Weldability Evaluation of High-Cr Ni-Base Filler Metals using the Cast Pin Tear. Test THESIS

Transcription

1 Weldability Evaluation of High-Cr Ni-Base Filler Metals using the Cast Pin Tear Test THESIS Presented in Partial Fulfillment of the Requirements for the Degree Master of Science in the Graduate School of The Ohio State University By Eric Przybylowicz Graduate Program in Welding Engineering The Ohio State University 2015 Master's Examination Committee: Boian Alexandrov, Advisor John Lippold

2 Copyright by Eric Thomas Przybylowicz 2015

3 Abstract High chromium, nickel-base filler metals have been commonly used throughout the nuclear power industry for the weld overlay repair of dissimilar metal welds (DMW) affected by primary water stress corrosion cracking (PWSCC). These alloys provide optimum resistance to PWSCC in nuclear power plant cooling systems. However some of these nickel alloys present weldability challenges including susceptibility to solidification cracking and ductility dip cracking. There is a current need to incorporate the evaluation of weldability into the alloy development process. The Cast Pin Tear Test may provide a viable means for evaluating alloy susceptibility to solidification cracking in a timely and economical manner. The first objective of this study was to optimize the new generation CPTT procedure in order to improve the reproducibility and reliability of the test. The second objective was to generate solidification cracking susceptibility rankings using the CPTT in a series of high-cr Ni-base filler metals: ERNiCrFe-7A (52M), ERNiCrFe-13 (52MSS), and ERNiCr-3 (82) filler metals, including two heats of 52M, two heats of 52MSS, and one heat of filler metal 82 and 690Nb. The effect of dilution with cast stainless steel substrate on solidification cracking was then investigated in alloys 52MSS-E and 52i-B at levels of 10 and 40 dilution. ii

4 An optimized testing procedure was developed for the CPTT. Procedural improvements in mold and sample cleanliness, the purge procedure, and the casting procedure have resulted in improved testing reproducibility. A reproducibility study was conducted on the CPTT using alloy 52M (ERNiCrFe-7A). The new generation CPTT was capable of successfully casting inch diameter pins ranging from inches in length and grams in mass. From the reproducibility study low standard deviation at pin lengths by which alloy susceptibility to solidification cracking are ranked (max. pin length with 0 cracking, min. pin length with 100 cracking) indicates good test reproducibility. The generation of solidification cracking susceptibility curves by visual evaluation of circumferential cracking produced the following ranking of the tested filler metals, from least to most susceptible: 82(A), 52i-B [187775], 52MSS-E [HV1500] and 52M [NX7206TK] and TG-SN690Nb [FBH2280], 52M [NX0T85TK], 52MSS-C [NX77W3UK]. The rankings produced using the new generation CPTT are in agreement with those obtained using the previous generation: Filler metal 82; least susceptible to solidification cracking followed by filler metals 52M, and 52MSS-C. A new procedure was developed for more reliable evaluation of cracking response in cast pin samples. Solidification cracking susceptibility curves were developed through tensile testing of CPTT samples. A new ranking criterion, UTS threshold reduction, was found to produce results similar to those determined by visual evaluation of circumferential cracking. The ranking of the tested filler metals by tensile inspection, from least to most susceptible to solidification cracking, is: 52MSS-E [HV1500], 52i-B iii

5 [187775], TG-SN690Nb [FBH2280], 52M [NX7206TK], 52MSS-C [NX77W3UK]. The difference between inspection method rankings lies in the more resistant ranking of alloy 52MSS-E which was determined by tensile inspection. It was concluded that tensile inspection offers a more objective and comprehensive evaluation of cracking in cast pin samples due to the inability of visual circumferential cracking measurement to account for crack depth. Dilution of alloys 52MSS-E and 52i-B at 10 CF8A resulted in a slight increase in solidification cracking susceptibility, as indicated by CPTT results. A dilution level of 40 resulted in a dramatic increase in cracking susceptibility in both alloys. Thermodynamic simulation results suggest that as the dilution level increases in alloys 52i-B and 52MSS-E from 0 to 40 the partitioning coefficient of Nb decreases due to an elevated nominal iron content. This increase in Nb partitioning results in increased eutectic formation in the form of laves phase. Among all the alloys studied it was determined that a larger amount of eutectic promotes solidification cracking resistance in the presence of a sufficiently high NbC/laves ratio. Metallurgical characterization using light optical microscopy, scanning electron microscopy (SEM), and energy dispersive spectroscopy (EDS) in the SEM was performed in order to identify solidification cracking mechanisms, and to study the effect of liquid film formation and backfilling on cracking susceptibility. Nb-rich eutectic type constituents were identified at crack tips in each of the studies alloys. Eutectic backfilling was observed most abundantly in alloys 52MSS-C and 52MSS-E by a Mo and Nb-rich eutectic constituent. iv

6 Dedication This document is dedicated to my family. v

7 Acknowledgments I would like to acknowledge Dr. Boian Alexandrov and Dr. John Lippold, my advisors, for supporting and guiding my research here at The Ohio State University. Tim Luskin for the work he did in creating this new generation cast pin tear test, providing me with a basis for my research. I would also like to thank Geoff Taber and Ken Copley for their help and support in conducting my research and troubleshooting testing equipment with me. I would also like to acknowledge the Electric Power Research institute for sponsoring this research. Specifically I would like to thank Steve McCracken for his time and direction over the last two years of my research. vi

8 Vita June R. L. Thomas High School B.S. Welding Engineering, The Ohio State University M.S. Welding Engineering, The Ohio State University Field of Study Major Field: Welding Engineering vii

9 Table of Contents Abstract... ii Dedication... v Acknowledgments... vi Vita... vii Table of Contents... viii List of Tables... xi List of Figures... xiii Chapter 1: Introduction... 1 Chapter 2: Background & Objectives... 6 Welding Metallurgy of Ni-Base Alloys... 6 Weldability... 8 Ductility-Dip Cracking... 8 Solidification Cracking Weldability Testing The Varestraint Test The Transvarestraint Test viii

10 The Cast Pin Tear Test The Threaded Copper Mold Test The Grooved Copper Mold Test The Cast Pin Tear Test (OSU) Summary Objectives Chapter 3: Materials & Experimental Procedure Introduction Materials Weldability Testing Procedure Material Preparation Button Melting Cast Pin Tear Testing Visual Analysis of Cracking Response in CPTT Samples Metallurgical Characterization Computational Modeling of Solidification Chapter 4: Results & Discussion Reproducibility of the CPTT Results Weldability Testing ix

11 CPTT Visual Crack Evaluation CPTT Tensile Testing Evaluation Thermodynamic Simulation Modeling Filler Metal Simulation Results Dilution Simulation Results Metallurgical Characterization Fractography Metallography Compositional Analysis Chapter 5: Summary and Conclusions Chapter 6: Recommendations for Future Work References x

12 List of Tables Table 1: Grooved copper mold crack evaluation weighing factors Table 2: Second generation CPTT procedure (39) Table 3: Material compositions (wt) Table 4: Mass requirements for Ni-based alloys at each pin length Table 5: Parameters for the button melting apparatus Table 6: CPTT results for 52M [NX7206TK] Table 7: Visual crack examination results for CPTT Ni-based weld filler metals Table 8: CPTT visual inspection results of circumferential cracking (52i-B dilution study) Table 9: CPTT visual inspection results of circumferential cracking (52MSS-E dilution study) Table 10: CPTT tensile inspection results for 52i-B [187775] Table 11: CPTT tensile inspection results Table 12: Comparison of solidification cracking susceptibility rankings determined using tensile inspection and visual inspection Table 13: CPTT tensile inspection results (52i-B dilution study) Table 14: CPTT tensile inspection results (52MSS-E dilution study) xi

13 Table 15: Thermo-CALC (Schiel-Gulliver) simulation results for undiluted filler metals. Elements that partition during solidification (k<1) are highlighted in red Table 16: Thermo-Calc simulation results comparing eutectic components in laves forming alloys Table 17: Thermo-CALC (Schiel-Gulliver) simulation results for 52i-B [187775] diluted with CF8A at 10, 25, and Table 18: Thermo-Calc simulation results comparing alloy 52i-B dilution series eutectic components Table 19: Thermo-CALC (Schiel-Gulliver) simulation results for 52MSS-E [HV1500] diluted with CF8A at 10, 25, and Table 20: Thermo-Calc simulation results comparing alloy 52MSS-E dilution series eutectic components Table 21: Thermo-CALC (Schiel-Gulliver) simulation results for Niobium, Titanium, and Molybdenum partitioning coefficients (k) vs. iron (Fe) and niobium (Nb) content for all simulated compositions Table 22: EDS compositional analysis vs nominal heat composition xii

14 List of Figures Figure 1: Cross section of typical 52M weld overlay (WOL) on a vessel nozzle-to-pipe dissimilar metal connection (1)... 3 Figure 2: Boundaries present in single phase austenite weld metal (16)... 7 Figure 3: Ductility-Dip Cracking formation and the influence of grain boundary tortuosity (21) Figure 4: The shrinkage-brittleness theory based on a eutectic phase diagram (25) Figure 5: Basic concepts of strain theory: Total amount and rate of strain concentrated on the liquid films determines the development of hot tearing in a casting (26) Figure 6: The stages of solidification according to Borland s generalized theory (25) Figure 7: Features of solidification crack surface (SUS 310S weld metal) (29) Figure 8: Modified generalized theory of cracking susceptibility during solidification (28) Figure 9: Stages of modified generalized theory (25) Figure 10: Technological strength theory (25) Figure 11: Technological strength theory proposed by Prokhorov (25) Figure 12: Schematic of the longitudinal-varestraint test (34) Figure 13: Schematic of Transvarestraint test (36) Figure 14: CPTT Apparatus developed by F.C. Hull (14) xiii

15 Figure 15: CPTT molds used by F.C. Hull (14) Figure 16: CPTT test on stainless steel by F.C. Hull (14) Figure 17: Threaded Copper Molds (37) Figure 18: Grooved copper mold (left) Finished grooved copper mold castings (right).. 32 Figure 19: First iteration of the second generation CPTT (39) Figure 20: Quantitative criteria for evaluation of solidification cracking susceptibility using the original second generation CPTT (39) Figure 21: Third Generation Cast Pin Tear Test Apparatus (15) Figure 22: Filler wire cutting aparatus Figure 23: Button melter apparatus Figure 24: New Generation CPTT apparatus (left) Water cooled copper CPTT coil (right) Figure 25: Schematic of third generation CPTT apparatus (15) Figure 26: Quartz insert for the CPTT. Large (>13g): A=23 mm, B=15 mm, C=10 mm Small (<13.5g): A=23 mm, B=13 mm, C=13 mm Figure 27: Drawing of copper relief disc (left). Photograph of copper relief disc (right): green depicts area of acceptable flashing; red depicts area of unacceptable flashing Figure 28: Stack up inside CPTT mold retainer: Mold half, Cu disc, spring, and spacer (left-right) Figure 29: CPTT LabVIEW control screen Figure 30: Schematic of CPTT purge procedure Figure 31: Scatter plot of CPTT results for 52M [NX7206TK] xiv

16 Figure 32: Scatter plot for CPTT results for tested alloys. (a) Filler metal 82(A) (b) 52i-B [187775] (c) 52M [NX7206TK] (d) TG-SN690Nb [FBH2280] (e) 52MSS-E [HV1500] (f) 52M [NX0T85TK] (g) 52MSS-C [NX77W3UK] Figure 33: Bar chart of visual ranking for tested alloys; Red bars represent the minimum pin length at which 100 circumferential cracking occurred. Grey bars represent the maximum pin length at which 0 cracking occurred Figure 34: CPTT visual inspection results of circumferential cracking (52i-B dilution study) Figure 35: CPTT visual inspection results (52MSS-E dilution study) Figure 36: Tensile force vs. elongation, or ram position, during tensile testing of 52i-B CPTT samples Figure 37: CPTT visual (red) and tensile (blue) evaluation method comparison for 52i-B [187775] Figure 38: Bar charts of the maximum (a), average (b), and minimum (c) reduction in UTS for several filler metals at pin lengths 0.75 inches to 1.25 inches. The average of the UTS reduction over these pin lengths is labeled for each alloy (orange bar) Figure 39: Visual ranking of CPTT samples using the average of maximum circumferential cracking from 0.75 inch to 1.25 inch pin lengths (ranking criteria: orange bar) Figure 40: Bar charts of the maximum reduction in UTS for the alloy 52i-B dilution series at pin lengths 0.75 inches to inches. The average of the UTS reduction over these pin lengths is labeled for each alloy (orange bar) xv

17 Figure 41: Visual ranking of alloy 52i-B dilution series CPTT samples using the average of maximum circumferential cracking from 0.75 inch to inch pin lengths (ranking criteria: orange bar) Figure 42: Bar charts of the maximum reduction in UTS for the alloy 52MSS-E dilution series at pin lengths 0.75 inches to inches. The average of the UTS reduction over these pin lengths is labeled for each alloy (orange bar) Figure 43: Visual ranking of alloy 52MSS-E dilution series CPTT samples using the average of maximum circumferential cracking from 0.75 inch to inch pin lengths (ranking criteria: orange bar) Figure 44: Bar chart of Thermo-CALC (Schiel-Gulliver) simulation results. The NbC (red) and Laves (green) start temperatures are depicted within the solidification temperature range (blue) Figure 45: Bar chart of Thermo-CALC simulation results showing the percentage of primary FCC solidification within the solidification temperature range Figure 46: Bar chart of Thermo-CALC (Schiel-Gulliver) simulation results. Total mole percent eutectic (NbC + Laves) is depicted in blue (labeled) with NbC and Laves (labeled) depicted in red and green respectively Figure 47: Bar chart of Thermo-CALC (Schiel-Gulliver) simulation results for 52i-B dilutions. The NbC (red) and Laves (green) start temperatures are depicted within the solidification temperature range (blue) xvi

18 Figure 48: Bar chart of Thermo-CALC (Schiel-Gulliver) simulation results for 52i-B dilutions. Total mole percent eutectic (NbC + Laves) is depicted in blue with NbC and Laves depicted in red and green respectively Figure 49: Bar chart of Thermo-CALC (Schiel-Gulliver) simulation results for 52MSS-E dilutions. The NbC (red) and Laves (green) start temperatures are depicted within the solidification temperature range (blue) Figure 50: Bar chart of Thermo-CALC (Schiel-Gulliver) simulation results for 52MSS-E dilutions. Total mole percent eutectic (NbC + Laves) is depicted in blue with NbC and Laves depicted in red and green respectively Figure 51: Bar chart of Thermo-CALC (Schiel-Gulliver) simulation results for Niobium, Titanium, and Molybdenum partitioning coefficients (k) as a function of iron content (Fe) for all simulated compositions Figure 52: Fracture surface of 100 cracked pins for each tested alloy: (a) Filler Metal 82(A) (b) 52i-B [187775] (c) 52M [NX7206TK] (d) TG-SN690Nb [FBH2280] (e) 52MSS-E [HV1500] (f) 52M [NX0T85TK] (g) 52MSS-C [NX77W3UK] Figure 53: Photomicrographs comparing crack "backfilling" in alloys (a) 52MSS-C and (b) 52M Figure 54: Micrographs depicting the relative differences in the magnitude of crack "backfilling" among tested alloys Figure 55: EDS analysis at solidification crack tip: 52M [NX7206TK], 15g, 1.875in, 65 circumferential cracking xvii

19 Figure 56: EDS analysis at solidification crack tip: 52i-B [187775], 13.5g, 1.5in, 72 circumferential cracking Figure 57: EDS analysis at solidification crack tip: 52MSS-C [NX77W3UK], 13g, 1.375in, 53 circumferential cracking Figure 58: EDS analysis at solidification crack tip: FM82(A), 15g, 1.875in, 77 circumferential cracking Figure 59: EDS analysis at solidification crack tip: TG-SN690Nb [FBH2280], 13g, 1.375in, 53 circumferential cracking Figure 60: EDS analysis at solidification crack tip: 52MSS-E [HV1500], 13.5g, 1.5in, 42 circumferential cracking xviii

20 Chapter 1: Introduction Ni-base alloy 600 filler metals, Alloy 182 and 82, have been commonly used in pressurized water reactors (PWRs) to join stainless steel piping to low alloy steel (LAS) components (1). The LAS components, typically vessel nozzles or valve bodies, have a coefficient of thermal expansion (CTE) which differs significantly from that of stainless steel; Alloys 182 and 82 provide a CTE which bridges this mismatch (2) (3). In addition to its thermo-mechanical properties, Alloy 600 was selected for this application due to its inherent corrosion resistance, although over the last 30 years primary water stress corrosion cracking (PWSCC) has been observed in dissimilar metal welds of reactor cooling system components welded with this alloy. Stress corrosion cracking of metals requires three factors to occur: environment, material conditions, and stress (2). The susceptibility of Alloy 600 and its filler metals are dependent on factors which include: chemical composition, metallurgical condition, and heat treatment. The element which appears to influence SCC most significantly is Chromium, which effects the formation of a protective Cr oxide layer and carbide precipitation during material processing (4) (5). Other factors such as grain size, degree of cold work, and residual and operating stresses play a role in Alloy 600 PWSCC susceptibility. In order to eliminate PWSCC one of the three controlling factors must be addressed. The stress component can be addressed through the application of a weld 1

21 overlay covering the length of the nozzle connection. An overlay such as this would apply a compressive stress at the crack tips which already exist in the dissimilar metal weld due to thermal contraction upon cooling; inhibiting further crack propagation and initiation. Another approach to mitigating this problem is through the resistance to PWSCC of the alloy which is used to overlay the connection. By increasing the Chromium from 22 (wt) to 30 (wt) the susceptibility of the overlay to PWSCC can be significantly reduced if not eliminated. This increase in Cr content has been shown to produce oxides with better mechanical properties and higher passivation rates as well as reduce Cr depletion after carbide precipitation (6). Alloy 600 requires a hour period at elevated temperature in order to replenish the chromium depleted regions created during carbide precipitation whereas Alloy 690, having higher chromium content requires significantly less time, and is virtually PWSCC immune if the carbon content is maintained below 0.02 (wt) (7) (8) (3). Repair and replacement of these connections since the 1980s have mainly utilized Alloy 690 and its filler metals (7). These high-cr Ni-base filler metals, 52 (ERNiCrFe-7) and 52M (ERNiCrFe-7A), have been implemented in nuclear applications where PWSCC resistance is required. The repair is done by depositing a multi-pass overlay across the entire dissimilar metal joint from the low alloy steel nozzle to the connecting stainless steel pipe, Figure 1. 2

22 Figure 1: Cross section of typical 52M weld overlay (WOL) on a vessel nozzle-to-pipe dissimilar metal connection (1). While PWSCC resistant, weld filler metals of alloy 690 have been found to exhibit a higher susceptibility to ductility-dip cracking (DDC) then filler metal 82 during the multipass weld stages of this overlay (9) (10). This can be attributed to a reduction in carbide and secondary phase forming elements, such as Niobium and Titanium, which have been shown to pin grain boundaries enhancing their tortuosity and therefore their resistance to DDC. Filler metal 52MSS is the next generation of high-cr Ni-base filler metal containing additions of 2.5 (wt) Niobium and 3.5 (wt) Molybdenum. Strain-to-fracture testing of this alloy has shown that these elemental additions significantly improve the materials resistant to DDC (9). This improvement is a result of a NbC eutectic which forms along grain boundaries during solidification, this inhibits grain boundary migration which is necessary for the occurrence of DDC. Single sensor differential thermal analysis (SSDTA) as well as previous work performed on the cast pin tear test (CPTT) indicates that the additions of niobium result in a widened solidification temperature range and an 3

23 increased susceptibility to solidification cracking. This can be attributed to low melting point NbC eutectic which forms in high Niobium bearing Ni-base alloys (11) (12) (13). It is evident that throughout the development and evolution of these alloys that compositional changes made in order to mitigate one type of cracking problem have results in the creation of another. When developing alloys the weldability of the alloy needs to be considered in order for the material to be put in service for most applications. One way to incorporate weldability into the alloy development process is through the use of weldability testing. There are a variety of weldability tests which test material susceptibilities to a variety of different cracking phenomena, the focus of discussion here is on solidification cracking as evaluated by the cast pin tear test. The cast pin tear test was first developed by F.C. Hull in the 1950 s to evaluate solidification cracking in austenitic stainless steels (14). The test is carried out by casting pins with an enlarged head and foot, intended to provide restraint upon solidification. As the pin solidifies both the expansion of the mold and the contraction of the solidifying material result in the accumulation of strain in the sample. This strain is concentrated in the last liquid to solidify and if sufficient strain accumulates cracking occurs as it would in a weld. A new generation of cast pin tear test has been developed at The Ohio State University by T. Luskin which utilizes Hull s concept of induction levitation melting charges which are subsequently cast into pins of varying length (15). The goal of this work is to establish a repeatable procedure for the new generation CPTT that allows for reproducible results, establish solidification cracking susceptibility 4

24 rankings for existing high-cr Ni-base filler metals, and continue to investigate other aspects of CPTT process improvement. 5

25 Chapter 2: Background & Objectives Introduction This chapter contains information pertaining to the welding metallurgy of nickel-base alloys. Weld solidification is discussed as well as several cracking phenomena associated with these alloys. Welding Metallurgy of Ni-Base Alloys The alloys used throughout this study are solid-solution strengthened Ni-base alloys which are often used when a combination of moderate strength and excellent corrosion resistance is required. Ni-base alloys solidify as austenite and are fully austenitic when solidification is complete. During solidification segregation may occur in these alloys which can result in the formation of secondary phases during the terminal stages of solidification. Segregation results in compositional variation at boundaries within the austenitic microstructure. Three different types of boundaries can be observed in an austenitic microstructure, Figure 2. 6

26 Figure 2: Boundaries present in single phase austenite weld metal (16). (1) Solidification Sub grain Boundaries (SSGBs) are present between cells and dendrites. SSGBs have a low degree of misorientation and are the finest structure that can be resolved under an optical microscope. A compositional variation occurs at SSGBs relative to the dendrite cores after solidification is complete due to microscopic solute redistribution (17). (2) Solidification Grain Boundaries (SGBs) are created when packets, or groups, of sub-grains intersect. Each packet of sub-grains has a different growth direction and orientation which results in a SGB upon intersection. These boundaries have a high degree of misorientation which can result in dislocation formation. Compositional variations occur at SGBs due to macroscopic solute redistribution during solidification; as a result secondary phases typically form here. Formation of low melting point secondary phases can result in solidification cracking along SGBs in single phase austenitic alloys (17). (3) Migrated Grain Boundaries (MGBs) are the crystallographic components of a SGB that has migrated away from the compositional component. When a SGB 7

27 forms it has a compositional and crystallographic components, in some cases the crystallographic boundary can reduce its energy by straightening (17). These boundaries/interfaces present in the microstructure are important because they are often the location of cracks and other defects in nickel-base weld metal. The following section describes in detail some of the defects which can exist at these different boundary types. Weldability Several different factors interact with one another to influence an alloys cracking susceptibility during welding: material composition, non-equilibrium solidification, and the evolution of stresses. Different forms of cracking are defined by the temperature range they occur in. Hot cracking, which will be the focus of discussion here, is a common problem experienced during welding and occurs in the temperature range where liquid films are present along grain boundaries or elsewhere in the solidification structure (18). Ductility dip cracking, which occurs along MGBs, and solidification cracking, which occurs along SGBs, are two of the predominate weldability problems associated with the nickel-base alloys investigated in this study. Ductility-Dip Cracking Ductility-dip cracking is a solid state cracking phenomenon that has presented problems when multi-pass welding high-cr Ni-base filler metals. Materials that are susceptible to DDC experience a drop in ductility between the solidus temperature (TS) and one half of 8

28 the solidus temperature (0.5TS) (17). When welding enough thermal strain arises and can exceed the ductility of the material in this temperature range resulting in the formation of a ductility-dip crack. Several different theories have been proposed to explain the mechanisms behind DDC (19) (20) (21) (22). One of the more prominent theories, proposed by Ramirez and Lippold, is grain boundary sliding. Ductility dip cracking as mentioned previously occurs along migrated grain boundaries; as a result migration of a grain boundary must occur before this cracking phenomenon can. As a result the DDC susceptibility of a material is affected by the tortuosity of grain boundaries, which is a function of the formation of precipitates and secondary phases along grain boundaries during solidification. These formations inhibit grain boundary migration effectively eliminating one of the conditions necessary for DDC to occur. Without this precipitation and secondary phase formation DDC becomes more prevalent, which is evident when comparing strain-to-fracture results of high purity alloys and alloys with an abundance of secondary phase forming elements (21) (17). The grain boundary sliding effect is represented schematically in Figure 3. 9

29 Figure 3: Ductility-Dip Cracking formation and the influence of grain boundary tortuosity (21). Work done by Lippold and Nissley showed that DDC susceptibility is greatly reduced by the addition of Niobium in high-cr Ni-base alloys due to the precipitates it forms at the end of solidification (23). These finding have provided some of the motivation for the development of alloys 52i, 52MSS, and other alloy 52 variations. 10

30 Solidification Cracking Weld solidification cracking is an intergranular cracking mechanism that occurs along solidification grain boundaries. This type of cracking is often associated with the presence of liquid films at grain boundaries during solidification. There are many different theories which attempt to explain solidification cracking phenomena. The shrinkage-brittleness theory, the strain theory, the generalized and modified generalized theory, and the technological strength theory all propose mechanisms to explain solidification cracking and the mechanisms behind it. The Shrinkage-Brittleness theory of solidification cracking was developed as a result of experimentation done with aluminum castings. The shrinkage-brittleness theory defines a coherency temperature, which is the temperature at which a coherent dendrite structure is first formed during solidification (24). Above the coherence temperature the ratio of liquid/solid is large, upon cooling the ratio of liquid/solid decreases until the coherency temperature is reached. At this point the solid formations begin to interact with one another forming a rigid network of solid-solid bridges that is capable of accumulating strain due to shrinkage; this is the onset of the Effective Interval or Brittle Temperature Range (Figure 4). Fracture according to this theory occurs at theses solidsolid interfaces due to strain accumulation over the effective interval. 11

31 Figure 4: The shrinkage-brittleness theory based on a eutectic phase diagram (25). The effective interval lies between the coherency temperature and effective solidus temperature. According to the shrinkage-brittleness theory solidification cracking occurs in this effective interval and the susceptibility to cracking increases with interval size. The size of the effective interval is influenced primarily by alloy composition as well as the amount and distribution of eutectic liquid within the solid network. Meaning a cracked alloy that undergoes a eutectic reaction at the end of solidification may be capable of crack healing if sufficient liquid is present to backfill the crack. Fracture according to this theory occurs at solid-solid interfaces due to shrinkage strain accumulation. The strain theory was developed using a casting method with chilled/restrained ends designed to develop solidification cracking in the sample. The theory defines a mushy stage and a film stage during solidification (26). As a welding arc travels across a work 12

32 piece it produces a pool of molten liquid which extends behind the arc until what Pellini defines as the mushy stage is reached. During the mushy stage cracking is not possible due to the presence of sufficient liquid and lack of a solid network. When the weld metal reaches the film stage strain can begin to develop along liquid films in the solid. At the onset of the film stage the liquid films are thick and continuous requiring very little strain to develop separation of the liquid but a great deal of separation to develop a fracture. Conversely later on the films decrease in thickness and continuity resulting in increased strain required for separation and reduced separation required for fracture to occur. The occurrence of fracture, or solidification cracking, is determined by the rate of strain imposed on the liquid films during the film stage. Figure 5 shows the relationship between temperature and time of film life under segregate and non-segregate conditions. 13

33 Figure 5: Basic concepts of strain theory: Total amount and rate of strain concentrated on the liquid films determines the development of hot tearing in a casting (26). Alloys with very low impurity content, under non-segregate conditions, pass through the film stage very quickly and are fairly resistant to solidification cracking. Under segregate conditions, specifically the presence of sulfur, the liquid film stage occurs over a larger temperature range. This depression of the solidus allows for the accumulation of more strain during the film stage and an increased susceptibility to solidification cracking. According to the Strain theory solidification cracking occurs at solid liquid interfaces therefore the fracture surface will exhibit smooth dendritic fracture morphology (25). 14

34 The generalized theory modifies and builds upon both the shrinkage-brittleness theory and the strain theory to explain how liquid quantity and distribution during solidification impact cracking susceptibility (27). Borland describes the solidification process as such: Stage 1: (Primary dendrite formation occurs) The solid phase is dispersed in continuous liquid and capable of relative motion (mushy stage). Stage 2: (Dendrite interlocking occurs) Continuous solid phases exist and continuous liquid phases move freely between interlocked dendrites. Stage 3: (Grain boundary development) Semi-continuous network of solid dendrites restricts liquid movement, eliminating relative motion between phases. Stage 4: (Solidification) Remaining liquid has solidified. These stages are provided on a binary phase diagram in Figure 6. Figure 6: The stages of solidification according to Borland s generalized theory (25). 15

35 The generalized theory bears a striking resemblance to the shrinkage-brittleness theory in that it the idea of a coherency temperature where the onset of solid-solid bridging begins (the onset of stage 2). Generalized theory modifies the effective interval proposed by Pumphrey and Jennings, separating it into two stages (stage 2 & 3). During stage 2 failure of the interlocking dendrite matrix can be tolerated due to the presence of sufficient liquid films, which facilitate crack healing. Stage 3 is referred to as the critical solidification range the onset of which is known as the critical temperature, during this stage cracking susceptibility is highest. During stage 4 cracking is not possible due to the lack of liquid present. Modified generalized theory was developed in by Matsuda and colleagues at JWRI in the 1980 s. Matsuda utilized an optical microscope to observe and document solidification cracking phenomena in carbon steels, stainless steels, and Inconel alloy during welding (28). The Fractographic features of a solidification crack provide important information regarding solidification phenomena. The authors of this theory have described a solidification crack surface as having three characteristic regions (Figure 7): Type D: Dendritic fracture surface. Type D-F: Transient dendritic to flat fracture surface. Type F: Flat fracture surface. 16

36 Figure 7: Features of solidification crack surface (SUS 310S weld metal) (29). The most notable modification to the generalized theory is that region susceptible to solidification cracking is shifted up into a higher temperature zone near the liquidus temperature. This contradicts the notion, that the solid phase gradually advances as the temperature decreases, discussed in Pellini s strain theory (26). Through dynamic observation it was shown that the solid phase rapidly advances during the early stages of solidification. Figure 8(a) shows the rapid growth of secondary dendrites and the distribution of liquid during the initial stage of solidification. Networks of dendrites are present during the liquid film stage here which occurs near the onset of solidification. This varies from the liquid film stage associated with strain theory in that the solid is not fully covered with a continuous liquid film. The critical temperature described by Borland is represented by the boundary between the liquid mass stage and the liquid film stage. Figure 8(b) describes how solidification crack initiation and propagation occurs according to the modified generalized theory. Cracking occurs in the liquid film stage and propagates toward the liquid mass and liquid drop stages with propagation occurring slower in the liquid mass stage due to crack healing (28). 17

37 Figure 8: Modified generalized theory of cracking susceptibility during solidification (28). According to Matsuda stage 1 and 2 are similar to the stages described previously in generalized theory although the temperature range over which they occur is much smaller and closer to the liquidus temperature. Stage 3 has been divided into two parts: Stage 3(h): (Liquid Film Stage) this stage is susceptible to crack initiation as well as crack propagation (Type D-F). Stage 3(l): (Liquid Droplet Stage) this stage is only susceptible to crack propagation not initiation (Type F). 18

38 Figure 9: Stages of modified generalized theory (25). The figure above (Figure 9) provides a graphical representation of the stages of modified generalized theory on a binary phase diagram. Dr. Prokhorov studied the mechanical behavior of materials during welding and in 1962 introduced the technological strength theory (30). This theory takes into account the technological strength of a material during welding or casting which can be evaluated by comparing the minimum ductility/deformability of a material with the deformation accumulated while passing through the brittle temperature range. The brittle temperature range, similar to theories described previously, represents a dip in ductility during solidification. Above the brittle temperature range deformation does not cause any irreversible changes as the alloy behaves like a liquid, within the brittle temperature range 19

39 deformation is able to accumulate and can result in cracking. The total deformation in the BTR is the summation of the shrinkage deformation and mechanical deformation, represented in Figure 10 by line A-C. Line A-B represents the shrinkage deformation and line A-D the critical deformation to cause cracking (25). Figure 10: Technological strength theory (25). When the total deformation exceeds the ductility curve, and enters the BTR, the accumulated deformation is sufficient to exhaust the ductility of the material and cause solidification cracking. This is illustrated in Figure 11, were three different strain rates have been superimposed over a ductility curve. 20

40 Figure 11: Technological strength theory proposed by Prokhorov (25). Strain rate C does not intersect the BTR and therefore will not be susceptible to cracking. Strain rates A and B are both susceptible to cracking as they pass through the BTR, although strain rate A is susceptible to over a much larger temperature range (SCTR) than strain rate B. The technological strength theory does not take into account microstructural aspects of a material rather cracking response is determined by the competition between accumulated deformation and ductility recovered in a material during solidification. Weldability Testing A variety of tests which evaluate a materials susceptibility to solidification cracking exist today. These test methods create strain conditions by self-restraint (intrinsic) or externalrestraint (extrinsic). Self-restrained solidification cracking tests use the thermomechanical reaction of the material to induce strain accumulation during solidification. In other words the shrinkage of the specimen itself results in strain accumulation during 21

41 solidification. One example of a self-restrained solidification cracking test is the Cast Pin Tear Test. Externally-restrained tests utilize external mechanical restraint such as bending or applying a tensile load. It is important to note that test methods are designed to be either self or externally restrained not a combination of the two; this simplifies quantification of strain levels during testing. Some examples of externally-restrained solidification cracking tests are: The Hot tensile test, the Programmable Deformation Rate test (PVR), and the Transvarstraint test (31) (32). The Varestraint Test The Variable restraint, or Varestraint test, was first introduced in 1965 by Savage and Lundin to study hot cracking (33). Varestraint testing is done by creating an autogeneous bead-on-plate weld on the surface of a test coupon which is then bent over a die block of a fixed radius. Bending results in straining of the specimen, referred to as augmented strain, and can be varied by the radius of the die block. The strain in a sample can be calculated by the relationship provided below; where εt is the augmented strain, t is the sample thickness, and R is the die block radius (Equation 1). Equation 1: Varestraint strain calculation. t t 2R The original test setup is provided below in Figure

42 Figure 12: Schematic of the longitudinal-varestraint test (34). During Varestraint testing cracking is observed in both the fusion zone and the heataffected zone (HAZ) adjacent to the weld pool. Several methods have been proposed to quantify the cracking response of a sample: Threshold strain (εth); this is the minimum strain at which cracking is observed, so a lower threshold strain implies lower resistance to crack formation. Total crack length (TCL); can be considered in the fusion zone or heat-affected zone, larger TCL values in the fusion zone indicate greater susceptibility to solidification cracking whereas larger TCL values in the HAZ indicate greater susceptibility to HAZ liquation cracking. Maximum crack length (MCL); The maximum crack length is typically observed in the fusion zone and correlates well with a materials solidification temperature range, larger MCL values are indicative of wider solidification temperature ranges and an increased susceptibility to cracking (33). 23

43 The Transvarestraint Test Varestraint testing gained momentum as a viable hot cracking evaluation method and in 1970 the first variation of the test was introduced by McKeown (35). McKeown employed a concept similar to that of Savage and Lundin (33) with the major difference being in the orientation of the weld with respect to the axis of stress. The weld in this case is made parallel to the radius of the die block and the axis of bending is perpendicular or transverse to the welding direction. The strain in samples produced using this method can be calculated using the relationship provided below; where εt is the augmented strain, t is the sample thickness, and R is the die block radius (Equation 2). Equation 2: Transvarestraint strain calculation. t t 2R t This relationship differs from that of the longitudinal, or varestraint, test because the transvarestraint test represents a simply supported beam rather than a cantilever beam. The schematic of a typical transvarestraint test is provided below in Figure

44 Figure 13: Schematic of Transvarestraint test (36). This orientation (transverse) tends to produce cracking solely within the fusion zone of the weld allowing this test method to focus on solidification cracking along the centerline of a weldment. The Cast Pin Tear Test The essence of the Cast Pin Tear Test is that as molten material is poured into the room temperature mold the heating of the mold material results in expansion while the cooling of the charge material results in contraction. As a result tensile strain accumulates in the sample as it solidifies. If this strain exceeds the ductility of liquid films present along grain boundaries during solidification, cracking will result. In other words longer pin lengths result in larger accumulations of tensile strain which result in larger cracking responses. Alloys are tested over a series of pin lengths and the cracking response at each length is the observed and recorded. Material ranking is done according to the maximum pin length at which 0 cracking consistently occurs and the minimum pin length at 25

45 which 100 cracking occurs. Using these criteria alloys can be ranked and their relative susceptibilities to solidification cracking compared. This test method offers a quick and economical means for determining different material compositions solidification cracking susceptibilities. The speed and small amount of material required to rank an alloy makes the Cast Pin Tear Test a viable evaluation method for ranking existing materials, studying the effects of dilution, and the development of new materials. The cast pint tear test was first developed by F. C. Hull in the 1950 s as a means to rank alloys by their susceptibility to hot cracking (14). For this test induction levitation melting is used to melt 19 gram samples within an inert atmosphere. This method of charge heating reduces the likelihood of contamination by eliminating the need for crucibles to contain the molten material. The molten sample is then cast into the shape of a tapered pin in a series of copper molds with varying pin length. Pouring for casting is controlled as the coil behaves like a magnetic funnel with an adjustable pour rate. Restraint at the ends of these pins results in the accumulation of tensile stress on the sample as the mold expands upon heating and the pin contracts upon cooling. The pin geometry is designed to produce some hot cracking response, dependent on the pin length, which provides a means of ranking alloy susceptibility. The testing apparatus used by Hull consisted of a large vacuum-tight cylindrical glass chamber which contained a copper induction coil, a series of up to 18 molds, and a manipulator, Figure

46 Figure 14: CPTT Apparatus developed by F.C. Hull (14). Hull found that the cracking response in a pin decreased as the pin length shortened or as the outside mold-diameter increased. Using this knowledge a series of molds were selected that produced comparable changes in each interval of pin length/diameter ratio, with each length/diameter ratio having its own arbitrary number assignment. The original design consisted of a split mold, a split restraining lock, and a bottom plate. The dimensions of the selected mold geometry are provided below in Figure

47 Figure 15: CPTT molds used by F.C. Hull (14). The pin volume is kept constant at 2.4 cc (equivalent to 19 grams of stainless steel) and the distance from the levitation coil to the mold maintained at in. for consistency from test to test. The surface of the pins is examined at up to 30 times magnification for cracking using a device for support and rotation which allowed for angular measurement of the cracking. Interdendritic cracks usually extend to the surface of the pins and propagate circumferentially so the sum of this crack type is totaled. Since the cracking response 28

48 often exceeds 100 of the pin s circumference Hull used an index to compare tested alloys. The Cracking Index is the sum of angular cracking measurements as a percentage of the circumference, which reaches a maximum when 100 circumferential cracking is reached, Figure 16. Figure 16: CPTT test on stainless steel by F.C. Hull (14). It was found during testing that the data produced by the CPTT was subject to a certain amount of scatter. This can be attributed to the fact that cracking involves the interaction of chemical, mechanical, thermal, and geometric effects that are not under simultaneous control, meaning a sufficient number of tests must be run to ensure a representative cracking response is achieved for a given length. The apparatus was used by Hull to test and rank over 100 different stainless steel alloys as well as to study the iron corner of the Fe-Ni-Cr system. In this study Hull used the 29

49 CPTT to show that the composition of a sample, or the amount of ferrite present, can be related to the phase diagram and correlated with the amount of cracking. His results were found to be in agreement with general welding experience; a small percentage of delta ferrite is beneficial. The Threaded Copper Mold Test The Threaded Copper Mold Weldability Test was developed by Talento in the 1960s (37). This version of the CPTT induction levitation melted 37 gram samples within an argon purged chamber similar to Hull s method. Once molten the samples are cast into threaded copper molds and a count of surface fissures was made on the as-cast specimen at 7X magnification, Figure 17. Figure 17: Threaded Copper Molds (37). 30

50 The threaded copper mold consisted of a cylindrical shaped cavity with 20 threads per inch on the inside with inch root diameter and inch top diameter. The thread pitch was adjusted in order to develop a stress which gives a meaningful index of fissure sensitivity when samples are cooled from casting temperature and the cooling-rate was controlled by mold wall thickness. Three different molds with varying external wall thickness were used, in order to determine which cooling rate produced the greatest sensitivity in fissuring response. The mold with the lowest cooling rate, or lowest wall thickness, was selected due to its resolution in identifying solidification cracking susceptibility differences. This mold design was very advantageous because it allowed one casting to represent a full range of experiments; additional experimentation was only done in order to improve the statistical accuracy of the results. Typically 30 pins were cast of each alloy. As mentioned previously the fissures were counted on each sample at 7X magnification and totaled over the 30 samples. This version of the CPTT was used to evaluate the cracking susceptibility of Ni-base alloys (ErNiCr-3) and their dilution with alloy 600. The Grooved Copper Mold Test The Grooved Copper Mold Test, developed by Armao and Yeniscavich, similar to the cast pin tear test developed by F. C. Hull (14) is a method for evaluating relative differences in hot cracking susceptibility between alloys (38). The test apparatus used is very similar to that employed by Hull. The test consists of five specimens weighing 32 grams each which are induction levitation melted within an argon purged chamber. The 31

51 samples are then cast into grooved copper molds and examined at 25X magnification for cracking, Figure 18. Figure 18: Grooved copper mold (left) Finished grooved copper mold castings (right) The procedure for evaluating of the grooved copper mold castings involves counting the number of cracks on each and multiplying it by the weighing factor corresponding to its size category, Table 1. Table 1: Grooved copper mold crack evaluation weighing factors 32

52 The weighted totals for each category are then added to give a fissure index for each sample. The fissuring indexes of the five samples are then totaled. Previous work done using this set-up required 30 test samples but due to test consistence the sample number has been reduced to 5, as a result the total fissuring index (GCM) of the 5 samples is multiplied by 6. Alloy 600 as well as EN82 was tested using the grooved copper mold setup. It is shown in the results of this work that the test is sensitive to the effects of alloying additions found in filler materials which reduce weld metal crack sensitivity. The Cast Pin Tear Test (OSU) Several iterations of the Cast Pin Tear Test have been developed at The Ohio State University. The original second generation test utilized a retractable gate in the bottom of a water-cooled copper hearth on which a sample was melted using a gas-tungsten arc welding torch and released into the mold below. The CPTT apparatus and half of a mold are provided in the figure below (Figure 19).Figure 19: First iteration of the second generation CPTT. 33

53 Figure 19: First iteration of the second generation CPTT (39). The molds used with this version of the CPTT resemble those used by Hull, although in this case a constant diameter was maintained and pin length was varied. Pin lengths ranging from 0.5 in. to 2 in. were cast with each respective length requiring a specific mass to fill the entire volume. The stress level in a sample is controlled by the mold material and pin length (39) (12). The enlarged head and foot of the pin impose restraint on the sample during solidification allowing stress to accumulate in the last liquid to solidify, longer pin lengths and fast cooling rates typically result in the accumulation of more stress. For this testing the molds were composed of a copper-beryllium alloy with a measured cooling rate of 245 C/s (15). The parameters used for this version of the CPTT are provided in Table 2. 34

54 Table 2: Second generation CPTT procedure (39). Cast pins were examined at 10-70X magnification and circumferential cracking was measured for each pin with representing of the pin. Based on the recorded results a response curve of Maximum Circumferential Cracking (MCC) is plotted vs. pin length. The maximum pin length with 0 cracking and the minimum pin length with 100 cracking are the quantitative criteria by which alloys were ranked, Figure

55 Figure 20: Quantitative criteria for evaluation of solidification cracking susceptibility using the original second generation CPTT (39). The newest generation of the CPTT, developed at The Ohio State University by T. Luskin, incorporates induction levitation melting into the casting process. A combination of the original test introduced by Hull and the previous iteration of the second generation CPTT developed at OSU, the new generation CPTT utilizes induction levitation melting to cast samples into molds of constant pin diameter with varying pin length. An image of the third generation CPTT apparatus is provided in Figure 21 36

56 Figure 21: Third Generation Cast Pin Tear Test Apparatus (15). The process begins with the preparation of charges with masses corresponding to the desired pin lengths. The charge is then placed into a cooled copper coil and levitation melted at a predefined AC current at 434 hertz. Once the sample is completely liquid and a set temperature is registered by the optical pyrometer above the sample, the current in the coil is ramped down and the sample descends into the mold beneath. Summary The structural integrity of dissimilar metal welds between low alloy steel nozzles and stainless steel safe ends has been shown to be degraded by primary water stress corrosion cracking (PWSCC) (40). These dissimilar metal welds were originally made using alloy 82/182 (ERNiCr-3) which contains roughly 20 wt Cr. Since the discovery of this 37

57 degradation, due to PWSCC, alloy 52/152 (ERNiCrFe-7) has been developed with about 30 wt Cr which has been shown to be significantly more resistant to PWSCC (7). Several variations on alloy 52 have been developed in order to solve DDC challenges which have arisen during multi-pass welding. This improved DDC resistance has been achieved primary through additions of niobium which forms carbides during solidification (21). While additions of niobium appear to increase DDC resistant they also widen the solidification temperature range in which can promote solidification cracking (17). In addition to the individual weldability of these alloys the dilution of the low alloy and cast stainless steel substrate into the welding filler metal must be considered for this application. McCracken and Smith investigated the hot cracking susceptibility of filler metal 52M on cast austenitic stainless steel during the field application of a weld overlay repair (40). They discovered that solidification cracking occurred only in the first layer of alloy 52M which was deposited over a buffer layer of ER308L. This solidification cracking was considered to be a result of dilution into the first 52M pass. The diluted 52M was found to contain high levels of iron which has been shown to increase the partitioning of niobium during solidification. This increase in niobium partitioning widens the solidification temperature range and can increase the fraction eutectic, both of which can lead to solidification cracking (41). In this continuous process of alloy development small compositional changes that mitigate one cracking problem often have unforeseen consequences and result in other cracking issues. Acceptable levels of dilution and the influence of dilution on cracking 38

58 susceptibility is often overlooked during the alloy development process and discovered during material application. Significant time and effort goes into the development and production of these alloys and the current methods for determining alloy weldability and testing acceptable levels of dilution are very costly and time consuming. The cast pin tear test provides a potential solution to this problem. Requiring a relatively small amount of material and time the cast pin tear test can be used to produce relative material and dilution solidification cracking susceptibility rankings in order to allow the weldability of prospective alloys to be investigated during the alloy development process. Objectives The objective of this study is to evaluate and improve the reliability of the New Generation Cast Pin Tear Test and apply the latter for evaluation and ranking the solidification cracking susceptibility in a series of high-cr Ni-base filler metals. Susceptibility rankings are generated for seven undiluted welding filler metals and three dilute filler metals at several levels of dilution with cast stainless steel (CF8A): 52M (NX7206TK) (0/10/40 dilution), filler metal 82(A), 52MSS-E (HV1500) (0/10/40 dilution), TG-SN690Nb (FBH2280), 52M (NX0T85TK), 52MSS-C (NX77W3UK), and 52i-B (187775) (0/10/40 dilution). In addition to establishing solidification cracking susceptibility rankings, fractographic, compositional, and thermodynamic analyses are performed in order to better understand the root cause for the variation in cracking susceptibility. The overarching objectives are listed below: 39

59 1) Optimization of the CPTT procedure for improved reproducibility and reliability of test results. 2) Evaluation of the effect of alloy composition and dilution with cast stainless steel substrate on the solidification cracking susceptibility in a series of high-cr Nibase alloys. 3) Generation of solidification cracking susceptibility rankings in a series of high-cr Ni-base alloys. 40

60 Chapter 3: Materials & Experimental Procedure Introduction The following chapter contains information pertaining to the tested materials and experimental methods used throughout this study. Materials Seven Ni-base filler metals were evaluated in this study: one heat of (ERNiCr-3) alloy 82, two heats of (ERNiCrFe-7A) alloy 52M, two heats of alloy (ERNiCrFe-13) 52MSS, one heat of (ERNiCrFe-15) alloy 52i-B, and one heat of (ERNiCrFe-7A) alloy TG- SN690Nb. The compositions of these alloys are provided in Table 3. 41

61 Table 3: Material compositions (wt). Alloy Compositions (weight ) Alloy FM82 * 52i-B 52M 52M ( ) (NX7206TK ) (NX0T85TK ) TG- SN690Nb (FBH2280 ) 52MSS-C (NX77W3UK ) 52MSS- E (HV1500) Al CF8 A B C 0.1 max Co Cr Cu Fe 3 max M n Mo Nb Ni 67 min P 0.03 max S max Si Ta Ti 0.75 max Zr *Composition from AWS 5.14 ERNiCr-3 42

62 Weldability Testing Procedure Material Preparation The material to be tested goes through three steps prior to being melted into a charge in the button melting apparatus. (1) First the material is cut, often from a spool of welding wire, into pieces small enough to be accommodated by the hearth of a button melter (approximately 1 in. in length). This is done using a rotating blade in conjunction with a Miller automatic wire-feed system (Figure 22). Figure 22: Filler wire cutting aparatus. (2) Once cut the material is ultrasonically cleaned in ethanol for a minimum of 30 minutes. From this point in the process on nitrile gloves are worn while handling materials in order to avoid contamination. (3) The clean material is then weighed out to the mass corresponding to the desired pin length (Table 4), typically six buttons for each pin length. 43

63 Table 4: Mass requirements for Ni-based alloys at each pin length. Mass (g) Length (in) Button Melting The button melting apparatus consists of a GTA torch positioned over a water cooled copper hearth within a cylindrical chamber (Figure 23). The GTA torch is powered by a Miller Dynasty 300LX constant current power supply with high frequency arc initiation. 44

64 Figure 23: Button melter apparatus. Prior to using the button melting apparatus the copper hearth is polished using 800-grit sandpaper, clean thoroughly with ethanol, and dried in order to ensure cleanliness and avoid sample contamination. Once the hearth of the button melting apparatus has been cleaned the material to be melted is placed on the copper hearth and the melting chamber reassembled. Prior to melting a charge, the chamber is purged with argon using the following procedure: (1) Open inlet valve to allow argon flow into button melting apparatus. (2) Adjust flow rate to 20 CFH. (3) Close upper exhaust valve until chamber pressure builds above 10 psi. (4) Open upper exhaust valve until chamber pressure is reduced below 5 psi. (5) Repeat steps 3 and 4 five times 45

65 Oxygen concentration measurements are taken periodically following this purge procedure using a Purgeye 500 weld purge monitor. The average oxygen concentration recorded over five different purge cycles was found to be 10.4 parts per million (ppm) with a maximum value of 13 ppm and minimum of 9 ppm. After the chamber has been purged the sample is melted using a GTA torch with a variable foot pedal set to a maximum current of 130A. For charges of a single alloy the material is melted once until round and symmetrical. For charges containing several alloys or a dilute composition the material is melted once, allowed to cool then flipped over, re-purged and melted again to ensure uniform composition. Once melting is complete the sample is allowed to cool within the inert atmosphere until completely solidified (approximately 2 min.) before removal from the chamber. The specific parameters for the button melting procedure are provided below, Table 5. Table 5: Parameters for the button melting apparatus. Button Melting Apparatus Parameters Charge Mass (g) Shielding Gas Argon (99.998) Gas Flow Rate (CFH) 20 Build Pressure (psi) 10 Release Pressure (psi) 5 Purge Cycles (#) 5 Current (A) 130 Melt Time (sec)

66 Cast Pin Tear Testing The new generation Cast Pin Tear Test (Third generation) apparatus consists of a water cooled copper coil within a purged cylindrical chamber. The coil is powered by a 10 kw EasyHeat Li industrial power supply and work head which operates at 224 khz. The coil is composed of rectangular copper tubing and consists of seven turns in one direction and two turns in the opposite direction to maintain the stability of the levitating sample (Figure 24). Figure 24: New Generation CPTT apparatus (left) Water cooled copper CPTT coil (right). The majority of the purging and casting process has been automated using LabVIEW software, an optical pyrometer, several valves, and a proportional-integral-derivative controller (PID), the operation of which will be described in the following procedure (Figure 25). 47

67 Figure 25: Schematic of third generation CPTT apparatus (15). Prior to placing a charge in the CPTT apparatus the lens of the optical pyrometer is cleaned with ethanol to ensure the accuracy of the measured charge temperature during casting. A quartz insert, in the shape of a funnel, is then placed within the coil in order to isolate the sample from the coil prior to levitation. Two different size funnels are used depending on the size of the charge being melted, for charges 13g or less a small insert is used (23 mm upper dia. 13 mm lower dia. 13 mm height), for charges greater than 13g a large insert is used (23 mm upper dia. 15 mm lower dia. 10mm height) (Figure 26). The motivation for this is that the smaller charges expand enough prior to levitation causing them become stuck in the large insert, while the larger charges can become separated into two parts or contaminated during casting when descending through the smaller insert. 48

68 Figure 26: Quartz insert for the CPTT. Large (>13g): A=23 mm, B=15 mm, C=10 mm Small (<13.5g): A=23 mm, B=13 mm, C=13 mm. After ensuring the pyrometer is clean and the appropriate quartz insert is in place the charge can be placed in the funnel through the hatch at the top of the chamber. The mold retainer can then be assembled using the following procedure: (1) Select the mold corresponding to the mass of the charge, rub felt string impregnated with 12 µm diamond paste over gauged section to remove any oxide on mold surface, clean head and foot of mold with copper bristled brush. Next ultrasonically clean the mold in ethanol for 30 seconds, then dry thoroughly ensuring pin holes used for alignment are free of any residual liquid. Once completely dry and clean insert alignment pins into the mold, put the mold halves together, and install a rubber O-ring around the smaller outer mold diameter. (2) Select a spacer and spring that will allow for adequate pressure of the copper relief disc against the bottom of the mold during casting. The amount of pressure 49

69 can be determined experimentally by flashing produced at the foot of the pins on the surface of the copper relief disc (Figure 27). Figure 27: Drawing of copper relief disc (left). Photograph of copper relief disc (right): green depicts area of acceptable flashing; red depicts area of unacceptable flashing. If the flashing exists only in the channels of the copper disc and does not flow from one to the next then adequate pressure was used during casting. If the flashing flows from one channel of the copper disc to the next freely this is an indication that the disc was depressed during casting and an inadequate amount of pressure was used. The pressure against the spring can be increased by spacer height, spring length, and spring constant (k). (3) Once steps 1 and 2 have been completed the mold retainer can be assembled by placing the spacer, spring, copper disc, and mold into the mold retainer in that order, Figure 28. Make sure to place the mold seam perpendicular to the set screws so as they re tightened the mold is forced together. 50

70 Figure 28: Stack up inside CPTT mold retainer: Mold half, Cu disc, spring, and spacer (left-right). The assembled mold retainer is then screwed into the bottom of the CPTT chamber, the gas outlet hose connected, and the top hatch closed. Once the CPTT chamber has been sealed with a charge/mold inside the purge parameters can be set and the purge procedure initiated. The purge parameters; Initial Purge Time, Pressurize Time, and Release Time can be input at the bottom of the control screen (Figure 29), once the parameters have been input the purge procedure will initiate once the Start Purge button has been pressed. Figure 29: CPTT LabVIEW control screen. 51

71 When the Start Purge button has been pressed the purge procedure will execute the following steps: (1) Initial purge begins; Hi-flow valve and top release valve remain open for designated amount of time (typically 80 seconds). Argon enters and fills CPTT chamber from the bottom and releases out of the top (Figure 30). (2) Upon completion of the initial purge the top release valve closes, building pressure, for a designated amount of time (typically 4 seconds) and then opens, releasing pressure for a designated amount of time (typically 4 seconds). This occurs out of the top release valve three times and then subsequently out of the bottom release, or variable flow valve, three times. (3) Once pressurize and release cycles have been completed the high flow valve closes and the flow of argon into the chamber is diverted through the 0.5 psi regulator and the Lo-Flow valve. At this point the variable flow valve, or bottom release valve, is set such that there is a flow rate of 2.5 CFH out of the bottom of the mold retainer (can be read on the LO-Flow meter) and a pressure of 0.3 psi in the CPTT chamber (can be read on the pressure gauge). Once this has occurred the Purge Complete indicator will become green indicating completion of the purge cycle. 52

72 Figure 30: Schematic of CPTT purge procedure. After the purge cycle is complete a set temperature, casting current, ramp-down current and ramp-down time must be selected before the casting process can be initiated. The set temperature is typically C above the liquidus temperature of the material being tested, in this case a set temperature of 1430 C-1450 C was used for Ni-base alloys. Casting current dictates how high the sample levitates in the coil and how quickly it heats to the set temperature, for this work a casting current of 385A was maintained. The rampdown current and time influence the descent of the material into the mold, for this study a ramp-down current of 275A was selected and a ramp-down time of 0.85 seconds. Once all of the parameters have been input Start Melt can be selected on the control screen and the following casting procedure will occur: 53

73 (1) Following Start Melt the coil will ramp up to the designated Casting Current over 10 seconds. Once the casting current is reached it is maintained until the Set Temperature is registered by the optical pyrometer. (2) When the optical pyrometer above the levitating charge registers the Set Temperature the current in the coil is ramped down from the Casting Current to the Ramp-down Current over the Ramp-down Time. This allows the charge to descend into the mold in a more controlled manner than if the current is cut off or reduced more abruptly. It is important to note here that the coil remains energized during the ramp-down cycle. The ramp-down cycle does not occur until the set temperature has already been reached, as a result peak sample temperatures typically exceed the set temperature by approximately 50 C, which can vary depending on sample size and ramp-down cycle. (3) Casting is complete. The power supply can be turned off and the CPTT sample removed from the mold for analysis. Visual Analysis of Cracking Response in CPTT Samples The strain level which accumulates in the sample is controlled by the pin length, and at some pin length solidification cracking occurs (typically near the head of the pin). Upon further increase in pin length cracking will increase until 100 circumferential cracking is observed. The pins are examined using a binocular microscope at a magnification of up to 70x. Crack length is measured circumferentially around the pin in degrees and calculated using the equation below (Equation 3). 54

74 Equation 3: Circumferential cracking equation for CPTT evaluation. L T: Total length of all cracks measured on pin surface Alloy susceptibility to solidification cracking is then characterized primarily by the maximum pin length with 0 cracking and also by the minimum pin length with 100 cracking. Rankings between alloys are generated by plotting the maximum circumferential cracking response versus pin length. The maximum circumferential cracking (MCC) response is the largest measured circumferential cracking value in at least four samples at a given pin length. Metallurgical Characterization Metallography samples were prepared using the following procedure. Cast Pins were sectioned in both the longitudinal and transverse directions using a Techcut 5 precision sectioning saw. After sectioning samples were mounted in Beuhler conductive resin with a LECO PR-36 mounting press. Once mounted the samples were polished using 180, 240, 600, and 800 grit silicon carbide paper, followed by 9, 6, 3, and 1 µm diamond paste on a LECO Spectrum System In between each polishing step samples are cleaned ultrasonically in ethanol. The Mounted and polished samples, when required, were etched electrolytically using a 10 volume chromic acid solution (CrO3). Samples were placed in a glass dish and submerged in the chromic solution; the current and voltage of the constant voltage DC 55

75 power supply were then set to 1A and 5V respectively. The tungsten anode was then placed in direct contact with the sample and an cathode in the form of a stainless steel foil placed in solution above the sample. The anode and cathode remained in place for 5-10 seconds to allow etching of the sample to occur, after etching the samples were rinsed in a bath of water and ultrasonically cleaned in ethanol. Light optical microscopy (LOM) and scanning electron microscopy (SEM) were used in the evaluation of Cast Pin Tear Test samples. An Olympus GX51 metallograph was used to capture images and analyze cracking in etched metallography samples. Two scanning electron microscopes were used for sample analysis. First, the Quanta-200 general purpose SEM, and second the XL-30F ESEM field emission gun SEM. For both of the previously mentioned SEMs secondary electron (SE) mode was used for fractographic analysis of the samples, backscatter electron (BSE) mode for analysis of secondary phases such as low melting point eutectics, and energy dispersive spectroscopy (EDS) detectors for elemental identification and composition sample analysis. Typical working distance for all sample analysis was roughly 10mm, with an accelerating voltage of KV, and a spot size of 3-5. Computational Modeling of Solidification Computational modeling of the solidification process in the tested alloys was conducted using the Scheil-Gulliver module with thermodynamic software; Thermo-Calc software. The Thermotech Ni-Data V.7 database was used to calculate solidification temperature ranges, eutectic solidification start temperatures, mole percent eutectic, and elemental 56

76 partitioning coefficents. The calculations were executed assuming the following: equilibrium of the solid-liquid interface, no diffusion occurs in the solid and complete mixing of the liquid phase. Alloy simulations descended from 2000 C in 1 C intervals until the solid phase fraction reached 98. Laves, liquid, and FCC phases were included in this simulation (NbC has an FCC crystal structure). 57

77 Chapter 4: Results & Discussion Reproducibility of the CPTT Results The reproducibility of the new generation Cast Pin Tear Test was evaluated using filler metal 52M [NX7206TK]. Pin lengths 0.75 inches to inches were cast; the visual inspection results of circumferential cracking are provided in Table 6. Table 6: CPTT results for 52M [NX7206TK]. 52M NX7206TK Length (in) # Min. Cracking Max. Cracking Avg. Cracking Standard Deviation

78 Circumferential Cracking Six or more pins were cast at lengths of 0.75 inches, inches, and 1.75 inches and the standard deviation of their cracking response calculated. The pins cast in the 0.75 inch molds were found to have a standard deviation of 2.1, the inch molds 1.60, and 1.75 inch molds These results indicate that at the threshold pin lengths, maximum pin length with 0 cracking and minimum pin length with 100 cracking, percent circumferential cracking is a reproducible criterion by which alloy susceptibility to solidification cracking can be ranked. Fewer tests were performed at the intermediate pin lengths due to a large amount of scatter in the cracking response. Figure 31 is a graphical representation of the data in Table 6. The circumferential cracking response of each individual pin (red) and the averages of all the pins (blue) at each given length are provided. 52M (NX7206TK) Average Pin Length (in) Figure 31: Scatter plot of CPTT results for 52M [NX7206TK]. 59

79 Figure 31 shows quite clearly the reduced deviation in cracking response at lower magnitudes of cracking. For this reason the remainder of the Cast Pin Tear Testing focus is on identifying the cracking threshold, or maximum pin length with 0 cracking, in order to rank relative alloy susceptibilities to solidification cracking. The secondary ranking criteria, or minimum pin length with 100 cracking, is also determined and utilized in the event that the primary ranking criteria does not provide adequate resolution between alloy susceptibilities. Weldability Testing The following section contains results obtained using the Cast Pin Tear Test on pure welding filler metals and several filler metals diluted with cast stainless steel CF8A. A visual inspection of circumferential cracking is performed on each pin in order to develop solidification cracking susceptibility rankings. Following visual inspection pins are subject to tensile inspection and ranking. The results of the two different inspection methods are compared in order to evaluate the viability of the tensile testing inspection method. CPTT Visual Crack Evaluation Filler Material Studies Six alloys were tested following 52M [NX7206TK]: filler metal 82(A), 52MSS-E [HV1500], TG- SN690Nb [FBH2280], 52M [NX0T85TK], 52MSS-C [NX77W3UK], and 52i-B [187775]. Three 60

80 successful casts was produced at each pin length for comparison and a minimum of four for each length to be used as a ranking criteria. The results of the visual crack inspection are provided below, Table 7 and Figure

81 Table 7: Visual crack examination results for CPTT Ni-based weld filler metals. Alloy : FM 82 (A) 52i-B (187775) 52M (NX7206TK) TG-SN690Nb (FBH2280) 52MSS-E (HV1500) 52M (NX0T85TK) 52MSS-C (NX77W3UK) Cracking Cracking Cracking Cracking Cracking Cracking Cracking Pin Length (in) # Avg Max # Avg Max # Avg Max # Avg Max # Avg Max # Avg Max # Avg Max

82 Circumferential Cracking Circumferential Cracking Circumferential Cracking Circumferential Cracking Circumferential Cracking Circumferential Cracking Circumferential Cracking FM82 (A) 52i-B Pin Length (in) Pin Length (in) (a) 52M (NX7206TK) (b) TG-SN690Nb (FBH2280) Pin Length (in) Pin Length (in) (c) 52MSS-E (d) 52M (NX0T85TK) Pin Length (in) Pin Length (in) (e) (f) 52MSS-C Pin Length (in) (g) Figure 32: Scatter plot for CPTT results for tested alloys. (a) Filler metal 82(A) (b) 52i-B [187775] (c) 52M [NX7206TK] (d) TG-SN690Nb [FBH2280] (e) 52MSS-E [HV1500] (f) 52M [NX0T85TK] (g) 52MSS-C [NX77W3UK] 63

83 Pin Length (in) Each tested alloy has been ranked primarily by its maximum pin length with 0 cracking, Figure 33. Minimum pin length with 100 cracking, a ranking criteria used in the past, has been included as a secondary ranking criteria if seperation between alloys is not provided by the primary ranking criteria. Min Pin Length 100 Cracking Max Pin Length 0 Cracking FM82 52i-B (187775) 52M (NX7206TK) TG-SN690Nb (FBH2280) 52MSS-E (HV1500) 52M (NX0T85TK) MSS-C (NX77W3UK) Figure 33: Bar chart of visual ranking for tested alloys; Red bars represent the minimum pin length at which 100 circumferential cracking occurred. Grey bars represent the maximum pin length at which 0 cracking occurred. The results of the CPTT visual inspection of circumferential cracking from most resistant to solidification cracking to most susceptible rank the tested alloys as follows: Filler metal 82(A), 52i-B [187775], 52M [NX7206TK] and TG-SN690Nb [FBH2280] and 52MSS-E [HV1500], 52M [NX0T85TK], 52MSS-C [NX77W3UK]. The lower ranking for 52M [NX0T85TK] is due to the secondary criterion, minimum pin length with 100 cracking, which is used when 64

84 two rankings have an equal maximum pin length with 0 cracking. The same heats of filler metal 52M [NX7206TK] and 52MSS-C [NX77W3UK] were tested using a previous generation of the CPTT in addition to a different heat of filler metal 82. The results match the ranking obtained using the new generation CPTT with the same ranking criteria: Filler metal 82 < 52M < 52MSS-C (42). Dilution Studies Several of the tested alloys were diluted at several levels with cast stainless steel CF8A in order to determine the effect of dilution on solidification cracking susceptibility over the stainless steel nozzle component of the DMW described previously. Alloy 52M [NX7206TK], 52i-B [187775], and 52MSS-E [HV1500] were each diluted with 10 and 40 CF8A. Pin lengths in the vicinity of the threshold, or maximum pin length with 0 cracking, were tested in order to determine the role of dilution in cracking susceptibility. 52i-B [187775] Circumferential cracking CPTT results for the alloy 52i-B dilution series are provided in Table 8. The average, maximum, and standard deviation in circumferential cracking response is provided for pin lengths 0.75 inches to inches with a minimum of four pins at each length. 65

85 Maximum Circumferential Cracking Table 8: CPTT visual inspection results of circumferential cracking (52i-B dilution study). Alloy: 52i-B 52i-B + 10 CF8A 52i-B + 40 CF8A Pin Length (in) # Avg Cracking Max Std. Dev. # Avg Cracking Max Std. Dev. # Avg Cracking Max Std. Dev Figure 34 provides a graphical representation of the maximum circumferential cracking vs. pin length. 52i-B 52i-B+10 CF8A 52i-B+40 CF8A Pin Length (in) Figure 34: CPTT visual inspection results of circumferential cracking (52i-B dilution study). 66

86 According Table 8 and Figure 34 the dilution of alloy 52i-B with CF8A (from 0-40) results in an increase in solidification cracking susceptibility. In the case of 52i-B + 10 and 40 CF8A the primary ranking criteria (maximum pin length with 0 cracking) is not able to be determined due to a significant cracking response at the shortest pin length. As a result a relative solidification cracking susceptibility ranking can be determined by comparing the maximum circumferential cracking response of the dilution series across the threshold pin lengths (maximum pin length with 0 cracking and minimum pin length with 100, cracking, for the pure filler metal). 52MSS-E [HV1500] Visual CPTT results for the alloy 52MSS-E dilution series are provided in Table 9. The average, maximum, and standard deviation in circumferential cracking response is provided for pin lengths 0.75 inches to inches with a minimum of four pins at each length. Table 9: CPTT visual inspection results of circumferential cracking (52MSS-E dilution study). Alloy: 52MSS-E 52MSS-E + 10 CF8A 52MSS-E + 40 CF8A Pin Length # Avg Cracking Max Std. Dev. # Avg Cracking Max Std. Dev. # Avg Cracking Max Std. Dev

87 Maximum Circumferential Cracking Figure 35 provides a graphical representation of the maximum circumferential cracking vs. pin length. 52MSS-E 52MSS-E+10 CF8A 52MSS-E+40 CF8A Pin Length (in) Figure 35: CPTT visual inspection results (52MSS-E dilution study). Similar to 52i-B, visual circumferential cracking inspection of the 52MSS-E dilution series shows increasing dilution with CF8A (up to 40) results in increased solidification cracking susceptibility. The addition of 10 CF8A results in only a slight increase in cracking response (~7) while the addition of 40 CF8A results in a dramatic increase (>30). CPTT Tensile Testing Evaluation Since the inception of the test visual inspection of circumferential cracking has been the standard evaluation method for cast pin samples. Using the visual analysis of cracking 68

88 response in CPTT samples crack measurements can vary from person-to-person on the same cast pin. As a result users can often only utilize the data they themselves have generated in order to have an accurate and consistent comparison of cracking response between pins and alloys. So in addition to a visual examination, the pins in this study were subject to tensile testing in order to determine if it might be a valid and repeatable evaluation method. After being visually examined the pins are placed into the tensile jaw inserts, which have been designed such that the pin is gripped by the head and the foot, not the gauged section. Using a Gleeble 3800, thermo-mechanical simulator, the pins are pulled at a rate of 1 inch per second until the head and foot are separated (typically a stroke length of 2 inches). The reason for a fast stroke rate was due to the large number of samples which needed to be tensile tested. The extension rate was kept constant throughout testing in order to ensure comparability of results. From this testing the ultimate tensile strength (UTS) is determined for each pin. Once separated the surfaces of the pin are examined for voids or casting defects that may have reduced the measured UTS; if any are discovered the pin is rejected as an invalid sample. Once a series of pins has been tensile tested each pin is then represented by a UTS value rather than a circumferential cracking percentage. For initial evaluation of tensile testing on CPTT samples data generated by tensile pulling alloy 52i-B is examined. For each tensile test preformed a record of elongation, time, and force is maintained. Tensile testing force vs. elongation curves for several cast pins are provided below in Figure 36. Due to unequal tension on each sample when placed into 69

89 Force (lbf) the Gleeble initially for tensile testing the elongation, or position of the ram, was normalized to zero when positive force values began to be consistently recorded. 0, 0.75" 0, 0.75" 30, 1.25" 45, 1.5" 71, 1.375" Elongation (in) Figure 36: Tensile force vs. elongation, or ram position, during tensile testing of 52i-B CPTT samples. From Figure 36 it is clear that increases in observed cracking response result in a lower maximum tensile force required for failure. Both 0.75 inch pins with 0 observed circumferential cracking result in maximum force of roughly 3330 lbf, whereas the 30, 45, and 71 circumferentially cracked pins result in maximum forces of: 3174 lbf, 2992 lbf, and 2972 lbf respectively. For each tested pin the maximum achieved force is recorded. From these values the ultimate tensile strength (UTS) can be calculated for each pin length; by dividing the maximum tensile force by the cross-sectional area of an un-cracked cast pin ( in 2 ). In order to compare the results of the tensile testing 70

90 between alloys with different ultimate tensile strengths a baseline UTS must be established for each alloy in order to normalize the generated data for accurate comparison. The baseline UTS is determined by averaging the UTS values of all the 0 circumferentially cracked pins of a particular alloy. This baseline UTS value was then used to calculate the minimum (Equation 4), average (Equation 5), and maximum (Equation 6) reduction in UTS at each pin length. Equation 4: Calculation of minimum UTS reduction percentage using the maximum UTS per pin length. Equation 5: Calculation of average UTS reduction percentage using the average UTS per pin length. Equation 6: Calculation of maximum UTS reduction percentage using the minimum UTS per pin length. The tensile testing results for filler metal 52i-B [187775] are provided below in Table 10 as well as calculations using the equations above at each pin length. The average circumferential cracking, determined by visual circumferential cracking inspection, is also included in order to compare the tensile and visual evaluation methods. 71

91 Table 10: CPTT tensile inspection results for 52i-B [187775]. 52i-B [187775] Baseline UTS: 110 kip Pin Length (in) # Min. Tensile Force (lbf) Avg. Tensile Force (lbf) Max. Tensile Force (lbf) Min. UTS Reduction Avg. UTS Reduction Max. UTS Reduction Standard Deviation Avg. Circumferen tial Cracking Standard Deviation From Table 10, as pin length increases the reduction in UTS also increases. When comparing this evaluation method to the visual circumferential cracking inspection method it follows the same general trend as the measured circumferential cracking response. A scatter plot comparing the two evaluation methods for each pin is provided in Figure

92 Visual Inspection [Circumferential Cracking] Tensile Inspection [Reduction in ultimate tensile strength] Pin length (in) Figure 37: CPTT visual (red) and tensile (blue) evaluation method comparison for 52i-B [187775]. While there is some deviation in the values obtained between inspection methods it appears as though they follow the same trend. One reason for some of the observed discrepancies may be due to the inability for visual circumferential cracking inspection to account for crack depth. When a visual evaluation is preformed if a crack is visible at the surface of the pin it is included in the cracking measurement regardless of the depth that crack may penetrate beyond the surface. As a result cracks counted during visual inspection may penetrate only slightly beyond the pin surface, or all the way to the core. When the cast pins undergo tensile inspection, stress is concentrated in the pin where the cross-sectional area has been most reduced and the tensile strength is then a function of the remaining, un-cracked, cross-sectional area. These results indicate that the tensile inspection method offers a more comprehensive evaluation as to the magnitude of cracking in the cast pin samples. Alloy 52i-B, which has elevated levels of niobium, 73

93 would be expected to form a eutectic phase during the terminal stages of solidification giving it the propensity to backfill solidification cracks. As a result of backfilling one would expect the visual circumferential cracking inspection method, which lacks the ability to evaluate crack depth, to produce a larger and more dramatic increase in circumferential cracking response per pin length than the tensile inspection method. This expectation is confirmed in Figure 37 where the visual circumferential cracking inspection method identifies a more significant increase in cracking per pin length. Tensile testing of CPTT samples appears to be a viable evaluation method when compared to visual inspection, but there is a need to determine the criteria by which alloys will be ranked via tensile testing. Minimum, average, and maximum reduction in UTS were selected for evaluation in order to determine which would be utilized to develop a ranking criteria for solidification cracking susceptibility. The following section Filler Metal Studies contains the tensile testing results for the remaining filler metals, as well as the steps taken in developing a susceptibility ranking criteria for this pin evaluation method. Filler Material Studies The tensile inspection results for the remaining filler metals are provided in Table 11. Minimum, average, and maximum reduction in UTS have been calculated for each alloy at pin lengths ranging from 0.75 inches to 2 inches. 74

94 # # # Table 11: CPTT tensile inspection results Alloy: 52i-B (187775) 52MSS-E (HV1500) 52M (NX7206TK) TG-SN690Nb (FBH2280) 52MSS-C (NX77W3UK) Pin Length (in) Baseli ne UTS: Min. UTS Reduction 110 Kip Avg. UTS Reduction Max. UTS Reduction Std. Dev. Baselin e UTS: Min. UTS Reduction 107 Kip Avg. UTS Reduction Max. UTS Reduction Std. Dev. Ultimate Tensile Strength (UTS) Baselin e UTS: Min. UTS Reduction 99 Kip Avg. UTS Reduction Max. UTS Reduction Std. Dev. Baselin e UTS: # Min. UTS Reduction 96 Kip Avg. UTS Reduction Max. UTS Reduction Std. Dev. Baseli ne UTS: # Min. UTS Reduction 109 Kip Avg. UTS Reduction Max. UTS Reduction Std. Dev

95 Table 11 represents a smaller subset of pins than the visual circumferential cracking inspection table (Table 7) due to voids and defects discovered in some samples during tensile inspection. Standard deviation was calculated from the percentage reduction in UTS from the baseline UTS for each pin at a given pin length. At the lower pin lengths (less than inches), where the pin count is larger for most of the alloys, the standard deviation is fairly low. Due to the low standard deviation and repeatability at these lower pin lengths the focus in developing a ranking criterion for this evaluation method will be here ( ). The following figure provides a graphical representation of the maximum, average, and minimum reduction in UTS for pin lengths 0.75 inches to 1.25 inches (Figure 38). The average reduction in UTS across pin lengths 0.75 inches to 1.25 inches has been calculated and displayed for each alloys maximum, average, and minimum UTS reduction (orange bars). The average reduction across these pin lengths was included as a potential ranking criteria for this set of alloys. 76

96 Minimum Reduction in UTS Avgerage Reduction in UTS Maximum Reduction in UTS AVG (0.75"-1.25") MSS-E [HV1500] 52i-B [187775] TG-SN690Nb [FBH2280] (a) Max. UTS reduction M [NX7206TK] 52MSS-C [NX77W3UK] " 0.875" 1" 1.125" 1.25" AVG (0.75"-1.25") MSS-E [HV1500] 52i-B [187775] TG-SN690Nb [FBH2280] 52M [NX7206TK] 52MSS-C [NX77W3UK] (b) Avg. UTS reduction AVG (0.75"-1.25") TG-SN690Nb [FBH2280] MSS-E [HV1500] 52i-B [187775] 52M [NX7206TK] 52MSS-C [NX77W3UK] (c) Min. UTS reduction Figure 38: Bar charts of the maximum (a), average (b), and minimum (c) reduction in UTS for several filler metals at pin lengths 0.75 inches to 1.25 inches. The average of the UTS reduction over these pin lengths is labeled for each alloy (orange bar). 77

97 Using the average reduction in UTS over pin lengths 0.75 inches to 1.25 inches as the ranking criteria for susceptibility to solidification cracking the maximum and average reduction in UTS provide the same ranking: 52MSS-E [HV1500], 52i-B [187775], TG- SN690Nb [FBH2280], 52M [NX7206TK], 52MSS-C [NX77W3UK] (Figure 38(a)(b)). The minimum reduction in UTS provides a different ranking: TG-SN690Nb [FBH2280], 52MSS-E [HV1500], 52i-B [187775], 52M [NX7206TK], 52MSS-C [NX77W3UK] (Figure 38(c)). The only difference between the rankings produced using this method is that the minimum reduction in UTS moved TG-SN690Nb up, indicating an increased resistance to solidification cracking, compared to the other alloys. Taking a closer look at the minimum reduction in UTS; the average reduction in UTS across pin lengths, which was used as the ranking criteria, only varies from 1 to 5 with TG-SN690Nb, 52MSS-E, and 52i-B having average values of 1, 1, and 2 respectively (Figure 38(c)). When compared to the range of average UTS reduction values calculated for the maximum and average reduction, 4-29 and 2-11 respectively, the minimum UTS reduction does not provide adequate resolution between alloys using this ranking criterion to develop distinctive solidification cracking susceptibility rankings. Examining the maximum reduction in UTS; a large range of ranking criteria values are represented in this set of data (4-29) as well as a susceptibility ranking which matches the average UTS reduction ranking. When comparing the tensile testing evaluation to the visual evaluation the maximum reduction in UTS corresponds most closely to the maximum circumferential cracking utilized by the visual examination for ranking. This is due to the fact that the most significantly 78

98 circumferentially cracked pin at a particular pin length should have the most reduced cross-sectional area. This reduction in area requires a reduced tensile force for failure compared to other pins of the same length and results in the determination of the lowest UTS value, or the maximum reduction in UTS. As a result the solidification cracking susceptibility rankings generated using the tensile testing evaluation method rely on the maximum reduction in UTS determined at each pin length. Rankings are produced by averaging together the maximum reduction in UTS at 0.75, 0.875, 1, 1.125, and 1.25 pin lengths. This average reduction in UTS from will be referred to as the UTS threshold reduction. The filler metals have been ranked by their threshold UTS reduction and compared to the rankings generated via visual circumferential cracking inspection. Table 12: Comparison of solidification cracking susceptibility rankings determined using tensile inspection and visual inspection. Solidification Cracking Susceptibility Tensile Inspection [UTS Threshold Reduction] Visual Inspection Least 52MSS-E [4] 52i-B Susceptible 52i-B [6] TG-SN690Nb [20] 52M(NX7206TK),TG-SN690Nb, 52MSS-E Most Susceptible 52M(NX7206TK) [21] 52MSS-C [29] 52MSS-C The solidification cracking susceptibility rankings generated using the tensile inspection method from least to most susceptible are: 52MSS-E [HV1500] < 52i-B [187775] < TG- SN690Nb [FBH2280] < 52M [NX7206TK] < 52MSS-C [NX77W3UK]. These rankings generally 79

99 match those generated using the visual circumferential cracking inspection method (52i-B [187775] < 52M [NX7206TK], TG-SN690Nb [FBH2280], 52MSS-E [HV1500] < 52MSS-C [NX77W3UK]) with the exception of a more resistant ranking of 52MSS-E and slight separation in the ranking of TG-SN690Nb and 52M [NX7206TK] (20 and 21). According to tensile inspection 52MSS-E is as, or more, resistant to solidification cracking than alloy 52i-B having a UTS threshold reduction of 4 compared to 52i-Bs 6. Visual evaluation of circumferential cracking determined 52i-B to be more resistant to solidification cracking with a maximum pin length exhibiting 0 cracking of inches compared to a pin length of inches for 52MSS-E. This discrepancy in susceptibility rankings between the two methods of inspection may be attributed to the relatively large amount of eutectic formed in 52MSS-E (see Thermodynamic Simulation Modeling) and lack of crack depth consideration during visual circumferential cracking inspection. In order to further investigate the proposed ranking criteria of UTS threshold reduction an attempt was made to apply this concept to the circumferential cracking response determined by visual inspection. For pin lengths 0.75 inches through 1.25 inches the maximum circumferential cracking (MCC) response has been averaged in order to produce a criterion for ranking which mimics the UTS threshold reduction (Figure 39). 80

100 Maximum Circumferential Cracking AVG (0.75"-1.25") i-B [187775] 12 TG-SN690Nb [FBH2280] MSS-E [HV1500] 52M [NX7206TK] 52MSS-C [NX77W3UK] 51 Figure 39: Visual ranking of CPTT samples using the average of maximum circumferential cracking from 0.75 inch to 1.25 inch pin lengths (ranking criteria: orange bar). The ranking produced using the average of MCC from 0.75 inch to 1.25 inch pins is: 52i- B [187775] < 52M [NX7206TK], TG-SN690Nb [FBH2280], 52MSS-E [HV1500] < 52MSS-C [NX77W3UK] (Figure 39). This solidification cracking susceptibility ranking is identical to the ranking determined using the maximum pin length with 0 cracking and minimum pin length with 100 cracking. This indicates that there is a fundamental difference in the magnitude of cracking identified by the two different inspection methods and confirms that the ranking criteria itself (UTS threshold reduction) is not the reason for the difference in solidification cracking susceptibility rankings of the tested alloys. 81

101 # Dilution Studies 52i-B [187775] The tensile test results for this CPTT dilution set are provided in Table 13. Minimum, average, and maximum reduction in UTS is provided for each alloy and pin length. Table 13: CPTT tensile inspection results (52i-B dilution study). 52i-B (187775) 52i-B + 10 CF8A Baseline UTS: 110 Kip 52i-B + 40 CF8A Alloy: Min. UTS Reduction Avg. UTS Reduction Max. UTS Reduction Std. Dev. # Min. UTS Reduction Avg. UTS Reduction Max. UTS Reduction Std. Dev. # Min. UTS Reduction Avg. UTS Reduction Max. UTS Reduction Std. Dev The maximum reduction in UTS per pin length is used for ranking alloys evaluated by tensile testing, as determined in the previous section. Due to the limited pin lengths incorporated into this dilution study the UTS threshold reduction is calculated over pin lengths 0.75 inches to inches rather than the 0.75 inches to 1.25 inches used previously. Figure 40 provides a graphical representation of the maximum reduction in 82

102 Maximum Reduction in UTS UTS for pins 0.75 inches to inches in length as well as the labeled UTS threshold reduction (orange) " 0.875" 1" 1.125" AVG (0.75"-1.125") i-B (187775) 52i-B + 10 CF8A 52i-B + 40 CF8A Figure 40: Bar charts of the maximum reduction in UTS for the alloy 52i-B dilution series at pin lengths 0.75 inches to inches. The average of the UTS reduction over these pin lengths is labeled for each alloy (orange bar). 52i-B dilute with 10 cast stainless steel (CF8A) results in a 14 increase in UTS threshold reduction while a dilution level of 40 results in a 59 increase. This would indicate that increasing dilution of 52i-B by CF8A from 0-40 results in increasing susceptibility to solidification cracking. A similar trend was observed when the cast pins were evaluated visually for circumferential cracking. Similar to the UTS threshold reduction the average of MCC values from 0.75 inch to inch pins was calculated for this dilution series and is presented in Figure 41 below. Here 52i-B dilute with 10 83

103 Maximum Circumferential Cracking CF8A results in a MCC threshold increase of 20 and an increase of 61 at a dilution level of " 0.875" 1" 1.125" AVG (0.75"-1.125") i-B (187775) 52i-B + 10 CF8A 52i-B + 40 CF8A Figure 41: Visual ranking of alloy 52i-B dilution series CPTT samples using the average of maximum circumferential cracking from 0.75 inch to inch pin lengths (ranking criteria: orange bar). The dilution of alloy 52i-B with CF8A at levels of 10 and 40 results in an increase in solidification cracking susceptibility. For both inspection methods a level of 10 dilution results in only a slight increase in susceptibility while a dilution level of 40 results in roughly three times the increase when UTS threshold reduction and MCC threshold increase are considered. 84

104 52MSS-E [HV1500] Tensile inspection results for this dilution series are in agreement with the visual circumferential cracking inspection, showing an increase in cracking susceptibility with increasing dilution, Table 14. Minimum, average, and maximum UTS reduction values are provided for each alloy at pin length ranging from 0.75 inches to inches. Table 14: CPTT tensile inspection results (52MSS-E dilution study). 52MSS-E (HV1500) 52MSS-E + 10 CF8A 52MSS-E + 40 CF8A Baseline UTS: 107 Kip Alloy: # Min. UTS Reduction Avg. UTS Reduction Max. UTS Reduction Std. Dev. # Min. UTS Reduction Avg. UTS Reduction Max. UTS Reduction Std. Dev. # Min. UTS Reduction Avg. UTS Reduction Max. UTS Reduction Std. Dev The maximum reduction in UTS per pin length is provided in the figure below (Figure 42). Similar to the rankings determined for alloy 52i-B the UTS threshold reduction for the 52MSS-E dilution series is calculated as the average of the maximum UTS reduction for pin lengths 0.75 inches to inches (labeled and provided in orange). 85

105 Maximum Reduction in UTS 0.75" 0.875" 1" 1.125" AVG (0.75"-1.125") MSS-E (HV1500) 52MSS-E + 10 CF8A 52MSS-E + 40 CF8A Figure 42: Bar charts of the maximum reduction in UTS for the alloy 52MSS-E dilution series at pin lengths 0.75 inches to inches. The average of the UTS reduction over these pin lengths is labeled for each alloy (orange bar). Similar to the previous dilution series the addition of 10 CF8A results in a slight increase in solidification cracking susceptibility with an increase in UTS threshold reduction of 13. The addition of 40 CF8A results in a 66 increase in the UTS threshold reduction, five times the increase resulting from 10 dilution. For comparison to the results of the visual circumferential cracking inspection the MCC threshold increase has been calculated for this dilution series and is provided below in Figure 43 (labeled, orange bars). 86

106 Maximum Circumferential Cracking 0.75" 0.875" 1" 1.125" AVG (0.75"-1.125") MSS-E (HV1500) 52MSS-E + 10 CF8A 52MSS-E + 40 CF8A Figure 43: Visual ranking of alloy 52MSS-E dilution series CPTT samples using the average of maximum circumferential cracking from 0.75 inch to inch pin lengths (ranking criteria: orange bar). Visual circumferential cracking inspection of circumferential cracking determines that 52MSS-E dilute with 10 CF8A results in a 20 MCC threshold increase and 52MSS-E dilute with 40 CF8A a 62 increase. For the visual inspection MCC threshold increase for 40 dilution is three times the increase for 10 dilution, whereas for tensile inspection the UTS threshold reduction for 40 dilution is five times larger than 10 dilution. Both inspection methods show a slight increase in solidification cracking susceptibility when 52MSS-E is dilute with 10 CF8A and a dramatic increase when dilute with 40 CF8A. 87

107 Thermodynamic Simulation Modeling The results of the thermodynamic simulations, conducted using the Thermo-Calc software, are provided in the following section. The Schiel-Gulliver module within the Thermo-Calc software was used for all simulations, the specifics of which are provided in the experimental procedure. Discussion on solidification temperature ranges, partitioning coefficients, and secondary phases is provided in and attempt to explain the solidification cracking susceptibility rankings developed in the previous section (weldability testing). Filler Metal Simulation Results The results of the thermodynamic simulations run within the Schiel-Gulliver module of the Thermo-Calc software for the undiluted filler metals are provided below (Table 15). The partitioning coefficients (k) are calculated using the composition of the first solid to form (FCC_A1#1) divided by the nominal input composition. The other values such as liquidus temperature, solidus temperature, eutectic temperatures, and mole percent eutectic are obtained from the data output by the simulation. 88

108 Table 15: Thermo-CALC (Schiel-Gulliver) simulation results for undiluted filler metals. Elements that partition during solidification (k<1) are highlighted in red. Partitioning Coefficient (K) Alloy: Element FM82 52i-B [187775] 52MSS- E [HV1500] TG- SN690Nb [FHB2280] 52M [NX7206TK] 52M [NX0T85TK] 52MSS- C [NX77W3UK] Al B C Co Cr Cu Fe Mn Mo Nb Ni Si Ti Liquidus ( C) NbC Start ( C) Laves Start ( C) Solidus ( C) STR ( C) Mole NbC Mole Laves Mole NbC + Laves The elements with partitioning coefficients less than 1, meaning they segregate strongly to the liquid during solidification, have been highlighted in red (Table 15). 89

109 Temperature ( C) In order to shed some light on the data presented in the table above and its relation to the solidification cracking susceptibilities of these alloys they are presented in terms of their relative susceptibility to solidification cracking, increasing in susceptibility from left to right. Note that the ranking order in the table above corresponds to the visual ranking; this is due to the fact that more alloys were inspected visually than were tensile tested. As a result it is important to bear in mind that there is some question as to whether alloy 52MSS-E should be ranked more (determined via visually inspection) or less (determined via tensile inspection) susceptible to solidification cracking than alloy 52i-B. Figure 44 shows the solidification temperature range (STR) for each alloy as well as the NbC and/or Laves start temperature within this solidification temperature range. STR ( C) NbC [Start-Solidus] ( C) Laves [Start-Solidus] ( C) FM i-B [187775] MSS-E [HV1500] TG-SN690Nb [FHB2280] M [NX7206TK] M 52MSS-C [NX0T85TK] [NX77W3UK] Figure 44: Bar chart of Thermo-CALC (Schiel-Gulliver) simulation results. The NbC (red) and Laves (green) start temperatures are depicted within the solidification temperature range (blue). 90

110 It is clear from Figure 44 that no direct correlation can be drawn between solidification temperature range and solidification cracking susceptibility. It also does not appear as though NbC and laves solidification start temperature have any direct relationship on cracking susceptibility, although the temperature range over which only primary FCC solidification occurs may. If the NbC Solidification start temperature is subtracted from the STR for each alloy this temperature range can be obtained. The primary FCC STR divided by the entire STR is then the fraction of the STR in which only primary FCC solidification occurs. For filler metal 82 through 52MSS-C in the order presented above (Figure 44) the values for the primary FCC STR percentage of the entire STR are: 35, 40, 57, 68, 71, 68, and 63. These values follow the visual solidification cracking susceptibility rankings fairly well, bearing in mind that alloys 52MSS-E, TG- SN690Nb, and both heats of 52M share an equal ranking. If 52MSS-E is moved up in the solidification cracking resistance ranking, ahead of alloy 52i-B, as determined by the tensile testing inspection the primary FCC STR appears to have less of a direct relationship on solidification cracking susceptibility (Figure 45). 91

111 Percentage Primary FCC Solidification STR Primary FCC/STR FM i-B [187775] 57 52MSS-E [HV1500] 68 TG-SN690Nb [FHB2280] 71 52M [NX7206TK] M 52MSS-C [NX0T85TK] [NX77W3UK] Figure 45: Bar chart of Thermo-CALC simulation results showing the percentage of primary FCC solidification within the solidification temperature range. The solidification susceptibility ranking determined visually would suggest that a lower fraction of primary FCC solidification within the STR promotes resistance to solidification cracking. The tensile ranking of these alloys would move 52MSS-E (57 primary FCC) ahead of 52i-B (40 primary FCC) in solidification cracking resistance weakening the argument that a lower fraction of primary FCC solidification increases solidification cracking resistance. Figure 46 represents graphically the mole percent NbC and Laves phase formed by each alloy. This is meant to be used as an approximation of the relative amounts of fraction eutectic formed in each tested material. The labels in the figure represent total mole percent eutectic (NbC + Laves) depicted by the blue bar and the mole percent laves (green bar) when applicable, the mole percent NbC is unlabeled (red bar). 92

112 Mole NbC + Laves Mole NbC Mole Laves FM82 52i-B [187775] 52MSS-E [HV1500] TG-SN690Nb [FHB2280] 52M [NX7206TK] 52M [NX0T85TK] 52MSS-C [NX77W3UK] Figure 46: Bar chart of Thermo-CALC (Schiel-Gulliver) simulation results. Total mole percent eutectic (NbC + Laves) is depicted in blue (labeled) with NbC and Laves (labeled) depicted in red and green respectively. The results in Figure 46 are again presented in the ranking order determined visually, with solidification cracking susceptibility increasing from left to right. This ranking indicates that a larger amount of eutectic, or NbC and Laves phase, is most desirable in order to produce a material resistant to solidification cracking, with the exception of 52MSS-E and 52MSS-C. When the evaluating the relationship between total amount of eutectic and the rankings developed by tensile inspection 52MSS-E moves ahead of 52i- B in cracking resistance, leaving 52MSS-C as the only exception. When looking at the information provided in both Figure 44 and Figure 46 an explanation can be provided for the ranking order of solidification cracking susceptibilities for these alloys. First let s assume that a larger fraction, or mole, percentage eutectic results in an increased 93

113 resistance to solidification cracking due to the ability of the material to backfill cracks that form during solidification. According to this assumption both cast pin inspection methods have miss-ranked alloy 52MSS-C, which should be most resistant to solidification cracking, and the visual circumferential cracking inspection method has miss-ranked alloy 52MSS-E. While it is possible that 52MSS-E is more resistant to solidification cracking than 52i-B (as determined by tensile testing) but inspection methods show conclusively that 52MSS-C is the most susceptible among these alloys. In order to understand why 52MSS-C is the most susceptible to solidification cracking, while having the largest amount of eutectic, the composition of the eutectic must be investigated. Thermodynamic simulations predict the formation of NbC in each of the alloys studies and laves in alloys 52MSS-E, 52i-B, and 52MSS-C. Of these laves forming alloys 52i-B and 52MSS-E have been ranked on the upper end of solidification cracking resistance while 52MSS-C has been ranked as most susceptible. Taking a closer look at what comprises the eutectic constituent in these alloys the relative percentages of NbC and laves within the total eutectic of each alloy are examined (Table 16). Table 16: Thermo-Calc simulation results comparing eutectic components in laves forming alloys. Eutectic mole Alloy (NbC + NbC () Laves () Laves) 52i-B [187775] MSS-E [HV1500] MSS-C [NX77W3UK]

114 52i-B forms 0.64 mole eutectic; composed of 66 NbC and 34 laves phase, 52MSS- E forms 0.89 mole eutectic; composed of 35 NbC and 65 laves phase, 52MSS-C forms 1.94 mole eutectic; composed of 10 NbC and 90 laves phase (Figure 46, Table 16). From the data provided above it is clear that the ratio of NbC to laves phase in alloy 52MSS-C is very small (0.11) when compared to the same ratio in 52i-B (1.94) and 52MSS-E (0.53). The presence of such a significant amount of laves phase may explain the ranking of alloy 52MSS-C as most susceptible to solidification cracking. As indicated by Figure 44 Laves phase forms over a narrow temperature range during the terminal stages of solidification. This narrow solidification temperature range and the relatively large amount of eutectic have a negative impact on solidification cracking resistance in 52MSS-C. Alloy 52MSS-C forms 1.74 mole Laves phase which is more than triple that of any other alloy studied and 90 of the eutectic formed in this alloy. Alloys 52i-B and 52MSS-E both undergo the same solidification sequence as 52MSS-C although the initial amount of eutectic formation (NbC) is much greater in magnitude. This increased magnitude of solidification at higher temperatures may allow for crack formation to occur earlier in the solidification process in alloys 52i-B and 52MSS-E. As a result when these alloys reach the secondary eutectic solidification start temperature (laves solidification) There is sufficient liquid remaining to fill the cracks that formed earlier during the solidification process. For alloy 52MSS-C the initial amount of eutectic solidification (NbC) is small in comparison to the other alloys, this results in the persistence of a relatively large amount liquid into the terminal stages of solidification. This suppression in solidification may reduce to ability of 52MSS-C to form a coherent solid network at 95

115 elevated temperatures. As a result cracks will not form at elevated temperatures, rather the total strain accumulated during solidification will be concentrated in the large amount of liquid present at the terminal stage of solidification. If the strain accumulation exceeds the ductility of this liquid film solidification cracking will occur. Due to the rapid solidification of the laves phase over a very narrow temperature range the cracks formed during its solidification may not have sufficient time to backfill before the solidus temperature is reached. The Schiel simulation results for the undiluted filler metals indicate that a larger fraction, or mole percent, two component eutectic can be correlated with an increased resistance to solidification cracking. Alloys 82, 52MSS-E and 52i-B have been determined to be the most resistant to solidification cracking. The rankings developed using tensile testing cast pin inspection appears to provide a more objective and comprehensive analysis of solidification cracking. As a result the Solidification susceptibility rankings from most resistant to most susceptible are: Filler metal 82(A), 52MSS-E [HV1500], 52i-B [187775], TG- SN690Nb [FBH2280], 52M [NX7206TK], 52M [NX0T85TK], 52MSS-C [NX77W3UK] Dilution Simulation Results Scheil simulation results obtained using alloy 52i-B [187775] diluted with CF8A at 10, 25, and 40 are provided in Table 17. The partitioning coefficients (k) are calculated using the composition of the first solid to form (FCC_A1#1) divided by the nominal input composition. The other values such as liquidus temperature, solidus temperature, eutectic 96

116 temperatures, and mole percent eutectic are obtained from the data output by the simulation. Table 17: Thermo-CALC (Schiel-Gulliver) simulation results for 52i-B [187775] diluted with CF8A at 10, 25, and 40. Partitioning Coefficient (K) Alloy: Element 52i-B [187775] 52i-B [+10CF8A] 52i-B [+25CF8A] 52i-B [+40CF8A] Al B C Co Cr Cu Fe Mn Mo Nb Ni Si Ti Liquidus ( C) NbC Start ( C) Laves Start ( C) Solidus ( C) STR ( C) Mole NbC Mole Laves Mole NbC + Laves

117 Temperature ( C) Table 17 and the following figures the alloys have been presented in terms of increasing solidification cracking susceptibility from left to right, as determined by the CPTT. One important note is that 52i-B dilute with 25 CF8A was not evaluated using the CPTT, the results of the thermodynamic simulation have been included in order to provide insight into the dilution gap between 10 and 40 dilution. Figure 47 shows the solidification temperature range (STR) for each dilution level as well as the NbC and/or Laves start temperature within this range. 350 STR ( C) NbC [Start-Solidus] ( C) Laves [Start-Solidus] ( C) i-B [187775] i-B [+10CF8A] 52i-B [+25CF8A] 52i-B [+40CF8A] Figure 47: Bar chart of Thermo-CALC (Schiel-Gulliver) simulation results for 52i-B dilutions. The NbC (red) and Laves (green) start temperatures are depicted within the solidification temperature range (blue). Similar to Figure 44 here in Figure 47 no direct correlation exists between STR, NbC start temperature and solidification cracking susceptibility. The percentage of primary FCC solidification for this series, unlike the undiluted filler metals, does not show any 98

118 correlation to cracking susceptibility: 40, 38.5, 38.4, and 41 respectively. This is due to the relatively similar STR and NbC start temperatures for these compositions. Laves start temperature does seem to have a direct relationship with the solidification cracking susceptibility, with an increasing laves start temperature resulting in an increased susceptibility to solidification cracking (15 C-37 C). Figure 48 represents graphically the mole percent NbC and Laves phase formed by each dilution level. The total mole percent eutectic (NbC + Laves) is depicted by the blue bar the red (NbC) and green (Laves) bars show the composition of this total in respect to the two eutectic components. Mole NbC + Laves Mole Laves Mole NbC i-B [187775] 52i-B [+10CF8A] 52i-B [+25CF8A] 52i-B [+40CF8A] Figure 48: Bar chart of Thermo-CALC (Schiel-Gulliver) simulation results for 52i-B dilutions. Total mole percent eutectic (NbC + Laves) is depicted in blue with NbC and Laves depicted in red and green respectively. 99

119 Figure 48 shows very clearly that susceptibility to solidification cracking increases with the mole eutectic. When the composition of this eutectic is examined the NbC mole remains fairly constant across each composition ( ), while the mole laves steadily increase ( ). This would suggest that the increase in laves phase is responsible for the increase in susceptibility to solidification cracking. Taking a closer look at what comprises the eutectic constituent in these alloys the relative percentages of NbC and laves within the total eutectic of each alloy are examined (Table 18). Similar to the trend observed with the undiluted filler metals as the ratio of NbC/laves decreases solidification cracking susceptibility increases: 52i-B (1.94), 10 CF8A (0.92), 25 CF8A (0.56), 40 CF8A (0.54). Table 18: Thermo-Calc simulation results comparing alloy 52i-B dilution series eutectic components. Alloy Eutectic mole (NbC + Laves) NbC () Laves () 52i-B [187775] i-B + 10 CF8A 52i-B + 25 CF8A 52i-B + 40 CF8A Increasing dilution of 52i-B with stainless steel CF8A from 0 to 40 results in an increase in total mole eutectic, an increase in the laves fraction of this eutectic, and an increasing susceptibility to solidification cracking. When the partitioning coefficients are examined for this dilution series it is evident that increasing levels of dilution result in 100

120 increased partitioning of niobium, molybdenum, and titanium. Cast stainless steel (CF8A) is composed of 69.3 wt iron which is very different from alloy 52i-B which contains 2.55 wt. Work done by DuPont on the Ni-Fe-Nb system has shown that increasing levels of iron result in increased partitioning of niobium (43). Laves phase formation has been shown to be promoted by increased niobium partitioning in the absence of significant levers of carbon (17). As a result of the high iron content in CF8A, increasing dilution of 52i-B results in increasing amounts of Nb-rich low melting point eutectic which results in an increasing susceptibility to solidification cracking. Schiel simulation results obtained using alloy 52MSS-E [HV1500] diluted with CF8A at 10, 25, and 40 are provided in Table 19. The partitioning coefficients (k) are calculated using the composition of the first solid to form (FCC_A1#1) divided by the nominal input composition. The other values such as liquidus temperature, solidus temperature, eutectic temperatures, and mole percent eutectic are obtained from the data output by the simulation. 101

121 Table 19: Thermo-CALC (Schiel-Gulliver) simulation results for 52MSS-E [HV1500] diluted with CF8A at 10, 25, and 40. Partitioning Coefficient (K) Alloy: Element 52MSS-E [HV1500] 52MSS-E [+10CF8A] 52MSS-E [+25CF8A] 52MSS-E [+40CF8A] Al C Cr Fe Mn Mo Nb Ni Si Ti Liquidus ( C) NbC Start ( C) Laves Start ( C) Solidus ( C) STR ( C) Mole NbC Mole Laves Mole NbC + Laves Like the previous section the alloy 52MSS-E dilution series has been provided in order of solidification cracking susceptibility, increasing from left to right (52MSS-E dilute with 25 CF8A CPTT evaluation was not conducted). Figure 49 shows the solidification temperature range (STR) for each dilution level as well as the NbC and/or Laves start temperature within this range. 102

122 Temperature ( C) STR ( C) NbC [Start-Solidus] ( C) Laves [Start-Solidus] ( C) MSS-E [HV1500] 52MSS-E [+10CF8A] 52MSS-E [+25CF8A] 52MSS-E [+40CF8A] Figure 49: Bar chart of Thermo-CALC (Schiel-Gulliver) simulation results for 52MSS-E dilutions. The NbC (red) and Laves (green) start temperatures are depicted within the solidification temperature range (blue). Similar to 52i-B the modeling results for 52MSS-E, Figure 49, do not show a correlation between STR or NbC start temperature and solidification cracking susceptibility. Nor does the percentage of primary FCC solidification: 57, 53, 53, and 57 respectively. Increasing laves start temperature on the other hand does show good correlation to increasing susceptibility to solidification cracking, increasing from 13 C- 30 C. Figure 50 represents graphically the mole percent NbC and Laves phase formed by each dilution level. The total mole percent eutectic (NbC + Laves) is depicted by the blue bar the red (NbC) and green (Laves) bars show the composition of this total in respect to the two eutectic components. 103

123 Mole NbC + Laves Mole Laves Mole NbC MSS-E [HV1500] 52MSS-E [+10CF8A] 52MSS-E [+25CF8A] 52MSS-E [+40CF8A] Figure 50: Bar chart of Thermo-CALC (Schiel-Gulliver) simulation results for 52MSS-E dilutions. Total mole percent eutectic (NbC + Laves) is depicted in blue with NbC and Laves depicted in red and green respectively. 52MSS-E exhibits a trend very similar to 52i-B, Figure 48, as susceptibility to solidification cracking increases from left to right the mole eutectic increase. Within this eutectic the NbC mole remains almost constant ( ) while the laves phase mole increases consistently ( ) with increasing cracking susceptibility. This again suggests that an increase in the amount of laves phase formation may be responsible for this increased susceptibility to solidification cracking. Similar to previous discussion; as the ratio of NbC/laves decreases solidification cracking susceptibility increases: 52MSS-E (0.54), 10 CF8A (0.43), 25 CF8A (0.35), 40 CF8A (0.33) (Table 20). 104

124 Table 20: Thermo-Calc simulation results comparing alloy 52MSS-E dilution series eutectic components. Eutectic mole (NbC + Alloy NbC () Laves () Laves) 52MSS-E [HV1500] MSS-E + 10 CF8A 52MSS-E + 25 CF8A 52MSS-E + 40 CF8A The iron content of CF8A (69.3 wt)is incredibly high when compared to that of 52MSS-E (0.03 wt). Increasing dilution levels of 52MSS-E by CF8A leads to significant increases in iron content which, as discussed previously, results in increased niobium partitioning during solidification. Due to the low carbon content in 52MSS-E (and the other alloys studied) this increased partitioning of niobium results in low melting point laves phase formation, having a negative impact on solidification cracking resistance. The increase in laves phase formation across the filler metals and their series of dilutions can be explained by the niobium and iron content in each tested composition. In order to form laves phase in the Ni-Nb system sufficient Nb must be present with sufficiently low carbon, if too much carbon is present NbC formation will dominate (43). Each of the tested filler metals contained sufficiently low carbon for the formation of laves phase, but not all contained sufficient niobium. From Figure 46, laves phase formed in alloys; 52i-B, 52MSS-E, and 52MSS-C due to their elevated niobium contents (~2.5wt). With 52i-B and 52MSS-E forming a relatively small amount compared to 52MSS-C. This 105

125 information coupled with the knowledge that a larger amount of eutectic constituent can promote crack backfilling, or healing, provides sound justification for the solidification cracking susceptibility rankings determined using the tensile evaluation of the cast pin tear test samples. Table 21 shows the Nb, Mo, and Ti partitioning coefficients for each simulated composition as well as the iron and niobium wt. Table 21: Thermo-CALC (Schiel-Gulliver) simulation results for Niobium, Titanium, and Molybdenum partitioning coefficients (k) vs. iron (Fe) and niobium (Nb) content for all simulated compositions. Alloy Partitioning Coefficient (k) Niobium Content Iron Content K (Nb) K (Mo) K (Ti) Laves (Mole ) (wt Nb) (wt Fe) 52MSS-E [HV1500] i-B [187775] FM MSS-E [+10CF8A] TG-SN690Nb [FHB2280] M [NX7206TK] MSS-C [NX77W3UK] M [NX0T85TK] i-B [+10CF8A] M [+10CF8A] MSS-E [+25CF8A] i-B [+25CF8A] M [+25CF8A] MSS-E [+40CF8A] i-B [+40CF8A] M [+40CF8A] As iron content increases across the simulated compositions the partitioning coefficients for Nb, Mo, and Ti consistently decrease. For alloys where sufficient niobium is present this increased partitioning results in increased laves formation which typically has a negative impact on solidification cracking resistance. This is further reinforced by Figure 106

126 Partitioning Coefficient (K) 51, below, which is a plot of the iron content of all simulated compositions vs. the partitioning coefficient (K) for Nb, Ti, and Mo. K (Nb) K (Ti) K (Mo) Fe Content (weight ) Figure 51: Bar chart of Thermo-CALC (Schiel-Gulliver) simulation results for Niobium, Titanium, and Molybdenum partitioning coefficients (k) as a function of iron content (Fe) for all simulated compositions. It is clear from this figure that as Fe content increases the partitioning coefficient for these three elements decreases, meaning they segregate more strongly from the solid phase during solidification. This increased segregation results in the promotion of laves phase formation, which is shown very well in Figure 48 and Figure 50, where 52i-B and 52MSS-E (2.55 wt Fe, 0.03 wt Fe) are diluted with CF8A (69.30 wt Fe). Increasing dilutions result in an increasing Fe wt, a reduced niobium partitioning coefficient, an increased mole laves phase, and an increased susceptibility to solidification cracking. 107

127 Metallurgical Characterization The following section contains some general characterization of the materials included in this study. Fracture surfaces of 100 cracked cast pins are analyzed in order to confirm the presence of solidification cracking. EDS scans are conducted on both the bulk material as well as at solidification crack tips in order to verify material compositions and attempt to identify eutectic phase compositions. Fractography Analysis was performed on 100 cracked cast pins from each alloy to ensure the presence of solidification cracking,. The crack surfaces exhibit dendritic egg-crate morphology, which is associated with solidification cracking. This confirms the ability of the new generation CPTT to replicate conditions which result in solidification crack formation. 108

128 (a) Filler Metal 82(A) (b) 52i-B [187775] (c) 52M [NX7206TK] Figure 52: Fracture surface of 100 cracked pins for each tested alloy: (a) Filler Metal 82(A) (b) 52i-B [187775] (c) 52M [NX7206TK] (d) TG-SN690Nb [FBH2280] (e) 52MSS-E [HV1500] (f) 52M [NX0T85TK] (g) 52MSS-C [NX77W3UK]. 109

129 Figure 52 continued (d) TG-SN690Nb [FBH2280] (e) 52MSS-E [HV15000] (f) 52M [NX0T85TK] Continued 110

130 Figure 52 continued (g) 52MSS-C [NX77W3UK] Metallography Cross-sectioned cast pins were mounted, polished, and etched electrolytically with chromic acid (see Weldability Testing procedure). Microstructures of the as-cast pins are evaluated as well as the magnitude of observed backfilling in each alloy. The figure below (Figure 53) shows very clearly the difference in crack backfilling observed between the two extremes of the alloys studied: 52MSS-C and 52M. 111

131 (a) 52MSS-C [NX77W3UK] (b) 52M [NX0T85TK] Figure 53: Photomicrographs comparing crack "backfilling" in alloys (a) 52MSS-C and (b) 52M. Thermo-Calc simulation results predict the formation of 1.94 mole eutectic formation in 52MSS-C and only 0.17 mole in 52M (Figure 46).These predictions are in agreement with the observed relative magnitude of secondary phase(s) depicted in Figure 53. Simulation results indicate that niobium, molybdenum, and titanium segregate strongly during the solidification of these alloys resulting in the formation of eutectic 112

132 constituents. As a result 52MSS-C, which contains significantly more niobium (2.51 wt ) and molybdenum (3.51 wt ) than 52M (0.87 wt, 0.05 wt ), forms a much larger amount of eutectic than 52M. When comparing the solidification cracking susceptibilities of these alloys both the visual and tensile inspection methods determined 52MSS-C to be more susceptible to solidification cracking. Although some amount of eutectic can offer increased resistance to solidification cracking, due to the occurrence of crack backfilling, eutectic formation in 52MSS-C appears to be detrimental to its susceptibility. This increased susceptibility to solidification cracking in the presence of an increased amount of eutectic compared to the other alloys can be explained by the elevated iron content in tandem with the niobium and molybdenum additions. Thermodynamic simulation results show a reduction in partitioning coefficient (k) for both molybdenum and niobium with increasing iron content. The increased partitioning of these elements results in the predicted formation of predominately laves phase which leads to the increase in solidification cracking susceptibility seen in 52MSS-C. Micrographs depicting crack tips of each tested alloy are provided in Figure 54 below. Each image was taken at the same magnification in order to allow for visual comparison of the relative amounts of secondary phases present and any eutectic backfilling. 113

133 (a) 52MSS-C [NX77W3UK] (b) 52MSS-E [HV1500] (c) 52M [NX0T85TK] (d) TG-SN690Nb [FBH2280] (e) 52i-B [187775] (f) FM82 [A] Figure 54: Micrographs depicting the relative differences in the magnitude of crack "backfilling" among tested alloys. 114

134 From Figure 54 the relative magnitude of eutectic present at crack tips, determined visually, from the most to least is: 52MSS-C, 52MSS-E, FM82(A), 52i-B, TG-SN690Nb and 52M[NX0T85TK]. This ranking of the relative magnitude of eutectic present at crack tips is in agreement with the eutectic (mole NbC + Laves) amounts predicted by the thermodynamic simulation results. Compositional Analysis The following section contains information pertaining to the measured compositions of sectioned and mounted cast pin samples using an ESEM and EDS equipment. Bulk Composition A minimum of three EDS scan were conducted on cross-sectioned pins at locations far removed from cracking. The results of these scans are provided in Table 22 below. 115

135 Table 22: EDS compositional analysis vs nominal heat composition. Filler Metal EDS Comparison [wt] 52M [NX7206TK] 52M [NX0T85TK] 52MSS-C [NX77W3UK] TG-SN690Nb [FHB2280] 52i-B [187775] 52MSS-E [HV1500] FM82(A) Element Nom. Comp. Avg EDS Nom. Comp. Avg EDS Nom. Comp. Avg EDS Nom. Comp. Avg EDS Nom. Comp. Avg EDS Nom. Comp. Avg EDS Nom. Comp. Avg EDS 116 Nb Mo Ti max - Cr Fe max 1.1 Ni min

136 The nominal and measured compositional values are in close agreement. Slight variations exist; these are likely a result of the non-equilibrium cooling and subsequent segregation experienced by the pins and/or the error associated with elemental identification using EDS equipment. This analysis confirms that the composition of the material is generally maintained throughout the button melting and cast pin procedure. Composition of Crack Healing Constituents Cracked cast pins from each tested alloy were sectioned, mounted, polished and examined using an ESEM and EDS equipment. A multitude of spot scans were taken at the crack tips of each alloy in an attempt to identify secondary constituents and their compositions. The solidification sequence of nickel-base niobium-bearing alloys can be described generally as such: (1) Primary solidification L γ during this stage of solidification inter-dendritic liquid becomes enriched in niobium and carbon (2) Eutectic reaction L (γ+nbc) which depletes the interdendritic regions of Carbon and (3) Secondary eutectic reaction L (γ+laves) which occurs during the final stage of solidification (43). This solidification sequence generally describes the formation of the microstructures shown below (some alloys will not form laves phase due to low Nb content). Figure 55 depicts a crack tip in alloy 52M [NX7206TK] along a table of EDS spot scan results at several locations surrounding the crack tip. Any scan pointing to more than one location represents the average of two scans that revealed very similar compositions. 117

137 1 2 3 Figure 55: EDS analysis at solidification crack tip: 52M [NX7206TK], 15g, 1.875in, 65 circumferential cracking. The phases present at crack tips in alloy 52M are enriched in niobium, molybdenum, titanium, and manganese and slightly depleted of chromium, iron, and nickel. Based on the low partitioning coefficients determined for these elements (via thermodynamic simulation) it is no surprise that they have been found in higher concentrations, than the nominal composition, at this crack tip. According to thermodynamic simulation results only 0.16 mole NbC should form as a secondary phase in this alloy (Table 15). The spot scan results show a very large concentration of niobium in this region, 24.2 wt compared to 0.8 wt nominal, which indicates NbC formation may have occurred. The results of several EDS spot scans along a backfilled solidification crack in alloy 52i-B [187775] are presented in Figure 56. Any scan pointing to more than one location represents the average of two scans that revealed very similar compositions. 118

138 1 3 2 Figure 56: EDS analysis at solidification crack tip: 52i-B [187775], 13.5g, 1.5in, 72 circumferential cracking. Figure 56 shows a crack tip of alloy 52i-B enriched in niobium, titanium, and manganese and somewhat depleted of chromium, iron, and nickel, this matches the thermodynamic simulation results discussed previously. Partitioning coefficients for Nb, Ti, and Mn were found to be less than one; indicating strong segregation to the liquid, whereas Cr, Ni, and Fe were determined to be greater than one. In alloy 52i-B thermodynamic simulations predict the formation of both NbC and laves phase. From the EDS spot scans in Figure 56 three different compositions can be seen with varying levels of niobium, titanium, chromium, and manganese. It is possible that two of these compositions are representative of NbC and laves phase and that the third scan was taken in a region of compositional transition between the primary matrix and one of the eutectic constituent phases. 119

139 EDS analysis of a crack tip in alloy 52MSS-C revealed elevated levels of niobium and molybdenum with deficient levels of chromium, iron, and nickel (Figure 57). Any scan pointing to more than one location represents the average of two scans that revealed very similar compositions. The compositional analysis for 52MSS-C matched the predicted results of the thermodynamic simulation. 52MSS-C was predicted to form 1.94 mole eutectic, 90 of which being laves phase and 10 of which being NbC. Four scans conducted at this crack tip found a relatively uniform composition in the backfilled interdendritic region. 1 2 Figure 57: EDS analysis at solidification crack tip: 52MSS-C [NX77W3UK], 13g, 1.375in, 53 circumferential cracking. Since alloy 52MSS-C is predicted to form predominately laves phase as a eutectic constituent it can be concluded that the composition provided above in Figure 57 is relatively close to the composition of laves formation in this alloy. Unlike alloys 52i-B 120

140 and 52M, which form significant amounts of NbC, none of the scans on 52MSS-C reveal a niobium content of over 20 wt. Compositional analysis in alloy 82(A) is in general agreement with the thermodynamic simulation results revealing elevated levels of niobium, titanium, and manganese with somewhat depleted levels of chromium, iron, and nickel interdendritically at crack tips (Figure 58). Any scan pointing to more than one location represents the average of two scans that revealed very similar compositions Figure 58: EDS analysis at solidification crack tip: FM82 (A), 15g, 1.875in, 77 circumferential cracking. The results of EDS analysis here are also in agreement with the partitioning coefficient predictions made from the thermodynamic simulations. FM82 (A) is predicted to form only NbC which is likely near the composition of Scan 2. Scans 2 and 3 are likely scans taken in regions of compositional transition or on unpredicted laves phase formations in this alloy. 121

141 Alloy TG-SN690Nb shows enrichment at crack tips of niobium, titanium, and manganese with depletion of chromium, iron, and nickel similar to the other alloys (Figure 59). Any scan pointing to more than one location represents the average of two scans that revealed very similar compositions. 1 2 Figure 59: EDS analysis at solidification crack tip: TG-SN690Nb [FBH2280], 13g, 1.375in, 53 circumferential cracking. The results of the EDS analysis here are also in agreement with the partitioning coefficient predictions made via thermodynamic simulation. TG-SN690Nb is predicted to form a small amount of NbC, although unlike the other alloys predicted to form NbC 690Nb does not register an enrichment in niobium of the same magnitude (~20 wt ). This may be a result of the spot scan placement when evaluating this crack tip resulting in the identification of another unidentified eutectic constituent or a region of compositional transition. Alloy 52MSS-E reveals enrichment of niobium, titanium, and manganese and depletion of chromium, iron, and nickel at crack tips (Figure 59). EDS scans have identified several 122

142 distinct eutectic constituents with varying levels of niobium, titanium, and nickel enrichment Figure 60: EDS analysis at solidification crack tip: 52MSS-E [HV1500], 13.5g, 1.5in, 42 circumferential cracking. Thermodynamic simulations predict the formation of both NbC and laves phase in alloy 52MSS-E. Scan 2 shows potential in identifying NbC with scans 1 and 3 representing either laves phase, other eutectic constituents, or a region of compositional transition. It is evident that that the cracks in all the tested alloys propagate and in some cases heal interdendritically through Nb-rich and in the case of 52MSS Mo-rich eutectics. This is caused by the segregation of these elements during solidification; the amount of eutectic that forms during the solidification process varies based on alloy composition. One feature to note is the varying amounts of secondary phases present at the crack tips. Alloy 52MSS-C appears to have the greatest amount; this is due to the relative amounts of niobium and molybdenum in the alloys. 52M contains wt Nb and less than 0.05 wt Mo, these low alloying additions result in little segregation during solidification. 123