METALLURGICAL CHARACTERIZATION AND WELDABILITY EVALUATION OF FERRITIC AND AUSTENITIC WELDS IN ARMOR STEELS THESIS

Transcription

1 METALLURGICAL CHARACTERIZATION AND WELDABILITY EVALUATION OF FERRITIC AND AUSTENITIC WELDS IN ARMOR STEELS THESIS Presented in Partial Fulfillment of the Requirement for the Degree Master of Science in the Graduate School of The Ohio State University By Matthew James Duffey Graduate Program in Welding Engineering The Ohio State University 2016 Master's Examination Committee: Professor Boian Alexandrov, Advisor Professor Antonio Ramirez

2 Copyright by Matthew James Duffey

3 ABSTRACT Armored steels are used in construction of armored vehicles for military applications in order to protect soldiers. Currently, these steels are welded using low-alloy consumables; however, welding with these consumables has led to hydrogen-induced cracking (HIC) under certain conditions. It is believed that welding these armored steels with austenitic stainless steel filler metals will mitigate the HIC issue. However, it is uncertain as to which filler metal will perform the best in terms of weldability in such welds. Three stainless steel filler metals were chosen for welding of these armored steels: Sandvik AXT, ER309LHF, and ER312. In order to determine which filler metal will produce the best welds, computational modeling, metallurgical characterization, and weldability testing was performed. Thermodynamic simulation modeling was performed to determine the solidification mode(s) that each weld exhibits. Metallographic examination was performed to determine microconstituents that exist, measure hardness of these features, and determine compositions (EDS analysis) of these features. The Tekken Test and Controlled Thermal Severity (CTS) Test was used to evaluate the weldability of the joints with respect to HIC. The combination of modeling and EDS analysis determined that only Armox 440 and Sandvik AXT weld combination exhibited austenitic solidification in a dilution that resulted in a cellular dendritic morphology. All other weld combinations either did not ii

4 experience austenitic solidification or experienced it in a planar growth region. Therefore, welds made with Armox 440 steel and Sandvik AXT filler metal is, in theory, the only weld that may be susceptible to solidification cracking. Each weld combination contained a coarse grain heat affected zone (CGHAZ), planar growth region, cellular dendritic weld metal, and weld swirls. The hardness of the CGHAZ exceeded the critical hardness for HIC to occur. The highest hardness' in each weld combination was located in the weld swirls, which could have trapped hydrogen in them. All weld combinations could experience HIC if sufficient hydrogen is present in weld swirls. Cracking occurred in many weldability samples; however, all welds with stainless steel consumables only experienced solidification cracking and not HIC. A majority of cracks were found in the weld metal along solidification grain boundaries and did not initiate at the location of maximum stress concentration of the particular weldability test. These cracks were identified as solidification cracks based on the dendritic morphology of the fracture surface. The extent of cracking was reduced by applying a preheat prior to welding. The CTS Tests that resulted in cracking had to be considered "invalid" because the weld crack lengths exceeded 5% of the throat thickness. Although relevant information was not acquired on the HIC susceptibility, the results revealed that in the case of highly restrained welds of armored steels, solidification cracking in the stainless steel weld metal can be a bigger problem than HIC. Although not predicted by modeling, all stainless steel weld combinations exhibited solidification cracking likely due to the high restraint. After thorough analysis, it has iii

5 been determined that ER309LHF is the most suitable (of the three tested filler metals) for welding of both Armox 440 and RHA. This stainless steel filler metal experienced the least severe cracking (if cracking occurred). In welds made with the low alloy steel welding consumables ER70S-6 and ER100, both HIC and solidification cracking occurred. It was found that ER100 is the more suitable consumable for welding Armox 440. iv

6 DEDICATION To my loving parents, Joseph and Marsha Duffey, who have unquestionably supported me and encouraged me to achieve my goals. v

7 ACKNOWLEDGEMENTS I would like to acknowledge and thank my advisor, Dr. Boian Alexandrov, for his guidance, advice, and encouragement throughout this project. I would also like to thank Dr. Antonio Ramirez for serving on my Master's Examination Committee. I am very grateful for all of the fellow graduate students and WJMG members, both past and present. Their friendship, advice, and wisdom have been vital to my success and I am very blessed to have had them through this entire project. I would also acknowledge American Engineering and Manufacturing Inc. (AEM) and Ma 2 JIC for sponsorship of this project. In particular, my sincere gratitude goes to John Lawmon for his advice and direction for this project. vi

8 VITA December 30, Born May B.S. Welding Engineering The Ohio State University Columbus, OH U.S.A 2014 to present... Graduate Research Associate The Ohio State University Columbus, OH U.S.A. Major Field: Welding Engineering FIELDS OF STUDY vii

9 TABLE OF CONTENTS ABSTRACT... ii DEDICATION... v ACKNOWLEDGEMENTS... vi VITA... vii FIELDS OF STUDY... vii TABLE OF CONTENTS... viii LIST OF TABLES... xiii LIST OF FIGURES... xiv CHAPTER 1: INTRODUCTION... 1 CHAPTER 2: BACKGROUND Introduction Armored Steels Alloying Elements Properties Welding Stainless Steels viii

10 2.3.1 Alloying Elements Austenitic Stainless Steels Constitution Diagrams Welding Metallurgy Weld Regions Weld Metal Boundary Types Dissimilar Metal Welds Solidification Weldability Solidification Cracking Hydrogen-Induced Cracking Carbon Equivalence Weldability Testing Welding Armored Steels with Stainless Steel Consumables Consumable Selection Joint Design Mechanical Properties Weldability CHAPTER 3: OBJECTIVES ix

11 3.1 Thermodynamic Prediction Metallurgical Characterization Weldability Testing CHAPTER 4: EXPERIMENTAL PROCEDURES Introduction Materials Thermodynamic Simulation Metallographic Characterization Sample Preparation Optical Microscopy Hardness Mapping Energy Dispersive Spectroscopy Weldability Testing Tekken Test CTS Test CHAPTER 5: RESULTS AND DISCUSSION Thermodynamic Simulations Armox Sandvik AXT Armox ER309LHF x

12 5.1.3 Armox ER RHA + Sandvik AXT RHA + ER309LHF RHA + ER Thermodynamic Simulations: Summary and Result Analysis Metallurgical Characterization Armox Sandvik AXT Armox ER309LHF Armox ER RHA + Sandvik AXT RHA + ER309LHF RHA + ER Metallurgical Characterization: Summary and Result Analysis Weldability Testing Armox Sandvik AXT Armox ER309LHF Armox ER Armox ER70S Armox ER xi

13 5.3.6 RHA + Sandvik AXT RHA + ER309LHF RHA + ER RHA + ER70S RHA + ER Weldability Testing: Summary and Result Analysis CHAPTER 6: SUMMARY AND CONCLUSIONS Thermodynamic Simulations Metallurgical Characterization Weldability Testing General Conclusions CHAPTER 7: RECOMMENDATIONS FOR FUTURE WORK REFERENCES xii

14 LIST OF TABLES Table 1: Minimum hardness requirements [2]... 7 Table 2: Minimum Charpy V-notch impact requirements [2]... 8 Table 3: Typical composition ranges for austenitic stainless steels Table 4: Solidification Type, Reaction, and Microstructure [7] Table 5: Parent metal composition ranges [42] Table 6: Hydrogen scales to be used with any arc welding process [42] Table 7: Range of validity for the CET method [42] Table 8: Susceptibility index grouping as a function of HD and P cm [42] Table 9: Minimum preheat and interpass temperature [42] Table 10: CTS test piece dimensions and tolerances [10] Table 11: Chemical Compositions of the Armored Steels (wt%) Table 12: Chemical compositions of the welding consumables (wt%) Table 13: Solidification Mode Summary Table 14: Summary of Hardness Values Table 15: Summary of Tekken Test Results Table 16: Summary of CTS Test Results Table 17: Average Weld Metal Dilutions from Each Weldability Test xiii

15 LIST OF FIGURES Figure 1: Relationship of solidification type to the pseudobinary phase diagram [7] Figure 2: Weld solidification cracking susceptibility as a function of composition based on Varestraint data [7] Figure 3: Suutala diagram [11] Figure 4: Diffusion coefficient of hydrogen as a function of temperature [12] Figure 5: Schaeffler Diagram of 1949 [14] Figure 6: Illustration of weld regions in a heterogeneous weld [15] Figure 7: Intermediate mixed zone features [20] Figure 8: Weld Metal Boundaries [22] Figure 9: Normal fusion boundary vs. dissimilar fusion boundary with Type II boundaries [26] Figure 10: Effect of temperature gradient in the liquid and solidification rate on solidification mode [23] Figure 11: Diffusion of hydrogen from weld metal to HAZ [35] Figure 12: Combined effect of stress intensity factor and hydrogen content on fracture surface [40] Figure 13: Effect of preheating on HIC [41] Figure 14: Carbon Equivalence vs. Combined Thickness to determine the Preheat Temperature [42] xiv

16 Figure 15: CE N correction with respect to the weld metal hydrogen content [42] Figure 16: CE N correction with respect to the weld heat input and CE IIW [42] Figure 17: Master curve for minimum preheat for the Tekken Test [42] Figure 18: Correction of necessary preheat of welding practice [42] Figure 19: Zone classification of steels [42] Figure 20: Tekken test setup [10] Figure 21: Measuring crack length and bead height [10] Figure 22: CTS Test setup [10] Figure 23: Jig used to position test assembly [10] Figure 24: Water bath arrangement [10] Figure 25: Sectioning of CTS test assembly [10] Figure 26: Armox Sandvik AXT Solidification Diagram Figure 27: Schaeffler Diagram connecting Armox 440 and Sandvik AXT Figure 28: Armox ER309LHF Solidification Diagram Figure 29: Schaeffler Diagram connecting Armox 440 and ER309LHF Figure 30: Armox ER312 Solidification Diagram Figure 31: Schaeffler Diagram connecting Armox 440 and ER Figure 32: RHA + Sandvik AXT Solidification Diagram Figure 33: Schaeffler Diagram connecting RHA and Sandvik AXT Figure 34: RHA + ER309LHF Solidification Diagram Figure 35: Schaeffler Diagram connecting RHA and ER309LHF Figure 36: RHA + ER312 Solidification Diagram xv

17 Figure 37: Schaeffler Diagram connecting RHA and ER Figure 38: Schaeffler Plot with all DMW combinations Figure 39: Armox Sandvik AXT [A] CGHAZ [B] Weld Swirl along top leg [C] Planar Growth [D] Weld Metal Figure 40: Armox Sandvik AXT [A] Macro View [B] Hardness Map overlaid on the Macro View Figure 41: Dilution determined using EDS Line Scan in Armox Sandvik AXT with a [A] Optical Microscopy Image [B] DIC Image [C] Dilution and Solidification Temperature Range vs. Distance Chart Figure 42: Armox Sandvik AXT Solidification Diagram overlaid with EDS results Figure 43: Armox ER309LHF [A] CGHAZ [B] Weld Swirls in the weld root [C] Planar Growth [D] Weld Metal Figure 44: Armox ER309LHF [A] Macro View [B] Hardness Map overlaid on the Macro View Figure 45: Armox ER309LHF Solidification Diagram overlaid with EDS results Figure 46: Armox ER312 [A] CGHAZ [B] Weld Swirl in the top leg [C] Planar Growth along the fusion boundary [D] Fusion Boundry and Weld Metal Figure 47: Armox ER312 [A] Macro View [B] Hardness Map overlaid on top of Macro View Figure 48: Armox ER312 Solidification Diagram overlaid with EDS results xvi

18 Figure 49: RHA + Sandvik AXT [A] CGHAZ [B] Weld Swirl along the top leg [C] Planar Growth [D] Weld Metal Figure 50: RHA + Sandvik AXT [A] Macro View [B] Hardness Map overlaid on Macro View Figure 51: RHA + Sandvik AXT Solidification Diagram overlaid with EDS results Figure 52: RHA + ER309LHF [A] CGHAZ and Weld Swirl in the weld root [B] Weld Swirl along the top leg [C] Planar Growth [D] Weld Metal Figure 53: RHA + ER309LHF [A] Macro View [B] Hardness Map overlaid on Macro View Figure 54: RHA + ER309LHF Solidification Diagram overlaid with EDS results Figure 55: RHA + ER312 [A] CGHAZ [B] Weld Swirl along top leg [C] Planar Growth [D] Fusion Boundary and Weld Metal Figure 56: RHA + ER312 [A] Macro View [B] Hardness Map overlaid on Macro View Figure 57: RHA + ER312 Solidification Diagram overlaid with EDS results Figure 58: Hardness Distribution Figure 59: EDS and Solidification Mode Comparison Figure 60: [A] Crack locations in Tekken samples [B] Crack locations in CTS samples Figure 61: Armox Sandvik AXT [A] Tekken Test Macro (No Preheat) [B] Tekken Test Crack (No Preheat) [C] CTS Test Macro (No Preheat) [D] CTS Test Crack (No Preheat) xvii

19 Figure 62: Armox ER309LHF CTS Test with No Preheat [A] Macro [B] Crack 129 Figure 63: Armox ER312 [A] Tekken Test Macro (No Preheat) [B] Tekken Test Crack (No Preheat) [C] CTS Test Macro (No Preheat) [D] CTS Test Crack (No Preheat) Figure 64: Armox ER70S-6 Tekken Test with No Preheat Figure 65: Armox ER100 Tekken Test with No Preheat Figure 66: RHA + Sandvik AXT CTS Test with No Preheat [A] Macro [B] Crack Figure 67: RHA + ER309LHF [A] Tekken Test Macro (No Preheat) [B] Tekken Test Crack (No Preheat) [C] CTS Test Macro (No Preheat) [D] CTS Test Crack (No Preheat) Figure 68: RHA +ER312 Tekken Test (No Preheat) [A] Macro [B] Cracks Figure 69: RHA + ER70S-6 Tekken Test with No Preheat Figure 70: RHA + ER100 [A] Tekken Test with No Preheat [B] CTS Test with No Preheat Figure 71: [A] Armox Sandvik AXT (No Preheat) WM crack [B] Fracture surface of Armox Sandvik AXT WM crack [C] RHA + ER309LHF (No Preheat) WM crack [D] Fracture surface of RHA + ER309LHF WM crack [E] RHA + ER312 (No Preheat) WM crack [F] Fracture surface of RHA + ER312 WM crack Figure 72: [A] RHA + ER70S-6 (No Preheat) WM crack [B] Fracture surface of RHA + ER70S-6 (No Preheat) WM crack [C] RHA + ER100 (Preheat) WM crack [D] Fracture surface of RHA + ER100 (Preheat) WM crack xviii

20 Figure 73: Crack Fracture Surface Change in Armox ER70S-6 [A] Intergranular [B] Intergranular transitioning to quasi-cleavage at the fusion boundary [C] Quasicleavage near the weld swirl [D] Microvoid Coalescence in the WM xix

21 CHAPTER 1: INTRODUCTION Armored vehicles utilize high-strength armor steel for strength and impact resistance. Conventionally, these armored steels have been welded using low-alloy consumables. However, these welded joints were susceptible to hydrogen-induced cracking (HIC). For HIC to occur, three conditions are required: 1. sufficient restraint 2. susceptible microstructure and 3. sufficient hydrogen. If one of these conditions can be eliminated, then HIC can be avoided. These vehicles are extremely thick and experience high stresses, which makes it difficult to minimize the high levels of restraint. These vehicles are also very large and a post-weld heat treatment (PWHT) cannot be applied, which leads to the inability to eliminate the susceptible microstructure produced from welding. The first two conditions cannot be avoided in these vehicles and it, therefore, means that hydrogen must be eliminated to avoid HIC. Low hydrogen practices can be implemented to minimize the hydrogen present. This includes baking and drying welding consumables, preheating the base materials, and cleaning plate surfaces prior to welding. However, there still may be sufficient hydrogen that diffuses to the heat-affected zone (HAZ) and leads to HIC. One way to minimize the hydrogen that reaches the HAZ is to use an austenitic stainless steel welding consumable. Hydrogen diffuses very slowly in austenite and becomes trapped in the weld metal prior to diffusing to the HAZ. Austenitic stainless steels are 1

22 not susceptible to HIC and, since hydrogen cannot diffuse to the HAZ, the armored steel is now no longer susceptible to HIC. Although austenitic stainless steels can assist in avoiding HIC, other weldability issues, such as solidification cracking, can arise from these consumables. The correct austenitic stainless steel must be chosen that has minimal susceptibility to solidification cracking and assists in eliminating HIC. This research has an underlying goal to determine which stainless steel consumable is the most effective in welding of two different armored steels. The three stainless steel consumables that were tested included Sandvik AXT (~ER307), ER309LHF, and ER312. The two armored steels that were tested include Armox 440 and RHA (Mil Standard Class 1). To evaluate which stainless steel, thermodynamic simulations, metallurgical characterization, and weldability testing techniques were performed. The thermodynamic simulations evaluated the solidification mode and solidification temperature range of various weld dilution levels. Metallurgical characterization techniques included optical microscopy, energy dispersive spectroscopy (EDS), and hardness mapping. These techniques aimed to evaluate compositions, hardness's, and characteristic features of different structures found in each weld. Weldability testing was comprised of Tekken and Controlled Thermal Severity (CTS) testing to assess the susceptibility to HIC. 2

23 CHAPTER 2: BACKGROUND 2.1 Introduction The military utilizes armored steels when making vehicles because of the high strength and the blast resistance of the material. Welding of these steels can often times lead to failure in the weld region. Improper consumable selection and/or weld parameters are the leading causes of such failures. Along with making good welds, it is also vital that properties of the armored steel are maintained through the welding process. Hydrogen-induced cracking (HIC) is a common issue when joining armored steel. Austenitic stainless steel filler metals are often used to join high-strength steels to mitigate the susceptibility to HIC. However, the use of austenitic stainless steel filler metals may present other weldability challenges, such as solidification cracking and/or reduced strength. Currently, there is a need to determine which stainless steel consumable is the most efficient in joining of the armored steels. In this chapter, fundamental concepts related to armored steels, stainless steels, and weld metallurgy will be presented. The mechanisms which cause weld cracking will also be discussed and various methods of evaluating the weldability of an alloy system will be described. 3

24 2.2 Armored Steels This section contains the current knowledge of armored steels. Understanding armored steels is vital in joining them while still being able to maintain their properties Alloying Elements High-strength low alloy (HSLA) steels, such as armored steels, were developed to have high strength-to-weight ratios. This can be achieved by selecting the proper alloying elements, which slows the austenite transformation upon cooling. This results in a lower pearlite transformation temperature and the pearlite will be very fine [1]. This section discusses the role each alloying element plays on the properties of armored steels. Carbon Carbon is the primary strengthener in HSLA steels. Carbon increases the tensile strength and hardness of steel. Typical Carbon contents range from wt%. Phosphorus Phosphorus is added to strengthen the ferrite in HSLA steels. This can be done in low alloy steels without inducing cold brittleness. The combined Phosphorus and Carbon content should not exceed 0.25 wt% or cold brittleness could be an issue [1]. Phosphorus contents range from wt%. Molybdenum Molybdenum can be effective as a strengthening element. It can also be used to increase hardenability, reduce susceptibility to temper embrittlement, and enhance 4

25 elevated temperature properties in quenched and tempered grades of HSLA steels [1]. Molybdenum contents range from wt%. Manganese Manganese is a major strengthening element when more than 1 wt% is present. Manganese is also added as a Sulfur "getter" and forms MnS. Usual potency of Manganese in HSLA steels is wt% [1]. Chromium Chromium is also a good strengthener. Chromium, in combination with other elements, can assist in preventing corrosion by forming an oxide layer on the surface of the steel. However, the Chromium alone does not prevent corrosion due to the low amount [1]. Chromium contents range from wt%. Silicon and Aluminum Silicon and Aluminum are added to HSLA steels to promote fine grain size and kill the steel. Sometimes, Silicon is added to promote ferrite formation as well [1]. Typical Silicon and Aluminum contents range from wt%. Copper Copper increases the steel's resistance to corrosion. Often times, Copper and Phosphorus are used in combination to further enhance corrosion resistance [1]. Typical Copper contents range from wt%. 5

26 Nickel Nickel plays a similar role as Copper. Nickel strengthens and increases corrosion resistance. However, this is not to the same degree as Copper. The amount of Nickel ranges from wt%. Other elements Other elements, such as Vanadium, Titanium, and Zirconium, are added to HSLA steels to promote deoxidation. Boron is often added to improve hardenability Properties Army standards divide armored steels into four separate classes based on their performance. Class 1 requires resistance to penetration. Class 2 requires resistance to shock. Classes 1 and 2 are steels that have a carbon equivalence (CE) less than 0.80 wt% for plates less than 4 inches thick and 0.90 wt% for plates greater than 4 inches thick. Class 3 is for evaluation only and cannot be used on military vehicles. Class 3 does not have a CE limit. Class 4 is for maximum resistance to penetration and is broken into two categories. Class 4a steels have a CE exactly the same as classes 1 and 2. Class 4b has a CE less than 0.60 wt% [2]. The most important property of armored steels is ballistic protection, which is the ability to stop a projectile at a specific velocity [3]. The ballistic protection level is directly related to the plate thickness (and weight). If the ballistic level is very high, then the weight is in excess and the armored steel may not be suitable for certain applications. It is important to achieve a reasonable weight to cost ratio while still maintaining the ballistic performance. 6

27 Other general properties of armored steels include high hardness and low toughness [3]. The high hardness is the principal characteristic of classifying a material as armor. Table 1 displays the minimum hardness requirements for the class of armored steel and plate thickness and Table 2 provides the minimum Charpy V- notch impact requirements for a given hardness and plate thickness [2]. These are usually results of the composition of the steel. However, these properties can lead to a variety of structural issues, with the most common being structural cracking. Welding of these steels can be difficult due to these properties. Table 1: Minimum hardness requirements [2] 7

28 Table 2: Minimum Charpy V-notch impact requirements [2] Welding Armored steels in warfare have been traced back to the fifteenth century. Prior to World War II, armored vehicles were joined using bolts and rivets. However, two issues arrived from these joints: 1. the flying bolt and rivet heads were extremely dangerous for the crew and 2. the riveted vehicles allowed bullets to enter the vehicle at the joint line [4]. By welding the armored vehicles, these issues were eliminated in addition to increasing labor efficiency, decreasing machining requirements, and reducing weight. 8

29 Common welding processes used to join armored steels include gas tungsten arc welding, gas metal arc welding, and shielded metal arc welding. In order to avoid cracking during welding, a high preheat and interpass temperature is often times employed. These steels often lose some of their strength after welding and require a post-weld heat treatment (PWHT). However, a majority of these steels are utilized for high energy armor and light weight applications for defense purposes. A PWHT is not possible in these applications. Therefore, the only PWHT possible is a low temperature tempering for stress relieving [5]. One of the primary issues with welding armored steels is the possibility of HIC in the weld heat-affected zone (HAZ). This is the result of high (tensile) restraint, a susceptible microstructure, and the presence of hydrogen from welding. There are multiple aspects that can be controlled in order to avoid HAZ cracking. One aspect is the plate composition, which results in a certain carbon equivalent. If the carbon equivalent is kept below 1, then HIC is unlikely to occur. A second aspect is the hydrogen content, which, if below 5 ml/100 g, generally does not result in HIC. A third aspect is the welding process, which generally utilizes a preheat in order to reduce residual stresses, eliminate moisture from the joint, and decrease the hardness of the steel [4]. One susceptible microstructure for HIC in steels is martensite. While martensite forms upon welding, the martensite start (M s ) temperature can be the deciding factor if cracking will occur or not [6]. The M s temperature is driven by the steel composition. If M s is low (generally below 325 C), then the resultant microstructure 9

30 is often acicular martensite and spheroidal carbides. If M s is high (generally above 395 C), then the resultant microstructure is often coarse martensite. As the M s temperature decreases, the hardness of the steel weldments increases and the susceptibility to delayed cracking increases [6]. Thus, a higher M s is desirable in order to increase the odds of avoiding any delayed cracking. Solidification cracking is typically not an issue due to the ferritic (BCC) solidification. Solidification cracking or weld metal HIC can be an issue in the weld metal depending on the consumable chosen. However, these issues can be avoided by proper consumable selection and process variables. Consumables with compositions that result in a microstructure that has a mixture of austenite and ferrite are generally not susceptible to solidification cracking. Process variables include the joint setup and weld pool profile, both of which can induce high strain on the weld and lead to solidification cracking [4]. 2.3 Stainless Steels This section contains the current knowledge of stainless steels, which are commonly used in industry. To weld these materials successfully, it is vital to understand their metallurgy and properties Alloying Elements Stainless steels are iron-based alloys, with the Iron content ranging from wt% of the composition [7]. To consider a steel "stainless", a minimum of approximately 10.5 wt% chromium must make up the composition. In ferritic and martensitic grade stainless steels, the major alloying elements are Chromium and 10

31 Carbon. The addition of Nickel is common in austenitic and duplex grade stainless steels [7]. Most stainless steels contain a combination of Manganese, Silicon, Molybdenum, Niobium, Titanium, Aluminum, Copper, Tungsten, Nitrogen, and others. This section discusses the role each alloying element plays on the properties of stainless steels. Chromium Chromium is added to the composition mainly to increase the corrosion resistance of the steel. Chromium forms a stable (Fe,Cr) 2 O 3 on the steel surface, which protects the steel from oxidizing environments. The addition of Chromium promotes ferrite formation. From the Iron-Chromium phase diagram, if the amount of chromium exceeds 12 wt%, then the alloy will be fully ferritic [7]. Increasing the Chromium content will also aid in retaining ferrite in non-ferritic grades, such as martensitic, austenitic, and duplex stainless steels. Chromium is also a strong carbide former. A common Chromium-rich carbide is M 23 C 6, where "M" is primarily Chromium but can contain Iron and Molybdenum [7]. Other carbides, such as Cr 7 C 3 and M 23 (C,N) 6 are possible along with some nitrides, such as Cr 2 N [8]. Nickel Nickel is commonly found in higher quantities in austenitic grades because it is an austenite promoter. The austenitic phase field (on the Fe-C phase diagram) is expanded with highly quantities of nickel, which leads to the austenite phase being stable at room temperature and below [7]. 11

32 Manganese Manganese was first added to stainless steels to avoid hot shortness during casting. This is a version of solidification cracking linked to the formation of Iron-Sulfide eutectic at low temperatures. Manganese combines with Sulfur much easier than Iron and forms a Manganese Sulfide that eliminated the hot shortness problem [7]. Manganese is an austenite promoter, although it is not a strong promoter compared to other elements. Silicon Small amounts of Silicon are usually present in stainless steels (usually wt%). The primary role of silicon is to act as a deoxidizing agent during melting. Aluminum can be used as a deoxidizing agent in place of Silicon if desirable. Some austenitic grade stainless steels have poor fluidity when in the molten state. The addition of Silicon has also been shown to improve fluidity of molten steel. Therefore, Silicon is usually added to weld filler metals [7]. Molybdenum Molybdenum plays a different role in stainless steels depending on the grade. In ferritic, austenitic, and duplex grades, Molybdenum is added to improve corrosion resistance, specifically pitting and crevice corrosion. In austenitic grades, Molybdenum improves the strength of the alloy at elevated temperatures [8]. In martensitic grades, Molybdenum forms carbides, which leads to secondary hardening of the alloy. Molybdenum is a ferrite promoter, and is generally undesirable in 12

33 martensitic grades because ferrite can reduce the toughness and ductility at room temperature [7]. Other Elements Chromium and Molybdenum are not the only carbide forming elements found in stainless steels. Elements, such as Niobium, Titanium, Tantalum, and Vanadium, are also added to form carbides. Niobium and Titanium form carbides that reduce the susceptibility of intergranular corrosion in austenitic grades. Tungsten, Tantalum, and Vanadium form carbides that increase the strength of the steel at elevated temperatures. Some elements are added to promote precipitation reactions that harden the steel. These elements include Aluminum, Titanium, Copper, and Molybdenum. Carbon and Nitrogen are the two most common interstitial elements found in stainless steels Austenitic Stainless Steels There are more than 250 different stainless steels. In order to differentiate the alloys, they have been divided up into five grades based on their crystal structure. These five groups include martensitic, austenitic, ferritic, duplex, and precipitationhardened stainless steels. Each grade experiences special mechanical properties. When welding with stainless steels, the application determines which grade (and specific alloy) to use. The most commonly used group of stainless steels is austenitic. Composition Austenitic stainless steels are comprised of elements that promote austenite and are thermo-mechanically processed so that the primary microstructure is austenite. 13

34 Typical compositions of austenitic stainless steels are found in Table 3. Austenitic stainless steels are comprised of the 200 and 300 series alloys as designated by the American Iron and Steel Institute (AISI) [7]. Table 3: Typical composition ranges for austenitic stainless steels Element Composition Range (wt%) Cr Ni 8-20 Mn 1-2 Si C Mo 0-2 N Ti + Nb Microstructure Morphology The microstructure evolution of austenitic stainless steels depends on the Cr eq to Ni eq ratio (i.e. depends on the ratio of ferrite to austenite promoting elements). There are four distinct solidification and solid-state transformation possibilities exhibited by austenitic stainless steels: A, AF, FA, and F. These reactions are summarized in Table 4 and are related to the Fe-Cr-Ni phase diagram as shown in Figure 1 [7]. 14

35 Table 4: Solidification Type, Reaction, and Microstructure [7] Figure 1: Relationship of solidification type to the pseudobinary phase diagram [7] Type A solidification occurs when the microstructure is fully austenitic at the end of solidification and remains that way upon cooling to room temperature. Type AF solidification occurs when the some ferrite forms at the end of the primary austenite solidification via a eutectic reaction [7]. The microstructure is characterized by a 15

36 majority of austenite with ferrite along the grain boundaries. Type FA solidification occurs when some austenite forms at the end of primary ferrite solidification. Upon cooling to room temperature, the ferrite becomes unstable and is partially consumed by the austenite, which results in a skeletal or lathy morphology [9]. Type F solidification occurs when solidification is completely ferrite. Upon cooling below the ferrite solvus, austenite will form within the microstructure and the final microstructure will have an acicular ferrite or ferrite with Widmanstätten austenite morphology [9]. Solidification Cracking Austenitic stainless steels are often utilized in welding; however, they possess a wide range of weldability issues if not welded properly. One issue includes solidification cracking. The susceptibility to solidification cracking is a function of composition (shown in Figure 2), which determines the solidification type. Austenitic stainless steels with type A solidification have high susceptibility to solidification cracking whereas austenitic stainless steels with type FA solidification (and a skeletal ferrite morphology) have low susceptibility to solidification cracking. Alloys that solidify as primary ferrite are the least susceptible to solidification cracking because of the presence of a two-phase austenite + ferrite mixture along the solidification grain boundaries at the end of solidification. This minimizes wetting of liquid films and results in a more tortuous boundary along which cracks must propagate [10]. 16

37 Figure 2: Weld solidification cracking susceptibility as a function of composition based on Varestraint data [7] Another attribute that increases susceptibility to solidification cracking is the impurity content, specifically Sulfur and Phosphorus. As the impurity content increases, the susceptibility to solidification cracking increases in type A and AF solidification mode alloys [11]. As solidification proceeds, the impurities segregate to the liquid and lower the solidus temperature. Thus, liquid remains at a lower temperature and the likelihood for solidification cracking increases. To assist in predicting solidification cracking susceptibility, a Suutala diagram (Figure 3) can be used. This diagram depicts the importance of composition and compares Cr eq /Ni eq to the impurity content. 17

38 Figure 3: Suutala diagram [11] Hydrogen-Induced Cracking Austenitic stainless steel filler metals are often used when welding high strength steels because of the low diffusion coefficient of hydrogen in the austenite [12]. Figure 4 illustrates the diffusion coefficient of hydrogen in austenitic materials (as a function of temperature) is much lower than in ferritic materials. The hydrogen is trapped in the austenitic weld metal and cannot diffuse to the HAZ to cause cracking. Additionally, the good ductility of the austenitic weld metals does not allow residual stresses to accumulate, which reduces the susceptibility of hydrogen-induced cracking [13]. 18

39 Figure 4: Diffusion coefficient of hydrogen as a function of temperature [12] Constitution Diagrams Knowing the microstructure of the weld metal makes alloy selection and properties much easier to determine. With stainless steels, constitution diagrams are used to determine the microstructure based on the austenite and ferrite stabilizing elements. Anton Schaeffler [14] developed a constitution diagram, known as the Schaeffler Diagram, which utilizes Chromium and Nickel equivalent formulas to determine the microstructure. The Nickel equivalent is calculated with austenite stabilizing elements using equation 1. The Chromium equivalent is calculated with ferrite stabilizing elements using equation 2. 19

40 1) Ni eq = Ni + 30C + 0.5Mn 2) Cr eq = Cr + Mo + 1.5Si + 0.5Cb Once the chromium and nickel equivalents have been calculated, the stainless steel alloy can be plotted on the Schaeffler Diagram (Figure 5) [14]. The area which the alloy is plotted shows the phase(s) that should be present. The diagram is mostly intended for stainless steel weld metal. However, the parent material can be plotted on the diagram and a tie-line between the two pure alloys can be made. All of the potential dilutions lie along this line. Any phase regions that the line intersects represents a potential phase that could be found in welds between the two alloys. Figure 5: Schaeffler Diagram of 1949 [14] 20

41 Since 1949, many constitution diagrams have been developed to more accurately predict the microstructure of stainless steel welds. Although the Schaeffler diagram is less accurate than the recently developed diagrams, it is much broader than the new diagrams. This allows for more weld combinations and potential dilutions to be analyzed. 2.4 Welding Metallurgy This section contains the current knowledge of welding metallurgy as it pertains to dissimilar metal welds. The metallurgy background is vital to understand the weldability of dissimilar metal welds and the weldability testing methods Weld Regions Welds can be divided into multiple regions based on their microstructural characteristics. These regions were originally investigated by Savage et al. [15] The regions of a weld include a true heat-affected zone (HAZ), a partially melted zone (PMZ), an unmixed zone (UMZ), and a composite region. The schematic in Figure 6 shows each weld region and its location in the weldment. 21

42 Figure 6: Illustration of weld regions in a heterogeneous weld [15] The true HAZ is the region in the base metal that does not melt, but undergoes microstructural changes in the solid-state due to the high temperature of welding. The true HAZ starts in base metal that experienced the A 3 temperature and ends in base metal that experienced the A 1 temperature. The PMZ contains a mixture of solid and liquid. Due to compositional variations, especially at the grain boundaries, localized melting can occur below the equilibrium solidus. The UMZ is a thin region located in the fusion zone. This region comprises of melted base metal that does not mix with the filler metal. The composite region makes up the remaining fusion zone, and comprises of mixture of melted base and filler metal. In dissimilar metal welds, there is a transition region that exists between the UMZ and the composite region. The transition region has a compositional gradient from 100% base metal to 100% filler metal. The width of the region varies depending on the alloys used, heat input, and welding process. 22

43 Heat-Affected Zone The heat affected zone (HAZ) is defined as the portion of the base metal whose properties and/or microstructure have been altered by the heat of welding [16]. All fusion welds and most solid-state welds have HAZ's due to high heat inputs. Carbon steels undergo allotropic transformations that control their microstructural alterations. The composition of each steel determines the precise liquidus and solidus temperatures, along with the allotropic phase transformation temperatures. In welding of steels, the HAZ is generally the base metal region that experiences temperatures ranging from the Ac 1 temperature to the solidus temperature [16]. HAZ's are generally divided into two sections: 1. coarse grain heat-affected zone (CGHAZ) and 2. fine grain heat-affected zone (FGHAZ). The CGHAZ is the section of the HAZ adjacent to the partially melted zone. This section experiences the highest temperatures without melting and has grains that are very large and coarse. The FGHAZ is the section of the HAZ furthest from the partially melted zone and adjacent to the unaffected base metal. This section experiences the lowest temperatures required for allotropic transformations and has grains that are very fine. Unmixed Zone The unmixed zone (UMZ) is the region of the fusion zone directly along the fusion boundary [17]. UMZ are often found in arc welds as opposed to high-energy density welds likely due to higher heat inputs, shallower fusion boundary temperature gradients, and less weld pool stirring. In regions of arc welds that have fluid flow that 23

44 is strong, the UMZ is "mixed" into the weld metal. Other regions, that experience lethargic fluid flow, the UMZ is often distinct. The temperature gradient at the fusion boundary affects the width of the UMZ because it affects the distance over which liquid base metal is present [17]. Often, due to the UMZ being so narrow (relative to the other regions), it is assumed that the UMZ is irrelevant. However, in many dissimilar metal welds, the mechanical and corrosion properties of the UMZ vary from the base and filler metals. In dissimilar metal welds between carbon steel and stainless steel, Doody [18] found that an intermediate mixed zone (IMZ) existed on the carbon steel side. Illustrated in Figure 7, the IMZ contained high hardness (martensite) features known as beaches, peninsulas, and islands. Rowe et al. [19] found that hydrogen from the arc can be trapped in these thin martensitic features, which can lead to HIC. These features possessed an intermediate composition between the carbon steel and bulk weld metal, although the composition was primarily made up of the carbon steel. These features are made up primarily of the carbon steel because the liquid carbon steel along the fusion boundary either did not or only partially mixed with the bulk weld metal [20]. 24

45 Figure 7: Intermediate mixed zone features [20] Partially Mixed Zone The partially melted zone (PMZ) is a transition region between 100% melting in the fusion zone (UMZ at the fusion boundary) and 100% solid region of the weld (HAZ) [21]. For pure metals, there is no PMZ because there is no solidification temperature range. For isotropic alloys (one with composition variations due to segregation, etc.), the PMZ represents the temperature range between the liquidus and solidus of the alloy. Segregation increases the "effective" melting temperature range of the base material. The difference in the liquidus and "effective" solidus temperatures generally describes the extent of the PMZ. In reality, solute distribution is not uniform in the base metal and additional segregation will likely occur during the weld thermal cycle [17]. This leads to local composition variations in the HAZ 25

46 adjacent to the fusion boundary. This region will experience melting at temperatures below that of the bulk microstructure Weld Metal Boundary Types A variety of different boundaries exist in austenitic weld metal during and following solidification. The weldability of austenitic stainless steels has a strong dependence on the characteristics of these boundaries as various cracking mechanisms occur along these boundaries. The three boundary types are solidification grain boundaries (SGBs), solidification subgrain boundaries (SSGBs), and migrated grain boundaries (MGBs) [22]. These boundaries are depicted in Figure 8. Figure 8: Weld Metal Boundaries [22] 26

47 Solidification from welding usually results in epitaxial nucleation of solid material from the base metal, which leads to columnar and dendritic grains in the weld metal [13]. SSGBs are considered low angle boundaries because of the small misorientation between dendrites [23]. SSGBs transpire in between the growing dendrites. SGBs are often found along the centerline of welds due to competitive growth of dendrites in different directions. Considerable crystallographic misorientation of SGBs is possible due to the dendrites growing from different directions. MGBs are SGBs that have migrated at high temperatures below the solidus. Migration of the SGB occurs in order to reduce the grain boundary free energy, which is accomplished by reducing the grain boundary area. The crystallographic misorientation across MGBs is usually in excess of 30 [22] Dissimilar Metal Welds Welds made with filler metals with different composition than the base metal are known as dissimilar metal welds (DMW's). DMW's are often used in applications where a transition in mechanical properties and/or performance in service is required. Although physical and/or chemical properties can be achieved, macrosegregation occurs near the fusion boundary in DMW's and, often times, degrades the weld quality [24]. When using stainless steels as the filler metal, the first pass of the weld is the most critical in terms of weld metal constitution because dilution effects are the greatest in this pass. It is desirable to obtain a dilution that results in a microstructure of mostly 27

48 austenite with a small amount of ferrite. This microstructure results in a low susceptibility to solidification cracking and the weld metal will be sufficiently ductile [7]. A transition regions lies along the fusion boundary of DMW's. This is the region that lies between the base metal and bulk weld metal. The microstructure in this region may vary from the bulk weld metal. In order to predict the microstructure of the weld metal and transition region in DMW's involving stainless steels, constitution diagrams, such as the Schaeffler Diagram, are used. Microstructure evolution along the fusion boundary of dissimilar metal welds can be complex. In cases where, at the melting temperature, the base metal and weld metal have different crystallographic orientations, then normal epitaxial growth may be suppressed and the formation of a Type II boundary can occur [7]. These boundaries run approximately parallel to the fusion boundary [25] as shown in Figure 9 [26]. Upon solidification, the Type II boundary grows some distance, becomes unstable, and then breaks down into cellular and cellular dendritic morphology. 28

49 Figure 9: Normal fusion boundary vs. dissimilar fusion boundary with Type II boundaries [26] Kou et al. [24] proposed two new mechanisms for macrosegregation at the fusion boundary of DMW's: one where the weld metal liquidus (T LW ) is greater than the base metal liquidus (T LB ) and one where the T LB is greater than T LW. In both mechanisms, it was found that the fusion boundary contains beaches, peninsulas, and islands that are approximately the same composition as the base metal. These features increased in thickness as the difference in T LW and T LB increased. For T LW > T LB, it was found that the fusion boundary contained thick, continuous beaches with random-oriented peninsulas and islands. This is due to the layer of liquid base metal being below T LW and, thus, cooler than the liquid weld metal. The liquid weld metal is pushed by convection into the layer of liquid base metal. This 29

50 liquid weld metal then freezes quickly without much mixing with the liquid base metal. The liquid base metal then solidifies as thick, continuous beaches [24]. For T LW < T LB, it was found that the fusion boundary contains thin, discontinuous beaches with peninsulas and islands roughly parallel to the fusion boundary. This is due to the region of liquid weld metal ahead of the T LW solidification front is below T LB. Thus, the liquid base metal is swept into that region and freezes quickly as peninsulas and islands. The thin and discontinuous liquid base metal left behind then solidifies as a thin and discontinuous beach [24] Solidification As metal alloys solidify, redistribution of solute occurs due to segregation [27]. Each alloying element has a different partitioning coefficient, which leads to the concentration of that alloying element in the liquid to either be enriched or depleted. The varying partitioning coefficients of the alloying elements lead to compositional variations in the fusion zone, predominantly at SGBs and SSGBs [23]. Segregation occurs both microscopically and macroscopically in weld metal solidification [23]. A compositional difference (due to segregation) at SGBs is considered macroscopic solidification, which is where solid diffusion is negligible and liquid diffusion is limited. Segregation that occurs between cells and dendrites is considered microscopic solidification, which is where solid diffusion is also considered negligible; however, complete mixing of the liquid takes place. Solute distribution from cell cores to boundaries can be described using the Scheil equation [27]. 30

51 Solidification conditions dictate the morphology and size of grains in the fusion zone. After the planar solidification front breaks down, multiple solidification modes exist which are caused by constitutional supercooling [27]. The ratio of the temperature gradient in the solidifying weld metal at the solid/liquid interface and the local solidification rate determines the solidification mode. This is depicted in Figure 10 [23]. Along the weld pool interface, these parameters vary which leads to multiple solidification modes in the same weld. Generally, the solidification mode varies from the planer front growth at the fusion boundary to equiaxed dendritic in the center. Figure 10: Effect of temperature gradient in the liquid and solidification rate on solidification mode [23] 31

52 2.5 Weldability Weldability is the ability to produce welds that are free of defects. The defects are the result of cracking mechanisms. In this study, cracking mechanisms, which include solidification cracking and hydrogen-induced cracking, were studied to evaluate the weldability of each weld combination Solidification Cracking Weld metal solidification cracking is a common weld defect in austenitic stainless steel due to the austenitic solidification mode. Solidification cracking is a hot cracking mechanism (i.e. it occurs above the effective solidus temperature) and occurs along solidification grain boundaries [28]. There is no concrete explanation for solidification cracking; however, multiple theories have been developed to explain the mechanism. Each theory explains the presence of liquid films along solidification grain boundaries that rupture with the application of tensile stresses [13]. The different theories include the Shrinkage-Brittleness Theory [29], Strain Theory [30], Generalized Theory of Super-Solidus Cracking [31], Technological Strength Theory [32], and the Modified Generalized Theory [33]. The Shrinkage-Brittleness Theory was proposed and developed by Bochvar [29]. This theory explains that there is an effective temperature interval (below a coherent temperature) that allows for solidification cracking to occur. Above the coherent temperature, molten material is free to flow. When the material reaches the coherent temperature upon cooling, an adequate solid network has formed to resist molten material to flow. 32

53 The Strain Theory was proposed and developed by Pellini [30]. This theory divides solidification into a mushy stage and a film stage. The mushy stage is early in solidification sequence and, similar to the Shrinkage-Brittleness Theory, solidification cracking is not an issue early in the solidification sequence. As solidification proceeds, the mushy stage transitions into a film stage, where liquid films are present along solid grains at (or around) the solidus temperatures. If these liquid films exist below the equilibrium solidus temperature, the susceptibility to solidification cracking increases. Tensile strain will accumulate along the wet boundaries due to solidification shrinkage. As the solidification temperature range increases, the time in film stage increases and the tensile strain builds up in the system. Thus, a larger temperature range increases the susceptibility to solidification cracking. The Generalized Theory of Solidification Cracking was proposed and developed by Borland [31]. This theory is a modified version of the Shrinkage-Brittleness Theory as it considers four stages of solidification. Stage 1 involves primary dendrite formation with continuous liquid present. Both the liquid and solid phases are capable of relative movement. In stage 2, the dendrites interlock each other, which prevent solid movement. However, the liquid is still able to move and healing of crack is possible. Stage 3 involves an intricate solid network that inhibits liquid movement. This stage is critical in solidification cracking because the liquid is no longer able to heal the cracks. Complete solidification occurs in stage 4. Solidification cracking cannot occur in stage 4 because there is no remaining liquid in the system. 33

54 The Modified Generalized Theory was proposed and developed by Matsuda et al [33]. This theory is nearly identical to the Generalized Theory proposed by Borland. There are only two differences in the theories. The first difference is that the temperature range in stage 1 (the mushy stage) is smaller. The second difference is that stage 3 is divided into a film and a droplet stage. In the film stage, liquid films allow for crack initiation and propagation. In the droplet stage, only crack propagation can occur. The Technological Strength Theory was proposed and developed by Prokhorov [32]. This theory examines the ductility and strain accumulation of the weld metal during solidification. If the strain exceeds the ductility of the material prior to solidification, then solidification cracking will occur. This can occur over any temperature in the solidification temperature range. All of these theories require liquid along the solidification grain boundaries during the final stages of solidification. The liquid is unable to support the strain induced from shrinkage and mechanical strain that leads to separation of these boundaries. This leads to solidification cracking. To identify solidification cracks, the fracture surface is often analyzed. The fracture surface is usually has a dendritic morphology with evidence of liquid films along the boundary [34]. If a sufficient amount of liquid is present upon final solidification, often time cracks are "healed" and back-filling of the cracks occurs [13]. Often times this effect is apparent via optical microscopy. There are various methods to prevent solidification cracking. The primary method is to control the solidification mode when possible. Austenitic solidification (A 34

55 mode) in steels is undesirable because austenite boundaries are easily wet by the liquid. Ferritic solidification (F mode) is will reduce the susceptibility to solidification cracking. If the solidification mode cannot be controlled (such as in austenitic stainless steels), minimizing impurity elements (S, P, and B) is helpful [17]. Other things that can be done to prevent solidification cracking are to reduce the restraint on the weld and to control the volume fraction of eutectic Hydrogen-Induced Cracking Hydrogen-Induced Cracking (HIC) is a type of cold cracking that occurs after welding has occurred. There are four components necessary for HIC to occur: hydrogen in the weld metal, high stresses, susceptible microstructure, and low temperature [13]. If one (or more) of these components can be avoided, then HIC will not occur. The sources of hydrogen include moisture in the flux, grease, paint & coatings, dirt, and rust. Hydrogen makes its way into the material through the weld metal as illustrated in Figure 11 [35]. T F indicates the austenite to ferrite/pearlite temperature and T B austenite to martensite transformation temperature. When the weld metal transforms from austenite to ferrite, the hydrogen goes to the ferrite due to the higher solubility of hydrogen in ferrite. Usually, the base metal has higher carbon than the weld metal, which means that the austenite to ferrite/pearlite transformation (in the weld metal) will occur before the austenite to martensite transformation (in the base metal). The build-up of hydrogen in the weld metal ferrite causes the hydrogen to diffuse across the fusion boundary into the HAZ austenite [13]. Since ferritic 35

56 materials have higher diffusion coefficients for hydrogen than austenitic materials [12], then hydrogen diffusion is promoted from the ferrite in the weld metal to the austenite in the base metal. The low diffusion coefficient in the austenite traps the hydrogen as the austenite transforms to martensite [13]. Figure 11: Diffusion of hydrogen from weld metal to HAZ [35] The high stresses necessary for HIC to occur are developed during cooling by solidification shrinkage and thermal contraction [13]. Martensite is a susceptible microstructure to HIC. The harder the martensite (higher carbon), the greater the susceptibility to HIC. Martensite formation generally occurs at lower temperature, which is why HIC occurs at low temperatures [13]. The mechanism for HIC to occur is not clearly understood; however, multiple theories have been developed. Troiano [36] proposed the decohesion theory, which states that the high stress present at a stress concentration, such as a crack tip, attracts 36

57 dissolved atomic hydrogen and reduces the cohesive strength. If the hydrogen exceeds a critical value, then a small crack forms and propagates. Petch [37] developed the surface adsorption theory. The basis of this theory is that hydrogen promotes crack growth by reducing the surface energy of the crack. The surface energy is lowered when hydrogen is absorbed onto the crack surface. According to Griffith's criterion [38], the fracture stress is lowered by the square root of the surface energy. Cracks can now propagate under low stresses. The planar pressure theory was developed by Zapffe and Sims [39], which is related to the development of hydrogen bubbles in the microstructure. The atomic hydrogen diffuses to defects in the microstructure, such as grain boundaries and interfaces. Once there is sufficient hydrogen built up at these sites, the atoms combine to form molecular hydrogen (H 2 ) bubbles with high internal pressure. If the internal pressure exceeds the yield strength of the material, then these internal voids grow, link together, and eventually lead to the formation of cracks. Beachem's stress intensity model [40] relates the stress intensity factor and hydrogen concentration at the crack tip with fracture behavior. This model suggests that the microscopic deformation ahead of the crack tip is related to the concentrated hydrogen dissolved in the lattice. The microscopic plasticity decreases by decreasing the stress intensity factor at the crack tip. Therefore, it was determined that the stress intensity factor and hydrogen concentration affects the fracture behavior. Figure 12 illustrates how varying both factors changes the fracture surface. At low stress intensity factors and hydrogen concentrations, HIC is not expected to occur. As both 37

58 factors increase, the fracture surface transitions from intergranular (IG) to quasicleavage (QC) to microvoid coalescence (MVC). Figure 12: Combined effect of stress intensity factor and hydrogen content on fracture surface [40] The two primary methods to prevent HIC is to choose the proper welding consumable and to control the welding parameters. A lower hydrogen welding process, such as gas tungsten arc welding or gas metal arc welding reduces the amount of hydrogen in the weld metal. Along with the low hydrogen process, consumables should be kept in sealed containers and (if wet) should be baked to dry them out. Generally, lower strength filler metals or austenitic stainless steel filler metals are used to avoid HIC. The lower strength filler metals help reduce stress 38

59 levels in the HAZ and the austenitic filler metals trap hydrogen in the weld metal as discussed in section [13]. The welding parameters that can be controlled include preheat (and interpass) temperature and postweld heat treatment (PWHT). A PWHT allows for stress relaxation, hydrogen to diffuse out of the workpiece, and austenite to transform into a less susceptible microstructure than martensite [13]. Figure 13: Effect of preheating on HIC [41] The use of a proper preheat and interpass temperature can diminish HIC as shown in Figure 13 [41]. To determine the proper preheat temperature, empirically derived tables on the material, which list the preheat temperature, should be utilized. Another option is to utilize the carbon equivalence method [13]. 39

60 2.5.3 Carbon Equivalence The carbon equivalent concept was first developed to give an indication of the hardenability of a steel based on the alloying elements. Over time, the concept was extended to determine the susceptibility to HIC based on the composition. Currently, the purpose of the carbon equivalence is to determine welding procedures that avoid HIC [42]. Many methods have been developed and they can vary considerably. All of the methods take into consideration steel composition, welding heat input, joint geometry, material thickness, hydrogen level, and preheat temperature. The four, most prominent, methods are the CE, CET, CE N, and AWS methods. These methods cannot be used for all ferrous materials, as they each have their own compositional limitations shown in Table 5 [42]. The method to be utilized for a specific application is determined by the industry. Table 5: Parent metal composition ranges [42] 40

61 CE method The British CE method is based on the concept of a critical hardness necessary to avoid HIC in the HAZ [42]. This method is broken down into seven steps: 1. The carbon equivalent value (%) is calculated using equation 3. 3) CE IIW = C + Mn + Cr+Mo+V + Ni+Cu The hydrogen scale of the welding process and consumable needs to be determined. The different scales are broken down in Table 6. These scales are based on the weld diffusible hydrogen content. Table 6: Hydrogen scales to be used with any arc welding process [42] 41

62 3. The joint type (fillet or butt weld) needs to be determined. 4. The appropriate graph (from the standard) must be selected for the hydrogen scale and carbon equivalent. If a graph is not available for a specific hydrogen scale and carbon equivalent combination, the next highest carbon equivalent value should be utilized. 5. The heat input should be calculated for the welding job. 6. The combined thickness of the fillet or butt weld should be calculated. 7. The heat input and combined thickness should be plotted on the graph that was determined in the third step. The preheat temperature is determined by reading the preheat line directly above or to the left of the point. If no preheat temperature is desired, then (using the same graph) determine the minimum heat input for no preheat by looking at the 20 C line [43]. An example of one of the graphs is shown in Figure

63 Figure 14: Carbon Equivalence vs. Combined Thickness to determine the Preheat Temperature [42] CET method The German CET method is based on results from the Tekken Test. This method is only valid for parameters found in Table 7 [42]. The carbon equivalent is calculated using equation 4. 4) CET = C + Mn+Mo 10 + Cr+Cu 20 + Ni 40 43

64 Table 7: Range of validity for the CET method [42] The CET method combines the carbon equivalent, plate thickness (d), hydrogen content (HD), and heat input (Q) into an equation (equation 5) that calculates the appropriate preheat temperature ( C). 5) T p = 697(CET) tanh ( d ) (HD0.35 ) + (53(CET) 32)Q 328 CE N method Like the CET method, the Japanese CE N method is based on the results of the Tekken Test. This method is broken down into nine steps. 1. The two different carbon equivalents need to be calculated using equations 6 and 7. 6) CE N = C + f(c) [ Si 24 + Mn 6 + Cu 15 + Ni + Cr+Mo+V 20 5 f(c) = tanh(20(c 0.12)) 7) CE IIW = C + Mn + Cr+Mo+V + Ni+Cu ] 44

65 2. The hydrogen content of the weld metal should be estimated and ΔCE N (hydrogen) should be found using Figure 15 [42]. Figure 15: CE N correction with respect to the weld metal hydrogen content [42] 3. The heat input of the welding job should be calculated and ΔCE N (heat input) should be found using Figure 16 [42]. 45

66 Figure 16: CE N correction with respect to the weld heat input and CE IIW [42] 4. The relevant CE N should be calculated, which is the sum of CE N, ΔCE N (hydrogen), and ΔCE N (heat input). 5. The plate thickness should be determined, which is that of the thicker plate (or pipe) in a butt or fillet weld geometry. 6. The necessary preheat temperature to avoid cracking during with the Tekken Test should be found using Figure 17 [42]. 46

67 Figure 17: Master curve for minimum preheat for the Tekken Test [42] 7. The yield strength of the weld metal should be determined. 8. The eighth step is to find the preheat temperature correction using Figure 18 [42]. The final step is to calculate the necessary preheat temperature by adding the preheat temperature (found in the sixth step) and the preheat temperature correction (found in the eighth step). This new preheat value should be rounded upwards to the nearest 5 C. 47

68 Figure 18: Correction of necessary preheat of welding practice [42] 9. The necessary preheat temperature should be calculated by adding the preheat temperature (found in the sixth step) and the preheat temperature correction (found in the eighth step). This new preheat value should be rounded upwards to the nearest 5 C. AWS method The American AWS method utilizes two separate techniques for estimating the weld conditions to avoid cracking: HAZ hardness control and hydrogen control. The HAZ hardness control technique is based on the assumption that cracking will not occur in the HAZ if the hardness is kept below a critical value [42]. To manage the hardness, the cooling rate is kept below a critical value, which is dependent on the 48

69 hardenability of the steel. Factors that assist in determining the critical hardness include steel type, hydrogen level, restraint, and service conditions. The hydrogen controlled technique is based on the assumption that cracking will not occur if the amount of hydrogen remaining in the joint after cooling to roughly 50 C does not surpass a critical value, which is dependent on the steel composition and restraint level [42]. To determine which technique to use, the carbon equivalence (equation 8) is calculated and the CE value is plotted vs. carbon content (Figure 19) to determine which zone the steel is classified. If the steel lies within zone 1, then cracking is unlikely to occur; however, cracking is possible if there sufficient restraint or high hydrogen. Thus, the hydrogen controlled technique should be used. If the steel lies within zone 2, then the hardness control technique should be used to determine the minimum heat input without preheat. If the heat input is impractical, then the hydrogen controlled technique should be used to determine the preheat. If the steel lies within zone 3, then the hydrogen controlled technique should be used. 8) CE = C + Mn+Si 6 + Cr+Mo+V 5 + Ni+Cu 15 49

70 Figure 19: Zone classification of steels [42] If the hardness controlled technique is required, then four more steps are needed to avoid cracking: 1. The heat input of the welding process should be determined. 2. Based on the material thickness, heat input, and leg length, the correct figure (from the standard) should be utilized. 3. Based on the cooling rate, estimate the hardness. 4. The results should be analyzed. If the expected hardness is less than 350 HV, then the high-hydrogen practice may be used. If the expected hardness is between HV, then the low-hydrogen practice should be used. According to ISO 3690, low-hydrogen practice (diffusible hydrogen content < 10 ml/100g) applies for SAW, SMAW, GMAW, and FCAW [44]. 50

71 If the hydrogen controlled technique is required, then four more steps are needed to avoid cracking: 1. The first step is to determine the composition parameter value using equation 9. 9) P cm = C + Si 30 + Mn 20 + Cu 20 + Ni 60 + Cr 20 + Mo 15 + V B 2. The estimated hydrogen content should be determined. There are three hydrogen levels in this technique: H1 extra-low hydrogen (diffusible hydrogen < 5mL/100g), H2 low hydrogen (diffusible hydrogen < 10mL/100g), and H3 hydrogen not controlled [42]. 3. The index grouping based off of the composition parameter and hydrogen content should be found using Table 8 [42]. 51

72 Table 8: Susceptibility index grouping as a function of HD and P cm [42] 4. The minimum preheat and interpass temperatures for the restraint level is estimated based on Table 9 [42]. 52

73 Table 9: Minimum preheat and interpass temperature [42] Critical Hardness Hardness is a major component in HIC susceptibility. Above a critical hardness value, HIC should (in theory) be an issue and occur. Many factors play a role in the critical hardness: material composition, plate thickness, joint type, heat input, preheat temperature, etc. Suzuki developed a formula which calculates the maximum hardness in the HAZ of steels [45]. However, Duren has modified the formula to improve accuracy of the hardness prediction [46] Weldability Testing A great deal of research has been done in the development of tests to evaluate the weldability of alloys. 53

74 Tekken Test The Tekken test, also known as the Y-groove Test, was originally developed by the Technical Research Institute of the Japanese National Railways. The Tekken test is a self-restrained test that is used to determine the susceptibility to HIC in weld metal and the HAZ of arc welds. This test is also used to optimize the preheating temperature of the weldment. The basic setup for the Tekken Test is shown in Figure 20 [10]. The setup consists of two, 200 mm base plates (>10 mm in thickness) that are coupled with anchor welds at the ends. The anchor welds must be made from a consumable with a yield strength equal to or greater than the base plate material [10]. The central 80 mm of the base plates forms a V-groove containing a 2 mm gap, which forms a "Y" shape. If anchor welds are not desired or are difficult to produce, it is possible to use one (200 x 150 mm) base plate and utilize a water jet to cut out the middle Y-groove in the plate. 54

75 Figure 20: Tekken test setup [10] A single-pass weld is made in the top V-portion of the Y-groove. The gap in the bottom of the plate acts as a stress concentrator and allows for cracking initiation in the weld or HAZ. After welding, a minimum of 48 hours is required prior to inspection. Two methods of inspection are to be performed: visual examination and metallographic examination. Visual examination is performed only if a crack broke through to the surface. For visual examination, surface cracks are visually examined and a crack ratio is calculated (equation 10). The total length of the surface cracks (Σ l f in mm) divided by the length of the test bead (L in mm) multiplied by 100 results in the crack ratio (C f in %). 10) C f = l f L 100 For metallographic examination, sections should be cut according to EN The surfaces should be prepared so that the weld metal and HAZ can be examined at 55

76 a minimum magnification of 50x. Crack lengths (H c in mm) are measured and divided by the test bead height (H in mm) and multiplied by 100 in order to calculate a crack section ratio (equation 11). A visual representation of measuring the crack section is shown in Figure ) C s = H c H 100 Figure 21: Measuring crack length and bead height [10] Controlled Thermal Severity Test The Controlled Thermal Severity Test, also known as the CTS Test, is a selfrestraint test to determine the susceptibility to HIC in the HAZ of arc welds. This test can be used with shielded metal arc welding, gas metal arc welding, and gas tungsten 56

77 arc welding. High current arc welding processes, such as submerged arc welding, cannot be used in this test. The setup for the CTS Test is shown in Figure 22 and the dimensions of the test pieces are in Table 10. The setup is comprised of a small block that is bolted to a plate, with the thickness of both being equal. Anchor welds are made between the small and large block on the side of the block perpendicular to the rolling direction. The anchor welds must be made from a consumable with yield strength equal to or greater than the base plate material [10]. After making the anchor welds, allow the samples to sit for 12 hours before making the test welds. Figure 22: CTS Test setup [10] 57

78 Table 10: CTS test piece dimensions and tolerances [10] Test welds should be made using the jig in Figure 23. The jig is required so that the test welds are symmetrically in the flat positions across the full width of the block in a single direction and single pass [10]. The test welds must not exceed the ends of the block. Within 60 seconds of completion of the first test weld, the test assembly should be placed in a water bath as shown in Figure 24. The test assembly should be kept in the water bath for a minimum of 48 hours prior to making the second test weld. 58

79 Figure 23: Jig used to position test assembly [10] Figure 24: Water bath arrangement [10] 59

80 The second test weld should be made and cooled in the same fashion as the first test weld. Allow the test assembly to sit for 48 hours after completion of the second test weld before examination. To examine the assembly, the assembly should be cut into equal size samples shown in Figure 25. The cut faces should be prepared for micro-examination at a minimum magnification of 50x. The sample will be considered "cracked" if there are HAZ cracks longer than 5% of the leg length. If there is a weld metal root crack that exceeds 5% of the throat thickness, then the test piece will be considered invalid. Figure 25: Sectioning of CTS test assembly [10] 60

81 2.6 Welding Armored Steels with Stainless Steel Consumables This section contains pertinent information and case studies on welding armored steels with stainless steels consumables. Included in this section is the weldability, properties, and reliability of these welds. It is important to understand previous work with these welds in order to avoid failure in future welds Consumable Selection HIC is a potential issue when welding armored steels. HIC results from the collision of three factors: stress, susceptible microstructure, and presence of hydrogen. Stress is nearly unavoidable due to restraint, distortion, and residual stresses. Martensite is a susceptible microstructure; however, armored steels are often designed to be martensitic. The presence of hydrogen is the factor that can be controlled the most, which can be accomplished by choosing a proper welding consumable. Kuzmikova et al. [47] studied the amount of hydrogen present in armored steel welds using a ferritic vs. austenitic filler metal. In order to measure diffusible hydrogen levels, hydrogen was collected over mercury (known as the Reference Test Method). In order to measure the residual hydrogen levels, the Inert Gas Melt Extraction Method was utilized. From the Reference Test Method, it was found that welds with the ferritic consumable resulted in a diffusible hydrogen level approximately 18x greater than that produced by the austenitic consumable. It is believed that this is due to hydrogen pick-up during prolonged storage of the consumable [47]. The residual hydrogen content in the welds made with the austenitic consumable was 5.8 ml/100g whereas 61

82 the ferritic welds had 0.9 ml/100g. This demonstrates that the hydrogen is "locked" in the austenitic weld metal [47]. Further testing on the effect of preheat and type of consumable took place for the austenitic consumable. It was found that, as the preheat temperature decreased from 80 C to 7 C, the diffusible hydrogen level increased and the residual hydrogen level decreased. This implies that a higher preheat increases the cooling time and, therefore, decreases the diffusible hydrogen [47]. Three types of consumables (fluxcored, metal-cored, and solid wire) were tested to analyze the difference in hydrogen content. The results show that the flux-cored wire had the highest diffusible and residual hydrogen and the solid wire had the lowest. This is due to the flux being a source of hydrogen [47] Joint Design Weld geometry and weld groove angle varies with material type, material thickness, and welding method. As the material thickness increases, the weld groove angle increases and requires additional material. This leads to increased cost, increased number weld defects, and higher residual stresses. İpek et al. studied the effects of welding groove angle and geometry on the mechanical properties of armored steel [48]. The study utilized gas metal arc welding (GMAW) of armor steel (0.272 C, 0.92 Cr, Ni, Mo, Mn in wt%) with an austenitic stainless steel (0.066 C, Si, Mn, P, S, Cr, Ni in wt%). X and V type geometries were tested along with groove angles of 48, 54,

83 Tension and compression tests were carried out to see the effects of each joint under different load types. In the tension tests, it was found that the V-groove geometries performed the best [48]. For equivalent groove angles, the V-groove geometries experienced higher tensile strengths than the X-groove geometries. As the weld groove angle decreased, the tensile strength increased in the V-groove geometries, which was expected. In the X-groove geometries, the tensile strength decreased as the weld groove angle decreased. This implies that, for welds in tension, the V-groove geometry with a small groove angle should be utilized [48]. In the compression tests, it was found that the X-groove geometries performed the best [48]. For equivalent groove angles, the X-groove geometries experienced higher compressive strengths than the V-groove geometries. In both geometries, the 54 groove angle experienced the highest compressive strength. This implies that, for welds in compression, the X-groove geometry with a mid-size groove angle should be utilized [48]. In both the tension and compression tests, it was found that the yield strength of both weld geometries increased with decreasing groove angle Mechanical Properties It is not difficult to increase the strength of armored steels. However, as the strength (and hardness) increases, the steel becomes more brittle and is more prone to failure. Welding of these high strength and high hardness steels can increase the susceptibility to failure. Therefore, a strong balance must be found between strength and toughness of welded armored steels. 63

84 Mitelea et al. [49] investigated butt joint properties of welds made with two armored steels and two austenitic stainless steels using pulsed MIG welding. Static mechanical tests were performed in order to determine the location of fracture and mechanical resistance. In all tests, the fracture occurred in the weld metal. Despite the location of fracture, the mechanical resistance was 50-90% higher in the weld metal than in pure base metal. The Charpy V-notch test was performed to determine the brittle fracture tendency of the joints. The V notch was placed in the weld, in the HAZ, or in the base metal. Testing occurred at multiple temperatures ranging from -50 C to +50 C. It was found that the fracture toughness of both the weld metal and HAZ was higher than the base metal at all temperatures. The weld metal had the highest fracture toughness readings. Characterization of the weld found that the weld metal consisted of austenite, martensite (from base metal dilution), and δ ferrite [49]. Reddy et al. [50] tested the mechanical properties of joining an HSLA steel with 309L stainless steel and 18Cr-8Ni-6Mn stainless steel. Weld joints tested were a double pass, 60 V-groove butt joint. Tensile testing was done to get strength and ductility measurements while Charpy V-notch testing was carried out (at room temperature) to get toughness measurements. Tensile testing revealed that both filler metals have low tensile strengths compared to the HSLA steel, which was expected. Therefore, fracture occurred in the weld metal. Fracture surfaces revealed that both filler metals failed by ductile fracture, with 309L exhibiting more ductility. This is due to the increased Ni content of 309L 64

85 and the tendency of 18Cr-8Ni-6Mn to form martensite with increasing dilution [50]. The yield strength of the 18Cr-8Ni-6Mn stainless steel was greater than the 309L. Charpy V-notch testing showed that the toughness of the HAZ and base metal are roughly the same. It also showed that the toughness of 309L exceeds that of the HSLA steel by nearly 1.5x, whereas the 18Cr-8Ni-6Mn had toughness below the HSLA steel. This is due to the cellular solidification structure of 18Cr-8Ni-6Mn, which suggests that there is weld metal segregation that could lead to lower toughness [50]. Weld metal segregation is greater in the first weld pass due to an increased dilution level, which is why the first pass of each weldment exhibited a lower toughness than the second pass Weldability Any DMW is going to have weldability issues. Welding of armored steels with stainless steel consumables is no different. Any potential issue that one of the alloys can sustain could occur in these welds. It is important to investigate potential issues and determine ways to mitigate (or eliminate) the issues. Reddy et al. [50] tested the weldability of joining an HSLA steel with 309L stainless steel and 18Cr-8Ni-6Mn stainless steel. The T-joint test was carried out to analyze the susceptibility to solidification cracking. The Tekken test was carried out to analyze the HIC susceptibility of the steel. For both tests, no cracking was found at 100x magnification, which implies that neither weld combination are prone to solidification cracking or HIC. 65

86 Alkemade [51] tested the HIC susceptibility of welding Bisalloy 500 (armored steel) with an austenitic stainless steel (310), duplex stainless steel (312), and ferritic steel (CIGWELD Autocraft S6). The Tekken test was utilized to test the susceptibility to HIC and the welding parameters were varied in order to see any possible changes. Two types of cracks were found. The first type of crack was an under bead crack, which initiated at the weld root and traveled through the HAZ. These cracks were only found in the ferritic welds with a heat input of 0.5 kj/mm and preheat at or below 75 C. This type of cracking did not occur when the heat input was raised to 1.2 kj/mm regardless of preheat. The second type of crack was a solidification crack in the middle of the weld metal. These cracks were only found in the duplex stainless steel welds with a heat input of 0.5 kj/mm and preheat at or below 75 C. This type of cracking did not occur when the heat input was raised to 1.4 kj/mm (and no preheat) or 1.2 kj/mm (150 C preheat). No cracking was found when using austenitic stainless steel 310 regardless of the heat input or preheat. Based on these results, it is recommended to have a heat input of kj/mm and preheat of C regardless of the filler metal. 66

87 CHAPTER 3: OBJECTIVES The primary objective of this work is to determine which stainless steel consumable(s) is the most effective in solving HIC problems and avoiding solidification cracking problems when welding both armored steels Armox 440 and RHA. Currently, low alloy steel consumables are used to weld these armored steels; however, they are not effective in avoiding HIC. It is desirable to test and compare three stainless steels with two low alloy steel consumables. From this testing, the goal is to determine which stainless steel is the most efficient when welding both armored steels. To achieve these goals, multiple steps need to be taken. 3.1 Thermodynamic Prediction The Schaeffler Constitution Diagram has been utilized for years in order to predict the microstructure of stainless steel welds. Despite some inaccuracies, the Schaeffler Constitution Diagram is a widely accepted tool for prediction of weld microstructure. New advancements in thermodynamic simulation have led to more accurate predictions of phase transformations and microstructures of metallic alloys, along with various metallurgical reactions that cannot be predicted by the Schaeffler Constitution Diagram. For this study, ThermoCalc will be used to predict the nonequilibrium solidification range, solidification mode, and phases present upon solidification. The results from both the Schaeffler Constitution Diagram and 67

88 thermodynamic simulations will be compared to and validated with experimental results. 1. Plot each filler metal-base metal combination on the Schaeffler Constitution Diagram to see the predicted microstructure throughout the whole base metal filler metal dilution range. 2. Thermodynamic simulations will be performed to determine the solidification mode, solidification temperature range, and microstructure throughout the whole base metal filler metal dilution range in each filler metal-base metal combination. 3.2 Metallurgical Characterization The objective of this study is to evaluate the microstructure of welds made with stainless steel consumables and armored steels. 1. Hardness mapping will be performed along the fusion boundary of each filler metal-base metal combination to determine the local hardness of microstructural constituents. 2. Optical microscopy will be utilized to examine the microstructure of the weld metal, fusion boundary, and heat affected zone (HAZ). 3. Energy-dispersive spectroscopy (EDS) will be utilized to determine composition gradients and the composition of each constituent. 3.3 Weldability Testing The objective of this study is to evaluate the weldability of using stainless steel consumables to weld armored steels. Low alloy steel welding consumables will also 68

89 be tested to provide a basis for comparison. Susceptibility to hydrogen-induced cracking will be examined in each filler metal-base metal combination. 1. Tekken and CTS Tests will be performed on each filler metal-base metal combination to evaluate and rank the susceptibility to hydrogen-induced cracking. 2. The test results of all filler metal-base metal combinations will be compared and analyzed to recommend welding consumable(s) for each base metal. 69

90 CHAPTER 4: EXPERIMENTAL PROCEDURES 4.1 Introduction This chapter includes the materials and experimental techniques employed during this project. The experimental techniques contain a description of the equipment, experimental procedure, and goal of the experimental technique. Testing was designed toward understanding the metallurgy, microstructure, and weldability of joining armored steels with stainless steel consumables. 4.2 Materials This project examined joining armored steels with stainless steel consumables. Table 11 contains the composition of the two armored steels that were tested. Table 12 contains the composition of the three stainless steel consumables that were tested. All compositions were provided by AEM. Included in Table 12 are the compositions of the low alloy consumables that were also tested as reference welds. 70

91 Table 11: Chemical Compositions of the Armored Steels (wt%) Material Armox 440 RHA Mil Standard Class 1 Fe C P Mn Si S Ni Cr Mo Cu Al Sn Ti B Nb As Pb Zr

92 Table 12: Chemical compositions of the welding consumables (wt%) Material Element ER309LHF Sandvik AXT (~ER307) ER312 ER70S-6 ER100 Fe C P Mn Si S Ni Cr Mo Cu Al Sn Ti B Nb As Pb Zr Thermodynamic Simulation This project utilized ThermoCalc thermodynamic software in order to determine the solidification range, solidification mode, and microstructure upon solidification. The TCFE5 database and Scheil module was utilized for all weld combinations. The assumptions made in the simulations were that carbon and boron were fast diffusing elements. Simulations were stopped at 98% solidification. 72

93 4.4 Metallographic Characterization This section contains the preparation and examination procedures for metallographic samples Sample Preparation Due to the increased difficulty in sample preparation, conventional metallography sample preparation was modified in this project. Sections, previously cut by AEM, were mounted in Buehler Konductomet conductive resin with a Leco PR-32 mounting press. Mounted samples were polished using 240, 400, 600, 800, and 1200 grit SiC metallographic paper. During each step, microid diamond compound extender, not water, was used to lubricate the paper. Following the 1200 grit paper, polishing continued with 6, 3, and 1 μm paste suspension while applying the microid diamond compound extended to each diamond pad. Polishing concluded with a 0.5 colloidal silica suspension for 2-3 minutes. Between each step in the polishing process, samples were sprayed with ethyl alcohol, cleaned in an ultrasonic bath for 3-5 minutes, sprayed with acetone, and dried with the heat gun for 2 minutes. Etching of samples was divided into two steps: 1. etching of the base metal and 2. etching of the filler metal. The etching of the armored steel (base metal) was performed first using a 5% Nital (nitric acid + ethyl alcohol) solution. The sample was flipped upside down and immersed roughly 2-3 mm into the solution for seconds. Following this step, the sample was sprayed with ethyl alcohol, cleaned in an ultrasonic bath for 3-5 minutes, sprayed with acetone, and dried with the heat gun for 2 minutes. Etching of the filler metal was performed next using a 10 vol.% CrO 3 73

94 solution. Samples were placed in a glass dish and fully submerged in the CrO 3 solution. The samples were electrolytically etched using a constant voltage DC power supply set at 5 volts. A tungsten electrode (positive) touched the sample and a piece of stainless steel foil (negative) was held over the sample. Both were immersed in the CrO 3 solution. The samples were etched for seconds. Following this step, the samples were washed in two, separate ethyl alcohol baths, sprayed with acetone, and dried using the heat gun for 2 minutes Optical Microscopy Samples were evaluated using an Olympus GX51 with image acquisition capabilities. This microscope was used to examine samples that were prepared metallographically Hardness Mapping Hardness mapping was utilized to determine areas of varying microstructure and dilution. Hardness mapping was performed on a Leco AMH43 Automatic Micro/Macro-indentation Hardness Testing System. Prior to hardness mapping, samples were mounted and polished from 240 grit paper down to the 0.5 colloidal silica solution using the same procedure in section The sample was placed in a mount, made flat by a level bar across the mount, and locked in place with a set screw. Finally, the sample was placed under the indenter and microscope. Using the 2.5x turret, the surface was mapped and an indentation path was created. The indentation path is the area of the sample that will undergo indentation and the hardness will be read. The indentation path was made along the fusion boundary, 74

95 going into both the weld metal and HAZ, for each sample. Indents were inserted in the indentation path 100 x 100 μm apart from each other. The load of each indent was set to 100 grams. Prior to the start of indenting, the sample was focused using the auto-focus at 2.5x, 10x, and 50x magnification. Once focused, indenting started followed by indent measurements. Upon completion of indent measurements, results were saved and exported to an excel file, which contained the hardness and location (x and y coordinates) of each indent. With the results from each sample, a scatter plot of the indents was created. The plot painted a picture of the location of various hardness values throughout the sample. Hardness ranges varied by color to show hardness differences in the sample Energy Dispersive Spectroscopy Energy Dispersive Spectroscopy (EDS) was utilized to qualitatively evaluate the composition of various weld regions. From the compositions, the appropriate dilution was calculated using the content of the larger elements (Fe or Cr) for each point along the EDS traverse. Based on the calculated dilutions, the susceptibility to cracking could be estimated. The width of regions was determined by overlaying the EDS traverse over an etched image of the examined area and measuring the distance the traverse traveled through each region. During optical microscopy, areas of interest along the fusion boundary were determined for EDS. Specific characteristics of the fusion boundary for EDS analysis include weld swirls, planar growth regions, transition region, cellular and cellular dendritic regions in the weld metal, and the HAZ. In a region that required an EDS line scan, hardness indents were placed at the 75

96 ends of the desired line. Samples were re-polished and EDS was performed using the XL-30F ESEM field emission gun SEM. EDS lines were made between the previously made hardness indents. 4.5 Weldability Testing Two tests were utilized to determine the weldability, specifically the susceptibility to HIC, of armored steel joints made with stainless steel consumables. The Tekken test was performed to determine the susceptibility to HIC in both the weld metal and the HAZ. The CTS test was performed to determine the susceptibility to HIC in the HAZ Tekken Test The Tekken test was used to evaluate the susceptibility to HIC in each weldment. The Tekken test experiments were carried out according to ISO (2003) [10]. A single-run weld bead was deposited into a V butt joint that was designed to allow for cracking in the HAZ or WM. The anchor welds were placed on both ends of the weld to develop the necessary restraint. The stresses that developed in the weld were the result of shrinkage restraint. After allowing the test welds to sit for a minimum of 48 hours, the test welds were sectioned into four equal sections. Each section was polished and etched as described in section Examination of each section was done in accordance to ISO (2003) [10]. Further examination included metallographic characterization using an optical microscope and fractography using an SEM. 76

97 To examine the effect of preheat temperature on these weldments, two sets of Tekken tests were made. One set was made with each filler metal-base metal combination and no preheat. The second set was made with each filler metal-base metal combination and preheat of 450 F (250 C). This preheat temperature was selected because it is currently used in service CTS Test The CTS test was used to evaluate the susceptibility to HIC in each weldment. The CTS test experiments were carried out according to ISO (2003) [10]. The anchor welds were placed on both sides of the top block to develop the necessary restraint. The stresses that developed in the weld were the result of shrinkage restraint. The first test weld was made 12 hours after completion of the restraint welds. The first test weld was cooled in a water bath for 48 hours prior to making the second test weld. The second test weld was cooled the same as the first test weld prior to examination. The test welds were cut into four equal sections. Each section was polished and etched as described in section Examination of each section was done in accordance to ISO (2003) [10]. Further examination included metallographic characterization using an optical microscope and fractography using an SEM. The effect of preheat temperature may not be as apparent in the CTS test compared to the Tekken test. However, the effect of preheat temperature was examined in the CTS test as a comparison to the Tekken test results. Similar to the Tekken test 77

98 weldments, two sets of CTS tests were made, one without preheat and one with preheat of 450 F (250 C). 78

99 CHAPTER 5: RESULTS AND DISCUSSION 5.1 Thermodynamic Simulations Computational modeling allows for efficient analysis of the weld metal microstructure without hand calculations and minimizes testing, which saves time and money. The following section portrays the results of modeling each DMW combination using the ThermoCalc Scheil Module. Simulations were performed for each DMW from 0% dilution to 100% dilution in increments of 5%. From the results, the liquidus temperature, solidus temperature (determined at 98%solid fraction), and temperature(s) at which solidification mode changed were plotted onto a solidification diagram. A solidification diagram is a plot of temperature vs. dilution and depicts how dilution affects solidification mode and solidification temperature ranges. This gives a rough estimate of the solidification cracking susceptibility. Along with modeling, the microstructure in each DMW was predicted using the Schaeffler Diagram. The solidification mode and microstructure of each DMW combination, as well as a brief analysis, will support the conclusions that were derived based on the results of this study Armox Sandvik AXT The solidification diagram for Armox Sandvik AXT (~ER307) is shown in Figure 26. Below 50% dilution and above 75% dilution in this combination, the 79

100 Temperature ( C) solidification mode is ferrite-austenite (FA). Between 50-75% dilution, the solidification mode is austenitic (A). Between 45-50% dilution, the solidification mode is austenite-ferrite (AF). It is well established that austenitic solidification mode in austenitic stainless steel welds can result in solidification cracking. This simulation result implies a potential risk of weld metal solidification cracking in the A solidification mode dilution range (mentioned above) of Sandvik AXT filler wire with Armox 440 steel. The largest solidification temperature range (difference between liquidus and solidus temperatures) is 111 C and it occurs at pure Sandvik AXT (0% dilution) and at 55% dilution. The smallest solidification temperature range is 101 C, which occurs at 95% dilution and pure Armox 440 (100% dilution) L + F L + F + A Liquid L + A Solid Sandvik AXT Dilution (%) Armox 440 Figure 26: Armox Sandvik AXT Solidification Diagram 80

101 Figure 27 is a Schaeffler Diagram with a tie line connecting Armox 440 and Sandvik AXT. Each marker indicates a 10% increase in dilution starting from 0% dilution (pure Sandvik AXT) up to 100% dilution (pure Armox 440). Sandvik AXT is predicted to be almost fully austenitic, with less than 5% ferrite. Although there is mixing with the base metal, weld metal dilutions are expected to be less than 50%. Assuming dilutions in the 30-40% range, the microstructure is projected to be a combination of austenite and martensite. For the pure base metal (a.k.a. HAZ), the Schaeffler diagram predicts a martensitic structure. As the composition deviates from pure base metal to the filler metal, a martensitic structure is still expected down to (roughly) 50% dilution. This implies that a portion of the transition region from pure base metal to weld metal should be martensitic. 81

102 Ni Eq. = %Ni +30*%C + 0.5*%Mn F + M A+M Armox 440 Armox Sandvik AXT A M M+F Sandvik AXT No Ferrite A + F Cr Eq. = %Cr + %Mo + 1.5*%Si + 0.5*%Cb F Figure 27: Schaeffler Diagram connecting Armox 440 and Sandvik AXT Armox ER309LHF The solidification diagram for Armox ER309LHF is shown in Figure 28. Below 50% dilution and above 77% dilution in this combination, the solidification mode is ferrite-austenite (FA). Between 65-77% dilution, the solidification mode is austenitic (A). Between 50-65% dilution, the solidification mode is austenite-ferrite (AF). This simulation result implies a potential risk of weld metal solidification cracking in the A solidification mode dilution range (65-77%). The largest 82

103 Temperature ( C) solidification temperature range is 101 C, which occurs at pure Armox 440 (100% dilution). The smallest solidification temperature range is 56 C, which occurs at pure ER309LHF (0% dilution) Liquid L + A 1430 L + F 1410 L + F + A Solid ER309LHF Dilution (%) Armox 440 Figure 28: Armox ER309LHF Solidification Diagram Figure 29 is a Schaeffler Diagram with a tie line connecting Armox 440 and ER309LHF. The pure filler metal (ER309LHF) is predicted to be austenite with a small amount of ferrite (~10%). Assuming weld metal dilutions in the 30-40% range, the weld metal microstructure is projected to be entirely comprised of austenite. For the pure base metal (a.k.a. HAZ), the Schaeffler diagram predicts a martensitic 83

104 Ni Eq. = %Ni +30*%C + 0.5*%Mn structure. As the composition deviates from pure base metal to the filler metal, a martensitic structure is still expected down to (roughly) 60% dilution. This implies that a portion of the transition region from pure base metal to weld metal should be martensitic. Armox 440 F + M A+M Armox ER309LHF M M+F A No A + F Cr Eq. = %Cr + %Mo + 1.5*%Si + 0.5*%Cb F Figure 29: Schaeffler Diagram connecting Armox 440 and ER309LHF 84

105 Temperature ( C) Armox ER312 The solidification diagram for Armox ER312 is shown in Figure 30. At 75% dilution, the solidification mode is austenitic (A). All other dilutions experience a ferrite-austenite (FA) solidification mode. This simulation result implies a potential risk of weld metal solidification cracking only at 75% dilution. The largest solidification temperature range is 188 C, which occurs at pure ER312 (0% dilution). The smallest solidification temperature range is 84 C, which occurs at 80% dilution Liquid 1450 L + F L + A L + F + A 1300 Solid ER312 Dilution (%) Armox 440 Figure 30: Armox ER312 Solidification Diagram 85

106 Figure 31 is a Schaeffler Diagram with a tie line connecting Armox 440 and ER312. The pure filler metal (ER312) is predicted to be a mixture of austenite and ferrite (roughly 40% ferrite). Assuming weld metal dilutions in the 30-40% range, the weld metal microstructure is projected to be comprised of austenite and ferrite (5-10% ferrite). For the pure base metal (a.k.a. HAZ), the Schaeffler diagram predicts a martensitic structure. As the composition deviates from pure base metal to the filler metal, a martensitic structure is still expected down to (roughly) 65% dilution. This implies that a portion of the transition region from pure base metal to weld metal should be martensitic. 86

107 Ni Eq. = %Ni +30*%C + 0.5*%Mn Armox 440 F + M A+M M Armox ER312 A M+F No Ferrite A + F Cr Eq. = %Cr + %Mo + 1.5*%Si + 0.5*%Cb F Figure 31: Schaeffler Diagram connecting Armox 440 and ER RHA + Sandvik AXT The solidification diagram for RHA + Sandvik AXT is shown in Figure 32. All dilution levels experience ferrite-austenite (FA) solidification mode. This simulation result implies that, since there is no A solidification, solidification cracking is not an issue between RHA and Sandvik AXT (in theory). The largest solidification temperature range is 123 C, which occurs at pure RHA (100% dilution). The smallest solidification temperature range is 110 C, which occurs at 15% dilution. 87

108 Temperature ( C) 1500 Liquid 1450 L + A 1400 L + F L + F + A 1350 Solid Sandvik AXT Dilution RHA Figure 32: RHA + Sandvik AXT Solidification Diagram Figure 33 is a Schaeffler Diagram with a tie line connecting RHA and Sandvik AXT. The pure filler metal (Sandvik AXT) is predicted to be nearly pure austenite. Assuming weld metal dilutions in the 30-40% range, the weld metal microstructure is projected to be comprised of austenite and martensite. For the pure base metal (a.k.a. HAZ), the Schaeffler diagram predicts a martensitic structure. As the composition deviates from pure base metal to the filler metal, a martensitic structure is still 88

109 Ni Eq. = %Ni +30*%C + 0.5*%Mn expected down to (roughly) 50% dilution. This implies that a portion of the transition region from pure base metal to weld metal should be martensitic. RHA F + M A+M RHA + Sandvik AXT M M+F A Sandvik AXT No Ferrite A + F Cr Eq. = %Cr + %Mo + 1.5*%Si + 0.5*%Cb F Figure 33: Schaeffler Diagram connecting RHA and Sandvik AXT RHA + ER309LHF The solidification diagram for RHA + ER309LHF is shown in Figure 34. All dilution levels experience ferrite-austenite (FA) solidification mode. This simulation 89

110 Temperature ( C) result implies that, since there is no A solidification, solidification cracking is not an issue between RHA and ER309LHF (in theory). The largest solidification temperature range is 123 C, which occurs at pure RHA (100% dilution). The smallest solidification temperature range is 56 C, which occurs at pure ER309LHF (0% dilution) Liquid L + F L + F + A L + A Solid ER309LHF Dilution (%) RHA Figure 34: RHA + ER309LHF Solidification Diagram Figure 35 is a Schaeffler Diagram with a tie line connecting RHA and ER309LHF. The pure filler metal (ER309LHF) is predicted to be austenite with a small amount of 90

111 Ni Eq. = %Ni +30*%C + 0.5*%Mn ferrite (~10%). Assuming weld metal dilutions in the 30-40% range, the weld metal microstructure is projected to be almost fully austenitic with a small amount of martensite. For the pure base metal (a.k.a. HAZ), the Schaeffler diagram predicts a martensitic structure. As the composition deviates from pure base metal to the filler metal, a martensitic structure is still expected down to (roughly) 60% dilution. This implies that a portion of the transition region from pure base metal to weld metal should be martensitic. A+M RHA F + M M RHA + ER309LHF M+F A No Ferrite A + F Cr Eq. = %Cr + %Mo + 1.5*%Si + 0.5*%Cb F Figure 35: Schaeffler Diagram connecting RHA and ER309LHF 91

112 Temperature ( C) RHA + ER312 The solidification diagram for RHA + ER312 is shown in Figure 36. All dilution levels experience ferrite-austenite (FA) solidification mode. This simulation result implies that, since there is no A solidification, solidification cracking is not an issue between RHA and ER312 (in theory). The largest solidification temperature range is 188 C, which occurs at pure ER312 (0% dilution). The smallest solidification temperature range is 101 C, which occurs at 85% dilution ER312 Liquid L + F L + A L + F + A Solid Dilution (%) RHA Figure 36: RHA + ER312 Solidification Diagram 92

113 Figure 37 is a Schaeffler Diagram with a tie line connecting RHA and ER312. The pure filler metal (ER312) is predicted to be a mixture of austenite and ferrite (roughly 40% ferrite). Assuming weld metal dilutions in the 30-40% range, the weld metal microstructure is projected to be a mixture of austenite and ferrite (~5-10%). For the pure base metal (a.k.a. HAZ), the Schaeffler diagram predicts a martensitic structure. As the composition deviates from pure base metal to the filler metal, a martensitic structure is still expected down to (roughly) 65% dilution. This implies that a portion of the transition region from pure base metal to weld metal should be martensitic. 93

114 Ni Eq. = %Ni +30*%C + 0.5*%Mn RHA F + M A+M M RHA + ER312 A M+F No Ferrite A + F Cr Eq. = %Cr + %Mo + 1.5*%Si + 0.5*%Cb F Figure 37: Schaeffler Diagram connecting RHA and ER Thermodynamic Simulations: Summary and Result Analysis Table 13 provides a summary of the solidification modes as a function of dilution for each DMW. In all Armox 440 steel welds, there are dilutions that experience A solidification mode and are potentially susceptible to solidification cracking. These dilution ranges include 50-75% (Armox Sandvik AXT), 65-77% (Armox ER309LHF), and 75% (Armox ER312). The only solidification mode experienced in all RHA steel welds is FA. However, there are dilutions in some 94

115 welds (50-60% in RHA + Sandvik AXT and 55-65% in RHA + ER309LHF) that have nearly 90-95% austenite upon final solidification. Although the FA mode generally does not lead to solidification cracking, the high amount of austenite in these dilution ranges could result in cracking. This assumption is supported by the presence of solidification cracks in test welds of RHA steel. Base Metal Armox 440 RHA Table 13: Solidification Mode Summary Dilution Filler Metal Sandvik AXT FA AF A FA ER309LHF FA AF A FA ER312 FA A FA Sandvik AXT FA ER309LHF FA ER312 FA Figure 38 is a Schaeffler Diagram with tie lines connecting Armox 440 and RHA with the three stainless steel consumables. For weld metal, lower dilutions (closer to the filler metal composition) are expected. Assuming dilutions in the 30-40% range for each weld combination, then the following microstructures would result in the weld metal: 1. Armox Sandvik AXT: mixture of austenite and martensite. 95

116 2. Armox ER309LHF: primarily austenite with minor amounts of martensite. 3. Armox ER312: austenite with [roughly] 5-10% ferrite. 4. RHA + Sandvik AXT: mixture of austenite and martensite. 5. RHA + ER309LHF: primarily austenite with minor amounts of martensite. 6. RHA + ER312: austenite with [roughly] 5-10% ferrite. For the pure base metal (a.k.a. HAZ), the Schaeffler diagram predicts a martensitic structure for both Armox 440 and RHA. As the composition deviates from pure base metal to the filler metal, a martensitic structure is still expected for both armored steels down to (roughly) 50-65% dilution. This implies that a portion of the transition region from pure base metal to weld metal should be martensitic. 96

117 Ni Eq. = %Ni +30*%C + 0.5*%Mn A+M A No Ferrite Sandvik AXT 10% Ferrite ER RHA Armox 440 F + M M M+F A + F Cr Eq. = %Cr + %Mo + 1.5*%Si + 0.5*%Cb Armox ER309LHF Armox ER312 Armox 440 +ER307 RHA + ER309LHF RHA + ER312 RHA + ER307 Figure 38: Schaeffler Plot with all DMW combinations F Metallurgical Characterization Metallurgical characterization is an important aspect of this study because it illustrates the microstructures present in each DMW combination. With a clear picture of the microstructure, a better understanding of the susceptibility to HIC and/or solidification cracking will be developed. The following section is dedicated to thoroughly explaining each microconstituent that is present in each DMW combination. The explanation of the formation of each DMW microstructure, as well 97

118 as a brief analysis, will support the conclusions that were determined based on the results of this study. The primary focus of these DMW microstructures includes the fusion boundary region, which encompasses the CGHAZ, planar growth region, cellular/cellular dendritic region, and weld swirls. The analysis of each DMW consists of a description of the microstructure with the aid of light optical microscopy (LOM), hardness mapping, and energy-dispersive x-ray spectroscopy (EDS) Armox Sandvik AXT The resulting microstructures produced in the Armox Sandvik AXT (~ER307) weld are summarized in Figure 39. The microstructure consisted of a CGHAZ, planar growth region, cellular dendritic weld metal, and weld swirls. Planar growth was found along the fusion boundary. As solidification occurred, the planar growth region formed, grew some distance, broke down, and cellular/cellular dendritic weld metal began to solidify. The Armox Sandvik AXT weld consisted of austenite and very small amounts of delta ferrite. The white phase that makes up a majority of the weld metal is austenite and the dark, inter-dendritic phase is the delta ferrite. The minimal amount of delta ferrite leads to poor visibility of the austenite dendrites at high magnification. Weld swirls were present in every weld along the weld legs and in the weld root, with larger swirls experienced along the weld legs. These swirls are islands and peninsulas of molten base metal that is swept into the weld pool and is partially mixed with the weld metal [20]. The structure within the weld swirls varied depending on the level of melting experienced by the particular weld swirl. 98

119 Figure 39: Armox Sandvik AXT [A] CGHAZ [B] Weld Swirl along top leg [C] Planar Growth [D] Weld Metal Figure 40 shows the hardness map of the fusion boundary, including the weld metal and CGHAZ, in the Armox Sandvik AXT weld. This weld had a heat input of 1.6 kj/mm. The highest hardness recorded was 483 Vickers, which was experienced in the weld swirl in the top leg. The weld swirls located in the top and bottom legs experienced the highest hardness values in the entire weld. The average 99

120 hardness in the weld metal was 252 Vickers and the average hardness in the CGHAZ was 428 Vickers. Figure 40: Armox Sandvik AXT [A] Macro View [B] Hardness Map overlaid on the Macro View EDS analysis was performed in order to determine the compositional gradients across the fusion boundary and in the weld metal, along with the compositions in typical microstructual constituents (i.e. cellular dendritic, planar growth zones, and weld swirls). The composition was then used to calculate the dilution. The EDS lines were placed on preselected microstructural constituents and the composition was then used to calculate the dilution. An example of this analysis is shown in Figure 41, which displays an EDS line scan in the Armox Sandvik AXT weld 100

121 combination. The results were plotted on the solidification diagram for each DMW. Such plots provide a direct correlation of the microstructural constituents in the fusion zone to the solidification temperature range, solidification mode, and dilution. These plots can be used for prediction of potential susceptibility to weld solidification cracking in dissimilar metal welds. Figure 42 depicts the Armox Sandvik AXT solidification diagram (Figure 26) overlaid with the dilutions that correspond to the solidification morphology and fusion boundary microstructural constituents from the EDS analysis. The planar growth region encompasses dilutions ranging from 53-96% and the cellular dendritic region encompasses 23-53% dilution. The average width of the planar growth region was 13.8 μm. Within the planar growth region, two solidification modes are possible: A (53-75% dilution) and FA (75-96% dilution). Within the cellular dendritic region, three solidification modes are possible: AF (45-50% dilution), A (50-53% dilution), and FA (23-45% dilution) solidification mode according to the solidification simulations. With A solidification mode being possible in the cellular dendritic region, solidification cracking is a possibility in this weld combination due to the low solubility of impurity elements in austenite. 101

122 Figure 41: Dilution determined using EDS Line Scan in Armox Sandvik AXT with a [A] Optical Microscopy Image [B] DIC Image [C] Dilution and Solidification Temperature Range vs. Distance Chart 102

123 Figure 42: Armox Sandvik AXT Solidification Diagram overlaid with EDS results Armox ER309LHF The resulting microstructures produced in the Armox ER309LHF weld are summarized in Figure 43. The microstructure consisted of a CGHAZ, planar growth region, cellular dendritic weld metal, and weld swirls. Planar growth was found along the fusion boundary. As solidification occurred, the planar growth region formed, grew some distance, broke down, and cellular/cellular dendritic weld metal began to solidify. The Armox ER309LHF weld consisted of austenite and delta ferrite. The white phase that makes up a majority of the weld metal is austenite and the dark, inter-dendritic phase is the delta ferrite. Weld swirls were present in every weld along the weld legs and in the weld root, with larger swirls experienced along the weld legs. 103

124 Figure 43: Armox ER309LHF [A] CGHAZ [B] Weld Swirls in the weld root [C] Planar Growth [D] Weld Metal Figure 44 shows the hardness map of the fusion boundary, including the weld metal and CGHAZ, in the Armox ER09LHF weld. This weld had a heat input of 1.7 kj/mm. The highest hardness recorded was 525 Vickers, which was experienced in the weld swirl in the top leg. The weld swirl located in the top leg experienced the highest hardness values in the entire weld. The average hardness in 104

125 the weld metal was 210 Vickers and the average hardness in the CGHAZ was 453 Vickers. Figure 44: Armox ER309LHF [A] Macro View [B] Hardness Map overlaid on the Macro View Figure 45 depicts the Armox ER309LHF solidification diagram (Figure 28) overlaid with the dilutions that correspond to the solidification morphology and fusion boundary microstructural constituents from the EDS analysis. The planar growth region encompasses dilutions ranging from 54-96% and the cellular dendritic region encompasses 22-54% dilution. The average width of the planar growth region was 15.5 μm. Within the planar growth region, three solidification modes are possible: AF (54-65% dilution), A (65-77% dilution), and FA (77-96% dilution). 105

126 Within the cellular dendritic region, two solidification modes are possible: AF (50-54% dilution) and FA (22-50% dilution) solidification mode according to the solidification simulations. With A solidification mode not present in the cellular dendritic region, the possibility of solidification cracking is reduced because the impurity elements are more soluble in ferrite. Figure 45: Armox ER309LHF Solidification Diagram overlaid with EDS results Armox ER312 The resulting microstructures produced in the Armox ER312 weld are summarized in Figure 46. The microstructure consisted of a CGHAZ, planar growth region, cellular dendritic weld metal, and weld swirls. Planar growth was found 106

127 along the fusion boundary. As solidification occurred, the planar growth region formed, grew some distance, broke down, and cellular/cellular dendritic weld metal began to solidify. The Armox ER312 weld consisted of austenite and delta ferrite. The white phase that makes up a majority of the weld metal is austenite and the dark, inter-dendritic phase is the delta ferrite. Weld swirls were present in every weld along the weld legs and in the weld root, with larger swirls experienced along the weld legs. Figure 46: Armox ER312 [A] CGHAZ [B] Weld Swirl in the top leg [C] Planar Growth along the fusion boundary [D] Fusion Boundry and Weld Metal 107

128 Figure 47 shows the hardness map of the fusion boundary, including the weld metal and CGHAZ, in the Armox ER312 weld. This weld had a heat input of 1.9 kj/mm. The highest hardness recorded was 542 Vickers, which was experienced in the weld swirl in the top leg. The weld swirl located in the top leg experienced the highest hardness values in the entire weld. The average hardness in the weld metal was 236 Vickers and the average hardness in the CGHAZ was 453 Vickers. Figure 47: Armox ER312 [A] Macro View [B] Hardness Map overlaid on top of Macro View 108

129 Figure 48 depicts the Armox ER312 solidification diagram (Figure 30) overlaid with the dilutions that correspond to the solidification morphology and fusion boundary microstructural constituents from the EDS analysis. The planar growth region encompasses dilutions ranging from 52-95% and the cellular dendritic region encompasses 23-52% dilution. The average width of the planar growth region was 19.7 μm. Within the planar growth region, two solidification modes are possible: A (75% dilution) and FA (52-95% dilution except 75%). Within the cellular dendritic region, only FA solidification mode is predicted according to the solidification simulations. With FA solidification mode having a low susceptibility to solidification cracking and ferrite having a high solubility for impurity elements, it is unlikely that solidification cracking is an issue in this weld combination. Figure 48: Armox ER312 Solidification Diagram overlaid with EDS results 109

130 5.2.4 RHA + Sandvik AXT The resulting microstructures produced in the RHA + Sandvik AXT (~ER307) weld are summarized in Figure 49. The microstructure consisted of a CGHAZ, planar growth region, cellular dendritic weld metal, and weld swirls. Planar growth was found along the fusion boundary. As solidification occurred, the planar growth region formed, grew some distance, broke down, and cellular/cellular dendritic weld metal began to solidify. The RHA + Sandvik AXT weld consisted of austenite and small amounts of delta ferrite. The white phase that makes up a majority of the weld metal is austenite and the dark, inter-dendritic phase is the delta ferrite. Weld swirls were present in every weld along the weld legs and in the weld root, with larger swirls experienced along the weld legs. 110

131 Figure 49: RHA + Sandvik AXT [A] CGHAZ [B] Weld Swirl along the top leg [C] Planar Growth [D] Weld Metal Figure 50 shows the hardness map of the fusion boundary, including the weld metal and CGHAZ, in the RHA + Sandvik AXT weld. This weld had a heat input of 1.0 kj/mm. The highest hardness recorded was 650 Vickers, which was experienced in the weld swirl in the top leg. The weld swirls located in the top and bottom legs experienced the highest hardness values in the entire weld. The average hardness in the weld metal was 263 Vickers and the average hardness in the CGHAZ was 535 Vickers. 111

132 Figure 50: RHA + Sandvik AXT [A] Macro View [B] Hardness Map overlaid on Macro View Figure 51 depicts the RHA + Sandvik AXT solidification diagram (Figure 32) overlaid with the dilutions that correspond to the solidification morphology and fusion boundary microstructural constituents from the EDS analysis. The planar growth region encompasses dilutions ranging from 55-96% and the cellular dendritic region encompasses 10-55% dilution. The average width of the planar growth region was 23.0 μm. In both the planar growth region and cellular dendritic region, only FA solidification mode is expected according to the solidification simulations. With FA solidification mode having a low susceptibility to solidification cracking and ferrite having a high solubility for impurity elements, it is unlikely that solidification cracking is an issue in this weld combination. However, while FA solidification mode is present between 50-55% dilution, there is less than 5% ferrite present upon 112

133 final (98%) solidification. There amount of ferrite may be insufficient to mitigate solidification cracking. Figure 51: RHA + Sandvik AXT Solidification Diagram overlaid with EDS results RHA + ER309LHF The resulting microstructures produced in the RHA + ER309LHF weld are summarized in Figure 52. The microstructure consisted of a CGHAZ, planar growth region, cellular dendritic weld metal, and weld swirls. Planar growth was found along the fusion boundary. As solidification occurred, the planar growth region formed, grew some distance, broke down, and cellular/cellular dendritic weld metal began to solidify. The RHA + ER309LHF weld consisted of austenite and delta ferrite. The white phase that makes up a majority of the weld metal is austenite and 113

134 the dark, inter-dendritic phase is the delta ferrite. Weld swirls were present in every weld along the weld legs and in the weld root, with larger swirls experienced along the weld legs. Figure 52: RHA + ER309LHF [A] CGHAZ and Weld Swirl in the weld root [B] Weld Swirl along the top leg [C] Planar Growth [D] Weld Metal Figure 53 shows the hardness map of the fusion boundary, including the weld metal and CGHAZ, in the RHA + ER309LHF weld. This weld had a heat input of 114

135 1.1 kj/mm. The highest hardness recorded was 639 Vickers, which was experienced in the weld swirl in the bottom leg. The weld swirls located in the top and bottom legs experienced the highest hardness values in the entire weld. The average hardness in the weld metal was 250 Vickers and the average hardness in the CGHAZ was 535 Vickers. Figure 53: RHA + ER309LHF [A] Macro View [B] Hardness Map overlaid on Macro View Figure 54 depicts the RHA + ER309LHF solidification diagram (Figure 34) overlaid with the dilutions that correspond to the solidification morphology and fusion boundary microstructural constituents from the EDS analysis. The planar growth region encompasses dilutions ranging from 54-93% and the cellular dendritic region encompasses 17-55% dilution. The average width of the planar growth region 115

136 was 11.1 μm. In both the planar growth region and cellular dendritic region, only FA solidification mode is expected according to the solidification simulations. With FA solidification mode having a low susceptibility to solidification cracking and ferrite having a high solubility for impurity elements, it is unlikely that solidification cracking is an issue in this weld combination. However, while FA solidification mode is present between 55-65% dilution, there is less than 5% ferrite present upon final (98%) solidification. There amount of ferrite may be insufficient to mitigate solidification cracking. Figure 54: RHA + ER309LHF Solidification Diagram overlaid with EDS results 116

137 5.2.6 RHA + ER312 The resulting microstructures produced in the RHA + ER312 weld are summarized in Figure 55. The microstructure consisted of a CGHAZ, planar growth region, cellular dendritic weld metal, and weld swirls. Planar growth was found along the fusion boundary. As solidification occurred, the planar growth region formed, grew some distance, broke down, and cellular/cellular dendritic weld metal began to solidify. The RHA + ER312 weld consisted of austenite and delta ferrite. The white phase that makes up a majority of the weld metal is austenite and the dark, interdendritic phase is the delta ferrite. Weld swirls were present in every weld along the weld legs and in the weld root, with larger swirls experienced along the weld legs. 117

138 Figure 55: RHA + ER312 [A] CGHAZ [B] Weld Swirl along top leg [C] Planar Growth [D] Fusion Boundary and Weld Metal Figure 56 shows the hardness map of the fusion boundary, including the weld metal and CGHAZ, in the RHA + ER312 weld. This weld had a heat input of 1.1 kj/mm. The highest hardness recorded was 639 Vickers, which was experienced in the weld swirl in the top leg. The weld swirls located in the top and bottom legs experienced the highest hardness values in the entire weld. The average hardness in the weld metal was 260 Vickers and the average hardness in the CGHAZ was 517 Vickers. 118

139 Figure 56: RHA + ER312 [A] Macro View [B] Hardness Map overlaid on Macro View Figure 57 depicts the RHA + ER312 solidification diagram (Figure 36) overlaid with the dilutions that correspond to the solidification morphology and fusion boundary microstructural constituents from the EDS analysis. The planar growth region encompasses dilutions ranging from 51-96% and the cellular dendritic region encompasses 25-51% dilution. The average width of the planar growth region was 16.0 μm. In both the planar growth region and cellular dendritic region, only FA solidification mode is expected according to the solidification simulations. With FA solidification mode having a low susceptibility to solidification cracking and ferrite 119

140 having a high solubility for impurity elements, it is unlikely that solidification cracking is an issue in this weld combination. Figure 57: RHA + ER312 Solidification Diagram overlaid with EDS results Metallurgical Characterization: Summary and Result Analysis Hydrogen Induced Cracking Table 14 is comprised of a summary of hardness values from the hardness maps. The critical hardness for HIC to occur was calculated using the Duren equations [46] and was found to be 457 Vickers for Armox 440 and 513 Vickers for RHA. These calculation were based on the heat inputs commonly used ( kj/mm), an assumption of 5 ml/100 g of hydrogen, and an assumption that no preheat was 120

141 applied. The average HAZ hardness was close to or exceeded the critical hardness in every DMW. This result, along with the martensitic structure of the HAZ, implies that the HAZ could be susceptible to HIC if sufficient hydrogen and stress is present. To minimize the hardness, preheat should be applied to each weld. The Duren equations specify a preheat of 122 C for Armox 440 and 146 C for RHA should be sufficient to reduce the hardness below the critical hardness. The hardest regions of every weld were the weld swirls, which exceeded average HAZ hardness and the critical hardness for HIC to occur. The critical hardness criterion of Duren cannot be applied to the weld swirls because the composition of the critical hardness corresponds to pure base material and the weld swirls have partial mixing with the filler metal. However, the high hardness could lead to HIC if hydrogen is trapped in the weld swirls [19]. Furthermore, the weld swirls were located at the root and along the legs (near the toes) of the welds, which typically have an increased stress concentration. These weld swirls are susceptible regions for cracking to occur. 121

142 Table 14: Summary of Hardness Values Base Metal Armox 440 RHA Filler Metal ER307 ER309LHF ER312 ER307 ER309LHF ER312 Heat Input (kj/mm) Maximum Hardness (HV) Minimum Hardness (HV) Weld Metal Average Hardness (HV) Weld Swirl Average Hardness (HV) HAZ Average Hardness (HV) Predicted Critical HAZ Hardness for HIC to occur (HV) Figure 58 displays the distribution of hardness values found in each weld combination. The large curves toward the right side of the diagram are hardness values found in the HAZ. The small curves toward the left side of the diagram are hardness values found in the weld metal. As can be seen, the RHA welds experience higher hardness values than Armox 440 welds in the HAZ. This, in large part, is due to the increased carbon content of RHA (0.26 wt%) over Armox 440 (0.19 wt%). The higher hardness indicates that RHA welds are more prone to HIC over Armox 440 welds. While RHA welds experience higher hardness' than Armox 440 welds, Armox 440 has better hardenability than RHA due to the higher Cr, Ni, and Mo content. Therefore, if higher heat inputs were used, it is expected that Armox 440 welds would exhibit higher hardness values than RHA welds. 122

143 Figure 58: Hardness Distribution Solidification Cracking Figure 59 compares the EDS results of all the DMW's. The green region represents the dilution levels that resulted in planar growth and the yellow region represents the dilution levels that resulted in a cellular/cellular dendritic morphology. The hashed areas correspond to dilutions that had A solidification (according to the Scheil solidification simulations) and the dark lines outline dilution levels found in weld swirls. The transition from planar growth to cellular dendritic is fairly consistent in these DMW's. The planar growth begins between 93-98% dilution (close to pure base metal) and breaks down between 52-56% dilution. The cellular 123

144 dendritic region begins at 52-56% dilution. As the distance from the planar growth region increases, the cellular dendritic region would plateau between 33-40% dilution in all DMW's. In the Armox ER309LHF and Armox ER312 welds, the A solidification zone lies in the planar growth region. Since segregation along solidification grain boundaries does not occur in planar growth structures, solidification cracking is unlikely to occur despite the A solidification mode. This implies that, in theory, these two welds should be immune to solidification cracking. In the Armox ER307 (Sandvik AXT) weld, the A solidification zone lies in the planar growth region and cellular dendritic region. Interdendritic segregation (along solidification grain boundaries) can occur in the cellular dendritic region, thus making solidification cracking a potential issue in this DMW combination. In the RHA welds, no solidification mode outside of FA mode was found. However, as stated previously, some dilutions ranges could be susceptible to solidification cracking due to the high volume fraction austenite formed during solidification. Both planar growth and cellular dendritic structures were the result of the 50-60% dilution range in RHA + ER307 welds and 55-65% dilution range in RHA + ER309LHF. Both dilution ranges contained a high amount of austenite. With these susceptible dilution ranges resulting in a cellular dendritic structure, segregation of impurities that form low melting constituents can occur and solidification cracking is a potential issue. 124

145 In every DMW, the weld swirls were found at dilutions ranging from 28 to 98%. The optical microscopy has revealed that the weld swirls contained various structures such as planar growth, cellular/cellular dendritic, and coarse grains resembling the CGHAZ. This implies that the extent of melting & mixing in each swirl varied and changed the microstructure of the swirl. Portions of some weld swirls completely melted and resulted in a cellular dendritic structure, which could be susceptible to solidification cracking. Other portions of some weld swirls had very little melting and could be susceptible to HIC because hydrogen could be trapped in the partially melted section of the swirl. Figure 59: EDS and Solidification Mode Comparison 125

146 5.3 Weldability Testing Weldability assumptions based on theoretical knowledge can be made from the modeling and characterization results. However, further testing is required to verify these assumptions and give quantitative results. The purpose of this section is to present the results of both the Tekken and CTS Testing of both preheated and nonpreheated welds. The results of both tests either verify or challenge the assumptions made in the previous sections. Although a plethora of results are presented, further testing needs to be conducted in order to support all conclusions. In the Tekken Test specimens, cracking occurred in 5 different locations of the weld, which are highlighted in Figure 60. In the CTS Test specimens, cracking is supposed to occur in the HAZ (crack location #3 in Figure 60B). If cracking occurs in the weld metal, then the test is considered invalid if the crack exceeds 5% of the throat thickness. All cracks found in CTS Test welds were located in the weld metal at locations #1 and #2 in Figure 60B. Figure 60: [A] Crack locations in Tekken samples [B] Crack locations in CTS samples 126

147 5.3.1 Armox Sandvik AXT In the Tekken Tests, cracking occurred in both preheat and no preheat condition. Cracking occurred in the middle of the weld metal (location #1 in Figure 60A) in the no preheat condition as shown in Figure 61. In the preheated condition, cracking was observed in the middle of the weld metal and near weld swirls along the top leg (location #3 in Figure 60A). The extent of cracking was similar in both conditions, as the crack section ratio was 8% in the preheat condition and 6% in the no preheat condition. From Figure 61B, the crack occurs along a solidification grain boundary. Similar to the Tekken Test results, cracking occurred in both preheat and no preheat condition in the CTS Test. In both conditions, cracking occurred in the weld metal near the weld swirls (location #1 in Figure 60B) as shown in Figure 61. While it's difficult to determine if the crack initiated at the weld swirl, the weld swirl may be an area of high stress upon solidification and could have led to the formation of the crack. The crack clearly goes through the weld swirl into the weld metal. The crack exceeded 5% of the throat thickness in this weld, which makes the test considered "invalid". 127

148 Figure 61: Armox Sandvik AXT [A] Tekken Test Macro (No Preheat) [B] Tekken Test Crack (No Preheat) [C] CTS Test Macro (No Preheat) [D] CTS Test Crack (No Preheat) Armox ER309LHF In the Tekken Tests, no cracking occurred in both preheat and no preheat condition. In the CTS Tests, cracking only occurred in the no preheat condition. Cracking in this test occurred in the weld metal (location #2 in Figure 60B) as shown in Figure 62. The crack is extremely small and is located on a solidification grain boundary indicating that it is potentially a solidification crack. 128

149 Figure 62: Armox ER309LHF CTS Test with No Preheat [A] Macro [B] Crack Armox ER312 In the Tekken Tests, cracking occurred in both preheat and no preheat condition. Cracking occurred in the middle of the weld metal (location #1 in Figure 60A) in both conditions. In the no preheat condition, cracking was also observed along the fusion boundary initiating from the stress concentrator (location #5 in Figure 60A). Figure 63A is the no preheat condition and shows both a crack in the middle of the weld metal and along the fusion boundary. The extent of cracking was much less in the preheat condition (2% crack section ratio) than the no preheat condition (18% crack section ratio). From Figure 63B, the weld metal crack occurs along a solidification grain boundary. 129

150 In the CTS Tests, cracking only occurred in no preheat condition and the crack was located in the weld metal (location 2 in Figure 60B) as shown in Figure 63C. The crack occurred along a solidification grain boundary. Under high magnification, some crack healing is observed at the end of the crack nearest the stress concentrator. This indicates that this is a solidification crack. The crack exceeded 5% of the throat thickness, which indicates that this test is considered "invalid". Figure 63: Armox ER312 [A] Tekken Test Macro (No Preheat) [B] Tekken Test Crack (No Preheat) [C] CTS Test Macro (No Preheat) [D] CTS Test Crack (No Preheat) 130

151 5.3.4 Armox ER70S-6 In both the Tekken and CTS Tests, cracking occurred in both preheat and no preheat condition. In both CTS Test welds and the Tekken Test weld (with preheat), cracking occurred in the middle of the weld metal and near weld swirls (locations #1 and #2 in Figure 60B, locations #1 and #3 in Figure 60A). The crack section ratio in the Tekken Test with preheat was 8%. The cracks that occurred in the CTS tests exceeded 5% of the throat thickness, which made both tests considered "invalid". In the Tekken Test weld with no preheat, cracking initiated at the stress concentrator, propagated through the HAZ, crossed over into the weld metal, and broke through the weld surface (crack section ratio = 100%) as shown in Figure 64A. Cracking initiated at the stress concentrator, which is what is supposed to happen in the Tekken Test (as opposed to cracking in the middle of the WM). With the crack propagating through both the HAZ and WM, the test shows that both are susceptible to HIC. 131

152 Figure 64: Armox ER70S-6 Tekken Test with No Preheat Armox ER100 In the Tekken Tests, cracking occurred only in the no preheat condition. Cracking initiated at the stress concentrator, propagated through the HAZ, crossed over into the weld metal, and broke through the weld surface (crack section ratio = 100%) as shown in Figure 65 (location #4 in Figure 60A). Cracking initiated at the stress concentrator, which is what is supposed to happen in the Tekken Test (as opposed to cracking in the middle of the WM). With the crack propagating through both the HAZ and WM, the test shows that both are susceptible to HIC. No cracking occurred in the CTS test welds. 132

153 Figure 65: Armox ER100 Tekken Test with No Preheat RHA + Sandvik AXT In the Tekken Tests, no cracking occurred in both preheat and no preheat condition. In the CTS Tests, cracking only occurred in the no preheat condition. Cracking in this test occurred in the weld metal (location #2 in Figure 60B) as shown in Figure 66. The crack lies along a solidification grain boundary indicating that it's potentially a solidification crack. 133

154 Figure 66: RHA + Sandvik AXT CTS Test with No Preheat [A] Macro [B] Crack RHA + ER309LHF In the Tekken Tests, cracking occurred in only the no preheat condition. Cracking occurred in the middle of the weld metal (location #1 in Figure 60A) as shown in Figure 67A. The extent of cracking is very small, as the crack section ratio was 3% in this test. The weld metal crack occurs along a solidification grain boundary. Similar to the Tekken Test results, cracking occurred in only the no preheat condition in the CTS Test. Cracking again occurred in the weld metal (location #2 in Figure 60B) as shown in Figure 67C. The crack, again, occurred along a solidification grain boundary. 134

155 Figure 67: RHA + ER309LHF [A] Tekken Test Macro (No Preheat) [B] Tekken Test Crack (No Preheat) [C] CTS Test Macro (No Preheat) [D] CTS Test Crack (No Preheat) RHA + ER312 In the Tekken Tests, cracking occurred in only the no preheat condition. Cracking occurred in the middle of the weld metal and along the fusion boundary (locations #1 and #5 in Figure 60A). Figure 68 is the no preheat condition and shows both a crack in the middle of the weld metal and along the fusion boundary. The crack section ratio was measured to be 27%. The crack in the middle of the weld metal occurs 135

156 along a solidification grain boundary and has evidence of crack healing, which indicates that this is a solidification crack. No cracking occurred in the CTS Test welds. Figure 68: RHA +ER312 Tekken Test (No Preheat) [A] Macro [B] Cracks RHA + ER70S-6 In the Tekken and CTS Tests, all welds experienced cracking in both preheat and no preheat conditions. In the Tekken Test with preheat and both CTS Tests, cracking occurred in the weld metal (locations #1, #2, and #3 in Figure 60A, location #1 in Figure 60B). The cracks in the CTS Tests exceeded 5% of the throat thickness making the test considered "invalid". In the Tekken Test with preheat, cracking is observed from a weld swirl in Figure

157 In the Tekken Test with no preheat, cracking occurred in location #4 in Figure 60A similar to the Armox ER70S-6 Tekken Test with no preheat. Cracking initiated at the stress concentrator, which is what is supposed to happen in the Tekken Test (as opposed to cracking in the middle of the WM). With the crack propagating through both the HAZ and WM, the test shows that both are susceptible to HIC. Figure 69: RHA + ER70S-6 Tekken Test with No Preheat RHA + ER100 In the Tekken and CTS Tests, all welds experienced cracking in both preheat and no preheat conditions. In the Tekken Test with preheat and both CTS Tests, cracking occurred in the weld metal (locations #2 and #3 in Figure 60A, location #1 in Figure 60B). The cracks in the CTS Tests exceeded 5% of the throat thickness making the 137

158 test considered "invalid". Figure 70B shows a crack along the top leg of the CTS Test with no preheat. In the Tekken Test with no preheat, cracking occurred in location #4 in Figure 60A similar to the RHA + ER70S-6 Tekken Test with no preheat. This is shown in Figure 70A. Cracking initiated at the stress concentrator, which is what is supposed to happen in the Tekken Test (as opposed to cracking in the middle of the WM). With the crack propagating through both the HAZ and WM, the test shows that both are susceptible to HIC. Figure 70: RHA + ER100 [A] Tekken Test with No Preheat [B] CTS Test with No Preheat Weldability Testing: Summary and Result Analysis Table 15 is a summary of the results from the Tekken Tests. The filler metal that experienced the least severe cracking was ER309LHF in both Armox 440 and RHA. 138

159 Only one ER309LHF Tekken Test contained a crack (RHA, no preheat); however, it had a crack section ratio of 3% and was a small crack compared to the other stainless steel filler metals. Sandvik AXT also experienced minor cracking, as the other Tekken tests that cracked had a crack section ratio < 10%. ER70S-6 and ER100 experienced the most severe cracking, as both cracked 100% in the no-preheat condition. Table 15: Summary of Tekken Test Results Table 16 is a summary of the results from the CTS Tests. The filler metal that experienced the least severe cracking was ER312 in both Armox 440 and RHA. Only 139

160 one ER309LHF Tekken Test cracked (Armox 440, no preheat) and it was verified to be a solidification crack. Sandvik AXT and ER70S-6 experienced the most severe cracking, as all tests experienced cracking. Both ER309LHF and ER100 experienced cracking in two tests and no cracking in two tests. Table 16: Summary of CTS Test Results In all cracked CTS Test samples, all cracks were located in the weld metal and no cracks were found in the HAZ. The weld metal crack lengths exceeded 5% of the throat thickness in each test, making each test invalid to HIC testing. Similarly, all Tekken Test samples made with a stainless steel consumable cracked in the weld 140

161 metal. While this does not mean the Tekken tests in invalid, it means that HIC did not occur because the crack initiation was not at the stress concentrator. This could have occurred for two reasons: 1. stress was relieved due to solidification cracking and/or 2. hydrogen is trapped in the austenitic weld metal. If cracks initiated at the stress concentrator, it would indicate HIC according to ISO [10]. Otherwise, the cracking may be of a different form and no data on the susceptibility to HIC was acquired. Tekken Test samples made with low-alloy steel consumables did crack in the HAZ and weld metal after initiating at the stress concentrator and were confirmed to be HIC. Although the stainless steel weldability tests provide no data on the susceptibility to HIC, the results do provide valuable information. The results of these tests show that in highly restrained conditions, welds produced with stainless steel filler metals are more susceptible to solidification cracking than HIC. To verify that these cracks are solidification cracks, characterization and fractography was performed on these samples. Figure 71 displays an example of a crack and fracture surface of a crack in a weld of each stainless steel filler metal. The microstructure around the crack leads to the notion of solidification cracking due to the cracks being on solidification grain boundaries. Some cracks had evidence of crack back-healing, which is common in solidification cracks. The fracture surfaces shown in Figure 71 have a dendritic morphology, which secures the conclusion that these cracks are solidification cracks. The fracture surface of the cracks was consistent in all cracked samples. 141

162 Figure 71: [A] Armox Sandvik AXT (No Preheat) WM crack [B] Fracture surface of Armox Sandvik AXT WM crack [C] RHA + ER309LHF (No Preheat) WM crack [D] Fracture surface of RHA + ER309LHF WM crack [E] RHA + ER312 (No Preheat) WM crack [F] Fracture surface of RHA + ER312 WM crack 142

163 Filler metals ER70S-6 and ER100 are currently used to weld Armox 440 and RHA. Similar to the stainless steel filler metals, cracking often occurred in the weld metal in the weldability tests. Figure 72 displays examples of cracks in welds made with ER70S-6 and ER100 and the resulting fracture surfaces. These cracks occur along solidification grain boundaries. However, unlike the cracks found in the stainless steel filler metals, these cracks do not exist in the middle of the weld. If cracking occurred in the weld metal of a weld made with a low-alloy consumable, the crack occurred along the top leg near a weld metal swirl (shown in Figure 72). The weld swirls can be points of high restraint and, upon final solidification, can lead to solidification cracking. The fracture surfaces of these cracks revealed liquid along the fracture surface, leading to the conclusion that these were solidification cracks. 143

164 Figure 72: [A] RHA + ER70S-6 (No Preheat) WM crack [B] Fracture surface of RHA + ER70S-6 (No Preheat) WM crack [C] RHA + ER100 (Preheat) WM crack [D] Fracture surface of RHA + ER100 (Preheat) WM crack Solidification cracks occurred in the welds made with low-alloy steels because of the increased sulfur and phosphorus content. From Table 12 in section 4.2, ER70S-6 and ER100 contain more sulfur and phosphorus than all three of the stainless steel consumables in the pure form. These impurity elements segregate to the grain boundaries during solidification and result in a low melting liquid present along the boundaries. If sufficient restraint is present, these liquid films result in solidification 144

165 cracking. Table 17 shows the measured [average] dilutions in each weldability test with the appropriate composition of that dilution. All welds made with ER70S-6 and ER100 contained greater amount of sulfur and phosphorus, which confirm the notion of greater impurity levels in these welds. Table 17: Average Weld Metal Dilutions from Each Weldability Test Some Tekken Test welds made with ER70S-6 and ER100 propagated through the entire weld (crack section ratio = 100%). The fracture surface of the crack changed 145

166 as the crack propagated as shown in Figure 73. The fracture surface was intergranular in the HAZ (Figure 73A), then adjusted to a cleavage surface (Figure 73B and C), and, finally, the crack ended with a ductile surface (Figure 73D). The change in fracture surface, along with the resulting fracture surfaces, indicates that this weld experienced HIC [40]. This is expected as it has been seen in practice with this weld combination. Similar fracture surfaces were observed in other weld combinations involving ER70S-6 and ER

167 Figure 73: Crack Fracture Surface Change in Armox ER70S-6 [A] Intergranular [B] Intergranular transitioning to quasi-cleavage at the fusion boundary [C] Quasi-cleavage near the weld swirl [D] Microvoid Coalescence in the WM In both the CTS and Tekken Test, application of preheat reduced the severity of the cracking for all weld combinations. While preheat does not have a big effect on solidification cracking, it is believed this occurred for two reasons. First, preheat was applied after the anchor welds were made in both tests. The preheat assists in 147