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the effects of welding heat input on the usability of high
Markku Pirinen THE EFFECTS OF WELDING HEAT INPUT ON THE USABILITY OF HIGH STRENGTH STEELS IN WELDED STRUCTURES Thesis for the degree of Doctor of Science (Technology) to be presented with due permission for public examination and criticism in Auditorium 1381 at Lappeenranta University of Technology, Lappeenranta, Finland, on the 25th of May, 2013, at noon. Acta Universitatis Lappeenrantaensis 514 Supervisor Professor Jukka Martikainen Faculty of Technology Department of Mechanical Engineering Lappeenranta University of Technology Finland Reviewers Professor Victor Karkhin Department of Welding and Laser Technologies St.Petersburg State Polytechnical University 29 Polytechnicheskaya, St. Petersburg 195251 Russia Professor emeritus Algirdas Bargelis (Honorary Doctor of Lappeenranta University of Technology) Faculty of Mechanical Engineering and Mechatronics Department of Manufacturing Technologies Kaunas University of Technology Kęstučio St. 27, LT-44025 Kaunas Lithuania Opponents Professor Victor Karkhin Department of Welding and Laser Technologies St.Petersburg State Polytechnical University 29 Polytechnicheskaya, St. Petersburg 195251 Russia Professor emeritus Algirdas Bargelis (Honorary Doctor of Lappeenranta University of Technology) Faculty of Mechanical Engineering and Mechatronics Department of Manufacturing Technologies Kaunas University of Technology Kęstučio St. 27, LT-44025 Kaunas Lithuania ISBN 978-952-265-399-4 ISBN 978-952-265-400-7 (PDF) ISSN-L 1456-4491 ISSN 1456-4491 Lappeenrannan teknillinen yliopisto Yliopistopaino 2013 2 ABSTRACT Markku Pirinen The effects of welding heat input on the usability of high strength steels in welded structures. Lappeenranta 2013 174 pages plus 4 appendices at 4 pages Acta Universitatis Lappeenrantaensis 514 Diss. Lappeenranta University of Technology ISBN 978-952-265-399-4 ISBN 978-952-265-400-7 (PDF) ISSN-L 1456-4491, ISSN 1456-4491 High strength steel (HSS) has been in use in workshops since the 1980s. At that time, the significance of the term HSS differed from the modern conception as the maximum yield strength of HSSs has increased nearly every year. There are three different ways to make HSS. The first and oldest method is QT (quenched and tempered) followed by the TMCP (thermomechanical controlled process) and DQ (direct quenching) methods. This thesis consists of two parts, the first of which part introduces the research topic and discusses welded HSS structures by characterizing the most important variables. In the second part of the thesis, the usability of welded HSS structures is examined through a set of laboratory tests. The results of this study explain the differences in the usability of the welded HSSs made by the three different methods. The results additionally indicate that usage of different HSSs in the welded structures presumes that manufacturers know what kind of HSS they are welding. As manufacturers use greater strength HSSs in welded structures, the demands for welding rise as well. i Therefore, during the manufacturing process, factors such as heat input, cooling time, weld quality, and more must be under careful observation. Keywords: high strength steel, usability, heat input, cooling time, high strength steel filler metal UDC 678.029.43:621.791:624.078.45:624.014.2 ii ACKNOWLEDGEMENTS This thesis has been carried out in the Department of Mechanical Engineering at Lappeenranta University of Technology. I would like first to thank Professor Jukka Martikainen for his guidance throughout this process. Your support in the major point of my work gave me bottom line that I can clarify in this journey. I want to express my utmost gratitude to Dr. Paul Kah, Dr. Mika Lohtander and Professor Timo Björk. You have given me a positive example to follow and great advice to help me to finish this thesis. Timo, you always supported me in my endeavors despite that fact that you were often very busy. I offer my sincere thanks to my colleagues for their friendly support and for our pleasant working atmosphere. Special thanks go to Harri Rötkö, Antti Heikkinen, Antti Kähkönen and Esa Hiltunen. You have done great work in the laboratory during test processes. I also wish to thank the department secretaries, Ms. Kaija Tammelin and Anna-Kaisa Partanen, for all their support in administrative issues. I also cannot forget the work of all the steel structures laboratory staff. You are all professional and I am proud that I have had opportunity to research with you. There are also many other people from Lappeenranta University of Technology that have not been mentioned, but I believe they know their contribution to this dissertation. Thank You. I thank my proofreader Miss Jennifer Riley. You have worked hard to correct my thesis into flowing English. Despite the distance between our homes, my children, their spouses, and my grandchildren are always on my mind. Your comments and lovely support during this process have been the power which has seen me through this work. iii My dearest Pirjo- thank you for your affection and patience during this journey. Without you, this never would have been possible. iv CONTENTS ABSTRACT ACKNOWLEDGEMENTS TABLE OF CONTENTS LIST OF ABBREVIATIONS AND SYMBOLS Standards ......................................................................................................... xii 1. INTRODUCTION .......................................................................................... 14 1.1. Background............................................................................................ 14 2. STATE OF ART ............................................................................................ 16 2.1. What is HSS? ........................................................................................ 17 2.2. Effects of alloying elements in HSS and in its weld ............................... 19 2.2.1. Aluminium and Silicon...................................................................... 22 2.2.2. Niobium............................................................................................ 23 2.2.3. Vanadium......................................................................................... 24 2.2.4. Titanium ........................................................................................... 25 2.2.5. Zirconium ......................................................................................... 27 2.2.6. Boron and Copper............................................................................ 27 2.2.7. Manganese and Nickel .................................................................... 28 2.2.8. Rare-earth elements ........................................................................ 28 2.3. Microstructure of welded HSS structure ................................................ 29 2.3.1. Microstructure and physical features of the HAZ ............................. 31 2.3.2. Microstructure of weld ...................................................................... 34 2.4. Undermatched, matched and overmatched filler metal.......................... 37 2.5. Heat input and cooling time ................................................................... 42 3. SCOPE OF THE RESEARCH ...................................................................... 47 4. AIM OF THE RESEARCH............................................................................. 50 5. RESEARCH METHODS ............................................................................... 52 6. EXPERIMENTAL INVESTIGATIONS ........................................................... 53 6.1. Experimental arrangement..................................................................... 53 6.1. Joint geometries and preparation .......................................................... 55 6.3. Test set up ............................................................................................. 58 6.4. Material properties ................................................................................. 61 6.5. Standard tests........................................................................................ 69 6.6. Additional material test .......................................................................... 71 6.6.1. CTOD test ........................................................................................ 72 6.6.2. Compared microstructure examination ............................................ 81 7. RESULTS AND DISCUSSION...................................................................... 83 7.1. Visual test .............................................................................................. 83 7.2. Macro photography ................................................................................ 83 7.3. Micro photography ................................................................................. 92 7.4. Radiographic tests ............................................................................... 103 7.5. Surface crack detection ....................................................................... 103 7.6. Transverse tensile test ......................................................................... 104 7.7. Transverse bend test ........................................................................... 112 7.8. Impact test ........................................................................................... 115 7.9. Hardness test....................................................................................... 123 7.9. CTOD tests .......................................................................................... 129 v 7.11. Additional microstructure tests ........................................................... 135 7.11.1. Microstructure of the base material .............................................. 136 7.11.2. Microstructure of weld metal ........................................................ 137 7.11.3. Microstructure of HAZ of QT and TMCP HSS .............................. 138 7.11.4. Comparison of HAZ microstructure of steels QT and TMCP ....... 147 7.11.5. Microstructure study of CTOD samples after simulated welding thermal cycle............................................................................................ 150 8. DIVERGENCE IN MANUFACTURERS’ HSS’s WITH DIFFERENT HEAT INPUTS ........................................................................................................... 152 9. CONCLUSIONS.......................................................................................... 154 10. FUTURE WORK ....................................................................................... 157 11. SUMMARY................................................................................................ 158 References...................................................................................................... 160 vi LIST OF ABBREVIATIONS AND SYMBOLS Abbreviations Explanation 9R Cu A copper particle type A Ampere A Austenitising A5 Elongation at break % AC1 The temperature at which austenite starts to form when heated. Ac3 In hypoeutectoid steel, the temperature at which the transformation of ferrite into austenite is completed. AcC Accelerated-Cooled AF Acicular Ferrite AHSS Advanced High Strength Steel Al Aluminium APFIM Atom Probe Field Ion Microscopy ASTM American Society for Testing and Materials a/W Overall crack depth/ specimen width B Boron BH Bake Hardenable Bs Temperature where bainite starts to form C Carbon CCT Continue-Cooling-Temperature (diagram) CEV Carbon Equivalent Value (IIW) CET Carbon Equivalent Value (SEW 088) CGHAZ Coarse-Grain Heat Affected Zone CJP Complete Joint Penetration CMn Carbon Manganese Cr Chromium CTOD Crack-Tip Opening Displacement Cu Copper DP-CP Dual Phase or Complex Phase DQ Direct Quenching vii DQ&T Direct Quenching and Tempering E Welding Energy EN European Standard exp Exponent Fe Iron FCAW Flux-Cored Arc Welding FGHAZ Fine Grane Heat Affected Zone FSP Ferrite Site Plate GBF Grain Boundary Ferrite GMA Gas Metal Arc GMAW Gas Metal Arc Welding HAZ Heat Affected Zone HB Brinell Hardness HBW Brinell Hardness specifies the use of a tungsten carbide ball indenter HIZ Heat Impact Zone HSLA High Strength Low Alloy HSS High Strength Steel HV Vickers Hardness HY High Yield Strength ICCGHAZ Intercritically reheated Coarse-grain Heat Affected Zone ICHAZ Inter-Critical Heat Affected Zone IF-HS High Strength Interstitial Free IIW International Institute of Welding IS Isotropic ISO International Standard Organization J Joule lHAZ/e HAZ width to sample thickness K Kelvin kg Kilogram kJ/mm Kilo Joule/ millimeter M Thermomechanically rolled viii M-A, M/A Martensite-Austenite MAG Metal Active Gas (welding) Mg Magnesium MIG Metal Inert Gas min Minute ML Lath Martensite mm Millimeter Mn Manganese MnS Manganese Sulphate Mo Molybdenum MPa MegaPascal MS Martensitic Ms Temperature where martensites start to form N Nitrogen N Normalized N Newton Nb Niobium, Columbium NDT Non-destructive Testing Ni Nickel Nital HNO3 + ethanol O Oxygen P Phosphorus P180, P400 Degree of coarseness PCM Carbon equivalent formula according to Ito-Bessyo pf polygonal ferrite PF Pearlite and Ferritic PJP Partial Joint Penetration ppm Parts per million pWPS Preliminary Welding Procedure Specification Q Quenched Q Heat amount QL Quenched and Tempered+ Low notch toughness temperature ix QT Quenched and Tempered S Sulphur s Second s Plate Thickness SA-Weld Submerged Arc Weld SE(B) Three point bend specimen SFS Finnish Standard Association Si Silicon SiC Silicon carbide SMA Submerged Arc (Welding) Sn Tin StPSPU St. Petersburg State Polytechnic University T Tempered t8/5 Cooling time from 800 °C to 500 °C ∆t8/5 Cooling time from 800 °C to 500 °C Ta Tantalum TEM Transmission Electron Microscopy Ti Titanium TiN Titanium Nitride TiO Titanium Oxide TM Thermomechanical TMCP Thermomechanical Controlled Process Tp Peak Temperature TRIP Transformation-Induced Plasticity TTT Time-Temperature-Transformation (diagram) U.S.Navy United State Navy V Vanadium V Voltage W Watt W Tungsten Wf Windmanstatten ferrite WM Weld Metal WPS Welding Procedure Specification x wt% Mass fraction X-ray Röntgen radiation Zr Zirconium YAG Yttrium-Aluminium-Garnet–laser ε-Cu epsilon copper μm micrometer α alpha α ferrite ɣ gamma ɣ austenite π pi μ mu δ delta λ lambda σ sigma ∞ Infinite °C degrees Celsius, degrees centigrade % percent ∆ delta η eta η arc heat efficiency xi Standards ASTM E 1290-2 Standard test method for crack-tip opening displacement (CTOD) fracture toughness measurement ASTM E 112-10 Standard Test Methods for Determining Average Grain Size SEW 088:1993 German standard, weldable fine grained steels; guidelines for processing, particular for fusion welding SFS-EN 10204 Metallic products. Types of inspection documents SFS-EN 1321 Destructive tests on welds in metallic materials. Macroscopic and microscopic examination of welds SFS-EN 1435 Non-destructive examination of welds. Radiographic examination of welded joints SFS-EN 571-1 Non destructive testing. Penetrant testing. Part 1: General principles SFS-EN ISO 148-1 Metallic materials. Charpy pendulum impact test. Part 1: Test method (ISO 148-1:2009) SFS-EN ISO 15164-1Specification and qualification of welding procedures for metallic materials. Welding procedure test. Part 1: Arc and gas welding of steels and welding of nickel and nickel alloys. SFS-EN ISO17637 Non-destructive testing of welds. Visual testing of fusionwelded joints (ISO 17637:2003) SFS-EN ISO 23277 Non-destructive testing of welds. Penetrant testing of welds. Acceptance levels (ISO 23277:2006) SFS-EN ISO 5173 Destructive tests on welds in metallic materials. Bend tests (ISO 5173:2009) SFS-EN ISO 4063 Welding and allied processes. Nomenclature of processes and reference numbers (ISO 4063:2009, Corrected version 2010-03-01) SFS-EN ISO 4136 Destructive tests on welds in metallic materials. Transverse tensile test (ISO 4136:2001) xii SFS-EN ISO 6057-1Metallic materials. Vickers hardness test. Part 1: Test method (ISO 6507-1:2005) SFS-EN ISO 6892-1Metallic materials. Tensile testing. Part 1: Method of test at room temperature (ISO 6892-1:2009) xiii 1. INTRODUCTION Welding is the most commonly used method to join different types of structures. In many respects, joints are the most critical components of a load-bearing steel structure. In order for the final product to be properly developed, a number of factors must be considered when manufacturing individual components, including design, processes, inspection and quality control of structure. At low service temperatures, questions about the ductility of the welded joint can arise, as the welded structure tends to low transition temperatures. This is especially the case if the joint is produced from high strength steels (HSSs). HSS has been in use in workshops since the 1980’s. At the time, the significance of the term HSS differed from the modern conception because maximum yield strength of HSSs has increased nearly every year. During the 1980’s, the maximum yield strength of weldable HSS was 500 MPa, whereas today it is at least 1000 MPa or more. In the beginning, only a few manufacturers had HSS, which was represented through a limited assortment of products. Today, HSS is constructed worldwide with most of the modern global production consisting of structural steel which is measured in tons with an approximate yield strength 355 MPa. 1.1. Background The need of utilization of HSS grows continuously. Currently, HSSs are used more frequently and in a diverse number of industries. Primarily, HSS was just used in the car industry, but today the material is used in a more diverse assortment of industries and locations including the arms of cranes and the frames of lumber carriers, although this list is by no means extensive. 14 To date, HSS has not been formally standardized. At the lower end, structure steels have a yield strength in the range of 235-355 MPa. Recent literature has stated that strong steels should have yield strength of at least 460 MPa, while steels with a yield strength of more than 550 MPa should be categorized as ultra HSSs. Today, the yield strength of some steel has increased to 1100 MPa, while in the commercial sector, steel with a rating of up to 1300 MPa (1500 MPa) is sold. There are three different ways to make HSS. First, the oldest method is the QT method (quenched and tempered method), followed by the TMCP (thermomechanical controlled process) and finally, the last method is direct quenching (DQ). The common goal of all of these above mentioned production methods is to create a steel of high yield strength and good ductility. All the steels that are created using one of these three different methods (QT, TMCP or DQ) have a bainite and/or martensite small microstructure in the main structure. TMCP steel can also have a ferrite-bainite main structure. This small microstructure is created through the alloying of various microelements such as niobium, titanium, vanadium, and boron, which in turn make inclusions like carbides and nitrides. Together with fast cooling and tempering, the resulting microstructure is small and the hardness of structure is high despite the small content of carbon. Some manufacturers have developed DQ steel to replace QT steel using this new method (Porter 2006). Additionally, chromium, nickel, molybdenum, aluminium, carbon, magnesium, silicon, phosphorus and other alloying elements are added (or are not taken away during the manufacturing process) to iron to make HSS. It is typical of HSSs to have a low carbon content which gives the steel a lower CEV (Carbon Equivalent Value) and good weldability. Before starting to use HSS in old structures, the entire structure must be redesigned. Simply thinning the structures is not enough as buckling, springing, or bending can easily occur. In their publication from GMA-welded AHSS structure, Kaputska et al. (2008) explained that it is important for designers and 15 manufacturing engineers to understand the factors that may be affected in these performances. As there are a large variety of manufacturers that make HSSs using different methods, it is important to clarify differences between these steels. Sampath (2006) explained that manufacturers must exercise extreme caution when transferring allowable limits of certified secondary construction practices from one type of HSS plate steel to another, even for same plate thickness. 2. STATE OF THE ART A large number of scientific reports and design guidelines have been published regarding the welding of HSSs (Zeman 2009, Shi & Han 2008, Liu et al. 2007, Pacyna & Dabrowski 2007, Yayla et al. 2006, Juan et al. 2003, Keehan et al. 2003, Miki et al. 2002, Zaczek & Cwiek 1993). Special attention has been devoted to welding HSSs with matching filler material, however, only a limited number of publications consider welding HSSs with undermatching filler material (Rodriques et al 2004a). In the 1980s HSS was pioneered in Japan and organized so that individual manufacturers had their own research projects on specific steels. As a result of this rigorous research, today’s steels are of much better caliber and quality. There are three different popular and widely available HSSs on the market including those manufactured through the QT, TMCP and DQ processes. QT has been available the longest and DQ HSS has only recently been developed and acquirable on the market. Consequently, most of the research has focused on QT steels, however DQ steel research has emerged in the 2000s and recently, comparing all three HSSs has been an emerging field of investigation. 16 2.1. What is HSS? The term HSS is variable concept. Today, HSSs are steels with a yield strength greater than 550 MPa. Classifying steels according to their yield strength allows for the correct comparison between different types of steels. Fig. 1 (World Auto Steels 2009) depicts the classifications of different HSS types. Conventional HSSs (HSS) have a yield strength lower than 550 MPa. Included in this group of steels are IF-HS (High Strength Interstitial Free) steels, BH (Bake Hardenable) steels, IS (Isotropic steels), CM (Carbon Magnanese) steels, and HSLA (High Strength Low Alloy) steels (World Auto Steel 2009). Advanced HSSs (AHSS) have yield strengths greater than 550 MPa. Some steels that fit into this category are TRIP (Transformation-Induced Plasticity) steels, DP-CP (Dual Phase or Complex Phase) steels, and MS (Martensitic) steels. MS steels are used in many different industries and can be found in cranes, earth-movers, harvesters, and more. Traditional HSSs, such as high-strength low-alloy (HSLA), have more than three decades of shop experience upon which to build a technology base. In contrast, users of AHSS demanded a fast track accumulation of knowledge and dissemination as they implemented these new steels. A considerable challenge arises along the total elongation and yield strength axes, as the trend shows that higher strengths steels have decreasing total elongation percentages. Manufacturers are currently looking for ways to maintain the total elongation percentages with steels of increased yield strength. 17 Figure 1. Relationship between yield strength and total elongation for various types of steels (World Auto Steel 2009). Fig. 2 depicts the developmental history of HSS for commercial use. The first HSS, S355, was developed in the 1940s with a yield strength of 355 MPa. By the 1970s, HSSs with a yield strength of up to 690 MPa had been created. By 1990, the maximum MPa had been increased to 960 MPa, and currently, HSSs with a yield strength of up to 1300 MPa can be found (Kömi 2009). History of Ultra High Strength steels Yield Strength, MPa Hardness, HBW Figure 2. The history of ultra HSS (modified from Jukka Kömi figure 2009, Rautaruukki Ltd). 18 HSSs have been used in the war industry since 1946. The U.S. Navy has used high yield (HY) strength steel, including HY-80, HY-100, and HY-130 steels (Moon et al. 2000 according to Holsberg, P.W. et al. 1989). However, these steels were originally quite expensive to make and additionally, the knowledge of this new generation of steel was kept within the government and therefore the private sector was, for a time, excluded from this new industry. The HYstrength steel corresponds with the ISO system, where the tensile strength of HY-70 (70 ksi) corresponds to 490 MPa, HY-80 (80 ksi) corresponds to 700 MPa, HY-100 (100 ksi) corresponds to 780 MPa, HY-120 (120 ksi) corresponds to 840 MPa and HY-130 (130 ksi) corresponds to 910 MPa. 2.2. Effects of alloying elements in HSS and in its weld Alloying elements are used in HSSs to reduce the phase microstructure. There are many appropriate alloying elements that can be used when making HSSs, including Cr, W, Mo, V, B, Ti, Nb, Ta, Zr, Ni, Mn and Al. Every alloy or blend of alloys has a different effect on the steel. These elements compose inclusions and precipitations such as nitrides, carbides, carbonitrides and composites in the HSS and inhibit grain growth. In order to create a HSS with a small grain size an alloy or combination of alloys should be used, and additionally planned rolling can contribute to the creation of a steel with the above mentioned desired characteristics. To prevent the growth of austenite grains, a maximum temperature, which is dependent on the alloying element, where carbides and nitrides will dissolve to austenite, must not be exceeded. Fig. 3 shows how carbide and nitride inclusions quickly dissolve into austenite once these temperatures have been exceeded. 19 Figure 3. The effects of microalloying on Al, Zr and Ti to austenite grain growth starting temperature (modified from Harri Nevalainen figure 1984). Titanium, niobium, zirconium, and vanadium are also effective grain growth inhibitors during reheating. However, for steels that are heat treated (QT, TMCP and DQ steels) these four elements may have adverse effects on hardenability because their carbides are quite stable and difficult to dissolve in austenite prior to quenching (Metal Handbook 1990). In many research projects alloying elements of HSSs and its welds have been under examination. For example, Kou (2003) reported that increasing the alloying content of weld metal increases its hardenability by pushing the nose of continuous cooling curves to longer times. Moon et al. (2000) noticed that the microhardness variations in the weld and HAZ areas can be examined to correspond with the microstructure of the weldment. At the same time they concluded that the HAZ of the base metal was the hardest region in each weldment examined, regardless of filler metal type, base metal, or heat input. Maximum hardness was reached about midway through the HAZ of each 20 weldment studied. Fig. 4 describes hardness areas with different heat inputs (4.33 kJ/mm, 2.17 kJ/mm and 1.18 kJ/mm) using HSSs HSLA100 and HY80. Figure 4. Microhardness maps of welds made with three different filler metals and different welding parameters. The corresponding microhardness scale is included at the bottom of this figure (Moon et al. 2000). Hamada (2003) reported that it is necessary to combine the values of the constituents in the steel material and the welding conditions after taking into account the necessary joint properties. In their research, they used five different HSSs, HT50, HT60, HT80 and two HT100. They concluded that the properties of the weld HAZ, especially those of the coarse grain HAZ and fine grain HAZ 21 heated to more than the AC3 transformation point, are determined by the composition of the steel along welding conditions, as seen in fig 5. Figure 5. Structural distribution within multi-layer welded joint HAZ (Hamada 2003 according to Shishida et al. 1987). Toughness deterioration is one of worst things that can happen when welding HSSs. Caballero et al. (2009) investigated HS bainite steel and concluded that a high degree of microstructural banding, as a result of an intense segregation of manganese during dendritic solification, leads to a dramatic deterioration in toughness in these advanced bainitic steels. They concluded that the stress concentration associated with heterogeneous hardness distribution in the microstructure can be considered a possible factor contributing to premature crack nucleation. 2.2.1. Aluminium and Silicon Aluminium (Al) is widely used as a deoxidizer and it was the first element used to control austenite grain growth during reheating. When Al or silicon (Si) reacts with oxygen, soft oxides are formed. These soft oxides do not create crack initiations of growth similar to what is seen in TiO precipitations (Vähäkainu 2003). However, in HSSs it has been noticed that niobium (Nb) and titanium (Ti) are more effective grain refiners than Al (Metal Handbook 1990). High Al 22 content weakens the toughness of steel, as it promotes the formation of preferred orientation of ferrite and upper bainite. Free Al promotes forming local areas which contain high contents of carbon, which are known as M-A islands. This mechanism prevents carbon diffusion and the formation of carbides (Matsuda et al. 1995). With regard to Al, Kaputska et al. (2008) have also observed that while Al has many effects in steel making, the CEV does not consider Al in its calculation. Si is one of the principal deoxidizers used in steel making. Killed steels may contain moderate amounts of Si, from 0 to a 0.6 % maximum (Metal Handbook 1990). Low-alloy steels are reinforced by Si, but Si does not affect the features of low carbon steels (Harrison & Wall 1996). 2.2.2. Niobium As an alloying element, Nb has an important role in HSS. The effects of niobium on steel and HAZ are not solely derived from niobium. Niobium affects steel and HAZ when it is combined with other alloying elements, such as Ti and V, and precipitations. In the welded joints of HS steels, the effects of niobium depend upon the heat input. If welding and using a low heat input, this will increase impact toughness, while if a high heat input is used it will decrease the impact toughness in the HAZ. In these HSSs, as carbon content increases, there in an inverse relationship as the impact toughness decreases (Tian 1998; Hatting & Pienaar 1998). In certain amounts, Nb (0.02-0.05 wt.%) increases austenite recrystallization temperature, provides strengthening by forming thermally stable, Nb(C,N) and Nb,Ti(C,N) precipitates. During fusion welding, the precipitates limit austenite grain growth in the weld HAZ, and thereby limit hardenability or improve weldability. Excessive amounts of Nb (>0.05 wt.%) can potentially impair HAZ toughness in high heat input weldments (Sampath 2005). Small additions of Nb 23 increase the yield strength of carbon steel. The addition of 0.02 % Nb can increase the yield strength of medium-carbon steel from 490 MPa to 700 MPa. This increased strength may be accompanied by considerably impaired notch toughness unless special measures are used to refine grain size during rolling. Grain refinement during rolling involves special thermomechanical processing techniques such as controlled rolling practices, low finishing temperatures for final reduction passes, and accelerated cooling after rolling is completed (Metal Handbook 1990). In HSLA steel with niobium, granular bainite is dominant within a wider cooling rate range. In addition, martensite is observed at high cooling rates with Nb 0.026 %, but is not produced in the same steel without Nb (Zhang et al. 2009). Zhang also reports that at lower cooling rates, under 32 °C/s, Nb addition suppresses grain boundary ferrite transformation and promotes the formation of granular bainite. Li et al. (2001) have reported that the addition of 0.031 % Nb to low carbon micro alloyed steel produced the largest size and greatest area of M-A phase. 2.2.3. Vanadium Vanadium (V) increases the austenite recrystallization temperature in HS steels. It provides room temperature strengthening by forming VN, V(C,N) and (V,Ti)N precipitates in ferrite (Sampath 2005). V also strengthens HSLA steels in two ways. First, the precipitation hardens the ferrite and secondly, the precipitation refines the ferrite grain size. The precipitation of V carbonitride in ferrite can develop a significant increase in strength that depends not only on the rolling process used, but also on the base composition. Carbon content above 0.13 to 0.15 % and Mn content of 1 % or more enhances the precipitation hardening, particularly when nitrogen content is at least 0.01 %. Grain size refinement depends on thermal processing (hot rolling) variables, as well as V content (Metal Handbook 1990). 24 Chen et al. (2006) have reported that there is a correlation between V content and the size of M-A particles. This is a direct correlation as the size of M-A particles increase with increased V content from 0 % to 0.151 %. When increasing V content, there is a decrease in the impact toughness in HSS. The coarse austenite and ferrite grain and M-A constituent were thought to be the main factors resulting in impact toughness deterioration. Both Chen et al. (2006) and Zhang et al. (2009) reported after their experiments on that the concentration of V should be limited to a low level, near 0.05 %. If the V content is 0.1 % or more, this results in a greater area fraction of the M-A phase, larger average and maximum sizes of M-A particles, and deterioration in toughness. 2.2.4. Titanium When considering the welding of steel, Ti is most important micro alloying element. Stable Ti nitrides that form in high temperatures inhibit grain growth in the HAZ. Consequently, because of this grain size CGHAZ cannot grow destructively (Liu & Liao 1998). Ti is unique among common alloying elements, because it provides both precipitation strengthening and sulfide shape control. Small amounts of Ti (<0.025 %) are also useful in limiting austenite grain growth in HSSs. However, it is only useful in fully killed (aluminium deoxidized) steels because of its strong deoxidizing effect. The versatility of Ti is limited because variations in O, N, and S affect the contribution of Ti as a carbide strengthener (Metal Handbook 1990). In controlled amounts (0.01-0.02 wt.%) Ti acts as a grain refiner, increases rerystallization temperature, fixes solute nitrogen as TiN, and provides strengthening by forming thermally stable, complex Ti(C,N) precipitates. During fusion welding, TiN precipitates limit austenite grain growth in the weld HAZ, thereby limiting hardenability and improving the HAZ strength and toughness. 25 Precipitation of TiN invariably reduces the HAZ toughness, especially at low temperatures (Sampath 2005). Ti can react with nitrogen in liquid condition. Large TiN precipitates will grow in steel and their formation is easier when the Ti/N ratio is large. These kinds of precipitates cannot prevent grain growth as the precipitates which form in lower temperature. Precipitates which are big and angular can nucleate cracks and decrease fatigue durability (Lee & Pan 1995). The size of some inclusions are explained in fig. 6. Figure 6. The nucleation ability of various inclusions (Lee & Pan 1995). Ti improves HAZ microstructure and toughness of welded structure with three inter-related mechanism. Those mechanism are refining of ferrite grains by the pinning effect of thermally stable Ti-nitride and Ti-oxide particles which are distributed in austenite, by formation of pure Ti-nitride and Ti-oxide particles which disperse in austenite at high temperature and then this particles can be as nucleation sites for acicular ferrite during the ɣ-α transformation. Third mechanism is formation of fine nitrides which decrease the detrimental effect of soluble nitrogen in ferrite (Rak et al. 1997). 26 2.2.5. Zirconium Zirconium can also be added to killed high-strength low-alloy steels to improve inclusion characteristics. This occurs with sulfide inclusions, where the changes in inclusion shape improve ductility in transverse bending (Metal Handbook 1990). 2.2.6. Boron and Copper Boron (B) is added to fully killed steel to improve hardenability. The average B content in steels ranges from 0.0005 to 0.003 %. When B is substituted in part for other alloys, it should be done only to alter the hardenability. The lowered alloy content may be harmful for some applications; however B is most effective in lower carbon steels (Metal Handbook 1990). According to Moon et al. (2008), the addition of B to high strength low alloy plate steel makes a fine martensite microstructure, which increases hardenability by making the prior austenite grain boundary more stable. Vickers hardness of base steels and CGHAZ increasing Cu and B content, solid-solution hardening as uncovered by Moon et al. (2008) investigation. In the same investigation, it was also noticed that Charpy V-notch toughness showed an opposite tendency. This is mainly due to the formation of the hard phase by increasing hardenability with Cu and B addition and where toughness in the CGHAZ is decreased as compared to base steels. The results published by Moon et al. (2008) indicate that Cu addition is not useful to improve the toughness of the HAZ in high strength low alloy plate steel. Hwang at al. (1998) studied that the structure of low-carbon (C 0.04 %) copper-bearing (Cu 1.8 %) alloy steel plate manufactured by the DQ&T process has been transformed into a fine structure with high dislocation density. During tempering, fine NbC and ɛ-Cu particles are precipitated in large amounts, which 27 do not get coarsened even when the tempering temperatures rise, resulting in excellent mechanical properties. The results of Hwang et al. (1998) indicate that the addition of alloying elements and the application of the DQ&T process to low-carbon alloy steel plates contribute to the production of plates with excellent strength and toughness. 2.2.7. Manganese and Nickel Manganese (Mn) improves the strength of steel without decreasing its impact toughness and is commonly used in steel making. Mn reacts with oxygen and sulphur quite easily and makes precipitations and is important because all non hopeless effects are outclosed. The use of Mn needs to limited to under 1.5 % as steel with over 1.5 % Mn content can be brittle (Vähäkainu 2003, Lindroos at el. 1986). Excessive amounts Mn increase hardenability and reduce weldability (Sampath 2005). In his study, Keehan (2004) investigates the effects of Ni and Mn in weld metal. TEM investigations in conjunction with APFIM (Atom Probe Field Ion Microscopy) concluded a mixed microstructure of martensite, bainitic and retained austenite at an alloying level where a fully martensite microstructure would normally be expected. For increased levels of Mn, a harder and more brittle mainly martensite microstructure formed. At lower levels of Mn a softer, tougher and more easily tempered microstructure with greater amount of bainite is formed. Ni reduction with Mn levels at 2 wt% lead to an increase in toughness. Hardness results showed that lower Mn and Ni levels lead to a softer weld metal (Keehan 2004). 2.2.8. Rare-earth elements Rare-earth elements, principally cerium, lanthanum, and praseodymium, can be used to provide shale control of sulphide inclusions. Sulphide inclusions, which are plastic at rolling temperatures and thus elongate and flatten during rolling, 28 adversely affect ductility in the short transverse (through-thickness) direction. The chief role of rare-earth additives is to produce rare-earth sulfide and oxysulphide inclusions, which have negligible plasticity at even the highest rolling temperatures. Excessive amount of cerium (>0.02 %) and other rareearth elements lead to oxide of oxysulphide stringers that may affect directionally. Treatment with rare-earth elements is seldom used because they produce relatively dirty steels. Treatment with calcium is preferred, because it helps with sulphide inclusions shape control (Metal Handbook 1990). 2.3. Microstructure of welded HSS structure The structure of the base metal of HSS is homogenous and the grain size is small and regular, fig. 7. When the steel is heated during welding, the homogenous microstructure changes immediately. The heat input in the HAZ is different depending on how far the area is from the fusion zone. Many features, such as hardness, ductility and impact toughness change radically, and in many cases, to defective direction. The main structure in the base of HSS is tempered martensite and/or bainite. In addition, there are other phases such as ferrite and M-A constituent. Other important parts of structure are segregations of inclusions and precipitations such as nitrides, carbides, carbonitrides and composites. Figure 7. Microstructure of TMCP HSS (own image 2010). Aspect ratio is 1:500. 29 Fig. 8 shows the schematic description from the HAZ area temperature during steel welding. The width of the HAZ depends on heat input and cooling time. A large proportion of inclusions and precipitations dissolve when the temperature is high. When this happens, there are no nether inclusions or precipitations in or near the fusion line. 1. Weld metal, 2. Fusion line, 3. CGHAZ, 4. FGHAZ, 5. Partly austenite zone 6. ICHAZ. T curve describe maximum temperature of base material during welding. Temperature Liquid Liquid + γ Austenite A1-boundary 723 °C Martensite Bainite Weld 20 HAZ area HAZ area Figure 8. Maximum temperature of base material during welding and HAZ microstructure after welding in steel (modified from Hitsaajan opas 2003). Inclusions and precipitations are important in HSS making, as they are processes which constrain the grain growth. The same texture, inclusions and precipitations, occur in HSS weld metal. Inclusions of different shapes and textures, including spherical and faceted, and agglomerations of particles were observed in the weld metals when welding HSSs with matching filler metal. The inclusions core mainly consists of a mixture of oxides of Ti, Mn, Si, and Al in different proportions, reflecting a very complex deoxidation product. Additionally, phases rich in either Mn and S, Si or Zr, C, and N, which indicates the presence of Mn sulphides, Si, or Zr carbonitrides, were also observed (Ramirez 2008). 30 2.3.1. Microstructure and physical features of the HAZ Near the fusion zone, the phase structure of base metal is coarse as a result of the high temperature of the base metal during welding. In multi-run welding, ICCGHAZ (intercritically reheated coarse grained heat affected zone) is the worst area in the base metal (Li et al. 2001; Kim et al. 1991; Davis & King 1993). Both heat input and t8/5 (cooling time from 800 °C to 500 °C) time change the microstructure of the welded base metal and these two factors must be under control while welding. There are numerous recommendations from manufacturers regarding heat input and t8/5 time. The main differences between recommendations relate to preheating and post-heating. In specifications, however, there are also differences in spotheating temperature. Using recommended values, it is possible to successfully weld HSS. In the study done by Kaputska et al. (2008), it was concluded that the fusion zone microstructure and hardness were found to be affected by the base metal chemistry, the cooling rate conditions, and the filler metal composition. The elongation of the welded structure decreases as the yield strength of HSS grows. Yasuyama et al (2007) compares steels with yield strengths ranging from 270 MPa to 980 MPa. In the study, steels were welded by the YAG laser, mash seam, and plasma arc methods. It was confirmed that the elongation of the weldment declined compared to that of the base metal, regardless of the base metal strength. This was determined by conducting a tensile test both parallel and perpendicular to the weld line. It was therefore concluded, that the elongation is very low in high strength welded structures (Yasuyama et al. 2007). 31 Lambert et al. (2000) studied the microstructure of the martensite-austenite constituent in HAZ of HSLA steel welds in relation to toughness properties. The material used in the research was HSLA steel, with a yield strength of 433 MPa. Charpy impact test results indicated that the correlation between the toughness and microstructure of low carbon steel simulated HAZs is rather complex. The amout of M-A constituents and coarseness of bainite are major metallurgical factors affecting the impact properties (Lampert et al. 2000). In the same study, Lampert et al. (2000) also noticed that retained austenite and low carbon transformed martensite have significantly different influences on cleavage fracture and impact properties of simulated HAZ microstructure, where freshly transformed high carbon martensite is much more deleterious than retained austenite. Metallographic investigations demonstrated the existence of different M-A constituents. In the most brittle zones (the ICCGHAZ), retained austenite was mostly located between bainitic packets, whereas blocky martensite and mixed M-A constituents were located at prior austenite grain boundaries. In mixed M-A constituents, austenite was distributed at the periphery, while martensite was located at the centre. This distribution of retained austenite could be a result of chemicals and/or the mechanical stabilization mechanism (Lambert et al. 2000). Furthermore, through TEM, Lambert et al. (2000) found a constituent retained austenite at room temperature. The presence of constituent may influence the thermal stability of retained austenite, as they propagate before transformation. These observations constitute preliminary investigations of the transformation mechanism of retained austenite islands. Moon et al. (2000) compared two new ultra-low-carbon matching filler metals, with HY steel (High yield, quenched and tempered, steel) of HSLA steel. Despite the low heat input, 1.2 kJ/mm, the fusion zone hardness of two of the new ultra-low-carbon matching filler metals are comparable to the base metal hardness. The results were achieved through researching the microhardness variations in the weld and HAZ areas and corresponding this with the 32 microstructure of the weldment. In addition, the heat affected zone of the base metal was the hardest region in each weldment examined, regardless of filler metal type, base metal, or heat input. The maximum hardness occurs about midway through the HAZ of each weldment studied, rather than adjacent to the fusion boundary (Moon et al. 2000). Additionally, Moon et al. (2000) studied that the fusion zone consists predominantly of lath ferrite with varying amounts (depending on location) of untempered fine lath martensite, as well as small amounts of interlath retained austenite and oxide inclusions. No polygonal ferrite or solid-state precipitates such as carbides or carbonitrides were observed in the fusion zone. The local variations in microhardness correlate well with the local variations in the microstructure. Research carried out to study the research done by Mohandas et al. (1999) has displayed that the high Ms and Bs temperatures of steel are also responsible for low softening tendency. Steel, which has longer critical cooling time for full martensite transformation, exhibited greater resistance for softening with high heat inputs. In the investigation of heat input it was realized that the number and morphology of the ML (lath martensite) in the HAZ had some variations under different weld heat inputs (E= 0.92 ~ 1.86 kJ/mm). The carbon gathers near the grain boundary and then becomes a carbide with Fe, Mn, Mo etc. so that the impact toughness decreases. The carbide has strong direction bonds with the lath microstructure which provides the low energy passage for the impact fracture and increases brittle crack sensitivity. The fine precipitate distributed inside the grain or at the boundary is favorable to improve toughness. By controlling weld heat input (E ≤ 2.0 kJ/mm), the presence of carbides in the HAZ can be removed, and therefore the impact toughness in this zone can be assured. It was also indicated from the test results of Juan at al. (2003) that the cooling time (t8/5) should be controlled (t8/5 10-20 s) to improve toughness in the HAZ. This is so, because the cooling time increases with larger weld heat 33 inputs, which increases the potential for the deterioration of impact toughness in the HAZ (Juan et al. 2003). When welding ultra HSS, with a yield strength of more than 900 MPa, with MAG welding, it is important to precisely and accurately control heat input to the lowest possible temperatures. Zeman (2009b) examined ultra HSS, with a yield strength of 1100 MPa. In the case of the joint made by the MAG method, the weld is characterized by its bainitic structure. In the HIZ (Heat Impact Zone), Zeman observed a purely martensite structure or mixture of bainite and martensite structures (Zeman 2009b). In the same study, Zeman (2009b) noticed that ultra HSS requires the linear energy of welding to be precisely adjusted. If the linear energy of welding is too low, there could be excessive hardening of the HIZ, which increases the risk of cold cracking, whereas if the linear energy of welding is too high, the strength properties can decrease. 2.3.2. Microstructure of weld The microstructure of the weld in welded HSSs should be small and homogeneous. Alloying elements are used to make inclusions in the weld and these inclusions are the beginnings of solidifications. The inclusion density tends to be quite high but the volume fraction is comparatively small. Ramirez (2008) found in his research that in the HSS filler metal the volume fraction of nonmetallic inclusions in most deposit HSS weld metals ranged from 0.2 to 0.6 %. In a few welds, the volume fraction was from 0.8 to 1.1 %. The inclusion density observed in the welds ranged from 1.2 x 108 to 5.4 x 108 particles per mm3, while the average inclusion diameter ranged from 0.3 to 0.6 μm and the maximum inclusion diameter from 0.9 to 1.7 μm. O and S levels correlate with the inclusion size and higher levels of O and S increase the inclusion size. The average inclusion size does not drastically change with combined O and S levels up to about 400 ppm. However, above 34 400 ppm, the average inclusion size increases with an increase of both O and S levels in the weld metal (Ramirez 2008). Ramirez (2008) has stated that there are dozens of different inclusions in HSS filler metal. Table 1 describes these inclusions, while fig. 9 a and b show the acceptable shape of spherical and angular inclusions, respectively. Finally, Fig. 10 shows the phase structure of one inclusion. The chemical composition of the inclusion in region a (fig.10) is 32.2O-0.5Al-1.3Si-0.9S-51.4Ti-13.7Mn (TiO2), in region b (fig.10) MnS, and in region c (fig.10) Ti-oxide. a) b) Figure 9. Inclusion of the weld in HSS (a) Spherical, (b) Angular (Ramirez 2008). Figure 10. Composites of inclusion (Ramirez 2008). The weld microstructure can be formed from many starting values. Fig. 11 illustrates those elements which must be taken into consideration when estimating the microstructure of a weld. Additionally, Mistra et al. (2005) researched different types of inclusions as seen in Table 1. 35 CHEMISTRY HARDENABILITY ELEMENTS JOINT DESIGN PARAMETERS • • • • WELD METAL MICROSTRUCTURE • Plate thickness Heat input Thermal diffusivity Joint geometry • • • Grain Boundary Ferrite Pearlite Ferrite Ferrite Site Plate Acicular Ferrite INCLUSION CONTENT HEAT INPUT • • • Current Speed Voltage Figure 11. The various factors that play a role in deciding weld microstructure (Modified from Basu & Raman 2002). 36 Table 1. Characteristic of nonmetallic inclusions (modified from Ramirez 2008). INCLUSION 1 2 3 4 INCLUSION CHARACTERISTIC CHEMICAL COMPOSITION Region a — 50.1O-0.7Mg-1.6Al-3.9Si-2.8S19.6Ti-21.4Mn Region b — 48.2O-0.9Mg-1.6Al-3.4Si-2.3S22.2Ti-21.4Mn 51.4O-1.4Al-4.5Si-1.7S-18.1Ti-22.8Mn Region a — 32.2O-0.5Al-1.3Si-0.9S-51.4Ti13.7Mn (Ti-O2) Region b MnS, Region c Ti-Oxide Region a — 32.3O-1.5Al-0.7Si-50.4Ti-15.1Mn Region b — 35.4O-3.2Al-6.1Si-0.8S-26.5Ti28.0Mn Region c — 35.3O-4.4Al-9.6Si-1.4S-3.6Ti45.8Mn DESCRIPTION O, Al, Si, S, Ti, Mn rich O, Al, Si, S, Ti, Mn rich Composite inclusion Ti-Mn oxide Table 2. Classification of precipitates of HSS with a yield strength 770 MPa into type I – IV based on size and morphology (Misra et al. 2005). 2.4. Undermatched, matched and overmatched filler metal Filler metal also has quite a considerable effect on the welded structure of HSS depending of the yield strength of filler metal corresponding with the yield strength of base metal on the filler wire used. The filler metal can be classified as either undermatched, matched or overmatched. The filler metal is undermatched when the yield strength of the filler metal is below the yield strength of the base metal. Matched filler metals have the same yield strength 37 as base metals, and overmatched filler metals have yield strength greater than the base metals. Generally, HSSs are welded by undermatched or matched filler metal, and overmatched filler metal is infrequently used as confirmed by Porter (2006). Welding HSS requires a high quality welding process, however, it is not economical to use overmatched filler metal for HSS as it does not garnish any additional benefits. Structural steels, whose yield strength is between 235 MPa and 460 MPa, are usually welded with overmatched or matched filler material. The yield strength of structural steels is lower compared to HSSs, and there are more possibilities when welding these steels. The flexibility has allowed for a greater variety of filler material research to be carried out with regards to structural steels. Only a few research projects have used undermatching filler metal when welding HSSs. A maximum undermatching valve of 10 % can be accepted for class 690 MPa yield strength HSS (Toyota 1986, Satoh & al. 1975). Pisarski and Dolby (2003) found out that in assessing the toughness of softened HAZs, the test specimen must match the practical situation in terms of yield strength, mismatch between weld deposit, and parent metal. They explained that the fracture toughness of softened HAZ regions depended on the mismatch in strength between the weld deposit and parent plate. Their research confirmed that the worst case fracture toughness of softened HAZs occurred when the HAZ undermatched in strength both the weld deposit and parent metal. Higher toughnesses were measured when either the weld metal or parent steel undermatched the HAZ in strength. Their conclusions also elaborated that the tolerance to flaws in softened HAZs critically depends on the fracture toughness of the HAZ region where tolerance reduces rapidly in a situation where the cleavage is the dominant failure mechanism (Pisarski and Dolby 2003). In a study carried out by Umekuni and Masubuchi (1997), the tensile strength test showed that the tensile strength of the undermatched weld increases due to 38 restraint by surrounding matched welds and the base metal. Results of fatigue testing showed that both undermatched and matched welds exhibited a similar relationship between crack growth rates and the stress intensity factor. Undermatched welds have proven to be effective with HSSs, reducing the need for preheating. Undermatched welds lead to lower residual stresses than matched welds, which has the potential to reduce crack initiation. The properties of the weld metal are also a factor in the effectiveness of undermatched welds on HSSs (Umekuni & Masubuchi 1997). The results of restraint cracking tests indicated that the application of undermatched welds to HSSs leads to the reduction of minimum preheating temperatures and thus preventing cold cracking on the weld metal. It is necessary to consider not only the strength of weld metal, but also its ductility, fracture toughness, and hydrogen content when selecting weld metals for undermatching (Umekuni & Masubuchi 1997). Undermatched welds have similar fatigue characteristics to matched welds, where both undermatched and matched welds have similar crack propagation rates (Umekuni & Masubuchi 1997). Additionally, with a WM undermatched yield strength level 12 %, the concentration of plastic flow in the weakest zone increased, while the strength and ductility of the weld loaded in tension decrease. This experiment was conducted with two different heat inputs (2.0 kJ/mm and 5.0 kJ/mm) on a 25 mm thick piece of 700 MPa HSS, yield strength 700 MPa. Mismatching yield strength grade between WM / BM was 0.815, when heat input was 2.0 kJ/mm and 0.765 when heat input was 5.0 kJ/mm (Loureiro 2002). Welding high strength and high hardness QT steel involves HAZ softening and is a characteristic feature of fusion welding processes and consumables used (Rodrigues et al. 2004b). 39 Initiating a simulation is one possible way to evaluate the features of a welded structure in HSS. Rodrigues et al. (2004b) used this method and concluded that the tensile strength of the soft zone determines the overall strength of the joint. In fact, independent to the level of the undermatched yield stress, the joints achieved the base plate strength in all overmatched tensile strength situations. For matched and undermatched cases, the strength of joint was strongly dependent on the HAZ dimensions. For the cases in which the ratio width of the HAZ to sample thickness was less than 1/3, the loss of strength never exceeded 10 %, even in cases of extreme strength undermatch. However, the joint strength decreased linearly with increased HAZ widths. In almost all the cases, mismatch lead to a decrease in joint ductility, which varied depending on HAZ dimensions and hardening values (Rodrigues et al. 2004b). Rodrigues et al (2004b) also wrote that the mechanical behaviour of the overall joint depends on the plastic distribution inside the HAZ. They noticed that the large undermatched tensile strength promotes strain localization in the HAZ from the start of deformation. When the HAZ dimension is very small (width < 1/3 of the thickness), it was found that the soft material can achieve the base plate strength. They further stated that if the undermatched level of yield stress is large and the HAZ width is equal to the sample thickness, then the constraint promotes premature failure in the soft zone and the overall strength of the joint decrease even more. In the matched situation of tensile strength, the HAZ constraint induces deformation to spread to the adjacent material, whereas the soft HAZ material avoids deformation. There is an apparent increase in the material strength in almost all the undermatched cases and for lHAZ/e (HAZ width to sample thickness) ratios lower than unity, which is due to constraint (Rodrigues et al. 2004b). In Complete Joint Penetration (CJP), where matching filler metal is required, one recommendation stipulates that there should be groove welds in the tension application. Duane (1997) wrote that using undermatched filler metal is useful in welds such as Partial Joint Penetration (PJP) groove welds and filler welds. In these situations, using undermatched filler material is a cost-effective and 40 desirable alternative to matched welds. Duane (1997) also explained that when welding higher strength steels with undermatched weld metal, it is important that the level of diffusible hydrogen in the deposit weld metal is appropriate for the higher strength steel that is being welded. An analysis of the microstructure and the resulting fusion zone hardness indicated that dilution of the filler metal by the base metal does play a role in weld metal microstructure evolution. Hardness traverses indicated that the weld has regions of significant hardening and softening depending on the base metal grade, filler metal type, and cooling rate conditions. The location of greatest hardening in the near HAZ (adjacent to the fusion boundary), is where the far HAZ experienced softening. The potential implications of the hardness increased in the near HAZ region are not well understood (Kapustka et al. 2008). In dynamic tensile test results of the 780 MPa butt joint and of the DP780 steels, all of these specimens failed in the softened region of the HAZ (Kapustka et al. 2008). It is clear from a large amount of research that the lower the weld strength mismatching, the higher the fracture toughness of the HAZ (Shi et al. 1998). When undermatched filler metal is used in welding HSS, a number of items must be taken into consideration. First of all, heat input and t8/5 time are two of the most important aspects to consider. These two elements depend on a number of factors, including thickness of steel, preheating, current, voltage, and the speed of welding. Some of these factors can be altered while others cannot. For example, metallurgic and chemical effects depend on base and filler material and predescrible the effects in the weld. 41 2.5. Heat input and cooling time Welding HSS is considerably more complex than welding lower yield strength structural steels. When welding HSSs, a number of quantity modifications are made during the heating process. The HAZ area has many different phase zones, and the CGHAZ is quite often the worst zone in HSS after welding. The phase structure depends on the thermal cycle, which in turn depends on heat input, work piece geometry, material properties, etc. In earlier research (Vilpas et al. 1985) low heat input was under 2.0 kJ/mm, but today low heat input correspond to values 0.5 kJ/mm or lower. When welding ultra HSSs heat input must be very low according to the recommendations of manufacturers. HSS has been studied in a number of research using different consumables and welding processes. Nevasmaa et al. (1992b) researched AcceleratedCooled (AcC) high strength TMCP steel X80 and noticed that those steel do not need to be preheated in the arc energy range from 1.5 to 5.0 kJ/mm. They also concluded that in SA-weld metals, the toughness requirement of 40 J at -40 °C was exceed throughout the arc energy range from 2.0 to 5.0 kJ/mm. Magudeeswaran et al. (2008) researched QT steel of two different types; (1) consumable made from austenitic stainless steel, and (2) low hydrogen ferritic steel. Welding with different heat inputs and two different methods (GMAW and FCAW), they concluded that the alloying content of manganese and nickel are important in the solidification process of HSS weld metals. They also noticed that the SMAW process is more useful for welding HSSs than the FCAW process. The joints produced by using the SMAW process exhibited superior tensile and impact properties and lesser degree of CGHAZ softening compared to their FCAW counterparts. 42 Wang et al. (2003) and Juan et al. (2003) researched heat input of HSS and the test results indicated that implementing a cooling time (t8/5 =10 - 20 s) improves toughness in the HAZ (when corresponding weld heat input is 1.31 - 1.86 kJ/mm). This is true, because the larger the weld heat input, the longer the cooling time and the easier it is for the deterioration of impact toughness in the HAZ. In another study carried out of Shi and Han (2008) on 800 MPa yield strength HSLA steel it was reported that the presence of allotriomorphic ferrite, bainitic ferrite and martensite exists for simulated HAZ of the test steel. This happens, because at a temperature range of 800-1300 °C, the austenite decomposes to various ferrite morphologies. In the subsequent cooling process from 800 °C to 300°C, the austenite decomposes to various ferrite morphologies. The austenite to ferrite decomposition starts with the formation of allotriomorphic ferrite at prior austenite boundaries and eventual coverage of these boundaries. With the continued cooling, the side plate ferrite may nucleate at the ferrite/ austenite boundaries and extend into the untransformed austenite grain interiors. Further cooling to even lower temperatures increases the possibility of bainitic ferrite or acicular ferrite formation. When carbide-free bainitic ferrite is formed, the remaining austenite is enriched into carbon and becomes stable. The carbon content of remaining austenite may reach 0.5 – 0.8 wt%. With further cooling as the temperature settles to room temperature, the remaining austenite may completely or partially transform to martensite (Shi & Han 2008). As the M/A constituent forms in the HAZ during bainite transformation, the carbon-enriched, untransformed regions will partially transform into martensite at low temperatures. The carbon-enriched austenite regions are formed by the rejection of carbon from ferrite to austenite following the transformation of bainite ferrite. The transformation of M/A constituent leads to the deterioration of toughness in the HAZ (Shi & Han 2008). Shi & Han (2008) also noticed that when the cooling time in simulated 800 MPa yield strength HSS is 18 s, the fracture toughness in the simulated HAZ is 43 highest. Additionally, when the value of t8/5 is 45 s or longer, the toughness of the weld deteriorates. A remarkable decrease in toughness is observed with the increased size of austenite grain and the volume fraction of the M/A constituent. The fact that the fracture toughness deteriorated drastically for the partially phase transformed HAZ may be related to the formation of a mixed microstructure, in which the M/A constituent is a distributed shape of networks (Shi & Han 2008). Liu et al. (2007) noticed in double thermal experiments that the impact toughness decreases dramatically and obvious brittlement happens in the intercritical region of CGHAZ. They investigated copper-bearing steel with a tensile strength of no less than 685 MPa. The decreased toughness and brittlement occurred, because pearlite is formed on the interface of original austenite and coarse granular bainite, which can reduce the impact toughness. The higher heat input, the more serious brittlement becomes. Thus, during multilayer welding, it is proposed to strictly control heat input. Single thermal cycle experiments show that the copper-bearing steel has a narrow range of heat-input and brittlement can easily occur in the region of CGHAZ with higher heat-input. Granular bainite transformed from austenite leads to brittlement, and the softening starts when t8/5 time is more than 7 s. The dissolution of ε-Cu and coarse lath bainite and more ferrite can cause softening of the CGHAZ. Many HSSs, particularly copper-bearing steels, have a narrow range of heat input when welding. The effective measure to avoid or reduce the softening phenomenon of CGHAZ is to limit or control the heat input during welding. During the welding thermal cycle, with increasing heat input, lath bainite becomes coarser and the amount of ferrite increases. Coarse lath bainite decreases dislocation density and ferrite is in a soft phase. Therefore, coarse lath bainite and more ferrite can cause the softening of CGHAZ (Liu, W-Y. 2007). The features of steel can vary with the cooling rate. Pacyna and Dabrovski (2007) investigated CEV 0.39 low-C, Mn-Mo, Al killed steel using different 44 cooling time in the manufacturing process. They noticed that depending on the rate of cooling, and within the air to water cooling temperature range, the new steel can attain a tensile strength between 504 MPa and 1122 MPa. The corresponding proof stress range is from 286 MPa to 478 MPa and the structure of the air cooled steel consists of ferrite, pearlite, and bainite. This research concluded that a low carbon equivalent allows for good weldability under any conditions. Depending on the welding current and travel speed combination used, significantly different dependencies on all the influencing parameters were observed even though the heat input was same. This can be attributed to differences in the weld bead morphologies. Different weld bead morphologies are likely to lead to different weld cooling rates that will affect the microstructure by itself and also different microstructural features, such as austenite grain size, inclusion parameters, which in turn, will further contribute to the final AF content (Basu & Roman 2002). The increase to the heat input increases the yield and undermatched tensile strength of the WM, and also produces an undermatched HAZ (Loureiro 2002). When the heat input is greater (4.5 kJ/mm), the weld metal can undermatch, despite the use of matching filler material (Nevasmaa & al. 1992a). If undermatching is 10 % or less, then a maximum heat input 2.0 kJ/mm can be accepted according to Nevasmaa et al. (1992a). An example of the microstructure of HSS is in fig. 12, which illustrates QT steel with a yield strength of 690 MPa or more and the CCT-diagram shows cooling curves from 1000 °C to room temperature, and together with table 3, it shows the main microstructure and hardness for this steel after different cooling times. This type of CCT-diagram can be used to describe the microstructure of high strength QT steels with a standard yield strength 690 MPa. That microstructure will form in different zones of QT steels HAZ (yield strength 690 MPa) after cooling. 45 CCT diagram QT steel 690 MPa Figure 12. CCT-diagram of QT steel which yield strength is 690 MPa or more (Modified from Dillinger Hüttenwerke AG 2008). Table 3. Example of microstructure, austenite grain size and hardness for QT steel (yield strength 690 MPa) after different maximum heating temperatures when t8/5 is 20 s (Modified from Dillinger Hüttenwerke AG 2008). PHASE STRUCTURE, HARD- 800 900 1000 1100 1200 1350 NESS °C °C °C °C °C °C Martensite % 5 10 35 50 60 70 Bainite % 55 80 60 50 40 30 Ferrite % 40 10 5 - - - HV10 227 223 275 313 328 319 11 12 10-11 6 6 2-3 Austenite grain size (ASTM) 46 3. SCOPE OF THE RESEARCH The research work reported in this thesis concerns 1. The microstructure and 2. Other features, such as hardness, yield strength, impact toughness e.g., of welded HSSs a. Using undermatching filler metal b. With varying welding heat input in different (QT, TMCP (and DQ)) HSSs. This study includes extensive experimental investigations of the HAZ of the HSS butt joint and material characterization. The joint testing portion of the research was performed at temperatures ranging from -40 to 20 °C. Some results were analyzed and assessed using CCT diagrams which are provided by material manufactures, while the CTOD test results were analyzed using equations from design guidance documents. Fracture mechanism (crack initiation and propagation) is not included in this research, because the function in this research was to compare welded steel structures made of different steels which were made with different manufacturing methods. This research solely looks at butt weld joints; fillet welds are excluded. Specifically, the study is focused on the V-joint and single-bevel butt welds. These specific joints were selected because they are widely used in many kinds of plate structures and the CTOD investigation of single-bevel butt welds were useful because there is a perpendicular fusion face research gap. In addition, Gleeble simulated tests were made to investigate CTOD in the CGHAZ. The materials used in this study were made with the 1. QT (Quenched and Tempered) method and 47 2. TMCP (Thermomechanical controlled process) method. 3. DQ (Direct Quenching) method has been limited to a theoretical discussion. Several parameters need to be considered when assessing the strength, toughness, and impact ductility of a butt welded steel structure. In the current study it was necessary to limit the number of heat input variations of the weld, so only three were selected; Q=1.0 kJ/mm, Q=1.3 kJ/mm and Q=1.7 kJ/mm. These heat input values resulted in different microstructure and mechanical properties in the HAZ area. All the welds were made using the MAG welding process, so the influence of welding processes is outside the scope of this work. The heat input during welding was controlled and good workmanship was applied in all construction phases as the same technician performed all welding operations. Thus, the potential influence in variations of weld quality was assumed to be excluded from this research. An extensive literature analysis was carried out during the preliminary period of this research project. This literature review covered the study of the different alloying elements that constitute the microstructure of different types of HSSs, earlier studies of heat input and cooling time and different levels of matching when welding various types of HSSs. Fig. 13 highlights the main ideas of this research project. All three different types of HSS must be carefully examined when planning and constructing welded structures using these steels. The main body of fig. 13 shows the main things that must be checked, such as cooling time, heat input, filler material, and both operating and manufacturing conditions. 48 QT HIGH STRENGTH STEELS HEAT INPUT - Between 0.5 to 1.7 kJ/mm TMCP HIGH STRENGTH STEELS t8/5 time - Between 5 to 20 s OPERATING CONDITIONS -Temperature - Loading -Working position -Filler metal, etc. DQ HIGH STRENGTH STEELS FILLER MATERIAL -Undermatched -Matched -Overmatched MANUFACTURE CONDITIONS - Climate - Machines and Equipment USABILITY OF QT, TMCP, DQ HIGH STRENGTH STEELS Figure 13. Fundamentals for usability of HSSs. 49 4. AIM OF THE RESEARCH The usability of HSSs in welded structures depends on a number of elements, including manufacturing methods, types of alloying elements, quantity of alloying elements, filler metals, heat input and t8/5 time, welding method, automation, and more. The number of variables are so high that not all of the characteristics of the welding can be explained, however, some of these elements should become clearer with this research. 1. The main aim of this research is to compare different HSS and their usability in welded structures. These steels have minimum yield strengths of 690 MPa (minimum yield strength of steel A was 650 MPa). The experimental portion of this study included eight pieces of different HSSs from six different factories, that were made through the QT, TMCP and DQ processes. (DQ HSSs are studied in theoretical part using earlier studies.) Some of the tests done on these HSS, such as impact energy and CTOD test, were carried out in temperatures as low as -40 °C. 2. The second aim of this research is to clarify the effects of three different heat input, 1.0 kJ/mm, 1.3 kJ/mm and 1.7 kJ/mm, with a tolerance level of ± 0.1 kJ/mm, in the HAZ on HSS which are made using QT, TMCP and DQ methods. One again, the DQ method is covered through a theoretical exercise. The hardened (martensite or/ and bainite) structure of QT HSS is more prone to heat input effects in welding. Also, TMCP HSS have different effects in welding, especially in the CGHAZ. The heat input is limited as a consequence of the base material is being welding. Additionally, the cooling time from 800 °C to 500 °C is important when discussing the microstructure of the HAZ in the base material. 3. Brittleness of the base material and the drop of transition temperature are two factors which will appear when the heat input and cooling time t8/5 are not correct. Problems in these areas lead to a decrease in ductility and impact toughness in the steel. Accordingly, one goal of this study was to investigate these base material modifications. 50 4. The target mismatch level between the weld metal and parent metal was 0.72, where mismatch is defined as the ratio of room temperature weld metal yield strength to parent steel yield strength. Therefore, the fourth aim was to clarify the influence of high level of mismatch between filler metal and base material to the usability of the welded structure. 51 5. RESEARCH METHODS Firstly, state of art was clarified using basic scientific knowledge and the newest scientific sources, including conference presentations and articles, journal articles, and books from HSS and welding. The experimental research was carried out using two research methods. In the first method, the structures were welded just as they are in normal manufacturing conditions, whereas in the second method the CGHAZ structure was simulated using Gleeble 3800 system. Standard SFS-EN ISO 15164-1, (Specification and qualification of welding procedures for metallic materials, welding procedure test, Part 1; Arc and gas welding of steels and arc welding of nickel and nickel alloys), was used as the research method in the welding tests. The Gleeble simulation was used for pieces being subjected to the CTOD test, while the CTOD tests were made using standard ASTM E1290-02, (Standard test method for crack-tip opening displacement (CTOD) fracture toughness measurements). 1. The tests using standard SFS-EN ISO 15164-1 included reconnaissance and radiographic inspections, a bending test, a tensile strength test, an impact (Charpy-V) test, hardness test, microfilming, and macro photography. A description of these tests is included and follows scientific standards. 2. Additionally, optical tests were performed to clarify the HAZ microstructure, while microhardness tests of the QT and TMCP steels with different heat input and cooling times were done as well. 3. To test the fracture mechanics of the HAZ area, CTOD tests were used. Welds were made for these tests and the CTOD test method has been used for testing fracture mechanics. 4. Gleeble simulation has been used to clarify the CGHAZ area in welded structures. Research pieces were made for simulation, and the simulation was done using the Gleeble 3800 machine, while the CTOD test method was used for testing fracture mechanics. 52 A study was also carried out on the different elements and metals related to HSSs on the basis of chosen relevance materials from various research. 6. EXPERIMENTAL INVESTIGATIONS Experimental investigations in this study were carried out to clarify main factors affecting the usability of high strength QT and TMCP steels. The steels used in this research have been picked out among worldwide common HSSs, also used in Finland today. On the eight different HSSs that were used as research material, all but one had a yield strength of 690 MPa. (This other steel was rated with a yield strength of 650 MPa.) These steels were made using different manufacturing methods including the QT and TMCP processes. These methods and the various steels have been elaborated on during the first part of this study. 6.1. Experimental arrangement All studied welds were made using the mechanization machine, as shown in fig. 14. It was made in the Laboratory of Welding Technology at the Lappeenranta University of Technology. This machine did not have any welding speed adjustment limits. The power source used in this study was a Kemppi Pro 5200 Evolution as shown in fig. 15; which is a modern machine that is commonly used in the industry. All of the welding data was collected and stored using Kemppi ProWeld Data computer software for research use. Fig. 16 shows the principle description of fastening of weldable pieces and the processes of welding. 53 Speed adjust Welding torch Frame of mechanizing machine Figure 14. Experimental mechanizing set up. Picture 15. Power source, Kemppi Pro 5200 Evolution. 54 WELDING TORCH MECHANIZING MACHINE PNEUMATIC FASTENERS WELDED TEST PIECE Figure 16. Fasten of welded pieces in mechanizing arrangement. 6.1. Joint geometries and preparation The plate thickness of all the studied HSSs was 8.0 mm (excluding one piece of 690 MPa yield strength QT HSS, steel G, which was only started being delivered to Finland in 2008 at a minimum thickness of 12 mm). A one side Vgroove as seen in fig. 17 was used. The groove angle was 60 degrees with an air gap of 1.5 mm and root edge of 1 mm. Test pieces with dimensions of 150 mm x 400 mm were welded together. Fig. 20 shows the preparation of the groove, where run-on and run-off plates were used as tacking. Tacks of welds were first welded to fasten test pieces together, with an advance angle of three degrees estimate angular distortion. Fig. 20 illustrates a good example of a complete penetration. Fiberglass tape was used as a backing ring, as shown in figs. 18 and 19. The groove was welded on one side with two welding beads, and the 12 mm HSS was welded with three welding beads. 55 Figure 17. One side V-groove preparation. Figure 18. Fiberglass tape was used as a backing ring. Picture 19. Glued backing ring. 56 Figure 20. Run-off plates pictured root of groove. Weldable pieces were fixed between holders in the mechanized machine. The welding torch was installed downright upon the welded groove. Fig. 21 shows the plate and completed torch installation. Figure 21. Fixed weldment in mechanized machine. 57 6.3. Test set up While the backing ring was primarily used to make sure that the root edges were completely melted, it is also important to consider that the backing ring shapes the surface of the backing weld. The surfaces of joint preparation were polished between weld passes and the interpass area between welds was not subjected to any heat input, and only experienced room temperature. Additionally, the amount of free wire, which depends on current levels and the pass that is being welded, was adjusted to appropriate lengths before welding. The gas run was also adjusted, and all of the welding parameters used, including pWPS’s appear appendices 1, 2 and 3. All the welding parameters used were collected and stored using Kemppi ProWeld Data computer software for research use. The equation used for t8/5 time was (two dimension heat conduction) according to standard SFS-EN 1011-2 𝑘 2 ·𝐸 2 t8/5=(4300-4.3·T0)·105· where 𝑑2 1 1 ·[(500−𝑇 )2 −(800−𝑇 )2]·F2 0 0 (1) t8/5 = cooling time between 800-500 °C (s) T0 = work temperature (°C) k = thermal efficiency (0.8 in MAG welding) E = welding energy (kJ/mm) d = thickness of welded piece (mm) F2 = Coefficient depending the type of joint in two dimensional heat conduction (it is 1.0 when the cooling curve (t8/5) in two dimensional heat conduction is in the oblique area) Used equation for heat input (Q) was 𝑄 = ɳE (2) 58 ɳ is 0.8 through 8 mm plate in MAG welding according to the standard SFS-EN ISO 4063. 60·𝑈·𝐼 𝐸 = 1000·𝑣 (3) where E = welding energy (kJ/mm) U = arc voltage (V) I = welding current (A) v = welding speed (mm/min) For thicker plates, a three dimensional equation will be used, as follows: 1 1 t8/5=(6700-5·T0)·k·E·[(500−𝑇 ) −(800−𝑇 )]·F3 0 0 (4) where; t8/5 = cooling time between 800-500 °C (s) T0 = work temperature (°C) k = thermal efficiency (0.8 in MAG welding) E = welding energy (kJ/mm) F3 = Coefficient depending on the type of joint in three dimensional heat conduction (it is 1.0 when the cooling curve (t8/5) in three dimensional heat conduction is in the oblique area) The welding parameters used in this study are shown below in table 4. 59 Table 4. MAG welding values in three different test procedures. ARC WIRE FLOW VOLTAGE WELDING FEED RATE ARC SPEED RANGE RANGE HEAT CURRENT RANGE INPUT [A] [V] [mm/min] [m/min] [l/min] root pass 220-225 22.3 243 5.8 16 1.0 225-230 25.5 275 6.8 16 1.3 260-270 29.0 270 8.0 16 1.7 260-270 30.9 230 7.6 16 The calculated cooling times for plates are in table 5. Table 5. Cooling times for heat inputs of 1.0 kJ/mm, 1.3 kJ/mm and 1.7 kJ/mm when plate thickness is 8 mm, 12mm or 15 mm. Heat input Heat input Heat input 1.0 kJ/mm 1.3 kJ/mm 1.7 kJ/mm 21 36 56 9 15 23 6 10 15 Cooling time t8/5 [s] Plate thickness 8 mm Cooling time t8/5 [s] Plate thickness 12 mm Cooling time t8/5 [s] Plate thickness 15 mm The root pass has a lower cooling time. The heat input of the root pass was 0.97 kJ/mm and it is the leader in cooling times with 8 mm plate in 17 s and with 12 mm plate in as low as 7 s. Using equation 4 the maximum heat input will be 5.0 kJ/mm and the cooling time will again be 21s – the same as was used in the 8 mm thick plate with a heat input of 1.0 kJ/mm. The three dimensional equation for heat input can be used if the plate thickness is more than 46 mm. This plate thickness can be 60 calculated so that equations 1 and 4 will be set even and then the thickness of plate will be calculated using a heat input of 1.0 kJ/mm. For example, if the plate thickness is 20 mm then the heat input can be 2.7 kJ/mm using t8/5 21 s. 6.4. Material properties HSSs, made by either the TMCP or QT manufacturing methods were the core materials of this study. The chemical properties of these steels are presented in table 6. The chemical properties are specified in the inspection certificate 3.1 (EN 10 204-3.1 2004) provided by the manufacturer. All manufacturers have stated that their steels are made according to the conditions specified in these certificates that were supplied for this study. The conditions under which the steels were created were carefully controlled. The mechanical properties of the steels in the research are not similar, as shown in table 7. The tensile strength of these steels, which have the required yield strength of 690 MPa, varies between 798 and 879 MPa. One of steels has a tensile strength of 769 MPa, but the standard yield strength value of it is 650 MPa. The change of highest tensile strength is 10 % compared to the lowest value, which is 798 MPa. Additionally, the elongation of HSS is lower than structural steel, at yield strength 235 and 355 MPa, respectively. The change of elongation in the steels used in the experiment is between 15 and 22 %. The lowest elongation percentage was seen in steels B and D, at 15 %, while the highest elongation percentage was found in steel F, at 22 %. The impact ductility of the steels in this investigation changed between 40 and 194 J, at a temperature of -40 °C, however steels A and C were tested at a temperature of -20 °C. An impact value of at least 27 J is needed for impact ductility. That means that all the reported values in the material certificates are quite exceptional compared material standards, however HSS’s have larger strength tolerance than structure steels. 61 Furthermore, table 6 presents the chemical properties of the steels that are tested with in this research and illustrates that there are varied amounts of alloying elements used in these different steels. For example, steels E and F have the most alloying elements, as Sn is found in steel F and Zr is found is steel E. Comparatively, steel H has much fewer alloying elements. The base elements in HSS are C, Si, Mn, P and S. In addition to these five core elements, steel H only includes two more elements, Cr and Mo. Mo is in all steels in this investigation, while Cr has been used in all QT steels, and Ni has been used in all irrespective of H steel. Carbon is used in the formation of all steel. Steels A and C had the lowest amount of carbon, each with 0.05 % C. The carbon content in the other steels used for the experiments was closer to 0.15 %. However, this can be explained by the fact that steels A and C are made by the TMCP method and all the others are made by the QT method. Steel B was produced through the quenched and tempered method but additionally has a low notch toughness temperature. The grade of this steel B was S690QL. These manufacturing specifications emphasize the tough features of Steel B in cold environments up to -40 ⁰C, according to standards SFS-EN 10025-6 + A1. Aluminium is also found in HSSs and is used in the deoxidation process. Of all the steels used in the scope of this research, only steels G and H do not have any Al. Furthermore, another element found in HSSs is nitrogen, which plays a role in making nitrides such as TiN. Of the steels used for this research, nitrogen is found in five of the eight steels; namely A, C, D, E and F. Boron is an important alloying element that aids to the hardness of the steel. Only small amounts of B are needed to do an adequate job, mostly under 0.005 %. B is found in steel B, D, E and F, and the hardness of steels in this investigations was between 270 HV5 and 290 HV5. Other micro alloying elements used in these steels were Nb, V, Cu and Ti. 62 8 QT QT QT QT QT D E F G H 8 12 8 8 0,160 0,140 0,140 0,137 0,130 0,049 0,159 0,052 C% 0,24 0,37 0,40 0,276 0,30 0,17 0,33 0,19 Si % 0,87 1,21 1,41 1,390 1,20 1,86 0,82 1,64 Mn % 0,011 0,013 0,011 0,013 0,009 0,008 0,008 0,010 P% 0,001 0,004 0,004 0,0013 0,002 0,004 0,001 0,003 S% + 𝑆𝑖 10 20 20 𝑁𝑖 + 60 + 𝑁𝑖 15 15 𝑀𝑜 𝑁𝑖+𝐶𝑢 + 40 + 𝐶𝑟+𝐶𝑢 𝑀𝑛+𝐶𝑢+𝐶𝑟 + 6 𝐶𝑟+𝑀𝑜+𝑉 𝑀𝑛+𝑀𝑜 6 𝑀𝑛 𝑃𝐶𝑀 = 𝐶 + 30 + 𝐶𝐸𝑇 = 𝐶 + 𝐶𝐸𝑉 = 𝐶 + QT= Quenched + Tempered + 𝑉 10 + 5𝐵 - - 0,037 0,061 0,044 0,025 0,049 0,029 Al % QL= Quenched and Tempered + Low notch toughness temperature M = TMCP 8 M C 8 8 M QL B Thickness mm A Delivery STEEL temper 0,35 0,07 0,02 0,052 0,26 - 0,3 - 63 - 0,001 0,02 0,066 0,04 - 0,05 - 0,22 0,11 0,002 0,029 0,148 0,008 0,223 0,009 Cr % Ni % Mo % - - 0,002 0,0021 0,002 - 0,0017 - B% Table 6. Chemical properties of various steels used in the research (wt %). - - 0,032 0,022 0,021 0,081 0,004 0,046 - 0,001 0,06 0,001 0,007 0,009 0,010 0,011 - 0,002 0,01 0,020 0,01 - 0,025 - - - 0,026 0,002 0,015 0,092 0,019 0,091 - - 0,0046 0,0050 0,004 0,005 - 0,006 Nb % V % Cu % Ti % N % - - 0,002 - - - - - Sn % - - - 0,0002 - - - - Zr % (3) (2) (1) 0,419 0,38 0,393 0,39 0,42 0,38 0,41 0,34 CEV 0,29 0,28 0,28 0,28 0,28 0,24 0,28 0,22 CET 0,24 0,22 0,24 0,23 0,23 0,15 0,25 0,14 PCM B C D E F G H M QL M QT QT QT QT QT OF OBSERAVA- RM A TEST MPa MPa % Av. J 650 700 15 40 -20° C CERTIFICATE 701 769 20 99 -20° C BROCHURE 690 770 14 30 -40° C CERTIFICATE 804 841 15 194 -40° C BROCHURE 700 750 15 40 -20° C CERTIFICATE 761 821 20 98 -20° C BROCHURE 700 780 14 27 -40° C FORMATION MATERIAL MATERIAL MATERIAL MATERIAL TION REH SOURCE OF IN- BROCHURE A IMPACT TEMPERATURE THICKNESS mm TEMPER STEEL DELIVERY Table 7. Mechanical properties of steels used in the research. 8 8 8 8 Rp 0,2 CERTIFICATE 818 852 15 47 -40° C BROCHURE 690 770 14 27 -40° C CERTIFICATE 793 835 16,3 103 -40° C BROCHURE 690 790 18 27 -45° C CERTIFICATE 740 798 22 40 -45° C BROCHURE 685 780 16 40 -40° C CERTIFICATE 840 879 20 145 -40° C BROCHURE 700 770 14 27 -40° C 822 864 16 156 -40° C MATERIAL MATERIAL MATERIAL MATERIAL CERTIFICATE 8 8 12 8 M = TMCP QL= Quenched and Tempered + Low notch toughness temperature QT= Quenched + Tempered The filler metal for all these steels was ESAB 12.51. The chemical analysis of which can be seen in table 8. It is an undermatched filler metal, because it has 64 a yield strength 470 MPa. The mechanical properties for this filler metal are in table 9 and a 1.2 mm fillet solid wire was used in the welding. Additionally, the shielding gas was an AGA mixing gas composed at 15 % CO2 and 85 % Ar. Table 8. Chemical Analysis of filler material OK AUTROD 12.51 (ESAB 2008). CHEMICAL Mn Cr Ni ANALYSIS C % Si % % P% S% % % Cu % N % Ti % OK AUTROD 12.51 0.07 0.89 1.45 0.012 0.02 0.05 0.04 <0.30 0.005 0.01 Table 9. Mechanical properties of filler material OK AUTROD 12.51 (ESAB 2008). YIELD TENSILE IMPACT MECHANICAL STRENGTH STRENGTH ELONGATION DUCTILITY PROPERTIES MPa MPa A5 % J COMMENT OK AUTROD 12.51 470 560 26 26 -30°C To know the real content of the weld metal, the area of the first and second pass must be measure from figure first and then calculated (figs. 22 and 23). Every different alloying element will be calculated one to one. Dilution will happen between the base material and the filler material. The first pass has a weld metal area of 52 mm x 54 mm= 2808 mm2 (the measurements 52 mm and 54 mm are measured from fig. 22). Smelted base material areas are 7 mm x 71 mm= 497 mm2 and 6 mm x 53 mm= 336 mm2. The sum of the smelted base material areas are 497 mm2 + 336 mm2 = 833 mm2. This is 30 % from all the weld area. A concentration of the alloy elements can be calculated: The concentration of QT HSS C of the first pass: Cweld = Cbase material * 0.3 + Cfiller material * 0.7 = 0.137 *0.3 + 0.07 * 0.7 = 0.0901%. 65 The same equation was applied for all the alloy elements in the first pass. The concentrations of the first pass to the welded TMCP HSS E are: Si = 0.7058%, Mn = 1.432%, P = 0.0123%, S = 0.0144%, Cr = 0.0506%, Ni = 0.0478%, Cu = 0.216%, N = 0.005% and Ti = 0.0076%. Other alloy elements are only in the base material. Then the content of the alloy elements in the weld is 30 % of the base material content. It is likely that the content of Mo was 0.3 x 0.029 % = 0.0087 % in weld and the content of Al was 0.0183%, Nb = 0.0066 %, V = 0.00006 % and B = 0.00063%. The second pass will be calculated between the base material, the first pass and the filler material. The second pass has the weld metal which will be calculated in four parts. The first area is 39 mm x 118 mm= 4608 mm2. Two triangles, (22 mm * 22 mm)/ 2 = 242 mm2 and (22 mm * 27 mm)/ 2 = 297 mm2 and second rectangle 13 mm * 22 mm = 286 mm2. The sum of weld metal is 5433 mm2. Smelted base material areas are 19 mm x 26 mm= 494 mm2 and 18 mm x 31 mm= 558 mm2. The sum of the smelted base material areas are 494 mm2 + 558 mm2 = 1052 mm2. Smelted first pass was 12 mm x 30 mm= 360 mm2. Filler metal was 5433 mm2 – 1052 mm2 – 360 mm2 = 4021 mm2. This is 74 % from all the weld area. Smelted base material was 19 % and smelted first pass was 7 % from all weld metal. A concentration of the alloy elements can then be calculated: The concentration of C of the second pass to the QT HSS E, Cweld = Cbase material * 0.19 + Cfiller material * 0.74 + Cfirst pass * 0.07= 0.137 *0.19 + 0.07 * 0.74 + 0.09 *0.901 = 0.086%. The same equation will be used for all alloy elements. All concentrations to second pass of welded TMCP HSS E are: 66 Si = 0.76%, Mn = 1.437%, P = 0.0122%, S = 0.016%, Al = 0.0129%, Cr = 0.0504%, Ni = 0.0455%, Mo = 0.00612%, B = 0.00044%, Nb = 0.0046%, V = 0.0002%, Cu < 0.241%, N = 0.005%, Ti = 0.0083% and Zr = 0.00004 %. Weld area Smelted base material Figure 22. Principle the figure to calculate the weld metal dilution of the first pass. Aspect ratio of 1:500. Surface of first pass after polishing Second pass Fusion line Weld metal First pass a) b) Smelted first pass Smelted base metal c) Figure 23. Principle figure to calculate weld metal dilution of the second pass. a) fusion line and surface of the polished first pass before welding, b) weld metal area, c) smelted base metal and first pass. Content of alloy elements in weld after welding are in table 10. 67 Table 10. Content of alloy element of QT HSS E in the first and second pass. Dilution between base material and filler material has happened in all QT HSSs in about the same proportion. This means that the content of alloy elements were at the same levels. In TMCP HSS content of C was less of than in QT HSS. It leads to smaller content of C in the weld of TMCP HSS. As in TMCP HSS C, the C content was in the first pass was 0.3 * 0.05% + 0.7 * 0.07% = 0.064 %, while in the second pass of TMCP HSS C, the C content was 0.05 *0.19 + 0.07 * 0.74 + 0.09 *0.064 = 0.067%. For all the other alloy elements, the content differences between TMCP and QT HSSs were not large. The content of other alloying elements in the weld was at same level in TMCP HSS as in QT HSS. 68 6.5. Standard tests A welding procedure test is an inclusive test for welded structures. Using this test, the usability of the welded structure can be examined. In the standard SFS-EN ISO 15164-1 welding procedure test, all of the applicable areas are tested. Testing includes both non-destructive testing (NDT) and destructive testing which shall be in accordance with the requirements of table 11. A description of these tests is provided in the enclosed standards, and all of the welding procedure tests done on all welded pieces were carried out by the chief researcher. The first test to be conducted was a visual examination of all of the pieces. Radiographic tests were made using an industrial X-ray machine, RUP-300. Additionally, penetrant testing was made to all pieces using red penetrating liquid and white development of dye. Table 11. Examination and testing of the test pieces (standard SFS-EN ISO 15164-1). 69 Metallographic specimens were polished and etched with 4 % Nital (HNO3 + ethanol) before being placed under a conventional light microscope. The polishing automat machine was a Struers TegraPol-31. Macro- and microscopic examinations were made to all the welded test pieces. The test machine for the macro photography was Wild M400 macroscope and an Olympus 4040 camera. In addition, microscopic examinations were made on all of the weldments including the HAZ area. Microfilming was made using a light microscope, Zeiss MC63, and the computer software was Isolution Lite. Additional microscopic test were done at St. Petersburg State Polytechnic University laboratory (StPSPU) using light metallographic microscope LEICA DMI5000M with magnification up to x1000 to clarify the exact microstructure of the HAZ. Additionally, the impact toughness test was measured using the standard Charpy V-notch impact test (standard SFS-EN ISO 148-1). The test temperature was -40 °C and test machine was model VEB Werkstoffpromachine Leipzig VBN with a load of 150 N. The 5 x 10 mm Charpy test pieces were shaped with a “V” notch of 2 mm depth with the notch tip in conformity with the standards of the HAZ and the weld. Vickers hardness tests were also performed on the welded specimens, to the SFS-EN ISO 6057-1 standard, using a 5 kg load. Test machine was a Zwick 3202. Four transverse bending tests were made using standard SFS-EN ISO 5173 to all welded structures, two from the weld surface and two from the root, and the machine used was a bend machine, WPN 20. The same WPN 20 machine was used to make tensile tests with an extensometer. The standard used with the tensile test was SFS-EN ISO 6892-1. The computer software used for this information was PicoLog for windows PLW recorder, and two tests were made to all welded structures. 70 When conducting a transverse bend test on HSSs, the diameter of pusher and opening of drums must be considerably larger than when testing lower yield strength steels. All of these values are found in standards SFS-EN ISO 5173 and SFS-EN ISO 15614-1. For example, an 8 mm thick plate must have a pusher diameter 45 mm and a drum opening of 65 mm, while 12 mm thick plate must have a pusher diameter 75 mm and a drum opening of 105 mm. According to the standard SFS-EN ISO 5173, the bending angle at which to conduct the bending test should be 180, however the bending machine that was used was limited to a maximum 150 angle. Equation is d=(100 x ts)/A-ts and (5) d+3 x ts ≥ l >(d+2 x ts) (6) where d= diameter of pusher A= minimum ultimate elongation of base metal ts= plate thickness l= opening of drums 6.6. Additional material test To confirm that standard SFS-EN ISO 15164-1 test has shown realistic results, an additional material test had to be conducted. CTOD tests and microstructure analysis, like analysing different faces and micro hardness, were done. HAZs were also calculated and CTOD tests were done to all HSS steels. Additional microstructure tests were performed on QT HSS steel (steel E) and TMCP HSS steel (steel C). 71 6.6.1. CTOD test In order to check if the impact toughness values were correct, a CTOD test was made on all welded structures. Fig. 24 shows the construction of the welded pieces with dimensions of 8 x 15 x 50 mm. The 8 mm in thick, 50 mm length and 15 mm lateral pieces were cut from the whole plate. Using tack welds these pieces were welded together and the fusion faces were machined. The gap was about 1.5 mm and root edge was about 1.0 mm, while a single-bevel (½-V) groove was used with a flank angle of 45 degrees. Fig. 25 illustrates the welded pieces before they were separated by saw. The beginning and end of the groove were made with assisting pieces. The test pieces with dimension 5 mm width, 10 mm high and 50 mm long were made by machining after cutting. Figure 25. CTOD test pieces after welding. Figure 24. Used one side single bevel (½-V) groove in CTOD tests. 72 In fig. 26 there is an etched CTOD test piece where the red line indicates the fusion line, the blue line indicates the start notch, the yellow line indicates the fatigue notch, and the green line indicates the test area. The groove was welded with three or four beads according to the heat input. The same three heat inputs (1.0; 1.3 and 1.7 kJ/mm) were used as in previous tests. Figure 26. Etched CTOD test piece. The CTOD test equipment was made by the Welding Technology Laboratory at Lappeenranta University of Technology. Fig. 27 illustrates the pusher and its counterpart, while fig. 28 is a picture of the actual machine used for the testing. 73 Figure 27. CTOD test components and test piece. PUSHER AND ITS COUNTERPART COMPUTER AND SOFTWARE COOLING UNIT FATIGUE TEST MACHINE Figure 28. CTOD test machine. 74 The testing temperature was -40°C. Ethanol was used to guarantee the constancy of the temperature, while and temperature adjustments were made with the application or removal of dry ice. Fig. 29 illustrates the equipment at the -40 °C test temperature. Figure 29. Isolated equipment at -40 °C and liquid intermediate test agent. As the size of the CGHAZ is quite small, a study of this region is particularly difficult in real welds. Therefore, a thermal simulation was used to generate a relatively large region of CGHAZ, which allowed the notch to be reliably located in the correct microstructure. The steels were subjected to a welding thermal simulation. Thermal simulation test blanks were cut from the surface position of each plate, with the test piece axis transverse to the rolling direction, in T-L direction. Fig. 30 shows the test blanks, 8 x 17 mm in size. After the thermal simulation, these blanks were machined down to a 5 x 10 mm size appropriate for CTOD test pieces. The weld HAZ thermal simulations were performed on a Gleeble 3800 simulator, as the one shared in fig. 31, which is owned by the StPSPU. 75 Figure 30. Test pieces proportion to rolling direction. DIGITAL CONTROL SYSTEM OPERATION CHAMBER FORM MOBILE CONVERT UNIT OPERATION CHAMBER MECHANICAL CABIN MECHANICAL CAPIN Figure 31. The Gleeble 3800 machine used in StPSPU laboratory. Current, Voltage, Welding Gross Net Net heat Cooling A V speed, power, power, input, time t8/5, mm s W W J mm-1 s cycle Thermal Table 12. Welding parameters and cooling time. -1 1 230 25.6 4.533 5888 4710 1039 21.0 2 268 29.0 4.517 7772 6218 1376 36.5 3 258 30.6 3.713 7895 6316 1701 55.6 The thermal cycles were calculated depending on the welding conditions (Table 12). When calculating of the temperature field, the following assumptions were made: a point heat source on the plate surface moves along the x-axis with 76 constant speed v, the origin of coordinates is fixed to the source, the plate surfaces are heat impermeable, and the plate is infinitely wide and long. Then the steady state of the temperature field T(x,y,z) in the moving reference frame is expressed by the following formula: T ( x, y, z ) = T0 + q 2πλ exp(− vR vx ∞ 1 )∑ exp(− n ) 2a n = −∞ Rn 2a (7) Rn = [ x 2 + y 2 + ( z − 2ns ) 2 ]1/ 2 (8) where T0 is the ambient temperature (T0 = 20°C), q is the net power, λ is the heat conductivity (λ = 0.035 W mm-1 K-1), a is the thermal diffusivity (a = 7.0 mm2 s-1), s is the plate thickness (s = 8 mm). The vertical z - axis is directed through the plate thickness and changes from coordinate x to time t is made according to the equation: t = - x/v. Then the thermal cycle of any point y, z at any time t can be calculated: T ( y, z , t ) = T0 + q 2πλ exp( vR v 2t ∞ 1 )∑ exp(− n ) 2a n = −∞ Rn 2a Rn = [(vt ) 2 + y 2 + ( z − 2ns ) 2 ]1/ 2 (9) (10) This formula was used to calculate the thermal cycle of the point having peak temperature Tmax = 1350°C at the top surface (z = 0). Three cycles are shown in Fig. 32 a - c. a) Q= 1.0 kJ/mm. b) Q= 1.3 kJ/mm. Figure 32. Welding thermal cycles. 77 c) Q= 1.7 kJ/mm. The first heat input was 1.0 kJ/mm and was applied to an 8 mm thick plate. This involved heating to a peak temperature (Tp1) of 1350 °C at a rate of approximately 450 °C/s and holding the peak temperature for less than 2 s, followed by a cooling time from 1350 °C to 800 °C for 10 seconds, between 800 °C to 500 °C (∆t8/5) in 20s, and from 500 °C to ambient temperature in 40 seconds. The second heat input was 1.3 kJ/mm and was applied to an 8 mm thick plate. This involved heating to a peak temperature (Tp1) of 1350 °C at a rate of approximately 450 °C/s and holding at the peak temperature for less than 2 s, followed by a cooling time from 1350 °C to 800 °C in 15 seconds, between 800 °C to 500 °C (∆t8/5) in 35 s and from 500 °C to ambient temperature in 65 seconds. Finally, the third heat input was 1.7 kJ/mm and was applied to an 8 mm thick plate. This involved heating to a peak temperature (Tp1) of 1350 °C at a rate of approximately 450 °C/s and holding at the peak temperature for less than 2 s, followed by cooling time from 1350 °C to 800 °C in 20 seconds, between 800 °C to 500 °C (∆t8/5) in 55 s and from 500 °C to ambient temperature in 80 seconds. In simulation, which occurred in a Gleeble 3800 machine between watercooled copper made grip jaws, the non-standard Gleeble specimen has been heated and cooled, as seen in figs. 33 a and b. a) b) Figure 33. The 5x10 grips jaws (a) and non-standard Gleeble specimen (b). 78 CTOD test pieces were produced from the thermal simulated test blanks with a 2.5 mm deep through-thickness notch cut in the sample. The position of the notch was in the center of the etched HAZ. The notch orientation was such that the crack propagation direction was parallel to the plate rolling direction, as seen in fig. 26, T-L direction. A fatigue crack of 2.5 mm nominal depth was then grown into the specimen, giving a nominal a/W (overall crack depth/ specimen width) value of 0.5. The CTOD samples were then tested at -40 °C, following ASTM E 1290-02 standard, to produce impact toughness. The equation in standard ASTM E 1290-02 for CTOD value δ is given as: 1 𝛿 = 𝑚𝜎 × 𝑌 𝐾2 (1−𝑣2 ) 𝐸+𝜂𝐴𝜌 𝐵(𝑊−𝑎0) (1+(𝛼+𝑧) (11) 0.8𝑎0 +0.2𝑊 Where δ = CTOD –value ν = Poisson’s ratio E = Young’s modulus at the temperature of interest Ap = Area under the plot of load versus plastic component of clip gage opening displacement vp corresponding to vc, vu or vm (see fig. 28) B = Thickness of test specimen W = Width of test specimen a0 = Average length of crack α = reference distance (α=0 in the case of the SEB specimen) z = distance of knife edge measurement point from front face (notched surface) on SE(B) specimen 𝜎𝑌 = 𝜎𝑌𝑆 +𝜎𝑇𝑆 2 where σY = effective yield strength at the temperature of interest σYS = yield or 0.2 % offset yield strength at the temperature of interest σTS = tensile strength at the temperature of interest 79 (12) 𝑌𝑃 𝐾 = 𝐵√𝑊 (13) where K= stress intensity factor P = force corresponding to Pc, Pu or Pm (See fig. 34) Y= Stress Intensity coefficient 𝑌= 𝑎 𝑎 𝑎 𝑎 𝑎 2 6� 0 ×�1.99− 0 �1− 0 ��×�2.15−3.93 0 +2.7� 0 � � 𝑊 𝑊 𝑊 𝑎 𝑎 3 �1+2 0 �×��1− 0 � 𝑊 𝑊 𝑊 𝑊 (14) Constraint m in equation 11: 𝑎 where 𝑎 𝑚 = 1.221 + 0.793 𝑊0 + 2.751(𝑛) − 1.418 � 𝑊0 � (𝑛) 𝑛 = 1.724 − 6.098 𝑅 8.326 + 𝑅2 − 3.965 (15) (16) 𝑅3 where 𝜎 𝑅 = 𝜎𝑇𝑆 (17) 𝑌𝑆 𝑎 𝑊 function η in equation 11: 𝑎 𝑎 𝜂 = 3.785 − 3.101 𝑊0 + 2.018 � 𝑊0 � 80 2 (18) Figure 34. Types of Force versus Clip Gage Displacements Records (ASTM E 1290-02). 6.6.2. Compared microstructure examination An additional test on the microstructure was conducted using a high-resolution microscope. The test results illustrate the microstructure differences between QT and TMCP steels. These tests also show the HAZ microstructure and the zone difference. QT HSS steel, steel E, has been investigated as a typical QT HSS steel and steel C has been investigated as a typical TMCP HSS steel. This test was conducted at StPSPU. Specimen preparation included following techniques: sectioning, mounting, grinding, polishing, etching. Abrasive cut-off machine Buehler Powermet 3000 was used for sectioning. Mounting was performed on Buehler Simplimet 1000 mounting press in Epomet and Transoptic mounting resins. Buehler Phoenix 4000 was used for grinding and polishing of the specimens. Grinding was undertaken with a set of SiC abrasive papers starting out with the roughest (P180) and gradually introducing the finest (P400). Polishing materials were the diamond suspensions with particles ranging from 9 to 1 μm, alumina suspension 0.01 μm. 81 Revealing of microstructure was conducted by etching of the specimens in nital e.g. 4% solution of HNO3 in ethanol. The examination of microstructure was made using the light metallographic microscope LEICA DMI5000M with magnification up to x1000. Acquisition of images was performed by digital camera LEICA DFC320 attached to the microscope, which has 3 MPix image sensor. LEICA Application Suite software was used for enhancement and analysis of captured images. Image analysis provided accurate means for determining grain size according to ASTM E112. Stereomicroscope LEICA Mz12.5 was used for examination of macrostructures of welded joints. Hardness measurement was conducted on Vickers hardness tester Wilson Wolpert 452SVD according to ISO6507. Microhardness of single phases or tiny constituents were measured by microhardness tester Wilson Wolpert 402MVD with diamond pyramid indenter under load of 0.0025 N. 82 7. RESULTS AND DISCUSSION All of the tests carried out on these steels were made according to standard SFS-EN ISO 15614-1 welding procedures. Additionally, CTOD tests were conducted using standard ASTM E 1290-2. An in depth explanation of the results of these tests is covered in this section. 7.1. Visual test The researcher conducted a 100 % visual test on all of the welds according to standard SFS EN ISO 17637. This step excluded all premature negative effects that are possible in destructive testing. A proper visual test was conducted, which included feeling the entire weld. No defects, such as undercut, high reinforcement, root concavity, root defect etc., were noticed in the welded structures, which may be due to the MAG welding methods which produce high quality welds. 7.2. Macro photography After etching, a macro photograph was taken of each of the welded joints. The test specimens were prepared and etched in accordance with standard EN 1321 on one side to clearly reveal the fusion line, the HAZ, and the build up of the runs. Fig. 35 shows the location of the different zones in the macro image. Tables 13 through 20 show and explain macro photographs from all of the welded steels. The influence of heat input is noticeable from the pictures, as the HAZ zone is wider whit higher heat inputs. All steels were welded with two passes, except for steel G, which was welded with three passes. Additionally, the thickness of steel H was 12 mm, while all other steels were 8 mm thick. The fill up run and the final run heat-treats the root pass, all of which can be seen in macro photographs. 83 Heat input has the ability to effect the weld, with bigger heat inputs displaying greater degrees of mixing between the base and filler metals. Similar Basu and Raman (2002), this study reports, that different weld cooling rates lead to different weld microstructure features and inclusion parameters, which further leads to different values in mechanical properties. This is clearly seen from different tensile test results that were included in this study. The fusion (mixing) zone is seen from macro photographs, however, micro photographs display this zone in much finer detail. Second pass 8 mm HAZ zone Fusion line First pass Figure 35. A macro photograph shows the different zones of a welded joint. 84 Table 13. Macro photographs of steel A and comments. Steel name Macro sections Heat Input (kJ/mm) 1.0 A A 1.3 A 1.7 85 Comments Narrow HAZ with clear zones. Backing ring has developed near lack of side weld fusion in the root pass. Very good root. Both welds are good. Heat input in capping run has changed the microstructure in the CGHAZ. Wide HAZ area. Great heat input has changed the microstructure and also the root pass area. Too wide capping run consequent on 1.2 mm filler metal and great heat input. Table 14. Macro photographs of steel B and comments. Steel name Macro sections Heat Input (kJ/mm) 1.0 B B 1.3 B 1.7 86 Comments Very clear zones in narrow HAZ area. Excess weld metal in root pass. Smooth capping run which is good in dynamic action. Very good joint between gapping run and base metal. Clear HAZ area. Wide HAZ area. Good joint in both surface and root sides. Fusion line is not clear as a consequence of good mixing. Table 15. Macro photographs of steel C and comments. Steel name Macro sections Heat Input (kJ/mm) Comments C 1.0 Some misalignment in the welded plates. Smooth joints between weld and base plates. Also, the CGHAZ is clear in root pass. C 1.3 Very good weld. All HAZ areas are evident. This kind of weld has good properties. When welding HSSs, this kind of weld is intended. 1.7 Good weld, only root opening is greater than 1 mm. Gapping run is wide because the groove was too full after the root pass. C 8 mm Steel C Heat input 1.7 kJ/mm 87 Table 16. Macro photographs of steel D and comments. Steel name Macro sections Heat Input (kJ/mm) 1.0 D D D 88 Comments In QT steel the CGHAZ area is not as clear as steels A and C which are TMCP steels. HAZ area is clear. Small heat input lead up to clear fusion line. 1.3 Smooth joint in both sites, top of preparation and root. Undermatched filler weld is distinguishable from base material. 1.7 Wide HAZ area as a result of high heat input. The fusion line is not as clear as in steel D with a 1.0 kJ/mm heat input. This happens because greater mixing occurs at higher heat inputs, the effects of which can be seen in micro photographs. Table 17. Macro photographs of steel E and comments. Steel name Macro sections Heat Input (kJ/mm) 1.0 E E 1.3 E 1.7 89 Comments The formation of the backing ring is important to the shape of the root pass. In this weld, the filler metal has spread over the base metal. When the fusion line is not completely melted, a lack of side weld fusion can occur. Zones in the HAZ are distinguished. This QT steel has a clear CGHAZ using 1.3 kJ/mm heat input. Wide HAZ as a consequence of high heat input. When the fusion line is not completely melted, a lack of side weld fusion can occur. Table 18. Macro photographs of steel F and comments. Steel name Macro sections Heat Input (kJ/mm) F F F 1.0 Smooth joint between weld and base metal. Clear HAZ area where individual zones can be seen. 1.3 Gapping run has tempered all root pass. Very good smooth joint between weld and base material. Using this kind of weld, welded HSS structure will be durable. 1.7 90 Comments Wide HAZ area. Shape of root pass is a little high. Wide CGHAZ is distinguished from HAZ. Table 19. Macro photographs of steel G and comments. Steel name Macro sections Heat Input (kJ/mm) 1.0 G Comments 12 mm width QT steel welded using three passes. With this heat input, the HAZ area is narrow. This is intended when welding HSSs. Not a much wider HAZ area than in 1.0 kJ/mm heat input. G 1.3 G 1.7 91 Wider HAZ area where CGHAZ is well seen. The HAZ of steel G is quite narrow compared to other steels tested but can be explained because steel G is 12 mm thick instead of 8 mm of the other tested steel. Table 20. Macro photographs of steel H and comments. Steel name Macro sections Heat Input (kJ/mm) 1.0 H H 1.3 H 1.7 Comments Good looking welded structure. Narrow HAZ is good in welded HSS structure. Base material is not excessively tempered. Wider HAZ area than in 1.0 kJ/mm heat input but very good looking welded HSS structure. Too wide HAZ area but otherwise good welded high strength QT steel structure. 7.3. Micro photography Micro photographs with an aspect ratio of 1:500 were taken of all of the welds and their HAZs. The micro photograph in fig. 36 shows different zones where pictures were taken, including the weld, fusion line (partially melted zone), CGHAZ, FGHAZ, ICHAZ, SCHAZ and base material. Additionally, tables 21 -28 show and explain micro photographs from all of the welded steels. The heat input moves the place of the FGHAZ, ICHAZ and SCHAZ further from the fusion line; however, the microstructure is the same throughout the entirety of the welded structure. 92 The base microstructure in the steels was either bainite-martensite or ferritebainite. Disparities occurred in the phase structure of the steels depending on the manufacturer. The weld structure was a ferrite-perlite microstructure, which is a typical microstructure when the filler material is ESAB OK 12.51. Initial columnar grains formed by epitaxial growth were detected by the presence of grains of polygonal ferrite and Widmanstatten ferrite along the former grain boundaries. However, the main constituent is an acicular ferrite, forming a "wicker basket" structure. The first phase forming on prior austenite grain boundaries during cooling below the AC3 temperature is referred to as polygonal ferrite. At relatively low undercooling temperatures, Widmstatten ferrite formation occurs. The ferrite plates grow rapidly with a high aspect ratio, resulting in parallel arrays. Widmanstatten ferrite plates grow directly from a prior austenite grain boundary or from polygonal ferrite at the grain boundaries. Acicular ferrite is recognized as an intragranular nucleated morphology of ferrite in which there are multiple impingements between grains. The acicular ferrite nucleates on inclusions inside the prior austenite grains during the γ→α transformation. Provided there is a high density of inclusions, a fine interlocking structure is produced. The microstructure of the fusion line was an alloy of filler material and base metal, the two of which mixed together. This zone is in partially melted state. The microstructure of FL is mixed and contains bainite and polygonal ferrite, fig.37. Near fusion line hardness in QT base metal started to grow fast and in the CGHAZ, hardness had reached its highest point. The microstructure in the CGHAZ of QT steels is martensite-bainite. The highest concentration of martensite was observed in the CGHAZ, however bainite was formed as well. The FGHAZ is the zone after CGHAZ, in which the microstructure is smaller than the latter. Bainite phase is predominant with small part of martensite phase. The last variable phase is the ICHAZ which has a 93 phase structure similar to the base material. The microstructure of the ICHAZ can have some changes in its carbide structure which can decrease its yield Line of photography 1 2 3 4 6 5 8 mm Weld Metal 2 mm strength compared to the base material. Fusion CGHAZ FGHAZ ICHAZ Line Base Metal Figure 36. Semantic photograph from welded structure showing the areas where the micro photographs had been taken. Figure 37. Optical microstructure of the fusion line of QT HSS E. 94 CGHAZ FUSION LINE Liquid base metal and weld metal has mixed together. Base metal alloying elements, like Nb, V, Ti, etc. have mixed with liquid filler metal. The strength of the weld has grown because of that. Microstructure of the base metal near fusion zone is bainite although ferritepearlite can occur. WELD METAL Solidified weld material is ferrite structure. Alpha ferrite, Windmannstätt ferrite and acicular ferrite occurs in the ferrite micro structures. Epitaxial crystal growth is well displayed (Lancaster 1980). In the CGHAZ zone, austenite had time to grow large. Cooling time has been fast and the microstructure after solidification is bainite through some pearliteferrite can occur. Size of grains has grown, but depends on t8/5 time (heat input). Inherent austenite grain size is seen in this figure. Width of the CGHAZ depends on heat input. A CGHAZ that is too wide can cause the welded structure to break under loading. Three different heat input 1.0, 1.3 and 1.7 kJ/mm had different CGHAZ widths and 1.7 kJ/mm had the widest CGHAZ. It is noticed that in this TMCP steel hardness does not grow in spite of fast cooling in the CGHAZ zone, because of the low carbon content of the base metal. 3 2 1 95 FGHAZ zone has austenitized during welding. Austenitizing had changed the micro structure and some phases are larger than in the base material. The main structure is pearliteferrite. FGHAZ area has the same strength than the base material or more. FGHAZ 4 In ICHAZ some carbides and nitrides had dissolved. Size of microstructure is same as base material. Main microstructure is bainite and ferrite. Difference of microstructure between ICHAZ and SCHAZ is difficult to see. ICHAZ and SCGAZ 5 Table 21. Micro photographs of welded TMCP steel A and comments. Aspect ratio of 1:500. Base microstructure of TMCP A steel was a bainite and ferrite microstructure. Microstructure was very small and homogeneous. Rolling direction has not any effect on steel A. BASE METAL 6 FUSION LINE A very clear fusion line is observed. Weld metal microstructure is the same as the weld. Base metal was molten in the fusion line. Base material mixes with melted filler material. Mixing can be clearly seen in the fusion line. Solidified base material grains have been directed towards the base metal. Microstructure is martensite-bainite near the fusion line of the base metal. Solidified weld material is ferrite structure. Alpha ferrite, Windmannstätt ferrite and acicular ferrite occur in the ferrite micro structures. Epitaxial crystal growth is well displayed (Lancaster 1980). 2 WELD METAL 1 Main microstructure is martensite and bainite in the CGHAZ. In Kaputska et al. (2008), the same microstructure was observed. Microstructure in CGHAZ has grown. Size of grains depends on t8/5 time (heat input). Grain size was largest when heat input was 1.7 kJ/mm. In all heat input 1.0, 1.3 and 1.7 kJ/mm this zone was most brittle in the HAZ. Width of CGHAZ is wider when heat input is greater and t8/5 time is longer. In literature the width of the CGHAZ area should be maximum 1/3 of thickness of the base metal. CGHAZ 3 96 Hardenability declines and softening takes place in the FGHAZ due to the miniaturization of the former austenite (Hamada 2003). Hamada (2003) concluded that toughness is generally high in the FGHAZ. Size of grains is mainly small, but some grain growth can occur. This HAZ of QT HSSs does not have any problem under loading. Strength and toughness are the same or better than in the base metal. FGHAZ 4 The agglomeration of spheroidized cementite particles at grain boundaries of SCHAZ is more noticeable than in ICHAZ. Concentration of the former austenite occurs in the ICHAZ and this hardened phase becomes a material ‘notch’ and the toughness deteriorates (Hamada 2003). In the ICHAZ, the base metal has tempered. Some carbides are sphere sometimes making the ICHAZ weaker than the base metal. ICHAZ and SCHAZ 5 Table 22. Micro photographs of welded QT steel B and comments. Aspect ratio is 1:500. Microstructure of steel B was tempered martensite and bainite. This quenched and tempered microstructure primarily consists of fine-lath martensite and significant amounts of coarse martensite (Moon et al 2000 according to Fonda et al. 1994). Steel B was QT HSS and this microstructure is typical to QT steel. The size of grains is small and texture is homogenous. This kind of microstructure gives good strength and toughness to steel. BASE METAL 6 2 FUSION LINE Melted base metal and liquid weld metal have mixed together. Base metal alloying elements, such as Nb, V, Ti, etc. have mixed with liquid filler metal, causing the weld’s strength to growth. The microstructure of the base metal near the fusion zone is bainite although ferrite- pearlite can occur. Great heat inputs near the fusion line have made large grains in the base metal. The inherent austenite grain size has grown near the fusion line because this zone has been the longest over the Ac3 point. 1 WELD METAL Solidified weld material is ferrite structure. Alpha ferrite, Windmannstätt ferrite and acicular ferrite occurs in the ferrite micro structures. Epitaxial crystal growth is well displayed (Lancaster 1980). It is noticed that in this TMCP steel C hardness does not grow in spite of fast cooling in the CGHAZ because of the low carbon content of the base metal. Width of the CGHAZ depends on heat input. A CGHAZ that is too wide can cause the welded structure to break under loading. Three different heat input 1.0, 1.3 and 1.7 kJ/mm had different CGHAZ widths and 1.7 kJ/mm had the widest CGHAZ. In CGHAZ austenite has time to grow large. Cooling time has been fast and microstructure after the solidification is bainite though some pearlite-ferrite can occur. The size of the grains has grown and the inherent austenite grain size can be seen in this figure. CGHAZ 3 97 FGHAZ zone has austenitized during welding. Austenitizing had changed micro structure and some phases are larger than in base material. Main structure is pearlite-ferrite. FGHAZ area has same strength than base material or more. FGHAZ 4 In ICHAZ some carbides and nitrides had dissolved. Size of microstructure is same as base material. Main microstructure is bainite and ferrite. Difference of microstructure between ICHAZ and SCHAZ is difficult to see. ICHAZ and SCHAZ 5 Table 23. Micro photographs of welded TMCP steel C and comments. Aspect ratio is 1:500. PCM of steel C was bigger than steel A. The carbon content was same but Mn and Nb contents were bigger. As results of these factors, steel C has greater mechanical features. Base microstructure of TMCP C steel was a bainite and ferrite microstructure. Microstructure was very small and homogeneous. Rolling direction has not had any effect on steel C. BASE METAL 6 A very clear fusion line is observed. Weld metal microstructure is the same as the weld. Base metal was molten in the fusion line. Base material mixes with melted filler material. Liquid metal has solidified towards the weld centre, along the temperature gradient. Solidified weld material is ferrite structure. Alpha ferrite, Windmannstätt ferrite and acicular ferrite occurs in the ferrite micro structures. Epitaxial crystal growth is well displayed (Lancaster 1980). Microstructure is martensite-bainite near the fusion line of the base metal. FUSION LINE WELD METAL This weld metal is undermatched with base metal. 2 1 FGHAZ 4 The main microstructure in this zone is martensite and bainite. Hardenability declines and softening takes place in the FGHAZ due to the miniaturization of the former austenite (Hamada 2003). Hamada (2003) concluded that toughness is generally high in the FGHAZ. Size of grains is mainly small, but some grain growth can occur. This HAZ of QT HSSs does not have any problem under loading. Strength and toughness are the same or better than in the base metal. 98 The main microstructure is martensite and bainite. Microstructure in CGHAZ has grown. Size of grains depends on t8/5 time (heat input). Grain size was largest when heat input was 1.7 kJ/mm. In all heat input 1.0, 1.3 and 1.7 kJ/mm this zone was most brittle in the HAZ. Width of CGHAZ is wider when heat input is greater and t8/5 time is longer. In literature the width of the CGHAZ area should be maximum 1/3 of thickness of the base metal. CGHAZ 3 The agglomeration of spheroidized cementite particles at grain boundaries of SCHAZ is more noticeable than in ICHAZ. Main microstructure is tempered martensite and bainite with cementite particles. Concentration of the former austenite occurs in the ICHAZ and this hardened phase becomes a material ‘notch’ and the toughness deteriorates (Hamada 2003). In the ICHAZ, the base metal has tempered. Some carbides are sphere sometimes making the ICHAZ weaker than the base metal. ICHAZ and SCHAZ 5 Table 24. Micro photographs of welded QT steel D and comments. Aspect ratio is 1:500. Microstructure of steel D was tempered martensite and bainite. Steel D was QT HSS and this microstructure is typical to QT steel. The size of the grains is small and texture is homogenous. This kind of microstructure gives good strength and toughness to steel. BASE METAL 6 A very clear fusion line is observed. Weld metal microstructure is the same as the weld. Base metal was molten in the fusion line. Base material mixes with melted filler material. Liquid metal has solidified towards the weld centre, along the temperature gradient. Solidified weld material is ferrite structure. Alpha ferrite, Windmannstätt ferrite and acicular ferrite occurs in the ferrite micro structures. Epitaxial crystal growth is well displayed (Lancaster 1980). Microstructure is martensite-bainite near the fusion line of the base metal. FUSION LINE WELD METAL This weld metal is undermatched with base metal. 2 1 The main microstructure in this zone is martensite and bainite. Microstructure in CGHAZ has grown. Size of grains depends on t8/5 time (heat input). Grain size was largest when heat input was 1.7 kJ/mm. In all heat input 1.0, 1.3 and 1.7 kJ/mm this zone was most brittle in the HAZ. Width of CGHAZ is wider when heat input is greater and t8/5 time is longer. In literature the width of the CGHAZ area should be maximum 1/3 of thickness of the base metal. CGHAZ 3 99 Main microstructure is martensite and bainite in that zone. Hardenability declines and softening takes place in the FGHAZ due to the miniaturization of the former austenite (Hamada 2003). Hamada (2003) concluded that toughness is generally high in the FGHAZ. Size of grains is mainly small, but some grain growth can occur. This HAZ of QT HSSs does not have any problem under loading. Strength and toughness are the same or better than in the base metal. FGHAZ 4 The agglomeration of spheroidized cementite particles at grain boundaries of SCHAZ is more noticeable than in ICHAZ. Main microstructure is tempered martensite and bainite with cementite particles. Concentration of austenite formers occurs in ICHAZ zone and this hardened phase becomes a material ‘notch’ and the toughness deteriorates (Hamada 2003). In ICHAZ zone base metal has tempered. Some carbides are sphered and it makes ICHAZ zone sometimes weaker than base metal. ICHAZ and SCHAZ 5 Table 25. Micro photographs of welded QT steel E and comments. Aspect ratio is 1:500. Microstructure of steel E was tempered martensite and bainite. Steel E was QT HSS and this microstructure is typical to QT steel. The size of the grains is small and the texture is homogenous. This kind of microstructure gives good strength and toughness to steel. BASE METAL 6 A very clear fusion line is observed. Weld metal microstructure is the same as the weld. Base metal was molten in the fusion line. Base material mixes with melted filler material. Liquid metal has solidified towards the weld centre, along the temperature gradient. Solidified weld material is ferrite structure. Alpha ferrite, Windmannstätt ferrite and acicular ferrite occurs in the ferrite micro structures. Epitaxial crystal growth is well displayed (Lancaster 1980). Microstructure is martensite-bainite near the fusion line of the base metal. FUSION LINE WELD METAL This weld metal is undermatched with base metal. 2 1 The main microstructure in this zone is martensite and bainite. FGHAZ 4 Main microstructure is martensite and bainite in that zone. Hardenability declines and softening takes place in the FGHAZ due to the miniaturization of the former austenite (Hamada 2003). Hamada (2003) concluded that toughness is generally high in the FGHAZ. Size of grains is mainly small, but some grain growth can occur. This HAZ of QT HSSs does not have any problem under loading. Strength and toughness are the same or better than in the base metal. 100 Microstructure in CGHAZ has grown. Size of grains depends on t8/5 time (heat input). Grain size was largest when heat input was 1.7 kJ/mm. In all heat input 1.0, 1.3 and 1.7 kJ/mm this zone was most brittle in the HAZ. Width of CGHAZ is wider when heat input is greater and t8/5 time is longer. In literature the width of the CGHAZ area should be maximum 1/3 of thickness of the base metal. CGHAZ 3 The agglomeration of spheroidized cementite particles at grain boundaries of SCHAZ is more noticeable than in ICHAZ. Main microstructure is tempered martensite and bainite with cementite particles. Concentration of austenite formers occurs in ICHAZ zone and this hardened phase becomes a material ‘notch’ and the toughness deteriorates (Hamada 2003). In ICHAZ zone base metal has tempered. Some carbides are sphered and it makes ICHAZ zone sometimes weaker than base metal. ICHAZ and SCHAZ 5 Table 26. Micro photographs of welded QT steel F and comments. Aspect ratio is 1:500. Microstructure of steel F was tempered martensite and bainite. Steel F was QT HSS and this microstructure is typical to QT steel. The size of the grains was small and the texture is homogenous. This kind of microstructure gives good strength and toughness to steel. BASE METAL 6 A very clear fusion line is observed. Weld metal microstructure is the same as the weld. Base metal was molten in the fusion line. Base material mixes with melted filler material. Liquid metal has solidified towards the weld centre, along the temperature gradient. Solidified weld material is ferrite structure. Alpha ferrite, Windmannstätt ferrite and acicular ferrite occurs in the ferrite micro structures. Epitaxial crystal growth is well displayed (Lancaster 1980). Microstructure is martensite-bainite near the fusion line of the base metal. FUSION LINE WELD METAL This weld metal is undermatched with base metal. 2 1 The main microstructure in this zone is martensite and bainite. FGHAZ 4 Main microstructure is martensite and bainite in that zone. Hardenability declines and softening takes place in the FGHAZ due to the miniaturization of the former austenite (Hamada 2003). Hamada (2003) concluded that toughness is generally high in the FGHAZ. Size of grains is mainly small, but some grain growth can occur. This HAZ of QT HSSs does not have any problem under loading. Strength and toughness are the same or better than in the base metal. 101 Microstructure in CGHAZ has grown. Size of grains depends on t8/5 time (heat input). Grain size was largest when heat input was 1.7 kJ/mm. In all heat input 1.0, 1.3 and 1.7 kJ/mm this zone was most brittle in the HAZ. Width of CGHAZ is wider when heat input is greater and t8/5 time is longer. In literature the width of the CGHAZ area should be maximum 1/3 of thickness of the base metal. CGHAZ 3 The agglomeration of spheroidized cementite particles at grain boundaries of SCHAZ is more noticeable than in ICHAZ. Main microstructure is tempered martensite and bainite with cementite particles. Concentration of austenite formers occurs in ICHAZ zone and this hardened phase becomes a material ‘notch’ and the toughness deteriorates (Hamada 2003). In ICHAZ zone base metal has tempered. Some carbides are sphered and it makes ICHAZ zone sometimes weaker than base metal. ICHAZ and SCHAZ 5 Table 27. Micro photographs of welded QT steel G and comments. Aspect ratio is 1:500. Microstructure of steel G was tempered martensite and bainite. Steel G was QT HSS and this microstructure is typical to QT steel. The size of the grains was small and the texture is homogenous. This kind of microstructure gives good strength and toughness to steel. BASE METAL 6 A very clear fusion line is observed. Weld metal microstructure is the same as the weld. Base metal was molten in the fusion line. Base material mixes with melted filler material. Liquid metal has solidified towards the weld centre, along the temperature gradient. Solidified weld material is ferrite structure. Alpha ferrite, Windmannstätt ferrite and acicular ferrite occurs in the ferrite micro structures. Epitaxial crystal growth is well displayed (Lancaster 1980). Microstructure is martensite-bainite near the fusion line of the base metal. FUSION LINE WELD METAL This weld metal is undermatched with base metal. 2 1 The main microstructure in this zone is martensite and bainite. FGHAZ 4 Main microstructure is martensite and bainite in that zone. Hardenability declines and softening takes place in the FGHAZ due to the miniaturization of the former austenite (Hamada 2003). Hamada (2003) concluded that toughness is generally high in the FGHAZ. Size of grains is mainly small, but some grain growth can occur. This HAZ of QT HSSs does not have any problem under loading. Strength and toughness are the same or better than in the base metal. 102 Microstructure in CGHAZ has grown. Size of grains depends on t8/5 time (heat input). Grain size was largest when heat input was 1.7 kJ/mm. In all heat input 1.0, 1.3 and 1.7 kJ/mm this zone was most brittle in the HAZ. Width of CGHAZ is wider when heat input is greater and t8/5 time is longer. In literature the width of the CGHAZ area should be maximum 1/3 of thickness of the base metal. CGHAZ 3 The agglomeration of spheroidized cementite particles at grain boundaries of SCHAZ is more noticeable than in ICHAZ. Main microstructure is tempered martensite and bainite with cementite particles. Concentration of austenite formers occurs in ICHAZ zone and this hardened phase becomes a material ‘notch’ and the toughness deteriorates (Hamada 2003). In ICHAZ zone base metal has tempered. Some carbides are sphered and it makes ICHAZ zone sometimes weaker than base metal. ICHAZ and SCHAZ 5 Table 28. Micro photographs of welded QT steel H and comments of it. Aspect ratio is 1:500. Microstructure of steel G was tempered martensite and bainite. Steel G was QT HSS and this microstructure is typical to QT steel. The size of the grains was small and the texture is homogenous. This kind of microstructure gives good strength and toughness to steel. BASE METAL 6 7.4. Radiographic tests Additionally, radiographic tests of standard SFS-EN 1435 were performed to all welds to examine porosity, cracks and inclusions. After the radiographic tests were performed, some extended gas pores were noticed in the welds, however the quantity and size of the gas pores were not a significant factor in the quality of the welds. As the gas pores are of an insignificant size and density, they have probably developed from the welding gun being held at 90° angle to the steel, resulting in paths for the gases to go away after each pass. Figs 38 and 39 show some typical samples of gas pores were found in the welds. As can be seen in these figures, the gas pores are round and are not collecting in groups. GAS PORES GAS PORES Figure 38. Sample figure of gas Figure 39. Sample figure of gas pores in steel A (heat input 1.3 pores in steel B (heat input 1.7 kJ/mm). kJ/mm). 7.5. Surface crack detection All visual surface crack detection tests were made using penetrant testing in accordance with testing standards SFS-EN 571-1 and SFS-EN ISO 23277. Any crack detections were observed and the size of the undercut was within in the limits of the standard as the welding had been conducted in a laboratory environment. 103 7.6. Transverse tensile test Two transverse tensile tests were performed on all welds in accordance with standard SFS-EN ISO 4136. The results of these tests are in figs 40 though 43. Fig. 40 represents all of the tensile test results that were collected and helps to illustrate that the tensile strength of the welded structure is lower when the heat input is bigger. The tensile strength of the filler material was 560 MPa, and the tensile strength of the base material, corresponding to its material standard, was between 700 and 770 MPa. All of the material certificates have actual values of tensile strength. When undermatched filler metal was used during welding, the real tensile strength of the undermatched welded structure was more than the tensile strength of filler material as a consequence of penetration and mixing between the base and filler materials. The tensile strength of the welded structure is near the tensile strength of the base material required by that steel’s standard. All of this can be seen in figs 40 through 42 and additionally all of the tested welded structures broke at their welding points as a result of the tensile test. MPa Tensile test values of structure 780 760 740 720 700 680 660 640 620 600 580 Heat input 1.0 kJ/mm A B Heat input 1.3 kJ/mm C STEELS D E Heat input 1.7 kJ/mm F G H Figure 40. Tensile strengths of welded joint made of different steels using three heat input. 104 TENSILE STRENGTH OF WELDED STRUCTURE 800 780 MPa 764 751 TENSILE STRENGTH MPa 750 735 732 739 721 709 700 690 680 672 683 678 667 666 650 629 600 Tensile strength of filler material 560 MPa 550 500 STEELS Heat input 1.0 kJ/mm Figure 41. Tensile strength of various steels using constant heat input 1.0 kJ/mm. TENSILE STRENGTH OF WELDED STRUCTURE 800 780 MPa 765 759 748 750 727 TENSILE STRENGTH MPa 715 704 700 688 682 708 685 671 661 650 645 651 597 600 Tensile strength of filler material 560 MPa 550 500 STEELS Heat input 1.3 kJ/mm Figure 42. Tensile strength of various steels using constant heat input 1.3 kJ/mm. 105 TENSILE STRENGTH OF WELDED STRUCTURE 800 780 MPa 750 739 718 TENSILE STRENGTH MPa 716 706 700 685 691 679 671 669 660 650 647 642 639 632 631 600 Tensile strength of filler material 560 MPa 550 500 STEELS Heat input 1.7 kJ/mm Figure 43. Tensile strength of various steels using constant heat input 1.7 kJ/mm. The mismatch level between filler metal and parent metal was 0.72, which is lower than the recommendations of many researchers (Toyota 1986, Satoh & et al. 1975). Within such as low mismatch value, it is clear that the weld is the weakest place in structure, especially when compared to the strengths of filler and base materials. Tensile strength values change when penetration and mixing between filler and base material occurs and figures 41 through 43 show that the strength values of the base material are higher than the filler material. A typical example of a broken tensile test bar is shown in figs 44 a, b, c and d. The fracture occurs in the weld at the point of reduction area, the failure of which arises in the HAZ and continues into the weld. 106 a)side picture c) root side b) face side d) broken tensile test bar Figure 44. Tensile test bar. Steel A has standard tensile test value 700 MPa and using the lowest heat input (1.0 kJ/mm), the values obtained from steel A were near the tensile strength of the base steel. The same happened in steel H when heat input was 1.0 and 1.3 kJ/mm, and nearly same situation occurred in steel D. In these situations, the tensile strength of welded structure was 4 % lower than the tensile strength of base material. The standard tensile strength of all steels with the exception of steel A was 780 MPa. In all welded structures, the failure started from the weld or the HAZ, however, in some instants the failure started from the fusion line between weld and HAZ. This happens because of the low yield strength of filler material but also because of possible deformation in the weld. Nearly all of the tensile test pieces failed starting at the fusion line and only few of them broke in the HAZ. When the failure began in the fusion line or the HAZ, the direction of the break was towards the weld at a traditional 45° angle. 107 There will always be differences between welded structures regardless of how the steel was welded or what the heat input of welded structure was. All of heat inputs, steel H had the best tensile strength values, closely followed by steel G. In all cases, tensile strength values were the lowest when the heat input was 1.7 kJ/mm, however steel B’s lowest tensile test occurred when the heat input was 1.0 kJ/mm. It is important to consider that tensile test values do not account for all feature of the welded structure, and this is why other tests were conducted within the scope of this research to determine other mechanical properties. Tables 29, 30 and 31 present tensile test values which are used in different heat inputs in welding. Regardless of heat input, all tensile test values are higher than the tensile test of the filler material, which was 560 MPa. The tensile test values of the base material was around 780 MPa or more (steel A had minimum tensile test value 700 MPa). In all steels, the real tensile test value was more than in manufactorer’s procedure. Heat input has lowering effect to tensile strength of structure. Manufactory method doesn’t effect to tensile strength. Also, TMCP and QT steels behaved equally when using different heat input in welding. In all structures, the elongation at the break was considerably smaller than the base material elongation. In 690 MPa class HSSs, the standards stipulate that the minimum A5 should be 15%. However, in this research the values for the elongation at the break were only half of the base material values. These discrepancies can be accounted for by the differences in elongation at the break between the base and filler material as seen in tables 29, 30 and 31. The gauge length was 85 mm (standard SFS-EN ISO 6892-1) while the length of the weld was around 25 mm. The yield strength of the filler material was 470 MPa while the yield strength of the base material was 690 MPa. As there was such a large difference between these yield strength values, most of the yielding occurred in the weld. As these steels were constructed under varying manufacturing methods, their resulting yield strength and elongation break 108 values differed from one another. These differences caused variations between elongation break values of the welded structures when using the same heat input. The amount of penetration and dilution that occurred between the base and filler materials led to a better tensile strength in the welded structure than in the filler material. There is a correlation between the tensile strength and elongation break value of HSS, where larger real tensile strength leads to smaller elongation break values. When dilution happens between the base material and the weld, alloy elements can mix together. Some alloys such as Nb mixes to the weld and increases the properties of the welded structure. The Metal Handbook (1990) explains that the yield strength of the carbon steel increases with small additions of Nb. The yield strength of carbon steel can increase from 490 MPa to 700 MPa when the addition of Nb is 0.02 %. Using fillet welds, it is possible to increase the size of the weld (effective throat thickness) which leads to a greater tensile strength in the welded structure. Aside from increasing the tensile strength, this method also has some negative side effects including a longer welding time, higher cost and decreasing productivity. As opposed to fillet welds, butt welds are limited and increasing the weld size is not possible. When using undermatched filler material, the welded structure will not have a strength matching its base material. 109 Table 29. Comparing the tensile strength and elongation at break of base material to the welded structure when heat input was 1.0 kJ/mm. Red font corresponds to the highest value while green font corresponds to the lowest. TEST SPECIMEN TESTED TENSILE TRENGTH MPa % BIGGER THAN FILLER MATERIAL A1 A2 B1 B2 C1 C2 D1 D2 E1 E2 F1 F2 G2 H1 H2 690 680 629 672 667 732 735 739 666 678 721 709 683 751 764 23.2 21.4 12.3 20.0 19.1 30.7 31.3 32.0 18.9 21.1 28.8 26.6 22.0 34.1 36.4 TENSILE STRENGTH OF BASE MATERIAL (from material certificate) MPa 769 844 821 852 835 798 879 865 % LOWER THAN BASE MATERIAL 10.3 11.6 ELONGATION AT BREAK A₅ % 11.2 22.3 13.2 11.7 8.5 7.9 4.8 5.2 7.0 5.4 6.4 6.4 4.3 4.6 5.5 5.4 7.0 6.3 6.3 15.4 6.1 25.5 20.4 18.8 10.8 13.7 13.3 20.2 18.8 9.6 MEAN VALUE 8.2 5.0 6.2 6.4 4.4 5.5 7.0 6.3 MEAN VALUE 25.2 Table 30. Comparing the tensile strength and elongation at break of base material to the welded structure when heat input was 1.3 kJ/mm. Red font corresponds to the highest value while green font corresponds to the lowest. TEST SPECIMEN A1 A2 B1 B2 C1 C2 D1 D2 E1 E2 F1 F2 G2 H1 H2 TESTED TENSILE STRENGTH MPa 645 651 715 688 682 685 748 727 661 671 704 708 597 765 759 MEAN VALUE % BIGGER THAN FILLER MATERIAL 15.2 16.3 27.7 22.9 21.8 22.3 33.6 29.8 18.0 19.8 25.7 26.4 6.6 36.6 35.5 TENSILE STRENGTH OF BASE MATERIAL (from material certificate) MPa 769 844 821 852 835 798 879 865 23.9 110 % LOWER THAN BASE MATERIAL 16.1 15.3 15.3 18.5 16.9 16.6 12.2 14.7 20.8 19.6 11.8 ELONGATION AT BREAK A₅ % 11.6 12.3 9.4 9.3 5.8 6.4 9.6 9.3 5.6 5.7 5.0 4.6 8.1 7.8 7.0 5.9 6.1 16.3 7.1 11.3 32.1 MEAN VALUE 9.3 6.1 9.4 5.7 4.8 7.9 7.0 6.0 Table 31. Comparing the tensile strength and elongation at break of base material to the welded structure when heat input was 1.7 kJ/mm. Red font corresponds to the highest value while green font corresponds to the lowest. TEST SPECIMEN A1 A2 B1 B2 C1 C2 D1 D2 E1 E2 F1 F2 G1 G2 H1 H2 TESTED TENSILE STRENGTH MPa % BIGGER THAN FILLER MATERIAL 660 647 685 691 17.9 15.5 22.3 23.4 671 706 716 642 639 669 679 632 631 739 718 MEAN VALUE 19.8 26.1 27.9 14.6 14.1 19.5 21.3 12.9 12.7 32.0 28.2 TENSILE STRENGTH OF BASE MATERIAL (from material certificate) MPa 769 844 821 852 835 798 % LOWER THAN BASE MATERIAL ELONGATION AT BREAK A₅ % 14.2 15.9 18.8 18.1 9.5 8.5 6.1 7.5 7.9 9.4 5.9 6.9 6.1 6.1 6.0 6.3 6.6 5.3 6.8 7.1 18.3 14.0 16.0 23.1 23.5 16.2 14.9 28.1 879 28.2 865 14.6 17.0 20.5 MEAN VALUE 9.0 6.8 8.6 6.4 6.1 6.1 6.0 7.0 7.0 This tensile test has proven that when heat input is bigger and consequence of that width of HAZ is consequently wider, the tensile properties of the welded structure are weaker than the base material. In the tensile tests, the weakest welded structure had the highest heat input. Rodriques et al. (2004a) came to the same conclusion in their study when they looked at matched and undermatched filler metal situations and determined that the strength of the joint is strongly depend on the HAZ dimension. It is therefore of utmost importance to use proper welding parameters when welding HSSs regardless of the filler material. 111 7.7. Transverse bend test Overall, four bend tests were carried out to determine the occurrence of cracks and unmelted fusion line among other issues. Two of these tests were carried out on the root of the groove, while the other two were carried out on the top of the groove. All bend tests were made according to standard SFS-EN ISO 5173. The transverse bend tests will show faults in welded structure, such as defective penetration or low mixture levels between base and filler material. The transverse bend tests that were done on these HSSs with undermatched filler material were much more demanding than normal transverse bend tests. The discrepancy between the tests occurs because the filler material has a lower yield strength than base material. In these tests, the first part to be bent was the welded structure and the base material. In the end of these tests, the weld yielded more than the base material and the bending angle was bigger in the weld than in the structure, as seen in figs 45 and 46. If the welded structure passes this bend test, the weld can then be considered of acceptable quality. Figure 45. Example from transverse bending test face side. 112 Figure 46. Example from transverse bending test root side. The results of the transverse bending tests are in table 32, where OK means that the weld passed the bending test. Of all the transverse bending tests, steel G got the worst results which can be explained through a number of factors. First of all, a thickness of 12 mm, the heat flow from the fusion line was faster than in other steels. During the solidification of the molten weld pool, the porosity could increase, and these porous areas will be the first to crack during bending tests. Additionally, dilution in fusion line could be too low for the same reasons. In steel G, all the root passes failed in the transverse bending test. This can potentially be explained by the fact that the cooling time of the root pass without being preheated is shorter in 12 mm thick plates than in 8 mm thick plates. If there are significant thickness discrepancies, it would be possible to use a three dimensional equation, however, the differences between 8 and 12 mm thickness (d) in equation 16 is 2.25 times (d2 in the equation). In addition to heat input, cooling time is another important component in the welding process. Cooling time is dependent factor that depends on heat input, but also plate thickness, workpiece geometry, material properties and more. The cooling time can be calculated, using equation 9. Equation 9 allows the cooling time to be calculated with allowance for thicker plate thickness. During the course of this research, 8 mm and 12 mm thick plates of steel displayed large differences in cooling time (fig. 46-1). Additionally, the cooling time of the root pass of QT HSS G was short, 7 s. A 113 short cooling time can lead to brittle martensite microstructure, which also has small ductile value. This is why the root pass of QT HSS G broke in the bending test. Figure 46-1. Cooling time t8/5 vs. plate thickness. Welding conditions are presented in Table 12. 114 Table 32. Results of tranverse bending tests. OK means acceptable test. MATERIAL A B C D E F G H WELD root 1 root 2 surface 1 surface 2 root 1 root 2 surface 1 surface 2 root 1 root 2 surface 1 surface 2 root 1 root 2 surface 1 surface 2 root 1 root 2 surface 1 surface 2 root 1 root 2 surface 1 surface 2 root 1 root 2 surface 1 surface 2 root 1 root 2 surface 1 surface 2 HEAT INPUT 1.0 kJ/mm OK OK OK OK OK OK OK OK OK OK OK OK OK OK OK OK OK OK OK OK OK OK OK Broken 51° Broken 57° Broken 26° OK broken 75° OK OK OK HEAT INPUT 1.3 kJ/mm OK OK OK OK OK OK OK OK OK OK OK OK OK OK OK OK OK OK OK OK OK OK OK Broken 39° Broken 28° OK OK OK OK OK OK HEAT INPUT 1.7 kJ/mm OK OK OK OK OK OK OK OK OK OK Broken 90° OK OK OK OK OK OK OK OK OK OK OK Broken 18° Broken 18° OK OK OK OK OK OK 7.8. Impact test Two sets of impact tests were conducted, each set including three test pieces. Standard SFS-EN ISO 148-1 was used and each piece was 5 x 10 x 55 mm and tested at a temperature -40 °C. A 2 mm V notch was cut into each test piece and its correct placement was ensured by etching the notch before machining. The place of Charpy-V impact test is in fig. 47, which figure clarifies the structure being tested. Dependent on welding heat input, the shape of weld 115 will curve more horizontally and it leads to different HAZs under the V-groove. As shown in fig. 47, the test area of Charpy-V test can include some weld metal, CGHAZ, FGHAZ, ICHAZ, SCHAZ and some base metal. Between first and second HAZ is the ICCGHAZ which earlier research (Liu at al. 2007, Hamada 2003, Li et al. 2001, Lambert et al. 2000, Matsuda et al. 1995, Lee et al. 1993) has shown to be the most fracture area in the HAZ. The brittle area of ICCGHAZ is small, but in some Charpy-V tests it can be under the test notch. Milled Charpy-V groove Second pass 5 mm First pass 8 mm 2 HAZ zone Fusion line Figure 47. Place of Charpy-V groove in test pieces. In earlier studies (Wang et al. 2003, Juan et al. 2003) it was noticed that lower toughness values occur because of a wide HAZ. The lowest toughness values were in CGHAZ and if the HAZ is wide all zones will be wider and then the Charpy-V test place is more in CGHAZ and fusion line. In this present research, the same results have been observed. The overall numbers of tests were small because of the testing standards, and some exceptional results are the consequence of statistical dispersion. Overall, the results of the impact tests were ambiguous. The test results from weld area, as seen in fig. 48 and table 33 were acceptable and these results show that undermatching weld metal has good impact ductility. This might be because the impact value was limited to 18 J for test bar 5 x 10 x 55 mm piece. As seen in table 33 steels A, E, F and G have a few results under 18 J, however 116 the vast majority of them are close to 18 J (16 J - 17 J) and they can be considered acceptable. Fig. 49 and table 34 display HAZ impact results. In earlier study by Shi et al. (1998) it was concluded that the lower the weld strength mismatching, the higher the fracture toughness of the HAZ. In this study, the mismatching value was very low at 0.72. There was not a great deal of consistency in HAZ impact test results, as some steels have good values for all three heat inputs while other steels had very low values. In fig. 49 the impact test values show great divergence between different HSSs. Additionally, the test values in fig. 49 are very low. Cells highlighted in yellow in table 34 indicate that the values are under standard recommendations. For example, steel H had a value 4 J twice when heat input was 1.7 kJ/mm and had poor values ranging from 8 - 13 J at 1.0 and 1.3 kJ/mm as well. Steels A, B, D and F also exhibited low impact test values, however there is no consistency in the results according to heat input. TMCP steels A and C have low C content. C content levels determine toughness properties in general and high C content is detrimental to toughness as Hatting and Pienaar (1998) have concluded. In this study, TMCP steels A and C have low C contents, whereas the C content in QT steels was considerable bigger. Accordingly, TMCP steels have good toughness values in the HAZ than most of the QT steels. As Tian (1998) and Hatting and Pienaar (1998) have researched, heat input has a direct effect on impact toughness in Nb added HSSs. When using a low heat input in welding, this will increase impact toughness, while if a high heat input is used in welding it will decrease the impact toughness in the HAZ. Six of eight tested steels had Nb as an alloying element in this study and the greater heat input led to the lower impact toughness. Ti precipitations have an impact to grain growth and they inhibit it very well. However, if the heat input is too high on welding, then the grains grow too much which leads to the coarse structure in the HAZ and consequently destroys the welded structure. Liu and Liao (1998) researched Ti nitrides and found that 117 those nitrides inhibit grain growth especially in high temperatures. However, when the temperature is too high for an extended period of time, the Ti nitrides also dissolve in the structure and their influence diminishes. This specific influence is seen in this study when using heat inputs 1.3 and 1.7 kJ/mm, where impact ductility values have decreased and grains have grown. Only steels G and H do not have Ti as an alloy element. As Rak et al. (1997) has concluded and is also clearly displayed in this research, the size and distribution of the Ti precipitates are important when studying the grain growth control and comparing it to the role of the chemical composition of the precipitates. It is important to keep the heat input as low as possible, because Ti precipitate dissolves in to base material at higher temperatures. When the heat input is kept low, there is no time for precipitates to dissolve and the properties of the welded structure remain satisfactory. In this research, the lowest heat input 1.0 kJ/mm gives the best results of impact ductility and strength test on the chemical composition and microstructure of HSS. In the undermatched weld structure, local mismatch can be the reason for lowered toughness. Dilution and alloying are not evenly distributed in undermatched welds and this leads to local mismatch. This study similarly clarifies the differences between impact test values as was in the Rak et al. (1995) study. 118 WELD Charpy V impact test values J 70 61 60 50 51 50 47 41 37 36 37 36 30 46 47 40 39 36 30 33 32 30 32 28 29 20 17 22 18 49 48 36 24 22 20 43 42 37 32 26 25 22 17 10 43 34 27 19 18 16 14 43 42 38 37 27 22 24 24 26 17 38 40 33 32 30 29 24 16 16 0 A B C D E WELD 1.0 kJ/mm HEAT F 1.3 kJ/mm G H 1.7 kJ/mm Figure 48. Impact test values to weld metal using different heat input when filler material was undermatched. HAZ Charpy V impact test values J 80 75 60 50 47 40 36 22 25 20 16 10 12 10 9 10,5 10 11 10 9 10 20 0 A B HEAT INPUT 32 58 53 49 24 24 18 C 44 37 37 9 7 44 28 23 23 18 18 14 12 9 8 D E WELDS 1.0 kJ/mm 45 31 27 23 18 16 17 10 7 11 6 7 F 1.3 kJ/mm 56 57 55 51 11 8 G 16 4 13 98 4 H 1.7 kJ/mm Figure 49. Impact test values to HAZ area structure using different heat input when filler material was undermatched. 119 In addition to previous research (Wang et al. 2003), this study confirms the influence of heat input to impact toughness in HSS welding. As the heat input grows, the deterioration of impact toughness in the HAZ of HSSs is quite clear. In this study, steels F and H had very low HAZ area impact values. To further bolster confidence in these impact test results, and uncover different implications, it was additionally determined to conduct CTOD tests. 120 Table 33. Impact test values from weld when filler material was undermatched. HEAT INPUT 1.0 1.3 1.7 kJ/mm kJ/mm 48 31 30 28 36 46 17 29 47 22 39 18 32 30 32 32 33 20 41 47 36 51 37 30 50 37 36 22 49 24 36 48 36 26 32 25 22 17 43 14 42 37 27 19 16 34 43 18 38 27 22 37 24 17 43 26 24 61 42 16 38 16 33 32 29 40 24 30 40 STEEL kJ/mm A B C D E F G H 121 Table 34. Impact test values from HAZ when filler material was undermatched. HEAT INPUT 1.0 1.3 1.7 kJ/mm kJ/mm 41 11 22 50 22 16 12 10 25 10 20 9 9 10.5 10 11 10 10 32 47 36 53 37 24 18 37 24 9 49 7 58 75 8 18 44 12 23 28 14 23 18 9 16 7 10 17 7 6 18 11 11 46 31 44 55 23 27 56 57 51 11 8 16 13 9 4 8 8 4 STEEL kJ/mm A B C D E F G H 122 7.9. Hardness test The welds were also subjected to Vickers hardness tests with SFS-EN ISO 6507-1 standards. The tests were conducted on the weld and HAZ areas at 0.5 mm intervals. Fig. 50 shows a hardness measurement sample, while figs 51, 52 and 53 show the hardness test results by varying heat inputs. In all of the tested pieces, the hardness values in the weld metal were the same, but when test moved though HAZ from fusion line to base material, hardness values were higher than in the weld. Test values differed depending on the base metal test material. In QT steels, the HAZ hardness curve is at its highest point in the CGHAZ, (just beyond the fusion line and the base material) and sinks down to its lowest point in the ICHAZ. As Loureiro (2002) has explained, a loss of hardness occurs in the ICHAZ because of carbide precipitation. Beyond the ICHAZ, the hardness levels rose until they reached the hardness level of the base material. The highest hardness values observed were the same or slightly higher than the hardness of the base material, a phenomenon that possibly be explained as an effect of quenching in the HAZ. Steels A and C which were made using TMCP method behaved quite different than the QT HSSs. The TMCP steels had a very straight hardness curve in the weld and the HAZ, the hardness curve gradually grew to the hardness of base material. These steels had C content 0.05 % while QT steels had C content close to 0.15 %. Additionally, the two types of steel have differences in their base material microstructure with TMCP steel having a ferrite-bainite mix and QT steel having martensite and bainite. Prior to welding, all steels had near same base metal hardness levels, 280-290 HV5. B was added as an alloying element to steels B, D, E and F, because B has been known to increase hardness in low carbon steels. Moon et al. (2008) researched that the hardness of the CGHAZ increases when B is added as 123 alloy element. In the same investigation, the researchers noticed that the impact toughness decreased at the same time. They additionally used Cu as an alloying element in their investigation. During the welding process, the microstructure of all the steels changed to austenite near the fusion line. Steels A and C, which have a low C content of 0.05%, did not quench, which must explain the lower martensite content in their CGHAZs. QT steels behaved quite differently, as after welding there was a change the martensite microstructure could be moved in the CGHAZ, thus increasing hardness to its highest point. A bigger C content of approximately 0.15 %, potentially contributed to the harder microstructure (martensite and bainite) observed in QT steels. C is the most important alloying element in the quenching process and as other studies (Kaputska et al. 2008) have concluded the peak hardness of QT steel is higher in the HAZ than in the base metal. Nb is an important alloying element in HSSs and it has an effect on the hardness as well. The content of martensite can depend on Nb as Zhang et al. (2009) has researched. They concluded that when the Nb content was 0.026 % and the cooling rate was high, martensite was observed, however when Nb was not in the steel no martensite was observed. An overall, maximum hardness of around 345 HV5, was observed when heat input was 1.0 kJ/mm in QT steel G. At the same heat input, QT steels F and H also had a maximum hardness that exceeded 300 HV5. Loureiro (2002 in accordance Yurioka et al. 1987) concluded that a totally martensite structure should have a maximum hardness of 444 HV10, while a non-martensite microstructure should have a hardness of 223 HV10. In this study the microstructure in the CGHAZ area of QT HSS was lower bainite and tempered martensite with a maximum hardness of 300 HV5 or more. In this study, it was clearly observed that slower cooling rates lead to lower hardness levels in the HAZ as was similarly concluded in the research conducted by Kaputska et al. (2008). This happens because the autotempered martensite has formed in the CGHAZ. In fig. 52, some curves end before base metal hardness, which means that these steels had wider HAZs than is able to 124 be read in the table. Our study proves that when welding HSSs, lower cooling rates tend to produce a wider HAZ, a phenomenon that has also been studied by Kaputska et al. (2008). FUSION LINE WELD METAL BASE METAL HAZ 2 mm 0.5 mm STEELS A B C D E F G BASE MATERIAL 6 BASE MATERIAL 5 BASE MATERIAL 4 BASE MATERIAL 3 BASE MATERIAL 2 BASE MATERIAL 1 HAZ 7 HAZ 6 HAZ 5 HAZ 4 HAZ 3 HAZ 2 HAZ 1 FUSION LINE WELD 4 WELD 3 HEAT INPUT 1.0 kJ/mm WELD 2 350 330 310 290 270 250 230 210 190 170 150 WELD 1 HARDNESS HV5 Figure 50. Sample from hardness measurement. H Figure 51. Hardness of the welded structure when the heat input was 1.0 kJ/mm. 125 WELD 1 WELD 2 WELD 3 WELD 4 FUSION LINE HAZ 1 HAZ 2 HAZ 3 HAZ 4 HAZ 5 HAZ 6 HAZ 7 BASE MATERIAL 1 BASE MATERIAL 2 BASE MATERIAL 3 BASE MATERIAL 4 BASE MATERIAL 5 BASE MATERIAL 6 HARDNESS HV5 350 330 310 290 270 250 230 210 190 170 150 HEAT INPUT 1.3 kJ/mm STEELS A B C kJ/mm. 126 D E F G H Figure 52. Hardness of the welded structure when the heat input was 1.3 STEELS HEAT INPUT 1.7 kJ/mm WELD 1 WELD 2 WELD 3 WELD 4 WELD 5 FUSION LINE HAZ 1 HAZ 2 HAZ 3 HAZ 4 HAZ 5 HAZ 6 HAZ 7 BASE MATERIAL 1 BASE MATERIAL 2 BASE MATERIAL 3 BASE MATERIAL 4 BASE MATERIAL 5 BASE MATERIAL 6 BASE MATERIAL 7 BASE MATERIAL 8 BASE MATERIAL 9 BASE MATERIAL 10 BASE MATERIAL 11 HARDNESS HV5 350 330 310 290 270 250 230 210 190 170 150 A B C D E F G H Figure 53. Hardness of the welded structure when the heat input was 1.7 kJ/mm. Fig. 54 shows the hardness results for the TMCP steels. This table shows that the hardness of the HAZ area does not grow until it reached the base material. Heat input also affects the width of the HAZ area, with greater heat inputs leading to wider HAZ areas. Of all the TMCP steels, the only one not displaying hardness growth at the base material was steel A when the heat input was 1.7 kJ/mm. Under these conditions, steel A is likely to have a wider HAZ than the other steels and hardness measurements from the base material of steel A were not captured within the scope of this test. Fig. 55 shows the hardness results of steel G which was 12 mm thick. It was remarkable that the width of the HAZ was the same regardless of heat input values. This could be attributed to the thickness of steel G which was greater 127 than all of the other steels. Another reason for this behaviour could be within the microstructure of steel G, which had neither Nb nor Ti. These microelements have a big influence upon the microstructure, where Ti inhibits grain growth in the HAZ, while Nb only has an effect upon the HAZ with the presence of other alloying elements. Overall, it is not very clear why the HAZ width is the same regardless of heat inputs in steel G. Another option to consider is whether the equations for two or three dimensional conduction of heat are still valid with Hardness HV5 HSSs. 350 300 TMCP steels A and C 250 Steel A 1.0 kJ/mm 200 Steel C 1.0 kJ/mm 150 Steel A 1.3 kJ/mm Steel C 1.3 kJ/mm Steel A 1.7 kJ/mm Steel C 1.7 kJ/mm Figure 54. Hardness of the welded structure of TMCP steels A and C. 128 STEEL G 1.0 kJ/mm 1.3 kJ/mm 1.7 kJ/mm WELD 1 WELD 2 WELD 3 WELD 4 WELD 5 FUSION LINE HAZ 1 HAZ 2 HAZ 3 HAZ 4 HAZ 5 HAZ 6 HAZ 7 HAZ 8 BASE MATERIAL 1 BASE MATERIAL 2 BASE MATERIAL 3 BASE MATERIAL 4 Hardness HV5 350 330 310 290 270 250 230 210 190 170 150 Figure 55. Hardness of the welded structure of QT steel G. In HSSs 780 and 980 DP, it was noticed that a greater reduction in base metal hardness occurs in the HAZ of 780 DP steel. This may be due the higher dislocation density present in the ferrite phase of this material producing a larger driving force for recovery (Kaputska et al. 2008), which is important to notice when planning steels structures using HSSs. 7.9. CTOD tests Even after the impact tests the fracture strength of welded structure was still unambiguous. As some results were not within the limit of the standards, CTOD tests were deemed necessary. These CTOD test were done according to standard ASTM E1290-2. The first CTOD test was conducted on the welded structure while the other test first used Gleeble simulation (as reported in experimental investigations 6.6.) before continuing with CTOD testing. 129 CTOD tests are trustworthy and give accurate measurements of material toughness. It is very important to clarify toughness in a welded structure, especially in the HAZ which is a critical area in relation to material toughness. The CGHAZ of the HAZ has been reported (Shi & Han 2007, Lee & al. 1993, Güran & al. 2007) to be the most brittle area where toughness is at its lowest. Depending on heat input, the CGHAZ can have different widths. Finding the CGHAZ during testing has proven to be quite difficult. Simulation has been developed to clarify the characteristics of different areas in the HAZ, and a Gleeble simulation was used in this research to clarify ductility in the CGHAZ. If the weld is welded with many passes, then the ICCGHAZ has been observed (Liu at al. 2007, Hamada 2003, Li et al. 2001, Lambert et al. 2000, Matsuda et al. 1995, Davis & King 1993, Lee et al. 1993) to be the worst impact ductility zone between two CGHAZs. This LBZ has a very brittle structure where the MA phase will destroy the impact ductility. This ICCGHAZ is narrow and discontinuous, and only 0.5 mm width (Davis & King 1993) depending on heat input. CTOD tests are better suited to find this kind of brittle areas than CharpyV tests, but in this study test place was unfortunately too far from the fusion line and ICCGHAZ LBZs were not under investigation. In Gleeble made test bars only one heat input was used. The very brittle microstructure proves that the CGHAZ is a weak area within the HAZ. In this situation, it is assumed that the CGHAZ is the weakest zone in welded structure. In real structures, there are many zones in the HAZ and the width of the CGHAZ is usually narrow. The total width of all zones in the HAZ depends on heat input. When the heat input is large, those zones are wider and the tensile strength and toughness properties of the structure go down. The microstructure of the CGHAZ can be composed of M-A constituents and this making the structure very brittle. A good example of this brittle structure is seen in fig. 57 which was taken of Gleeble simulated QT test bar. The main microstructure is martensite and the proportion of bainite is less than half. Additionally, the coarseness of bainite is a metallurgical factor affecting the impact properties as Lampert et al. (2000) have also studied. 130 CTOD test results from the welded structure and base material are presented in table 35 and in fig. 58. Overall, the base material has the lowest CTOD value. Only steel H exhibited different behaviour, as the base material of steel H had the highest CTOD value and only decrease by its higher heat inputs. This result was one of the hypotheses of this study. As seen in fig. 58, the highest results from this CTOD test were 0.2 or more. Five of the eight tested HSSs reached this value when the heat input was 1.7 kJ/mm. Steel A also reached this value with a heat input of 1.0 kJ/mm, but the value was too high as the result of a measurement mistake which is not clear. There were big differences between the base material CTOD test values. The lowest values, near 0.05, were seen in steels A, C, D and G, whereas the highest value, 0.2, was seen in steels B, F and H. With the heat input at 1.0 and 1.3, the measured values were not so unambiguous because the measured HSS, like steel B, had a low value when the heat input was 1.0 (0.15) and a high value (0.3) when the heat input was 1.3 and 1.7kJ/mm. Steel F had good values with all the welded structures. It is very difficult to find the weakest zone of the HAZ. It was expected that the CGHAZ would to be the weakest area, however, it is very difficult to find the CGHAZ from within the HAZ. The place of CTOD test was 2 mm from fusion line, the same measurement as was used in Charpy-V test. As shown in table 35, near all test results were higher than base materials results. This most likely means that these measurements were taken from a HAZ area other than the CGHAZ. In fig. 56 it can clearly be seen that the place of the test was not in the CGHAZ. Depending on heat input, this zone of the HAZ was so far from the fusion line that the test place was most likely in the ICHAZ or SCHAZ. When conducting the CTOD test on a welded structure made from HSS, the initial crack must not be more than 0.5 mm from fusion line. If this criterion is met, then the initial crack will be in the CGHAZ. Table 36 shows the results from Gleeble tested pieces. These test pieces were made to clarify the features of the CGHAZ microstructure from the tested HSSs. These CTOD test results are very low compared to CTOD test results from 131 welded structure, which means that all the CTOD test results from simulated structures were very brittle, as seen in fig. 59. In fig. 59, it is clearly explained that all of the results of this CTOD test were very low and within close value proximity to one another. Only steel E had one value over 0.05, however, this value was very low when compared to the base material. Fig. 58 additionally explains that the base material CTOD test values in all the tested HSSs were higher than in the Gleeble simulated test bars. 2 mm End of sawed crack starter notch Fusion line Crack extension using cyclic force Breakage after bending Figure 56. Broken QT steel CTOD test piece where the place of initial crack is well seen. When welding using undermatched filler material as was used in this study, it is clear that the weakest zone is in the weld. The rate of undermatching has a significant role in fracture toughness. When the rate of undermatching is low, the HAZ can have lower toughness than the weld or base material. This can encourage the toughness of the welded structure to decrease. Pisarski and Dolby (2003) concluded that the worst case fracture toughness of softened HAZs occurred when the HAZ undermatched in strength both the weld deposit 132 and the parent plate. In this study, the highest level of undermatching was in the weld, which makes the fracture toughness acceptable. 100 µm Figure 57. The CGHAZ of Gleeble simulated and CTOD tested QT steel. Aspect ratio is 1:500. Table 35. CTOD test values from the welded structure. HEAT INPUT 1.0 HEAT INPUT 1.3 HEAT INPUT 1.7 kJ/mm kJ/mm kJ/mm BASE MATERIAL VALUE VALUE VALUE VALUE STEEL mm CATEGORY mm CATEGORY mm CATEGORY mm CATEGORY A 0.05 c 0.42 m 0.12 m 0.12 u B 0.21 m 0.14 m 0.30 m 0.31 m C 0.06 u 0.07 u 0.12 m 0.27 m D 0.07 u 0.14 u 0.04 c 0.27 m E 0.10 m 0.11 u 0.14 m 0.21 m F 0.18 m 0.23 m 0.28 m 0.28 m G 0.08 u 0.10 u 0.12 m 0.11 m H 0.21 m 0.17 u 0.15 u 0.13 u c= critical u= unstable m=high tensile 133 Table 36. CTOD values (mm) of Gleeble simulated CGHAZ. Base Heat Heat Heat input input input 1.0 1.3 1.7 STEEL material kJ/mm kJ/mm kJ/mm A 0.05 0.022 0.015 0.027 B 0.21 0.024 0.025 0.018 C 0.06 0.012 0.011 0.008 D 0.07 0.013 0.016 0.008 E 0.10 0.031 0.036 0.065 F 0.18 0.021 0.010 0.016 G 0.08 0.011 0.014 0.022 H 0.21 0.021 0.017 0.010 0,45 CTOD values of welded HAZ structure 0,4 0,35 0,3 BASE MATERIAL 0,25 HEAT INPUT 1.0 kJ/mm 0,2 HEAT INPUT 1.3 kJ/mm 0,15 HEAT INPUT 1.7 kJ/mm 0,1 0,05 0 A B C D E F G H Figure 58. Compared CTOD values (mm) of welded HAZ structure. 134 0,25 CTOD values of simulated structure 0,2 0,15 Base material Heat input 1.0 kJ/mm 0,1 Heat input 1.3 kJ/mm Heat input 1.7 kJ/mm 0,05 0 A B C D E F G H Figure 59. Compared CTOD values (mm) of Gleeble simulated CGHAZ. When comparing the structure of a Gleeble made CTOD test bar and a welded test bar, the size and phase of microstructure of the CGHAZ is different. In Gleeble made test bars, the initial austenite grain size was greater, ranging in value from 3-4 (ASTM E112-10) than grain size of welded CGHAZ, ranging in value from 4-5 (ASTM E112-10). The same differences in size were observed in the initial austenite grains of both QT and TMCP HSSs. This is explained in more detail in 7.11.5. Additionally, the microstructure of Gleeble made test bar had more martensite than the welded CGHAZ. In this study, the fracture toughness between welded and simulated structure cannot be compared because the CTOD test of welded structure was not in the CGHAZ. 7.11. Additional microstructure tests The test results of the additional microstructure tests to the steels TMCP HSS C and QT HSS E clearly clarify differences between the QT and TMCP HSSs. Steels C and E well describe their own steel group and the results are characteristic of both their own steel group. 135 7.11.1. Microstructure of the base material As was studied earlier in this research, the microstructure of the base material of the QT and TMCP HSSs differ and also the HAZ microstructure changes are different. The types of differenced have been tested in additional microstructure tests. The microstructure of the QT HSS E base metal, fig. 60, consists of tempered martensite and bainite. The size of the initial austenite grain corresponds to 12 number according to ASTM E112-10, 5.6 µm. Microstructure of base metal is homogeneous, and through thickness inequigranularity was not observed. Limited carbon content up to 0.15 % in base metal allows obtaining lath martensite and avoiding formation of twinned martensite in order to increase toughness in combination with high strength. The microstructure of the steel C, TMCP HSS, base metal, fig. 61, consists of bainite (70%) and ferrite (30%). The effective grain size of the base metal corresponds to 14 number, ASTM E112-10, 3.0 µm. The optimum microstructure with a desired balance of mechanical properties are achieved through a suitably designed thermomechanical process. This includes heavy deformation of the austenite, carried out in the non-recrystallisation temperature region, which brings about significant refinement of the final transformation microstructure. Figure 60. QT HSS E microstructure: tempered martensite and bainite. 136 Figure 61. TMCP HSS C microstructure: bainite and ferrite. Both QT and TMCP HSS steels have exceptional working properties and although they have different microstructures, both HSSs are good to cold form, cut or machine. However, the welding these steels makes their properties quite different. Both HSS have bainite in their microstructure, but the tempered martensite microstructure of QT HSS forms differently than ferrite in TMCP HSS. Also TMCP HSS has more bainite (70%) than QT HSS, and additionally rolling TMCP HSS has worked its faces more parallel than the faces of QT HSS. 7.11.2. Microstructure of weld metal The weld metal does not differ between TMCP and QT HSS steel. Initial columnar grains formed by epitaxial growth are detected by the presence of grains of polygonal ferrite and Widmanstatten ferrite along the former grain boundaries. However, the main constituent is an acicular ferrite, forming a "wicker basket". Both base metal weld microstructures are illustrated in fig. 62. 137 a) b) Figure 62. Microstructure of TMCP (a) and QT HSS (b) weld. Wf (Windmanstatten ferrite), pf (polygonal ferrite) and af (acicular ferrite) are observed. 7.11.3. Microstructure of HAZ of QT and TMCP HSS Microstructure of the metal surrounding weld interface is influenced by heat while the weld joint is being formed. In the studied welded joint of QT HSS E and TMCP HSS C CGHAZ, FGHAZ, ICHAZ and SCHAZ are clearly recognized. The microstructure changes continuously depending upon the maximum temperature attained in each region of the HAZ. Close to the weld interface the metal is exposed to the temperatures between liquidus and solidus lines described as the fusion line (FL). This zone is in partially melted state. Microstructure of FL of QT HSS E has mixed microstructure which contains bainite and polygonal ferrite, fig. 63. 138 Figure 63. Optical microstructure of the fusion line of QT HSS E. QT HSS E CGHAZ borders the FL and refers to the HAZ subjected to peak temperatures above the grain coarsening temperature, the latter is 1300 oC for steels which have been Ti-treated to elevate their grain coarsening temperature (Eastling 1992). As the peak temperature exceeds the critical point, AC3, complete retransformation to austenite occurs, fig. 64. The extent of following grain coarsening depends on the peak temperature, the time above the grain coarsening temperature, the chemical composition of the steel and presence of undissolved nitride and carbonitride particles. When heated above 1300 oC, most of these particles, except the most stable such as TiN, dissolve (Mitchell et al. 1995). This results in reduction of pinning effect of the particles and following grain growth. At the same time long exposure of the HAZ to high temperature promotes homogenizing of austenite by alloying elements. So grain coarsening and homogenizing of austenite make it more stable. During cooling the grain coarsened austenite transforms to non-equilibrium transformation products depending on steel chemistry and cooling rate. 139 Figure 64. Scheme of CGHAZ formation. In both TMCP and QT HSS steel coarse grain microstructure of initial austenite grains is clearly revealed in CGHAZ. Austenite grains grew from 5.6 µm, number 12 (base metal) up to 75 µm, number 4-5 (according to ASTM E112-10) during welding heating. In QT HSS E during subsequent cooling coarse grains were divided into packets of a lath bainite and low-carbon martensite, which slightly refines the constituents of the structure and has a positive effect on the resistance to crack propagation (Lamberte-Perlade et al. 2004). In TMCP HSS C during subsequent cooling coarse grains were divided into packets of a lath and granular bainite. Both microstructures are seeing in fig. 65 a and b. a) b) Figure 65. Optical microstructure of CGHAZ of QT HSS E (a) and TMCP HSS C (b). Identification of structural constituents was derived from measuring their microhardness. Microhardness indentation was conducted by Vickers scale and 140 0.025 kgf loading. The hardness of the lath martensite in CGHAZ exceeds 300 HV and reaches 340 HV, fig. 66. Packets of bainite have hardness less than 300 HV. Tempered martensite of the base metal is characterized by a hardness of 270-280 HV. In the present investigation, martensite and bainite are distinguished by quite different etching susceptibilities as shown by optical micrographs, fig. 66. Since the bainitic transformation occurs at a higher temperature compared to the martensitic transformation, carbon can diffuse to a greater extent either to the remaining austenite islands or to the boundary between laths (Thewlis 2004). When this structure is etched, the boundaries of the retained austenite islands or its decomposition products etch deeply, giving the overall appearance of a plate shaped ferritic matrix with a superimposed dispersion of dark contrasting particles. The martensitic transformation is characterized by clusters of very fine ferritic laths which form at lower temperatures. Since the carbon distribution in the martensitic structure is more uniform, it etches more evenly. Figure 66. Microhardness measurement in CGHAZ of QT HSS Е. The microstructure of the CGHAZ of TMCP HSS C formed during weld thermal cycle consists of the products of bainite transformation of austenite, fig. 67. These microstructures are classified as bainite which may take many morphologies. Bainite-ferrite is one example of a microstructure which consists of a carbide-free ferrite matrix with well-defined islands of retained austenite or martinsite-austenite (M-A) constituent. The microstructure of granular ferrite 141 consists of dispersed retained austenite or M-A constituent in a featureless matrix which may retain the prior austenite grain boundary structure (Krauss G & Thompson 1995). Most prior austenite grain boundaries are clearly visible in CGHAZ of TMCP HSS C, allowing the mean austenite grain size to be measured. The mean austenite grain size at this size is 89.0 µm, 4 number (according to ASTM E11210). Within prior austenite grain several crystallographic packets with high misorientation angles between them, which slightly refines effective grain size, can be identified. Figure 67. Optical microstructure of CGHAZ of TMCP HSS C. As determined in CTOD and Charpy-V tests, a coarse microstructure decreases impact ductility. Charpy-V values of CGHAZ TMCP HSSs were good but some QT HSS steels had low impact ductility values. CTOD test values of Gleeble made CGHAZ test bars were very low. Impact ductility of bainite microstructure is higher than martensite microstructure. FGHAZ refers to HAZ regions which have been subjected to peak temperatures between the austenite grain coarsening temperature and the upper critical point AC3, typically between about 1300 and 910 °C (Eastling 1992). Both CGHAZ and FGHAZ are the zones which have become fully austenitic due to weld thermal cycle. The microstructures of these zones continuously change 142 according to the former austenite grain size. Consequently, it is difficult to precisely indicate the boundary between CGHAZ and FGHAZ. The reduction in peak temperatures in this zone implies that, following the α→γ transformation during heating, the austenite does not have time to develop properly, and the grain size remains small. In addition, nitrides and carbides may not be fully dissolved, fig. 68. During α→γ transformation γ grains nucleate heterogeneously at the boundaries prior γ grain and grow along them. Also the nucleation of γ grains occurs due to the dissolution of cementite, fig. 68. During γ→α transformation, the large grain boundary area tends to promote nucleation of fine ferrite grains. Figure 68. Scheme of FGHAZ formation. Along the HAZ of HSS QT steel, FGHAZ has the most fine grain structure with the mean grain size of 4.0 µm, 13 number (according to ASTM E112-10), fig. 69. There are more equilibrium transformation products, such as polygonal ferrite, and islands of granular bainite in this zone. Compared with tempered martensite of BM, microstructure constituents of FGHAZ have lower hardness. Hardness of ferrite equals 210 HV, granular bainite 230 HV, fig. 70. 143 Figure 69. Optical microstructure of FGHAZ of QT HSS E. Figure 70. Microhardness measurement in FGHAZ of QT HSS E. As a result of rapid heating and short exposure to high temperatures, the homogenization of austenite is not completed and some islands of retained austenite are enriched by carbon, that could promote formation of martensite or transformation to perlite in these islands. The most fine grain and uniform structure within the HAZ of TMCP HSS C is observed in FGHAZ, fig. 71. The microstructure contains mostly polygonal ferrite with a hardness of 220 HV and dispersed islands of granular bainite with a hardness of 240 HV. 144 Figure 71. Optical microstructure of FGHAZ of TMCP HSS C. ICHAZ refers to HAZ regions which have been subjected to peak temperatures between the upper and lower critical points AC3 and AC1, typically between 910 and 720 °C (Eastling 1992). In this region partial retransformation to austenite occurs during heating, the exact extent of which is governed by the peak temperature within the intercritical temperature range. During cooling, the austenite regions decompose to different extents and to various transformation products, fig. 72 (Matsuda et al. 1996). Figure 72. Scheme of ICHAZ formation. The microstructure in this region consists of a mixture of bainite, tempered martensite and perlite, fig. 73. Carbides, mainly cementite also experience a process of spheroidization and coagulation. 145 Figure 73. Optical microstructure of ICHAZ of QT HSS E. The SCHAZ is the region of HAZ that has been subjected to peak temperatures below the lower critical point AC1, below 720 °C (Eastling 1992). The processes of nucleation and spheroidization of carbides occurs in this zone, fig. 74. Black cementite conglomerates are clearly identified in fig. 75. The agglomeration of spheroidized cementite particles at grain boundaries and triple junctions emphasizes the role of grain boundaries as high diffusivity channels for carbon at these low temperatures. Figure 74. Scheme of SCHAZ formation. 146 Figure 75. Optical microstructure of SCHAZ of QT HSS E. The ICHAZ and SCHAZ regions of TMCP HSS, fig. 76 a and b, can be hardly distinguished unlike the HAZ of steel QT, fig. 73 and 75. This happens because the TMCP steel has a low carbon content and heating up to temperatures around critical point AC1 does not produce large scale nucleation of cementite and its coagulation. a) b) Figure 76. Optical microstructure of TMCP HSS C: a) ICHAZ, b) SCHAZ. 7.11.4. Comparison of HAZ microstructure of steels QT and TMCP Measurements of the microhardness in cross section of the studied welded QT and TMCP HSS joints were made. Distributions that were obtained are shown in Fig. 77. 147 Base metal microhardness of the QT and TMCP steels is similar, 265 and 273 HV respectively, fig. 77. The weld metal has a lower hardness (200-210 HV) in comparison with BM, while undermatching between the weld and base metal occurs. a) b) Figure 77 . Microhardness distribution in the weld joint: а) QT HSS E; b) TMCP HSS C. As it seen from fig. 77, the HAZ microhardness of the both steels varies over a wide range. Characteristic of the TMCP steel HAZ is a general decrease in hardness with respect to the base metal. In the HAZ of QT steel a decrease as well as increase in hardness is observed depending on the resulting microstructure. In the HAZ of QT welded joint the highest hardness reaches 290-317 HV and is observed in the CGHAZ close to the fusion line. This can be explained by the formation of bainite- martensite microstructure. Increased hardenability of steel at the CGHAZ, because of the increased carbon content in the base metal and a strong grain growth due to welding thermal cycle, is the cause of such microstructure. There is a gradual decrease in hardness with decreasing the fraction of martensite in the microstructure with the distance from the weld and the associated reduction of maximum heating temperature. 148 CGHAZ of the TMCP steel welded joint has microhardness 230-240 HV. The decrease in the hardness in relation to the base metal is explained by the full recrystallization of microstructure and its transformation to the austenite during heating. Optimum microstructure with a desired balance of mechanical properties and primary bainitic microstructure with a high density of dislocation are achieved through suitably designed thermomechanical process. Low hardenability of the TMCP steel is explained by a very low level of alloying elements and carbon. When heating exceeds the AC3 temperature the full recrystallization of the microstructure occurs and the more equilibrium products of transformation with lower density of dislocation are achieved. This is the main cause of a decreasing of the hardness at a considerable distance from the HAZ. The lowest hardness in the HAZ of the both steels corresponds to the FGHAZ. It is explained by the formation of the polygonal ferrite in this area. Austenite fine grain and insufficiently high cooling rates assist in transformation of the austenite into ferrite with low density of dislocations. SCHAZ of the both steels is characterized by the decrease in hardness due to tempering of the base metal. During heating between AC1 and AC3 (ICHAZ) austenite composition in the microstructure varies from 0 to 100% according to the local maximum temperature or in other words to the distance from the fusion line. Ferrite as a product of austenite decomposition determines the hardness of this region of the HAZ after cooling. So both TMCP and QT steels are characterized by the softening in the HAZ but the lowest hardness relates to the weld metal. Formation of the quenched structures in the HAZ of QT steel can lead to cold cracking during welding and deterioration of the toughness of CGHAZ. 149 7.11.5. Microstructure study of CTOD samples after simulated welding thermal cycle Fig. 78 shows the microstructure of QT HSS E Gleeble simulated and welded joint when heat input was 1.3 kJ/mm. The grain boundaries are depicted by the red lines. The Gleeble sample has a coarse microstructure in comparison with CGHAZ of the real welded joint. Austenite grains have a number 4-5 (according to ASTM E112-10) in the welded joint and 3.5-4 in the simulated sample, the differences of which can be explained by the effect of high temperature gradient in the welded joint. Additionally, the microstructural constituents of both samples are similar. a) b) Figure 78. Microstructure of CGHAZ HSS E of the Gleeble sample a) and a real GMAW welded joint b). Fig. 79 shows the microstructure of TMCP HSS C Gleeble simulated and welded joint when heat input was 1.3 kJ/mm. The Gleeble sample has a similar microstructure in comparison with CGHAZ of the real welded joint. Austenite grains have a number 4 (according to ASTM E112) in case of welded joint and 4-5 for simulated sample. Additionally, the microstructural constituents of both samples are similar. 150 a) b) Figure 79. Microstructure of TMCP HSS E CGHAZ of the Gleeble sample (a) and a real GMAW welded joint (b). When comparing the HAZ grain growth in the simulated weld and the real weld, this current study also came to the same conclusions of previous research (Easterling 1992) where it was observed that the maximum initial austenite grain size in real welds is less than what is seen in simulated welds. This trend has been mainly observed with medium heat input values. This phenomenon can possibly be explained by the fact that small grains hinder the grain growth of the large grains in real welds. This is shown in fig. 80, where it can be observed that the change in grain size is associated with a very steep temperature gradient. The grains can move other way, like grains which have a large temperature gradient across them tend to grow non-uniformly. Then it results change of shape from equiaxed to pear-shaped. A grain can also experience surface tension restrictions when adjacent grains are trying to expand at faster or slower rates (Easterling 1992). 151 Figure 80. Grain size in the HAZ as a function of the peak temperature and distance from the fusion line (Easterling 1992). 8. DIVERGENCE IN MANUFACTURERS’ HSS’s WITH DIFFERENT HEAT INPUTS This study began with the idea that the main structure of the base material is different when comparing TMCP, QT and DQ HSSs. It is especially important to consider that the microstructure of these steels are quite different from one another. The chemical properties of these steels are also different with some steels having a large variety of alloying elements compared to other HSSs. QT and DQ steels have same kind of tempered martensite and bainite microstructure, while the main microstructure of TMCP steels is ferrite-bainite. When comparing QT and DQ steels with TMCP steels, there is a distinct difference in the HAZ hardness as seen from the results of the hardness tests. When constructing steel structures using TMCP steels, the HAZ hardness must be taken into consideration. When using undermatching filler material this is not of the utmost importance, however when using matching filler material, the HAZ hardness should be closely monitored. The divergence between different manufacturers HSSs can be clearly observed in the width of the HAZ. Many researchers (Magudeeswaran et al. 2008, Shi & Han 2008, Liu et al. 2007, Liu W-Y 2007, Pavyna & Dabrovski 2007, Wang et al. 2003, Basu & Roman 2002, Louriero 2002, Nevasmaa et al. 1992a, Nevasmaa et al. 1992b, Vilpas et al. 1985) have examined the effects in the 152 HAZ under welding, especially the effects of the heat input and the cooling time. It has been clearly observed within this research that higher yield strengths HSSs require lower heat inputs and cooling time t8/5. In this study there were so many different HSS from the different manufacturers (eight steels from six manufacturers) that the observation was unambiguous regardless of the steel or manufacturer. Welding DQ HSSs required the lowest values in heat input. The microstructure and the hardness are the most susceptible areas of DQ steels, yet these steels have the highest yield strength of HSSs. The martensite and bainite microstructure of QT steels leads to a brittle CGHAZ structure, as seen in DQ steels, and therefore the heat input must be low, near 1.0 kJ/mm. In TMCP steels hardness decreases in the CGHAZ when compared to the base material. This must be taken into consideration, especially if the filler material is undermatched, because a soft HAZ can weaken the entire welded structure. If the filler material is matched and the heat input is as low as 1.0 kJ/mm, then despite reduced hardness in the HAZ, the welded structure will have as good strength values as the base material in TMCP HSS. At the same time that hardness is decreasing, the impact toughness decreases when the cooling time is longer. Similar to the work of other researchers (Shi & Han 2008, Liu at al. 2007, Wang et al. 2003, Juan et al. 2003), this study has also observed that greater heat input leads to decreasing impact toughness values, which can lead to damage in the welded structure especially in low temperatures such as -40 °C. The same phenomenon will happen despite the manufacturing method used to make the HSSs. This study has shown that the welding circumstances in the workshops, good professional skills, and needed WPSs are important when pursuing a good impact toughness in HSS welded structures. 153 9. CONCLUSIONS 1. It has been acknowledged in this study that when welding HSS with a minimum yield strength of 690 MPa, the heat input cannot be over 1.0 kJ/mm. If the heat input is greater than this, then the impact ductility, toughness, tensile strength and fatigue strength properties of the welded structure start to decrease and in the worst case scenario, the welded HSS will break unexpectedly because of the brittle structure in the HAZ. M-A grains in the HAZ can be the source of initial crack as an increased heat input results in more M-A grains in the HAZ. The heat input 1.0 kJ/mm in this study leads to t8/5 time 21 s if the plate thickness is 8 mm, however when the thickness of the welded plate is thicker, the heat input must be calculated again. Based on earlier studies, it was recommended that the heat input for HSS with a minimum yield strength of 690 MPa should be anywhere between 1.0 through 2.0 kJ/mm. This study has clearly indicated that these ranges are too high regardless of the method which was used to make the HSS; QT, TMCP or DQ. 2. The disappearance of nitrides and carbides in the CGHAZ during welding leads to a growth of initial austenite grains. The base metal temperature in the CGHAZ exceeds 1300 °C and this causes the microstructure to change. When the temperature of the CGHAZ decreases, the stable particles that gave the base material its small homogenous microstructure have disappeared and the microstructure consequently becomes coarse. It has been clearly shown in this research that the heat input must be low, under 1.0 kJ/mm attaining a narrow CGHAZ. The CGHAZ that is susceptible to cold cracking during welding due to the coarse hardenable martensite or bainite microstructure. One way to monitor excessive heat input is to use new welding methods that have been developed by welding machine manufacturers. These methods, which lower the energy during welding, offer a new way to lower heat input and they have more features to adjust welding. Nevertheless, 154 these methods were not utilized in the course of this research, although they can offer new solutions for welding HSSs. 3. Elongation at break in all the HSS welded structures was too low when compared to the standards of these steels. Values of 6.1 % were observed when heat input was 1.0 kJ/mm, 7.1 % when heat input was 1.3 kJ/mm and 7.0 % when heat input was 1.7 kJ/mm. These values are only half of the required 15 % necessary for HSSs. Big differences between the yield strengths of the weld and base materials meant that most of the yielding occurred in the weld. The same situation occurs in real welded structures when using undermatched filler material, main yield will happen in the weld. Designers of steel structures must consider that the majority of the yielding will happen in the weld. 4. The tensile strength of welded structures was good. Although the tensile strength of filler material was only 72 % of base material tensile strength, some welded structures had near the same tensile strength as the base material. The average value of weakness was 15.4 %, when heat input was 1.0 kJ/mm, 16.3 % when heat input was 1.3 kJ/mm and 18.7 % when heat input was 1.7 kJ/mm. Additionally, the tensile strength of welded structure was 25.2 % when heat input was 1.0 kJ/mm, 23.9 % when heat input was 1.3 kJ/mm and 20.5 % when heat input was 1.7 kJ/mm, higher than the tensile strength of the filler material. The fusion zone has experienced mixing during the welding process, most likely involving the mixture of alloying elements that make inclusions, such as carbides and nitrides. 5. Using undermatched filler material when welding HSSs with a yield strength of 700 MPa is a workable method. There are many benefits to using this method as have been previously discussed. Planning ahead careful welding is the best guarantee to ensure a good final result when welding HSSs with a yield strength of 700 MPa. Undermatched filler material survives as filler metal, too. 6. However, if the steel structure is loaded in low temperatures, from -20 °C to -40 °C, then the CGHAZ could be the place from which failure can occur. The CGHAZ is near the fusion zone and there is always undercut 155 between the weld and base material. The undercut is the initial crack near the weakest zone of the HAZ. It is important to repair it within the structure through grinding and polishing. This is important if the welded structure is dynamic loaded. 156 10. FUTURE WORK Understanding the use of matching filler material in welded structure is important. There are still many structures where the behaviour of the strength and ductility of welded structures using matching filler material need to be clarified. Of most importance will be research that looks into how the structure will behave using different heat inputs and t5/8 cooling time. Research using different steels made by TMCP, QT and DQ method with matching filler material is also needed. HSSs use has been growing in the steel industry. Many of those products will be in use in the winter in Arctic areas. It would be important to clarify more behaviours of welded HSS structures in -40 °C and -60 °C temperatures. All tests should be conducted at these lower temperatures to ensure that HSS will be able to endure in the demanding Arctic area. In this study, the CTOD test was only implement in the CGHAZ. It would be important to test all HAZ zones to test if the hypothesis that the CGHAZ is the weakest zone in the HAZ. This is important when QT steels and TMCP steels are in service in the same steel structure. Only two impact ductility tests have made in this study. There were such a big range of values that the mean values of some HAZs were not the real impact ductility value. To make sure which is the real impact ductility mean value in the HAZ more tests must be done. Together with CTOD tests it will give the best estimation of the structure. In this study micro photography has also been done. However, TEM testing of the microstructure gives a better description of the microstructure. Using this method would be an easy way to clarify content of inclusions, such as carbides, nitrides and carbonitrides. Additionally, the size and shape of inclusions could be clarified through TEM testing. 157 11. SUMMARY In this doctoral thesis the usability of HSS in welded structures has been researched. Welded QT, TMCP and DQ HSSs have been under examination. The use of these HSSs grows in many industrial areas and the need for knowledge of these steels structures manufacturing is in high demand. Today, HSS is manufactured using three different methods, QT, TMCP and DQ. The microstructure of these steels and HAZ area after welding, mechanical properties, usability, and other main discrepancies in the welded structures were researched. Only after carefully clarifying the research topic and discusses welded high strengths structures was experimental research done using different laboratory tests. These tests were all conducted with undermatched filler material with three different heat inputs, 1.0, 1.3 and 1.7 kJ/mm. The research carried out during this doctoral thesis had four key findings. 1) A clear implication of this study points out that when welding HSS, thickness 8 mm and butt joint, the heat input must be 1.0 kJ/mm or lower. HSS steels with a heat input of 1.0 kJ/mm have better HAZ microstructures and additionally superior tensile strength and impact test values than steels with a heat input of 1.3 or 1.7 kJ/mm. 2) When welding all three types of HSS (QT, DQ and TMCP), the CGHAZ was very brittle. This brittleness occurred because of the high heat input used during the welding process causes dissolve of carbides and nitrides and also growing of initial austenite grains. The CGHAZ is narrow using low heat input and in normal steels structures it does not significantly weaken the structure. However, if the steel structure is loaded in low temperatures, from -20 °C to -40 °C, then the CGHAZ could be the area from which failure can occur. 158 3) The tensile strength of the welded structures was acceptable. Although the tensile strength of the filler material was only 72 % of base material’s tensile strength, some welded structures had nearly the same tensile strength as the base material. 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Prequalified Welding Procedure Specification, pWPS, heat input 1.0 kJ/mm HSS, yield strength 700 MPa Base Materials 8 mm Thickness Type of joint preparation Outside diameter of pipe Weld pass sequence V-Groove 60 ° Welding process MAG Welding position PA Groove preparation machining Groove cleaning Fastening Edge fastener Accesory equipment Tack welding Back gouging Non Backing ring Fiberglass tape Electrode Cutting-edge angle 0° Filler material and shielding gas Classification of filler material Torch angle EN 440 SFA/AWS A5.18 Angle of tilt Distance from workpiece Working temperature Trade mark of filler material OK AUTROD 12.51 Elevated working temperature 20 °C Interpass temperature Powder Preheating temperature Ar + 15 % CO2 Shielding gas Measuring of working temperature 16 l/min Flow rate range Post-welding heat treathment Plasma gas Method Flow rate range Heating rate Backing gas Soaking temperature Flow rate range Soaking time DC Type of current Cooling rate + pole Polarity Finishing Notes: Backing ring was woven glass Bead Welding Process Filler material Ø 1 MAG 2 MAG Customer Date and author: 07.04.2009 MPirinen Flow rate range (A) Arc voltage range (V) Welding speed range (mm/min Wire feed range (m/min) Heat input range ( kJ/mm) Length of free wire ( mm ) Oscillation frequence ( Hz ) Amplitude ( mm ) 1.2 220225 22.3 243 5.8 1.0 15 - - Measured values Kemppi Data pro DLI10 1.2 225230 25.5 275 6.8 1.03 19 - - Measured values Kemppi Data pro DLI10 Accepted 7.4.2009 Markku Pirinen 169 Notes! Appendix 2. Prequalified Welding Procedure Specification, pWPS, heat input 1.3 kJ/mm HSS, yield strength 700 MPa Base Materials 8 mm Thickness Type of joint preparation Outside diameter of pipe Weld pass sequence V-Groove 60 ° Welding process MAG Welding position PA Groove preparation machining Groove cleaning Fastening Edge fastener Accesory equipment Tack welding Back gouging Non Backing ring Fiberglass tape Electrode Cutting-edge angle 0° Filler material and shielding gas Classification of filler material Torch angle EN 440 SFA/AWS A5.18 Angle of tilt Distance from workpiece Working temperature Trade mark of filler material OK AUTROD 12.51 Elevated working temperature 20 °C Interpass temperature Powder Preheating temperature Ar + 15 % CO2 Shielding gas Measuring of working temperature 16 l/min Flow rate range Post-welding heat treathment Plasma gas Method Flow rate range Heating rate Backing gas Soaking temperature Flow rate range Soaking time DC Type of current Cooling rate + pole Polarity Finishing Notes: Backing ring was woven glass Date and author: 07.04.2009 MPirinen Bead Welding Process Filler material Ø Flow rate range (A) Arc voltage range ( V) Welding speed range (mm/min Wire feed range (m/min) Heat input range ( kJ/mm) Length of free wire ( mm ) Oscillation frequence ( Hz ) Amplitude ( mm ) 1 MAG 1.2 220-225 22.3 243 5.8 1.0 15 - - Measured values Kemppi Data pro DLI10 2 MAG 1.2 260270 29.0 270 8.0 1.35 19 - - Measured values Kemppi Data pro DLI10 Customer Accepted 7.4.2009 Markku Pirinen 170 Notes! Appendix 3. Prequalified Welding Procedure Specification, pWPS, heat input 1.7 kJ/mm HSS, yield strength 700 MPa Base Materials 8 mm Thickness Type of joint preparation Outside diameter of pipe Weld pass sequence V-Groove 60 ° Welding process MAG Welding position PA Groove preparation machining Groove cleaning Fastening Edge fastener Accesory equipment Tack welding Back gouging Non Backing ring Fiberglass tape Electrode Cutting-edge angle 0° Filler material and shielding gas Classification of filler material Torch angle EN 440 SFA/AWS A5.18 Angle of tilt Distance from workpiece Working temperature Trade mark of filler material OK AUTROD 12.51 Elevated working temperature 20 °C Interpass temperature Powder Preheating temperature Ar + 15 % CO2 Shielding gas Measuring of working temperature 16 l/min Flow rate range Post-welding heat treathment Plasma gas Method Flow rate range Heating rate Backing gas Soaking temperature Flow rate range Soaking time DC Type of current Cooling rate + pole Polarity Finishing Notes: Backing ring was woven glass Date and author: 07.04.2009 MPirinen Bead Welding Process Filler material Ø Flow rate range (A) Arc voltage range ( V) Welding speed range (mm/min Wire feed range (m/min) Heat input range ( kJ/mm) Length of free wire ( mm ) Oscillation frequence ( Hz ) Amplitude ( mm ) 1 MAG 1.2 220-225 22.3 243 5.8 1.0 15 - - Measured values Kemppi Data pro DLI10 2 MAG 1.2 260270 30.9 230 7.6 1.75 19 - - Measured values Kemppi Data pro DLI10 Customer Accepted 7.4.2009 Markku Pirinen 171 Notes! Appendix 4. Table 1 Characteristic of nonmetallic inclusions (Ramirez 2008). Inclusion 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 Inclusion Characteristics Chemical Composition Region A — 50.1O-0.7Mg-1.6Al-3.9Si-2.8S-19.6Ti-21.4Mn Region B — 48.2O-0.9Mg-1.6Al-3.4Si-2.3S-22.2Ti-21.4Mn 51.4O-1.4Al-4.5Si-1.7S-18.1Ti-22.8Mn Region A — 32.2O-0.5Al-1.3Si-0.9S-51.4Ti-13.7Mn (Ti-O2) Region b MnS, Region c Ti-Oxide Region A — 32.3O-1.5Al-0.7Si-50.4Ti-15.1Mn Region B — 35.4O-3.2Al-6.1Si-0.8S-26.5Ti-28.0Mn Region C — 35.3O-4.4Al-9.6Si-1.4S-3.6Ti-45.8Mn 30.9O-1.8Si-26.5S-3.5Ti-37.3Mn 56.2O-1.3Al-5.5Si-2.1S-15.4Ti-19.5Mn 77.8O-0.9Si-1.3S-17.2Ti-2.8Mn 65.5O-0.5Si-1.4S-22.8Ti-9.8Mn 65.4O-2.5Si-13.0S-16.0Ti-3.1Mn 67.9O-3.5Si-4.4S-21.1Ti-3.1Mn 73O-1.9Al-6.9Si-1.0S-14.6Ti-2.7Mn 55.1O-4.0Al-17.6Si-1.6S-3.6Ti-18.2Mn 55.9O-4.2Al-17.6Si-1.8S-2.4Ti-18.1 Mn Region A — 57.9O-4.6Al-17.4Si-1.9S-2.8Ti-15.5Mn Region B — 60.2O-1.7Al-2.2Si-0.6S-24.5-10.8 33.7O-2.3Al-15.4Si-3.5S-6.5Ti-38.7Mn 53.7O-5.0Al-17.6Si-2.0S-4.1Ti-17.6Mn 68.6O-0.9Al-15.6Si-1.5S-2.9Ti-10.5Mn 80.6O-0.7Al-14.0Si-2.1S-2.6Ti Region A — 49.9O-10.9Si-1.1S-12.0Ti-26.2Mn Region B — 49.3O-13.4Si-3.8S-3.9Ti-29.6Mn 12.1O-1.2Si-32.9S-53.8Mn 62.0O-9.8Si-0.7S-10.5Ti-17.0Mn 56.8O-1.9Al-16.0Si-2.1S-2.2Ti-21.1Mn 47.0C-14.4N-10.9O-1.2Mg-2.0Al-24.5Zr Zr Region A — 23.0N-1.9Mn-7.9Al-66.6Zr-0.7Ti; Region B — 39.9N-23.4O-1.0Mg-30.3Al-5.4Zr 45.4C-14.6N-15.6O-0.8Al-23.9Zr 40.3C-13.0N-13.4O-1.7Mg-2.8Al-28.7Zr Region A — 40.4O-10.9Mg-23.0Al-25.7Zr Region B — 48.9O-15.0Mg-36.1Al Region A — 20.8N-33.2O-1.7Mg-1.6Al-42.7Zr Region B — 79.5O-20.5Zr Region A — 18.9N-29.2O-2.95Mg-3.0Al-46.0Zr Region B — 17.2N-40.8O-3.8Mg-13.9Al-24.3Zr 11.2N-50.6O-14.2Mg-19.6Al-4.4Zr 62.7O-3.4Mg-2.0Al-31.92Zr 56.0O-3.8Mg-29.5Al-10.8Zr 63.7O-36.3Si 59.3O-13.2Al-9.0.Si-6.1Ti-12.4Mn 65.0O-10.0Al-5.9Si-6.7Ti-12.5Mn 59.5O-10.4Al-13.8Si-2.3Ti-14.0Mn 63.7O-5.3Al-5.2Si-11.1Ti-14.7Mn Description O, Al, Si, S, Ti, Mn rich O, Al, Si, S, Ti, Mn rich Composite inclusion Ti-Mn oxide Mn, S, O rich Ti-Mn oxide O, Si, S, Ti, Mn rich Ti oxide O, S, Ti rich Ti Oxide O, Al, Si, Ti, Mn rich O, Si, Mn rich O, Al, Si, Mn rich Composite inclusion O, Al, Si, S, Ti, Mn O, Al, Si, Ti, Mn rich O, Al, Si, S, Ti, Mn rich O, Al, Si, S, Ti rich O, Si, S, Ti, Mn rich O, Si, S, Mn rich O, Si, S, Ti, Mn rich O, Al, Si, S, Ti, Mn rich Carbo-Nitride - Al2O3 Composite inclusion Zr Carbo-Nitride Zr Carbo-Nitride Composite inclusion Composite inclusion Composite inclusion Composite inclusion O, Mg, Al, Zr rich O, Al, Mg, Zr rich SiO2 O, Al, Si, Ti, Mn rich O, Al, Si, Ti, Mn rich O, Al, Si, Ti, Mn rich O, Al, Si, Ti, Mn rich ACTA UNIVERSITATIS LAPPEENRANTAENSIS 473. 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