Compositional Optimization, Mechanical Properties, and Tempering Response in Type 410 Stainless
Steel Welds
Thesis
Presented in Partial Fulfillment of the Requirements for the Degree Master of Science in the
Graduate School of The Ohio State University
By
Benjamin James Lawson
Graduate Program in Welding Engineering
The Ohio State University
2019
Thesis Committee
Dr. Boian T. Alexandrov, Advisor
Dr. Avraham Benatar
Copyrighted by
Benjamin James Lawson
2019
Abstract
Industry experience with Type 410 steel welds in downstream hydro-processing installations, using generic welding consumables, has shown difficulties in meeting the weld metal and heat affected zone (HAZ) hardness and toughness requirements. Recent research has pointed out the wide composition specifications of Type 410 base metal and welding consumables as the leading cause for significant hardness and toughness variations, related to exceeding the A1 temperature during post weld heat treatment (PWHT) and formation of fresh martensite, and to retention of significant amounts of delta ferrite both of which reduce material toughness. Previously developed predictive equations for the A1 temperature and the content of retained delta ferrite in
Type 410 steel were used for development of customized welding consumables with optimized chemical compositions. The predictive accuracy of these equations was validated through phase transformation analysis and metallurgical characterization using the customized welding consumables and commercially available base metals.
Test welds in two heats of Type 410 steel were produced using two customized metal core welding consumables and GTA manual welding process. The test welds were subjected to
PWHT at temperatures selected using the A1 temperature predictive equation and the ASTM recommended temperature range. Charpy impact testing on specimens extracted from the weld metal, the HAZ, and the base metal was performed at temperatures of 75F. Metallurgical characterization and hardness measurements were also performed. The test results met the ii NACE hardness and ASME toughness requirements for all base metal / welding consumable combinations, providing additional validation for the A1 temperature and the retained delta ferrite predictive equations.
The tempering response in weld metal of one customized consumable and in one base metal was evaluated using a series of heat treatment experiments. Holloman-Jaffe type parametric equations for prediction of weld metal and base metal hardness as a function of the tempering parameters were developed. It was demonstrated that selecting a PWHT temperature simply based on the
ASME recommended temperature range, without considering the A1 temperature of the particular welding consumable or base metal, could result in insufficient tempering or in formation of fresh martensite.
iii Dedication
To my mom and dad who have given me amazing support throughout my collegiate career at the
Ohio State University, and have shown me what it takes to make things happen throughout my
life.
To my brother who has always been a source of motivation and encouragement. His consistent
hard work is and always will be an inspiration to me.
To Leah Reynolds, for her encouragement, patience, and ability to brighten my day no matter what. I am excited for the adventures we will experience together as we tackle the next phase of
our life.
iv Acknowledgments
I would first like to thank my graduate adviser Dr. Boian Alexandrov for all his help and guidance in this project. He first brought me on as an undergraduate researcher in 2015, and the knowledge I have gained working with him since then has been insurmountable. The experience
I have gained in welding and metallurgy during my time with him will help me throughout my career as a welding engineer. I would also like to thank Dr. Avraham Benatar for taking the time to serve on my master’s committee chair and reviewing my thesis.
I would like to thank Dr. Jorge Penso from Shell for supporting this project. Thank you for your industry guidance, and input on the project. Thanks is also given to Joe Bundy, and David
Benson from Hobart Brothers Inc. for all their support in this project. All custom welding consumables, test welds, and impact tests were completed at Hobart under their supervision.
Their help was a need in being able to complete all the characterization I needed from the test welds, and custom consumables. Thanks to PhD candidate Jacob Wildofsky for his assistance on the light radiation furnace. Thanks is also given to the labs supervisor Ed Pfeifer, and PhD candidates Jim Rule and Tate Patterson for all their help and assistance on all other equipment used in this project.
v Vita
2013…………………………………….Worthington Kilbourne High School – Worthington, OH
2017……………………………………….B.S. Welding Engineering, The Ohio State University
2019……………………………...... M.S. Welding Engineering, The Ohio State University
Fields of Study
Major Field: Welding Engineering
vi
Table of Contents
Abstract ...... ii
Dedication ...... iv
Acknowledgments...... v
Vita ...... vi
List of Tables ...... xii
List of Figures ...... xiv
Chapter 1: Introduction ...... 1
Chapter 2: Literature Review ...... 4
2.1 Martensitic Stainless Steel ...... 4
2.1.1 Characteristics of Stainless Steel Alloys...... 4
2.1.2 Martensitic Stainless Steel ...... 5
2.1.3 Type 410 Martensitic Stainless Steel ...... 6
2.2 Phase Transformations ...... 7
2.2.1 Solidification ...... 7
vii 2.2.2 Austenite Martensite and Ferrite Phase Transformations ...... 9
2.2.3 Delta Ferrite ...... 13
2.3 Effects of Alloying Elements ...... 17
2.3.1 Carbon ...... 17
2.3.2 Chromium ...... 18
2.3.3 Copper ...... 18
2.3.4 Manganese ...... 19
2.3.5 Molybdenum ...... 19
2.3.6 Nitrogen ...... 19
2.3.7 Nickel ...... 20
2.3.8 Silicon ...... 20
2.4 Predictive Formulae and Composition Range Development for Type 410 Steel ...... 20
2.4.1 Background ...... 20
2.4.2 A1 Temperature Predictive Formula ...... 21
2.4.3 Delta Ferrite Retention Predictive Formula ...... 23
2.4.4 Optimized Composition Window ...... 26
2.5 Experimental A1 Transformation Analysis ...... 28
2.5.1 Dilatometry ...... 28
2.5.2 Single Sensor Differential Thermal Analysis (SS-DTA)...... 29
viii 2.6 Tempering Response of Type 410 Steels...... 31
2.6.1 Post Weld Heat Treatment ...... 31
2.6.2 Carbide Formation in Type 410 Stainless Steel ...... 32
2.6.3 Time-Temperature Tempering Response ...... 33
Chapter 3: Objectives ...... 36
Chapter 4: Materials and Procedure ...... 37
4.1. Materials ...... 37
4.2. Validation of A1 Temperature and Delta Ferrite Content Predictive Formulas ...... 38
4.2.1. Samples Preparation...... 38
4.2.2. Delta Ferrite Quantification ...... 41
4.2.3. Dilatometry ...... 43
4.2.4 Single Sensor Differential Thermal Analysis ...... 44
4.4 Test Welds ...... 45
4.5 Mechanical Testing of Weld Deposits ...... 47
4.5.1 Charpy V-Notch Tests ...... 47
4.5.2 Hardness Testing ...... 47
4.6 Metallurgical Characterization...... 47
4.6.1 Metallographic Sample Preparation ...... 47
4.6.2 Optical Metallography ...... 49
ix 4.7 Tempering Response Procedure ...... 49
Chapter 5: Results and Discussion of Predictive Formula Validation and Test Weld
Characterization ...... 52
5.1 Validation of A1 Predictive Formula ...... 52
5.2 Validation of the Delta Ferrite Predictive Formula ...... 56
5.3 Test Welds ...... 58
5.3.1 Impact Testing ...... 58
5.3.1.1 Alloy 4.2 Impact Tests ...... 58
5.3.1.2 Alloy 1.2 Impact Tests ...... 60
5.3.2 Hardness Traverses ...... 61
5.3.2.1 Alloy 4.2 Test Weld Hardness ...... 61
5.3.2.2 Alloy 1.2 Test Weld Hardness ...... 64
5.4 Metallurgical characterization ...... 65
5.4.1 Weld Metal...... 65
5.4.2 HAZ ...... 67
5.4.3 Base Metal ...... 69
Chapter 6: Tempering Response Study...... 72
6.1 Tempering Heat Treatment Results ...... 72
6.1.1 Alloy 4.2 Tempering Response...... 72
x 6.1.2 Base Metal B Tempering Response ...... 75
6.1.3 Holloman-Jaffe Parameter ...... 79
Chapter 7: Conclusions ...... 83
References ...... 86
Appendix A: Dilatometry Results ...... 90
Appendix B: SSDTA Results...... 94
Appendix C: As Cast Microstructures for Retained Delta Ferrite Quantification ...... 96
Appendix D: Impact Toughness Results...... 98
Appendix E: Tempering Hardness Data ...... 102
xi List of Tables
Table 1: Composition standards for Type 410 Stainless Steel ...... 7
Table 2: Mechanical standards for Type 410 steel for base metal and shielded metal arc welding electrodes ...... 7
Table 3 Chemical composition of custom made metal cored weld wire and commercially made base material (wt%). A1 and DF values based on predictive formulas. DF prediction based on
46oC/s cooling rate ...... 37
Table 4 Test weld plan ...... 46
Table 5 Test weld parameters including preheat temperature, voltage, amperage, travel speed and heat input ...... 46
Table 6 SS-DTA results for all alloys tested ...... 56
Table 7 All dilatometry results ...... 93
Table 8 Alloy 4.2 welded on Base Metal A Charpy V-notch toughness test results ...... 98
Table 9 Alloy 4.2 welded on Base Metal B Charpy V-notch toughness test results ...... 99
Table 10 Alloy 1.2 welded on Base Metal A Charpy V-notch toughness test results ...... 100
Table 11 Alloy 1.2 welded on Base Metal B Charpy V-notch toughness test results ...... 101
Table 12 Alloy 4.2 hardness data at 807oC tempering temperature at various times. *30 minute tempering time tempered at 820oC ...... 102
Table 13 Alloy 4.2 hardness data at 760oC tempering temperature at various times ...... 102 xii Table 14 Base Metal B hardness data at 766oC tempering temperature at various times ...... 103
Table 15 Base Metal B hardness data at 800oC tempering temperature at various times ...... 103
xiii List of Figures
Figure 1 Potential solidification modes of stainless steel [6] ...... 8
Figure 2 Phase stability in the Fe-Cr pseudo-binary phase diagram with varying carbon content
[1] ...... 10
Figure 3 Chromium’s effect on stabilizing the ferrite at a) 13% Cr and b) 17%Cr [1] ...... 11
Figure 4 Variation in A1 and A3 temperatures with increasing heating rates [10] ...... 12
Figure 5 Decreasing Ms temperature with decreasing cooling rate and varying austenitization temperatures [12] ...... 13
Figure 6 Charpy impact energy curves on super martensitic stainless steel with tempered martensite, fully martensitic, and 14% retained delta ferrite microstructures [17] ...... 16
Figure 7 Microstructures seen in a multi-pass bead on plate weld from a) the center of the weld and b) from the last pass of the weld [25] ...... 17
Figure 8 Possible A1 Temperature vs ASME B31.3 recommended PWHT with compositions within the ASTM A240 and AWS 5.4 and 5.9 requirements [28] ...... 21
Figure 9 Relative Alloy effect on A1 temperature [28] ...... 22
Figure 10 Pareto chart showing the elements with a significant effect of delta ferrite retention
[28] ...... 24
Figure 11 Relative effects of alloying elements on retained delta ferrite from DICTRA simulations [28] ...... 26 xiv Figure 12 Custom Type 410 steel composition ranges [28] ...... 27
Figure 13 Dilatometry curve showing austenite and martensite transformations in Fe-12Cr stainless steel [34] ...... 28
Figure 14 Phase transformations in E11018-MR weld metal detected by SS-DTA [36] ...... 30
Figure 15 An example of a) martensite transformation after heat treating above A1 and b) lack of martensite transformation when heated treated below A1 [30] ...... 31
Figure 16 Tempering temperature effect on toughness and hardness of Type 410 steel [38] ..... 33
Figure 17 Relation between measured hardness data and predicted hardness values using an
HJP with a C value of 18 [43] ...... 35
Figure 18 Nehrenberg HJP for high alloy steels using a C value of 20 [44]...... 35
Figure 19 Button melting apparatus ...... 38
Figure 20 a) funnel and mold used to cast cut weld wire into a cube b) SS-DTA samples before and after sectioning from cube cast ...... 40
Figure 21 Induction levitation melting device ...... 41
Figure 22 a) Copper mold used for dilatometry cast geometry b) 2” dilatometry specimen ...... 41
Figure 23 Copper mold used for delta ferrite study ...... 42
Figure 24 a) Dimension of the ceramic insert placed in copper mold b) Thermocouple placement and cast sample weight ...... 43
Figure 25 Dilatometry set up in the GleebleTM 3800 ...... 44
Figure 26 Light radiation furnace ...... 45
Figure 27 Tempering sample cross section and surface hardness indents ...... 50
Figure 28 Dilatometry curve showing A1 and A3 transformations in Alloy 1.1 ...... 53
xv Figure 29 Predicted A1 temperature vs. actual A1 temperature measured through dilatometry 54
Figure 30 SS-DTA curve for Alloy 4.1 heat treated below the predicted A1 temperature ...... 55
Figure 31 SS-DTA curve for alloy 4.1 heat treated above the predicted A1 temperature ...... 55
Figure 32 a) Microstructure of Alloy 4.1 cast and b) Microstructure of Alloy 4.2 cast used to quantify delta ferrite ...... 57
Figure 33 Actual and predicted retained delta ferrite content ...... 57
Figure 34 Impact toughness results for Alloy 4.2 test welds after 1 hour PWHT ...... 59
Figure 35 Impact toughness results for Alloy 1.2 at 2 hours PWHT ...... 61
Figure 36 Alloy 4.2 welded on Base Metal B hardness traverses (1kg load) taken in the cap passes and root pass of the weld in the as-welded condition ...... 62
Figure 37 Alloy 4.2 test weld hardness in the four PWHT conditions. Base Metal A welds 1hr
PWHT duration. Base Metal B welds 2hr PWHT duration...... 64
Figure 38 Alloy 1.2 test weld hardness data in the four PWHT conditions. PWHT duration of 2 hours ...... 65
Figure 39 a) Alloy 1.2 welded on Base Metal B 760oC PWHT bulk weld area low mag b) Alloy
4.2 welded on Base Metal B 760oC PWHT bulk weld area low mag c) Alloy 1.2 welded on Base
Metal B 760oC PWHT high mag d) Alloy 4.2 welded on Base Metal B 760oC PWHT bulk weld area high mag ...... 66
Figure 40 a) Alloy 1.2 welded on Base Metal B 760oC PWHT cap pass b) Alloy 4.2 welded on
Base Metal B 760oC PWHT cap pass ...... 67
Figure 41 Alloy 4.2 welded on Base Metal B 760oC PWHT delta ferrite growth in HAZ from base metal stringers ...... 68
xvi Figure 42 Alloy 4.2 on Base Metal B HAZ with delta ferrite growth in the a) CGHAZ and b) delta ferrite growth in stringers in the HAZ ...... 68
Figure 43 a) HAZ in Base Metal A at 770oC PWHT and b) HAZ of Base Metal B at 760oC PWHT
...... 69
Figure 44 Base Metal A 1 hr 800oC PWHT microstructure at a) low mag and b) high mag ...... 70
Figure 45 Base Metal A 2 hr 800oC PWHT microstructure at a) low mag and b) high mag ...... 70
Figure 46 Base Metal B 2 hr PWHT microstructure at a) 760oCPWHT and b) 800oC ...... 71
Figure 47 Alloy 4.2 tempering data at 760oC and 807oC tempering temperature. Red data point indicates samples tempered at 820oC for 30 minutes. As welded hardness taken as average from
Alloy 4.2 on Base Metal B test weld ...... 73
Figure 48 Microstructure of Alloy 4.2 tempered at 807oC for a.) 1 minute b.) 5 minutes c.) 2 hours d.) 4 hours ...... 74
Figure 49 Base Metal B tempering data at 766oC and 800oC tempering temperature ...... 75
Figure 50 SS-DTA Curves for Base Metal B at various heat treatments at 800oC ...... 77
Figure 51 Base Metal B microstructures a) as received b) 1 minute temper c) 30 minute temper d) 1 hour temper e) 2 hour temper f) 4 hour temper ...... 78
Figure 52 Range of R2 values with varying C values. Red data point indicates C-value with best fit ...... 80
Figure 53 Holloman-Jaffe relation for Alloy 4.2 ...... 81
Figure 54 Range of R2 values with varying C values for Base Metal B. Red data point indicates
C-value with best fit ...... 81
Figure 55 Holloman-Jaffe relation for Base Metal B ...... 82
xvii Figure 56 Alloy 3 dilatometry curve ...... 90
Figure 57 Alloy 4.1 dilatometry Curve ...... 91
Figure 58 Alloy 4.2 dilatometry curve ...... 91
Figure 59 Base Metal A dilatometry curve ...... 92
Figure 60 Base Metal B dilatometry curve ...... 92
Figure 61 Alloy 1.2 heat treated a) 10oC below 796oC and b) 10oC above 796oC ...... 94
Figure 62 Alloy 4.2 heat treated a) 10oC below 817oC and b) 10oC above 817oC ...... 94
Figure 63 Base Metal A heat treated a) 10oC below 786oC and b) 10oC above 786oC ...... 95
Figure 64 Base Metal B heat treated a) 10oC below 776oC and b) 10oC above 776oC ...... 95
Figure 65 Alloy 1.1 example cast microstructure. Predicted retained delta ferrite: 1.5 vol%.
Delta ferrite in this image 0 vol%. Experimental average 1.48 vol% ...... 96
Figure 66 Alloy 1.2 example cast microstructure. Predicted retained delta ferrite: 2.3 vol%.
Delta ferrite in this image 4.5 vol%. Experimental average 4.32 vol% ...... 97
Figure 67 Alloy 3 example cast microstructure. Predicted retained delta ferrite: 1.62 vol%.
Delta ferrite in this image 0 vol%. Experimental average 0.5 vol% ...... 97
xviii Chapter 1: Introduction
410 martensitic stainless steel was selected for use in some downstream hydro-processing installations, due to its increased hardenability, good resistance to sulfide corrosion and chloride stress corrosion cracking, and relatively low cost over austenitic stainless steels. Industry experience with Type 410 steel welds, using generic welding consumables, has shown difficulties in meeting the ASME B31.3 and NACE MR0103 hardness and toughness requirements in the weld metal and HAZ. Recent research has pointed out the wide composition specifications of Type 410 base metal and welding consumables as the leading cause for significant hardness and toughness variations, related to exceeding the A1 temperature during
PWHT and formation of fresh martensite, and to retention of significant amounts of delta ferrite both of which reduce material toughness in Type 410 stainless steel. With high levels of ferrite stabilizing elements excessive delta ferrite will be formed, but if the ratio of ferrite to austenite to stabilizing element is decreased the A1 temperature can be reduced to temperature below the
ASME B31.3 PWHT range. In addition to compositional influence on retaining delta ferrite, fast cooling rates have been proven to retain delta ferrite due to leaving insufficient time for ferrite to transform into austenite.
In order to combat this issue a two stage study was completed by David Stone, which included creating predictive formulas that predict A1 temperature, and retained delta ferrite for Type 410 stainless steel, and forming a custom composition range that would ensure the A1 temperature 1 would be above 800oC, and would keep retained delta ferrite below 20 vol%. A DOE approach was used to model the effect of composition on A1 temperature and delta ferrite stability.
ThermocalcTM and DICTRA were used to model the A1 temperature, and retained delta ferrite using thermodynamic and kinetic models. The relative effects of each element on the A1 temperature and retained delta ferrite were identified, as well as the relative effect on cooling rate on delta ferrite retention. The results from these models were used to create regression equations that were able to predict the A1 temperature and retained delta ferrite content. Using these formulas over 2.5 million theoretical compositions were used to identify a composition range that would ensure a favorable A1 temperature and level of retained delta ferrite.
The goal of this study was to validate the accuracy of the predictive formulas, to test weld metals made within the custom composition ranges developed by Stone, and to develop a tempering response parameter for these alloys. Five metal cored weld consumables were made in compliance with three of the custom composition ranges, and two commercial Type 410 steel plates were provided for this research. The A1 temperature predictive formula was validated through dilatometry, and single sensor differential thermal analysis (SS-DTA). The delta ferrite formula was validated through measuring retained delta ferrite using a point count method on cast samples whose thermal histories were measured in order to obtain experimental cooling rates close to that which is experienced during welding.
The custom Type 410 steel consumables were then used to produce test welds in two commercial base materials in order to characterize the hardness and impact toughness after PWHT. PWHT temperatures were selected using predicted A1 temperature values in addition to 800oC, the upper range of the ASME B31.3 PWHT range, to look for adverse effects when heat treatments
2 are run above the material A1 temperature. The test welds showed good hardness and toughness results indicating the optimized composition ranges were capable of producing good welding consumables.
A tempering response study was used to create a Holloman-Jaffe Parameter (HJP) that could predict the hardness of the material based on tempering temperature and time. The results of this study also showed the potential negative effects that could be experienced when heat treating at the upper and lower limits of the ASME PWHT range, indicating the usefulness of being able to heat treat a material based on the predicted A1 temperature.
3 Chapter 2: Literature Review
2.1 Martensitic Stainless Steel
2.1.1 Characteristics of Stainless Steel Alloys
Stainless steels are a group of high alloyed steels that are characterized on their stainless feature that comes from their Chromium content. In order for these steels to be considered stainless they generally need to contain at least 10.5wt% Cr and generally do not contain more than 30wt% Cr.
Cr forms a passive surface oxide, (Fe,Cr)2O3, which forms on the steel’s surface that prevents oxidation. Chromium’s presence in the steel at levels of 10.5wt% or higher required to be considered stainless increases the stability of the Cr2O3 oxide due to Cr having a higher affinity to oxygen than iron, making it less likely for iron the oxide to form. In more aggressive environments higher levels of Cr will be needed in order to form a protective Cr2O3 layer [1],
[2]. This protective layer also provides good high temperature oxidation resistance, and resistance to corrosion that low alloy steels are generally susceptible to. Stainless steels tend to be classified based on their microstructures: austenitic, ferritic, duplex which is a combination of austenite and ferrite, and martensitic. Additionally precipitation-hardenable (PH) martensitic and austenitic stainless steels are separately named as PH stainless steels. In order to acquire desired mechanical properties in stainless steels, control of the microstructure is crucial. The iron- chromium, iron-chromium-nickel, and iron-chromium carbon systems are used to help control
4 microstructure, and are often used with an added alloying element to better characterize the microstructure.
2.1.2 Martensitic Stainless Steel
Martensitic stainless steels are based on the Fe-Cr-C ternary system. They undergo the allotropic transformation to form martensite from austenite except if very slow cooling rates occur; for example, furnace cooling can prevent the formation of martensite on cooling. Compositional control of these alloys is necessary in order to limit ferrite and austenite retention in order to keep the desired mechanical properties. Martensitic stainless steels generally contain 11.5-18 wt% Cr with carbon content going as high as 1.2 wt% which are classified as group 3 martensitic stainless steels due to their higher susceptibility to hydrogen cracking [1], [3]. Allowable alloying content in martensitics is limited when compared to other classifications of stainless steels in order to avoid retained austenite while also limiting retained ferrite [3]. Welding using matching filler and base material is also common practice in welding martensitic stainless steel in order to have weld strength that matches the base material. Austenitic filler metals have improved corrosion resistance when they form dual phase austenite + ferrite microstructure that will be much softer than the base material when welded on martensitic stainless steels. Heat treatments of welds that are comprised of austenitic stainless steel welded onto martensitic stainless steel may result in brittle sigma phase formation due to sigma phase being able to form in an austenitic microstructure in the tempering phase range. Martensitic stainless steel’s lower cost, which is attributed to lower alloy inclusions, are also a motivation for avoiding dissimilar metal welding in these alloys [1], [4].
5 Martensitic stainless steels originally started off as a metal used for cutlery, but found new applications due to their good strength to weight ratio, good corrosion resistance, and low cost
[5]. Since then martensitic stainless steels have been used in steam, gas, and jet engine turbine blades, steam piping, large hydro-turbines, freshwater canal locks, and cladding for casting rolls
[1]. They have also been found to be a favorable material in the oil and petrochemical industry due to their resistance to transgranular stress corrosion cracking by chlorides, and resistance to attack by sulphuric compounds at elevated temperatures. In the oil industry they are often found in hydrocrackers, hydro-treaters, and distilling units for crude oil [4].
2.1.3 Type 410 Martensitic Stainless Steel
Like most martensitic stainless steels Type 410 requires PWHT, and are generally welded with matching consumables. Along with increasing fracture toughness and ductility PWHT’s are valuable in decreasing the risk of hydrogen induced cracking, which is also avoided by employing a 200-300o preheat and inter pass temperature [5]. Composition limits set by AWS and ASTM for Type 410 steels can be seen in Table 1. ASTM A240 specifies base metal where
AWS A5.4 and 5.9 specify composition limits for consumables used for SMAW, and bare electrodes and rods typically used for GTAW correspondingly. Other than Cr, alloying elements are kept at or under 1 wt%. Mo, N, and Cu do not have specified limits in ASTM, but only N for
AWS. Relative effects of alloying elements on phase transformations of Type 410 steel will be discussed later on. Mechanical property requirements for Type 410 base material, and weld consumables are outlined in Table 2. AWS only specifies mechanical properties for E410 with tension tests of an all weld metal section machined from a groove weld that need to meet properties outlined in AWS B4.0 or B4.0M. ER410 does not specify mechanical testing
6 standards, only the composition range is defined. ASTM A240 defines max hardness in Brinell and Rockwell B, in addition to tensile and yield strength and elongation requirements.
Requirements set for Charpy V-notch tests are set by the purchaser of the base material.
Table 1: Composition standards for Type 410 Stainless Steel
Table 2: Mechanical standards for Type 410 steel for base metal and shielded metal arc welding electrodes Tensile Elongation Hardness Strength, min Yield Strength Min Percent Classification ksi MPa Brinell Rockwell B ksi MPa AWS A5.4 75 450 - - 20% - - E410XX ASTM A240- 65 450 30 205 20% 217 96 S41000
2.2 Phase Transformations
2.2.1 Solidification
Stainless steels have four modes they can initially solidify as depending on composition as seen in Figure 1. Type 410 stainless steels primarily solidify as ferrite, while sometimes forming
7
Figure 1 Potential solidification modes of stainless steel [6]
some austenite before full solidification. Knowing the solidification path of the material is important for obtaining desirable properties [6].
Lippold outlined three main transformation paths for martensitic stainless steel, which is heavily dependent on composition [1]. The first path, Equation 1, is a common path for Type 410 steel and was classified for martensitic stainless steels with 11-14wt% Cr, and 0.1-0.25wt% C. In this transformation path the metal solidifies as ferrite, and enters a fully ferrite phase field, then proceeds to go through a solid state transformation of ferrite to austenite. The material will then form a fully austenitic microstructure, then transform into martensite on further cooling.
Another common transformation path for Type 410 steel, Equation 2, gives a scenario where ferrite does not fully transform into austenite, and is retained through austenite’s transformation to martensite. The last transformation, path, Equation 3, considers austenite and eutectic ferrite that is enriched with Cr and Mo forming with primary ferrite. This eutectic ferrite is thought to 8 reside along solidification grain boundaries, and subgrain boundaries. This austenite and ferrite mixture stays present through the primary ferrite’s transformation to austenite, then the eutectic ferrite is retained as the austenite transforms into martensite [1].
(1) 퐿 → 퐿 + 퐹 → 퐹 → 퐹 + 퐴 → 퐴 → 푀
(2) 퐿 → 퐿 + 퐹 → 퐹 → 퐹 + 퐴 → 푀 + 퐹
(3) 퐿 → 퐿 + 퐹 + (퐴 + 퐹푒 ) → 퐹 + 퐴 + 퐹푒 → 퐴 + 퐹푒 → 푀 + 퐹푒
2.2.2 Austenite Martensite and Ferrite Phase Transformations
The final microstructure of type 410 steels depend mostly on composition, and also on cooling rate. Balancing austenite and ferrite stabilizers is key to having the desired microstructure for martensitic stainless steels [1], [7]. After the initial delta ferrite solidification, two important phase fields Type 410 steels may enter are the γ-loop, austenite phase field, and the γ+δ-ferrite phase field seen in Figure 2. Fe-Cr-C pseudo-binary phase diagrams are important indicators of the presence of ferrite in the final microstructure. Figure 2 shows the effect of increasing carbon content on the γ-loop, which increases the temperature range of the loop and the stability of austenite with increasing Cr. The γ-loop starts to narrow around 10 wt% Cr, and generally closes in the range of 12 wt% Cr alloys. Generally 410 steels will cool through the δ+γ phase field, but will still solidify as a fully martensitic microstructure with an Mf temperature above room temperatures, and with limited ferrite promoters [1], [8].
With sufficient ferrite stabilizing elements, delta ferrite will be present in the final microstructure. Where eutectic ferrite is formed with excessive segregation of ferrite stabilizers
9
Figure 2 Phase stability in the Fe-Cr pseudo-binary phase diagram with varying carbon content [1]
to the grain boundary, delta ferrite will stay stable with a larger ferrite phase field due to a smaller γ-loop [8]. A good example of this can be seen through a study completed by Castro and
Tricot which showed an increase in the ferrite phase field with an increase in Cr, the most prominent ferrite promoter in 410 stainless steels [1]. Figure 3 shows a comparison of a 410 stainless steel with 13 wt% Cr, and a 17 wt% Cr ferritic stainless steel. At Cr concentrations below 12 wt% a full austenite transformation should occur on cooling resulting in a fully martensitic structure, but other elements will have an effect on this transformation. Being able to accurately predict delta ferrite retention, which cannot happen with pseudo-binary phase diagrams, becomes an important task.
Four important temperatures related to phase transformations that occur in 410 stainless steel are the A1 and A3 temperatures, which denote the temperatures that the austenite-ferrite phase transformation begins and ends, and the Ms, and Mf temperatures which indicate the martensite start and finish temperature. These transformation temperatures are dependent on material
10
Figure 3 Chromium’s effect on stabilizing the ferrite at a) 13% Cr and b) 17%Cr [1]
composition, and thermal cycles. The A1 and A3 temperatures are primarily diffusional transformations so in addition to the compositional effects on these temperatures, diffusion kinetics will also have an effect.
The A1 and A3 temperatures are known to increase with increasing heating rates due to the slow dissolution of carbides. In multiple studies high heating rates have shown to cause a diffusion- less shear transformation from martensite to austenite [9]–[11]. In the study by Apple et al heating rates of from 3oC/s to 28,000oC/s were used to test for a shear transformation to austenite vs. a thermally induced transformation to austenite in an Fe-Ni-C steel. The massive transformation associated with the shear transformation resulted in an austenite microstructure that obtained high dislocation densities and surface tilt characteristics associated with shear transformations. The thermally activated transformation to austenite when exceeding the A1 temperate at slower heating rates was attributed to short range diffusion leading to a lower dislocation density [9]. In the work completed by Lee et al on a Fe-3Si-13Cr-7Ni steel the A1 and A3 temperature was analyzed at varying heating rates. Figure 4 from this study shows that
11 once a heating rate of 10oC/s (10 K s-1) was exceeded the increase in the A1 temperature started to plateau showing a change from a diffusional transformation to a shear transformation, and no further increase of the A1 or A3 temperature was seen at heating rates above 50oC/s [10].
Figure 4 Variation in A1 and A3 temperatures with increasing heating rates [10]
The Ms transformation temperature, though diffusion less, shows variation based on austenitizing temperature and cooling rate [7], [8], [12]. This has been shown in a study on AISI
410 stainless steel completed by Tsai et al which was completed to better understand martensite transformation, and austenite retention in this material. An obvious decrease in the Ms temperature with increasing cooling rates, from 0.5oC/s to 10oC/s, was observed in dilatometry curves. These tests were run with sections of an AISI 410 steel bar, and it was observed that the samples with a higher austenitization temperature had lower Ms temperatures when run at the same cooling rate as different samples [12]. Figure 5 shows the effects of varying cooling rates and austenitization on Ms temperature. A suggested mechanism on the effect of a higher austenitization temperature on the Ms temperature is a high concentration of vacancies without 12 the precipitation of chromium carbides leads to an increased strength of austenite through dislocations and by the solid solution from carbon and chromium [12].
Figure 5 Decreasing Ms temperature with decreasing cooling rate and varying austenitization temperatures [12]
2.2.3 Delta Ferrite
Being able to control delta ferrite retention in 410 stainless steels is important due to excessive delta ferrite causing a degradation of toughness properties, but having small amounts can increase resistance to stress corrosion cracking, as well as increase yield, and tensile strength in some alloys [13], [14]. The cause of retained delta ferrite and its effect on mechanical properties has been studied for different classes of steel due to its unpredictable nature. There are two competing theories on retained delta ferrite’s influence on the reduction of toughness: the first theory suggests that excessive carbide growth along the grain boundaries between delta ferrite and martensite causes decohesion reducing impact toughness, the opposing theory suggests a lack of solid solution in the ferrite grains resulting in a preferred fracture path.
13 A study by Bhambri in 1986, and by Wang in 2010 showed delta ferrite to be responsible for degradation of impact properties in 13Cr steels with primarily martensitic structures. Bhambri’s work initially saw preferred intergranular fracture due to impurities at grain boundaries. These impurities were avoided through fast cooling rates during reversal treatments which lead to a grain boundary network of delta ferrite along. This still resulted in intergranular fracture in a primarily martensitic structure where transgranular fracture was usually the preferred fracture path. He believed this to be due to the ferrite being a brittle phase in comparison to tempered martensite [15]. Wang’s work led to conclusions that retained delta ferrite in the microstructure of primarily martensitic steels lead to a decrease of crack initiation and propagation energy in these steels while in the ductile to brittle transition temperature (DBTT) range, while also increasing this temperature range. This was attributed to easy ductile crack growth in the ferrite attributed to lower carbon levels in the ferrite the allowed for greater dislocation motion, which in turn propagated into the martensite causing a brittle failure mode. Despite this effect the upper and lower shelf energy of this material was seen to be the same between materials with and without retained delta ferrite, and was found that this degradation of impact energies was unrelated to carbides forming along the ferrite-martensite interface [16].
Opposing theories seen in three other studies showed that reduction in toughness in dual phase delta ferrite and martensite steels was could not only be attributed to ferrite presence, but a combination of retained delta ferrite and particles forming at the interface, as well as prior austenite grain size [17]–[20]. Carrouge found that the DBTT was strongly influenced by ferrite content in super martensitic stainless steels, but this did not result in a loss of the upper and lower shelf energies when compared to alloys containing less ferrite. Instead loss of impact
14 toughness in the HAZ of these materials was attributed to coarser grain size, which limited boundaries to crack propagation [17]. Large delta ferrite grains in the HAZ of SMA welds of
12Cr steel was observed to have a detrimental effect on impact toughness, so it was recommended to reduce heat input to keep grain size down [21]. Another study in dual phase ferrite-martensite steels excessive brittle carbides was the cause for low impact energies. This unfavorable carbide growth was attributed to service temperatures and heat treatments that led to the nucleation of coarse carbides in large random groupings [18]. Ultimately appropriate Schafer found that a varying mixtures of delta ferrite and dendritic carbides (M23C6) in the final microstructure results in either phase being the predominant reason for a reduction in toughness leading to the combating theories on toughness reduction when there is delta ferrite present.
It was also observed in this study that a pure delta ferrite formation increased ductility and toughness, but both the ferrite and dendritic carbides had a negative impact on strength [19]. In order to avoid unfavorable amounts of retained delta ferrite it is important to understand how welding effects the growth and retention of delta ferrite in martensitic stainless steel welds. The
Balmforth diagram was developed for predicting retained delta ferrite in martensitic stainless steel welds for the purpose of avoiding excessive ferrite [22]. This diagram was shown to be a good predictor of ferrite content using a range of heat inputs while melting samples with a
GTAW torch, but in some studies inhomogeneous microstructures have been observed after welding. Along with composition heat input and in turn cooling rate of the material have bene shown to have a strong impact on delta ferrite retention, but there are different findings on whether high or low heat inputs retain more delta ferrite. In a study on P91 and P92 steels it was
15
Figure 6 Charpy impact energy curves on super martensitic stainless steel wit h tempered martensite, fully martensitic, and 14% retained delta ferrite microstructures [17]
found that higher concentrations of delta ferrite were retained with higher heat inputs, and fast cooling rates [23] where fast cooling rates were generally associated with lower heat inputs in several other studies [14], [24]–[26].
In a study on 9-12Cr steels multi pass bead on plate welds were made to observe the effect of multiple weld passes on microstructure. Delta ferrite retention in these welds started off as polygonal in morphology, but after reheating from subsequent weld passes was retained along prior austenite grain boundaries seen in Figure 7. Without being reheated large polygonal delta ferrite grains were left in the final pass of the weld [25]. This study found that keeping the chromium equivalent below 13.5, was necessary to keep the microstructure completely martensitic. A study on reduced activation stainless steels also saw unexpected large grain delta ferrite growing in single pass bead on plate welds that grew from the weld interface attributed to
16 fast cooling rates and epitaxial growth of delta ferrite from partially melted ferrite grains in the
HAZ [14].
Delta Ferrite
Martensite
Figure 7 Microstructures seen in a multi-pass bead on plate weld from a) the center of the weld and b) from the last pass of the weld [25]
2.3 Effects of Alloying Elements
2.3.1 Carbon
Carbon is present in all steels, and is the key component to the strengthening mechanism in martensitic steels through interstitial strengthening. This effect is enhanced at higher temperatures due to carbon’s grater solubility in the matrix. Higher carbon content has the potential to reduce corrosion resistance due to its affinity to form carbides with chromium.
Common chromium rich M23C6 carbides contain roughly 16 times as much chromium as carbon leading to the reduction in corrosion resistance due to leaving that much chromium out of solution. Carbon is a very potent austenite stabilizers in steels. When sufficient carbon is
17 present and is supersaturated in the FCC crystal lattice it will be distorted into an HCP structure where BCC is usually stable [1].
2.3.2 Chromium
Chromium’s primary function in stainless steels is to provide corrosion resistance, and shows its greatest resistance in oxidizing environments such as nitric acid. As mentioned above the addition of 10.5 wt% Cr is what is required to consider a steel “stainless” which comes from the
(Fe,Cr)2O3 passive oxide it forms on the surface of these materials. In addition to the corrosion resistance, Cr is a strong ferrite promoter which is represented in the Creq equations used in constitution diagrams. Cr promotes ferrite stabilization in all stainless steel grades, so despite increasing corrosion resistance at higher levels, a higher retention of ferrite in the final microstructure of austenitic, duplex, and martensitic stainless steels may lead to unfavorable properties. Cr also is important for carbide formation in stainless steels [1], [5]. M23C6 carbides are a key component in tempering martensite to achieve desired mechanical properties. These will be discussed further below.
2.3.3 Copper
Copper is a weak austenite promoter in stainless steels, and is used as a precipitation hardening element in precipitation-hardenable martensitic stainless steels. It has been found to reduce grain size, and reduce growth of revered austenite in super martensitic stainless steel. In 15Cr steels copper was found to improve the machinability tested up to 1.9 wt% Cu, and maintains adequate corrosion resistance. The levels of Cu found in commercial 410 steels is generally below 0.1 wt% Cu which alleviates concern of unfavorable secondary phases forming.
18 2.3.4 Manganese
Manganese is generally considered an austenite promoting element, but can be dependent on the amount of manganese present, and the amount of nickel present. At lower temperatures is effective at stabilizing austenite, but in 304 steels is seems to have little effect in stabilizing austenite at higher temperatures [1]. F. C. Hull’s study showed manganese acted as an austenite promoting element below 6 wt%, but as ferrite promoter above these levels [13]. Manganese is present in virtually every steel, and in martensitic and ferritic steels is generally present below 1 wt%. Manganese also acts as a deoxidizer, and prevents hot shortness which is a form of solidification cracking that occurs as a result of iron sulfide which is a low melting eutectic.
2.3.5 Molybdenum
Molybdenum is a strong ferrite stabilizer, and provides additional corrosion resistance in stainless steels. It is effective at scavenging embrittling impurities, but this effect could be lost to carbide formation during tempering or high temperature service [27]. In a study on PWHT for
13Cr-Ni steels it was found that the reduction of silicon and carbon with added molybdenum up to 0.25wt% helped retard M23C6 carbides from the grain boundaries which improved the impact toughness [20].
2.3.6 Nitrogen
Nitrogen, which is generally in solution as an impurity in stainless steels, is the other strongest austenite promoter next to carbon. It is sometimes added to austenitic stainless steel in order to act as solid solution strengthener. Nitrogen levels are something that need to be monitored in martensitic stainless steels due to its ability to stabilize austenite, and its effect on lowering the
19 A1 temperature. Type 410 stainless steel does not have limits on nitrogen in place, which could lead to unfavorably low A1 temperatures based on PWHT temperatures [1].
2.3.7 Nickel
Nickels primary function in stainless steels is to retain austenite in the final microstructure. In martensitic stainless steels nickel content is general kept below 1 wt% in order to avoid retained austenite, as well as nickels susceptibility to stress corrosion cracking. The Copson curve is a useful tool for showing a reduction of SCC resistance in aggressive Cl containing environments.
Ni is known to increase toughness and ductility in martensitic stainless steels, and also increases strength through solid solution strengthening [1].
2.3.8 Silicon
Silicon is present in stainless steels as a deoxidizer in the range of 0.3 to 0.6 wt%. The low amounts of silicon in stainless steels is due to a range of iron sillicides it can create embrittling the microstructure, and it tendency to form a low melting eutectic with nickel that can lead to solidification cracking. One additional benefit of silicon is that it can improve the fluidity of molten steel, so it may be added to weld consumables to a higher degree [1].
2.4 Predictive Formulae and Composition Range Development for Type 410 Steel
2.4.1 Background
A study on Type 410 steels was started in order to avoid A1 temperatures that are below the
ASME B31.3 PWHT range, and in attempt to keep retained ferrite to a limit that does not negatively affect the toughness properties, which can occur due to a wide composition range in
ASTM and AWS specifications, Table 1. Work was started in 2015 by David Stone to formulate a way to predict A1 temperature, and retained DF in 410 stainless steels, and ultimately finding 20 composition limits that would create Type 410 weld consumables, and bulk metal that is guaranteed to meet desired A1 temperatures and retained DF levels [28].
2.4.2 A1 Temperature Predictive Formula
A design of experiment (DOE) methodology was used to run 82 theoretical compositions that were generated within Type 410 Steel composition range through ThermoCalcTM to generate phase volume fraction vs temperature diagrams in order to acquire A1 and A3 temperatures
[28]–[30]. From these runs it was found that the A1 temperature had a range of 203oC from 686-
889oC which is compared to the ASME PWHT range in Figure 8. The data was input into
Minitab software to statistically analyze the compositional effects on the A1 temperature.
Section 2.3 above outlines the effect of elements on A1 temperature,
Figure 8 Possible A1 Temperature vs ASME B31.3 recommended PWHT with compositions within the ASTM A240 and AWS 5.4 and 5.9 requirements [28]
21 and delta ferrite retention, but the work completed by Stone went into detail on which elements had the largest effects on suppressing and increasing the A1 temperature [28]–[30]. Three missing elements from ASTM A240, Mo, N, and Cu, were shown to have significant effects on suppressing the A1 temperature, N and Cu, and increasing the A1 temperature, Mo. Figure 9 outlines the relative effects on each element on increasing or decreasing the A1 temperature.
Figure 9 Relative Alloy effect on A1 temperature [28]
The analysis of the composition, and calculated A1 temperatures in Minitab was also used to create an equation predicting the A1 temperature based on a given composition. Equation 4 was generated through a multiple regression analysis using a surface response design including all significant reactions (P=0.05). The R2 value of this equation was 99.4% indicating a near perfect fit. An additional equation, Equation 5, was generated without interactions, and was found to have a lower prediction accuracy, but was found to have an R2 value of 97.8% which still shows 22 a great fit to the A1 temperatures calculated in ThermoCalcTM. By eliminating the interaction terms multicollinearity within the model was eliminated therefore indicating the actual magnitude of the elements effect on A1 temperature [28]–[30].
퐴1 = 712.54 + 211.1퐶 + 10.83퐶푟 − 15.77푁푖 + 14.73푀표 + 19.07푆푖 − 55.06푀푛 −
5.26퐶푢 − 281.0푁 + 10918푁2 − 22.71퐶×퐶푟 + 35.3퐶×푁푖 − 97.7퐶×푀표 −
3.87퐶푟×푁푖 + 1.97퐶푟×푀표 − 2.55퐶푟×퐶푢 − 43.8퐶푟×푁 − 12.12푁푖×푀표 − 10.20푁푖×푀푛 −
11.21푁푖×퐶푢 + 189.2푁푖×푁 + 3.50푀표×푆푖 − 9.54푀표×푀푛 −
130.3푀표×푁 − 5.0푆푖×퐶푢 − 12.19푀푛×퐶푢 + 81.6푀푛×푁 + 201.1퐶푢×푁
Equation 4 Full developed model for predicting the equilibrium A1 temperature within the composition range of Type 410 steel [28]
A1=772.66+6.5Cr+20.91Mo+18.5Si -90.5C-70.4Ni-65.4Mn-45.3Cu-242.6N
Equation 5 Full developed model for predicting the equilibrium A1 Temperature [28]
2.4.3 Delta Ferrite Retention Predictive Formula
The predictive formula for retained delta ferrite volume percent used ThermoCalcTM to measure the elements relative effect on the ferrite stability range. For this experimental run 17 theoretical compositions were used in a ¼ fractional DOE with the same 8 elements used in the A1 temperature study. The impact of alloying elements on ferrite stability was determines by integrating the stability range over these 17 compositions. Looking the pareto chart in Figure 10 it was shown that C, Cr, Mo, Ni, and Cu were the elements with a significant effect on the ferrite
23 stability range. From this study a regression equation, Equation 6, predicting the ferrite stability range with an R2 value of 96% indicating a good fit [28], [31].
Figure 10 Pareto chart showing the elements with a signifi cant effect of delta ferrite retention [28]
퐹푆푅 = −1 + 4퐶푟 + 7.9푀표 − 174.7퐶 − 9.12푁푖 − 11.34퐶푢
Equation 6 Regression equation describing the ferrite stability range area (FSR) , the integrated area under the delta ferrite curve [28]
A computational model of ferrite to austenite transformation, which occurs in the 1400-680oC range, in Type 410 steel was applied to the seventeen theoretical compositions with the three cooling rates of 1, 15, and 50oC/sec in DICTRA. These cooling rates were selected based on the
o o typical cooling rates for casting for a 1 C/sec cooling rate, the t8/5 15 C/sec cooling rate in button 24 melting cooling times, and the 50oC/s cooling rate found from thermal data recorded during the casting of 6g Type 410 steel samples in an induction levitation melting device. The DICTRA simulations covered the entire austenitic phase field from initial nucleation temperature to 680oC which was below 686oC, the lowest A1 temperature found in the study performed for Equations
4 and 5. From this experiment it was found that there was an insufficient driving force for alpha ferrite formation before the martensite transformation, which was backed up with a TTT diagram that showed alpha ferrite formation required a holding temperature above Ms for long times unrealistic to welding and post processing conditions [28], [31].
In order to get maximum simulation stability in DICTRA only phases found in equilibrium were included in the model, FCC and BCC. Since it becomes stable at lower temperatures, M23C6 was not included in the model which would result in an instability of the model. From the 17 theoretical compositions the predicted volume fraction of delta ferrite was calculated at the three cooling rates above. A cooling rate of 1oC/sec showed a volume fraction range of 0-77% delta ferrite, 15oC/sec had a range of 0-78% delta ferrite, and 50oC/sec had a range 0-92% delta ferrite, showing an increase of volume fraction of retained delta ferrite with increasing cooling rate [28],
[31]. Figure 11 shows the relative effects of elements and cooling rate on retained delta ferrite.
In some of the theoretical alloys it was found that the retained delta ferrite was lower at higher cooling rates. This was attributed to an FA solidification mode from liquid as opposed to an F solidification mode before entering the ferrite to austenite transformation temperature range.
From this information it was deduced that faster cooling rates could suppress ferrite retention given that there is austenite formed during solidification.
25
Figure 11 Relative effects of alloying elements on retained delta ferrite from DICTR A simulations [28]
In order to take solidification mode into account, elements with a potential FA solidification mode when simulated at 50oC/sec were taken out of the simulation and a regression was run on these data points for calculating the retained delta ferrite. With this data set Equation 7 was generated with an R2 value of 84% showing a good fit [28], [31].
훿 − 퐹푒푟푟푖푡푒 = −3.483 − 0.1578 푁푖 + 0.3234 퐶푟 + 16.63 퐶 + 0.221 푀표
+ 0.001590 [퐶표표푙푖푛푔 푟푎푡푒 (°퐶/푠)] – 1.601 퐶푟×퐶
Equation 7 Delta ferrite prediction model including a cooling rate parameter [28]
2.4.4 Optimized Composition Window
A study was run on finding an optimized composition window for Type 410 steel using the A1 temperature, and retained ferrite study in order to find a composition range that meets the ASTM and AWS composition range. This optimized composition was made in an effort to be able to 26 make A1 temperature above 800oC, and retained delta ferrite below 20 vol% by controlling the composition. This was completed by running 2,562,890,625 composition combinations with each element iterating through 10 different levels. Each element from this simulation was then graphed and analyzed to select a 20% bin that represented the best probability of having an appropriate A1 temperature and ferrite retention. This process was repeated by iterating through composition windows that fall within the newly made composition range until an optimized window is developed. Using this methodology the composition range was reduced from the
ASTM and AWS composition ranges to composition ranges that had near a 100% chance of achieving an A1 above 800oC and retained ferrite below 20%. Figure 12 contains the four composition ranges generated in this study [28].
Figure 12 Custom Type 410 steel composition ranges [28]
27 2.5 Experimental A1 Transformation Analysis
2.5.1 Dilatometry
Being able to detect phase transformations in stainless steels is important for many reasons stated above. Dilatometry is commonly used for in-situ monitoring of bulk phase transformations in stainless steel on heating and on cooling. Specific volume changes that occur during phase transformations due to a changing lattice structure are able to be detected through measured changes in length of the sample [32], [33]. These volumetric changes are easily detectable due to the mostly linear increase in expansion of materials attributed to the coefficient of thermal expansion. The change from martensite to austenite results in a negative change in volume on heating past the A1 temperature, where the transition from austenite to martensite at the Ms
Figure 13 Dilatometry curve showing austenite and martensite transformations in Fe -12Cr stainless steel [34]
28 results in a positive change in strain shown in a Fe-12Cr stainless steel in Figure 13. Quantifying the phase transformation temperatures can be done by plotting dilation against temperature. The
A1 temperature is quantified by marking the initial drop of dilation in the processed curve. It has been found that the A3 temperature is not as easy to identify. This was shown in a study by F.
Christien et al martensite was still present at temperatures higher than the A3 temperature measured using dilatometry. By using neutron diffraction and magnetization the A3 temperature was more accurately measured, and it was recommended that using dilatometry to measure the
A3 temperature should only be used as an approximation [35]. A study on the austenite to martensite transformation used dilatometry to effectively measure the A1 and Ms temperatures using a heating rate of 20oC/min and a controlled cooling rate of 10oC/min [34].
2.5.2 Single Sensor Differential Thermal Analysis (SS-DTA)
In 2005 a paper was published on a new methodology for in-situ detection of phase transformations during welding that was developed by Alexandrov, and Lippold called single sensor differential phase analysis. This was accomplished by recording thermal histories at fast sampling rates on the order of 500 to 1000 Hz in order to detect thermal effects caused by phase transformations. In order to accurately detect these phase transformation a software developed at
Ohio State compares the experimental thermal history to a simulated reference curve generated with a Rosenthal based equation to compare to the experimentally obtained cooling curve.
Deviations in temperature between the experimental thermal history and the reference curve, denoted ΔT, are caused by thermal effects of absorption or release of latent heat, and can be used to identify phase transformations depicted in Figure 14 [36].
29
Figure 14 Phase transformations in E11018-MR weld metal detected by SS-DTA [36]
Additional work was completed by Alexandrov and Lippold verifying the accuracy of SS-DTA when compared to other techniques such as dilatometry and differential thermal analysis (DTA).
It was found that SS-DTA had better accuracy when determining phase transformation start and finish temperatures, and was able to detect smaller enthalpy and volume changes such as the formation of grain boundary ferrite and small amounts of bainite that cannot be detected by dilatometry. SS-DTA was able to accurately detect phase transformations during welding, casting, and the heat treatments of steel and nickel alloys [37].
The work completed by David Stone on 410 stainless steel involved running SS-DTA during heat treatments to identify the A1, A3, Ms, and Mf temperature. In his work he found that detecting the A1 and A3 temperatures proved difficult due to the Curie transition of Type 410 steels being close to the A1 temperature on heating causing an overlap in endothermic effects on these transition temperatures. In order to validate the A1 temperature the samples were instead heated 10oC above and below the A1 temperature in order to test for a martensite transformation
30 from fresh austenite formed above the A1 temperature, and the lack of a martensite transformation below the A1 temperature indicated in Figure 15 [30].
Figure 15 An example of a) martensite transformation after heat treating above A1 and b) lack of martensite transformation when heated treated below A1 [30]
2.6 Tempering Response of Type 410 Steels
2.6.1 Post Weld Heat Treatment
Due to the hard martensitic microstructure associated with Type 410 steel in the as welded condition, a PWHT is required to decrease the hardness, which can be above 500HV, and achieve the desired ductility and toughness by tempering the martensite. It is essential that the weld cools to allow for complete martensite transformation before PWHT in order to successfully temper the material [1], [3], [8]. If austenite is still present before PWHT it will result in fresh martensite transformation upon cooling from the PWHT. This same concept applies when preheating, and having an inter-pass temperature below the MF which allows for the tempering of previous weld passes by being reheated [1]. The softening of martensite through tempering is achieved by carbide formation which will be discussed in detail in the next
31 section. Martensitic stainless steels can be heat treated in a range of 480-750oC, but for Type
410 steel being used in the application relevant to this work the PWHT range is 760-800oC set by
ASME B31.1. Having an appropriate heat treat range is important for not exceeding the A1 temperature, and avoiding tempering embrittlement [8].
2.6.2 Carbide Formation in Type 410 Stainless Steel
The possible carbide formation in martensitic stainless steels most commonly includes M23C6, and M2C, carbides, where M denotes the main element that forms the carbide with carbon, generally Cr or Fe, at typical tempering temperatures [38]. MC, M7C3, M5C and M5C2 carbides are also commonly found in found in martensitic stainless steels, but are often not associated with Type 410 due to the low alloying additions [8].
A recent paper by Chakraborty et. al. studied the effect of tempering temperature on tempering embrittlement compared to the normalized condition in Type 410 steel from 400-650oC. The main culprit of a reduction in toughness can be attributed to the presence of Fe2C carbides that primarily form around prior austenite grain boundaries, and while these carbides increase strength, there is a significant reduction in toughness [38]. From 450-500oC there was a steep decline in impact energies attributed to fine Fe2C carbides forming along martensite lathe boundaries, and within the lathe. The toughness then starts to increase again after 550o when the
Fe2C carbides dissolve and M23C6 carbides transition from Cr15.58Fe7.42C6 to the Cr dominate
Cr22.23Fe0.77C6 carbides which indicated progress towards an effective tempering cycle, Figure 16
[38]. Tempering embrittlement has also been attributed to impurity segregation to prior austenite grain boundaries, and has been detected in type 403 steels. This may be avoided if Mo is not bound to carbides due to its ability to scavenge embrittling impurities, namely phosphorus.
32
Figure 16 Tempering temperature effect on toughness and hardness of Type 410 steel [38]
Embrittlement due to P segregation was not detected in this study, but was detected in Type 410 base material when brazed with Nicobraze filler material [38], [39]. Type 420 steels that were heat treated in the range of 400-500oC experienced secondary hardening due to the formation of
M7C3 carbides. Tempering at higher temperatures results in the coarsening of these carbides until they eventually transform into M23C6 carbides past a certain temperature. This results in a softer microstructure and more favorable properties [7], [40].
2.6.3 Time-Temperature Tempering Response
In 1945 Holloman and Jaffe developed a tempering relationship for carbon steels, and some low alloy by relating tempering time, temperature, and carbon content to hardness during isothermal heat treatments. They made assumption that there was a relationship between hardness and the diffusion equation, Equation 9, due to the relation of time and temperature to tempering. They found that this equation did not vary smoothly with the hardness, so they formulated Equation 8, known as the Holloman-Jaffe equation, which is also commonly referred to as the Larson-Miller equation due to their work relating this equation to creep performance [41], [42]. The constant C
33 HJP = T*(C+log(t))
T=Time (K)
C=Carbon Constant
t=Time (hr)
Equation 9 Hollomon-Jaffe tempering parameter [41]
in this formula is used to relate carbon content to hardness, but Holloman and Jaffe found this value to not be as critical as the tempering time and temperature to for a good relationship in this equation. A lot of work was done to build off of the Holloman-Jaffe equation, and their equation was used to relate different variables effects on hardness such as individual element’s effects on the change of hardness completed by Grange and Baughman. The work completed by Grange and Baughman showed a Holloman-Jaffe relationship could accurately predict experimental hardness data for a range of steel alloys using a C value of 18, represented in Figure 17 [43].
Nehrenberg developed a Holloman-Jaffe relation based on tempering data collected from Type
410, 416, and 420 steels, Figure 18. In his work he found a C value of 20 to give the best fit for stainless, and other high alloy steels [44].
34
Figure 17 Relation between measured hardness data and predicted hardness values using an HJP with a C value of 18 [43]
Figure 18 Nehrenberg HJP for high alloy steels using a C value of 20 [44]
35 Chapter 3: Objectives
The first objective of this study was to validate the accuracy of the A1 temperature and retained delta ferrite formulas made for Type 410 steel through experimental testing on provided Type
410 custom weld consumables and commercially made base material. Additional validation was done through characterizing the microstructure, hardness, and impact toughness of test welds made with these materials in order to test the effectiveness of selecting PWHT temperature based on the A1 temperature predictive formula, and the potential negative effects of selecting PWHT temperature within the ASME B31.3 PWHT range without taking the materials A1 temperature into consideration. The second objective was to develop a Holloman-Jaffe type tempering parameter that can be used in evaluation of the tempering response of Type 410 steel weld metals and base metals, and in selection of PWHT temperatures and durations.
36 Chapter 4: Materials and Procedure
4.1. Materials
The chemical composition of the 410 steel metal core filler wires and base metal plates tested in this study are listed in Table 3. The metal core filler wires, Alloys 1, 3, and 4, were custom made within three optimized compositional ranges, Range 1, 3, and 4, shown in Figure 12. These compositional ranges were designed to provide A1 temperature close to or above the upper limit of the PWTH temperature defined by ASTM (800 oC) and retained delta ferrite below 20% [30].
Table 3 Chemical composition of custom made metal cored weld wire and commercially made base material (wt%). o A1 and DF values based on predictive formulas. DF prediction based on 46 C/s cooling rate A1 DF C Cr Ni Mo Si Mn Cu N S P oC Vol % Alloy 1.1 0.11 11.04 0.3 0.078 0.36 0.18 0.11 0.026 0.013 0.01 795 1.5
Alloy 1.2 0.15 11.65 0.31 0.079 0.34 0.18 0.11 0.022 0.008 0.009 796 2.3
Alloy 3 0.11 11.1 0.34 0.07 0.7 0.3 0.12 0.022 0.011 0.011 792 1.6
Alloy 4.1 0.121 11.87 0.11 0.21 0.97 0.19 0.19 0.029 0.012 0.01 820 17.1
Alloy 4.2 0.15 11.63 0.12 0.19 0.89 0.18 0.16 0.022 0.009 0.01 817 7.6
Metal A 0.105 12.11 0.319 0.016 0.405 0.498 0.039 0.016 0.0038 0.0162 787 17.0
Metal B 0.106 12.40 0.37 0.053 0.385 0.614 0.048 0.0266 0.0006 0.021 776 21.2
The predicted probabilities of Ranges 1, 3, and 4 meeting the desired A1 temperature and delta ferrite content were correspondingly 96%, 94%, and 99% [28]. The initially made Alloy 1.1 and
37 4.1 filler metals had chromium content outside the Range 1 and 4 specifications. Two additional filler wires, Alloy 1.2 and Alloy 4.2, were produced that met the desired chromium contents.
4.2. Validation of A1 Temperature and Delta Ferrite Content Predictive Formulas
4.2.1. Samples Preparation
The metal cored welding filler wires and the two base materials were used in validation of the
A1 temperature and delta ferrite content predictive formulas. The A1 temperature was validated through dilatometry and single sensor differential thermal analysis (SS-DTA), and the delta ferrite content was validated through a point counting using the ASTM E562 procedure. For all testing procedures, the samples were prepared using the button melting apparatus shown in
Figure 19, which utilizes a GTAW torch to melt material placed in a copper crucible. The
Figure 19 Button melting apparatus
38 welding filler wire was sectioned into approximately 1” long sections and weighed out based on the test being performed. The sectioned weld wire was then placed in a copper crucible located in the base of the button melting apparatus. A glass cylinder was mounted in the base of the apparatus and a stand with the GTAW torch attached to it was assembled over the glass and clamped to the base of the button melter to enclose the material in an airtight seal. The chamber was filled with technically pure argon (9.4) pressurizing the chamber to 10 psi. The argon was purged by quickly releasing the gas from the chamber, and allowing it to fill back up to 10 psi 3 times to ensure a pure argon atmosphere. The crucible holding the material to be welded was water cooled using a chiller attached to the base. The GTAW torch used to melt the material was controlled through a Lincoln Electric Aspect 375 power supply using a foot pedal to manually adjust current from 0-170A to melt the material. The melts were repeated twice to ensure homogeneity in the button.
For single sensor differential thermal analysis 20 g of material was melted using the above described procedure. A third melt was done in a copper funnel attached to a cube mold in place of the copper crucible, Figure 20a. The same purging procedure was repeated for this melt.
These cube samples provided enough material to create three separate 4mm thick samples,
Figure 20b. These samples were cut using an abrasive Allied Tech Cut 5 saw. A 5/64” hole was then drilled into the side of the sample halfway into the middle of the sample for thermocouple placement.
Dilatometry samples of 12 g were produced using the same button melting procedure. The induction levitation device shown in Figure 21 was used to melt and cast the 12 g buttons into 5 mm diameter 2 in. long cylindrical samples. The button sample was placed in a glass funnel at 39 the bottom of an induction coil to hold in place before melting. A program that controls the levitation device was set for chamber pressure, cast temperature, and purge time. The chamber was purged for 80 seconds with technically pure argon. The cast temperature was set to 1560oC to ensure full melting and the power supply was turned on to start the melt. The induction coil
Figure 20 a) funnel and mold used to cast cut we ld wire into a cube b) SS-DTA samples before and after sectioning from cube cast
creates eddy current that results in resistive heating, levitation, and melting of the sample. The induction levitation device automatically turns off and the molten charge falls into a 2” copper mold, Figure 22a, to create the desired geometry for dilatometry samples, Figure 22b.
40
Figure 21 Induction levitation melting device
Figure 22 a) Copper mold used for dilatometry cast geometry b) 2” dilatometry specimen
4.2.2. Delta Ferrite Quantification
The samples used for delta ferrite quantification were 6 gram buttons prepared in the button melter. These samples were then cast a copper mold using an induction levitation melting device. The copper mold, Figure 23, was fitted with ceramic inserts to slow the cooling rate of 41 the sample. A type-C thermocouple was inserted within the melt area through a hole at the bottom of the copper mold. Diagrams of the mold, and thermocouple set-up, Figure 24a and 24b give further representation of thermocouple placement, and mold geometry. Data acquisition of the thermocouple signal was started during heating to be able to record the fast cooling rates at high temperatures. The cooling curves were processed through a MatLab program in order to calculate the cooling rate of the cast sample from peak recorded temperature to 700oC, an estimated temperature range for the delta ferrite to austenite transformation. Micrographs were taken of cross sections of the cast material after being etched in order to perform point counts.
400 point grids were overlain on the micrographs using a MatLab program, and a GUI was used to indicate martensite, or delta ferrite at each point.
Figure 23 Copper mold used for delta ferrite study
42
Figure 24 a) Dimension of the ceramic insert placed in copper mold b) Thermocouple placement and cast sample weight
4.2.3. Dilatometry
Dilatometry was performed on the cast filler metal samples and on machined samples from both base materials using a 3800 GleebleTM thermo-mechanical simulator. The samples were cleaned in an ethanol ultrasonic bath to ensure they were free of grease and oils. Type K thermocouples were attached at the center of the samples for temperature control and measurements. Each sample was centered in copper jaws with a 1” free-span to allow enough room for proper heating. The dilatometry was performed using a set of low force jaws. The right jaw is secured so it does not move during the test, and the left jaw is allowed to move freely in order for the sample to expand uninhibited during heating and cooling which is important for acquiring accurate dilatometry data. The dilatometer was secured to the middle of the sample to measure the local strain at the thermocouple location, Figure 25. The chamber was closed and filled with technically pure argon, and purged twice to ensure an argon atmosphere. ASTM A1033-04 was followed for a heating rate to acquire the near equilibrium A1 temperature. The sample would be heated up to 700oC at a rate of 10o C/s and then heated up through the A1 and A3
43
Figure 25 Dilatometry set up in the Gleeble TM 3800
temperatures at 28o C/Hr from 700oC to 1000oC. The sample was allowed to cool freely from
1000oC to room temperature with the heat being extracted through the copper jaws and in the argon atmosphere. The strain and temperature data were recorded and processed using MatLab to determine the changes in strain relative to temperature. The transformation temperatures of the data were determined by local strain variations diverged from linear data. A line was drawn starting in the linear section of the strain – temperature curve, and continued over the stain deviation from linearity. A cursor was then placed at the initial change from the linear trend to indicate the phase transformation temperature.
4.2.4 Single Sensor Differential Thermal Analysis
SS-DTA was used to indirectly determine the A1 temperature by identification off martensite transformations after heat treatment near the predicted A1 temperature. The heat treatment of the samples prepared for SS-DTA, Figure 20b, was performed in an Ulvac-Riko MILA 5000
Light Radiation Furnace, Figure 26. The applied thermal histories consisted of 10oC/min heating rate to the desired holding temperature, holding for 2 hours, and free cooling. Holding
44
Figure 26 Light radiation furnace
temperatures of 10oC above the predicted A1 temperature of each tested alloy were used to confirm partial transformation to austenite during holding followed by martensitic transformation on cooling. Correspondingly, holding temperatures of 10oC below the predicted A1 temperatures were used to confirm absence of phase transformations during holding and cooling.
The thermal histories of the test samples were measured using Type-K thermocouples and recorded using a data acquisition system (DAQ) and a personal computer. Using the SS-DTA software, the recorded cooling curves were compared to software generated reference curves to identify deviations that would correspond to phase transformations. The martensitic transformations were identified as large exothermic (positive) deviations of the measured from the reference curve.
4.4 Test Welds
Test welds were produced using the two metal core filer wires made by Hobart, Alloys 1.2, and
4.2, and base metals A and B in the configurations shown in Tables 4a, and 4b. The welds were made using GTAW on a 1/2in. thick plate using a k-joint geometry that had a 37.5o bevel, and a
45 1/8 - 5/32in. root gap. A range of weld parameters is outlined in Table 5. The higher heat inputs correspond to the root passes in these welds, and the lower heat inputs tend to match up with cap passes, but not in all instances. Quicker travel speeds also indicate a lower heat input in the weld. A 400oF preheat and interpass temperature was used. The preheat temperature, voltage, and amperage range was consistent in each weld, with the lower amperages correlating with weld passes in the root. Each weld was heat treated at two different temperatures, one based on the predicted A1 temperature of the base material within code compliance, and at 800oC, the higher limit of ASME B31.3 PWHT temperature range. A heating rate of 200oC/hr to the hold temperature was used, and welds were held at temperature for 2 hours.
Table 4 Test weld plan
Table 5 Test weld parameters including preheat temperature, voltage, amperage, travel speed and heat input Preheat Voltage Amperage Travel Speed Heat Input Weld Materials (oF) (V) (A) (in/min) (J/in) 73963- Alloy 4.2 on Metal A 400 11 110-150 0.17-1.34 436788 33337- Alloy 4.2 on Metal B 400 11 110-150 0.35-2.97 209931 43190- Alloy 1.2 on Metal A 400 11 110-150 0.27-2.29 314775 45661- Alloy 1.2 on Metal B* 400 11 110-150 1.02-2.17 97178 46 4.5 Mechanical Testing of Weld Deposits
4.5.1 Charpy V-Notch Tests
All test welds listed in section 4.4 were subjected to impact testing after a PWHT at the temperatures specified in Table 4a and 4b. Standard Charpy impact specimens were extracted with notches located at the HAZ of the bevel flat side, at the weld centerline, and in the base material were produced from the test welds. The samples were extracted from the base plate/ weld metal mid thickness and oriented normal to the welding direction. The testing was performed at 15oF (-9.4oC) and 70oC (21.1oC).
4.5.2 Hardness Testing
Hardness testing was completed in order to ensure that the NACE MR0175 250HV10 maximum hardness requirement is not exceed. Hardness traverses were made across the weldment, covering the weld metal and the heat affected zones and base metals were on both sides of the weld. A spacing of at least three times the indent diagonal was used in order to avoid making indents in plastically deformed material leading to inaccurate readings. A Vickers hardness test with 10kg load was used in line with industry practice.
4.6 Metallurgical Characterization
4.6.1 Metallographic Sample Preparation
All sample were sectioned using an Allied Techcut 5 abrasive saw fixed with a bound diamond blade at 3000RPM before mounting. The test welds were initially sectioned with a carbide tipped band saw blade. The test welds were sectioned to ensure base metal, and HAZ were included
47 along with the weld metal in the metallography sample. The samples were sectioned so they could fit on a 1.25” and 1.5” diameter mounting presses. For optical microscopy the samples was mounted in Bakelite, and for SEM characterization conductive Bakelite was used.
Once mounted the sample’s surfaces were taken through a series of grinding steps generally starting with 400 grit sandpaper and grinding in one direction until the scratches in the sample were parallel. The sample was then rotated 90o and grinded with 600 grit sandpaper. This procedure was repeated for 800 and 1200 grit sandpaper. The sample would be rotated between each step in order to easily discern between scratches from the previous and current grinding steps. Between each grinding step the sample was cleaned in an ultrasonic bath in order to remove any residual debris left from the sandpaper, and to avoid contamination once on the next grinding step. Polishing followed the grinding steps starting with a 6µm diamond paste on a microfiber cloth. This step was performed for over 2 minutes, then followed by 3µm diamond paste for around 2 minutes, the followed up with 1µm for 1 minute. Each step was followed by cleaning off the sample with a spray of an ethanol from a squirt bottle followed by an ultrasonic bath for at least a minute to ensure the sample was free of contaminants.
For optical microscopy the sample was etched in order to characterize the microstructure using a delta ferrite color etchant. This was comprised of 10g of ammonium diflourite, 1g of potassium metabisulfite, and 50mL of water. The sample was submerged in the etchant for 10-20 seconds at a time then rinsed off in water to clean the sample off. Water was used clean the mounted samples due to this etchants tendency to react with ethanol which in turn compromises the quality of the etch.
48 4.6.2 Optical Metallography
An Olympus GX-51 light optical microscope was used to image the samples after etching in order to characterize the microstructures. Images were taken at 50x, 100x, 200x, 500x, and
1000x magnification and processed through a DLS software that was able to process the image by means of adjusting brightness and contrast in order to better characterize the microstructure.
Images taken at 500x and 1000x were used for point counts in order to quantify delta ferrite volume percent.
4.7 Tempering Response Procedure
Samples used in the tempering response study were prepared the same way as the SS-DTA samples and heat treatments were also performed in the same light radiation furnace used for SS-
DTA. Thermal histories were recorded with Type-K thermocouples using the same DAQ as before. Alloy 4.2 and Base Metal B were the two alloys used in the tempering response study.
These alloys were tested at two temperatures, and six different times. One tempering temperature was selected based on the ASME B31.3 PWHT range, and the second temperature was 10oC below the predicted A1 temperature.
Alloy 4.2 was heat treated at 807oC, 10oC below its A1 temperature of 817oC, and at 760oC, the lower limit of the ASME PWHT range to see if the lower temperature would have a negative impact on tempering response. Base Metal B which had a predicted A1 temperature of 776oC, so it was heat treated at 766oC, and at the upper limit of the ASME PWHT range, 800oC to test for the formation of fresh martensite when heat treating above the A1 temperature. The Base Metal
B samples tempered at 800oC were processed through the SS-DTA software in order to detect 49 martensite transformations on cooling. Six different tempering times of 1, 5, and 30 minutes and
1, 2, and 4 hours were used in order to create a logarithmic relationship between the different tempering times. The 2 hour tempering time was included to match required the PWHT time used in industry for Type 410 steel welds. Base Metal B was also characterized in the as received condition due receiving a normalizing and tempering treatment during processing which is described in Chapter 6.
Figure 27 Tempering sample cross section and surface hardness indents
After tempering the samples were sectioned as seen in Figure 27 in order to take hardness measurements through the sample cross section, and on the samples surface. Five hardness indents were taken through the center of the cross section, and five hardness indents were taken on the surface section of the sample all at a 10kg load. Four of the hardness indents made on the
50 surface of the sample were taken near the corners of the sample, and one was taken near the center. Microstructural characterization was done on the tempered samples in order to visually observe the tempering effect, and to look for fresh martensite or delta ferrite growth in the Base
Metal B samples tempered at 800oC.
51 Chapter 5: Results and Discussion of Predictive Formula Validation and Test Weld
Characterization
5.1 Validation of A1 Predictive Formula
Validation of the A1 predictive formula started with analyzing dilatometry curves of the custom- made welding consumables and commercial base materials. All five welding filler wires and the two base materials were analyzed to validate the accuracy of the predictive formula. Figure 28, a dilatometry curve from custom Alloy 1.1, shows the expected dip in strain that occurs while heating through the A1 temperature at 790oC. Multiple dilatometry tests were performed in order to validate that the A1 formula is able to accurately predict the A1 temperature. Three tests were ran on Alloy 3 in order to establish a working procedure, and test repeatability. Alloy 1.1, 4.1, and 4.2 were tested twice for repeatability, and the two base materials were tested once after repeatability was established. The average values of the experimentally determined A1 temperatures in all tested materials are compared to the corresponding predicted values in Figure
29. This histogram shows that Equation 4 was able to predict the A1 temperature with deviations of less than 5oC for all tested materials, with the largest deviations occurring in Alloy
4.1 and Alloy 4.2. Alloy 4.1 had a predicted A1 temperature of 820oC and an experimental A1 temperature of 815oC in the second run, and Alloy 4.2 had a predicted A1 temperature of 817oC, where the first dilatometry run for 4.2 had an experimental A1 temperature of 812oC. Slight variations in experimental A1 temperatures gathered for an individual alloy could be due to a
52 temperature gradient caused by the resistive heating method of the Gleeble 3800TM, and through slight variations in dilatometer placement from test to test. The tested samples were 2 in. long with a 1 in. free span, so a uniform peak temperature only covers a small area of the sample.
Additional dilatometry curves for each tested material are in Appendix A.
Figure 28 Dilatometry curve showing A1 and A3 transformations in Alloy 1.1
In previous work, SS-DTA was used to determine the A1 temperature on heating for Type 410 steel, but it was found that an error of ±10oC occurred due overlap with the thermal effect of the
Curie transition [28]. Due to these variations, an indirect method was used that was based presence or absence of martensitic transformation during cooling after heat treatments performed correspondingly 10oC above and below the predicted A1 temperature. The martensite phase
53 transformations results in a release of latent heat on cooling which is an effective indicator of the
A1 temperature being exceeded during
Figure 29 Predicted A1 temperature vs. actual A1 temperature measured through dilatometry
PWHT. Figure 30 shows SS-DTA cooling curve from Alloy 4.1 heat treated below its predicted
A1 temperature. There was no deviation between the cooling and reference curves (δT=0) within the temperature range where a martensite transformation would occur. After PWHT above the A1 temperature, the SS-DTA cooling curve in Alloy 4.1 showed a strong exothermic effect of latent heat release at 400oC corresponding to martensitic transformation, Figure 31.
The SS-DTA curves of all tested materials are shown in Appendix B and the results are summarized in Table 5. For all alloys, transformation to fresh martensite was absent when heat treated 10oC below the predicted A1 temperature and present after heat treatments 10oC above
54 the predicted A1 temperature. These results provided additional validation for the accuracy of the
A1 temperature predictive formula.
Figure 30 SS-DTA curve for Alloy 4.1 heat treated below the predicted A1 temperature
Figure 31 SS-DTA curve for alloy 4.1 heat treated above the predicted A1 temperature
55 Table 6 SS-DTA results for all alloys tested
5.2 Validation of the Delta Ferrite Predictive Formula
Delta ferrite retention quantified through point counts was able to verify the retained delta ferrite predictive formula. Welding filler metal samples in the as cast condition were imaged and had point counts taken in three different areas of the cast. Figure 32a and 32b show the microstructure and delta ferrite content of Alloys 4.1, and 4.2. Figure 33 summarizes the average delta ferrite content for all tested welding filler metals, showing that Equation 1 can predict delta ferrite content within 2% of the actual value. The micrographs used for delta ferrite quantification of all tested alloys are shown in Appendix C.
56
Figure 32 a) Microstructure of Alloy 4.1 cast and b) Microstructure of Alloy 4.2 cast used to quantify delta ferrite
Figure 33 Actual and predicted retained delta ferrite content
57 5.3 Test Welds
After validation of the predictive formulas, Alloy 1.2 and 4.2 consumables were selected for production of test welds in base metals A and B. Alloy 4.2 had an optimal combination of 817oC
A1 temperature (above the ASTM specified PWHT temperature range) and predicted delta ferrite content of 7.6 vol% (below the threshold value of 20% that would negatively affect impact toughness). Alloy 1.2 had an optimal predicted ferrite content of 2.3 vol%, but the predicted A1 temperature was 792oC which is within the ASTM PWHT range, but it is close to the 800oC upper limit of the PWHT range, so it was expected to still result in favorable properties. A K-joint preparation was used for the test welds in order to provide access for impact testing of the HAZ. Two PWHT temperatures were selected for each weld (as shown in
Table 2 in results and discussion section) below the predicted base metal A1 temperature but within the ASME B31.1 PWHT temperature range, and at 800oC, the upper limit of that temperature range.
5.3.1 Impact Testing
5.3.1.1 Alloy 4.2 Impact Tests
The test welds were subjected to 1 hour PWHT at each of the selected temperatures, followed by
Charpy impact testing in the base metal, HAZ, and WCL at 15oF and 70oF. The average impact energies for all tested conditions and notch locations are shown in Figure 34. Both test welds exhibited very good impact properties. The impact toughness at 70oF and at 15oF in all tested locations varied in the ranges between 73 and 165 ft.-lbs and 45 and 111 ft.-lbs correspondingly.
Weld A, which base metal had higher A1 temperature and lower delta ferrite content, had better base metal and HAZ impact toughness compared to Weld B. For both welds, the lower and
58 higher PWHT temperatures brought the 70oF HAZ toughness to the range of their base metal values, while 15oF HAZ toughness was significantly lower. At 70oF, both welds had similar
WCL impact toughness values, which were close to the HAZ impact toughness of weld B.
Testing at 15oF did not result in significant loss of WCL toughness in weld B.
Figure 34 Impact toughness results for Alloy 4.2 test welds after 1 hour PWHT
The 800oC PWHT resulted in higher impact toughness at all tested locations of both welds compared to the lower temperature PWHTs. This was expected result for the Alloy 4.2 weld metal toughness, since: 1) fresh martensite would not form when PWHT was performed below the A1 temperature, and 2) higher PWHT temperature would result in better tempering of the original weld metal martensitic microstructure. However, tempering at 800oC, which was above the A1 temperatures in both base metals, also resulted in better impact toughness of the base metal and HAZ in both welds. One hour PWTH at 800oC is more efficient in tempering the
59 original martensite in HAZ compared to the 760oC and 770oC PWHTs, but is apparently insufficient to generate significant amount of fresh martensite that would negatively affect impact toughness.
5.3.1.2 Alloy 1.2 Impact Tests
Alloy 1.2 impact test specimens were subjected to 2 hour PWHT to match industry practices, where the ASME B31.3 states a 1 hour minimum heat treatment is acceptable for a 0.5 inch thick weld. Impacts were completed at 70oF for the weld center line, and HAZ for these welds, and base metal impacts were completed at 15oF with the Alloy 1.2 test welds, and the results are given in Figure 35. The weld center line, and HAZ impacts showed great results ranging from
294.7 to 68.0 ft-lbs. The HAZ in the Base Metal A welds showed the highest impact energies with averages of 264.3 and 294.7 ft-lbs. All three impacts at the 800oC PWHT temperature stopped the hammer, and did not fully break, as well as two out of three impacts at the 770oC
PWHT temperature for Base Metal A welds. The HAZ impact energies were much lower in the
Base Metal B welds with averages of 98.0 and 68.0 ft-lbs for the 780oC and 800oC PWHT’s correspondingly, but these are still very good results. The Base Metal B weld showed better impact properties at a 780oC PWHT in the WCL and HAZ. The Base Metal A welds showed better impact properties at an 800oC PWHT temperature, where Base Metal B showed better impact properties at the 780oC PWHT compared to the 800oC treatment.
Base Metal B impact tests at 15oC showed acceptable average impact energies at 40.5 and 33.5 ft-lbs, with all impacts being above 20 ft-lbs. Base Metal A had low impact energies with an average of 11.5 ft-lbs with both impact energies below 20 ft-lbs minimum required for this application, and 35.5 ft-lbs with one test giving an impact energy of 7ft-lbs, but one at 64.0ft-lbs.
60
Figure 35 Impact toughness results for Alloy 1.2 at 2 hours PWHT
These results were unexpected due to the 70oF impact results for the base material were all above
100 ft-lbs. Looking at base metal microstructures of the different welds there was a microstructural difference which can be attributed to different PWHT times. 1 hour PWHT time was used for the Alloy 4.2 weld, and a 2 hour PWHT was used for the Alloy 1.2 welds. Base metal microstructures will be discussed below.
5.3.2 Hardness Traverses
5.3.2.1 Alloy 4.2 Test Weld Hardness
Hardness traverses across the cap passes and the root area of all Alloy 4.2 welds in the as welded and PWHT conditions are shown in Figure 36 and 37. The as received weld hardness was taken at a 1 kg load. In the as welded condition, the weld metal and HAZ hardness in the cap passes varied between 290 and 443 HV1. An immediate drop to hardness values around 250 HV1 was seen in the base metal. The hardness traverse in the root area had values below 250 HV1 except
61 for one value at 251 HV1. The hardness reduction in the as welded condition can be attributed to tempering in the root area by reheating from subsequent weld passes.
The PWHT for Alloy 4.2 welded on Base Metal A was held at 1 hour at both heat treatment temperatures based on the minimum acceptable PWHT duration in ASME B31.3, where the
Alloy 4.2 welded on Base Metal B was increased to 2 hours in order to
Figure 36 Alloy 4.2 welded on Base Metal B hardness traverses (1kg load) taken in the cap passes and root pass of the weld in the as-welded condition
match common industry practices. Hardness in the root passes were below 250 HV10 in all welds. The HAZ in the Base Metal B weld showed hardness values just below 250 HV10 which could be attributed to the shorter PWHT time of 1 hour. Lower hardness values were seen in the weld metal and HAZ of the samples heat treated at 800oC showing that despite being above the
62 A1 temperature of the base metal sufficient martensite was not formed to create a hard and brittle microstructure. Despite these results the tempering response study discussed below suggests that these PWHT temperatures above the A1 temperatures are not recommended. All hardness data for Alloy 4.2 welds in the PWHT condition can be seen in Figure 37.
The cap passes showed higher hardness values than the root pass. This was to be expected due to tempering in the root caused from reheating in subsequent weld passes as seen in the as welded Alloy 4.2 weld. Values above 250 HV10 were seen in the HAZ of Alloy 4.2 on base metal A heat treated at 770oC for 1 hour, and in the weld metal of Alloy 4.2 welded on Base
Metal A heat treated at 760oC for 2 hours. The high HAZ hardness in Base Metal B matches the higher hardness in the HAZ of the root of the 770oC PWHT of the Base Metal A weld. The high hardness seen in the weld metal of the 760oC PWHT can be attributed to a heat treat temperature
57oC below the A1 temperature of Alloy 4.2. The tempering data in the next chapter mirrors these results. The rest of the test welds showed favorable hardness below 250 HV10 to go along with the good impact properties.
63
Figure 37 Alloy 4.2 test weld hardness in the four PWHT conditions . Base Metal A welds 1hr PWHT duration. Base Metal B welds 2hr PWHT duration.
5.3.2.2 Alloy 1.2 Test Weld Hardness
All Alloy 1.2 test welds were heat treated for 2 hours. The Alloy 1.2 cap pass traverses showed
o hardness above 250 HV10 in the HAZ of Alloy 1.2 on Base Metal B heat treated at 800 C seen in
Figure 38. There were three additional hardness values above 250 HV10 in the HAZ and weld
o metal for Alloy 1.2 on Base Metal B at 760 , and 1 hardness value above 250 HV10 in the Base
Metal A HAZ heat treated at 770oC. The welds heat treated at 800oC had lower hardness values in the weld metal showing a stronger tempering effect at the higher PWHT temperature. The A1 temperature for Alloy 1.2 is 796oC, so the 800oC PWHT temperature only slightly exceeds this.
64 The HAZ hardness was higher than the weld metal hardness. The base metal saw higher hardness values after the 800oC heat treatment, with the largest difference seen at the two PWHT in Base Metal A.
Alloy 1.2 Test Weld Hardness Data PWHT Cap Passes 280
260
240
220 Alloy 1.2 on Base Metal A 770C 200 Alloy 1.2 on Base Metal A 800C Alloy 1.2 on Base Metal B 760C Hardness Hardness HV10 180
160 Alloy 1.2 on Base Metal B 800C
140 -10 -5 0 5 10 Relative Indent Location
Figure 38 Alloy 1.2 test weld hardness data in the four PWHT conditions. PWHT duration of 2 hours
5.4 Metallurgical characterization
5.4.1 Weld Metal
There were different microstructural features in the bulk and root of the weld versus the cap passes of Alloy 1.2 and Alloy 4.2 welds. The bulk weld metal was primarily tempered martensite with retained delta ferrite at prior austenite grain boundaries seen in Figure 39a, and b at a lower magnification, and 39c, and d for high magnification of Alloy 1.2 and 4.2. The cap passes showed significant retention of delta ferrite. The large concentration of large ferrite
65 grains was unexpected in these weld metals due to the predicted ferrite values for Alloys 1.2 and
4.2 were 2.3 and 7.6 vol% respectively at a cooling rate of 46oC/s which was expected to be similar to cooling rates seen in welding, Figure 40a and 40b. This could be attributed to a couple different factors. One study discussed in the literature review found that faster cooling rates in martensitic/ ferritic steels resulted in retention of ferrite in the final microstructure [14]. The heat input data for the test welds showed lower heat inputs in the cap passes which would result in faster cooling rates. Another cause could be segregation in these final passes. To characterize these passes further, chemistry analysis of the welds would be needed. Segregation of
Figure 39 a) Alloy 1.2 welded on Base Metal B 760 oC PWHT bulk weld area low mag b) Alloy 4.2 welded on Base Metal B 760 oC PWHT bulk weld area low mag c) Alloy 1.2 welded on Base Metal B 760oC PWHT high mag d) Alloy 4.2 welded on Base Metal B 760 oC PWHT bulk weld area high mag
66
Figure 40 a) Alloy 1.2 welded on Base Metal B 760oC PWHT cap pass b) Alloy 4.2 welded on Base Metal B 760 oC PWHT cap pass
chromium at the prior austenite grain boundaries could also be responsible for delta ferrite along the prior austenite grains seen in Figure 39.
5.4.2 HAZ
The HAZ in these welds mostly consisted of tempered martensite. Ferrite was present along isomorphic grains of tempered martensite with an absence of rolling stringers adjacent to the fusion boundary. Further away from the fusion boundary stringers were present again. The stringers in the base material contain some ferrite. Looking at Figure 41 larger amounts of ferrite were present in some areas of the HAZ due to these stringers being easy nucleation sites of ferrite when heated up to the γ-δ phase region during welding. Despite ferrite growth in the
HAZ favorable impact results were still achieved. In areas of the HAZ delta ferrite did not nucleate into larger proportions, but it can be seen along prior austenite grain boundaries in some coarse grain regions of the HAZ where the stringers recrystallize as polygonal grains of delta ferrite seen in Figure 42a. Further away from the fusion boundary stringers are present again with delta ferrite growth seen in Figure 42b. More typical HAZ microstructures for each base
67 material can be seen in Figure 43a and 43b where the stringers recrystallize, but delta ferrite growth isn’t seen.
Fusion Boundary
Stringer DF
CGHAZ DF
WM HAZ BM
Figure 41 Alloy 4.2 welded on Base Metal B 760 oC PWHT delta ferrite growth in HAZ from base metal stringers
Figure 42 Alloy 4.2 on Base Metal B HAZ with delta ferrite growth in the a) CGHAZ and b) delta ferrite growth in stringers in the HAZ
68 CGHAZ CGHAZ
Stringers
Figure 43 a) HAZ in Base Metal A at 770 oC PWHT and b) HAZ of Base Metal B at 760 oC PWHT
5.4.3 Base Metal
The base material saw a variation of microstructures based on heat treatment temperature, and duration. Stringers can be seen in the base metal that are oriented along the rolling direction.
Base Metal A saw a significant reduction in toughness when heat treated for 2 hours than when heat treated for 1 hour. Looking at Figure 44 the 1 hour PWHT in Base Metal A a fine grain tempered martensite microstructure is observed similar to Base Metal B microstructure in Figure
46a and b. Figure 45a and b shows the Base Metal A microstructure at a 2 hour PWHT. This base material has larger grain size than the 1 hour PWHT. Looking Figure 44b there is a large and even distribution of carbides in the microstructure. There is a white etched region, and an etched region that is slightly darker. It is hard to discern whether these different colored areas are different phases with this etchant, so further characterization needs to be done to characterize them. Larger proportions of ferrite could be present in this microstructure due to heating above the A1 temperature. Heating above the A1 temperature would result in the dissolution of carbides, leaving areas of free chromium in the microstructure which could promote ferrite
69 growth during PWHT. This effect was seen at both PWHT temperatures. Base Metal B has similar microstructures at both PWHT temperatures. The 800oC saw an increase in delta ferrite size in the stringers.
Figure 44 Base Metal A 1 hr 800oC PWHT microstructure at a) low mag and b) high mag
Figure 45 Base Metal A 2 hr 800oC PWHT microstructure at a) low mag and b) high mag
70
Figure 46 Base Metal B 2 hr PWHT microstructure at a) 760 oCPWHT and b) 800 oC
71 Chapter 6: Tempering Response Study
6.1 Tempering Heat Treatment Results
The objective of this study was to determine effects of the ASME B31.3 tempering range when heat treating close to the predicted A1 temperatures discussed above. Weld consumable Alloy 4.2, and Base Metal B were used to study tempering response 10oC below the A1 temperature, and to test the effects of tempering at the lower and upper limit of the ASME B31.3 PWHT range. Testing at the ASME PWHT limits would help show the importance of knowing the A1 temperature in order to effectively temper the material after welding. A table of all hardness data can be seen in Appendix E.
6.1.1 Alloy 4.2 Tempering Response
Alloy 4.2 was tempered at 807oC, 10oC below the predicted A1 temperature in order to test the tempering effect with a temperature based on the predictive formula, and at
760oC to test the bottom limit of the ASME B31.3 PWHT range. This was tested to see if the bottom of the heat treatment temperature range would be sufficient for tempering
Alloy 4.2 which has an A1 temperature 57oC above the lower required ASME PWHT temperature. Additionally a Holloman-Jaffe parameter was developed for this material in order to determine if a selected PWHT time and temperature would sufficiently temper the material in question. Before developing the H-J parameter the hardness values were observed against tempering time to compare tempering duration effect on the degree 72 tempering. Figure 47 shows the hardness data of Alloy 4.2 at 760oC and 807oC. This data indicates that tempering at 760oC does not provide sufficient tempering to get hardness levels below 250HV10. This shows importance of knowing the A1 temperature of the material before a heat treatment is performed. Even at 4 hours the average hardness for this sample was 258 HV10. At 2 hours which is the typical PWHT duration for this application the average hardness was 264 HV10. This insufficient tempering could be attributed to a tempering temperature 57oC below the predicted A1 temperature, which may be too low to efficiently form carbides for tempering.
Figure 47 Alloy 4.2 tempering data at 760 oC and 807 oC tempering temperature. Red data point indicates samples tempered at 820 oC for 30 minutes. As welded hardness taken as average from Alloy 4.2 on Base Metal B test weld
The 807oC heat treatments showed sufficient tempering at the 2 and 4 hour heat treatments with hardness values of 243 and 231 HV10 respectively. The 1 hour heat treatment gave an average hardness of 251 HV10 which was almost sufficient for meeting
73 the 250 HV10 requirement. It should be noted that the 30 minute heat treatment was
o performed at an 820 C hold temperature, and showed an average hardness of 247 HV10.
This shows a strong effect of temperature on hardness that was able to temper the specimen to an acceptable value. Figure 48 shows the microstructure of Alloy 4.2 at different heat treatment times. As time increases a more uniform dispersion of carbides can be observed at longer tempering times indicating a stronger tempering effect. The hardness values were able to create a logarithmic relationship with time at an R2 value of
0.9894 at 760oC, and 0.895 at 807oC showing a good fit. When the 30 minute tempering at a temperature of 820oC in the 807oC tempering data set is removed the R2 value slightly increases to 0.898.
Figure 48 Microstructure of Alloy 4.2 tempered at 807oC for a.) 1 minute b.) 5 minutes c.) 2 hours d.) 4 hours
74 6.1.2 Base Metal B Tempering Response
Base Metal B was a rolled base material that was normalized at 1000oC for 30 minutes, then tempered at 690oC for 1 hour minimum which was specified in the material test report. In the as received condition Base Metal B had an average hardness of 247 HV10.
Being received in a tempered condition with a hardness already blow 250 HV10 tempering below the A1 temperature should give acceptable hardness values at all heat treatment times. Like Alloy 4.2, Base Metal B was heat treated 10oC below its A1 temperature, at 766oC. Tempering temperature of 800oC was also used to test the effect of heat treating the material above its A1 temperature, but still within the ASME B31.3
o heat treat range. At 766 C for a 1 minute hold an average hardness of 237 HV10 was achieved which was 10 HV10 below the as received hardness seen in Figure 49. A 4 hour
Figure 49 Base Metal B tempering data at 766 oC and 800 oC tempering temperature
75 o temper resulted in a hardness of 203 HV10. The hardness data of the 766 C heat treatments showed a logarithmic relationship with time with an R2 value of 0.9312.
The 800oC tempering treatment resulted in an expected increase of hardness up to the 1 hour heat treatment, but the hardness started decreasing again from the 2 hour to 4 hour heat treatment. SS-DTA was performed on these heat treatments in order to detect martensite transformations on cooling. The 1 and 5 minute tempers resulted in hardnesses of 223 and 225 HV10. SS-DTA did not detect a martensite transformation in the 1 minute hold, but was detected in the 5 minute hold. The 30 minute hold resulted in an average hardness of 300 HV10. The max average hardness achieved was 405 HV10 during the 1 hour temper. During the 2 hour and 4 hour holds there was a decrease in hardness to 225 and 223 HV10 respectively. This decrease in hardness may be due to growth of delta ferrite at these longer tempering times. The SS-DTA curves in Figure 50 show a lack of martensite transformation in the 1 minute hold. The 5 minute, 1 hour, 2 hour and 4 hour hold all indicate martensite transformations at different degrees. The 5 minute and 2 hour hold show a peak thermal difference of just under 5 and 20oC respectively. The 1 hour hold had a more dramatic thermal difference of about 48oC.
Additional SS-DTA analysis on tempering above the A1 temperature may help indicate the degree of fresh martensite transformation in this material, and a switch from fresh austenite growth to delta ferrite growth at these hold temperatures.
Overall these tempering treatments shows the importance of knowing the welding consumable, and base metal A1 temperature before welding and heat treating. Too low of a tempering heat treatment may result in insufficient tempering within the ASME
76
Figure 50 SS-DTA Curves for Base Metal B at various heat treatments at 800 oC
B31.3 PWHT range, seen with hardness values above 250 HV10 in the Alloy 4.2 heat treatments at 760oC. The tempering of Base Metal B above the A1 temperature resulted in fresh martensite transformation, and potentially delta ferrite growth during longer tempering times. While the longer tempering times resulted in lower hardness due to delta ferrite growth, this could result in lower toughness in the base material of these welds. In the 30 minute and 1 hour tempering times fresh martensite is evident. The 2 hour and 4 hour tempering times do not demonstrate a fresh martensite microstructure, and it is difficult to determine the cause of the reduction in toughness based on microstructure alone. The as received microstructure, as well as the 1 minute tempering
77 time microstructure can be seen in comparison to the other tempering samples in Figure
51a-f. The darker etched region is delta ferrite present along rolling stringers.
Figure 51 Base Metal B microstructures a) as received b) 1 minute temper c) 30 minute temper d) 1 hour temper e) 2 hour temper f) 4 hour temper
78 6.1.3 Holloman-Jaffe Parameter
Alloy 4.2 was able to show a strong relationship between hardness, time, and temperature. Hollomon-Jaffe parameters are mainly dependent of time and temperature to create a relationship, but the constant, C, which was considered by Hollomon and Jaffe to be dependent on carbon content, also plays a role in obtaining a good fit, though not as much [41]. As discussed in the literature review Grange and Baughman found a C value of 18 to be the best value for the steel used in their study. Nehrenberg, whose study included 410 stainless steel, as well as other martensitic stainless steels and low alloy steels, found that a C value of 20 was the best fit for his data. Despite these findings it has also been discussed that determining the C value based on the experimental data gathered will result in a better fit equation. For Alloy 4.2 an initial C value of 20 was used to test the fit of Nehrenberg’s parameter. This resulted in a linear relationship with an R2 value of 0.8942, which is a good fit, but it was possible that a different C value could result in a better fit relationship. Figure 52 shows a range of C values used on the
Alloy 4.2 tempering data with their respective R2 values showing which C value resulted in the best fit. An immediate decrease in the R2 value was observed when decreasing the
C value below 20. When increasing C a steady increase w as seen in the relationship until it peaked at 35 with an R2 value of 0.9627. This parameter can be seen in Figure 53.
The resulting equation, Equation 8, can be used to determine if an effective time and temperature will result in the desired hardness results.
A Holloman-Jaffe Parameter was also developed for Base Metal B. Due to Base Metal B having a lower carbon content than Alloy 4.2, 0.106 vs. 0.15 wt% C, a lower C value was
79 expected. For this alloy, a C value of 30 was found to generate the best fit. Comparisons of predictive accuracy (R2) to the C value can be seen in Figure 54 for Base Metal B tempered at 766oC. The as received (normalized and tempered at 690oC) hardness of this alloy was plotted with the tempering data showing a good relationship between hardness and the Hollomon-Jaffe relationship in Figure 55. Equation 9 was developed from regression of this data that can determine the hardness of a sample based on the
Hollomon-Jaffe parameter like with Alloy 4.2.
Figure 52 Range of R2 values with varying C values. Red data point indicates C-value with best fit
When comparing the HJP of the two materials it shows the importance of received condition of the material, along with carbon content, tempering temperature, and time.
This hardness data from cast weld metal, and from a normalized and tempered base material give very different tempering results. Further studying the tempering effect in a 80
Figure 53 Holloman-Jaffe relation for Alloy 4.2
Hardness = -0.0125*HJP+717.31
Equation 8 Holloman-Jaffe relation for predicting hardness for Alloy 4.2
Figure 54 Range of R 2 values with varying C values for Base Metal B . Red data point indicates C-value with best fit 81 simulated HAZ of the base materials provided for this project would give further characterization of the tempering effects on Type 410 steel. This would also help determine the effects of PWHT temperature on Type 410 welds and base material with
A1 temperatures that have a large enough mismatch that could result in hardness above
250HV10.
Figure 55 Holloman-Jaffe relation for Base Metal B
Hardness = -0.014*HJP+652.91
Equation 8 Holloman-Jaffe relation for predicting hardness for Base Metal B
82 Chapter 7: Conclusions
1. The A1 temperature predictive formula was validated through dilatometry within 5oC
accuracy of the predicted value. SS-DTA detected martensite transformation when
samples were heat treated 10oC above the predicted A1 temperature, and no
martensite transformation was detected when the heat treat temperature was 10oC
below the A1 temperature. A1 temperatures of the custom weld consumables ranged
from 792 to 820oC.
2. Retained delta ferrite was predicted within 2% of the actual value in cast custom weld
consumable with delta ferrite contents ranging from 0.5-17 vol%.
3. Test welds made with the custom weld consumables showed great toughness results
in the weld metal and HAZ at both impact testing temperatures. All base metal
impacts testing results showed good impact energies except for Base Metal A tested
at 15oF.
4. With appropriate PWHT temperatures hardness values below 250HV10 were achieved
in the weld metal and HAZ accept for a couple high hardness (above 250HV10) values
in the HAZ. Some expected high hardness values occurred in Alloy 4.2 when heat
treated at 760oC which agreed with hardness values seen in the tempering response
study. Base Metal A experienced high hardness in the HAZ at the 800oC PWHT
temperature in the Alloy 1.2 weld. In order to avoid these variations in the degree of
tempering between the weld metal and HAZ selecting welding consumables and base
83 materials with a similar A1 temperature would be beneficial for ensuring hardness
values below 250HV10.
5. Delta ferrite retention in the weld metal was not homogeneous through the whole
weld with large polygonal delta ferrite grains present in the cap pass, but small
amounts of ferrite along prior austenite grain boundaries in the weld passes that
experienced reheating from subsequent weld passes. This same delta ferrite growth
was seen in J. Onoro’s work on 9-12%Cr steels [25]. The HAZ experienced some
delta ferrite growth which nucleated from the stringers caused by rolling, but not in
all areas of the HAZ, and the growth was not significant.
6. The tempering response study showed that acceptable hardness values will be
achieved in the weld metal at when the tempering temperature is selected based on
the predicted A1 temperature. The lower temperature of the ASME B31.3 PWHT
range proved to be unable to achieve hardness values below 250HV10, even with a
tempering time of 4 hours. Base Metal B in the as received condition (normalized
and tempered), obtained a hardness value below 250HV10, and achieved hardness as
o low as 200HV10 when tempered at 4 hours. When tempered at 800 C Base Metal B
showed an unpredictable microstructural evolution. Fresh martensite was detectable
at all tempering times above 1 minute, and an increase in average hardness was
observed up to the 1 hour temper, but an unexpected reduction in hardness was seen
in the 2 and 4 hour tempering times.
7. Based on the hardness results of the test welds, and results in the tempering response
study it is recommended that the PWHT temperature should be selected based on the
84 predicted A1 temperature, and selecting welding consumables and base materials
with a similar A1 temperature in order to achieve the most beneficial tempering
results. With these results and recommendations the ASME B31.3 PWHT range
proved to be inadequate in selecting an appropriate PWHT temperature for Type 410
stainless steel.
8. Due to the base material being required to be austenitized and tempered after
processing delta ferrite was only present in the stringers, and did not have a negative
effect on impact properties in the HAZ.
85 References
[1] J. C. Lippold and D. J. Kotecki, Welding metallurgy and weldability of stainless steels. Hoboken, NJ: John Wiley, 2005. [2] J. C. Lippold, “Introduction to the Selection of Stainless Steels,” in Welding, Brazing and Soldering, D. L. Olson, T. A. Siewert, S. Liu, and G. R. Edwards, Eds. ASM International, 1993, pp. 431–431. [3] J. R. Davis, “Selection of Wrought Martensitic Stainless Steels,” in Welding, Brazing and Soldering, D. L. Olson, T. A. Siewert, S. Liu, and G. R. Edwards, Eds. ASM International, 1993, pp. 432–442. [4] G. Arnik, C. M. Menken, and C. W. N. Veerling, “Welding Wrought 12% Chromium Martensitic Stainless Steel,” Metal construction and British Welding Journal, pp. 286–290, Aug. 1972. [5] A. W. Marshall and J. C. M. Farrar, “Welding of Ferritic and Martensitic 11-14% Cr Steel,” Welding in the World, vol. 45, no. 4, pp. 32–55, 2001. [6] T. Koseki, H. Inoue, H. Morimoto, and S. Ohkita, “Prediction of Solidification and Phase Transformation of Stainless Steel Weld Metals.” Nippon Steel Technical Report, Apr-1995. [7] T. G. Gooch and P. Woollin, “Welding Metallurgy of Low Carbon 13% Chromium Martensitic Steels,” Supermartensitic Stainless Steels, pp. 188–195, 1999. [8] R. L. Klueh and D. R. Harries, High-chromium ferritic and martensitic steels for nuclear applications. W. Conshohocken, PA: ASTM, 2001. [9] C. A. Apple and G. Krauss, “The effect of heating rate on the martensite to austenite transformation in Fe-Ni-C alloys,” Acta Metallurgica, vol. 20, no. 7, pp. 849–856, Jul. 1972. [10] Y.-K. Lee, H.-C. Shin, D.-S. Leem, J.-Y. Choi, W. Jin, and C.-S. Choi, “Reverse transformation mechanism of martensite to austenite and amount of retained austenite after reverse transformation in Fe-3Si-13Cr-7Ni (wt-%) martensitic stainless steel,” Materials Science and Technology, vol. 19, no. 3, pp. 393–398, Mar. 2003. [11] Q. Ma, Y. Shao, Y. Liu, Z. Gao, and L. Yu, “Martensite–austenite transformation kinetics of high Cr ferritic heat-resistant steel,” International Journal of Materials Research, vol. 104, no. 10, pp. 935–940, Oct. 2013. [12] M. C. Tsai, C. S. Chiou, J. S. Du, and J. R. Yang, “Phase transformation in AISI 410 stainless steel,” Materials Science and Engineering: A, vol. 332, no. 1–2, pp. 1– 10, Jul. 2002.
86 [13] F. C. Hull, “Delta Ferrite and Martensite Formation in Stainless Steels,” pp. 193– 203, May 1973. [14] S. Sam et al., “Delta ferrite in the weld metal of reduced activation ferritic martensitic steel,” Journal of Nuclear Materials, pp. 343–348, Jul. 2014. [15] S. K. Bhambri, “Intergranular fracture in 13 wt% chromium martensitic stainless steel,” Journal of Materials Science, vol. 21, no. 5, pp. 1741–1746, May 1986. [16] P. Wang, S. P. Lu, N. M. Xiao, D. Z. Li, and Y. Y. Li, “Effect of delta ferrite on impact properties of low carbon 13Cr–4Ni martensitic stainless steel,” Materials Science and Engineering: A, vol. 527, no. 13–14, pp. 3210–3216, May 2010. [17] D. Carrouge, H. K. D. H. Bhadeshia, and P. Woollin, “Effect of δ-ferrite on impact properties of supermartensitic stainless steel heat affected zones,” Science and Technology of Welding and Joining, vol. 9, no. 5, pp. 377–389, Oct. 2004. [18] A. F. Szewczyk and J. Gurland, “A Study of the Deformation and Fracture of a Dual-Phase Steel,” Metallurgical Transactions A, vol. 13, no. 10, pp. 1821–1826, Oct. 1982. [19] L. Schafer, “Influence of delta ferrite and dendritic carbides on the impact and tensile properties of a martensitic chromium steel,” Journal of Nuclear Materials, pp. 258–263, 1998. [20] Y. Iwabuchi and S. Sawada, “Metallurgical Characteristics of a Large Hydraulic Runner Casting of Type 13Cr-Ni Stainless Steel,” ASMT, pp. 332–354, 1982. [21] T. G. Gooch and J. Ginn, “Heat-Affected Zone Toughness of SMA Welded 12%Cr Martensitic-Ferritic Steels,” Welding Research Supplement, pp. 431s–440s, 1988. [22] M. C. Balmforth and J. C. Lippold, “A Preliminary Ferritic-Martensiitic Stainless Steel Constitutioin Diagram,” Welding Journal, p. 7, Jan. 1998. [23] C. Pandey, M. M. Mahapatra, P. Kumar, and N. Saini, “Dissimilar joining of CSEF steels using autogenous tungsten-inert gas welding and gas tungsten arc welding and their effect on δ-ferrite evolution and mechanical properties,” Journal of Manufacturing Processes, vol. 31, pp. 247–259, Jan. 2018. [24] Q. Wu, S. Zheng, S. Liu, C. Li, and Q. Huang, “Effect of post-weld heat treatment on the mechanical properties of electron beam welded joints for CLAM steel,” Journal of Nuclear Materials, vol. 442, no. 1–3, pp. 512–517, Nov. 2013. [25] J. Oñoro, “Martensite microstructure of 9–12%Cr steels weld metals,” Journal of Materials Processing Technology, vol. 180, no. 1–3, pp. 137–142, Dec. 2006. [26] S. A. David, S. S. Babu, and J. M. Vitek, “Welding: Solidification and microstructure,” JOM, vol. 55, no. 6, pp. 14–20, Jun. 2003. [27] L. W. Tsay, Y. M. Chang, S. Torng, and H. C. Wu, “Improved Impact Toughness of 13Cr Martensitic Stainless Steel Hardened by Laser,” Journal of Materials Engineering and Performance, vol. 11, no. 4, pp. 422–427, Aug. 2002. [28] D. J. Stone, “Optimal Composition Window of Type 410 Welding Consumables and Base Metals for Hydro-processing Applications,” Ohio State University, Columbus, M.S. thesis, May 2017. [29] D. J. Stone, B. T. Alexandrov, and J. A. Penso, “Post Weld Heat Treatment and Formation of Untempered Martensite in 410 Steel Welds,” in Volume 7:
87 Operations, Applications and Components, Vancouver, British Columbia, Canada, 2016, p. V007T07A007. [30] D. J. Stone, B. T. Alexandrov, and J. A. Penso, “Alloy Composition and Critical Temperatures in Type 410 Steel Welds,” Welding Journal, vol. 97, no. 10, pp. 286– 296, Oct. 2018. [31] D. J. Stone, B. T. Alexandrov, and J. A. Penso, “Control of Retained Delta Ferrite in Type 410 Stainless Steel,” in Volume 6A: Materials and Fabrication, Waikoloa, Hawaii, USA, 2017, p. V06AT06A011. [32] T. A. Kop, J. Sietsma, and S. Van Der Zwaag, “Dilatometric analysis of phase transformations in hypo-eutectoid steels,” Journal of Material Science, vol. 36, pp. 519–526, 2001. [33] S.-J. Lee, M. T. Lusk, and Y.-K. Lee, “Conversional model of transformation strain to phase fraction in low alloy steels,” Acta Materialia, vol. 55, no. 3, pp. 875–882, Feb. 2007. [34] M. A. Alvarado-Meza, E. García-Sanchez, O. Covarrubias-Alvarado, A. Salinas- Rodriguez, M. P. Guerrero-Mata, and R. Colás, “Effect of the High-Temperature Deformation on the M s Temperature in a Low C Martensitic Stainless Steel,” Journal of Materials Engineering and Performance, vol. 22, no. 2, pp. 345–350, Feb. 2013. [35] F. Christien, M. T. F. Telling, and K. S. Knight, “A comparison of dilatometry and in-situ neutron diffraction in tracking bulk phase transformations in a martensitic stainless steel,” Materials Characterization, vol. 82, pp. 50–57, Aug. 2013. [36] B. T. Alexandrov and J. C. Lippold, “A New Methodology for Studying Phase Transformations in High Strength Steel Weld Metal,” presented at the 7th International Trends in Welding Research Conference, 2005, p. 6. [37] B. T. Alexandrov and J. C. Lippold, “Single Sensor Differential Thermal Analysis of Phase Transformations and Structural Changes During Welding and Postweld Heat Treatment,” Welding in the World, vol. 51, no. 11–12, pp. 48–59, Nov. 2007. [38] G. Chakraborty et al., “Study on tempering behaviour of AISI 410 stainless steel,” Materials Characterization, vol. 100, pp. 81–87, Feb. 2015. [39] G. K. . L. Goh and L. C. Lim, “Embrittlement of brazed martensitic stainless steel,” Materials Science and Technology, vol. 14, no. 3, pp. 251–256, Mar. 1998. [40] A. N. Isfahany, H. Saghafian, and G. Borhani, “The effect of heat treatment on mechanical properties and corrosion behavior of AISI420 martensitic stainless steel,” Journal of Alloys and Compounds, vol. 509, no. 9, pp. 3931–3936, Mar. 2011. [41] J. H. Holloman and J. H. Jaffe, “Time-temperature relations in tempering steel,” Transactions of the american Institute of Mining and MEtallurgical Engineers, vol. 162, pp. 223–249, 1945. [42] L. C. F. Canale, X. Yao, J. Gu, and G. E. Totten, “A historical overview of steel tempering parameters,” International Journal of Microstructure and Materials Properties, vol. 3, no. 4/5, pp. 474–525, 2008.
88 [43] R. A. Grange and R. W. Baughman, “Hardness of tempered martensite in carbon and low alloy steels,” Transaction of American Society of Metals, vol. 48, pp. 165– 197, 1956. [44] A. E. Nehrenberg, “Master curves simplify stainless tempering,” Steel, vol. 127, pp. 72–76, Oct. 1950.
89 Appendix A: Dilatometry Results
Figure 56 Alloy 3 dilatometry curve
90
Figure 57 Alloy 4.1 dilatometry Curve
Figure 58 Alloy 4.2 dilatometry curve
91
Figure 59 Base Metal A dilatometry curve
Figure 60 Base Metal B dilatometry curve
92
Table 7 All dilatometry results Material Predicted A1 (oC) Dilatometry Results (oC) 794 Alloy 1.1 794 790 791 Alloy 3 792 789 794 815 Alloy 4.1 820 824 816 Alloy 4.2 817 821 Base Metal A 786 787 Base Metal B 776 779
93 Appendix B: SS-DTA Results
Figure 61 Alloy 1.2 heat treated a) 10 oC below 796 oC and b) 10oC above 796oC
Figure 62 Alloy 4.2 heat treated a) 10 oC below 817 oC and b) 10oC above 817oC
94
Figure 63 Base Metal A heat treated a) 10 oC below 786 oC and b) 10oC above 786oC
Figure 64 Base Metal B heat treated a) 10 oC below 776 oC and b) 10oC above 776oC
95 Appendix C: As Cast Microstructures for Retained Delta Ferrite Quantification
Figure 65 Alloy 1.1 example cast microstructure. Predicted retained delta ferrite: 1.5 vol%. Delta ferrite in this image 0 vol%. Experimental average 1.48 vol%
96
Figure 66 Alloy 1.2 example cast microstructure. Predicted retained delta ferrite: 2.3 vol%. Delta ferrite in this image 4.5 vol%. Experimental average 4.32 vol%
Figure 67 Alloy 3 example cast microstructure. Predicted retained delta ferrite: 1.62 vol%. Delta ferrite in this image 0 vol%. Experimental average 0.5 vol%
97 Appendix D: Impact Toughness Results
Table 8 Alloy 4.2 welded on Base Metal A Charpy V-notch toughness test results
Alloy 4.2 on Base Metal A Alloy 4.2 on Base Metal A
770C PWHT 1 hr 800C PWHT 1 hr Impact Energy (ft-lbs) Impact Energy (ft-lbs) Base Base Metal Metal 70F 70F 1 124 1 165
WCL WCL 70F 15F 70F 15F 1 88 1 64 2 50 2 117 3 87 3 121 4 93 4 117
HAZ HAZ 70F 15F 70F 15F 1 130 1 128 2 190 2 135 3 122 3 145 4 100 4
98 Table 9 Alloy 4.2 welded on Base Metal B Charpy V-notch toughness test results
Alloy 4.2 on Base Metal B Alloy 4.2 on Base Metal B
760C PWHT 1 hr 800C PWHT 1 hr Impact Energy (ft-lbs) Impact Energy (ft-lbs) Base Base Metal Test Temperature Metal Test Temperature 70F 70F 1 xxx 1 112
WCL WCL 70F 15F 70F 15F 1 99 1 108 2 77 2 128 3 90 3 87 4 92 4 82 5 64 HAZ HAZ 70F 15F 70F 15F 1 54 1 123 2 81 2 121 3 85 3 125 4 29 4 68 5 62 5 63
99 Table 10 Alloy 1.2 welded on Base Metal A Charpy V-notch toughness test results
Alloy 1.2 on Base Metal A Alloy 1.2 on Base Metal A
770C PWHT 2 hrs 800C PWHT 2 hrs Impact Energy (ft-lbs) Impact Energy (ft-lbs) Base Base Metal Metal 70F 15F 70F 15F 1 15 7 7 2 8 8 64 WCL WCL 70F 15F 70F 15F 1 105 1 136 2 116 2 130 3 115 3 144 4
HAZ HAZ 70F 15F 70F 15F 1 294 4 298 2 298 5 296 3 201 6 290 4
100 Table 11 Alloy 1.2 welded on Base Metal B Charpy V-notch toughness test results
Alloy 1.2 on Base Metal B Alloy 1.2 on Base Metal B
780C PWHT 2 hrs 800C PWHT 2 hrs Impact Energy (ft-lbs) Impact Energy (ft-lbs) Base Base Metal Test Temperature Metal Test Temperature 70F 15F 70F 15F 1 35 1 44 46 23 WCL WCL 70F 15F 70F 15F 1 126 1 108 2 122 2 81 3 113 3 100 4 4
HAZ HAZ 70F 15F 70F 15F 1 105 1 69 2 114 2 89 3 75 3 46 4 4
101 Appendix E: Tempering Hardness Data
Table 12 Alloy 4.2 hardness data at 807oC tempering temperature at various times. *30 minute tempering time tempered at 820oC Alloy 4.2 807C 4hr 2 hr 1hr 30* min 5min 1 min 233.4 239.4 251.1 245 255.7 265.4 229.1 244.7 252.2 246.3 264.1 266.5 231 247 250.1 245.2 261.1 267.8 228 241.5 249 254.8 265.2 265.7 Hardness 230.4 246.2 249.8 246.8 259.1 269.9 (HV10) 228.8 244.4 254.2 244.4 260.8 267.2 231.2 238.8 245.1 247.2 263.9 270.7 229.8 246.3 257.6 249.8 265.2 270.4 232.5 241.7 254.4 250.1 263.8 267.9 231.7 240.5 249.3 245.5 260.8 263.3
Table 13 Alloy 4.2 hardness data at 760oC tempering temperature at various times Alloy 4.2 760C 1 min 5 min 30 min 1 hr 2 hr 4 hr 286.2 281.1 275.1 265.3 264.5 256.4 287.7 280.5 267.2 269.3 264.8 258.8 293 283.8 271.6 264.8 264.7 260.3 289.3 279.5 270.3 263.2 265.1 258.6 Hardness 287.4 279.3 267.2 271.9 259.7 260.8 (HV10) 290 275.9 271.7 264.4 267.7 252 287.4 280.6 273.4 278.1 265.8 258.6 287.3 280.2 274 269.9 265.3 259.7 288 282.7 270 269.2 257.3 256.3 287.7 284.6 270.4 273.2 261.3 259.2
102 Table 14 Base Metal B hardness data at 766oC tempering temperature at various times Base Metal B 766C 1 min 5 min 30 min 1 hr 2 hr 4hr 231.4 231.4 228 216 206.1 202.9 235.9 232.9 227.5 215.6 208.2 202.1 236.9 237.5 223.8 216.9 210.3 198.9 235.4 232.9 225.2 218.2 210.8 200.5 Hardness 238.5 234.9 224.2 218.2 207.4 202.1 (HV10) 242.1 231.9 222.4 212.5 208.6 200.1 237.5 231.4 227.5 213.8 209.1 206.1 241.1 231.4 222.8 215.4 209.1 206.6 237.5 228 227.1 213.4 208.6 202.1 240 230.5 223.3 220.1 207 206.1
Table 15 Base Metal B hardness data at 800oC tempering temperature at various times Base Metal B 800C 1 min 5 min 30 min 1 hr 2 hr 4 hr 222.3 224.9 303.9 434.6 247.1 208 225.6 227.4 292.3 434.8 248.9 211.9 226.1 223 291.5 429.4 249.8 213.5 223.2 222.2 301 433.6 258.1 216.2 Hardness 222.8 221.7 296.8 448.7 252.5 219 (HV10) 221 230.2 303.2 378.3 243.1 213.7 221.8 223.9 308.3 362.8 254 218.5 219.9 226 296.1 378 248.9 210.3 224.5 226.5 308.3 374.6 251.4 214 226.4 226 303.1 372.9 245.7 213.9
103