DEVELOPMENT OF A GLEEBLE BASED TEST FOR POST WELD HEAT TREATMENT CRACKING IN NICKEL ALLOYS. A Thesis

Transcription

1 DEVELOPMENT OF A GLEEBLE BASED TEST FOR POST WELD HEAT TREATMENT CRACKING IN NICKEL ALLOYS A Thesis Presented in Partial Fulfillment of the Requirements for the Degree Master of Science in the Graduate School of The Ohio State University By Seth Jason Norton, B.S. ********* The Ohio State University 2002 Master s Examination Committee: Dr. John Lippold, Adviser Approved by Dr. Dave Dickinson Adviser Department of Welding Engineering

2 ABSTRACT The purpose of this work was to use a thermo-mechanical simulator to develop a simple test that will quantify susceptibility to the various forms of postweld heat treatment cracking, including stress-relief cracking and strain-age cracking. Materials evaluated include Waspaloy and Alloy 718. Samples are initially given an elevated temperature thermal exposure that simulates the weld HAZ. Upon cooling from elevated temperature, the sample is loaded such that yield strength magnitude stresses are present at room temperature. The sample is then heated to a selected PWHT temperature and held for up to 8 hours. Acquisition of tensile force data during PWHT shows an initial relaxation of stresses, followed by a rise in stress as precipitation reactions proceed. Hot ductility tests were performed at various PWHT temperatures and times. From the reduction in area measurements a mathematical model was developed to relate ductility to PWHT temperature and time. The alloys tested show a significant dip in ductility at elevated temperature. Waspaloy samples exhibited lower ductility at elevated temperatures than did Alloy 718. SEM fractography and optical microscopy were used to analyze the failed samples. Minimum ductility samples revealed ductile intergranular fracture paths in Waspaloy and a tendency for transgranular cracking in Alloy 718. Intergranular microii

3 cracks were evident in a coarse grained heat affected zone in both alloys. Grain size was shown to effect ductility and fracture mode. Smaller grained simulated HAZ samples exhibited more ductility and exhibited more transgranular ductile fracture. Larger grains caused an increase in intergranular fracture and a reduction in elevated temperature ductility. Alloy 718 showed evidence of bending of crystal lattices to accommodate stresses. Both alloys showed indications of an effect of crystal orientation on cracking susceptibility. iii

4 Dedicated to my parents who encouraged my curiosity and desire to learn. iv

5 ACKNOWLEDGEMENTS I would like to thank my adviser, John Lippold, for is patience, encouragement and intellectual support. The Edision Welding Institute receives my gratitude for the funding that made this work possible. Thanks to Shu Shi for helping with the metallographic preparation. Assistance with Gleeble problems from David Jacon and Dan Eaken must also be acknowledged. To Cameron Begg I must give thanks for his help with using the SEM in the Ohio State Universities campus electron optics facility. To my wife, Shiloh, and my son, Caleb, I apologize for my frequent absence while writing this thesis. Thank you for being so understanding. v

6 VITA November 29, Born Sanger, California B.S. Engineering, LeTourneau University Researcher/Lab Instructor, LeTourneau University Welding Engineering Dept present...Graduate Research Associate, The Ohio State University FIELDS OF STUDY Major Field: Welding Engineering vi

7 TABLE OF CONTENTS Page Dedication... iv Acknowledgments...v Vita... vi List of Figures...x List of Tables... xvi Abbreviations... xvii Chapters: 1. Introduction Background Superalloys Physical Metallurgy Solution Strengthening Precipitation Strengthening Grain Boundary Strengthening Other Phases Weldability Issues Solidification Cracking Liquation Cracking Ductility Dip Postweld Heat Treatment Cracking Residual Stress in Welds Gleeble Thermal System Mechanical System Hot Ductility Test...18 vii

8 2.6 Postweld Heat Treatment Cracking Tests Circular Patch Test Controlled-Heating-Rate Test Constant Load Rupture Test Previous Gleeble-based Tests Wu and Herfert Duval and Owczarski Franklin and Savage Dix and Savage Nawrocki et al Concluding Remarks Objectives Experimental Approach Experimental Procedure Materials Sample Preparation Equipment Environment Gleeble Tests HAZ Simulation Postweld Heat Treatment Simulation Hot Ductility Test Sample Numbering Metallography Grain Size Measurement Fractography Follow-up Testing Results Gleeble Feedback Acquisition HAZ Residual Force Stress Increase During PWHT Strength Ductility Waspaloy Characterization Waspaloy Microstructures Waspaloy Fractures Alloy 718 Characterization Alloy 718 Microstructures...89 viii

9 6.4.2 Alloy 718 Fractures Discussion Procedure Development Initiation of Stroke Peak HAZ Temperature PWHT Temperature Use of Hot Ductility to Assess PWHT Cracking Susceptibility Jaw Type Stroke Rate and Distance Residual Stress Repeatability Stress Relaxation Aging Response Stress Rise During PWHT Hot Strength Hardness Testing Ductility Ductility Model PWHT Temperature and Time Effects Grain Size Effect Ductility Recovery Ductility Comparison Fracture Paths Stress Accommodation Orientation Conclusions Suggestions for Future Work References Appendix A: Experimental Run Order Appendix B: Hot Ductility Measurements ix

10 LIST OF FIGURES Figure Page 2.1 Schematic illustration for PWHT cracking in nickel-base superalloys Relationship between estimated weldability and aluminum and titanium content Development of residual stresses in fusion welds Specimen held in jaws with thermocouples welded at center Temperature profile across the free span of a Gleeble sample Hot ductility of Waspaloy bar Circular patch weld restraint specimen used for evaluating resistance to PWHT cracking Comparison of ductility results for René 41 heats studied with controlled heating rate tensile test Cut away drawing of constant load rupture apparatus showing pre-notched specimen held in controlled atmosphere furnace Schematic representation of relation between constant load rupture graph and constant load rupture graph and stress relaxation in an actual weldment Time-temperature dependence of fracture in Waspaloy Effect of time and temperature on the isothermal stress relaxation behavior of solution annealed René x

11 2.13 Stress relaxation behavior at 1600 F (871 C) Effect of testing temperature on ductility at two different strain rates Effect of strain rate and temperature on ductility Schematic illustration of Lehigh stress-relief cracking test cycle Data acquired from preliminary PWHT cracking test Schematic of desired data from initial test As received Waspaloy bar microstructure As received Alloy 718 microstructure Schematic illustration of sample dimensions Schematic illustration of thermal and mechanical control for HAZ and PWHT simulation Schematic of reduction in area measurements Typical acquisition curves from HAZ and PWHT portion of test Typical stress vs. time curve for hot ductility portion of test Residual stress developed in Waspaloy HAZ simulations (1240 C peak, mm/min stroke rate) Residual stress developed in Alloy 718 HAZ simulations (1184 C peak, mm/min stroke rate) Increasing stress relaxation as PWHT temperature is increased in Waspaloy Delayed force response in Alloy 718 specimen Comparison of stress increase after relaxing from heating to PWHT temperature (Waspaloy C, Alloy C)...63 xi

12 6.8 Comparison of stress increase after relaxing from heating to PWHT temperature (Waspaloy 816 C, Alloy C) Comparison of stress increase on log time scale (Waspaloy C, Alloy C) Comparison of stress increase on log time scale (Waspaloy C, Alloy C) Yield strength of Waspaloy measured in hot ductility portion of test Yield strength of Alloy 718 measured in hot ductility portion of test Ductility surface plot from regression analysis of Waspaloy data Ductility surface plot from regression analysis of Alloy 718 data Comparison of ductility dips in Waspaloy and Alloy 718 regression models for no PWHT Comparison of ductility dips in Waspaloy and Alloy 718 regression models at 3 hours PWHT Actual sample data point plotted with regression model at 2.5 hours (Waspaloy 800 C for 2.5 hours) Actual sample data point plotted with regression model (Waspaloy 840 C for 3 hours) Actual sample data point for 2 second peak HAZ temperature plotted with regression model for 20 second peak HAZ temperature (Waspaloy 816 C for 1 hours) Actual sample data point for 60 second peak HAZ temperature plotted with regression model for 15 second peak HAZ temperature (Alloy C for 2 hours)...75 xii

13 6.21 Representative microstructure of preliminary tests Voids formed through centerline of preliminary sample Microcracks at grain boundaries in CGHAZ away from centerline of sample Cavities at grain boundaries in Waspaloy SEM Image of cavities in Waspaloy Wedge shaped cracks at boundaries in Waspaloy. (2 second peak HAZ, 816 C7PWHT) Note transgranular fracture of some grains shown by arrows SEM image of wedge crack at boundary in Waspaloy (boundaries highlighted by dashed lines) Wedge cracks and intergranular fracture in a preliminary Waspaloy test sample Crack at interface of a twin and a grain boundary in Waspaloy (895 C PWHT) Another crack at the intersection of a twin and grain boundary in Waspaloy (760 C PWHT) Grain elongation in tensile direction of Waspaloy sample tested at 895 C Intergranular fracture in preliminary tests conducted in Waspaloy Micro-ductility on intergranular fracture surface Intergranular fracture surface with micro-ductility in 760 C PWHT Waspaloy sample Evidence of both intergranular and intragranular fracture in 760 C PWHT Waspaloy sample Intergranular fracture surface in 816 C PWHT Waspaloy sample...87 xiii

14 6.37 Higher magnification of apparently smooth grain surface in figure 6.35 (note evidence of twin on fracture surface at arrows) Intergranular fracture surface in 895 C PWHT Waspaloy sample Markings of micro-ductility and macro-ductility in Waspaloy sample (895 C PWHT) Waspaloy sample held at peak HAZ temperature for 2 seconds showing both intergranular and transgranular fracture Cavities and wedges in Alloy 718. (718 C PWHT, 8 hours) Note lattice bending in some crystals Cavities along grain boundaries in Alloy 718 (718 C PWHT, 4 hours) Grain elongation in Alloy 718 after hot ductility test (sample tested at 818 C) Bent twins in deformed crystal of Alloy 718 sample tested at 818 C (note voids formed at globular NbC particles) Wedge crack in sample held at HAZ peak temperature for 60 seconds Fracture in sample held at HAZ peak temperature for 60 seconds Intragranular fracture path in Alloy 718 sample (718 C PWHT, 8 Hours) Detail of intragranular fracture in Alloy 718 sample High magnification of distortion in lattice at fracture surface in Alloy 718 (718 C PWHT, 4 hours) Ductile dimples on surface of Alloy 718 specimen tested at 818 C Detail of deep ductile dimples on surface of Alloy 718 specimen tested at 818 C...97 xiv

15 6.52 Alloy 718 fracture surface showing both intragranular ductile fracture and intergranular fracture (718 C PWHT, 8 hours) Different surface characteristics exhibited in Alloy 718 intergranular fracture (718 C PWHT, 4 hours) Different intergranular fracture characteristics on surface of single grain (718 C PWHT, 4 hours) Evidence of grain bending in Alloy 718 fracture (718 C PWHT, 4 hours) Fracture surface of Alloy 718 sample held for 60 seconds at peak HAZ temperature and PWHT at 718 C for 2 hours Detail of ductile fracture in center of figure Detail of intergranular fracture shown in figure Higher magnification of intergranular fracture in figure 6.57 revealing microductility Voids and cracks formed in HAZ simulation prior to PWHT simulation by stroking before sample cooled below DRT Data acquisition from running PWHT cracking test on 304 stainless steel bar Schematic representation of expected grain size effect on ductility Schematic of Zener s proposal for intercrystalline fracture. Sliding along the boundary mn relaxes the shear stress along the boundary and concentrates the stress at the grain corner. The stress concentration at a grain corner may be relieved by plastic deformation ahead of the sliding boundary xv

16 LIST OF TABLES Table Page 2.1 Composition and Chronology of Nickel-Base Superalloys Summary of reviewed Gleeble based PWHT tests Chemical compositions of Ni-base alloys tested Waspaloy PWHT temperature-time matrix Alloy 718 PWHT temperature-time matrix Standard conditions in HAZ simulation for both alloys Controlled variables for additional tests Waspaloy grain size measurements Alloy 718 grain size measurements...90 xvi

17 ABBREVIATIONS ANOVA ASM AWS BCT CCT CGHAZ CLR DDC DRT FCC HAZ HCP NDT NST PMZ PWHT RPI SEM SGB SSGB TCP TTT analysis of variance American Society for Metals American Welding Society body centered tetragonal constant cooling transformation coarse grained heat affected zone constant load rupture ductility dip cracking ductility recovery temperature face centered cubic heat affected zone hexagonally closed packed nil ductility temperature nil strength temperature partially melted zone post weld heat treatment Rensselaer Polytechnic Institute scanning electron microscope solidification grain boundary solidification subgrain boundary topologically closed packed time-temperature-transformation xvii

18 CHAPTER 1 INTRODUCTION Precipitation hardening materials are often very difficult to weld. Obtaining the desired properties from a weldment in such materials requires careful control of not only the welding parameters but also the pre-weld and postweld thermal treatments. Postweld heat treatment (PWHT) is often performed on welded structures to relieve stresses and to optimize mechanical properties. Under certain stress and temperature combinations, the weld HAZ or fusion zone may crack during PWHT. The problem of PWHT cracking was recognized early in the development of nickel-base superalloys and has continued to be a persistent problem. Attempts to better understand the phenomenon of PWHT cracking by many researchers have led to a variety of specialized tests. To date however, there is no standardized test for quantifying the susceptibility of an alloy to PWHT cracking. The development of a standard test will be useful not only for comparison of cracking susceptibility among different alloys, but would aid alloy designers in developing new superalloys for critical welded applications. The Gleeble system allows precise control of thermal and mechanical conditions and is an excellent candidate for the development of a test for quantifying PWHT cracking susceptibility. Through careful control of thermal conditions the Gleeble 1

19 may reproduce the different regions of a weld heat affected zone and simulate isothermal heat treatments. The Gleeble s mechanical system allows tensile loading and monitoring of load and deformation. The problem with previous Gleeble tests is that they did not accurately simulate the thermo-mechanical history of a weld. Most tests have not simulated the development of residual stresses in a weldment as it cools after the weld metal is deposited. Those tests that have developed residual stresses on cooling from the peak temperature achieved during welding have not allowed relaxation of the stresses in subsequent PWHT simulations. The approach used in this work was to more closely simulate the actual conditions experienced by the HAZ during welding. The Gleeble thermo-mechanical simulator was used to accomplish this. Elevated temperature ductility is assumed to correlate strongly with the potential for cracking. The effects of residual stress, PWHT temperature, and hold time on elevated temperature ductility are evaluated in order to quantify susceptibility to PWHT cracking. The technique developed also provides insight into the PWHT cracking mechanism in nickel-base superalloys. 2

20 CHAPTER 2 BACKGROUND 2.1 Superalloys The superalloys were developed in the 1940 s for commercial use in jet engine turbines [1]. Prior to this, some nickel-base precipitation-hardening alloys were invented but not made commercially available [2]. As engine design evolved, the materials used in the hot section of the engine were also being advanced. The characteristics emphasized in these new alloys were oxidation resistance, high temperature strength, and stress rupture properties. The development of alloys far outpaced the development of welding procedures for the alloys and resulted in problems for welding engineers [3]. Table 2.1 shows the evolution of the superalloys. 3

21 Alloy Date C Cr Co W Mo Al Ti Others Inconel Pre Fe-7 Inconel X Fe-7 Nb-1 Hastelloy X Fe-18.5 Waspaloy Zr-0.1 Udimet B-0.02 Inconel Fe-18 Nb-5 AF2-1DA Ta-1.5 Hf-1.5 Mar-M Nb-1 Hf-2 Table 2.1 Composition and Chronology of Nickel-Base Superalloys [1] 2.2 Physical Metallurgy Excellent strength and oxidation resistance of superalloys at high temperatures are the reasons for their use in turbine applications. There are three mechanisms for achieving strength: solid solution strengthening, precipitation strengthening, and grain boundary strengthening. Other constituents, besides the strengthening precipitates, may be present in the microstructure of the material, all of which will be described in the following sections. 4

22 2.2.1 Solution Strengthening Initial alloys created for use in jet turbines were nickel-chromium alloys. The nickel and chromium form a solid solution termed gamma (γ). Chromium enhances the oxidation/corrosion resistance and provides some solid solution strength. Other alloying additions such as aluminum, tungsten, and molybdenum also work as potent solution strengtheners. Iron, vanadium, titanium and cobalt may also be added to improve the strength of the alloy through solution strengthening [4] Precipitation Strengthening Around 1944, aluminum and titanium were added to the basic nickel-chromium alloy to produce superalloys that were precipitation hardenable [1]. These elements formed an ordered FCC phase in the normal γ-nickel matrix. The resulting intermetallic phase (Ni 3 (Al,Ti)) has been designated γ and is the cause for the improvement of the creep strength for most superalloys [1,2]. The precipitates are very small and are normally only observable with electron microscopy in the peak aged condition. The γ phase, like the γ phase, is a solid solution. The lattice parameter for both γ and γ depends on the alloy composition; therefore, lattice parameters may differ for different alloys and may vary within an alloy if there is non-homogeneity. According to Prager and Shira [2] the volume percent of γ in nickel-base superalloys exceeds 25% of the total, so even a small difference in lattice parameters between the two phases may yield considerable dimensional changes upon precipitation of the second phase. 5

23 The lattice mismatch between the surrounding γ matrix and the γ matrix has been shown to have an effect on the shape of the precipitate. Small lattice mismatch yields a spherical shaped precipitate while a larger misfit yields cubic precipitates [5]. An applied tensile stress may also affect the morphology of the γ precipitate. Tien and Copley [6] showed that a tensile stress along the [001] direction will convert cubic γ into plates lying on the (001) plane. The use of niobium as a substitute for aluminum in superalloys yielded a modified second phase. The new phase has been named γ and is an ordered body centered tetragonal (BCT) based on the Ni 3 Nb compound. Alloys such as Inconel 718 may form both γ and γ as second phase strengtheners. While the γ strengthened alloys are often inferior in strength to other heat resistant nickel alloys [7] they have better weldability, as will be discussed later Grain Boundary Strengthening Elements such as carbon, zirconium, and boron are added to achieve specific grain boundary characteristics and improve elevated temperature ductility [8]. Magnesium and hafnium are also grain boundary modifiers [1,5]. These elements are mismatched in atomic size with the γ matrix and segregate to grain boundaries. It has been suggested that these elements strengthen boundaries by filling vacancies and reducing grain boundary diffusion rates [1,5]. It is also thought that boron may facilitate slip accommodation at the boundaries by either promoting slip transmission across the boundary or enhancing dislocation generation at the boundaries [9]. 6

24 2.2.4 Other Phases There are other phases associated with the nickel superalloys that may be present in the microstructure. MC, M 6 C, and M 23 C 6 carbides are frequently observed in superalloys. The M in these formulas represents a metallic element and may actually be representative of several elements so that the formula for a M 23 C 6 carbide may actually be, for example, Cr 21 Mo 2 C 6. A less frequently reported carbide is M 7 C 3 [2,10]. MC type carbides are blocky FCC structures that form below the freezing temperature of the surrounding matrix. At elevated temperatures, MC carbides react with other phases to produce different carbides and phases. The following reactions demonstrate the possible changes [5]. MC + γ M 23 C 6 + γ MC + γ M 6 C + γ M 6 C + M M 23 C 6 + M Where M and M represent metallic elements in the matrix. In addition to the carbides, there may be nitrides (TiN) and an assortment of other phases, which have been labeled η, µ, σ, δ, and Laves. The µ, σ, and Laves phases are topologically closed packed (TCP) and generally form platelike structures that promote cracking. The σ phase is especially deleterious to nickel-base superalloys. It forms in the C temperature range and nucleates frequently on M 23 C 6 carbides [5]. The η phase is an HCP structure with the composition Ni 3 X. It forms from γ when the titanium, niobium, or tantalum content is high enough. Lastly, the δ phase is an orthorhombic 7

25 crystal structure of Ni 3 Nb. If the Nb content is sufficiently high, it may convert to the δ phase at temperatures above 650 C. 2.3 Weldability Issues Welding provides an economical means for joining subcomponents with little additional weight or reduction of service capabilities [3]. Fusion welding practices involve melting and resolidifying metal to form a joint. However, several cracking issues make welding of superalloys difficult: solidification cracking, liquation cracking, ductility dip, and post weld heat treatment cracking [11] Solidification Cracking Solidification cracking is associated with the fusion zone of weldments. The two conditions that are essential for this type of cracking to occur are a susceptible microstructure and restraint [12]. There are several theories about the mechanism behind solidification cracking [13,14,15,16] and most focus on the solid-liquid region of a phase diagram. Because welding is a non-equilibrium process, low melting point eutectics may be formed during cooling. Liquid films may form along solidification grain boundaries (SGB) and solidification subgrain boundaries (SSGB). Restraint caused by thermal contraction during cooling or from external loads may cause cracks to form by opening the boundaries which have liquid films along them. Once the material has cooled to the effective solidus temperature of the material, there may already be many cracks at SGB 8

26 and SSGB s. The fracture surface of a solidification crack normally exhibits a cellular or dendritic appearance but may be disguised by liquid films [12] Liquation Cracking Liquation cracking, as the name implies, is related to the presence of liquid films. The difference between liquation cracking and solidification cracking is that liquation cracking does not occur where the bulk material has been melted and is resolidifying. Rather, liquid films may wet grain boundaries in the HAZ that border the fusion zone. The heat of welding causes the formation of liquid films at the boundaries by one of two mechanisms [12]. Alloys with other particles or phases present in the matrix may experience constitutional liquation [17], which is the formation of a liquid film at the particle/matrix interface without melting the particle. The liquid film in turn pins and penetrates moving grain boundaries. This is known as the penetration mechanism. The other explanation for liquation cracking is the segregation of impurity elements to grain boundaries, which in turn form low melting point constituents that lower the melting point of the boundary compared to the surrounding matrix. Fracture surfaces created by liquation cracking are intergranular and may be decorated with resolidified liquid films. Liquation by both penetration and segregation mechanisms has been shown to occur in nickel base superalloys [18,19,20,21]. 9

27 2.3.3 Ductility Dip Ductility dip cracking (DDC) is caused by a severe loss of ductility below the solidus temperature of a material. The mechanism for this type of cracking is not understood and has not been defined. There are however several characteristics which have been documented [12]. DDC is associated with austenitic (fcc) microstructures and is always intergranular. The temperature range over which DDC occurs is between the solidus and half the melting temperature of the material. Also, DDC is associated with large grain size. Fractures caused by DDC may be found in both weld metal and the HAZ. The fracture surface may show ductile features, in which case the fracture is termed ductile intergranular. DDC may occur on-heating or on-cooling and is thought to be more common in high purity materials (low tramp elements) Post Weld Heat Treatment Cracking One of the main problems associated with the nickel base superalloys is the PWHT portion of the welding procedure. The rapid heating and cooling associated with welding causes residual stresses to be built up in weldments. A PWHT is used to relieve residual stresses and achieve maximum strength. This is done by heating to a solutionizing temperature followed by holding at an aging temperature. Solutionizing relieves the residual stresses from welding and puts alloying elements back into solution. Holding in the aging temperature range causes precipitates to come out of solution. A weldment must pass through the aging temperature range while being heated to the solutionizing temperature because the aging temperature range at which precipitates 10

28 come out of solution is lower than the solutionizing temperature. Precipitation prior to the relief of stresses is harmful and often results in PWHT cracking. Figure 2.1 gives a schematic illustration of the welding and PWHT thermal cycles for a nickel-base superalloy. Figure 2.1 Schematic illustration for PWHT cracking in nickel-base superalloys [29] It is thought that PWHT cracking is due to low ductility and high strain accumulation in the HAZ [1,2,22,23]. Several mechanisms have been proposed for the low ductility of the HAZ. The proposed mechanisms include embrittlement of the grain boundary by oxygen during heat treatment [24,25], embrittlement of the grain boundary 11

29 by liquation or solid state reactions during welding [26], and a change in deformation mode from transgranular slip to grain boundary sliding [21,23,27,28]. Figure 2.2 Relationship between estimated weldability and aluminum and titanium content [11] PWHT cracking in nickel-base alloys has often been called strain-age cracking because it occurs in highly restrained weldments when they are heated through their aging temperature range [29]. The strains in the HAZ may be caused by relaxation of residual welding stresses, stresses caused by thermal contraction or expansion, and the 12

30 dimensional changes associated with the formation of strengthening phases. There is some overlap between the temperatures at which precipitates form rapidly and stress relief begins. For this reason the hardener content of an alloy has been linked to its cracking susceptibility. Figure 2.2 shows the relationship between aluminum and titanium content and the weldability of several superalloys [30]. Alloy 718 is shown to be readily weldable in Figure 2.2. This alloy was designed specifically to resist post weld heat treatment cracking [1]. The γ precipitate is thought to precipitate at a slower rate in Ni- Al-Nb alloys than γ does in analogous Ni-Al-Ti alloys [7,22]. It is thought this difference in particle growth rates allows relaxation of residual stresses before the matrix becomes too hardened. Intergranular fracture paths characterize fracture surfaces caused by PWHT cracking. Fractures may exhibit both smooth and ductile intergranular surfaces [12]. The fractures normally occur in the HAZ very close to the fusion boundary of a weld and may even form in the PMZ. Cracks formed in the PMZ display the presence of liquid films [22]. Cracking may occur in the fusion zone also, but this is normally controlled through the use of ductile filler metals [28]. 2.4 Residual Stress in Welds The steep temperature gradients associated with welding are responsible for the development of residual stresses in the weld and surrounding material. The AWS welding handbook [31] gives an excellent model for the development of residual stresses 13

31 caused by thermal gradients. The weld develops stresses on the order of the yield stress and the surrounding material develops compressive stresses. Figure 2.3 shows the development of residual stresses caused by changing temperature distribution in a weld. Figure 2.3 Development of residual stresses in fusion welds [31] 14

32 2.5 Gleeble The Gleeble system was developed at Rensselaer Polytechnic Institute (RPI) for the simulation of heat affected zones [32] and later revised to include mechanical loading [33]. Over time, the Gleeble has evolved into a robust testing machine capable of performing a wide variety of functions. These systems have been used to simulate HAZ s, test hot ductility properties, create TTT and CCT diagrams, simulate thermal fatigue, and many other metallurgical studies. The operating principles of the Gleeble are not too complex and merit closer inspection Thermal System Samples are placed between two sets of jaws that are part of a high current circuit. As current passes through the sample, it is resistively heated. A thermocouple welded to the center of the sample is used in a feedback loop to monitor temperature and control current. A sketch of the jaws, sample, and thermocouple wires is shown in Figure 2.4. The temperature to which the sample is heated may be controlled manually or automated and controlled by a computer program. A chiller, which pumps water through the jaws and transformer, keeps the system from overheating and helps remove heat from the sample during cooling. 15

33 Free span Figure 2.4 Specimen held in jaws with thermocouples welded at center Temperature distribution in the sample between the jaws is not uniform and a parabolic temperature distribution (Figure 2.5) results with a peak temperature at the center of the sample. Using jaws of different geometry and conductivity can change the distribution. The distance between the jaws, or free span, may also affect the thermal gradient by making it steeper as the jaws are brought closer together. Typical jaws are made of copper and conduct heat away from the sample very quickly. Hot jaws are made from stainless steel and have very small area of contact between the jaws and the sample. The result of the lower thermal conductivity of the stainless steel and the small area of contact is a much flatter temperature gradient across the sample as shown in 16

34 Figure 2.5. The gradient has been mathematically modeled and documented recently by Brown et al [34]. DSI, the manufacturer of the Gleeble system, recommends using copper jaws for welding simulations because of the ability to produce a steep temperature gradient [35]. Figure 2.5 Temperature profile across the free span of a Gleeble sample [35] 17

35 2.5.2 Mechanical System The mechanical system consists of a hydraulic ram attached to one of the sets of copper jaws. The other set of jaws is fixed. By moving the ram, a sample can be placed in tension or compression along its center axis. The movement of the ram is part of a feedback loop, just as the temperature is for the thermal system. Ram movement may be monitored by distance traveled (stroke), force applied, or strain. Stroke and force are the two most common types of control. As with the thermal system, the mechanical system may be controlled manually or with a computer. Some older Gleebles and all new ones have the capability to switch between mechanical control modes while being computer controlled Hot Ductility Test One of the first uses of a Gleeble with a mechanical system was the study of elevated temperature properties of metals [33]. A hot ductility (or hot strength) test is often used to show a material s ductility and strength characteristics at elevated temperatures. A machined tensile test specimen is pulled uniaxially, while held at an elevated temperature, until failure occurs. Ductility is quantified by measuring the reduction in cross sectional area at the fracture. Three temperatures commonly reported in hot ductility testing are the nil ductility (NDT), nil strength (NST), and ductility recovery temperature (DRT) [20,30,36]. The NDT is the temperature at which the material exhibits no reduction in area before fracturing. It is associated with the presence of thin, continuous liquid films. At the NDT a material may still exhibit some strength. At 18

36 the NST the material is thought to have extensive, continuous liquid films along grain boundaries and no tensile strength is exhibited. The DRT is that temperature on-cooling from above the NDT at which the liquid films have solidified enough for the sample to display some ductility. On-heating tests may be conducted by heating to the desired testing temperature before pulling the sample and on-cooling tests heat to some predetermined temperature between the NDT and NST before cooling to the testing temperature where tension will be applied. The data from a series of tests can be used to plot the on-heating and oncooling ductility (normally in terms of reduction in area transverse) against the testing temperature. A sample of the characteristic hot ductility curve for a superalloy is shown in Figure

37 DRT NDT NST Figure 2.6 Hot ductility of Waspaloy bar [20] 20

38 2.6 Postweld Heat Treatment Cracking Tests A review of published literature shows a great deal of testing took place on superalloy weldability in the 1960 s and early 70 s [1,2,21,22,23,24,25,26,27,28,36,37]. A variety of PWHT cracking susceptibility tests were developed as part of this work. This section reviews some of the tests Circular Patch Test The circular patch test is a self-restraint test used in industry to screen materials and qualitatively evaluate welding variables affecting PWHT cracking. It is essentially a yes/no test for determining whether or not a specific heat of an alloy is susceptible to cracking. The geometry of the specimens is largely tailored to the user s specifications [2]. The basis for all the variations of the circular patch tests is the same however. A circular weld is made in the material to be tested. The weld may be joining a circular patch into a matching sized hole or a bead-on-plate or bead-in-groove with no joint. The surrounding material provides restraint for the weld and may be welded to a stiffener (strong-back) for additional restraint. An example of one circular patch geometry used by Boeing is shown in Figure 2.7. The entire specimen (patch, surrounding material, and stiffener) is given a postweld solution heat treatment. After this heat treatment the patch and surrounding material are examined for cracks. 21

39 Figure 2.7 Circular patch weld restraint specimen used for evaluating resistance to PWHT cracking [2] Controlled-Heating-Rate Test Researchers at Rocketdyne came to the conclusion that ductility could be related to strain-age cracking resistance [38]. They developed an elevated temperature tensile test that utilized a constant heating rate. Solution-annealed tensile specimens were heated in a split tube furnace to 1000 F (538 C) where they were held to allow the test fixtures to reach equilibrium. After reaching equilibrium the sample was heated at a rate of 25 to 30 F/min (14 to 17 C/min) into the aging temperature range. Tensile testing took place 22

40 immediately upon reaching the test temperature. Elongation was measured by means of a high temperature extensometer. Figure 2.8 shows the results of controlled-heating-rate tests on several heats of René 41. Heats with higher minimum ductility were considered to have higher resistance to PWHT cracking. Figure 2.8 Comparison of ductility results for René 41 heats studied with controlled heating rate tensile test. [38] 23

41 2.6.3 Constant Load Rupture Test In the early 1970 s a test based on fracture mechanics was developed at the Welding Institute in England [28]. Specimens were pre-treated to simulate situations such as aged, strained, or HAZ thermally cycled, before notching and initiating a fatigue crack. The notched sample was rapidly heated to temperature in a controlled atmosphere furnace and then loaded in tension. The test was named the constant load rupture test (CLR). A drawing of the apparatus is shown in Figure 2.9. Figure 2.9 Cut away drawing of constant load rupture apparatus showing pre-notched specimen held in controlled atmosphere furnace. [28] 24

42 The stress intensity applied to the root of the notch was calculated and plotted against time to failure. As the stress intensity factor decreased, time to failure increased until some minimum stress intensity factor below which failure did not occur. A curve, which represented the crack/no crack boundary, could be drawn on a graph with axes of stress intensity factor and time. The researcher postulated that materials that had slow stress relaxation rates would cross the CLR curve and cracking would occur. A schematic representation of the constant load rupture curve and the stress relaxation of a weldment is shown in Figure Fractographic analysis led the researcher to conclude PWHT cracking was caused by grain boundary sliding and was independent of grain boundary precipitation. Figure 2.10 Schematic representation of relation between constant load rupture graph and constant load rupture graph and stress relaxation in an actual weldment. Left (a) a failure situation; right (b) no failure [28] 25

43 2.6.4 Previous Gleeble-based Tests Much of the weldability testing incorporated the Gleeble and focused on the problems associated with welding superalloys described above. The following is an overview of previous work on PWHT cracking using Gleeble simulations Wu and Herfert In 1967, two research engineers from Northrop Corporation published a report on their tests of nickel base superalloy René 41 [27]. The focus of their strain-age cracking test was on the effect of both peak temperature in the heat affected zone and the threshold stress for initiating cracking. The samples used in the testing were 0.25 in. diameter bars. The bars were initially cycled through various peak temperatures and cooled. The bar then had a notch machined into it to develop a triaxial stress field at the center of the sample. After machining the notch, the samples were put into tension (32-34 ksi). Once loaded, the bar was heated to 1600ºF (871ºC) and the load was maintained at the initial level with no relaxation in stress. From these tests, the researchers concluded cracking occurs only in sections of the HAZ heated above 2200ºF (1204ºC). They cited dissolution of M 6 C above 2200 F and precipitation of M 23 C 6 at grain boundaries on cooling as a factor in the cracking phenomena. The authors of the paper noted large grains and grain boundary precipitates in samples heated through 2300 F and small grains with relatively clean boundaries in a sample heated to a peak HAZ temperature of 2015 C. 26

44 In a second set of tests, all the samples were heated to 2300ºF (1260ºC) for the initial HAZ simulation. Each sample was notched as in the previous tests but the applied load was increased with each successive sample. The researchers found that stresses below 24.6 ksi did not cause cracking and loads above 33.2 ksi caused complete fracture. They concluded that the combination of thermally induced stresses, internal stress caused by γ precipitate, and restraint of the joint design causes cracking through continuous M 23 C 6 films which developed at grain boundaries near the weld fusion line Duval and Owczarski The next Gleeble based test for superalloy weldment cracking during heat treatment was designed by Duval and Owczarski (Pratt & Whitney), who published their work in 1969 [22]. Their test included both simulated HAZ s and actual weldments of Waspaloy and Alloy 718. Some of the HAZ simulations were designed to produce a partially melted region while others produced solid HAZ regions. The cracking portion of the test heated specimens to the PWHT temperature under no mechanical load. Once the desired temperature was reached, the specimen was loaded in tension by moving the jaws a set distance and then holding that position. Specimens were periodically unloaded and rapidly cooled to check for crack formation on the surface. The results of the test included a time-temperature C-curve behavior for Waspaloy shown in Figure None of the Alloy 718 HAZ specimens cracked during the isothermal creep-relaxation tests. Duval and Owczarski concluded that ductility differences between regions of the Waspaloy HAZ were a primary cause of different 27

45 cracking tendencies. They found no indication of transient embrittlement during heat treatment and suggested γ precipitation, intergranular carbide precipitation, partial melting during welding, and possibly strain-induced intragranular precipitation contributed to ductility differences in the HAZ. Figure 2.11 Time-temperature dependence of fracture in Waspaloy [22] Franklin and Savage The test developed by Duval and Owczarski was modified slightly by Franklin and Savage [37] and used to test René 41. The modification to the test was a differential 28

46 screw that limited the travel of the movable jaw and thus the total strain introduced in the specimen. The purpose of Franklin and Savage s study was to determine the effect of preweld heat treatments on the susceptibility of the HAZ to strain-age cracking. They tested samples of two pre-weld heat treatments: solution annealed to preclude γ precipitation, and overaged to form large γ precipitates. In addition to the pre-weld heat treatment, some samples were given a simulated HAZ treatment in a Gleeble. Figure 2.12 Effect of time and temperature on the isothermal stress relaxation behavior of solution annealed René 41. [37] 29

47 The results of the test showed that the stress in the specimens that had been solution annealed or given an HAZ treatment initially began to drop and then increased. Increasing temperatures yielded greater stress relaxation prior to the increase in stress for these specimens (see Figure 2.12). In overaged samples without an HAZ thermal cycle the stress relaxed continuously for the duration of the test (see Figure 2.13). The researchers noted M 23 C 6 carbides at grain boundaries that failed. No mention of grain size was made. Figure 2.13 Stress relaxation behavior at 1600 F (871 C) A = overaged base metal specimen. B = solution annealed base metal specimen [37] 30

48 Several conclusions about post-weld isothermal stress relaxation behavior were drawn from this work. One conclusion was that the as-welded structure of the region of the HAZ subjected to temperatures of 2380 F (1304 C) and above was independent of pre-weld heat treatment. This led to the conclusion that overaging is not an effective means of increasing the capacity of the HAZ to accommodate strains associated with residual welding stresses. Intergranular failure was associated with M 23 C 6 grain boundary networks in HAZ cycled specimens. Finally, stress relaxation is interrupted by a period of stress increase as a result of aging contraction accompanying the precipitation of γ in solution annealed metal Dix and Savage After Duval and Owczarski published their work linking strain-age cracking to a reduction in ductility [22] a new test was developed to determine the mechanical response of a superalloy to tensile loads at elevated temperatures. Dix and Savage used a Gleeble at RPI to test Inconel X-750 [23]. The goal was to determine the deformation modes operating in the temperature range where PWHT cracking occurs in superalloys. The test began by heating samples to the desired testing temperature at a constant rate. The specimen was then uniaxially loaded at a constant rate until failure occurred. The temperature, load, and elongation were all recorded during testing. Two crosshead velocities, 0.16 and 16.0 ipm (0.4 and 40.6 cm/min) were used over a temperature range of F ( C) for one set of tests. A second set of tests examined intermediate crosshead velocities (strain rates) at 1300 and 1600 F. 31

49 The initial set of tests confirmed the reduction in ductility and showed that the ductility reached a minimum at approximately the same temperature Duval and Owczarski [22] had shown to be the nose of the time-temperature C-curve. Figure 2.14 shows the minimum ductility at elevated temperature is dependent on strain rate. In their testing of intermediate strain rates, Dix and Savage found the ductility was largely unaffected until the crosshead velocity exceeded 1 in/min. The results of these tests are shown in Figure Analysis of the fractured samples showed evidence of recrystallization and migrated grain boundaries in the samples tested at 1700 C and above. Samples tested between 1000 C and 1600 C showed wedge shaped voids at grain intersections. Some of the samples also showed evidence of bends in normally straight-sided annealing twins. This lead Dix and Savage to conclude that transgranular slip and grain boundary sliding are competing deformation processes in the temperature range tested. Grain boundary sliding becomes more prevalent as temperature increases. Predominate grain boundary sliding lowers ductility until recrystallization and boundary migration occur at high temperatures and restore ductility. 32

50 Figure 2.14 Effect of testing temperature on ductility at two different strain rates [23] Figure 2.15 Effect of strain rate and temperature on ductility [23] 33

51 Nawrocki et al More recently, researchers at Lehigh University developed a Gleeble based test for stress relief cracking in Cr-Mo ferritic steels [39]. The test begins with an HAZ simulation. As the sample cools from the peak HAZ temperature, a uniaxial load is applied. The mechanical system is operated in force control mode, so the jaw moves until the programmed force is reached. The applied load remains constant once applied. The sample continues to cool to room temperature under tensile loading. After cooling to room temperature, the sample is subjected to a simulated PWHT by holding at an elevated temperature. A schematic of the test cycle is shown in Figure Figure 2.16 Schematic illustration of Lehigh stress-relief cracking test cycle [39] 34

52 The ductility, measured as percent reduction in area, was used to compare the susceptibility of two ferritic steels to stress relief cracking. One of the two alloys had low ductility (<10% RA) at all temperatures tested and was deemed to be more susceptible to cracking during PWHT. The proposed mechanism for fracture was microvoid coalescence at prior austenite grain boundaries, with carbides acting as void nucleation sites. 2.7 Concluding Remarks The versatility of the Gleeble has allowed it to be used in a variety of ways for testing susceptibility of materials to PWHT cracking. Each of the Gleeble based tests reviewed above is different from the others and focuses on different aspects of PWHT cracking. Table 2.2 summarizes the control method and conclusions drawn from each of the reviewed tests. The first four Gleeble based tests reviewed above did not address the development of residual stresses on cooling from the peak HAZ temperature. The Lehigh test utilized load control to consistently develop residual stress while cooling from the peak HAZ temperature. The shortcoming of the Lehigh test is that force control does not allow stress relaxation during PWHT. A test that is representative of actual weldments should develop residual stresses while cooling from the peak HAZ temperature and allow stress relaxation during PWHT. 35

53 Researchers Wu and Herfert Duvall and Owczarski Franklin and Savage Dix and Savage Nawrocki et. al. Alloy system studied Ni-base superalloy Ni-base superalloy Ni-base superalloy Ni-base superalloy Cr-Mo ferritic steel Simulation controls Constant tensile load applied at all temperatures Heat to test temperature, stroke set distance and hold Heat to test temperature, stroke set distance and hold by means of differential screw Heat to test temperature at constant rate and load at constant stroke rate until failure occurs Load to constant force on cooling from peak HAZ, maintain load during PWHT simulation Conclusions Thermally induced stresses, stresses caused by precipitation, and joint restraint cause cracking through M 23 C 6 films at grain boundaries. Large grain size more susceptible to cracking. Differences in ductility between regions of HAZ are main reason for PWHT cracking. Aging contraction due to γ precipitation shown by force increase. Intergranular failure associated with M 23 C 6 carbide networks. Maximum reduction in ductility at 1600 C.Failure at minimum ductility associated with grain boundary sliding. Onset of recrystallization above 1600 C causes sharp ductility increase. Ductility is useful measure of susceptibility to PWHT cracking. Intergranular fracture by coalescence of microvoids nucleated at grain boundary carbides. Table 2.2 Summary of reviewed Gleeble based PWHT tests 36

54 The reviewed articles show a very strong correlation between ductility and PWHT cracking susceptibility. Also, many of the previous researchers claim that strains induced by precipitation reactions are responsible for the reduction in ductility. Alloys that have slow rates of precipitation are supposed to be more resistant to PWHT cracking than alloys with fast precipitation rates. A test for PWHT cracking should quantify ductility and precipitation rate. Also, because the rate of precipitation is temperature dependent, the test should include different temperatures. Thus, a robust test for PWHT cracking should attempt to quantify time (rate), temperature, and ductility and the relationship between the three. 37

55 CHAPTER 3 OBJECTIVES The three main objectives of this investigation were: 1. To develop a PWHT cracking test for Ni-based alloys (Waspaloy, Alloy 718) which replicates the conditions experienced by the HAZ of a weldment. 2. Acquire weldability and PWHT data using a standardized test technique. 3. Compare results to previous tests for PWHT cracking 38

56 CHAPTER 4 EXPERIMENTAL APPROACH It should be noted that the author was not the initial researcher for this undertaking. The material to be used in the development of a test had already been acquired, samples had been machined, and preliminary testing had begun when the reporting researcher took on the project. In this regard some aspects of the experiment had already been defined. Ideally a PWHT cracking test should switch between force and stroke control so that forces can be precisely controlled during portions of the test and stroke can be controlled and force monitored during others. The Gleeble used for this, however, was not capable of switching between methods of mechanical feedback and therefore only one feedback mode could be used. Stroke was chosen as the means of mechanical loading. The force output from a transducer in series with the sample could then be monitored for any changes in the applied load while controlling the stroke. The difficulty in using stroke as a means of control is repeatability. For the same amount of jaw movement, a sample with no load prior to initiation of a test may develop a lower peak stress than a sample that has a 200kg load at the initiation of the test. To rectify the 39

57 problem, all samples must be under the same load prior to any programmed thermal loading or jaw movement. Preliminary tests by the initial researcher were conducted to determine the elevated temperature tensile properties of simulated HAZ s. Values for the NST, NDT, and DRT were obtained from another researcher using the same material [20]. With this information the first series of PWHT cracking tests was conducted. The first attempt at a PWHT cracking test consisted of the following steps: HAZ simulation to a peak temperature above the NDT (1280 C) with no mechanical loading; free cooling to room temperature with no loading; heating to a PWHT temperature (700 C to 816 C); manual adjustment of stroke to induce a slightly positive load; and finally, programmed stroke control to load the sample to approximately 70% of the UTS of the material at that temperature. Figure 4.1 shows the data acquired from one such test. As can be seen in the Figure 4.1, the force reached a maximum upon initial programmed stroke and then relaxed. Following relaxation, the force began to rise as the volumetric change associated with precipitation took place. The problems associated with this early test included a small force response for the aging reaction and non-programmed movement of the jaws, which caused a significant decrease in the tensile load. The data shown in Figure 4.1 was actually the exception and not the norm. 40

58 Sample 2A(1) 1. 40E E E E E E E E E+00 T_Actual Force Stroke E E E E E E+02 Time (sec.) -3.00E-02 Figure 4.1 Data acquired from preliminary PWHT cracking test 41

59 The problems encountered in the early test gave cause for rethinking the test procedure. It was decided that building residual stress on cooling from the peak HAZ would better replicate actual conditions than inducing a stress by straining the sample after heating to the PWHT temperature. In order to accurately model the residual stresses developed on cooling from the peak HAZ temperature, the simulation must build stress on the order of magnitude of the yield stress for the material. To correctly model the PWHT, which is applied to relieve stress, the axial load after the HAZ simulation must not be augmented with mechanical loading from the Gleeble hydraulic system during the PWHT cycle. Any additional loading should be the result of strain induced by the formation of precipitates. Temperatures used for the PWHT portion of the simulation should be similar to those used in industry. The range of PWHT temperatures should also coincide with those reported by previous researchers for comparison to their work. It was hoped that monitoring the force response during PWHT in the test described above would reveal time to crack initiation and time to failure (Figure 4.2). This information could be used to develop a C-curve on a temperature versus time plot for initiation of cracking and time to failure. After conducting many such tests it was determined the time to failure measurements would require tests which were too long and the time to initiation of cracking was erratic. Previous work has shown PWHT cracking in nickel base superalloys to be linked to a reduction in ductility [22,23,38]. Hot ductility tests performed after several different PWHT times could show any relationship between aging time and ductility. Additionally the tests could be used to correlate ductility and temperature. By combining the two 42

60 controlled variables, time and temperature, with the measured result, reduction in area, it would be possible to create a three-dimensional matrix of time, temperature, and ductility. Stress relaxation Failure Force Crack initiation Time Figure 4.2 Schematic of desired data from initial test 43

61 CHAPTER 5 EXPERIMENTAL PROCEDURE 5.1 Materials The two alloys chosen for this study were Waspaloy and Alloy 718. Both alloys are used in jet turbine engines. Waspaloy is a nickel-chromium alloy while 718 is a nickel-chromium-iron alloy. Waspaloy is a γ strengthened alloy which falls close to the 3% Al-6% Ti line (Figure 2.2) and has proven in the past to be susceptible to PWHT cracking [1,21,22]. Alloy 718 is γ strengthened alloy and has been used as a comparison because of its lower PWHT cracking susceptibility [1,22,36]. The chemical compositions of the materials used in this investigation are listed in Table 5.1. The Waspaloy bar used was received in the solution-annealed condition and contained bands of carbides distributed along the rolling direction in a matrix of equiaxed grains. When the Gleeble samples were machined from the bar the carbide bands ran parallel to the axis of the samples. The average grain size was 16.1µm. The microstructure of the as received Waspaloy material is shown in Figure

62 Element Waspaloy, wt% Alloy 718, wt% Cr Co Al Ti Mo Nb Fe W Cu Zr Ta < Mn V C B Si S < P < Mg Ni Balance Balance Table 5.1 Chemical compositions of Ni-base alloys tested The Alloy 718 used in this study was in the form of a wrought plate and was also supplied in the annealed condition. The microstructure was characterized by a distribution of both intergranular and intragranular bar-shaped δ phase. The average grain size was 59.3µm. The microstructure of the as received condition for the alloy 718 used in the experiment is shown in Figure

63 Figure 5.1 As received Waspaloy bar microstructure Figure 5.2 As received Alloy 718 microstructure 46

64 5.2 Sample Preparation The materials were machined into 4 inch long reduced section Gleeble bars with threaded ends. A schematic of the sample with dimensions is shown in Figure 5.3. The surface finish was equivalent to an 800 grit polish. Type K (chromel-alumel) thermocouple wires were percussion welded at the center of the gage section for temperature monitoring Ø 0.25 Ø All radii 5/32 Figure 5.3 Schematic illustration of sample dimensions 5.3 Equipment The Gleeble thermo-mechanical simulator was chosen for use in developing the PWHT cracking test. The machine used was a Gleeble 1500 with an analog controller. Cooling of the Gleeble was performed with a chiller (Affinity RAA-007C ) running distilled water through the transformer and heating jaws. The controller was linked to a 47

65 PC by means of an A/D converter. The PC ran the Gleeble Programming Language (GPL) and operated the controller for the simulations. Acquisition of data from three channels was made possible by the use of another PC and an IOtech data acquisition system. Millivolt outputs from the Gleeble sensors were read by an IOtech DBK8 voltage input card before being sent to an IOtech DaqBook for conversion to digital signals. The IOtech software used a user-defined linear equation to convert the incoming signal to the correct magnitude and units. The data was recorded, after multiplication, directly into a Microsoft Excel spreadsheet for later evaluation. During the course of this project some control problems were experienced and some tests were run on a Gleeble 1000 with built-in DSI data acquisition system. The same programs could be run, but the acquired data had to be converted from a delimited ASCII file into an Excel spreadsheet. 5.4 Environment All samples were run in a chamber which was evacuated to a pressure of 5x10-3 torr and back filled with argon gas until only a slight vacuum (~-50kPa) remained. The process was repeated a second time before running the sample to ensure an inert environment. 5.5 Gleeble Tests The test procedure was developed to show the relationship between ductility and the time a simulated HAZ was held at various PWHT temperatures. For the two alloys 48

66 tested, the Gleeble was used to simulate a weld HAZ in the center of the Gleeble bars. A matrix was devised such that the HAZ s would be tested at four different PWHT temperatures and four times. Tables 5.2 and 5.3 show the matrices of the PWHT times and temperatures used for Waspaloy and alloy 718 respectively. After the prescribed time at the PWHT temperature the samples were pulled to failure at the PWHT temperature. Samples not given any PWHT time (noted as 0 hour in tables 5.2 and 5.3) were used as references for the hot ductility of the simulated HAZ. The order of the runs was randomized and can be found in Appendix A. PWHT Temp. Time at PWHT (hours) 760 C C C C Table 5.2 Waspaloy PWHT temperature-time matrix PWHT Temp. Time at PWHT (hours) 668 C C C C Table 5.3 Alloy 718 PWHT temperature-time matrix 49

67 5.5.1 HAZ Simulation The HAZ portion of the test contained a heating portion in which the sample had no mechanical load applied and a cooling portion in which the sample was loaded in tension as it cooled. Each test began by placing the sample into Gleeble jaws. The inner faces of the copper jaws were always set 1 inch apart from each other. When used, stainless steel jaws were always given ¾ between inner faces. The sample was then manually loaded in tension to ensure the jaws were tight in the Gleeble and to remove any slack in the system. Prior to the initiation of the Gleeble program, the sample was manually returned to the unloaded state. Once the program was initiated, the operator was allowed one minute to manually adjust the mechanical load to 30 kg before the initiation of the HAZ thermal cycle. This was done as a means of assuring repeatability. All samples were heated at 100 C per second to the peak HAZ temperature and held for a short time before controlled cooling to room temperature. After the samples had cooled below 1100 C the stroke began moving the jaws apart. A tensile load developed in the sample as a result of stroke movement during cooling. The HAZ portion of the test concluded with the bar at room temperature and a tensile load close to the yield strength of the material in its as received condition. Table 5.4 shows the standard conditions for the HAZ portion of the tests for both alloys. 50

68 Waspaloy Alloy 718 Jaw Blocks Copper Stainless steel Peak HAZ Temp C 1184 C Time at Peak Temp. 20 seconds 15 seconds Stroke 0.045cm 0.045cm Stroke Rate mm/min mm/min Table 5.4 Standard conditions in HAZ simulation for both alloys HAZ PWHT Stroke Time Figure 5.4 Schematic illustration of thermal and mechanical control for HAZ and PWHT simulation 51

69 5.5.2 Postweld Heat Treatment Simulation PWHT simulation followed the HAZ simulation for those samples receiving PWHT. The jaws did not move for the PWHT simulation but remained at the position reached at the end of the HAZ simulation. The samples were heated to the desired PWHT temperature, at 100 C per second, and held there for the prescribed time. Figure 5.4 gives a graphical representation of the HAZ simulation followed by a PWHT simulation Hot Ductility Test The hot ductility portion of the test was conducted immediately after the prescribed PWHT simulation. Following the PWHT portion of the simulation, the sample s diameter was measured with Vernier calipers at the center of the gage section without removing the sample from the Gleeble. The sample was reheated to its PWHT temperature and the moveable jaw was stroked at a rate of 1 mm/min until the sample failed. Upon sample failure the test was stopped and the diameter of each half was measured at the fracture. Each half of the fractured sample was measured at three places as shown in Figure 5.5. The average of the six readings was recorded as the reduced diameter of the sample. 52

70 Figure 5.5 Schematic of reduction in area measurements 5.6 Sample Numbering A numbering system was devised to make sample identification simple. Samples were numbered with a series of four numbers separated by underscores. The number scheme for a sample was as follows: PP_TTT_t_n Where PP is the time at peak temperature in the HAZ portion of the simulation, TTT is the PWHT temperature in degrees Celcius, t is the time, in hours, at the PWHT temperature, and n is representative of sample number run at that temperature. So that a specimen marked: 20_816_4_3, is representative of the third sample held for 20 seconds at the peak HAZ temperature, post weld heat treated for 4 hours. 53

71 5.7 Metallography The center portion (gage section) of the samples was cut from the 4 bar with an abrasive cut off saw. This piece was sectioned axially and mounted in a thermosetting mold. In samples pulled to failure, one half was sectioned and mounted and the other kept for fractography. Samples to be examined in a SEM were mounted in Konductomet, which is electrically conductive, and all other samples were mounted in bakelite. After mounting, samples were ground through 600 grit SiC abrasive before continuing through 1 µm diamond and finishing with 0.05 colloidal silica. Mounted and polished samples were electrolytically etched in 10% chromic acid. Light optical microscopy was performed on a Nikon Epiphot microscope and pictures were digitally acquired by means of a Hitachi HV-C20U camera and PAX-IT software. Pictures were taken with the tensile axis oriented in the horizontal direction, parallel to the micron bar scale placed in the picture. In addition to optical imaging techniques, electron microscopy was performed on a Philips XL30-FEG SEM at The Ohio State University s campus electron optics facility. 5.8 Grain Size Measurement Average grain diameters were calculated using the mean intercept length procedure outlined in the ASM handbook (ASTM E 112) [40,41]. Digital pictures and PAX-IT software were used for making the measurements. 54

72 5.9 Fractography Samples that fractured by the hot ductility test had half of the sample reserved for fractographic analysis in an SEM. The half was cut in a Leco CM-15 abrasive cut-off saw to shorten it for easier mounting in the SEM stage. The shortened fracture sample was ultrasonically cleaned in a Fisher Scientific FS20 for 5 minutes. Following ultrasonic cleaning the fracture surface was rinsed with ethyl alcohol and dried. Fractography was performed in the same Philips XL30-FEG SEM used for electron microscopy of mounted samples Follow-up Testing Several tests were conducted after analysis of the data from the initial tests described above. The variables for those additional tests are shown below in table 5.5. Material Peak PWHT PWHT Temp HAZ Temp (sec) Temp (hr) Waspaloy Waspaloy Waspaloy Alloy Table 5.5 Controlled variables for additional tests 55

73 CHAPTER 6 RESULTS 6.1 Gleeble Feedback Acquisition Three channels of feedback were recorded from the Gleeble during the HAZ and PWHT simulations. These channels were: temperature at the thermocouple, axial load, and stroke distance. The temperature and stroke distance were recorded to make certain the test ran as programmed. The force measured by the load cell was the most critical data acquired. The force acquisition showed the augmented residual stress after cooling to room temperature from the HAZ simulation, the minimum force after thermal expansion caused by heating to the PWHT temperature, and the change in force as the sample precipitation hardened. The axial force measured by the transducer was converted to engineering stress with units of megapascals (MPa). Figure 6.1 gives an example of a typical data acquisition plot for the HAZ and PWHT simulation of a Waspaloy bar. The same three channels were measured for the hot ductility portion of the test. The useful information taken from the acquired hot ductility data was the yield strength, the tensile strength, and the time to failure. Figure 6.2 shows a typical data acquisition plot for the stress in the hot ductility portion of the test. 56

74 20_816_4_ :00:00 0:10:00 0:20:00 0:30:00 0:40:00 0:50:00 1:00:00 1:10:00 1:20:00 1:30:00 1:40:00 1:50:00 2:00:00 2:10:00 2:20:00 2:30:00 2:40:00 2:50:00 3:00:00 Stress (MPa) Temp ( C) 3:10:00 3:30:00 3:40:00 3:50:00 4:00:00 3:20:00 Time 2.00E E-01 Stroke (cm) Figure 6.1 Typical acquisition curves from HAZ and PWHT portion of test E E E E-02 Stress T_act Stroke 57

75 20_816_0_ :00:50 0:01:00 0:01:10 0:01:20 0:01:30 0:01:40 0:01:50 Stress (MPa) 0:02:00 0:02:10 0:02:20 0:02:30 0:02:40 Time Figure 6.2 Typical stress vs. time curve for hot ductility portion of test 58 58

76 6.1.1 HAZ Residual Force The repeatability of the development of residual stress in the HAZ portion of the test was important for comparison of the effects of PWHT temperatures on cracking susceptibility. Figures 6.3 and 6.4 show the acquired data for stress in Waspaloy and Alloy 718 respectively. The peak HAZ temperature was the same in each of the samples tested for the respective alloys. The slight rise in stress recorded for the Waspaloy samples was an anomaly related to the programming of the test. In the Alloy 718 samples the changes in slope occurring in the first two minutes are associated with yielding of the samples. Waspaloy Stress Response Stress (MPa) _0_1 816_0_1 870_0_1 895_0_1 0:00:00 0:00:10 0:00:20 0:00:30 0:00:40 0:00:50 0:01:00 0:01:10 0:01:20 0:01:30 0:01:40 0:01:50 0:02:00 0:02:10 0:02:20 0:02:30 0:02:40 0:02:50 Time 0:03:00 0:03:10 0:03:20 0:03:30 0:03:40 0:03:50 Figure 6.3 Residual stress developed in Waspaloy HAZ simulations (1240 C peak, mm/min stroke rate) 59

77 Alloy 718 Stress Response :00:00 0:00:30 0:01:00 0:01:30 0:02:00 0:02:30 0:03:00 0:03:30 0:04:00 0:04:30 0:05:00 0:05:30 0:06:00 0:06:30 0:07:00 0:07:30 0:08:00 Stress (MPa) 0:08:30 0:09:00 0:09:30 668_0_1 718_0_1 768_0_1 818_0_1 Time Figure 6.4 Residual stress developed in Alloy 718 HAZ simulations (1184 C peak, mm/min stroke rate) Stress Increase During PWHT Waspaloy samples typically showed increasing amounts of stress relaxation on heating as the PWHT temperature increased. An example of this is shown in Figure 6.5. This shows that the relaxation is due mainly to thermal expansion of the sample as it is heated to the PWHT temperature. Many of the Alloy 718 bars relaxed to no load and did not show any response for between 30 minutes and an hour (see Figure 6.6). Most Alloy 718 samples showed some delay before showing a rise in the tensile stress. Figures 6.7 and 6.8 show comparisons between the stress increase of Waspaloy and Alloy

78 samples after heating to the PWHT temperature. For both alloys, the minimum relaxed load obtained after initially heating the sample to its PWHT temperature was used as the zero point Stress (MPa) :00:00 0:00:05 0:00:10 0:00:15 0:00:20 0:00:25 0:00:30 0:00:35 0:00:40 0:00:45 0:00: Time (sec) Figure 6.5 Increasing stress relaxation as PWHT temperature is increased in Waspaloy 61

79 15_768_2_1 1.40E E E E E E E E E+02 0:00:00 0:10:00 0:20:00 0:30:00 0:40:00 0:50:00 1:00:00 1:10:00 1:20:00 1:30:00 1:40:00 1:50:00 2:00:00 2:10:00 Time 2.00E E-01 Stroke (cm) 1.00E E E E-02 Force T_act Stroke Force (kg) Temp ( C) Figure 6.6 Delayed force response in Alloy 718 specimen 62 62

80 :00:00 0:10:00 0:20:00 0:30:00 0:40:00 0:50:00 1:00:00 1:10:00 1:20:00 1:30:00 1:40:00 1:50:00 2:00:00 2:10:00 2:20:00 2:30:00 Stress (MPa) Waspaloy Alloy 718 Time Figure 6.7 Comparison of stress increase after relaxing from heating to PWHT temperature (Waspaloy C, Alloy C) Stress (MPa) Waspaloy Alloy :00:00 0:10:00 0:20:00 0:30:00 0:40:00 0:50:00 1:00:00 Time Figure 6.8 Comparison of stress increase after relaxing from heating to PWHT temperature (Waspaloy 816 C, Alloy C) 63

81 The rate of stress increase in the samples decreased in magnitude with increasing time as was shown in Figures 6.7 and 6.8. Plotting the stress against a logarithmic scale for time yielded a linear slope. Figures 6.9 and 6.10 are for the same samples plotted in Figures 6.7 and 6.8 but the time scale is logarithmic. It is interesting to note that the slope of the stress response for Alloy 718 appears to be steeper than that of Waspaloy in Figures 6.9 and This may be due to a difference in the amount of lattice mismatch between precipitates and γ matrix in the two alloys Stress (MPa) Waspaloy Alloy Time (s) Figure 6.9 Comparison of stress increase on log time scale (Waspaloy C, Alloy C) 64

82 Stress Waspaloy Alloy Time (s) Figure 6.10 Comparison of stress increase on log time scale (Waspaloy C, Alloy C) Strength The yield strength of each sample was recorded from the data acquired in the hot ductility portion of the test. Ultimate tensile strength was also recorded. The values are reported in Appendix B. Figures 6.11 and 6.12 give graphical representations of the measured yield stress for the test specimens. 65

83 Waspaloy Yield Strength Strength (MPa) PWHT Time (hr) 1 PWHT Temp ( C) Figure 6.11 Yield strength of Waspaloy measured in hot ductility portion of test 66

84 Alloy 718 Yield Strength Stregth (MPa) PWHT Time (hr) 2 PWHT Temp ( C) Figure 6.12 Yield strength of Alloy 718 measured in hot ductility portion of test 67

85 6.2 Ductility The ductility of the materials, measured by reduction in cross sectional area, is shown in Appendix B. This data was used to develop a multivariate polynomial for calculating the ductility as a function of PWHT temperature and time. The factors used in the model were PWHT temperature and time of PWHT. The collected data appeared to have a parabolic curve, so the model was chosen to be a second order polynomial. A spreadsheet for both alloys was created with values of time, time squared, temperature, temperature squared, and time multiplied by temperature for each sample. Using the LINEST linear regression tool in Microsoft s Excel software, a line was fit to the reduction in area measurements. The resulting output gave the intercept and slope for each of the five variables as well as the coefficient of determination for the fit. The coefficients of determination (R 2 ) for the Waspaloy and Alloy 718 surface plot polynomials are 0.92 and 0.91 respectively. The linear fit was actually a second order polynomial because some of the variables used in the linear estimation were exponents of the other variables (ie. time and time squared) Figures 6.13 and 6.14 show the surface plot created from the formula developed for each alloy. Figure 6.15 gives the regression plots for no PWHT and Figure 6.16 shows a comparison of the regression models at 3 hours. A sufficient number of Waspaloy samples allowed testing of the model. Figures 6.17 to 6.19 show the results of ductility measurements from a few extra tests conducted to determine the accuracy of the model. One Alloy 718 sample was tested with an 68

86 excessive HAZ peak temperature hold time to grow very large grains in the CGHAZ. The fit of that data point to the Alloy 718 regression is shown in Figure

87 Time (hr) Temp ( C) % 50.0% 45.0% 40.0% 35.0% 30.0% 25.0% RA 20.0% 15.0% 10.0% 5.0% 0.0% 50.0%-55.0% 45.0%-50.0% 40.0%-45.0% 35.0%-40.0% 30.0%-35.0% 25.0%-30.0% 20.0%-25.0% 15.0%-20.0% 10.0%-15.0% 5.0%-10.0% 0.0%-5.0% Figure 6.13 Ductility surface plot from regression analysis of Waspaloy data 70

88 90.00%-95.00% 85.00%-90.00% 80.00%-85.00% 75.00%-80.00% Time (hr) Temp ( C) % 90.0% 85.0% 80.0% 75.0% 70.0% 65.0% 60.0% 55.0% RA 50.0% 45.0% 40.0% 35.0% 30.0% 25.0% 20.0% 15.0% 70.00%-75.00% 65.00%-70.00% 60.00%-65.00% 55.00%-60.00% 50.00%-55.00% 45.00%-50.00% 40.00%-45.00% 35.00%-40.00% 30.00%-35.00% 25.00%-30.00% 20.00%-25.00% 15.00%-20.00% Figure 6.14 Ductility surface plot from regression analysis of Alloy 718 data 71

89 Reduction in Area 80% 70% 60% 50% 40% 30% 20% 10% 0% Temperature ( C) Waspaloy Alloy 718 Figure 6.15 Comparison of ductility behavior in Waspaloy and Alloy 718 regression models for no PWHT 80% 70% Reduction in Area 60% 50% 40% 30% 20% 10% 0% Temperature ( C) Waspaloy Alloy 718 Figure 6.16 Comparison of ductility behavior in Waspaloy and Alloy 718 regression models at 3 hours PWHT 72

90 Reduction in Area 50% 45% 40% 35% 30% 25% 20% 15% 10% 5% 0% Temp ( C) 20_800_2.5_1 Regression Figure 6.17 Actual sample data point plotted with regression model at 2.5 hours (Waspaloy 800 C for 2.5 hours) 50% 45% 40% Reduction in Area 35% 30% 25% 20% 15% 10% 5% 0% Temp ( C) 20_840_3_1 Regression Figure 6.18 Actual sample data point plotted with regression model (Waspaloy 840 C for 3 hours) 73

91 Reduction in Area 50% 45% 40% 35% 30% 25% 20% 15% 10% 5% 0% Temp ( C) Regression 2_816_1_1 Figure 6.19 Actual sample data point for 2 second peak HAZ temperature plotted with regression model for 20 second peak HAZ temperature (Waspaloy 816 C for 1 hours) 74

92 80% 70% Reduction in Area 60% 50% 40% 30% 20% Regression 60_718_2_1 10% 0% Temp ( C) Figure 6.20 Actual sample data point for 60 second peak HAZ temperature plotted with regression model for 15 second peak HAZ temperature (Alloy C for 2 hours) 75

93 6.3 Waspaloy Characterization Waspaloy Microstructures Preliminary testing in the development of the test procedure utilized a temperature above the NDT for the alloy. When held at PWHT temperatures for extended periods of time some of these specimens failed, but not at the centerline of the sample. Figures 6.21, 6.22, and 6.23 show a representative microstructure of those preliminary tests. The samples produced with the procedure described in the previous chapter failed at the center of the sample in a coarse grained region. These samples were heated to a peak HAZ temperature below the NDT of the material. The grain sizes for the Waspaloy heat affected zones produced during testing are shown in table 6.1. Condition Grain size (µm) As Received 16.1 ± C 72.5 ± C 168 ±23 Table 6.1 Waspaloy grain size measurements Small cracks were evident in cross sections of all Waspaloy specimens. There were some oval cavities visible at high magnification such as those shown in Figures 6.24 and 6.25, but the majority of the cracks appeared as wedge shaped cracks like those in 76

94 Figures 6.26 and This corresponded well with the fractures and cracks in the CGHAZ of the preliminary tests (Figure 6.28). An interesting feature found in several of the Waspaloy samples was the appearance of cracks at the intersections of annealing twins and grain boundaries (Figures 6.29 and 6.30). The other notable feature found was elongation of grains in the direction of tensile loading in samples tested at 895 C. An example of this is shown in Figure

95 Figure 6.22 Figure Figure 6.21 Representative microstructure of preliminary tests 78

96 Figure 6.22 Voids formed through centerline of preliminary sample Figure 6.23 Microcracks at grain boundaries in CGHAZ away from centerline of sample 79

97 Figure 6.24 Cavities at grain boundaries in Waspaloy Figure 6.25 SEM Image of cavities in Waspaloy 80

98 Figure 6.26 Wedge shaped cracks at boundaries in Waspaloy. (2 second peak HAZ, 816 C PWHT) Note transgranular fracture of some grains shown by arrows. Figure 6.27 SEM image of wedge crack at boundary in Waspaloy (boundaries highlighted by dashed lines) 81

99 Figure 6.28 Wedge cracks and intergranular fracture in a preliminary Waspaloy test sample twin Grain boundary crack Figure 6.29 Crack at interface of a twin and a grain boundary in Waspaloy (895 C PWHT) 82

100 Figure 6.30 Another crack at the intersection of a twin and grain boundary in Waspaloy (760 C PWHT) Figure 6.31 Grain elongation in tensile direction of Waspaloy sample tested at 895 C. 83

101 6.3.2 Waspaloy Fracture Morphology Preliminary tests, which failed away from the centerline of the sample in the CGHAZ, exhibited intergranular fracture surfaces as shown in Figure Higher magnification revealed evidence of micro-ductility (see Figure 6.33). Once the peak HAZ temperature and PWHT temperatures were set to those used in the testing matrix described in the experimental procedure, the fractures continued to exhibit intergranular fracture paths just as in the preliminary testing samples. The intergranular fractures occurred at all temperatures. Micro-ductility was evident on the fracture surfaces at all temperatures. There was, however, some indication of intragranular fracture on both the lowest and highest PWHT temperature samples. Figures 6.34 through 6.39 show fracture surfaces for different PWHT temperatures. The fracture surface shown in Figure 6.40 is of the Waspaloy sample held at the peak HAZ time for only 2 seconds rather than the normal

102 Figure 6.32 Intergranular fracture in preliminary tests conducted in Waspaloy Figure 6.33 Micro-ductility on intergranular fracture surface 85

103 Figure 6.34 Intergranular fracture surface with micro-ductility in 760 C PWHT Waspaloy sample Figure 6.35 Evidence of both intergranular and intragranular fracture in 760 C PWHT Waspaloy sample 86

104 Figure 6.36 Figure 6.36 Intergranular fracture surface in 816 C PWHT Waspaloy sample Figure 6.37 Higher magnification of apparently smooth grain surface in Figure 6.35 (note evidence of twin on fracture surface at arrows) 87

105 Figure 6.38 Intergranular fracture surface in 895 C PWHT Waspaloy sample Figure 6.39 Markings of micro-ductility and macro-ductility in 895 C PWHT Waspaloy sample 88

106 Figure 6.40 Waspaloy sample held at peak HAZ temperature for 2 seconds showing both intergranular and transgranular fracture. 6.4 Alloy 718 Characterization Alloy 718 Microstructure The grain size measurements for the Alloy 718 as received material and simulated HAZ s are provided in Table 6.2. The Alloy 718 samples failed at the center of the sample in a coarse grained region. Small cracks were evident in the interior of the 718 samples just as they were in the Waspaloy specimens. There was a fairly even distribution of cavities along grain boundaries and wedge shaped cracks at triple points. The exception to this was the samples tested at 818 C, which had grains that appeared to have been stretched and very 89

107 few interior cracks at boundaries. The sample held at the peak HAZ temperature for 60 seconds showed more wedge cracks than cavities. The sectioned fractures revealed intragranular fracture paths and what appear to be bent crystals. Figures 6.41 to 6.49 show the microstructures found in Alloy 718 samples. Condition Grain size (µm) As Received 59.3 ± C 84.8 ± C 207 ±44 Table 6.2 Alloy 718 grain size measurements 90

108 Cavities Wedges Figure 6.41 Cavities and wedges in Alloy 718. (718 C PWHT, 8 hours) Note lattice bending in some crystals Figure 6.42 Cavities along grain boundaries in Alloy 718 (718 C PWHT, 4 hours) 91

109 Figure 6.43 Grain elongation in Alloy 718 after hot ductility test (sample tested at 818 C) Twins Figure 6.44 Bent twins in deformed crystal of Alloy 718 sample tested at 818 C (note voids formed at globular NbC particles) 92

110 Figure 6.45 Wedge crack in sample held at HAZ peak temperature for 60 seconds Figure 6.46 Fracture in sample held at HAZ peak temperature for 60 seconds 93

111 Figure 6.47 Intragranular fracture path in Alloy 718 sample (718 C PWHT, 8 Hours) Figure 6.48 Detail of intragranular fracture in Alloy 718 sample 94

112 Figure 6.49 High magnification of distortion in lattice at fracture surface in Alloy 718 (718 C PWHT, 4 hours) 95

113 6.4.2 Alloy 718 Fracture Morphology The Alloy 718 samples did not always exhibit the intergranular fracture with micro-ductility that was evident in the Waspaloy samples. Those samples tested at 818 C showed a ductile fracture surface with deep dimples as shown in Figures 6.50 and Fractures in samples tested at lower temperatures showed mixed modes of fracture (see Figure 6.52). In regions where intergranular fracture did occur, it was not always ductile. As shown in Figures 6.53 and 6.54 the intergranular fractures in Alloy 718 showed different characteristics, sometimes even on the same grain. SEM fractography also showed evidence of grain distortion such as that in Figure The sample that was held for 60 seconds at the peak HAZ temperature showed much more ductile intergranular fracture on its surface than the other Alloy 718 samples. Figures 6.56 through 6.59 show the fracture surface of the 60-second HAZ sample. 96

114 Figure 6.50 Ductile dimples on surface of Alloy 718 specimen tested at 818 C Figure 6.51 Detail of deep ductile dimples on surface of Alloy 718 specimen tested at 818 C 97

115 Figure 6.52 Alloy 718 fracture surface showing both intragranular ductile fracture and intergranular fracture (718 C PWHT, 8 hours) Figure 6.53 Different surface characteristics exhibited in Alloy 718 intergranular fracture (718 C PWHT, 4 hours) 98

116 Figure 6.54 Different intergranular fracture characteristics on surface of single grain (718 C PWHT, 4 hours) Figure 6.55 Evidence of grain bending in Alloy 718 fracture (718 C PWHT, 4 hours) 99

117 Figure 6.56 Fracture surface of Alloy 718 sample held for 60 seconds at peak HAZ temperature and PWHT at 718 C for 2 hours Figure 6.57 Detail of ductile fracture in center of Figure

118 Figure 6.58 Detail of intergranular fracture shown in Figure 6.56 Figure 6.59 Higher magnification of intergranular fracture in Figure 6.58 revealing micro-ductility 101

119 CHAPTER 7 DISCUSSION 7.1 Procedure Development Waspaloy was used for most of the preliminary procedure development because it has been shown to be susceptible to PWHT cracking and its cracking behavior has been reported on extensively in the literature [1,21,22] Initiation of Stroke As stated previously, the goal of developing room temperature stress levels of the magnitude of the yield stress of the base material was critical to the simulation of PWHT cracking. Most models of for PWHT cracking assume that this stress level is present in the HAZ prior to the PWHT cycle. To accomplish this the Gleeble must stroke as the sample cools from the peak HAZ temperature. The timing of the stroke as well as the distance traveled must be chosen so that the sample does not yield or crack prior to the PWHT simulation. Several tests were conducted to determine the effect of stroking the Gleeble while cooling from the peak temperature in the HAZ simulation. The goal of these tests was to develop a simulated HAZ with tensile loads on the order of magnitude of the yield 102

120 strength of the material without creating any cracks or yielding. Initially, the same peak HAZ temperature (1280 C) used by the prior investigator was selected. The diameter of the samples was measured before and after the tests to check for yielding. After the sample had been run in the Gleeble and measured, it was sectioned, mounted, and polished to examine the interior for the presence of cracks. Samples in which the jaws began to move before the temperature had fallen below the DRT exhibited voids and cracks in the center of the sample as shown in Figure 7.1. These are apparently liquation cracks, since the sample was strained while liquid films were still present along grain boundaries at temperatures above the DRT. The test was subsequently modified so that stroke was not applied until the sample had cooled below the DRT. This insured that the presence of pre-existing cracks would not influence the subsequent PWHT simulation. 103

121 Figure 7.1 Voids and cracks formed along grain boundaries during the HAZ thermal cycle prior to PWHT simulation that resulted from stroking before the sample cooled below the DRT Peak HAZ Temperature After the change in the initiation of stroke was implemented, another series of samples was given the HAZ simulation and then heated to a PWHT temperature of 816 C. These samples also showed voids at the center of the sample, similar to those found in samples stroked before cooling below the DRT from the peak HAZ temperature (Figure 6.22). These voids were associated with the presence of liquid films. However, when these samples were examined, there were also cracks observed in a grain-coarsened region away from the center of the simulated HAZ (Figure 6.23). Some of the samples given this treatment failed through this coarse grained HAZ (CGHAZ) during the PWHT 104