INVESTIGATION OF WELDABILITY IN HIGH-CR NI-BASE FILLER METALS THESIS

Transcription

1 INVESTIGATION OF WELDABILITY IN HIGH-CR NI-BASE FILLER METALS THESIS Presented in Partial Fulfillment of the Requirements for the Degree Master of Science in the Graduate School of The Ohio State University By Timothy Clark Luskin Graduate Program in Welding Engineering The Ohio State University 2013 Master s Thesis Committee: Professor John Lippold, Advisor Dr. Boian Alexandrov Professor Sudersanam Suresh Babu

2 Copyright by Timothy Clark Luskin 2013

3 ABSTRACT High-Cr Ni-base alloys have been used in the nuclear industry for weld overlay repair of dissimilar metal welds which have been affected by primary water stress corrosion cracking. During welding, defects may result from super-solidus and subsolidus cracking phenomena. The ability to evaluate and quantify the degree to which a given alloy is susceptible to each cracking mechanism is of great importance in alloy selection, alloy development and weld design. The first aim of this project was to improve the reliability of the cast pin tear test (CPTT) as a tool for evaluation of solidification cracking susceptibility. The second aim was to evaluate the weldability of a new heat of Inconel alloy 52MSS (ERNiCrFe-13) with reduced iron content as well as a single heat of Kobelco 690NB (ERNiCrFe-7A) comparing these materials with other available alloys of the same type. The incorporation of levitation melting has enabled precise control of the casting process in the CPTT. Active temperature monitoring during melting enables control of the casting temperature with a standard deviation of 8.8 C. Gradual reduction in the power applied to the coil has enabled focused transfer of the charge into the mold. ii

4 Thermocouple measurements indicated that the cooling rate during solidification in the CPTT was approximately twice that which is experienced in GTAW welding using parameters typical for use in weld overlay repairs. Finite element simulations indicated that a significant reduction in cooling rate of the casting could be obtained by altering the mold material, mold outer diameter or cast pin diameter. Prototype mold screening tests enabled design of a mold that provides a cooling rate of 160 C/s based on FEA modeling, compared with the previous cooling rate of 245 C/s, while maintaining feasibility of the CPTT. A longer pin length is required to cause cracking with the redesigned molds when compared with the previous molds. Also, significant crack healing with Nb-rich eutecticlike constituents was observed in pins cast from alloy 52MSS in the new molds. A decrease in strain rate most likely accompanies the decrease in cooling rate which would allow increased time for the flow of fluid to enable crack healing in alloys which exhibit a sufficient fraction eutectic. The solidification cracking susceptibility in the heat of alloy 52MSS with reduced iron content (52MSS-E) was evaluated and compared with alloy 690NB. 52MSS-E was determined to have a wider solidification temperature range based on thermodynamic modeling and single-sensor differential thermal analysis. This indicates that 52MSS-E may be more susceptible to solidification cracking than alloy 690NB. Transvarestraint testing showed that the maximum crack distance (MCD) was greater in 52MSS-E at 3%- 7% augmented strain. CPTT also indicated that alloy 52MSS-E is more susceptible to iii

5 solidification cracking than alloy 690NB. All tests, therefore, indicate that alloy 52MSS- E is more susceptible to solidification cracking than alloy 690NB. Alloys 690NB, 52MSS-E and 52i-B were evaluated based on their susceptibility to ductility dip cracking (DDC) using the strain to fracture test at 950 C. Alloys 690NB and 52MSS-E were found to have an equivalent threshold strain of 8% for DDC to occur, indicating similar resistance to DDC. Alloy 52i-B was found to have a threshold strain of 14% indicating superior resistance to DDC compared with 690NB and 52MSS-E. iv

6 DEDICATION To Erica, Carter and Seth. v

7 ACKNOWLEDGMENTS I would like to acknowledge and thank my advisor, Dr. Boian Alexandrov, for all of the guidance and advice that I have received from him in my research. I would also like to thank Professor John Lippold for being my academic advisor and Professor Suresh Babu for serving on my master s committee. I am indeed grateful for all of the members of the Welding and Joining Metallurgy Group at Ohio State for their friendship and enlightening discussion which has taught me so much during my time at OSU. I would like to acknowledge EPRI for sponsorship of this project. In particular, I would like to extend my gratitude to Steve McCracken for his advice and direction for my research. vi

8 VITA B.S. Mechanical Engineering, Brigham Young University, Provo, UT 2011-Present...Graduate Research Associate The Ohio State University, Columbus, OH FIELDS OF STUDY Major Field: Welding Engineering vii

9 TABLE OF CONTENTS Abstract... ii Dedication... v Acknowledgments... vi Vita... vii List of Tables... xii List of Figures... xv Chapter 1: Introduction... 1 Chapter 2: Background Introduction Welding Metallurgy Weld Regions Weld Metal Boundary Types Solidification Weldability... 9 viii

10 2.3.1 Solidification Cracking Ductility Dip Cracking Weldability Testing Cast Pin Tear Test Original Cast Pin Tear Test Threaded Mold Weldability Test Grooved Mold Test Casting Tests Determination of Strain in Weldability Tests Weldability of High-Cr Ni-Base Filler Metals Chapter 3: Objectives Cast Pin Tear Test Weldability Testing Chapter 4: Experimental Procedures Introduction Materials Weldability Testing Button Melting ix

11 4.3.2 Single Sensor Differential Thermal Analysis Cast Pin Tear Test Transvarestraint Test Strain to Fracture Test Metallography Sample Preparation Optical Microscopy Scanning Electron Microscopy Chapter 5: Development of the Third Generation Cast Pin Tear Test Introduction Apparatus Design Procedure Development Validation of Third Generation CPTT CPTT Automation Development CPTT Mold Optimization Cooling Rate Measurement in the Cast Pin Tear Test Cooling Rate Measurement in GTA Weld Pool FEA Modeling of the CPTT x

12 5.6.4 Determination of the Interfacial Heat Transfer Coefficient Effect of Mold Material on the Solidification Cooling Rate Effect of Mold Geometry on the Solidification Cooling Rate Optimization of the CPTT Mold Material and Geometry Prototype Mold Screening Chapter 6: High-Cr Ni-Base Filler Metal Weldability Thermodynamic Simulation Single-Sensor Differential Thermal Analysis Cast Pin Tear Test Transvarestraint Testing Strain to Fracture Testing Chapter 7: Summary and Conclusions Cast Pin Tear Test High-Cr Ni-base Filler Metal Weldability Chapter 8: Recommendations for Future Work References xi

13 LIST OF TABLES Table 1: Externally Loaded Solidification Cracking Test Characteristic Parameters Table 2: Grooved Mold Test Weighting Factors [34] Table 3: Chemical Compositions of High-Cr Filler Metals wt. % Table 4: Button Melting Parameters Table 5: Test parameters used for SS-DTA experiments Table 6: Sample mass used for Ni-Base materials for each pin length in the CPTT Table 7: Casting Temperature Used for CPTT of Various Nickel-base Welding Consumables Table 8: Transvarestraint Welding Parameters Table 9: Die block radius, strain and stroke used for transvarestraint testing with 0.25 in thick specimen Table 10: Parameters Used to Produce GTA Spot Weld on Strain to Fracture Samples. 58 xii

14 Table 11: Oxygen Level Readings in the Button Melting and CPTT Equipment Table 12: Test Material Composition (wt. %) Table 13: CPTT Validation Experiments Table 14: Casting Parameters Table 15: Chemical Composition of the Alloys used in Cooling Rate Measurements, wt.% Table 16: CPTT Test Parameters Table 17: Typical GTAW Procedure for Weld Overlay Repairs Table 18: Primary Dendrite Arm Spacing Measured in CPTT Pin Cast in C63000 Mold from Alloy 52M, Representative Bead on Plate Weld and Alloy 52M Transvarestraint Test Specimen Table 19: Results of C63000 Al-Bronze Mold Screening Tests Performed in Molds with 2 in length Table 20: Comparison of the Effect of Mold Material on Dendrite Arm Spacing in Pins Cast from Alloy 52MSS-C in the Third Generation CPTT Table 21: Calculated Solidification Temperature Ranges for Solid Solution Strengthened Nickel Based Filler Metals xiii

15 Table 22: Transformation Temperatures Determined Using SS-DTA for High-Cr Ni-Base Alloys during Solidification xiv

16 LIST OF FIGURES Figure 1: Weld Regions [5]... 5 Figure 2: Weld Metal Boundary Types [7]... 7 Figure 3: Effect of temperature gradient in the liquid and solidification rate on solidification mode [6]... 9 Figure 4: Illustration of factors controlling hot tearing according to the strain theory[12] Figure 5: Schematic Representation of the RDG Criterion for Hot Tearing Figure 6: Effect of grain boundary precipitation and tortuosity on DDC formation [27] 16 Figure 7: Schematic representation of transvarestraint test (left) and varestraint test (right) [33] Figure 8: Schematic of Programmable Deformation Cracking Test (PVR-Test) Procedure [31] Figure 9: Original CPTT Apparatus [4] xv

17 Figure 10: Original CPTT Molds [4] Figure 11: Threaded Copper Molds [35] Figure 12: Grooved Mold in Open Position [34] Figure 13: Button Melting Apparatus used for the Second Generation Cast Pin Tear Test Figure 14: Cast pin mold made of Cu-Be-Co alloy C17000 [37] Figure 15: Cast pins of the second generation CPTT. Arrows show the typical location of solidification cracks. [37] Figure 16: Quantitative criteria for evaluation of susceptibility to solidification cracking by CPTT. Circumferential cracking response curve in alloy René 77. The numbers on the plot indicate the number of samples tested for each mold length. [37] Figure 17: Hot Tearing Test Apparatus of Clyne and Davies [39] Figure 18: Self Restrained Casting Test Developed by Spittle and Cushway Figure 19: Experimental Setup for Acquisition of Thermal History in a Solidifying Spot Weld to be used for phase transformation analysis by SS-DTA Figure 20: Mold Geometry Used to Evaluate High-Cr Ni-Base Filler Materials Figure 21: Transvarestraint Testing Equipment xvi

18 Figure 22: Strain to Fracture Sample Preparation Figure 23: Weld Current During GTA Spot Welding of Strain to Fracture Samples Figure 24: Strain to Fracture Specimen After GTA Spot Welding [54] Figure 25: Third Generation CPTT Apparatus Design Figure 26: Custom Levitation Induction Coil Figure 27: Quartz Insert Placed inside Induction Coil to Hold Sample Prior to Levitation Figure 28: 1.75 in. C17000 Cu-Be CPTT Mold Half Used During Initial Investigation with the Third Generation CPTT Containing Cracked Pin of Material ERNiFeCr Figure 29: Thermal history in 10g sample of Alloy ERNiCrFe Figure 30: Heating Time for ERNiCrFe Figure 31: CPTT Pins in Filler Metal ERNiCrFe-13 Produced with the Third Generation CPTT Figure 32: Cracking Response in Filler Metal ERNiFeCr-13 Determined with the Third Generation CPTT Figure 33: Solidification Cracking Response in Filler Metal ERNiFeCr-13 Determined with the Second and Third Generation CPTT using C17000 Cu-Be Molds xvii

19 Figure 34: Surface Solidification Cracking in 1.75 in. Cast Pin of Alloy ERNiCrFe-13 with 100% Circumferential Cracking Produced with the Third Generation CPTT Figure 35:Dendritic Fracture Morphology of Surface Solidification Crack in 1.75 in. Cast Pin of Alloy ERNiCrFe-13 Produced with the Third Generation CPTT Figure 36: Fracture Morphology of Solidification Crack in 1.75 in Cast Pin of Alloy ERNiCrFe-13 Produced with the Third Generation CPTT Figure 37: Fracture Morphology of Solidification Crack in 1.75 in Cast Pin of Alloy ERNiCrFe-13 Produced with the Second Generation CPTT Figure 38: Cross Section of 1.75 in Cast Pin of Alloy ERNiCrFe-13 Produced with the Third Generation CPTT Figure 39: CPTT Control System Graphical User Interface Figure 40: Thermal History for Melting 9 Samples of Alloy 52MSS-C with Mass Ranging from 12-16g and Set Casting Temperature of 1500 C Figure 41: Second Generation CPTT Mold Design Figure 42: Thermocouple Locations in the CPTT Mold Figure 43: Measured Thermal Histories in the Cast Pin and at the External Mold Surface. Casting Temperature 1500 C, Cast Pin of Filler Metal 52M, Mold Material C xviii

20 Figure 44: Temperature Measurement in the Weld Pool of GTAW Cold Wire Process: a) Experimental Set-up, b) Weld Cross Section at the Thermocouple Location Figure 45: Comparison of Thermal Histories in CPTT and BOP GTA Welding Figure 46: Microstructure in Bead on Plate Representative GTA Weld Using Alloy 52m on Clad Stainless Steel Base Material Figure 47: FEA Models of the CPTT Mold Cast Pin Assembly (a) and of the Cast Pin (b) Figure 48: Temperature Dependent Material Properties of Filler Metal Inconel 52M Predicted Using JMatProTM Figure 49: FEA Predicted Solidification Sequence in a Pin of Filler Metal 52M Cast in Mold of Alloy C Figure 50: Microstructure in a Cast Pin of Filler Metal 52M: a) Upward Dendritic Solidification in the Cast Pin Foot, b) Transition from Fine Grained Chill Zone to Columnar-Dendritic Growth in the Cylindrical Part of the cast Pin Figure 51: a) Temperature Dependance of the IHTC between the Mold Wall of Alloy C1700 Mold and a Cast Pin of Filler Metal 52M Determined using Inverce Modeling in Pro-Cast TM, b) Validation of the IHTC by comparison of FEA Predicted and Measured Cooling Histories in Cast Pin of Filler Metal 52M Figure 52: Location of the FEA Node for Cooling Rate Predictions xix

21 Figure 54: Effects of the Mold Wall Thickness and Pin Diameter on the Solidification Cooling Rate in the CPTT Figure 55: Proposed Mold Design Figure 56: Comparison of the FEA Predicted Cooling Curve in a Cast Pin with Optimized Mold Design and the Measured Cooling Curve in BOP Weld Figure 57: CPTT Prototype Mold Designs 5 mm Inside Diameter (left) and 6 mm Inside Diameter (right) Figure 58: Microstructure of Pins Cast with the Third Generation CPTT in Alloy 52MSS- C A) C17000 Cu-Be Molds B) C63000 Al Bronze Molds Figure 59: Results of Initial Validation CPTT Validation of C63000 Al-Bronze Molds with Alloy 52MSS-C Figure 60: CPTT Cracking Response of IN52MSS-C in Redesigned molds and Previous Molds Figure 61: Crack Healing in Pin Cast from Alloy 52MSS-C in C63000 Al-Bronze Molds Figure 62: EDS Spectrum Comparing Interdendritic Region with Crack Healing (A) to the Dendrite Core (B) xx

22 Figure 63: Measured Thermal History of a Solidifying Weld Pool in Alloy 690NB Used for Single Sensor Differential Thermal Analysis Figure 64: Temperature Differential between Cooling Data and Calculated Reference Curve Used to Determine Solidus Temperature in Alloy 52MSS-E Figure 65: Solidification Temperature Range Determined Using Scheil Predictions and Single-Sensor Differential Thermal Analysis Figure 66: Maximum Circumferential Cracking Determined Using the Third Generation CPTT with C17000 Cu-Be Molds Figure 67: Maximum Circumferential Cracking for High-Cr Ni-Base Filler Materials Determined Previously Using the Second Generation CPTT [56] Figure 68: Comparison of Adjusted Maximum Pin Length without Cracking Determined using the Cast Pin Tear Test Figure 69: Solidification Cracking and Microstructure in CPTT Pins Cast in C17000 Cu- Be Molds A) Alloy 690NB B) Alloy 52MSS-E Figure 70: Fracture Surfaces on CPTT Pins cast in C63000 Al-Bronze Molds with 100% Circumferential Cracking Figure 71: Maximum Crack Distance in Alloys 52MSS-E and 690NB Determined in the Transvarestraint Test xxi

23 Figure 72: Solidification Cracking Which Occurred in the Transvarestraint Test at 5% Augmented Strain A) Alloy 52MSS-E B) Alloy 690NB Figure 73: DDC in Alloy 690NB Transvarestraint Specimen Tested at 5% Augmented Strain Figure 74: Maximum Crack Distance Determined Using the Transvarestraint Test for Several High-Cr Ni-Base Alloys [56] Figure 75: Maximum Crack Distance Determined in the Transvarestraint Test at 5% Strain Figure 76: DDC Cracking Response of High-Cr Ni-Base Filler Materials Determined Using the Strain to Fracture Test at 950 C Figure 77: Ductility Dip Cracking in Strain to Fracture Specimens Tested At 950 C A) 690NB, 9.9% Strain B) 52MSS-E, 8.0% Strain C) 52i-B, 15.7% Strain Figure 78: Comparison of Strain to Fracture Results with Several Previously Tested Alloys [28, 57] xxii

24 CHAPTER 1: INTRODUCTION During initial construction of many nuclear power plants, filler metal 82/182 was used for dissimilar metal welds for critical components of the reactor coolant system. [1] After years of service, these dissimilar metal welds have been affected by primary water stress corrosion cracking (PWSCC). This created a need for reliable repair procedures to keep the current fleet of plants in operation. Structural weld overlay repair techniques have been developed to mitigate the damage caused by PWSCC. Alloys with increased corrosion resistance are desirable for use in these weld overlay repairs. Nickel based filler metals with higher levels of Cr ( 30 wt%) are well suited for this application. The application of high Cr Ni-base filler metals in weld overlay repairs has been limited by several weldability issues. Several variants of filler metal 52 have been developed in an effort to reduce weld related defects. Alloys currently in use are susceptible to ductility dip cracking (DDC). [2] Efforts have been made to modify these alloys to form carbides at the end of solidification to allow cause boundary tortuosity which mitigates DDC. [3] A resulting increase in the solidification temperature range in these alloys has caused concern with solidification cracking. 1

25 The need to evaluate and quantify the susceptibility of alloys to various cracking phenomena has led to developments in the science of weldability testing. Researchers have developed many different weldability tests in an effort to compare and develop alloys. The cast pin tear test was first developed by Hull in the 1950 s to evaluate weld cracking in austenitic stainless steels. [4] This test is performed by casting pins with enlarged ends to anchor the pin during solidification. Solidification shrinkage and thermal contraction cause strains to build. The strain accumulates in the last remaining liquid to solidify. If the strain is sufficiently high, a solidification crack will form as it does in a typical weld. The first aim of this work is to improve the CPTT into a more viable test for weld solidification cracking. To achieve this goal, a new apparatus was designed and built for the CPTT with improved control of the casting process. A reliable procedure was developed for the CPTT which has enabled more definitive testing of alloys and testing of alloys which have been difficult to cast in the past. Finite element analysis of the casting process has also been done to evaluate the thermal conditions in the CPTT. The second aim of this work is to evaluate the weldability of several high-cr Nibase welding filler metals. Solidification temperature ranges were determined using thermodynamic simulation and single-sensor differential thermal analysis. Solidification 2

26 cracking susceptibility was evaluated using the transvarestraint test and the CPTT. Ductility dip cracking susceptibility was evaluated using the strain to fracture test. 3

27 CHAPTER 2: BACKGROUND 2.1 Introduction Nickel based alloys have been used for dissimilar metal welds between carbon steel pressure vessels and stainless steel piping in nuclear power plants. The use of Nialloys enables field construction without the need for post weld heat treatment. Stress corrosion cracking has become a problem in many of these dissimilar metal welds produced with Ni-based filler metal 82/182. In order to repair these welded joints and mitigate damage from stress corrosion cracking, weld overlay (WOL) techniques have been developed using variants of Ni-base filler metal 52 with higher Cr content to increase resistance to stress corrosion cracking. While several WOL repairs have completed successfully, weld cracking has presented a challenge in some situations. [1] In order to reduce the occurrence of such defects, it is necessary to have a fundamental understanding of the mechanisms which cause cracking to occur and a reliable way to measure cracking susceptibility in an alloy system. In this chapter, fundamental concepts related to the welding metallurgy of Ni-base alloys will be presented. The mechanisms which cause weld cracking will also be discussed and various methods of evaluating weldability of an alloy system will be described. 4

28 2.2 Welding Metallurgy This section contains the current understanding of fundamental concepts of welding metallurgy in austenitic alloys. This is extended into a discussion of weldability, being defined as the ability to produce welds free from defects. This background is necessary to understand the methods of weldability testing Weld Regions The microstructural regions in welds were investigated by Savage et al. [5] Terminology for several regions was proposed including the true heat-affected zone (HAZ), partially melted zone (PMZ) and the fusion zone (FZ), containing the unmixed zone and composite region. These regions are shown schematically in Figure 1. While some additional regions have been defined for specific alloy systems, these definitions are still valid today. Figure 1: Weld Regions [5] 5

29 The HAZ is the region of the base metal which does not melt, but microstructural changes do occur in the solid state. The PMZ is the transitional region between complete liquid and complete solid. Due to compositional variations within the material and metallurgical reactions between particles and the matrix, localized melting can occur below the equilibrium solidus. This further extends the PMZ into the base metal. The UMZ is a narrow region where the base material does not mix with the filler material due to fluid boundary layer existing in the weld pool. The composite region contains a uniform mixture of base and filler materials. [6] It is noteworthy that in dissimilar metal welds, a region of transition exists between the UMZ and composite region where a compositional gradient exists. This can lead to unique microstructural differences from the other weld regions Weld Metal Boundary Types The weldability of Ni based alloys is limited by cracking phenomena which occur along boundaries; therefore it is necessary to understand the different types of boundaries which form in austenitic weld metal during and following solidification. The three boundary types are the Solidification Grain Boundary (SGB), Solidification Subgrain Boundary (SSGB) and Migrated Grain Boundary (MGB). [7] These are shown schematically in Figure 2. 6

30 Figure 2: Weld Metal Boundary Types [7] Solidification during welding typically produces columnar or dendritic grains resulting from epitaxial nucleation of solid material from the base metal and subsequent competitive growth. [8] SSGBs occur in the regions between dendrites. The misorientation between dendrites across SSGBs is small due to the parallel growth directions of adjacent dendrites; therefore SSGBs are considered low angle boundaries. [6] SGBs generally occur along the centerline of welds due to impinging growth of dendrites from opposite directions. Since the dendrites are impinging from opposite directions, the crystallographic misorientation of the SGB can be significant. MGBs occur in the solid state due to the migration of SGBs at elevated temperatures. The driving force for grain boundary migration is the reduction of grain boundary free energy 7

31 by reduction of grain boundary area which is similar to the occurrence of grain growth. It has been reported that the crystallographic misorientation across MGBs is generally in excess of 30. [7] Solidification During solidification of metal alloys, segregation leads to the redistribution of solute. [9] Depending of the partition coefficient of each element in the alloy system, the concentration of that alloying element in the liquid will either be enriched or depleted. This leads to compositional variations in the FZ, particularly at SGBs and SSGBs. [6] Segregation occurs on both a microscopic and macroscopic level in weld solidification. [6] Segregation leading to compositional differences in the SGB is generally described by macroscopic solidification in which case solid diffusion is considered negligible and liquid diffusion is limited. Segregation between cells and dendrites is considered microscopic solidification in which case solid diffusion is also considered negligible, however complete liquid mixing is assumed. The distribution of solute from cell/dendrite cores to boundaries can therefore be described using the Scheil equation.[9] The morphology and size of grains within the FZ is controlled by solidification conditions. Several solidification modes exist due to the breakdown of the planar solidification front. This is caused by constitutional supercooling. [9] The solidification mode is controlled by the ratio of the temperature gradient in the solidifying weld metal 8

32 at the solid/liquid interface and the local solidification rate, as illustrated in Figure 3. [6] Since these parameters vary at different locations along the weld pool interface, it is common to see multiple solidification modes in the same weld generally progressing from planer at the outer edges of FZ to equiaxed dendritic in the center. Figure 3: Effect of temperature gradient in the liquid and solidification rate on solidification mode [6] 2.3 Weldability For this study, weldability will be considered as the ability to produce welds free of defects which result from various cracking phenomena. Several types of weld cracking can occur in Nickel based alloys including solidification cracking, liquation cracking and ductility dip cracking. [10] 9

33 2.3.1 Solidification Cracking Weld metal solidification cracking is a common weld defect in Ni-base alloys and other materials which solidify as austenite. This type of weld cracking occurs at temperatures above the effective solidus. Several theories have been developed in an effort to explain the mechanism behind solidification cracking. In essence, the various theories all describe an interlocking solid network separated by liquid films which are ruptured by tensile stresses. [8] These theories include the Shrinkage-Brittleness Theory [11], Strain Theory [12], Generalized Theory of Super-Solidus Cracking [13], Technological Strength Theory [14], Modified Generalized Theory [15] and RDG Criterion [16, 17]. The Shrinkage-Brittleness Theory was proposed by Bochvar [11]. This theory explains that there is some effective temperature interval for solidification cracking occurring at temperatures below a coherent temperature. When the material has reached the coherent temperature, a sufficient solid network has formed to prevent the flow of molten material. At temperatures above the coherency temperature, fluid flows freely to compensate for opening of the solid material resulting from shrinkage. The Strain Theory, proposed by Pellini, [12] explains that solidification of alloys concludes with a film stage. As in the Shrinkage-Brittleness Theory, cracking will not occur in the early, mushy stage of solidification. During the film stage, liquid films occur between solid grains at near solidus temperatures. Tensile strain resulting from solidification shrinkage will concentrate along the wet boundaries. Segregate liquid films 10

34 existing below the equilibrium solidus will increase the tendency for cracking to occur. The existence of a liquid film provides the condition where cracking may occur. The increase in temperature range and thus increased time in this susceptible state allows for increased strain to accumulate mechanically in the system. The two critical factors for hot tearing to occur are therefore the time that a solidifying material exists in the film stage of solidification and the rate of strain application as illustrated in Figure 4. Figure 4: Illustration of factors controlling hot tearing according to the strain theory[12] 11

35 Borland proposed the Generalized Theory of solidification cracking [13], which is a modification of the Shrinkage-Brittleness Theory. The generalized theory considers four stages of solidification. Stage 1 consists of primary dendrite formation with continuous liquid presence. At this point, solid and liquid phases are both capable of relative movement. During stage 2, dendrites interlock preventing relative solid motion; however liquid is still capable of movement and healing of cracks may occur. Stage 3 is the critical stage for solidification cracking because the solid network is too complex for fluid to flow and backfill cracks. Stage 4 is complete solidification where there is no liquid remaining in the system and therefore, solidification cracking cannot occur. Also, Borland proposed that if the liquid phase composition is much different than the solid, then wetting will not occur and cracking may be prevented. The Modified Generalized Theory was proposed by Matsuda et al based on actual observation of weld solidification cracking. [15] The stages of solidification are essentially the same as those proposed by Borland. The mushy stage, stage 1, of solidification was determined to occur over a small temperature range. Also, stage 3, where cracking may occur, was divided into a film stage and a droplet stage. During the film stage, liquid films allow both crack initiation and propagation. During the droplet stage, however, only crack propagation may occur. The Technological Strength Theory, proposed by Prokhorov, provides more insight into the mechanical aspects of solidification cracking. [14] The ductility of the solidifying weld metal is considered throughout the solidification range. The strain 12

36 accumulation is also considered throughout the solidification temperature range. If the strain exceeds the ductility of the material at any temperature, cracking will occur. Rappaz, Drezet, and Gremaud have developed a hot tearing criterion based on the ability of molten material to flow and fill interdendritic regions. [16, 17] This criterion, which has become known as the RDG criterion [18], is based on the level of resistance to mechanical deformation along wet boundaries being equal to the cavitation pressure. The cavitation pressure is a measurement of the pressure below which a void in the fluid will form which may subsequently develop into a solidification crack. If the pressure in the fluid remains above the cavitation pressure, void formation and thus crack formation cannot occur. The operating principle of the RDG criterion is represented schematically in Figure 5. Increase of the volume in interdentritic regions occurs as a result of solidification shrinkage and strain applied perpendicular to the growth direction. If the increased volume cannot be compensated by fluid flow because of the increasingly complex solid network, a low pressure region will form. The pressure cannot be reduced below the cavitation pressure and once this pressure is reached, void formation and possibly crack formation may occur. 13

37 Figure 5: Schematic Representation of the RDG Criterion for Hot Tearing Ductility Dip Cracking Ductility Dip Cracking (DDC) is a solid-state cracking phenomenon which has become a challenge to the welding industry with the increased use of Ni-base filler materials with high-cr. [10] Nickel based alloys have a dip in ductility through a range of 14

38 approximately 50-70% of the melting temperature. [19] Thermal strain which occurs following welding can exceed the ductility of the material within this temperature range resulting in DDC. Several different mechanisms have been proposed for DDC. [19-26] Ramirez and Lippold investigated the mechanism of DDC in Ni-base welding filler metals. [27] The dip in ductility and formation of DDC has been attributed to grain boundary sliding which is active within a range of sub-solidus intermediate temperatures. DDC occurs along the MGB. Therefore, the grain boundary tortuosity and formation of precipitates along grain boundaries affect the grain boundary migration and ultimately the DDC susceptibility of materials. This is represented schematically in Figure 6. Lippold and Nissley found that DDC susceptibility can be greatly reduced through the addition of Nb which forms precipitates at the end of solidification. [28] This is the strategy that is used in recently developed variants of filler metal 52. [3] 15

39 Figure 6: Effect of grain boundary precipitation and tortuosity on DDC formation [27] 16

40 2.3.3 Weldability Testing Much research has been done to develop methods to evaluate and quantify the susceptibility of alloys to various cracking phenomena. Many different tests have been developed to simulate, in a controlled way, the conditions which lead to defects. Some of the available weld cracking tests have been standardized and are well described in ISO [29-31]. Many variations of these tests and other tests exist and are generally performed using equipment and procedures developed within a specific lab. Weldability tests may be classified as externally-loaded or self-restrained. Externally loaded tests increase the level of strain within the weldment by applying an external tensile or bending load. In self-restrained tests, cracking is caused by strain which naturally accumulates due to solidification shrinkage and thermal contraction. Four types of externally loaded solidification cracking tests have been standardized in ISO : the Varestraint test, the Transvarestraint test, the flat tensile test (PVR test) and hot tensile test.[31] Other externally loaded tests are described in [18, 32]. The characteristic parameters of solidification cracking susceptibility determined by the standardized externally loaded tests are given in Table 1. 17

41 Table 1: Externally Loaded Solidification Cracking Test Characteristic Parameters Test Varestraint Transvarestraint PVR Test Hot Tensile Test Results Total Crack Length, Total number of Cracks, Maximum Crack Length, Solidification Cracking Temperature Range Total Crack Length, Total number of Cracks, Maximum Crack Length, Solidification Cracking Temperature Range Critical Strain Rate for Cracking BTR Varestraint, short for variable restraint, testing is performed by bending a test plate around a die while a weld is being performed. The varestraint, performed by bending perpendicular to the direction of travel, and transvarestraint, performed by bending along the weld, are represented schematically in Figure 7. Both versions of the varestraint test are typically performed with autogenous welds on specimens prepared from the material to be investigated. 18

42 Figure 7: Schematic representation of transvarestraint test (left) and varestraint test (right) [33] The augmented strain considered in the varestraint test is the level of strain present due to bending on the top surface of the test specimen. The level of strain (ε) controlled by the radius of the die block (r) and the sample thickness (t) as described by equation 1. (1) The test is repeated with a series of die blocks to vary the level of augmented strain. The specimen is inspected under magnifications of approximately 25x. Several results can be gathered from the versions of the varestraint test including Total Crack Length, Total number of Cracks, Maximum Crack Length, Solidification Cracking Temperature Range. 19

43 The programmable deformation cracking test, also known as the PVR test is a type of flat tensile test. In the PVR test, a flat test specimen is loaded in tension while a weld is performed on the top surface. The rate of deformation is increased linearly as shown in Figure 8. The specimen is examined for cracking under magnification of 25x and the critical tension speed, v cr, is determined where cracking first begins to appear as shown in Figure 8. This provides an indication of the critical strain rate for solidification cracking. Figure 8: Schematic of Programmable Deformation Cracking Test (PVR-Test) Procedure [31] 20

44 Various self-restrained tests have been developed. Three tests are described in ISO : T-joint weld cracking test, weld metal tensile test and longitudinal bend test.[30] The weld metal tensile test and longitudinal bend test are performed on previously welded specimens to inspect for cracks which occurred during welding. These standardized tests simulate actual welding conditions; however control of the level of strain in the weld metal is limited and the results are primarily qualitative. Many other variations of these tests have been developed which are described in [32]. 2.4 Cast Pin Tear Test First developed as a weldability test by Hull in the 1950 s, the CPTT is a solidification cracking test performed by casting small pins in rigid molds.[4] CPTT molds are designed with enlarged ends to limit shrinkage along the axis of the pin during solidification and cooling. This results in strain accumulation in the final solidifying liquid resulting from solidification shrinkage and thermal contraction. Some variations of this test have been used by other researchers to evaluate weld metal cracking.[34, 35] In recent years, the CPTT has been further developed at The Ohio State University.[36, 37] Original Cast Pin Tear Test The original cast pin tear test (CPTT) was developed by Hull in the 1950s and initially used for evaluation of the solidification cracking susceptibility in stainless steel weld metals [4]. This test was performed by levitation melting samples weighing 19g and casting them into tapered pins using a series of copper molds. 21

45 The testing apparatus consisted of a large chamber containing an induction coil, a series of molds and a manipulator, Figure 9. The chamber was evacuated and filled with helium or argon at a pressure of 1 atm. The induction coil was energized by a 10 kw industrial RF generator. Figure 9: Original CPTT Apparatus [4] Hull s original mold design consisted of a split mold, split restraining lock, bottom plate and pedestal, Figure 10. Hull found that the likelihood of cracking decreased as the pin length is decreased or the outer diameter of the mold is increased. In order to identify cracking susceptibility, pin length and mold outer diameter were varied, while 22

46 the volume of the mold was maintained at 2.4 cm 3. The tapered pin diameter was therefore increased by using a larger taper-pin reaming tool in pins of shorter length. The pedestal length was varied to maintain constant mold height. Molds were given arbitrary numbers based on geometry with higher numbers being less likely to form cracks. In order to quantify the results of the original CPTT, Hull devised a cracking index. The cracking index for a given pin was determined by systematically inspecting the pin surface under magnification of approximately 30 times. The pin was mounted in a device which allowed rotation and included a graduated dial allowing angular measurement of crack size. The cracking index was the sum of the angular measurement of all cracks taken as a percentage of the circumference up to 100%. 23

47 Figure 10: Original CPTT Molds [4] The original CPTT was used to test stainless steel alloys. In particular, the test found application in evaluation of the solidification cracking susceptibility in alloys of the Fe-Cr-Ni system. This was done by correlating the amount of ferrite present after solidification with the magnitude of solidification cracking which occurred. The effect of impurities on stainless steel solidification cracking was also studied with the test. The CPTT was also suggested for use in welding filler metal quality control Threaded Mold Weldability Test In the 1960s, Talento developed a modified version of the CPTT called the Threaded Mold Weldability Test [35]. Samples of material with 37 g mass were used. 24

48 The samples were levitation melted under argon in an apparatus, similar to that used by Hull [4]. This test utilized a threaded split copper mold of length 4.5 in, Figure 9. The threads had in. root (bottom) diameter and in. top diameter with 20 threads per inch. The thread pitch was considered analogous to the distance between points of restraint. Three different molds were developed with the same internal geometry and varying outer diameters to control the cooling rate during solidification. The mold which caused the lowest cooling rate was used for experimentation because when the test was performed with higher cooling rates, it was not successful in differentiating the solidification cracking susceptibility of different alloys. This mold design provided a great advantage in that each casting represented a full experiment set. The only need for casting multiple pins of the same material was to improve the statistical accuracy of the test data. Typically, thirty pins were cast with each material. 25

49 Figure 11: Threaded Copper Molds [35] Cracks were counted under magnification of seven times. No large cracks formed, but small fissures could be counted. The result of this test was the number of fissures counted in 30 cast pins. The threaded mold weldability test was used to evaluate the solidification cracking susceptibility of nickel-base alloys. In particular EN 82 (ErNiCr-3) welding filler metal diluted with Alloy 600 base metal was evaluated Grooved Mold Test In the 1970s, Armao and Yenicavich, developed another modification of the CPTT which utilized a grooved copper mold and samples with mass of 32 g [34]. This 26

50 test was performed with a melting apparatus similar to that used by Hull and was generally repeated three times for each material [4]. In this case, the mold length was approximately 5 in. The mold contained a single through hole of constant diameter with short circumferential grooves of larger diameter placed 1 in. apart to provide restraint, Figure 12. This mold geometry produced cast pins with smooth surfaces that were more easily inspected for cracks compared to the previous threaded mold design. The solidification cracking susceptibility was characterized using a fissuring index. The surface of the cast specimens was examined for cracks at 25 times magnification. The cracks were categorized by length and the number of cracks in each category was multiplied by a weighting factor, Table 2. The fissuring index was calculated as a sum of the number of cracks multiplied by the weighing factor in all categories. 27

51 Figure 12: Grooved Mold in Open Position [34] Table 2: Grooved Mold Test Weighting Factors [34] Category Size (in) Weighting Factor Small Medium Large Extra Large > The grooved mold test was also used to evaluate the solidification cracking susceptibility in Alloy 600 base metal and EN82 filler metal. In particular, different heats of material were compared as solidification cracking susceptibility in industry varied 28

52 widely with this material combination. The test was sensitive to variations in solidification cracking susceptibility between different heats of material. A second generation CPTT was developed at the Welding Engineering Laboratory of The Ohio State University in the 2000s [36-38]. The second generation CPTT utilized the original concept of Hull for causing solidification cracking in restrained cast pins by varying the magnitude of thermal strain during solidification. In this version of CPTT, a re-designed gas-tungsten arc button melting apparatus (Figure 13) is utilized to prepare button-shaped samples of the tested alloy and to cast these in copper mods. Figure 13: Button Melting Apparatus used for the Second Generation Cast Pin Tear Test 29

53 The pin casting procedure involves melting of a button over an open copper hearth. In this procedure, the molten charge is held in the open hearth by the surface tension of the molten alloy. The casting temperature is controlled by the level of superheating needed to reduce the surface tension of the tested alloy. The second generation CPTT utilizes split molds (Figure 14) that produce restrained cylindrical cast pins with a constant diameter of in, Figure 15. The mold outside diameter and the restraining pin head and foot geometry are held constant. The thermal strain applied to the pin is controlled by varying the mold / pin length between in. and in. The cooling rate is controlled by using nominally pure copper (C10100) or a copper-beryllium-cobalt alloy (C17000) as the mold material. The cast pins are examined for cracks using a binocular microscope at magnifications from 10x to 70x. The projected crack length in a plane that is perpendicular to the pin axis is measured in degrees and is converted into percentage of circumferential cracking, similar to the technique of Hull. Based on the CPTT results of particular alloy, a response curve of Maximum Circumferential Cracking (MCC) is plotted as a function of pin length, Figure 168. The maximum pin length of no cracking, the minimum pin length of 100% circumferential cracking, and the zero to 100% cracking range of pin lengths are used as quantitative criteria for ranking the susceptibility to solidification cracking. 30

54 Figure 14: Cast pin mold made of Cu-Be-Co alloy C17000 [37] Figure 15: Cast pins of the second generation CPTT. Arrows show the typical location of solidification cracks. [37] 31

55 Percent Circumferential Cracking 100% 90% 80% 70% Maximum circumferential cracking response curve Minimum pin length of 100% cracking 60% 50% 40% 0% to 100% cracking range 30% 20% Maximum pin length of no cracing 10% 1 0% Pin Length (in) mm Pin Length, in (mm) Figure 16: Quantitative criteria for evaluation of susceptibility to solidification cracking by CPTT. Circumferential cracking response curve in alloy René 77. The numbers on the plot indicate the number of samples tested for each mold length. [37] The CPTT has several advantages: the strain is nearly uniaxial in nature, the geometry is simple, the level of restraint can be well controlled by varying the length of the pin, and only a small amount of material is needed which can be produced as an experimental laboratory melt. 2.5 Casting Tests The development of the CPTT for weldability has many advantages as described previously. Hot tearing in castings is closely related to weld solidification cracking and may be described using many of the same theories. Many self-restrained casting tests 32

56 other than the CPTT have been developed and used to evaluate the susceptibility of alloys to hot tearing. While most of the investigations were involving aluminum alloys, the principles employed in the self-restrained casting tests are relevant to the development of the CPTT and could possibly be altered to evaluate weld metal. Pellini used a restrained mold to evaluate the temperature at which hot tearing occurred during casting of aluminum and steel alloys.[12] The mold used in this test consisted of a simple flat plate cast horizontally with flanged ends. Contraction of the flanges was limited by restraining bars and chills were placed against the flanges to accelerate cooling in these areas. Risers fed material to the center of the casting and this would be the final region to solidify. The temperature of the solidifying material was measured using a thermocouple. Radiographs were taken at different stages during solidification in order to identify the temperature at which hot tearing occurred. Clyne and Davies developed a quantitative solidification cracking test in which they used a novel resistance measurement technique to evaluate the extent of cracking.[39] Their casting apparatus, shown in Figure 17, utilized crucible melting and an imposed a temperature gradient in an open top mold through the use of resistance heating and water cooling. Restrained bars were cast in an open top steel mold and solidification terminated at the center due to heat extraction from the ends of the casting. Evaluation was done by measuring the resistance in the castings in both a non-cracked and cracked region. This enabled determination of the effective cross sectional area in the cracked region which occurred in the center of the cast bar due to the controlled heat 33

57 extraction. The resistance results were also compared with tensile test results where the reduction in tensile strength could be related to the reduction in effective cross section. Figure 17: Hot Tearing Test Apparatus of Clyne and Davies [39] Spittle and Cushway investigated the effect of superheat and grain structure on the hot cracking behavior in al-cu alloys by casting dog bone shaped samples.[40] In this work, the samples were cast horizontally with a feeder in the center of the mold, Figure 34

58 18. This mold was also designed to ensure that a feed supply was available to the hot spot in the casting, thus ensuring that cracking indeed occurs due to characteristics of solidification rather than a lack of available molten material to flow through the forming solid network. Copper chills were located at the ends of the bars with insulating sleeves over the central portions of the casting; solidification occurred progressively from the ends to the center of the casting. It was reported that solidification was unidirectional with columnar grains normal to the chill faces continuing to the center of the casting. Cracking was evaluated through the use of electrical resistance measurement as developed by Clyne et al. 35

59 Figure 18: Self Restrained Casting Test Developed by Spittle and Cushway. Countless hot tearing tests have been developed in an effort to quantify critical conditions for solidification cracking. While many of these tests have been developed using aluminum alloys, the principles of testing are relevant to the development of the cast pin tear test for weldability. Many other hot tearing tests for castings can be found in [41]. 2.6 Determination of Strain in Weldability Tests A challenge for researchers has been accurate quantification of the critical strain conditions which cause weld solidification cracking. For this reason, solidification cracking tests have been useful primarily as a means of comparison among materials. Externally restrained tests such as the Varestraint permit calculation of the average strain 36

60 imposed on the specimen surface during bending; however this is not an accurate measure of the local strain concentrated in the solid-liquid coexistent region.[42] Matsuda et al. developed the MISO technique which uses an in-situ optical measurement of the distance between features on the solidifying weld metal surface.[43] The change in distance between these features is used to determine strain as cracking occurs. Matsuda found that the minimum strain for solidification cracking to occur, as measured by MISO, was significantly higher than that determined by Varestraint testing, indicating that the strain induced by bending tends to concentrate along boundaries in the solidifying weld metal. Some researchers have turned to modeling to determine the local strain present in solidification cracking tests. Wei et al. modeled the Transvarestraint test using FEA.[42] Using this approach, the local strain at the trailing edge of the weld pool was calculated under various imposed bending strains. The local strain calculated using FEA was significantly higher than that calculated using traditional equations for the average strain on the sample surface. The simulated local strains were compared with MISO measurements and were in good agreement. It was proposed that Transvarestraint results be adjusted using the FEA model to account for local transverse strain. Shibahara et al. developed an FEA model with temperature dependent interface elements to predict solidification cracking in welds.[44] These elements are used between ordinary elements to account for high temperature brittleness. This is done by calculating 37

61 a temperature dependent critical stress based on the BTR and a scale parameter. The scale parameter is closely related to the minimum strain for solidification cracking to occur. The Transvarestraint and Houldcroft tests were modeled in order to determine these parameters. It was proposed that the scale parameter be used as a material constant representing solidification (hot) cracking susceptibility. Feng et al. modeled the Sigmajig weldability test using FEA in order to quantitatively evaluate the thermo-mechanical conditions leading to solidification cracking.[45] This model calculated the thermal and mechanical strain resulting from the solidification of the molten weld pool. Using simultaneous modeling and experiments, welding parameters and external loading of the test specimen were related to the stress distribution in the sample and ultimately to cracking behavior. 2.7 Weldability of High-Cr Ni-Base Filler Metals Primary water stress corrosion cracking (PWSCC) has degraded the structural integrity of dissimilar metal welds (DMW) between low alloy steel nozzles and stainless steel safe ends.[1] Originally, welds were performed using Inconel 82/182 (ERNiCr-3) with approximately 20 wt % Cr. Inconel alloy 52/152 has been developed with approximately 30 wt % Cr and has been shown to be far more resistant to PWSCC.[46] Several variants of Alloy 52 have been developed to solve DDC challenges which have occurred during welding. This has been done by alloying, particularly with Nb, to introduce carbides at the end of solidification which have been shown to reduce susceptibility to DDC.[27] While these alloys show great promise, wide solidification 38

62 temperature ranges due to eutectic phase formation pose a concern of solidification cracking.[10] Wu and Tsai investigated the hot cracking susceptibility of alloys 52 and 82 in Alloy 690 welds.[47] Differential thermal analysis and varestraint tests were performed on both alloys. It was determined that alloy 82 had a wider solidification temperature range than alloy 52. Also, alloy 82 was determined to be more susceptible to solidification cracking than alloy 52. Yushchenko et al performed a comparative analysis of welded joins in alloy 690 using Inconel 52 and 52MSS.[48] Gleeble-type hot tensile tests were performed on all weld metal specimens of both alloys at temperatures ranging from 20 to 1100 C. PVR tests were performed composite specimens with weld metal deposited into grooves in plates of alloy 690. The hot tensile tests revealed that the dip in high temperature ductility was greatly reduced in alloy 52MSS when compared with alloy 52. The PVR test results indicated that alloy 52MSS is more resistant to DDC than alloy 52, however alloy 52MSS is more susceptible to solidification cracking. Varestraint results also confirmed this finding. Cofie et al investigated the effectiveness of the application of a buffer layer prior to alloy 52M in weld overlay repair of alloy 82/182 weld in pressurized water reactors.[49] On-site testing revealed solidification cracking in the weld overlay applied on CF8M cast stainless steel piping. It was proposed that a buffer layer of ER308L be 39

63 applied to the cast material prior to the application of alloy 52M. 2 full scale mock-ups indicated that the buffer layer applied to the cast piping material eliminated solidification cracking in the weld overlay repair. It was noted that some small indications were present at the interface of the alloy 52M and the stainless steel which were likely DDC. McCracken and Smith evaluated the hot crack susceptibility of filler metal 52M on cast austenitic stainless steels in response to weld cracking during field application of a weld overlay repair.[1] Hot cracking occurred only in the first layer of alloy 52M which was applied over a buffer layer of ER308L. The cause for hot cracking was attributed to a high level of iron in the diluted alloy 52M which increased the amount of Nb to partition during solidification. This was believed to cause an increase in solidification temperature range and fraction γ/nbc eutectic which in turn led to solidification cracking. Alexandrov et al investigated the susceptibility to solidification cracking in High- Cr Ni-base alloys.[2, 36, 50] Alloys 82, 52M, 52MSS and 52i were evaluated. Alloy 52i is a modification of alloy 82 with improved corrosion resistance. It was determined that the solidification temperature range (STR) in Alloy 82 was wider than that in alloy 52i or 52MSS and alloy 52M had the narrowest STR. According to CPTT and transvarestraint tests, alloy 82 was less susceptible to solidification cracking than alloy 52M and alloy 52MSS was most susceptible. Alloy 52i was determined to be more susceptible to solidification cracking than alloy 52M and less susceptible than alloy 52MSS based on the transvarestraint test and cast pin tear test. One heat of alloy 52i demonstrated superior solidification cracking resistance in the CPTT 40

64 CHAPTER 3: OBJECTIVES 3.1 Cast Pin Tear Test The CPTT has shown promise in becoming a viable tool to evaluate the susceptibility of alloys to solidification cracking during welding as the mechanism for cracking is similar in both situations. The casting process, however, is not well controlled. The first aim of this work is to incorporate a controlled melting system into the CPTT to control the casting temperature and reduce process variability. The second aim in development of the CPTT is to optimize the mold design to more accurately reflect GTAW conditions. 1. Design third generation CPTT apparatus to incorporate levitation melting 2. Automate the melting process to ensure precise temperature control 3. Develop testing procedure for the third generation CPTT 4. Optimize mold design and select mold material to more accurately reflect the welding conditions experienced in structural weld overlay repairs 3.2 Weldability Testing Several variants of High-Cr Ni-Base filler metals have been developed to improve weldability. The objective in this weldability study is to evaluate the weldability of a low 41

65 iron heat of Inconel 52MSS-E (ERNiCrFe-13) as well as a heat of Kobelco 690NB (ERNiCrFe-7A). Susceptibility to solidification cracking and ductility dip cracking will be considered in both alloys. The susceptibility of alloy 52i to ductility dip cracking has not been studied and will be evaluated. The results will be compared with other alloys of the same type which have been tested previously. 1. Perform thermodynamic simulations to determine the solidification temperature ranges of 52MSS-E and 690NB 2. Perform single sensor thermodynamic analysis to experimentally determine solidification temperature ranges of 52MSS-E and 690NB 3. Perform transverse-varestraint testing of alloy 52MSS-E and 690NB to evaluate solidification cracking susceptibility 4. Perform cast pin tear testing on alloy 52MSS-E and 690NB to rank solidification cracking sensitivity 5. Perform strain to fracture testing on 52MSS-E, 690NB and 52i-B to evaluate susceptibility to ductility dip cracking 6. Compare weldability results between the alloys and with previous test results on alloys of the same type 42

66 CHAPTER 4: EXPERIMENTAL PROCEDURES 4.1 Introduction This chapter contains descriptions of tested materials and experimental methods used during this study. The experimental methods include a description of equipment and experimental procedures. Testing was aimed toward increasing understanding of the weldability of high-cr Ni-base welding consumables. 4.2 Materials During this study, the weldability of several heats of solid solution strengthened Ni-based alloys with high Cr concentration was evaluated. All of these materials were obtained commercially. Measured compositions for the filler metals which were tested are given in Table 3. 43

67 Table 3: Chemical Compositions of High-Cr Filler Metals wt. % Material Heat Al C Cr Fe Mn Mo Nb Ni P S Si Ti 52M NX0T85TK rem MSS-C NX77W3UK rem MSS-E HV rem i-B rem TG- SN690Nb FHB < rem Weldability Testing Several weldability tests were used to build a comprehensive understanding of the weld cracking susceptibility of some welding consumables. Phase transformation analysis was performed using button melting techniques with single-sensor differential thermal analysis (SS-DTA). Solidification cracking susceptibility was evaluated using the CPTT and transvarestraint tests. DDC susceptibility was evaluated using the strain to fracture (STF) and transvarestraint tests Button Melting Button shaped samples were used for phase transformation analysis with SS-DTA and in the CPTT. Buttons were prepared in a button melting apparatus, Figure 13. This apparatus consists of a cylindrical glass chamber, a water-cooled copper hearth and a GTA torch. The power used for button melting was provided by a Miller Dynasty 300 LX constant current GTAW power supply with high frequency arc initiation. 44

68 The hearth was polished using 800-grit sandpaper and cleaned with ethyl alcohol prior to button melting to avoid contamination of the sample. Welding wire was cleaned using ethyl alcohol, cut into 1 inch long segments and placed in the copper hearth. Nitrile gloves were worn while handling materials to avoid contamination. Prior to melting buttons, the chamber was purged with shielding gas using the following procedure: 1. Open the inlet valve to begin the flow of argon into the chamber 2. Adjust the gas flow rate to 20 cfh 3. Close the upper exhaust valve to build chamber pressure 4. Once the chamber pressure reaches 10 psi, open the upper exhaust valve 5. Repeat steps c and d 2 additional times 6. Repeat steps c and d 3 times releasing the pressure through the mold Thorough control of the atmosphere is accomplished by using high purity argon (99.998%). The oxygen concentration in the chamber atmosphere was measured after the purging cycle using a Purgeye 500 weld purge monitor produced by Huntingdon Fusion Techniques. This measurement was taken after each of 5 purge cycles and the average oxygen level was 11.8 ppm with a maximum of 16 ppm. The specific parameters used during button melting are given in Table 4. 45

69 Table 4: Button Melting Parameters Button Mass 1-20 g Shielding Gas Argon (99.998%) Gas Flow Rate Current Melt Time 20 CFH A 5-15 s Single Sensor Differential Thermal Analysis Phase transformation analysis was used for determination of the solidification temperature range which is related to the solidification cracking susceptibility of materials.[51] Phase transformation analysis was done by performing button melting experiments and analyzing thermal cooling data with SS-DTA. The use of the SS-DTA technique allows determination of the solidification temperature range in a solidifying weld pool under non-equilibrium conditions, similar to those experienced during actual welding. The experimental setup is shown in Figure 19. Both a 20 gram button and a 1 gram button of the test material were prepared. A hole was drilled in each 20 gram button using a 5/64 inch drill bit and countersunk on top at a 60 angle to a ¼ inch diameter. The 20 gram button was fastened to a copper platform. A type C thermocouple consisting of 46

70 0.005 inch diameter wire contained in a 4 passage ceramic insulating tube was inserted through the hole. The 1 gram button was placed in the countersunk region. SS DTA Software T R (t) T( S ) PC DAS T S (t) Figure 19: Experimental Setup for Acquisition of Thermal History in a Solidifying Spot Weld to be used for phase transformation analysis by SS-DTA A GTA torch is placed directly over the 1g button and a weld pool is made using the parameters shown in Table 5. Upon melting, the 1 gram button becomes part of a single weld pool within the 20 gram containing the thermocouple in the center. Timetemperature data was acquired at a sampling rate of 4 khz throughout the thermal cycle using an InstruNet model 100 data acquisition system. Three thermal cycles were recorded with each set of buttons and three sets of buttons were melted for each material. 47

71 Table 5: Test parameters used for SS-DTA experiments Weld Current (A) Weld Time (s) Arc Length (in) Shielding Gas Gas Flow Rate Ar (99.998%) 20 CFH The data acquired during the button melting experiments was analyzed using SS- DTA software which is capable of determining the temperatures where phase transformations occur. This is different from traditional differential thermal analysis (DTA). During DTA, a test specimen and a reference sample undergo a simultaneous heating or cooling cycle. The difference in temperature between the two samples provides an indication of latent heat either released or absorbed as a phase transformation occurs. Within the SS-DTA software, a reference curve is generated mathematically. As in DTA, deviations of the time-temperature data from the reference curve indicate a phase transformation. The liquidus and solidus temperatures were determined using this method. Also in alloys with reasonable fraction eutectic at the end of solidification, the start of the eutectic reaction can be identified Cast Pin Tear Test The cast pin tear test (CPTT) is used to evaluate the solidification cracking susceptibility of alloys. Details of the development of the third generation CPTT are contained in chapter four. A test alloy is cast into a cylindrical pin with enlarged ends to 48

72 anchor the casting during solidification, Figure 20. As the pin solidifies and cools, strain from solidification shrinkage and thermal contraction accumulate in the SGBs due to liquid films which occur in those regions. If the strain is sufficiently large, the SGB will be separated and a crack will form in the same manner that weld solidification cracks form. The pin length controls the level of thermal strain that accumulates during solidification. A series of pins in lengths from 1 inch to 2.5 inches are cast to identify the length of pin which causes cracking to occur in a given material. Figure 20: Mold Geometry Used to Evaluate High-Cr Ni-Base Filler Materials Buttons were prepared in the manner described previously. The mass of the button varies depending on pin length as given in Table 6. The button is placed inside the 49

73 coil and the chamber is sealed from the atmosphere. Shielding gas flows into the chamber and the chamber is purged using the following procedure: 1. Open the inlet valve to begin the flow of argon into the chamber 2. Open the upper exhaust valve 3. Adjust the gas flow rate to 30 cfh 4. Allow gas to purge for 30 s 5. Reduce the flow rate to 10 cfh 6. Close the upper exhaust valve to build chamber pressure 7. Once the chamber pressure reaches 3 psi, open the upper exhaust valve 8. Repeat steps c and d 2 additional times 9. Repeat steps c and d 3 times releasing the pressure through the lower exhaust valve 10. Adjust chamber pressure to 0.5 psi and flow rate to 5 cfh 11. Close lower exhaust valve Table 6: Sample mass used for Ni-Base materials for each pin length in the CPTT Pin Length (in) 50 Sample Mass (g)

74 Once the chamber has been purged, the coil is energized with a current of 430 A with a frequency of 230 khz to levitate and heat the sample. The temperature of the sample is monitored during the melting process. Temperature monitoring is accomplished by a 2 color optical pyrometer which produces a signal that is read by a National Instruments USB-6008 input/output device. This signal is used in a program produced in LabView to control the process. Once the sample reaches a set temperature, the current in the coil is decreased from 434 A to 50 A over the course of 1 sec. The reduction of current causes the sample to transfer directly into the mold in a controlled way. At the same time that casting is initiated, an electromagnetic valve is opened in the exhaust port at the bottom of the mold retainer to allow any gas contained in the mold to exit as the charge fills the mold cavity. The casting temperature set for each material is given in Table 7. Alloy 52MSS-C was tested at 1500 C during validation of the CPTT and these results were also used for comparison of materials. All other materials were tested at 100 C superheat. 51

75 Table 7: Casting Temperature Used for CPTT of Various Nickel-base Welding Consumables Casting Temp ( C) Liquidus ( C) 52M NB MSS-C MSS-E The cast pins were inspected for cracks under a binocular microscope at magnifications up to 80 times. This is done with the pin in a rotating fixture enabling measurement of the crack around the circumference in degrees. This value is taken as a percentage of the entire circumference and plotted vs. pin length as the circumferential cracking response curve. Pin surfaces are examined using a binocular microscope at 10 to 70 times magnification. The pin is held in a fixture which enables rotation and axial translation. Angular measurement is made possible by use of a dial which is part of the fixture. The angular length of each crack is measured. The percentage of circumferential cracking is calculated as shown in equation 2. % Cracking=100xLT/360, % (2) 52

76 where L T is the total length of all cracks on the pin surface measured in degrees. The maximum value calculated is 100%. The maximum circumferential cracking is plotted as a function of pin length. Alloys are ranked in terms of their susceptibility to solidification cracking as in the second generation CPTT. Ranking is done based on three criteria. These are the 1) maximum pin length without cracking, 2) the minimum pin length with cracking and 3) change in pin length during the transition from zero to 100% circumferential cracking Transvarestraint Test Transvarestraint testing was performed to determine solidification cracking susceptibly and also provide an indication of susceptibility to DDC. The transvarestraint test is performing an autogenous GTA weld on plates of the test alloy. During welding, the plate is bent over a die block to introduce strain into the sample. The transvarestraint machine used during this investigation is a custom design produced at The Ohio State University. This equipment consists of an automated GTAW system, a die block to support the sample and control the bend radius, bending rollers spaced 3.25 in apart to bend the specimen, Figure 21. This test is repeated with a series of die blocks which have varying radius to control the level of strain introduced into the sample. 53

77 Figure 21: Transvarestraint Testing Equipment The transvarestraint specimens were prepared by welding the test alloy in grooved plates of Inconel alloy 600 using GTAW. The plates were then machined to produce a 3 in x 8 in x 0.25 in specimen with three 1 inch wide sections of test alloy spaced 1 inch apart. Samples were cleaned thoroughly with ethyl alcohol prior to testing. To perform the test, the band of test alloy was centered over the die block and the bending roller was lowered onto the specimen surface to eliminate mechanical play in the system. A 2 inch autogenous GTA weld was performed with parameters given in Table 8. Parameters were determined from work done by Finton to optimize the transvarestraint test.[52] After the welding torch had traveled 1.5 inches, the rollers were lowered at a controlled rate of 2.5 in/s. The distance that the rollers were lowered depends on the radius of the die block. Die block radius, strain level and stroke are shown in Table 9. 54

78 Table 8: Transvarestraint Welding Parameters Weld Current Travel Speed Arc Length 180 A 5 in/min 0.08 in Table 9: Die block radius, strain and stroke used for transvarestraint testing with 0.25 in thick specimen Die Block Radius (in) Strain Stroke (in) % % % % % % % After performing the welds, the weld is examined under a Nikon binocular microscope with magnifications up to 80x to inspect for cracks. The distance from the instantaneous fusion boundary during bending to the end of the crack is measured. The 55

79 longest distance is taken as the MCD which is used as a comparative criterion for comparing alloys. The samples are also inspected for DDC Strain to Fracture Test The strain to fracture (STF) test was developed at The Ohio State University to evaluate the sensitivity of alloys to ductility dip cracking (DDC). [53] The STF is a Gleeble based test performed by heating a specimen of test material with a weld metal microstructure into the ductility dip temperature range and applying strain. The test is performed at various strain levels to determine the threshold level of strain for ductility dip cracking to occur. The standard specimen preparation procedure is shown in Figure 22. The specimens for these tests were prepared with a modified procedure as follows: 1) A ball mill was used to machine three 7/8 wide grooves in a 9 x16 x ½ Alloy 600 plate 2) The 600 plate was welded to a 2 CS strong back with two center bolts to keep it flat 3) TG-SN690NB, 52i, and lo fe 52MSS beads were then deposited in the grooves 4) Low dilution parameters and bead placement were used to minimize dilution 5) Alloy 600 plate with groove welds was cut off the strong back 6) Alloy 600 plate top surface was machined flat 7) Plate bottom surface was machined to achieve ¼ thickness 8) STF specimens were cut out by water jet 56

80 Figure 22: Strain to Fracture Sample Preparation Each STF specimen was placed in a copper fixture and a GTAW spot weld was performed in the center of the weld metal, Figure 24. The parameters used to produce the spot weld are given in Table 10. The current during application of the GTA spot weld is shown in Figure 23. This was done to produce a similar microstructure in each sample. The spot weld was surface ground using SiC metallographic paper from 180 to 800 grit. This was done to produce a smooth surface and enable inspection for cracking after the test was performed. Gauge marks were placed approximately 10 mm apart and the actual distance was measured using a microscope with image analysis. 57

81 Table 10: Parameters Used to Produce GTA Spot Weld on Strain to Fracture Samples Value Time Voltage 12.5 V Pre-Flow 20 CFH 10 s Start Level 20 A 0.1 S Initial Slope 5 s Weld 140 A 20 s Down Slope 12.7 s Final Level 20 A 0.1 s Post-Flow 20 CFH 15 s Arc Length Shielding Gas 0.08 in Argon 58

82 Figure 23: Weld Current During GTA Spot Welding of Strain to Fracture Samples Figure 24: Strain to Fracture Specimen After GTA Spot Welding [54] 59

83 Each sample was loaded into a Gleeble 3800 thermo-mechanical simulator. A type K thermocouple was welded to the gauge section for precise temperature control. The chamber was closed and a vacuum was pulled to 0.2 torr and purged with argon. This purge process was repeated a second time to reduce high temperature oxidation. The sample was resistively heated to 950 C and held for 10 seconds to ensure uniform temperature in the gauge section. The sample is then strained with a stroke rate of 0.6 mm/s. Once the sample has cooled, the gauge marks on the specimen were measured and the actual strain was calculated. Since gauge marks were placed on each edge, the average of the two strain measurements was used. The surface of the spot weld was inspected for cracks and the total number of cracks was recorded at each level of strain. 4.4 Metallography The techniques used to prepare and examine metallographic samples are presented in this section Sample Preparation Metallographic samples were prepared using the following procedure. Sections were cut from samples using an Allied Techcut 5 precision sectioning machine. The samples were mounted in Beuhler Konductomet conductive resin with a Leco PR-32 mounting press. Mounted samples were ground using 240, 400, 600 and 800 grit SiC metallographic paper. Following grinding with 800 grit paper, samples were polished using 9, 6, 3 and 1 µm diamond paste suspension. Between each stage of grinding and 60

84 polishing, samples were cleaned in an ultrasonic bath of ethyl alcohol. Final polishing was performed by placing the samples on a Buehler Vibromet 2 vibratory polisher using non crystallizing 0.02 µm colloidal silica suspension for 4 hours. Sample etching was performed using 10 vol. % CrO 3 solution. The samples were etched electrolytically by the following procedure. The sample was placed in a glass dish and submerged in the etchant solution. A tungsten cathode was placed on the sample surface and a stainless steel anode was placed in solution. A constant voltage DC power supply set at 5 volts was connected to the cell. The samples were etched for 5-10 seconds. After etching, the samples were rinsed in a bath of water, followed by a bath of ethyl alcohol and a final ultrasonic bath of ethyl alcohol. The samples were rinsed with ethyl alcohol and dried with hot air Optical Microscopy Samples were evaluated using two types of light optical microscopes (LOM). A Nikon SMZ1000 stereoscope with a Nikon digital camera was used to evaluate cracking in weldability test specimens. An Olympus GX51 with image acquisition capabilities was used to examine samples which were prepared metallographically Scanning Electron Microscopy Two scanning electron microscopes were used for analysis. The first is a Quanta 200 general purpose SEM with tungsten source. The second is a XL-30F ESEM field emission gun SEM. Both microscopes were used with secondary electron, backscatter 61

85 electron (BSE) and energy dispersive spectroscopy (EDS) detectors. Secondary electron mode was used for fractography as well as examination of etched samples. BSE was used to examine un-etched samples. The working distance was typically 10mm. The beam accelerating voltage was kev with spot size of 4-5. EDS was used to qualitatively evaluate the elemental segregation to solidification grain boundaries. 62

86 CHAPTER 5: DEVELOPMENT OF THE THIRD GENERATION CAST PIN TEAR TEST 5.1 Introduction The second generation cast pin tear test (CPTT) was developed at Ohio State University in order to evaluate the susceptibility of materials to weld solidification cracking. [36] The CPTT is performed by casting a series of cylindrical pins with enlarged ends to prevent overall contraction of the pin during solidification. Strain in the pin is caused by solidification shrinkage and thermal contraction. The thermal strain accumulates along boundaries in the final liquid to solidify. Longer pin length leads to an increase in the amount of thermal strain accumulation. When strain is sufficient to overcome the ductility of the solidifying metal, solidification cracking will occur. The length of pin which causes cracking in different materials can be compared to evaluate their relative solidification cracking susceptibility. This test has been successful in determining the relative susceptibility of cracking in weld filler metals. [50] Also, since the samples used in the CPTT are small ( 14g) laboratory melts, experimental alloys may be evaluated without need for producing bulk amounts of material. While this test has been used successfully, there are some drawbacks to the second generation CPTT. Melting of the charge is performed with GTA 63

87 melting over an open hearth. When the sample is molten and the surface temperature is sufficiently low, the sample falls through the hearth into an open-top mold. This made control of casting temperature and molten metal transfer challenging. The objectives in developing a third generation cast pin tear test are to incorporate levitation induction melting to control the casting temperature and transfer of molten material and to design molds which will more accurately reflect actual welding conditions. The design of the third generation CPTT apparatus and molds is described in this chapter. 5.2 Apparatus Design The third generation CPTT incorporates levitation induction melting, Figure 25. The apparatus consists of a chamber containing a levitation coil with a mold situated directly below the coil. A two color optical pyrometer is installed above the coil. 64

88 Figure 25: Third Generation CPTT Apparatus Design The induction coil for the third generation CPTT, Figure 26, has been designed and fabricated specifically for this application. The coil is fabricated from rectangular copper tubing. It consists of seven turns in one direction and two turns in the opposite 65

89 direction to maintain stability of the levitating sample. It is powered by a 10 kw industrial induction power supply and work head. The system operates at 224 khz with this particular coil. Within the coil is a quartz insert to isolate the button sample from the coil prior to levitation, Figure 27. Figure 26: Custom Levitation Induction Coil 66

90 Figure 27: Quartz Insert Placed inside Induction Coil to Hold Sample Prior to Levitation An optical pyrometer measures the sample temperature from above during melting. The ability to superheat materials and control casting temperature is of great importance. Once the sample reaches the set temperature, the casting process in initiated. Also, the thermal history of each sample on heating is recorded for later analysis if necessary. Melting is performed inside a sealed chamber that consists of a glass cylinder with a copper plate on bottom and Teflon plate on top. The mold is located directly below and in close proximity to the bottom of the coil. The molten charge is therefore able to transfer directly into the mold. An opening in the top of the chamber allows the sample to be loaded. The mold is removable from the bottom. This allows testing to be performed 67

91 without complete disassembly of the chamber, as was required in the second generation CPTT. A series of molds are used with varying length from in to 2.25 in, constant pin shaft diameter of in and outer diameter of in, Figure 28. Head and foot geometry was also held constant. The mold alloy used during initial investigation of the CPTT system was copper-beryllium-cobalt alloy (C17000) with thermal conductivity of 118 W/m-K. Figure 28: 1.75 in. C17000 Cu-Be CPTT Mold Half Used During Initial Investigation with the Third Generation CPTT Containing Cracked Pin of Material ERNiFeCr-13 Shielding gas flows into the chamber through a port in the bottom plate. Exhaust ports are placed in the top plate and below the mold. Exhaust ports are valved to allow redirection of the shielding gas for purging of the chamber and for casting. 68

92 5.3 Procedure Development Prior to heating the sample, the melting chamber must be purged to eliminate any existing air. Atmosphere quality is extremely important to avoid oxidation of the sample. Purging is performed by allowing gas to flow through the chamber exhausting from the top at 30 cfh for 30 seconds then pressurizing the chamber to 4 psi and releasing three times through the top vent and three times through the bottom vent. The gas used for purging the chamber is pure argon with high purity (99.998%). The purging process was validated by measuring the oxygen level in the exhausted gas to ensure adequate purity. Measurements were performed using a Purgeye 500 weld purge monitor produced by Huntingdon Fusion Techniques. The oxygen level was taken after each of 5 purging cycles. The same procedure was used to evaluate the oxygen level in the button melting apparatus. The oxygen readings for both pieces of equipment are given in Table 11. The average and maximum oxygen level are higher in the equipment used to prepare the samples for CPTT than they are in the CPTT apparatus indicating that the purge procedure provides an atmosphere of sufficient quality. Table 11: Oxygen Level Readings in the Button Melting and CPTT Equipment Button Melter CPTT Apparatus Average O x Level (PPM) Max O x Level (PPM)

93 It was found that instantaneously energizing the coil would cause the sample to lift out of the coil due to rapid acceleration. It is necessary to gradually energize to the coil to ensure stable levitation. Energizing the coil to the heating power of 434 A linearly over the course of 0.5 seconds was found to suffice. During heating and melting of the sample, the current is held at 434 A, as recommended by the manufacturer of the power supply and the coil. The power output of the power supply at this setting with a typical sample is approximately 9 kw. Initially, the time required to heat a sample to the desired temperature was determined experimentally. The charge was heated for a set amount of time and then the casting process was initiated. Later development of automation enabled a closed loop control system which initiates casting when the set temperature is reached. Typical thermal histories during heating in the levitation coil are shown in Figure 29. The rate of heating is approximately linear in the solid state. The temperature is constant when the charge is semi-solid due to the latent heat being consumed during the phase transformation. A linear increase in temperature occurs when the sample is completely molten. 70

94 Figure 29: Thermal history in 10g sample of Alloy ERNiCrFe-13 The time required for melting and heating of samples with different mass is shown in Figure 5. It is interesting to note that the heating time decreases as the sample mass increases. This effect has been attributed to the location of the sample in the levitation coil during melting. [55] Due to lower gravitational force, samples with lower mass will tend to levitate higher in the coil where the heating effect is not as strong. 71

95 Figure 30: Heating Time for ERNiCrFe-13 Once the casting process is initiated, the coil is de-energized and will no longer support the molten charge. The liquid metal is transferred to the mold by gravity. Some gas flow is necessary through the bottom of the mold to enable evacuation of residual gas and complete mold filling. It was found that instantaneously de-energizing the coil led to splashing of the charge on the induction coil and out of the mold. A gradual reduction in power focusses the molten charge to transfer in a stream-like manner. This reduces contact with the induction coil and directing transfer into the mold. A linear reduction in current applied to the coil from 434 A to 50 A over the course of 1.0 sec was found to be ideal. 72

96 5.4 Validation of Third Generation CPTT The third generation CPTT has been validated by testing a single heat of high-cr Ni-base filler metal ERNiFeCr-13 (INCONEL 52MSS). The purpose of this validation test is to ensure that the length of pin controls the cracking behavior in the third generation CPTT. The measured composition of this filler metal is shown in Table 12. The validation experiments are summarized in Table 13. Three pins were cast in each length throughout the range where cracking occurred and at two lengths above and below the cracking range. Two pins were cast in each of the other lengths. Pin length of 1.75 in. was determined to be the transition to 100% cracking. Ten pins with 1.75 in. length were cast to determine cracking response at the transition length. This was performed prior to the development of the automated control system using a timed heating cycle with heating times given in Table 13. The casting parameters are shown in Table 14. One set of cast pins from the validation test is shown in Figure 31. The maximum circumferential cracking was measured, Figure 32, along with cracking response for each of the samples. 100% cracking occurred in all samples of length in. and longer. As the mold length increases, the thermal strain also increases. The thermal strain reaches a critical value which cannot be supported by the liquid film present at the end of solidification. At this point cracking occurs as it would in a weld. 73

97 Table 12: Test Material Composition (wt. %) Al C Cr Fe Mn Mo Nb Ni Si Ti rem Table 13: CPTT Validation Experiments Pin Length (in) Sample Mass (g) Heating Time (s) Number of Tests Performed

98 Table 14: Casting Parameters Shielding Gas Chamber Pressure Gas Flow Rate Ramp Up Time Ar (4.8 purity) 3.45 kpa (0.5 psi) 10 CFH 500 ms Heating Time See Table 3 Ramp Down Time Coil Current Mold Alloy Coil Frequency 500 ms 434 A Cu-Be 224 khz Figure 31: CPTT Pins in Filler Metal ERNiCrFe-13 Produced with the Third Generation CPTT 75

99 Figure 32: Cracking Response in Filler Metal ERNiFeCr-13 Determined with the Third Generation CPTT The maximum circumferential cracking for the second and third generation CPTT was compared using the same heat of material, Figure 33. The same sharp transition to 100% cracking was found using both systems. Both the second and the third generation CPTT generated maximum circumferential cracking curves with zero to 100% cracking range of 0.75 in. However, the maximum circumferential cracking curve generated with the third generation CPTT was shifted to longer pin lengths. There is a considerable amount of scatter at the transition length of 1.75 in. At this length, cracking may or may 76

100 not occur due to noise in the process. One possible source of this scatter is the exact casting temperature, because these pins were cast after a set heating time prior to development of automated temperature control. Figure 33: Solidification Cracking Response in Filler Metal ERNiFeCr-13 Determined with the Second and Third Generation CPTT using C17000 Cu-Be Molds The reason for the shift in length required to cause cracking to occur in the redesigned system invites question of what factor control the solidification cracking behavior in the CPTT. In the previous system utilizing gas tungsten arc melting, the amount of superheat in the material was not well controlled or well known. It is most 77

101 likely that the material was superheated to a greater degree in the redesigned system with induction heating. An increase in the casting temperature would lead to a decrease in the cooling rate during solidification. Assuming that the rate of strain accumulation is proportional to the cooling rate, a decreased cooling rate would lead to a decreased strain rate. The decrease in strain rate would allow more time for fluid to flow and prevent cracking. Also, less strain would accumulate in the last liquid to solidify as a result of thermal contraction in the region of the pin which had already solidified. Solidification cracking on the surface of a 1.75 in. cast pin of Generation 3 CPTT was examined using scanning electron microscopy, Figure 34. Higher magnification of a surface crack in this cast pin reveals dendritic morphology that is typical of solidification cracking, Figure 35.[10] The fracture surface of a completely separated pin was also examined, Figure 36. This surface shows typical egg crate solidification morphology observed with the presence of liquid films. The fracture surface in a cast pin of the same heat of filler metal ERFeNiCr-13 produced with the second generation CPTT is shown in Figure 37. The fracture morphology in both cases is identical. 78

102 Figure 34: Surface Solidification Cracking in 1.75 in. Cast Pin of Alloy ERNiCrFe- 13 with 100% Circumferential Cracking Produced with the Third Generation CPTT Figure 35:Dendritic Fracture Morphology of Surface Solidification Crack in 1.75 in. Cast Pin of Alloy ERNiCrFe-13 Produced with the Third Generation CPTT 79

103 Figure 36: Fracture Morphology of Solidification Crack in 1.75 in Cast Pin of Alloy ERNiCrFe-13 Produced with the Third Generation CPTT Figure 37: Fracture Morphology of Solidification Crack in 1.75 in Cast Pin of Alloy ERNiCrFe-13 Produced with the Second Generation CPTT 80

104 The solidification cracking morphology was also examined in a sectioned 1.75 in. cast pin with 100% circumferential cracking, Figure 38. It is apparent that the cracking occurred in the interdendritic regions similar to solidification cracking during welding. Also, backfilling of cracks with low melting temperature constituents is evident at the crack tips. Cracks which have been completely healed by a backfilling mechanism are also present. 81

105 Figure 38: Cross Section of 1.75 in Cast Pin of Alloy ERNiCrFe-13 Produced with the Third Generation CPTT. 5.5 CPTT Automation Development Precise control of the superheat in the material is desired for repeatable casting conditions. Casting based on a timed sequence of events results in parameters dependent 82

106 on the mass of the sample. Also, some variation in the location and orientation of the sample leads to variation in the sample temperature during casting. In order to eliminate this variable, a system for automating the casting based on real time temperature readings was developed. Standard PID temperature control methods are not suitable for levitation melting as the power applied to the coil must remain at a high level to suspend the sample. A National Instruments LabView PC based control system has been developed utilizing a national instruments USB-6008 I/O device. The sample temperature is read into an analog input from the optical pyrometer, Omega IR2P at a frequency of 1 khz. The temperature is recorded during casting and output to a file for further analysis if necessary. The temperature reading is also used to initiate the casting process once the superheated material reaches the set temperature. A graphical user interface was developed, shown in Figure 39, was developed to enable simple user control. The set temperature is entered into a field. A stop button initiates the casting process to enable controlled casting if the test must be aborted prior to automatic trigger due to sufficient sample heating. The thermal history of the sample is displayed and a button is used to end the record at which point the user is prompted for a file path to save the data. 83

107 Figure 39: CPTT Control System Graphical User Interface The reproducibility of heating was evaluated by analyzing the temperature records for 9 casting cycles with samples ranging in mass from 12 g to 16 g. The casting temperature was set to 1500 C. The thermal histories for these tests are plotted in Figure 40. The average peak temperature was 1521 C with a standard deviation of 8.8 C. The automation system has therefore enabled a high level of sample reproducibility. 84

108 Temperature ( C) time (s) Figure 40: Thermal History for Melting 9 Samples of Alloy 52MSS-C with Mass Ranging from 12-16g and Set Casting Temperature of 1500 C Also, during casting it is necessary to allow residual gas in the mold to vent and enable mold filling. An electromagnetic valve was placed on the lower exhaust vent in the lower mold holder. This valve is closed during melting and no gas is able to flow from the chamber. The valve is opened upon initiation of the casting process allowing gas to vent. The pressure and flow rate are set prior to casting by adjusting the gas regulator and a needle valve attached to the outlet of the electromagnetic valve. This enables a slightly higher chamber pressure to enable mold filling while reducing the chamber pressure during solidification to the set level. 85

109 5.6 CPTT Mold Optimization In this study, the thermal history of the test material in the CPTT was recorded to evaluate the solidification cooling rate (SCR). The thermal history during GTA welding was measured to obtain a target value for the CPTT to represent actual welding conditions. The CPTT pin and mold were modeled using finite element analysis (FEA). The interfacial heat transfer coefficient between the cast pin and the mold wall was obtained through inverse modeling. A series of FEA simulations have been performed to evaluate the effect of mold material and various aspects of mold design on the SCR in the region where cracking typically occurs in the CPTT. The results of these simulations have been used to select a suitable material for the CPTT mold and to develop an optimized mold design to replicate the cooling conditions in GTA welding Cooling Rate Measurement in the Cast Pin Tear Test In order to perform the test, a small charge is levitation melted in the coil with the temperature monitored optically. Once the charge reaches a pre-determined casting temperature, the coil is de-energized and the sample is transferred into the mold by gravity. The design of the initial CPTT mold is shown in Figure

110 Figure 41: Second Generation CPTT Mold Design A mold with an opening for a thermocouple to pass through the mold wall was used for the temperature measurement. The mold material was alloy C17000 (Cu-1.7Be) and the cast material was filler metal 52M. A type C thermocouple (W-W/Re) protruding from a ceramic insulating sleeve was inserted 1mm deep into the mold. A type K thermocouple (Ni-Ni/Cr) was welded to the outer surface of the mold. The thermocouple locations are shown in Figure 42. The thermal histories in the cast pin and at the mold external surface were measured at a sampling rate of 4 khz using a fast sampling data acquisition system and recorded with a personal computer. The chemical compositions of the mold material and of the tested filler metal are given in Table 2. The parameters of the CPTT casting procedure are presented in Table 3. 87

111 Figure 42: Thermocouple Locations in the CPTT Mold Table 15: Chemical Composition of the Alloys used in Cooling Rate Measurements, wt.% CPTT Mold C Mn Si Cr Ni S P Mo Fe Nb Cu Be C rem* 1.70 Cast Pin / Filler Metal Inconel 52M rem Overlay Substrate 303 stainless steel clad with ER308LSi rem *Co+Ni >0.20, Co+Ni+Fe<

112 Table 16: CPTT Test Parameters Coil Current 434 A Casting Temperature 1500 C Casting Material Mold Material Mold Length Sample Mass Shielding Gas Gas Flow Rate Filler Metal Inconel 52M C17000 (Cu-1.7Be) 38 mm 12 g Argon 5.6 L/min The measured thermal histories of the cast pin and at the mold external surface are shown in Figure 43. Because the CPTT is used to evaluate solidification cracking which occurs while liquid is still present, the cooling rate during solidification is of greatest interest. The liquidus and solidus temperatures of the cast pin material were determined using a Scheil Simulation in JMatPro thermodynamic modeling software. The liquidus and solidus temperatures were predicted to be 1394 C and 1185 C respectively. The cooling rate in this temperature range (solidification cooling rate) was 245 C/s. 89

113 Temperature ( C) Liquidus Solidus Cast Pin Mold Surface Time (s) Figure 43: Measured Thermal Histories in the Cast Pin and at the External Mold Surface. Casting Temperature 1500 C, Cast Pin of Filler Metal 52M, Mold Material C Cooling Rate Measurement in GTA Weld Pool In order to evaluate how accurately the CPTT represents actual welding conditions, it was necessary to measure the GTAW thermal history under conditions similar to those experienced in practice. Solid solution strengthened Ni-base alloys are commonly used for weld overlay repairs in the nuclear power industry. Such weld overlays are produced using GTAW process with addition of cold wire in the weld pool. In this study, weld overlay repair was simulated using bead on plate (BOP) welding with 90

114 filler metal 52M. The weld beads were deposited on Type 303 stainless steel base plate which had been previously clad with 308LSi at approximately 50% dilution. Type C thermocouple was plunged in the molten weld pool during welding. The weld cooling history was measured at a sampling rate of 4 khz using a fast sampling data acquisition system and recorded with a personal computer. The chemical compositions of the filler metal and the clad substrate are given in Table 15. The welding was performed with parameters that are typically used in weld overlay repairs, Table 4. The length of the test weld was 110 mm. The experimental set up and a cross sectioned weld bead indicating the thermocouple location are shown in Figure 44. Table 17: Typical GTAW Procedure for Weld Overlay Repairs Base Material Clad Stainless Steel Plate* Filler Metal Inconel 52M Plate Dimensions 150mm x 110mm x 20mm Weld Current 240 A Weld Voltage 9.8 V Travel Speed 10.2 cm/min Deposition Rate 0.41 kg/hr *Type 303 Stainless Steel Plate Clad with ER308LSi (~50% Dilution) 91

115 a) b) Figure 44: Temperature Measurement in the Weld Pool of GTAW Cold Wire Process: a) Experimental Set-up, b) Weld Cross Section at the Thermocouple Location. The thermal history measured in the GTAW BOP experiment is shown together with the CPTT thermal history in Figure 45. The average solidification cooling rate (SCR) for the BOP weld was 114 C/s which, approximately half the cooling rate of 245 C/s found previously in the CPTT. In order to relate the results of the CPTT to welding design, the cooling conditions during solidification must be similar. This difference in the cooling rates motivated the application of FEA modeling and simulations to evaluate the factors that control the cooling rate in the CPTT and to optimize the CPTT mold material and geometry in order to better reproduce the cooling conditions in actual GTA welding. 92

116 Temperature ( C) Liquidus Cast Pin Tear Test GTAW Bead on Plate Solidus Time (s) Figure 45: Comparison of Thermal Histories in CPTT and BOP GTA Welding The microstructure in the BOP weld was examined to measure the primary dendrite arm spacing for comparison with the CPTT. Photomicrographs of the BOP weld are shown in Figure 46. Several solidification cracks were present in the BOP representative welds. The primary dendrite arm spacing was measured in the pin cast from alloy 52M to obtain a baseline cooling rate as well as a transvarestraint test specimen of alloy 52M from a previous investigation. Five measurements were taken across several dendrite arms and the average was coompared. The DAS measurements are shown in Table 18. The DAS is larger in the BOP weld and the transvarestraint test specimen when compared with the cast pin. This provides further evidence that the 93

117 cooling rate during solidification is higher in the CPTT than GTAW under these specific conditions. Figure 46: Microstructure in Bead on Plate Representative GTA Weld Using Alloy 52m on Clad Stainless Steel Base Material Table 18: Primary Dendrite Arm Spacing Measured in CPTT Pin Cast in C63000 Mold from Alloy 52M, Representative Bead on Plate Weld and Alloy 52M Transvarestraint Test Specimen Primary DAS (µm) σ Alloy 52M Cast Pin Bead on Plate Weld Transvarestraint Test

118 5.6.3 FEA Modeling of the CPTT The CPTT pin and mold were modeled using Pro-Cast FEA software. The CPTT geometry was imported into Pro-Cast using Parasolid file format. A mesh was created within the pin and the mold using four node tetrahedral elements with length of 1mm. The FEA models of the cast pin and the mold are shown in Figure 8. The cast pin and mold materials used in the FEA model were filler metal 52M and copper-beryllium alloy C17000 respectively. The chemical composition of filler metal 52M is given in Table 2. The temperature dependent properties used for thermal analysis in Pro-Cast TM are thermal conductivity, density, fraction solid and enthalpy. The temperature dependent properties of filler metal 52M were obtained using the Scheil function contained within JMatPro thermodynamic modeling software, Figure 48. These predicted properties of filler metal 52M were used in the FEA simulations. All parameters of the FEA model are summarized in Table 5. The interfacial heat transfer coefficient between the cast pin and the mold was determined using inversed modeling based on experimentally determined cooling histories and will be described in a subsequent section. 95

119 a) b) Figure 47: FEA Models of the CPTT Mold Cast Pin Assembly (a) and of the Cast Pin (b). 96

120 Figure 48: Temperature Dependent Material Properties of Filler Metal Inconel 52M Predicted Using JMatProTM. 97

121 Model Geometry and Mesh Table 5: Parameters of the FEA Model of CPTT Cast Pin / Mold Design Geometry and Dimensions Shown in Figure 3 CAD File Format Mesh Element Type Mesh Element Length Parasolid 4-node tetrahedral 1 mm Initial Conditions Mold Temperature 20 C Molten Charge Temperature 1500 C Boundary Conditions All External Surfaces Air-Cooling (10 W/m 2 -K, ambient temperature 20 C) Interfacial H.T. Coefficients Mold to Bottom Plug 1000 W/m 2 -K Mold to Casting Temperature Dependent, Determined using Inverse Modeling Mold Material Properties Alloy Density Specific Heat Capacity Thermal Conductivity C17000 Cu-1.7Be 8.26 g/cc J/g- C 118 W/m-K Cast Pin Material Properties Alloy Thermophysical Properties Filler metal Inconel 52M Temperature Dependent, Predicted using JMatPro TM 98

122 The developed FEA model of the CPTT was used to simulate the solidification sequence in a cast pin of filler metal 52M. Contour plots of the fraction solid at four timesteps generated in these simulations are shown in Figure 49.Figure 49: FEA Predicted Solidification Sequence in a Pin of Filler Metal 52M Cast in Mold of Alloy C17000 The FEA predicted solidification sequence in the CPTT was evaluated qualitatively by comparison to the solidification microstructure in a cast pin of alloy 52M (shown in Figure 50). Figure 50a shows the microstructure in the foot of the pin with vertical columnar-dendritic growth. Figure 50b shows the transition from a fine grained chill zone which forms on the mold surface to columnar dendritic growth further in the pin. The growth direction in the cast pin microstructure is in full agreement with the solidification sequence predicted using the developed FEA model. 99

123 Time 0.00 s 0.44 s 0.76 s 0.92 s Fraction Solid Figure 49: FEA Predicted Solidification Sequence in a Pin of Filler Metal 52M Cast in Mold of Alloy C17000 a) b) Figure 50: Microstructure in a Cast Pin of Filler Metal 52M: a) Upward Dendritic Solidification in the Cast Pin Foot, b) Transition from Fine Grained Chill Zone to Columnar-Dendritic Growth in the Cylindrical Part of the cast Pin. 100

124 5.6.4 Determination of the Interfacial Heat Transfer Coefficient The interfacial heat transfer coefficient (IHTC) between the mold wall and the cast pin surface was determined using an inverse modeling calibration technique which is incorporated in the Pro-Cast software. The FEA model defined in Table 5 was used in these inverse simulations to predict the thermal history in a particular location of interest in the cast pin. This location coincided with thermocouple location in the cast pin temperature measurement experiments shown in Figure 4. The experimentally measured cast pin cooling curve (Figure 5) was used as a target cooling curve in the inversed modeling. An initial estimate for the IHTC was also defined. The temperatures where the IHTC was calculated and optimized were 200 C, 1100 C, 1250 C and 1410 C. The inverse modeling procedure compares the predicted thermal history at the defined cast pin location with the experimentally measured thermal history. At each iteration of inverse modeling, the IHTC is optimized based on the residual differences between the predicted and the measured thermal histories. This process is repeated until convergence is reached between the measured and the simulated thermal histories. The temperature dependence of the IHTC determined using the inverse modeling procedure is shown in Figure 51a. The predicted IHTC has strong temperature dependence and decreases significantly as the cast pin solidifies. The IHTC was validated by comparison of the FEA predicted and the experimentally measured cooling curves. Figure 12b shows that the cooling curve predicted using the derived temperature dependence of the IHTC is in good agreement with the measured cooling curve. 101

125 Heat Transfer Coefficient (W/m2*K) Temperature ( C) Heat Transfer Coefficient 4000 Solidus Liquidus IHTC Validation Liquidus CPTT Measurement FEA Simulation Solidus Temperature ( C) Time (s) a) b) Figure 51: a) Temperature Dependance of the IHTC between the Mold Wall of Alloy C1700 Mold and a Cast Pin of Filler Metal 52M Determined using Inverce Modeling in Pro-Cast TM, b) Validation of the IHTC by comparison of FEA Predicted and Measured Cooling Histories in Cast Pin of Filler Metal 52M Effect of Mold Material on the Solidification Cooling Rate The effect of the mold material thermal conductivity on the cast pin cooling rate in the solidification temperature range was evaluated by performing a series of eight FEA simulations. The FEA model defined in Table 5 was used in these simulations. The SCR was predicted in the cast pin location where solidification cracking typically occurs. The FEA node of interest is located 8 mm below the top of the pin and is coincident with the pin central axis, Figure 52. The mold thermo-physical properties were considered based on available copper mold alloys, Table

126 Figure 52: Location of the FEA Node for Cooling Rate Predictions. Table 6: Thermo-physical Properties of Commercially Available Copper Mold Alloys at Room Temperature Material Thermal Conductivity (W/m- K) Density (g/cc) Specific Heat Capacity (J/g-K) Cu Mold Max SC Mold Max V Mold Max LH Mold Max HH C Mold Max XL C

127 The mold material properties used in the FEA simulations are listed in Table 7. The heat conductivity was varied in a range between 5 and 385 W/m-K that covers most of the available copper mold alloys. The specific heat capacity and density were held constant to identify the effect of mold thermal conductivity alone. Due to unavailable temperature dependences of the thermal conductivity in copper mold alloys, room temperature values were assumed in this investigation. Table 7: Mold Material Properties used to Evaluate the Effect of Thermal Conductivity on Cast Pin Cooling Rate Simulation Thermal Conductivity Density (g/cc) Specific Heat Capacity (J/g- C) (W/m-K) The predicted effect of mold material thermal conductivity on the cooling rate through the solidification temperature range is shown in 104

128 Solidification Cooling Rate C/s Figure 53. The conducted FEA simulations show that mold thermal conductivity has the greatest effect on the SCR below thermal conductivity values of 100 W/m-K. The lowest thermal conductivity in the considered copper mold alloys was 38 W/m-K of alloy C Using this alloy as a mold material, the SCR in the cast pin can be only reduced by 45 W/m-K. The results of these FEA simulations show that further reduction in the cooling rate of the cast pin aiming to better replicate the conditions of GTAW would require optimization in the mold geometry Thermal Conductivity W/m-K Figure 53: Effect of Mold Thermal Conductivity on the Solidification Cooling Rate in a Cast Pin of Filler Metal 52M Cast in Mold with Properties Specified in Table

129 5.6.6 Effect of Mold Geometry on the Solidification Cooling Rate In order to optimize the mold geometry, the effects of the mold volume and of the cast pin surface to volume ratio on SCR were evaluated. The FEA model defined in Table 5 was used in these simulations with the mold geometry shown in Figure 41. The SCR was predicted in the cast pin location where solidification cracking typically occurs, Figure 13. The mold volume affects the SRC through the heat capacity of the mold. In a series of three FEA simulations, the mold volume was varied by reducing the mold wall thickness at constant cast pin diameter of 4.8 mm (3/16 in) and mold length of 38 mm (1.5 in). The outer diameter of the mold was decreased to reduce the wall thickness. The simulations were performed with wall thicknesses of 8.5 mm (nominal), 6.0 mm, and 3.5mm. The cast pin surface to volume ratio affects the SRC through the rate of heat extraction at the cast pin / mold contact surface. In a series of three FEA simulations, this ratio was varied by simultaneously increasing the pin diameter and decreasing the pin length to maintain a constant pin volume. The FEA simulations were performed at pin diameters of 4.8 mm (nominal), 6.4 mm and 8.0 mm. The mold outer diameter was held constant at the nominal value of 22 mm (7/8 in), which resulted in reduction of the mold volume with increasing the pin diameter. The simultaneous variation of the cast pin surface to volume ratio and of the mold volume represented the combined effect of the rate of heat extraction and mold heat capacity on the SCR. 106

130 Solidification Cooling Rate C/s The results of these simulations are summarized in Figure 54. Decreasing the mold heat capacity by reducing the mold wall thickness decreased the SCR by 82 o C/s. The combined effect of reducing the rate of heat extraction and of decreasing the mold heat capacity by increasing the pin diameter resulted in a reduction of the SCR by 187 o C/s. It is possible that increasing the pin diameter could provide an added gain as a result of a larger air gap formation as the pin diameter decreases due to thermal contraction Mold Wall Thickness mm Figure 54: Effects of the Mold Wall Thickness and Pin Diameter on the Solidification Cooling Rate in the CPTT 107

131 5.6.7 Optimization of the CPTT Mold Material and Geometry The results of the performed FEA simulations have indicated that the targeted cooling rate during the cast pin solidification cannot be achieved by a separate optimization of any of the evaluated controlling factors. Among these controlling factors the pin surface to volume ratio had the greatest effect in reducing the cast pin SCR, followed by the mold volume and the mold material thermal conductivity. A combination of these factors was utilized to design a mold which replicates the cooling rate experienced during solidification in structural weld overlays and may be applied in the current apparatus. A new mold was designed based on the results of this study. The goal was to reduce the cast pin cooling rate by reducing the rate of heat extraction at the pin / mold contact surface, and the heat capacity and thermal conductivity of the mold. The geometry of the new mold is shown in Figure 55. The mold design was optimized to reduce the pin surface to volume ratio and the mold volume. This was achieved by enlarging the pin diameter and reducing the mold wall thickness, while maintaining unchanged the nominal pin volume and mold outer diameter. The external dimensions of this mold were kept the same as the current design shown in Figure 3 to allow direct application into the CPTT apparatus. The cylindrical portion in the foot was eliminated to reduce the total pin volume and avoid formation of shrinkage porosity in the foot. Alloy 6300, which had the lowest thermal conductivity of the available copper mold alloys, was selected as the mold material on the newly designed CPTT mold. The 108

132 chemical composition of this alloy is given in Table 8 and its thermophysical properties are listed in Table 6. The FEA model defined in Table 5 was used to predict the solidification cooling rate in a pin of filler metal 52M cast with the newly designed mold. The SCR was predicted in the cast pin location where solidification cracking typically occurs, Figure 13. The FEA predicted cooling curve in the cast pin is compared to the measured cooling curve in the weld pool of cold wire GTAW in Figure 17. The predicted SCR of 129 C/s in the cast pin nearly matches the target value of 114 C/s, which validates the optimized mold design. Figure 55: Proposed Mold Design 109

133 Temperature ( C) Table 8: Chemical Composition of Alloy C63000 wt.% Cu Sn Zn Fe Ni Al Mn Si balance Weld Measurement CPTT Simulation Time (s) Figure 56: Comparison of the FEA Predicted Cooling Curve in a Cast Pin with Optimized Mold Design and the Measured Cooling Curve in BOP Weld Prototype Mold Screening Two prototype molds were produced to test the effect of pin diameter on the cracking response in various materials. One of the major advantages of the CPTT is the ability to perform the tests with small laboratory melts. Larger diameter molds will 110

134 increase the amount of material necessary for testing. Drawings of the two molds produced are shown in Figure 57. The new mold design incorporates a larger head to retain material during filling and a smaller foot to reduce the required sample volume. The length of mold produced was 2 in. The molds were produced from C63000 aluminum bronze. Figure 57: CPTT Prototype Mold Designs 5 mm Inside Diameter (left) and 6 mm Inside Diameter (right) In order to evaluate the cracking sensitivity in the molds, three materials were tested. Previous CPTT with the second generation system identified Rene 142 as the most susceptible material tested to date, Alloy 600 as the least susceptible and Inconel 52MSS as an intermediate material. Also, Inconel 52MSS is the type of material considered in 111

135 this investigation. One pin was cast in each mold from each material. Five pins of alloy 52MSS-E were cast in the 5mm mold. The experiments were performed at a casting temperature of 1500 C. The chamber pressure and gas flow rate were 0.3 psi and 3 cfh respectively. The results of these screening tests are summarized in Table 19. In this case, Inconel 52MSS was the least susceptible. This is likely due to an increase in the amount of crack backfilling with reduced cooling rate. Table 19: Results of C63000 Al-Bronze Mold Screening Tests Performed in Molds with 2 in length Cast Pin Diameter 5mm 6mm Rene 142 Separated Separated 52MSS HV % Separated 75 % In the mold with larger diameter, the cracking susceptibility was reduced. This would require longer length molds in order to have a valid test. Also, more material would be required for molds of the same length. Use of larger diameter molds would require samples with mass above the limit of the casting apparatus. 5 mm was selected as the mold inner diameter to cause cracking to occur with a reasonable mass requirement. 112

136 An FEA simulation was performed to evaluate the cooling rate in the 5 mm mold using the modeling procedure described previously. The predicted cooling rate in the C63000 molds was 160 C/s. The microstructure of pins cast in the redesigned molds fabricated from C63000 Al-bronze was compared with pins cast in the original C17000 Cu-Be molds. Light optical micrographs of pins cast from Alloy 52MSS-C in both mold types are shown in Figure 58. The microstructure in the original CPTT molds consists of a fine-grained equiaxed chill zone adjacent to the mold followed by a region of columnar-dendritic growth with an equiaxed dendritic region in the center of the cast pin. The solidification structure is similar in pins cast with the new molds. The region of columnar growth in pins cast in the redesigned molds extends further toward the center than that region in the original molds. 113

137 A) B) Figure 58: Microstructure of Pins Cast with the Third Generation CPTT in Alloy 52MSS-C A) C17000 Cu-Be Molds B) C63000 Al Bronze Molds Dendrite arm spacing in solidification structures provides an indication of the cooling rate during solidification.[8] The dendrite arm spacing (DAS) was compared in pins cast with the original and redesigned CPTT molds. Six measurements of primary and secondary DAS were taken in each sample. This was done by measuring the distance across several dendrite arms and dividing that distance the number of dendrite arms. The average of these measurements as well as the standard deviation for each is shown in Table 20. The primary and secondary dendrite arm spacing is larger in the redesigned Albronze mold. This indicates that the cooling rate has been reduced. 114

138 Table 20: Comparison of the Effect of Mold Material on Dendrite Arm Spacing in Pins Cast from Alloy 52MSS-C in the Third Generation CPTT Average Dendrite Arm Spacing (µm) Mold Material Primary σ p Secondary σ S C17000 Cu-Be C63000 Al Bronze A validation test was performed in the new molds by casting 3 pins at each of 9 lengths. These experiments were performed at 1500 C with a chamber pressure of 0.3 psi and a flowrate of 3cfh. The results of this test are shown in Figure 59. Crack free pins resulted in each length and no conclusive trend could be observed. 115

139 Circumferential Cracking 100% 90% 80% 70% 60% 50% 40% 30% 20% 10% % Pin Length Figure 59: Results of Initial Validation CPTT Validation of C63000 Al-Bronze Molds with Alloy 52MSS-C Pins were cast from several high-cr Ni-base filler metals in the 2.5 inch mold to ensure that cracking is possible in the redesigned molds. It was observed that the chamber pressure has a significant effect on cracking response in the redesigned molds. Pins could be cast from Inconel 52MSS without cracking in all lengths with a chamber pressure of 0.3 psi. When the chamber pressure was reduced to 0.1 psi, 100% cracking occurred in alloys 52M, 52MSS-C, 52MSS-E, 52i-B and 690NB. The pressure in the casting chamber was reduced to 0.1 psi. Three pins were cast in alloy 52MSS-C at 2.5 inches and 100% cracking resulted. Also, complete cracking occurred at 2.5 in in pins of Alloys 52M, 52MSS-C, 52MSS-E, 690NB and 52i-B. The 116

140 chamber pressure clearly plays a role in cracking response. Also, the scatter in the validation data can be explained by the level of pressure control in the current system. Given the apparent sensitivity of the test to small variations in pressure, the current system does not have the capability for that level of control. Improvements in the system are needed to measure and control the pressure and flowrate of gas during casting. To obtain an indication of the cracking response in the redesigned molds compared with the previous version, initial testing was performed with the current system. The increment in mold length was in, ranging from 1.5 to 2.5 in. Longer mold provide a higher level of restraint. Experiments were performed with a casting temperature of 1500 C, chamber pressure of 0.1 psi and flowrate of 2 cfh. The experiment consisted of two pins cast in each length and the maximum value of circumferential cracking was considered at each length. The results of this initial investigation are shown in Figure

141 Figure 60: CPTT Cracking Response of IN52MSS-C in Redesigned molds and Previous Molds The length of pin required to induce cracking is shifted to longer lengths in the new molds. This is likely due to a reduced strain rate leading to more time for fluid to flow and backfilling of strained boundaries where cracking may occur otherwise. Also, the length of columnar-dendritic growth region is longer in the new molds, leading to a microstructure which could enable increased fluid flow. Cracking was examined using SEM. The effect of crack healing is prominent in the redesigned molds, as shown in Figure 61. Cracks with complete healing are visible, as well as a large partially healed surface crack. Closer investigation shows the morphology 118

142 of the backfilling constituent which is indicative of the presence of laves phase in the terminal stages of solidification. Qualitative EDS indicated an enrichment of Nb along the strained boundaries exhibiting backfilling as shown in Figure 62. Figure 61: Crack Healing in Pin Cast from Alloy 52MSS-C in C63000 Al-Bronze Molds 119

143 A) B) Figure 62: EDS Spectrum Comparing Interdendritic Region with Crack Healing (A) to the Dendrite Core (B) It is possible that the reduced cooling rate in the CPTT has a synergistic effect with increased pressure on the melt in the flow of fluid to heal solidification cracks. This phenomenon follows the RDG theory.[16] Pressure loss in the final liquid to solidify occurs due to the constricting solid network and strain applied to the boundaries. If there is not sufficient pressure in the liquid along wet boundaries to avoid cavitation and void formation, cracking may result. In the CPTT with reduced cooling rate, the strain rate in most likely reduced as well. This enables more time for fluid to flow and thus a reduced fluid velocity and pressure loss. This possible synergistic effect of cooling rate and pressure on the molten pool is worthy of further investigation. Perhaps this could be applied in actual welds. Adjustment 120

144 to welding parameters could be made to reduce the cooling rate and thus the rate of strain between impinging dendrites due to solidification shrinkage. At the same time, perhaps, a trailing shield with jets of cover gas directed to the trailing edge of the weld pool could increase the pressure on the liquid in the final stages of solidification leading to increased crack healing. This would likely be most effective in alloys with a significant fraction eutectic, which are also generally more resistant to DDC. 121

145 CHAPTER 6: HIGH-CR NI-BASE FILLER METAL WELDABILITY 6.1 Thermodynamic Simulation Solidification temperature ranges (STR) were determined under non-equilibrium solidification conditions using Scheil simulations. The solidification temperature range was considered from the liquidus of the material to 0.98 fraction solid. The results of these simulations are given in Table 21. Table 21: Calculated Solidification Temperature Ranges for Solid Solution Strengthened Nickel Based Filler Metals Designation ERNiCrFe-13 ERNiCrFe-7A Material 52M 690NB 52MSS-C 52MSS-E T L T s,eq T eutectic T s ΔT eq ΔT primary fcc ΔT L-0.98 S A wider STR was predicted in the ERNiCrFe-13 alloys than in the ERNiCrFe-7A type alloys. The primary difference between these alloys is the addition of Nb to improve resistance to DDC in ERNiCrFe-13. Nb partitions strongly to the liquid during solidification forming γ/nbc eutectic and possibly Laves phase during the terminal 122

146 stages of solidification. This will effectively depress the solidus of the material and increase the solidification temperature range. In the ERNiCrFe-7A type alloys, 690NB exhibited a STR of 184 C compared with 215 C in 52M. In the ERNiCrFe-13 type alloys, 52MSS-E was predicted to have a reduced STR of 236 C when compared with 52MSS-C which has a STR of 251 C. 6.2 Single-Sensor Differential Thermal Analysis Single-sensor differential thermal analysis (SS-DTA) was used to experimentally determine the STR. In SS-DTA, the thermal history of a solidifying weld pool is measured. One measured thermal history for alloy 690NB is shown in Figure 63. Due to latent heat release during phase transformations, a knee can be seen in the cooling data. A reference cooling curve is calculated based on a portion of the thermal history where no transformation takes place. The difference between the actual thermal data and reference cooling curve is plotted to determine transformation temperatures as shown in Figure 64 for determination of the solidus temperature in alloy 52MSS-E. 123

147 Temp (deg.c) Temp (deg.c) Temperature 1800 Ch4Vin Time (secs) Figure 63: Measured Thermal History of a Solidifying Weld Pool in Alloy 690NB Used for Single Sensor Differential Thermal Analysis 1140 Ch4Vin+ Ch4Vin X: Y: X: Y: Fit Range: o C Rough Order Estimate: 1000 Tss = Tp = T85 = T (deg.c) 1060 Fit Range: o C Rough Order Estimate: 1040 Tss = Tp = T85 = T (deg.c) Figure 64: Temperature Differential between Cooling Data and Calculated Reference Curve Used to Determine Solidus Temperature in Alloy 52MSS-E 124

148 The STR for 690NB and 52MSS-E determined experimentally using SS-DTA is shown in Table 22 together with data which was acquired previously using this method [56]. The eutectic reaction could not be detected in either ERNiCrFe-7A type alloy, most likely due to the decreased fraction eutectic present during solidification of these alloys. The solidification temperature range in alloy 690NB was found to be 170 C compared with 241 C in alloy 52MSS-E. This is in reasonable agreement with data obtained previously which indicates a wider STR in ERNiCrFe-13 alloys compared with ERNiCrFe7A alloys. The solidification temperature ranges are shown along with those predicted using Scheil calculations in Figure 65. The alloys are ranked according to their STR from greatest to least as follows: 82, 52i-B, 52MSS-E, 52i-A, 52MSS-B, 52MSS-A, 690NB, 52MSS-C, 52M. 125

149 Figure 65: Solidification Temperature Range Determined Using Scheil Predictions and Single-Sensor Differential Thermal Analysis Table 22: Transformation Temperatures Determined Using SS-DTA for High-Cr Ni-Base Alloys during Solidification Alloy Designation T L T rate change T eutectic T s ΔT eutectic ΔT L-S 82 ERNiCr ± ± ± ± i-A ERNiCrFe- 1343± ± ± ± i-B ± ± ± ± MSS-A ERNiCrFe- 1351± ± ± ± MSS-B ± ± ± ± MSS-C 1332± ± ± ± MSS-E 1344± ± ± ± M ERNiCrFe- 1358± ± ± NB 7A 1380± ± ±

150 Maximum Circumferential Cracking 6.3 Cast Pin Tear Test The third generation cast pin tear test (CPTT) was used to evaluate the solidification cracking susceptibility of alloys 690NB and 52MSS-E. Alloys 52M and 52MSS-C were also evaluated to provide a standard for comparison. The maximum circumferential cracking which occurred in these tests is shown in Figure 66. Longer pin lengths required to cause cracking indicate lower susceptibility in the material. 100% CPTT Maximum Circumferential Cracking 52mss-c 80% 60% 52m 690nb 52mss-e 40% 20% 0% Pin Length (in) Figure 66: Maximum Circumferential Cracking Determined Using the Third Generation CPTT with C17000 Cu-Be Molds 127

151 The maximum pin length where no cracking occurred was in 52MSS-C, 52MSS-E and 690NB. Crack free pins were cast in alloy 52M up to in. This indicates that alloy 52M is the least susceptible to solidification cracking. The maximum pin length required to cause complete circumferential cracking in alloy 52MSS-C was 1.75 in compared with the other alloys which all cracked completely at in. This indicates that Alloy 52MSS-C is more susceptible than alloy 690NB and alloy 52MSS-E. Alloy 52MSS-E showed increased cracking response at and 1.75 in. indicating that it is more susceptible than alloy 690NB. Based on this reasoning, the relative cracking susceptibility ranking from most susceptible to least susceptible is: 52MSS-C, 52MSS-E, 690NB, 52M. Several other similar materials have been tested using the second generation CPTT. The maximum circumferential cracking response is shown in Figure 67. As noted in the description of CPTT development efforts, direct comparison is not possible between the second and third generation CPTT results. 128

152 Maximum circumferential Cracking, % Maximum Circumferential Cracking (MCC) Response Curves iA 52i-B M 52MSS-A 52MSS-B 52MSS-C MSS-A 52M 82 52i-B i-A Pin Length, in Figure 67: Maximum Circumferential Cracking for High-Cr Ni-Base Filler Materials Determined Previously Using the Second Generation CPTT [56] Alloys 52M and 52MSS-C have been tested using both systems and the maximum length of pin cast without cracking was shifted by in in both cases. In order to compare the alloys tested in this investigation with other previous alloys, the maximum pin length without cracking was reduced by in, shown in Figure 68. The adjusted maximum crack free pin length was 0.75 in in alloys 52MSS-E, 52MSS-C, 52MSS-B, 52MSS-A and 690NB indicating similar solidification cracking susceptibility in these alloys. The other alloys considered were less susceptible. The ranking in terms of solidification cracking susceptibility from most to least susceptible is 52i-A, 52M, 82, 52i-B. 129

153 Adjusted Crack Free Pin Length in Figure 68: Comparison of Adjusted Maximum Pin Length without Cracking Determined using the Cast Pin Tear Test Partially cracked pins cast from 690NB and 52MSS-E were cross sectioned longitudinally for analysis, Figure 69. Cracking in both pins is interdendritic occurring along boundaries. A larger fraction of eutectic type constituents are present in alloy 52MSS-E, which is typical in this type of alloy due to the additions of Nb which promote γ/nbc and Laves phase in the terminal stages of solidification. Crack healing is apparent surrounding the cracking in alloy 52MSS-E, unlike alloy 690NB. 130

154 A) B) Figure 69: Solidification Cracking and Microstructure in CPTT Pins Cast in C17000 Cu-Be Molds A) Alloy 690NB B) Alloy 52MSS-E The fracture surface of pins produced in the third generation CPTT using C63000 Al-Bronze molds were evaluated using SEM, Figure 70. All fracture surfaces exhibited morphology indicative of liquid films present when cracking occurred, confirming solidification cracking as the mode of failure. A transition in fracture surface morphology was observed from the edge of the pin to the center. The fracture surface at the edge of the pin contained smaller droplets compared with the egg crate morphology visible at the center of the cast pin. This is likely due to an increased amount of liquid present at the center of the pin due to the temperature gradient across the pin causing the outer edges to exist at a later stage of solidification. 131

155 Center Edge 52M 690NB 52MSS-C 52MSS-E Figure 70: Fracture Surfaces on CPTT Pins cast in C63000 Al-Bronze Molds with 100% Circumferential Cracking 132

156 Maximum Crack Distance (mm) 6.4 Transvarestraint Testing The maximum crack distance (MCD) was determined at each level of strain in Alloys 690NB and 52MSS-E, Figure 71. Solidification cracks which occurred at 5% augmented strain during the transvarestraint test are indicated by black arrows in Figure 72. Cracking initiated at 1% augmented strain in alloy 690NB compared with 2% in alloy 52MSS-E. At strain levels greater than 2%, the extent of cracking was more severe in alloy 52MSS-E. This indicates that alloy 690nb is more susceptible to solidification cracking under low restraint conditions and 52MSS-E is more susceptible under high restraint conditions MSS-E 690NB % 1.0% 2.0% 3.0% 4.0% 5.0% 6.0% 7.0% Augmented Strain (%) Figure 71: Maximum Crack Distance in Alloys 52MSS-E and 690NB Determined in the Transvarestraint Test 133

157 Figure 72: Solidification Cracking Which Occurred in the Transvarestraint Test at 5% Augmented Strain A) Alloy 52MSS-E B) Alloy 690NB A possible and likely explanation for the relative behavior of these materials in the transvarestraint test is the effect of crack healing with eutectic constituents at the terminal stages of solidification. Under low restraint conditions, alloy 52MSS-E has sufficient liquid remaining at the end of solidification to compensate for strain along solidifying boundaries. At high levels of restraint, the strain cannot be compensated and cracking is more severe. The transvarestraint specimens were also inspected for DDC, which occurs in the solidified weld when strain is applied. Small cracks were observed in the Transvarestraint specimens of 52MSS-E and 690NB at augmented strain levels of 5% and greater. DDC was severe in alloy 690NB at 5% augmented strain as shown in Figure

158 Figure 73: DDC in Alloy 690NB Transvarestraint Specimen Tested at 5% Augmented Strain The MCD determined in the transvarestraint test was compared with results from previous alloy testing in Figure 74. The previous testing was performed on a different piece of equipment. Validation testing on the new equipment comparing results with those obtained on the previous equipment indicated reasonable correlation above saturation strain of approximately 2%. The MCD at 5% strain was considered to obtain a comparative ranking of the cracking susceptibility in these alloys, Figure 75. The ranking order of solidification cracking susceptibility based on MCD from least to greatest is 52M, 82, 690NB, 52MSS-E, 52i-B, 52MSS-B, 52i-A, 52MSS-C, 52MSS-A. In general, the alloys with significant Nb additions, which are beneficial for resistance to DDC, were determined to be more susceptible to solidification cracking. 135

159 MCD, mm This is due to the depression of the solidus temperature resulting from low melting temperature eutectics forming in the terminal stages of solidification. This introduces the challenge in alloy development as the solution to one problem creates a new challenge EN52MSS MLTS/EN52i EN82H 82 52i-A 52i-B 52M 52MSS-A 52MSS-B 52MSS-C 52MSS-E (Low Iron) TG-SN 690NB Strain, % Figure 74: Maximum Crack Distance Determined Using the Transvarestraint Test for Several High-Cr Ni-Base Alloys [56] 136

160 Figure 75: Maximum Crack Distance Determined in the Transvarestraint Test at 5% Strain 6.5 Strain to Fracture Testing Strain to fracture (STF) testing was performed at 950 C to evaluate the DDC resistance of alloys 690NB, 52MSS and 52i-B. The total number of cracks occurring at each strain level is plotted in Figure 76. Examples of DDC in each alloy tested are shown in Figure 77. These results are compared with other alloys which have been tested previously in Figure 78. The threshold strain level for DDC as determined by the STF was 8% in alloy 690NB and alloy 52MSS-E. Alloy 52i-B had a threshold strain level of 14%. 137

161 Number of cracks Response to DDC at 950 o C 52MSS-E 690NB 52i-B Strain, % Figure 76: DDC Cracking Response of High-Cr Ni-Base Filler Materials Determined Using the Strain to Fracture Test at 950 C. Figure 77: Ductility Dip Cracking in Strain to Fracture Specimens Tested At 950 C A) 690NB, 9.9% Strain B) 52MSS-E, 8.0% Strain C) 52i-B, 15.7% Strain 138