire set of pin lengths. The data further above the LCT shows much scatter, indicating that this transition region from 0% cracking to 100% cracking is not an accurate measurement of susceptibility to solidification cracking. This scatter best observed in the

Maximum/Minimum Circumferential Cracking versus Pin Length (Figure 41).

Table 9: Cracking response characteristics of Heat A Heat AB8573

Pin Length Count Min Max Ave. (in) 2013 2015 2013 2015 2013 2015 2013 2015 0.75 6 5 0% 0% 0% 0% 0% 0% 0.875 7 6 0% 0% 34% 5% 11% 1% 1 6 5 0% 0% 14% 28% 4% 10% 1.125 6 5 0% 0% 16% 27% 4% 5% 1.25 6 6 5% 12% 35% 46% 12% 27% 1.375 6 1 5% 20% 100% 60% 49% 40% 1.5 5 1 42% 100% 100% 100% 77% 100% 1.625 6 1 26% 100% 100% 100% 79% 100% 1.75 1 1 100% 100% 100% 100% 100% 100% 1.875 1 1 100% 100% 100% 100% 100% 100%

Figure 40: CPTT complete curve (A) and average circumferential cracking percentage (B) versus pin length comparing two studies of Heat A

Figure 41: CPTT maximum cracking (A) and minimum cracking percentage (B) versus pin length comparing two studies of Heat A

5.1.2 Heat Y: Low Susceptibility to Solidification Cracking

Heat Y is a heat of FM 82 which exhibits good solidification cracking resistance as assessed by the CPTT but poor solidification cracking resistance based on actual welding experience—the reason for this discrepancy is unknown. The comparative study for this heat proved to verify the same lower cracking threshold as the previous study of this heat—exactly 1.125 inches. Thus, a zero-pin deviation was established and the study-to- study correlation was an excellent match. Table 10 compares the CPTT data obtained for this heat in both experiments. Figure 42 demonstrates the full range of cracked pins tested

(A) and the average Circumferential Cracking versus Pin Length (B) which determines the cracking susceptibility ranking of the material. The Maximum/Minimum

Circumferential Cracking versus Pin Length is shown in Figure 43.

Table 10: Cracking response characteristics of Heat Y Heat YB8908 Count Min Max Ave. Pin Length (in) 2013 2015 2013 2015 2013 2015 2013 2015

0.875 2 2 0% 0% 0% 0% 0% 0% 1 2 2 0% 0% 0% 0% 0% 0% 1.125 3 2 0% 0% 0% 0% 0% 0% 1.25 6 2 0% 0% 15% 17% 4% 9% 1.375 5 7 0% 0% 65% 76% 33% 16% 1.5 4 4 0% 12% 100% 81% 43% 31% 1.625 5 4 10% 0% 93% 45% 36% 20% 1.75 5 5 0% 0% 97% 100% 51% 39% 1.875 7 2 15% 75% 100% 86% 68% 81% 2 4 1 100% 100% 100% 100% 100% 100% 2.125 2 1 100% 100% 100% 100% 100% 100% 2.375 1 1 100% 100% 100% 100% 100% 100% 2.5 2 1 100% 100% 100% 100% 100% 100%

Figure 42: CPTT complete curve (A) and average circumferential cracking percentage (B) versus pin length comparing two studies of Heat Y

Figure 43: CPTT maximum cracking (A) and minimum cracking percentage (B) versus pin length comparing two studies of Heat Y

5.1.3 Heat B: Low Susceptibility to Solidification Cracking

Heat B is a heat of FM 82 which also exhibits good solidification cracking resistance.

The comparative study for this heat also showed excellent LCT correlation with the results of the previous study using the CPTT. The lower cracking threshold in the both experiments was 1.125 inches, although the most recent study exhibited just 1% cracking at the next pin length. Similar to the other heats, a zero-pin deviation shows excellent correlation between the studies. Table 11 compares the CPTT data obtained for this heat in both experiments. Figure 44 demonstrates the full range of cracked pins tested (A) and the average Circumferential Cracking versus Pin Length (B) which determines the cracking susceptibility ranking of the material. The Maximum/Minimum Circumferential

Cracking versus Pin Length is shown in Figure 45.

Table 11: Cracking response characteristics of Heat B Heat B8142 Count Min Max Ave. Pin Length 2013 2015 2013 2015 2013 2015 2013 2015 (in) 0.875 2 2 0% 0% 0% 0% 0% 0% 1 2 5 0% 0% 0% 0% 0% 0% 1.125 5 3 0% 0% 0% 0% 0% 0% 1.25 4 4 0% 0% 2% 34% 1% 10% 1.375 5 3 0% 0% 40% 32% 17% 19% 1.5 4 2 0% 2% 100% 30% 41% 16% 1.625 3 3 37% 15% 55% 89% 44% 40% 1.75 5 2 35% 34% 100% 75% 67% 55% 1.875 3 1 40% 87% 100% 87% 80% 87% 2 3 2 55% 80% 100% 100% 79% 90% 2.125 4 1 100% 100% 100% 100% 100% 100% 2.375 2 1 100% 100% 100% 100% 100% 100% 2.5 1 1 100% 100% 100% 100% 100% 100%

Figure 44: CPTT complete curve (A) and average circumferential cracking percentage (B) versus pin length comparing two studies of Heat B

Figure 45: CPTT maximum cracking (A) and minimum cracking percentage (B) versus pin length comparing two studies of Heat B

5.1.4 CPTT Dilution Studies

Two dilution experiments were conducted to develop a proper understanding of the interaction of different heats by mixing a heat that exhibits good resistance to solidification cracking (Heat B) with a heat that exhibits poor resistance (Heat A). The two dilution mixtures were Heat B diluted with 10wt% and 25wt% of Heat A.

The 90-10wt% mixture exhibits poor solidification cracking behavior when considered in the ranking with the other heats of this study. The lower cracking threshold was identified at the 1 inch pin length. When comparing the data for the superior heat

(Heat B), the small addition of the poor-performing heat (Heat A) proved to have a detrimental effect on the solidification cracking resistance. Table 12 compares the CPTT data obtained for this dilution after crack analysis. Figure 46 plots both the full range of cracked pins tested (A) and the average Circumferential Cracking versus Pin Length (B).

Table 12: Cracking response characteristics of 10% Dilution of Heat B with Heat A 10%Dilution of B8142 with AB8573 Pin Length Count Min Max Ave. (in) 0.75 2 0% 0% 0% 0.875 5 0% 0% 0% 1 7 0% 0% 0% 1.125 2 9% 63% 36% 1.25 3 0% 100% 46% 1.375 2 30% 75% 53% 1.5 2 50% 100% 75% 1.625 2 43% 63% 53% 1.75 2 100% 100% 75% 1.875 2 100% 100% 100% 2 2 100% 100% 100% 2.125 2 100% 100% 100%

Figure 46: CPTT complete curve (A) and average circumferential cracking percentage (B) versus pin length of Heat B diluted with 10wt% Heat A

The 75-25wt% mixture exhibits even worse solidification cracking behavior, similar to that of Heat A. The lower cracking threshold was determined to be 0.75 inches.

Therefore, an inclusion of just 25wt% of a susceptible heat into a mixture exhibited similar cracking behavior to that of 100wt% of that heat. Thus, dilution of this amount will cause detrimental behavior in the form of solidification cracking. Table 13 compares the CPTT data obtained for this dilution after crack analysis. Figure 47 plots the full range of cracked pins tested (A) and the average Circumferential Cracking versus Pin

Length (B). The full cracking responses comparison between the two dilution heats can be seen in Figure 48.

Table 13: Cracking response characteristics of 25% Dilution of Heat B with Heat A 25%Dilution of B8142 with AB8573 Pin Length Count Min Max Ave. (in) 0.75 2 0% 0% 0% 0.875 5 0% 39% 13% 1 3 22% 100% 62% 1.125 3 0% 84% 28% 1.25 4 39% 100% 66% 1.375 2 25% 50% 38% 1.5 3 13% 51% 35% 1.625 2 51% 51% 51% 1.75 2 45% 53% 49% 1.875 5 31% 100% 70% 2 2 63% 100% 82% 2.125 2 100% 100% 100% 2.375 1 100% 100% 100% 2.5 1 100% 100% 100%

Figure 47: CPTT complete curve (A) and average circumferential cracking percentage (B) versus pin length of Heat B diluted with 25wt% Heat A

Figure 48: CPTT cracking response curves comparing 10wt% and 25wt% dilutions of Heat B by Heat A

After testing was completed, three different criteria of the cracking responses of each heat were observed. They were:

[1] Maximum average of 0% cracking

[2] Minimum average of 100% cracking

[3] Transition Range

Previous to this work, different heats and alloying systems were compared using three other, but similar, criteria. Those being:

[1] Last pin length to exhibit 0% cracking

[2] First pin length to exhibit 100% cracking

[3] Difference between [1] and [2]

The newer method was developed to create more realistic comparisons to determine trends between heats. Due to the human error variability that comes with counting cracks, average values should be considered more reliable than actual, single data points. Given this observation, the first pin to exhibit 100% cracking was not considered when comparing data. Variables such as pressure differences, solidification patterns, and/or human error all may have an effect on solidification cracking in CPTT specimens. With this,

Table 14 and Figure 49 represent the three new criteria for crack analysis for each heat tested (A, Y, B, 10wt% dilution, 25wt% dilution). The LCT value is the last pin length that exhibited 0% cracking. The minimum average of 100% cracking (UCT) value is the first pin length that exhibited 100% cracking across all trials. The transition range is simply the difference between the other two criteria. This transition range contains the largest deviations (or scatter) across any study.

Table 14: Cracking response criteria for each heat Transition Heat LCT UCT Region AB8573 0.75 1.5 0.625 25wt% Dilution 0.75 2.125 1.375 10wt% Dilution 1 1.875 0.875 YB8908 1.125 2 0.875 B8142 1.125 2.125 1

Figure 49: Graphical representation of cracking criteria in each heat

5.2 Thermodynamic Simulations

Thermodynamic simulations using the Scheil-Gulliver module in 2015a software were performed to model the solidification behavior of two heats of FM 82 to see if any discernable characteristics could be made between a “resistant” and “susceptible” heat.

Discussion on STR, secondary phase formation, and elemental partitioning coefficients is provided below. All calculations were made at 99% solid formation. Each simulation was made using the parameters in Table 15. It was calculated that austenite (FCC_A1#2) and

MC-type carbides (FCC_A1#1) form on solidification in both heats. Although Laves phase was included in the program script, it was unseen in all trials.

The elements chosen for the simulations include C, Cr, Fe, Mn, Nb Ni, and Ti. Al and Si were not included due to the detrimental effects they had on Scheil solidification in preliminary trials. Additionally, elements whose weight percent was approximately zero were not used to avoid the “formation” of phases that are unseen in real-world SEM analysis. Such elements include Co (0.009wt%), Cu (0.007wt%), Mo (0.005wt%), Ta

(0.002wt%), V (0.006wt%) and W (0.004wt%). Table 16 contains chemical compositions that were input into the simulations. For comparison simplicity, just one

“good” heat (B) and one “bad” heat (A) were modeled.

Simulations were performed to 99% solid to fully capture the STRs of each individual phase that forms. Visual representations of Scheil solidification calculations are represented in Figure 50. Composition profiles for primary FCC solidification (Figure

51) and secondary FCC solidification (Figure 52) are also shown. Profiles for austenite solidification in both cases show typical matrix compositions that are high in Ni and Cr,

with spikes in Nb and Mn at the end of solidification, indicating segregation to the last- to-solidify liquid. Curves plotting mass fractions of each element versus mass fraction of

MC-carbides were calculated. The graphs show that these carbides are indeed NbC that also have high contents of Ti, indicating the possible presence of (Nb,Ti)(CN) or MX- type precipitates.

Table 15: Parameters Used in Thermodynamic Simulations using Thermo-Calc 2015a Software Simulation Parameter Input Module Scheil-Gulliver Database TCNI8 Major Element Ni Alloying Elements See below Temperature 2000 C Phases Rejected All Liquid Phases Restored FCC_A1 C14_LAVES Miscibility Gap Check? FCC_A1 Cr, Fe, Mn, Major Constituents Nb, Ni, Ni Minor Constituents C Temperature Step 1 C Default Stop Point? Yes Fast Diffusing C Components

Table 16: Chemical Compositions used for Computational Modeling Chemical Composition (wt %) Heat C Cr Fe Mn Nb Ni Ti B (good) 0.027 19.68 1.05 2.84 2.37 Rem 0.36 A (bad) 0.038 19.27 1.11 2.84 2.37 Rem 0.35

Figure 50: Scheil STRs of Heats B (A) and A (B)

Figure 51: Composition profiles for primary FCC (austenite) solidification for Heats B (left) and A (right)

Figure 52: Composition profiles for secondary FCC (MC-carbide) solidification for Heats B (left) and A (right)

Additionally, the composition of the solid (CS) and the nominal input composition

(CL) were predicted at 99% solid and are shown in Table 17. The partitioning coefficients, k, for each element can be calculated by CS / CL to understand how the element behaves during solidification. Elements with partitioning coefficients less than

one indicate that they segregate to the liquid during solidification. Due to the small variances in composition between the two heats, it can be seen that k-values for each element are nearly the same. Table 18 summarizes the entirety of the results from the solidification of Heats B and A using the Scheil model.

To graphically represent the data shown in Table 18, the overall STRs are broken down into individual phases in Figure 53 . They are presented in order of their susceptibility to solidification cracking from left (resistant) to right (susceptible). The susceptibility ranking order is based on both in-service ranking as well as CPTT verification results. Since only two alloys are plotted, it is difficult to determine the relevance of trends between them. Simulations for Heat B predicted larger solidification temperature ranges in both the γ and γ/NbC eutectic phase fields, and thus a larger overall

STR as well. In general, it is understand that smaller STRs are more beneficial for solidification cracking resistance. This discrepancy can be described by pointing out that computational modeling does not take eutectic crack healing into account, which is very prevalent in FM 82, especially Heat B. Work conducted by E. Przybylowicz on Ni-based alloys (including FM 82) showed that no direct correlation was drawn between solidification temperature range or NbC STR and solidification cracking susceptibility

[13]; the same trends hold true in this experiment.

100

Table 17: CS, CL and k (CS / CL) at the End of Solidification (99% solid) for Heats B and A Heat Element B8142 AB8573

CS CL K CS CL K C 2.724E-05 2.700E-04 1.009E-01 3.903E-05 3.817E-04 1.023E-01 Cr 1.992E-01 1.968E-01 1.012E+00 1.948E-01 1.927E-01 1.011E+00 Fe 1.109E-02 1.050E-02 1.056E+00 1.173E-02 1.110E-02 1.057E+00 Mn 1.814E-02 2.840E-02 6.386E-01 1.814E-02 2.845E-02 6.377E-01 Nb 1.06E-02 2.37E-02 4.465E-01 1.06E-02 2.38E-02 4.480E-01 Ni 7.58E-01 7.36E-01 1.030E+00 7.62E-01 7.40E-01 1.030E+00 Ti 2.29E-03 3.60E-03 6.351E-01 2.24E-03 3.51E-03 6.394E-01

Table 18: Thermo-Calc Scheil Simulation Results for Heats B and A Heat: Heat B Heat A Element Resistant Susceptible

C 0.1009 0.1027 Cr 1.0124 1.0108 Fe 1.0562 1.0565 Mn 0.6386 0.6385 Nb 0.4465 0.4487

Ni 1.0297 1.0300 PartitioningCoefficient Ti 0.6351 0.6387 Liquidus (°C) 1378 1379 NbC Start (°C) 1298 1306 Laves Start (°C) - - Solidus (°C) 1162 1176 STR (°C) 216 203 NbC [Start-Solidus] (°C) 136 130 Primary FCC [Liquidus-NbC 80 73 Start] Primary FCC Fraction 0.37 0.359 Mole % Austenite 98.8015 98.72 Mole % NbC 0.1985 0.2822

101

Figure 53: Phase STRs for Heats B and A

5.3 Metallurgical Characterization of Sample Wire Heats

Metallurgical and chemical analyses were performed on pins from Heats A (susceptible) and B (resistant). Metallography was performed using the Olympus microscope at OSU while microscopy and fractography were performed using the Quanta 200 scanning electron microscope (SEM), the XL-30 environmental scanning electron microscope

(ESEM), and the Sirion SEM at the Center for Electron Microscopy and Analysis

(CEMAS). EDS scans are used to determine bulk matrix chemical composition as well as that of crack tips, secondary phase constituents, and backfilling.

5.3.1 Metallography

Metallographically, several features were taken into consideration in order to compare different heats from solely an observation standpoint. Those features were: crack

102

initiation sites, grain size and appearance, precipitation dispersion, visible liquid around crack surfaces, and eutectic backfilling and crack healing.

Initially it was necessary to verify that cracks were initiating and propagating at and along SGBs. Evidence of this in both heats is shown in Figure 54. The cracks observed in these materials all showed classic signs of solidification cracks including: location, tortuosity, and, in some cases, the presence of eutectic liquid around the crack.

Fractographic images of fracture surfaces are shown in Fractography with evidence of solidification crack dendritic morphology. Figure 55 demonstrates that grain size and precipitation dispersion differences from a metallographic view were proven to be negligible between heats and thus, deemed to have no correlation to crack susceptibility.

The only potential difference correlating secondary phase formation and grain size would be differing crack tortuosity, which does play a significant role in crack susceptibility— this, however, was not deemed to play a role.

More importantly, the presence of last-to-solidify eutectic around cracks is certainly significant in regards to backfilling and crack healing. Figure 56 demonstrates the trend seen across all trials between these two heats where much liquid surrounds the cracks in the “good” heat whereas the opposite is true in the “bad” heat. It is necessary to have sufficient liquid at the end of solidification to flow into the cracks and heal any regions where cracking occurs before solidification ends. The same trend is followed regarding crack healing of any kind as shown in Figure 57. Heat A demonstrates no observable crack backfilling, whereas there is complete crack healing in Heat B. These trends correlate well with the solidification crack susceptibility ranking and suggest backfilling contributes to resistance to solidification cracking in FM 82.

103

Figure 54: Evidence of cracking along SGBs in heats: A (A) and B (B); 1000X magnification

104

Figure 55: Grain size and secondary phase dispersion in heats: A (A) and B (B); 1000X magnification

105

Figure 56: Amount of eutectic liquid surrounding crack in heats: A (A) and B (B); 1000X magnification

106

Figure 57: No crack backfilling in Heat A (A), and complete crack healing in Heat B (B); 500X and 1000X magnification respectively

107

5.3.2 Fractography

Fracture surfaces were analyzed in each heat to compare and contrast major and minor features of Heats A and B. Analyses confirm the presence of solidification cracking which are observed as dendritic, or "egg-crate", fracture morphology. Since this filler metal undergoes a eutectic reaction, liquid coverage of the dendrites occurs and can make it difficult to observe secondary precipitation. A clear image of last-to-solidify liquid covering a dendrite is shown in the bottom left of Figure 58. Evidence of “egg-crate” morphology fracture surfaces of each heat are depicted in Figure 59. Additionally, evidence of titanium-rich precipitates of cuboidal or “blocky” morphology are shown in

Figure 60. Although not analyzed by transmission electron microscopy (TEM), these are identified as TiN in the literature.Further fractography work shows the apparent difference in precipitate morphology between heats that exhibit “good” solidification cracking behavior versus that of a heat exhibiting “poor” performance (Figure 61).

The poor performing heat, Heat A, exhibits more rod-like and cruciform-shaped precipitates when observing regions near cracks in cast pin samples. Heat B, which exhibits resistance to solidification cracking, produces precipitates of the cuboidal and spherical forms. The differences in morphologies point out that “good” and “bad” heats may be taking slightly differing solidification paths which may lead to the varying cracking responses when being exposed to harsh environments.

108

Figure 58: Liquid coverage of a dendrite in Heat A fracture surface

109

Figure 59: Fracture surfaces of Heats: A (A) and B (B) showing evidence of solidification cracking by “egg-crate” dendritic fracture surface

110

Figure 60: Titanium-rich precipitates in Heats: A (A) and B (B)

111

Figure 61: Precipitate morphologies in Heats: A (A)—rod-like; B (B)--spherical

112

5.3.3 Compositional Analysis

This section discusses the compositional analysis of the three main wire heats studied in this research. EDS point scans were used to obtain the data shown.

Bulk matrix point scans of metallographic pin samples were taken in near-crack regions for verification of correlation between the EDS capability of the SEM and measurements obtained by an outside vendor. The comparison between the ESEM EDS and vendor compositions were deemed acceptable despite minor discrepancies with Nb and Si in particular. These deviations may stem from the particular areas chosen to conduct EDS on depending on the spot size used and the available counts registered in the EDS system. Average compositions of three bulk matrix point scans of each heat are shown in Table 19. Elements that were discarded for excessive error % were: C, Mo, P,

S, Cu, Co, Pb, V, Ta, and W. Vendor matrix compositions for each heat are also displayed in Table 19. In terms of solidification cracking ranking, no notable compositional data was taken from these analyses. Further compositional work using transmission electron microscopy must be done to find any notable discrepancies that may lead to an understanding of differing cracking susceptibilities among heats.

Although not pictured, chemical analyses on backfilled regions and crack tips showed higher-than-nominal concentrations of both Nb and Mn along with lower-than- nominal concentrations of Ni and Cr in each heat. This is not surprising due to the partitioning coefficients of both Nb and Mn indicating they segregate strongly to the liquid, which is present in the form of eutectic at the end of solidification along SGBs.

113

Table 19: EDS/Vendor Chemical Composition Comparison SEM EDS Matrix Point Scan Averages Heat Al Cr Fe Mn Nb Ni Si Ti Heat A 0.23 19.50 0.94 3.09 1.86 73.71 0.33 0.33 Heat Y 0.21 19.54 1.28 3.30 1.72 73.10 0.23 0.63 Heat B 0.20 19.40 1.40 3.20 1.62 73.50 0.16 0.54 BWXT Vendor Composition Averages Heat Al Cr Fe Mn Nb Ni Si Ti Heat A 0.260 19.270 1.110 2.840 2.370 73.800 0.170 0.350 Heat Y 0.205 19.520 1.140 2.815 2.470 73.400 0.160 0.370 Heat B 0.510 19.680 1.050 2.840 2.373 73.500 0.096 0.360

5.4 Effects of Nitrogen Additions

Weld mockups were made with intentional nitrogen additions in the argon shielding gas to observe how nitrogen affects the cracking susceptibility of FM 82. A heat of FM 82 known for its resistance to solidification cracking both in-service, as well as in the CPTT, was used for testing and analysis. The chemical composition of Heat YT0159 (Heat YT) is shown in Table 20.

Table 21 shows the nitrogen addition weld samples that were used for testing.

Samples were later sectioned, polished, melted, tested using the CPTT, and prepared for microscopy. The CPTT cracking response curves for each individual gas mixture can be found in the Appendix. Figure 83 represents the plot of the full range of cracked pins tested (top) and the plot of the average Circumferential Cracking versus Pin Length

(bottom) in Heat YT with no N additions.

114

Table 20: Heat YT Chemical Composition Heat YT0159 Al C Co Cr Cu Fe H Mg Mn Mo 0.055 0.055 0.004 20.39 0.004 1.14 0.0008 <.001 3.25 0.003 N Nb Ni O P S Si Ti V W 0.0016 2.29 72.3 0.0011 <.005 <0.001 0.11 0.42 0.007 <.001

Table 21: Nitrogen Addition Weld Sample Gas Mixtures and Content Nitrogen Weld Sample Content, LECO (ppm) 100% Argon ~19

99.8% Ar / 0.2% N2 ~140 99.6% Ar / 0.4% N2 ~200 99% Ar / 1% N2 ~600

95% Ar / 5% N2 ~1300

5.4.1 Heat YT: 140 ppm N

The first group of nitrogen-bearing samples was welded with 0.2 vol% N2 in the

shielding gas, which translates to approximately 140 ppm N in the weld metal. It was

immediately observed that any small increase in N content in the weld metal increased

the cracking susceptibility of this alloy. The cracking threshold associated with this N

content was reduced from the original LCT of Heat YT (1.375 in) to 0.75 in—a 5 pin-

length reduction. Figure 83 represents the plot of the full range of cracked pins tested

(Top) and the plot of the average Circumferential Cracking versus Pin Length (Bottom);

the latter determines the cracking susceptibility ranking of the material.

In the case of these samples, precipitates were observed using SEM/EDS

analyses. Compared the original heat, an increase in niobium and titanium-rich particles

taking on irregular morphologies are observed to decorate the matrix both intergranularly

115

and intragranularly. Examples of such precipitates are shown in Figure 62-Figure 64. It is well-documented for this material that Ti-rich nitrides form at very high temperatures in the liquid and niobium-rich carbides form at lower temperatures as the product of a γ/

NbC eutectic reaction. Lastly, many “micro” solidification cracks were observed throughout this particular sample. Evidence of this is shown in Figure 65.

116

Figure 62: (Nb,Ti)-rich particle observed in sample containing 140 ppm N

117

Figure 63: NbC of irregular morphology

118

Figure 64: Irregular NbC

Figure 65: Evidence of "micro" cracks in 140 ppm N sample

119

5.4.2 Heat YT: 200 ppm N

The second group of nitrogen-bearing samples was welded with 0.4 vol% N2 in the shielding gas, which translates to approximately 200 ppm N in the weld metal. This additional increase in N content in the weld metal caused a further reduction in the LCT to 0.625 in—a 6 pin-length reduction from the original resistant heat. Figure 85 demonstrates the plot of the full range of cracked pins tested (top) and the plot of the average Circumferential Cracking versus Pin Length (bottom).

In the case of these samples, similar observations were made compared with the

140 ppm N sample. As expected, an increase the in weld metal N content caused an increase in niobium-rich (Figure 66) and titanium-rich particles. Along with the same

“irregular”-shaped precipitates, there was a clear transition of the Ti-rich particles to more cuboidal solidification morphology. Examples of such precipitates are shown in

Figure 67, Figure 68, and Figure 69. Additionally, in Figure 67 and Figure 68, it is apparent the nucleation or continual growth of Nb-rich precipitates from the Ti-rich precipitates.

120

Figure 66: Nb-rich particles decorating the matrix of a 200 ppm N sample

121

Figure 67: Evidence of Nb-rich precipitation or continual growth from existing Ti-rich constituent leading into crack surface

122

Figure 68: (Ti, Nb)-rich and Nb-rich particles

123

Figure 69: (Nb,Ti)-rich particle

124

5.4.3 Heat YT: 600 ppm N

The third group of nitrogen-bearing samples was welded with 1 vol% N2 in the shielding gas, which translates to a considerably high amount of N in the weld metal— approximately 600 ppm. This large increase in N further reduced the LCT to below the lowest testable pin length of 0.5 in—a7 pin-length deviation from the original resistant heat. Figure 86 demonstrates the plot of the full range of cracked pins tested (top) and the plot of the average Circumferential Cracking versus Pin Length (bottom).

In the case of samples sectioned from this high-N mockup, clear differences are seen compared to the former heats with lower nitrogen content regarding precipitate presence and morphology. As expected, an increase in nitrogen in the weld metal resulted in a higher amount of both Ti- and Nb-rich precipitation. Ti-enriched particles, however, were only typically observed as clear, cuboidal morphology and with large amount of Nb either surrounding or within the particle. Examples of Ti and Nb-rich particles are shown in Figure 70 and Figure 71.

125

Figure 70: Cuboidal (Nb,Ti)-enriched precipitates in 600 ppm N sample

126

Figure 71: Intergranular Nb-rich precipitate

127

5.4.4 Heat YT: 1300 ppm

In this group, cuboidal and triangular titanium-rich particles are present and are observed to heavily decorate the matrix both intergranularly and intragranularly. Examples of such particles are evident in Figure 72. Long segmented Nb-rich constituents are also observed, especially at GBs. These carbides form at the very end of solidification as a eutectic constituent in GBs where the last to solidify liquid is present.

Being welded in 5 vol%N2, a large amount of N is transferred to the weld metal via the arc and thus forming a high amount of secondary phase enriched in Ti. In the case of such high amounts of N, there is a presence of large, agglomerated titanium-rich

(assumed to be nitrides) that are found in groups around the matrix. Figure 74 and Figure

75 show that nitrides on the order of 2-6 µm come together and coalesce in groups. It was observed that in these large groups, porosity is often present in the center of the nitrides.

This can be explained perhaps by the fact that with such high N2, the system may be approaching its solubility limit for N. Thus, welding with such high N contents is difficult, causing much spatter as the N tends to want to escape the weld metal. This behavior results in porosity, which is also observed in Figure 74 and Figure 75. Figure 81 also shows the dendritic morphology of shrinkage porosity observed the cast pin sample.

Using SEM and EDS methods, it was observed that niobium-rich constituents do, in fact, heterogeneously nucleate or continuously grow off TiN. Figure 76, Figure 77,

Figure 78, and Figure 79 demonstrate this precipitation transition with EDS mapping. In all cases the square-shaped Ti-rich cores are surrounded with Nb-rich precipitation in the form of a “swirl”, “tail”, or “blocky” geometry. Thus, it is possible that these particles originate as TiN with Ti-enriched cores, and later nucleate as NbC where the surrounding

128

edges are enriched in Nb. Figure 80 even displays a distinct interface separating the two particles. Similar observations were made in work done by Ramirez et al. where cubic precipitates observed in the microstructure of A625 consisted of a core with TiN surrounded by a niobium titanium carbide (Nb,Ti)C [40].

Figure 72: Intragranular and intergranular titanium-rich precipitation

129

Figure 73: EDS of intergranular Nb-rich carbides

A B

Figure 74: Large, intragranular, agglomerated titanium-rich particles surrounding a pore

130

Figure 75: Evidence of agglomeration of large titanium nitrides

TiN

NbC

NN Figure 76: Nb-rich precipitation and/or continual growth from existing TiN

131

Figure 77: (2) Nb-rich precipitation and/or continual growth from existing TiN

132

Figure 78: (3) Nb-rich precipitation and/or continual growth from existing TiN

133

Figure 79: (4) Nb-rich precipitation and/or continual growth from existing TiN

134

Figure 80: Interface between Ti-rich precipitate and NbC

Figure 81: Dendritic morphology of shrinkage porosity

135

5.4.5 Summary of Nitrogen Additions

To summarize, the effect on solidification cracking susceptibility of intentional nitrogen additions in gaseous form during welding is clear; the reason for why, however, is not. To support what has been seen in-service, it was experimentally shown that an increase in cracking susceptibility results from increasing nitrogen content in the weld metal. CPTT results for this study are displayed in

Table 22.

Figure 82 plots the lower cracking thresholds of Heat YT0159 and the four heats with additional nitrogen tested using the CPTT. In a former experiment at OSU, it was determined that Heat YT exhibited a resistant LCT of 1.375”. With just 140 ppm N in the weld metal, the LCT decreased 5 pin lengths to 0.75”. With 200 ppm N, the LCT was reduced an additional pin length to 0.625”. Lastly, in high-N groups of 600 and 1300 ppm

N, cracking was observed at the lowest possible tested pin length—an actual threshold could not be determined.

The results of this experiment have many implications not only on testing of interstitials in Ni-based alloys, but potentially in other systems as well. Much more needs to be done to characterize these samples and develop a fundamental understanding for the decreased resistance to cracking, but the experimental data shows promise in solving the overall issue at hand. Perhaps there is an interaction between interstitials at hand (i.e

N+C), and those combinations need to be tested to find a more steady correlation from heat-to-heat. As of now, it is unknown whether NbC nucleate from the TiN or whether they evolve from continual growth. Further studies must be done to understand this solidification phenomenon as well. This could be achieved by testing filler metals with

136

higher C content to see what trends, if any, is produced regarding solidification cracking susceptibility.

137

Table 22: CPTT Results for Nitrogen Addition Study Pin Length (in)

0.5 0.625 0.75 0.875 1 1.125 1.25 1.375 1.5 1.625 1.75 1.875 2

0 0% 61% 87% 80% 100% 97%

140 0% 7% 5% 20% 17% 43% 49%

200 0% 7% 2% 4% 11% 13% 25% (ppm)

600 4% 6% 11% 6% 6% 38% trogen Content Content trogen

Ni 1300 2% 5% 8% 5% 22% 64%

138

Figure 82: CPTT results for the nitrogen addition study

Table 23: Nitrogen Study Summary Table Nitrogen Weld LCT Precipitate Content Sample (in) Morphology (ppm) Base ~19 1.375 irregular Base + 0.2N ~140 0.75 irregular Base + 0.4N ~200 0.625 irregular Base + 1N ~600 <0.5 cuboidal Base + 5N ~1300 <0.5 cuboidal

139

CHAPTER 6: CONCLUSIONS

The conclusions of this research can be divided into two sections to represent the respective phases of the project. Conclusions have been drawn regarding both:

[1] Optimization and reliability of the CPTT, and

[2] Solidification cracking performance and optimization of FM 82

6.1 CPTT

1. The third generation CPTT developed at OSU is capable of successfully ranking an

alloy’s (or heat’s) susceptibility to solidification cracking resulting from small

variations in chemical composition

2. Low standard deviations of cracking percentage at the Lower Cracking Threshold

values are a good indicator of ranking solidification cracking susceptibility. Other

high standard deviation pin lengths such as with the Upper Cracking Threshold are

not as good for ranking susceptibility

3. Average Circumferential Cracking versus Pin Length is best criteria for assessing

cracking response. Maximum and minimum cracking curves show too much scatter,

especially at higher pin lengths

4. The CPTT is a reproducible testing method for Ni-based alloys. 0-pin deviations were

found at the cracking thresholds when retesting three heats of FM 82. Additionally,

140

common cracking performance trends are obvious looking at each heat’s Average

Circumferential Cracking versus Pin Length curves

6.2 FM 82 Cracking Performance

5. The mode of failure for each pin analyzed was confirmed to be solidification cracking

based on classic dendritic fracture morphology

6. CPTT results determined the following ranking of solidification cracking resistance

from left to right of three heats of FM 82 and two dilution heats:

A = 25%A-75%B< 10%A-90%B < Y = B

7. Heat-to-heat cracking response variations may be related to eutectic crack healing

8. SEM EDS did not declare many notable composition variances between “good” and

“bad” heats

9. SEM fractography studies showed clear difference in precipitate morphology among

“good” and “bad” heats. Heat B exhibited Nb-rich constituents of spherical and

cuboidal morphologies while Heat A exhibited those of rod-like morphology

10. Intentional nitrogen additions introduced through the shielding gas increased the

susceptibility to solidification cracking in an otherwise resistant FM 82 heat

11. An increase in nitrogen content in the weld metal lead to an increase in observed

intergranular and intragranular secondary phase formation in the form of TiN and

(Nb,Ti)(CN) (MX-type)

12. An increase in nitrogen content in the weld metal led to an increase in observed

porosity

141

13. Based on CPT testing and weld mockups, an increase in nitrogen content in the weld

metal lead to a decrease in surface cleanliness in cast pins

142

CHAPTER 7: RECOMMNEDATIONS FOR FUTURE WORK

This chapter encompasses the author’s suggestions for future work to aid this project as well as future research. Recommendations are made for both the Cast Pin Tear Test method and the metallurgical evaluation of FM 82.

7.1 CPTT

1. Verification of weight loss during button melting process

Button-melting is a time-saving and practical way to simulate weld solidification behavior of steels, Ni-based alloys, aluminum alloys and other materials. On the other hand, it is necessary to create a study that observes the macro and micro effects of this process.

From a macro perspective, bulk weight loss of the material can affect the CPTT results where it is essential to maintain a consistent procedure for comparative reasoning.

It needs to be addressed whether button mass is as important as the whether or not the material fills the mold completely. If it is not, a standard deviation must be determined for each pin length that is acceptable to create reproducible results. Any fluctuation in mass may affect the cracking susceptibility when subjected to a certain strain level, and thus, producing skewed results.

Next, micro effects of melting can influence different characteristics of solidification such as elemental segregation and formation of secondary phases (eutectic reactions)—

143

both of which affect the cracking susceptibility of an alloy system. Many studies focus on the addition or variation of an element to observe the how the system reacts upon welding. For elements in small amounts, variation studies heavily rely on the accuracy of measurement. It is necessary to conduct pre/post experiments using energy dispersive spectroscopy to analyze the loss of alloying or impurity elements due to melting effects on vaporization, oxidation, etc.—which is especially important for interstitials and high vapor pressure elements . A study must be performed to observe the effects of one and two-stage melting on these elements.

2. Verification of appropriate mixing and optimization of time to complete mixing

in third generation CPTT

To best simulate the fusion welding process, complete mixing of the liquid phase upon melting is critical to allow for proper segregation of alloying elements and to accomplish a homogenous composition across the microstructure. If complete mixing could be achieved without the button melting process, this would be advantageous since minimal evaporation would occur in each specimen. Additionally, it would eliminate one step from the preparation process. With this thought, a different coil design may be possible to melt and cast material that does not enter the chamber as a button shape.

Furthermore, time to achieve complete mixing varies depending on the size of the button. It may be more advantageous to make the mixing process a time-dependent one.

Although the optical pyrometer is used to read a set temperature and communicate to the

DAQ system once that temperature has been reached, manual casting is often the default option if the alloy evaporates material that covers the pyrometer lens. A method to protect the lens should be proposed so automatic casting can happen more consistently.

144

3. Test the effect of internal shrinkage voids in pins

Sometimes, upon sectioning a sample, shrinkage voids running the length of the pin are noticed. Although this does not occur in most pins, in may be a cause for concern if they go unnoticed because they could be a cause for cracking, resulting in false data. An optimization of parameters to avoid such voids should be implanted to discount any concern for these defects which are unobservable without destructive testing.

4. Modeling of casting defects

The importance of avoiding shrinkage porosities and other casting defects can be examined using modeling techniques. FEA models can be created to observe crack propagation behaviors in different locations on a pin upon solidification with changing temperatures. Perhaps defects are more/less important in crack resistant alloys than susceptible ones.

5. Retainer assembly combinations (spacers, springs, copper stops) should be

standardized for each pin length

If standard retainer assemblies were created, it would ensure consistent spring pressure for each sample within a pin set. It is important to keep this pressure constant because it affects the amount of material that is left in the mold after casting. Too little a pressure and flashing will occur, yielding less material in the mold.

145

7.2 Metallurgical Evaluation of FM 82

6. Electron microprobe for analyzing segregation. Equipped with wavelength

dispersive spectroscopy to determine precipitate composition

For this research where composition accuracy is so critical, this technique aids in determining nitride and carbide composition. It would be interesting and relevant to this work to understand the composition and morphologies of the CN that form in FM 82 to better understand if one solidification path is more beneficial for crack resistance.

7. Viscosity effects on eutectic backfilling

The fluid characteristics of the last-to-solidify eutectic play an important role in eutectic backfilling and therefore crack healing. The characteristics that most likely affect this phenomenon that affect crack susceptibility in a major way especially in Ni-based alloys are viscosity and wettability. A suggested experiment to conduct would be observing the liquid motion through a ceramic tube held at constant temperature using samples with increased weld metal N and C content, using gravity as the force of motion. It is predicted that with increased N and C, increased secondary phase will form and inhibit the fluid flow. With increased viscosity, backfilling and complete crack healing would be less likely to occur, decreasing the solidification crack resistance.

146

REFERENCES

[1] E. Przybylowicz, Unpublished research performed at the Ohio State University,

2014.

[2] S. McCracken and R. Smith, "Overlays on Cast Austenitic Stainless Steel Base

Materials," in Hot Cracking Phenomena in Welds III, J. C. Lippold, T. Bӧllinghaus

and C. E. Cross, Eds., Springer Berlin Heidelburg, 2011, pp. 333-352.

[3] Electric Power Research Institute, "Materials Reliability Program: Technical Basis

for Preemptive Weld Overlays for Alloy 82/182 Butt Welds in PWRs (MRP-169),"

EPRI, Palo Alto, CA, 2005.

[4] J. C. Lippold, W. A. Clark and M. Tumuluru, "An Investigation of Weld Metal

Interfaces," The Metal Science of Joining, pp. 141-146, 1992.

[5] J. N. Dupont, J. C. Lippold and S. D. Kiser, Welding Metallurgy and Weldability of

Nickel-base Alloys, Hoboken: John Wiley and Sons Inc., 2009.

[6] J. C. Lippold, Welding Metallurgy and Weldability, Hoboken: John Wiley & Sons

Inc., 2015.

[7] D. A. Porter, K. E. Easterling and M. Y. Sherif, Phase Transformations in Metals

and Alloys, 3rd ed., Boca Raton, FL: CRC Press, 2009.

[8] S. A. David, S. S. Babu and J. M. Vitek, "Welding: Solidification and

147

Microstructure," JOM, vol. 55, no. 6, pp. 14-20, 2003.

[9] S. Kou, "Solidification and Liquation Cracking Issues in Welding," The Journal of

the Minerals, Metals, and Materials Society, vol. 55, no. 6, pp. 37-42, 2003.

[10] Bochvar, A.A.S and Z.A, "Failures in Castings Caused by Shrinkage Stresses During

Crystallization," Izvstiva Akademik Nauk SSSR, vol. 3, pp. 349-355, 1947.

[11] W. I. Pumphrey and P. H. Jennings, "A Consideration of the Nature of Brittleness at

Temperatures above the Solidus in Castings and Welds in Aluminum Alloys," J.

Inst. Metals, vol. 75, pp. 235-257, 1948.

[12] W. S. Pellini, "Strain Theory of Hot Tearing," The Foundry, vol. 80, no. November,

pp. 125-133, 1952.

[13] E. Przybylowicz, Master's Thesis: Weldability Evaluation of High-Cr Ni-base Filler

Metals using the Cast Pin Tear Test, Columbus, OH, 2015.

[14] T. Luskin, Master's Thesis: Investigation of Weldability in High-Cr Ni-base Filler

Metals, Columbus, OH, 2013.

[15] J. C. Borland, "Generalized Theory of Super-Solidus Cracking in Welds (and

Castings)," British Welding Journal, no. August, pp. 508-512, 1960.

[16] F. Matsuda, H. Nakagawa and K. Sorada, "Dynamic Observation of Solidification

and Solidification Cracking during Welding with Optical Microscope (I)," Trans

JWRI, vol. 2, no. September, pp. 67-77, 1982.

[17] F. Matsuda, H. Nakagawa, S. Ogata and S. Katayama, "Fractographic Investigation

on Solidifiaction Cracks in the Varestraint Test of Fully Austenitic Stainless Steel,"

148

Trans JWRI, vol. 7, no. 1, pp. 59-70, 1978.

[18] N. N. Prokhorov and N. N. Prokhorov, "The Technological Strength of Metals while

Crystallising During Welding," Transaction of the Japan Welding Society, vol. 2, no.

2, pp. 109-117, 1971.

[19] F. C. Hull, "Cast-Pin Tear Test for Susceptibility to Hot Cracking," Welding

Journal, vol. 38, pp. 176-179, 1959.

[20] H. Herold, A. Hübner and M. Zinke, "Investigations on the Use of Nitrogen

Shielding Gas in Welding and its Influence on the Hot Crack Behaviour of High-

Temperature Resistant Fully Austenitic Ni-and Fe-base Alloys," Welding in the

World, vol. 49, pp. 50-63, 2005.

[21] A. J. Ramirez, J. W. Sowards and J. C. Lippold, "Improving the ductility-dip

cracking resitance of Ni-base alloys," Journal of Materials Processing Technology,

vol. 179, pp. 212-218, 2006.

[22] J. Unfried-Silgado and A. J. Ramirez, "Modeling and characterization of as-welded

microstructure of solid solution strengthened Ni-Cr-Fe alloys resistant to ductility-

dip cracking Part II: Microstructure characterization," Met. Mater. Int., vol. 20, no.

2, pp. 307-315, 2014.

[23] K. H. Tseng and C. P. Chou, "The study of nitrogen in argon gas on the angular

distortion of austenitic stainless steel weldments," Journal of Materials Processing

Technology , vol. 142, pp. 139-144, 2003.

[24] A. A. Kodentsov, J. H. Gülpen, C. Cserháti, J. K. Kivilahti and F. J. Van Loo,

149

"High-Temperature Nitridation of Ni-Cr Alloys," Metallurgical and Materials

Transactions A, vol. 27A, pp. 59-69, 1996.

[25] C. Kowanda and M. O. Speidel, "Solubility of nitrogen in liquid nickel and binary

Ni-X alloys (X= Cr, Mo, W, Mn, Fe, Co) under elevated pressure," Scripta

Materialista, vol. 48, pp. 1073-1078, 2003.

[26] A. B. Murphy, M. Tanaka, S. Tashiro and J. J. Lowke, "A computational

investigation of the effectiveness of different shielding gas mixtures for arc

welding," Journal of Physic D: Applied Physics, vol. 42, pp. 1-13, 2009.

[27] H.-Y. Huang, "Effects of shielding gas composition and activating flux on GTAW

weldments," Materials and Design, vol. 30, pp. 2404-2409, 2009.

[28] T. Ogawa, K. Suzuki and T. Zaizen, "The Weldability of Nitrogen-Containing

Austenitic Stainless Steel: Part II--Porosity, Cracking and Creep Properties,"

Welding Research Supplement, pp. 213s-223s, 1984.

[29] B. Nabavi, M. Goodarzi and V. Amani, "Nitrogen Effect on the Microstructure and

Mechanical Properties of Nickel Alloys," Welding Journal, vol. 94, pp. 53s-60s,

2015.

[30] C. D. Penna, "Development of New Nitrided Nickel-Based Alloys for High

Temperature Applications," in Superalloys 2000, Zurich, Switzerland, 2000.

[31] X. Huang, Y. Zhang and Z. Hu, "Effect of Small Amounts of Nitrogen on Properties

of a Ni-Based Superalloy," Metallurgical and Materials Transactions A, vol. 30A,

pp. 1755-1761, 1999.

150

[32] E. R. Cutler, A. J. Wasson and G. E. Fuchs, "Effect of minor alloying additions on

the solidification of single-crystal Ni-base superalloys," Journal of Crystal Growth,

vol. 311, pp. 3753-3760, 2009.

[33] X. Huang, Y. Zhang, Y. Liu and Z. Hu, "Effect of small amount of nitrogen on

carbide characteristics in unidirectional Ni-base superalloys," Metall. Mater. Trans.

A, vol. 28A, pp. 2143-2147, 1997.

[34] K. L. Zieslel-Mashl and B. J. Pletka, "Segregation during solidification,"

Superalloys, pp. 175-184, 1992.

[35] H. Wada and R. D. Pehlke, "Nitrogen Solution and Titanium Nitride Precipitation in

Liquid Fe-Cr-Ni Alloys," Metall. Trans. B, vol. 28B, pp. 675-682, 1997.

[36] U. Krupp and H. J. Christ, "Internal Nitridation of Nickel-Base Alloys. Part I.

Behavior of Binary and Ternary Alloys of the Ni-Cr-Al-Ti System," Oxidation of

Metals, vol. 52, pp. 277-298, 1999.

[37] U. Krupp and H. J. Christ, "Internal Nitridation of Nickel-Base Alloys. Part II.

Behavior of Quaternary Ni-Cr-Al-Ti Alloys and Computer-Based Description,"

Oxidation of Metals, vol. 52, pp. 299-320, 1999.

[38] H. J. Christ, S. Y. Chang and U. Krupp, "Thermodynamic characterisitcs and

numerical modeling of internal nitridation of nickel base alloys," Materials and

Corrosion, vol. 54, no. 11, pp. 887-894, 2003.

[39] Special Metals, "INCONEL Filler Metal 82," [Online]. Available:

http://www.specialmetals.com/assets/smw_docs/fm82.pdf. [Accessed 12 November

151

2015].

[40] H. Cordeiro de Miranda, M. Ferreira Motta, J. Pereira Farias, C. Ramos Moreira

Afonso, A. Jose Ramirez and C. Carvalho Silva, "New insight on the solidification

path of an alloy 625 weld overlay," Journal of Materials Research and Technology,

vol. 2, no. 3, pp. 228-237, 2013.

152

APPENDIX: Nitrogen Study CPTT Results

153

Figure 83: CPTT complete curve (A) and average circumferential cracking percentage (B) versus pin length of Heat YT with 19 ppm N [12]

154

Figure 84: Figure 68: CPTT complete curve (A) and average circumferential cracking percentage (B) versus pin length of Heat YT with 140 ppm N

155

Figure 85: CPTT complete curve (A) and average circumferential cracking percentage (B) versus pin length of Heat YT with 200 ppm N

156

Figure 86: CPTT complete curve (A) and average circumferential cracking percentage (B) versus pin length of Heat YT with 600 ppm N

157

Figure 87: CPTT complete curve (A) and average circumferential cracking percentage (B) versus pin length of Heat YT with 1300 ppm N

158