metallurgical characterisation of simulated heat

METALLURGICAL CHARACTERISATION OF SIMULATED HEAT AFFECTED ZONE IN BORON MODIFIED P91 STEEL Submitted in partial fulfilment of the requirements of the degree of Master of Technology in Industrial Metallurgy By MODASSIR AKHTAR (Roll no. 155513)

Supervisors: Dr. T. Jayakumar, MMED NIT Warangal Dr. Rajneesh Kumar, National Metallurgical Laboratory Jamshedpur

DEPARTMENT OF METALLURGICAL AND MATERIALS ENGINEERING NATIONAL INSTITUTE OF TECHNOLOGY WARANGAL - 506004 (June – 2017)

National Institute of Technology Warangal

National Institute of Technology Warangal

2017

Metallurgical Characterisation of Simulated Heat Affected Zone in Boron Modified P91 Steel Modassir Akhtar National Institute of Technology Warangal

Recommended citation A. Modassir, Metallurgical Characterisation of Simulated Heat Affected Zone in Boron Modified P91 Steel, M.Tech. dissertation, Department of Metallurgical and Materials Engineering, National Institute of Technology Warangal, 2017.

To the loving memory of Prophet Mohammad (PBUH)

National Institute of Technology Warangal – 506004

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Declaration

I declare that this written submission represents my ideas in my own words and where others' ideas or words have been included, I have adequately cited and referenced the original sources. I also declare that I have adhered to all principles of academic honesty and integrity and have not misrepresented or fabricated or falsified any idea/data/fact/source in my submission. I understand that any violation of the above will be cause for disciplinary action by the Institute and can also evoke penal action from the sources which have thus not been properly cited or from whom proper permission has not been taken when needed.

Modassir Akhtar Roll no. 155513 Date: June 2017


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Abstract Modern ferritic/martensitic 9% Cr steels are alloys with improved mechanical properties at elevated temperatures. This study is based on microstructural investigations before and after post weld heat treatment of Gleeble weld HAZ simulated subzones of boron containing modified P91 creep resistant steel (P91B) which is designed for boiler tube applications in fossil-fueled power plants. P91B steel is generally received in normalised and tempered conditions having enhanced oxidation resistance and creep resistance. Welding of this steel is quite difficult due to high carbon equivalent and the welding process results in degraded properties at heat affected zone (HAZ). The HAZ varies from a few hundred microns to a few millimetres depending upon the welding technique employed. Such narrow zone further comprises of several sub-zones which create difficulties in conducting metallurgical as well as mechanical investigations. In this work, the Gleeble thermo-mechanical simulator was employed to replicate separate sub-zones of weld HAZ i.e. CGHAZ (coarse grain heat affected zone), FGHAZ (fine grain heat affected zone), ICHAZ (inter-critical heat affected zone) and CG + FGHAZ (coarse grain and fine grain heat affected zone). Hardness for different zones of simulated weld HAZ confirm that softening arises in the region close to base metal which is ICHAZ. The degree of softening in ICHAZ makes this subzone as the weakest link which may be susceptible to Type IV cracking. In the real-time welding, preheating and PWHT are essential to minimise gradient in hardness and microstructural features during both single and multi-pass welding. Therefore, to understand the minimization of non-uniformity in both microstructure and hardness among Gleeble simulated HAZ subzones of P91B creep resistant steel, PWHT was done for 7600C/3h which resulted in tempered martensite with a uniform distribution of an optimum number of fine secondary precipitates. Keywords: CGHAZ, FGHAZ, Gleeble, ICHAZ, JMatPro, M23C6, microstructure, multi-pass welding, MX, P91B, PWHT and Thermocalc.


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Acknowledgements

All praises belong to Allah, Lord of all the worlds, Glorious and Merciful. Who sent his “beloved Prophet Mohammad (PBUH)” for the welfare of all humanity & universe. I have loving memory of my father and can’t fulfil the level of gratitude towards my mother, who she really honoured, also thankful to my elder brother and younger sister. I would like to express my honorary gratitude to distinguished Prof. T. Jayakumar and Dr. Rajneesh Kumar for giving me FOR (Freedom Of Research) and constant encouragement that allows me to think and work whichever I wanted to do. It was pleasure to be part of CSIR-NML Jamshedpur, special thanks to the Director of CSIRNML Dr. I. Chattoraj and Dr. N. Narasaiah, MMED, NITW for allowing me to complete my dissertation Apart from them, my thankful gratitude towards some more people to provide me with the seed of knowledge and efforts to compile this program like Prof. M. K. Mohan, and Dr. Ajoy Kumar Pandey and Scientists of CSIR-NML Jamshedpur Dr. Mainak Ghosh and Dr. J. Swaminathan. Special thanks to research scholar of NIT Jalandhar, Mr. Akhil Khajuria and scientist of IGCAR Kalpakkam, Dr. Saroja Saibaba. My colleagues like Manish Pandey from O.P. Jindal University, Suneel Kumar from MMED, NITW, Mayur Pratap Singh from NIT Jalandhar and V. Santosh project assistant NML Jamshedpur also need to be thanked.

Modassir Akhtar


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Table of Contents

Dissertation approval for M.Tech………………………………………………………….... ii Certificate…………………………………………………………………………………… iii Declaration............................................................................................................................... iv Abstract ..................................................................................................................................... v Acknowledgements .................................................................................................................. vi Table of Contents .................................................................................................................... vii List of figures ............................................................................................................................ x List of tables ........................................................................................................................... xiv Abbreviations .......................................................................................................................... xv Chapter 1: Introduction............................................................................................................ 1 1.1: Background and motivation ....................................................................................... 1 1.2: Objective of this study ................................................................................................ 2 1.3: Thesis Scope ............................................................................................................... 2 1.4: Outline of this thesis ................................................................................................... 3 Chapter 2: Literature review .................................................................................................... 5 2.1: History and codes ......................................................................................................... 5 2.2: Physical Metallurgy ..................................................................................................... 6 2.2.1: Martensite ................................................................................................................ 6 2.2.2: Heat treatments ........................................................................................................ 7 2.3: Mechanical properties................................................................................................ 10 2.3.1: Factors affecting Creep strength ............................................................................ 10 2.4: Key elements for Alloy design of P91B creep resistant steel .................................. 10 2.5: HAZ ............................................................................................................................. 14 2.6: Cracking ...................................................................................................................... 15 2.6.1: Recent failures ....................................................................................................... 16 2.6.2: Type IV cracking ................................................................................................... 16


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2.7: Strengthening mechanism ......................................................................................... 18 2.7.1: Solid solution strengthening .................................................................................. 18 2.7.2: Dislocation strengthening ...................................................................................... 19 2.7.3: Grain/subgrain boundary strengthening ................................................................ 19 2.7.4: Precipitation strengthening .................................................................................... 19 Chapter 3: Experimental Set-up and Methodologies ............................................................ 26 3.1: As Received Material (AR) of P91B steel ................................................................. 26 3.2: Gleeble Simulation and PWHT ................................................................................. 26 3.2.1: HAZ Simulation by Gleeble 3800 ......................................................................... 26 3A.2: Post Weld Heat Treatment (PWHT) ...................................................................... 28 3.3: Hardness testing ......................................................................................................... 28 3.4: X-ray diffraction (XRD) ............................................................................................ 30 3.5: Metallography ............................................................................................................ 31 3.5.1: Optical Microscopy (OM) ..................................................................................... 31 3.6.2: Scanning electron microscopy and Energy Dispersive Analysis of X-ray spectroscopy .................................................................................................................... 32 3.6.3: Electron backscatter diffraction (EBSD) ............................................................... 32 3.6: Thermodynamic and kinetic modelling.................................................................... 32 3.6.1: JMatPro (Java-based Materials Property) ............................................................. 32 3.6.2: Thermocalc (Thermodynamic Calculations) ......................................................... 33 Chapter 4. Experimental Results and Discussions................................................................ 34 4.1. Thermal profiles of Gleeble simulated HAZ ............................................................ 34 4.2: Hardness evaluation ................................................................................................... 36 4.2.1: Bulk hardness on surface of the sample ................................................................ 36 4.2.2: Micro-hardness at cross section of the sample ...................................................... 38 4.3: Calculation of dislocation density by XRD .............................................................. 43 4.4: Microstructural assessment ....................................................................................... 45 4.4.1: By OM before and after PWHT treatment ............................................................ 45 4.4.2: PAG size estimation through OM microphotographs ........................................... 50 4.4.3: SEM microstructures analysis before and after PWHT treatment ........................ 52


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4.4.4: Precipitates finding and its analysis before and after PWHT ................................ 57 4.4.5: EBSD of As received material .............................................................................. 59 4.5: Computational diagrams by thermodynamic and kinetic calculations ................ 63 4.5.1: CCT and other diagrams by JMatPro .................................................................... 63 4.5.2: Calculation of thermodynamic diagrams by Thermocalc ..................................... 63 4.6: Other tests ................................................................................................................... 74 4.6.1: Impact energy for CGHAZ D04 ............................................................................ 74 4.6.2: Dilatometry for Ac1, Ac3, Ms and Mf .................................................................... 74 4.6.3: Furnace heat treatment of P91B steel .................................................................... 77 4.7: Effect of peak temperature on microstructure ........................................................ 79 4.7.1: AR microstructure interpretation ........................................................................... 79 4.7.2: CGHAZ samples ................................................................................................... 79 4.7.3: FGHAZ samples .................................................................................................... 80 4.7.4: ICHAZ samples ..................................................................................................... 81 4.7.5: CG + FGHAZ D01 ................................................................................................ 81 4.9: Boron effect on microstructure ................................................................................. 82 Chapter 5: Summary and Conclusions .................................................................................. 83 Future work and scope ..................................................................................................... 85 Bibliography and references .................................................................................................. 87 Appendices ................................................................................................................................ 1 Brief introduction to HAZ simulation ............................................................................... 1 Metallographic preparations and microstructural analysis by Light and Electron Microscopy ........................................................................................................................... 1


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List of figures Figure 2.1: Plot of (a) CO2 reduction with an increase in thermal efficiency and (b) comparison of thickness reduced of P91 steel with other steels [3]. ............................................................ 5 Figure 2.2: Evaluation of ferritic steel with 105 h creep rupture strength at 600 oC [8]. ........... 5 Figure 2.3: CCT curve for P91 steel [5]. ................................................................................... 8 Figure 2.4: Micro-hardness variation of welded sample reported by Wang et al. [34]. ......... 10 Figure 2.5: Schematic representations of (a) single pass and (b) multipass welded bead [41].14 Figure 2.6: Several types of cracking in a weld joint [3, 42]. ................................................. 15 Figure 2.7:

Microstructure of different precipitates which chiefly observed in 9-12%Cr

containing steels [22]. .............................................................................................................. 22 Figure 3.1: Plate of P91B steel received in normalised and tempered conditions……………..27 Figure 3.2: Gleeble® 3800 simulator, left side is digital control console, center is Gleeble load unit and the right one mobile conversion unit. ........................................................................ 27 Figure 3.3: Test specimen of length 78mm with 11X11 mm2 cross section after simulation and mid portion showing thermocouple weld. ............................................................................... 27 Figure 3.4: (a) Vickers macro-hardness testing machine, Economet VH-50MD and (b) Vickers micro-hardness testing machine, Leica VMHT Auto. ............................................................. 29 Figure 3.5: Schematic diagrams of (a) Principle diagram of Vickers hardness and (b) macroindentation tested surfaces. ...................................................................................................... 29 Figure 3.6: Leica DM 2500M Optical Microscope. ................................................................ 31 Figure 4.1: Gleeble thermal profiles of single and multipass welding HAZ simulation of P91B steel…………………………………………………………………………………………...35 Figure 4.2: Hardness plots of various simulated samples including AR material before PWHT. Whereas SL11= hardness profile of the first surface with the first line, SL12= hardness profile


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of the first surface with line 2nd, SL21= hardness profile of the second surface with the first line and SL22=hardness profile of the second surface with the second line. .......................... 40 Figure 4.3: (a) Summary of micro-hardness by taking standard deviation as an error bar and (b) to (n) contour mapping of micro-indentation testing with average, standard deviation and hardness colour bar indicator. .................................................................................................. 42 Figure 4.4: XRD pattern of (a) AR material and HAZ simulated samples before PWHT, (b) high-resolution scanning of AR material, CGHAZ D05, FGHAZ D02, ICHAZ D01 and CG + FGHAZ D01 and (c) HAZ simulated samples after PWHT. .................................................. 44 Figure 4.5: AR material (BM) etched with 10% aqua regia in distilled water. ....................... 46 Figure 4.6: Microstructures from the optical microscope of various samples prior to PWHT with their corresponding etchants. ........................................................................................... 49 Figure 4.7: Light microscope micrographs after PWHT, etched with 5–7 % nital + Vilella’s reagent of 4 g picric acid to CGHAZ D05 at (a) 50X and (b) 1000X, FGHAZ D02 at (c) 50X and (d) 1000X, ICHAZ D01 at (e) 50X and (f) 1000X and CG + FGHAZ D01 at (g) 50X and (h) 1000X................................................................................................................................. 50 Figure 4.8: Grain size distributions in the form of the histogram with normal distribution for different samples...................................................................................................................... 52 Figure 4.9: SEM microstructures of various samples prior to PWHT, AR material (a) and (b), simulated samples like CGHAZ D01 (c)and (d), CGHAZ D05 (e) and (f), FGHAZ D02 (g) and (h), ICHAZ D01 (i) and (j), ICHAZ D03 (k) and (l) and CG + FGHAZ D01 (m) and (n). .... 55 Figure 4.10: SEM microstructures of various samples after post weld heat treatment at 760 oC, 3 h, CGHAZ D05 (a) and (b), FGHAZ D2 (c) and (d), ICHAZ D01 (e) and (f) and CG + FGHAZ D01 (g) and (h). ....................................................................................................................... 56 Figure 4.11: EDAX plots of (a) Fe-rich M23C6 carbide (b) Cr-rich M23C6 carbide and (c) Mo6C carbide. .................................................................................................................................... 57 Figure 4.12: SEM microphotographs showing several types of precipitates before PWHT of samples (a) AR material, (b) CGHAZ D05 (c) FGHAZ D02 (d) ICHAZ D1 (e) ICHAZ D03 and (f) CG + FGHAZ D01. ..................................................................................................... 58 National Institute of Technology Warangal – 506004

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Figure 4.13: SEM microphotographs displaying several types of precipitates after PWHT of at 760 0C/3h (a) CGHAZ D05, (b) FGHAZ D02 (c) ICHAZ D01 and (d) CG + FGHAZ D01. 59 Figure 4.14: Inverse pole figure (IPF) in Z [001] direction (perpendicular to the screen) shows the grain orientation distribution for AR material. .................................................................. 60 Figure 4.15: Grain size distribution of AR material in terms of area fraction. ....................... 60 Figure 4.16: Distribution of high angle and low angle grain misorientation for AR material. 62 Figure 4.17: Distribution of KAM for AR material. ............................................................... 62 Figure 4.18: JMatPro diagrams. .............................................................................................. 66 Figure 4.19: Thermocalc diagrams for P91B steel, (a) Fe-C diagram, (b) Fe-Cr diagram, (c) FeMo diagram, (d) Fe-V diagram, (e) Fe-N diagram and (f) Fe-N diagram. .............................. 67 Figure 4.20: Binary diagram of P91B steel where 1* region represents no boron nitride. ..... 67 Figure 4.21: Property diagram and amount of components in a phase diagrams for P91B steel by Thermocalc. ........................................................................................................................ 69 Figure 4.22: Mean radius (bulk nucleation sites) of various precipitates at different PWHT conditions. ............................................................................................................................... 70 Figure 4.23: Critical radius of various precipitates at different PWHT conditions. ................ 71 Figure 4.24: Driving force of various precipitates at different PWHT conditions. ................. 72 Figure 4.25: Number density of various precipitates at different PWHT conditions. ............. 73 Figure 4.26: Volume fraction of various precipitates at different PWHT conditions. ............ 74 Figure 4.27: Plots of dilatometry measurements and thermal histories. ................................. 76 Figure 4.28: Thermal profiles of (a) FGHAZ D04 and (b) ICHAZ D04. ............................... 76 Figure 4.29: Microstructure of (a) FGHAZ D04 etched with a mixture of 28% nital and Vilella’s reagent and (b) ICHAZ D04 etched with 28% nital. ............................................................... 77 Figure 4.30: Thermal profile given to P91B steel during furnace treatment. .......................... 77 National Institute of Technology Warangal – 506004

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Figure 4.31: Microphotographs of furnace treated samples (a) and (b) 865 oC/1h, (c) and (d) at 1240 oC/1h, micro-indentation of (e) soft phase and (f) hard phase........................................ 79 Figure A1.1: (a) Thermocouple welder (b) Thermocouple Push Posts that used to couple welded thermocouple to generate thermal/mechanical profiles and (c) weld beads that permit correct simulation……………………………………………………………………………...4 Figure A1.2: Thermal profile of AISI 1018 steel with respect to different grips. ..................... 5 Figure A2.1: Microstructures of non-simulated region (a) CGHAZ D01 etched with 4% picral, (b) CGHAZ D04 etched with 28% nital, (c) CGHAZ D05 etched with Vilella’s etchant, (d) FGHAZ D02 etched with 28% nital, (e) FGHAZ D02 etched with Marshall reagent, (f) FGHAZ D02 etched with Vilella’s reagent, (g) FGHAZ D02 etched with 6% picral, (h) FGHAZ D03 etched with 6% picral, (i) ICHAZ D01 etched with 6% picral and (j) ICHAZ D03 etched with Vilella’s reagent………………………………………………………………………………..3


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List of tables

Table 2.1: International specifications for 9Cr-1Mo ferritic/martensitic steel [20]. ................. 6 Table 3.1: Chemical composition of P91B steel in wt.%.........................................................26 Table 4.1: Thermal histories of Gleeble Simulated HAZs. Where, EHT= Exact holding time..36 Table 4.2: Dislocation densities of both AR material and simulated HAZ samples before and after PWHT.............................................................................................................................. 43 Table 4.3: Amount of RA calculated from classic Koistenen and Marburger equation. ........ 45 Table 4.4: (a) PAG distribution in terms of area fraction and (b) misorientation angle in terms of number fraction. .................................................................................................................. 60 Table 4.5: Dilatometry results for P91B steel. ........................................................................ 75 Table A1.1: Comparison of heating methods based on different testing variables……………..3 Table A2.1: Effects of different etchants tried to reveal microstructural features on Gleeble simulated weld HAZ samples of P91B steel ............................................................................. 4


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Abbreviations Ac1= Lower critical temperature Ac3= Upper critical temperature AR = As received ASME = American society of mechanical engineers ASTM = American society for testing and materials CCT = Continuous cooling transformations CG + FGHAZ = Coarse grain and fine grain heat affected zone also known as super critically reheated coarse-grained Heat Affected Zone CGHAZ = Coarse grain heat affected zone CTE = Coefficient of thermal expansion EBSD = Electron backscattered diffraction FGHAZ = Fine-grained heat affected zone HAZ = Heat affected zone GB = Grain boundary ICHAZ = Inter-critical heat affected zone Ms = Martensite start temperature Mf = Martensite finish temperature OM = Optical microscopy PAG = Prior austenite grain PAGB = Prior austenite grain boundaries PWHT= Post weld heat treatment RA= Retained austenite SEM = Scanning electron microscope SGB= Sub-grain boundary National Institute of Technology Warangal – 506004

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Chapter 1: Introduction 1.1: Background and motivation Continuous production of CO2 gases from fossil fueled power plants has an adverse effect on the environment. So to minimize CO2 emissions and to increase operating temperature and pressure of thermal power plants, resulting in increased thermal efficiency and to decrease per capita cost in terms of ₹/kWh, there is a need to develop new generation ferritic/martensitic steels which can withstand high thermo-mechanical conditions in thermal power units [1, 2, 3, 4]. Grade 91 steel is ferritic/martensitic steel and was originally developed by Oak Ridge National Laboratory (ORNL), USA in the 1970s, for fast breeder reactors [5]. Ferritic/martensitic steels have high thermal conductivity, low thermal expansion and high resistance to thermal fatigue than austenitic steels but poor creep strength. Though sufficient strengthening has been achieved in P91 grade steel by the addition of Nb and V for high temperature pressure components to enhance creep resistance, research on aspects like an alloy development, heat treatment and via different precipitations etc. is still going on [6, 7]. Now, modified P91 steel is also being used in supercritical and ultra-supercritical plants as it has good thermo-mechanical properties like elevated temperature mechanical properties, weldability and fabricability etc. To increase thermal efficiency beyond 46 % for advanced ultra super critical power plants with a decrease in CO2 emissions at working temperature of 700 oC, austenitic steels and nickel base alloys are being preferred but at low-temperature zones still, P91 steels are being used [8, 9]. The high heat input from the fusion welding of power plant components for both similar and dissimilar joints produces heat affected zone (HAZ) near the fusion line with evolved complex microstructures, different creep strengths and detrimental mechanical properties across the zone [10, 11]. Based on the order of peak temperature and its corresponding phase transformation, HAZ is further classified into four subzones from the fusion line towards the base metal (BM) i.e. CGHAZ, FGHAZ, ICHAZ and over tempered zone. The formation of FGHAZ and ICHAZ introduces metallurgical notch resulting in high triaxiality which promotes cavitation at higher service temperatures [8, 12]. Such cavities coalesce to form creep voids over to due course of service exposure and thus the weldment ruptures by a phenomenon referred as type IV cracking [8]. However, the countries like India and China where large numbers of power generation units are recently built to fulfil energy demands, thus there is an extensive scope of monitoring, detecting, inspecting and repairing of type IV cracking [3]. The addition of controlled boron with low levels of nitrogen to modified

P91 steel seems to provide an alternative for enhancing creep strength across the weld HAZ due to increase in coincident lattice sites thereby increasing the targeted design life of components like boiler headers, tubes, turbine rotors etc. Segregation of controlled boron at prior austenite grain boundary (PAGB) after normalising is known to provide grain as well as subgrain boundary (SGB) hardening by retarding the coarsening of M23C6 carbides, delaying tertiary creep and refining MX precipitate [8, 4, 13, 14, 15, 16]. It is also understood that 90-130 ppm levels of boron are helpful in enhancing creep rupture life of P91 steel welds but excess boron can form boro-nitrides during normalising. However, mechanisms pertaining to such improvements and mitigation of type IV cracking in P91 steel weld HAZ on the addition of boron (P91B steel) are still unclear [17]. Additionally, weld HAZ contains complex microstructures due to phase transformations, dissolution of precipitates and recrystallization in a confined region depending on the welding process employed thus affecting the hardness uniformity across the HAZ. Hence, post-weld heat treatment (PWHT) at 740-780 o

C for 2 to 3 hours is recommended for the P91 steel weld to minimise hardness variation and

microstructural gradient over the weld HAZ [10, 18]. The preparation of specimens for mechanical testing from this narrow region of weld HAZ is a challenging task. To replicate the real-time welding conditions on a larger volume and to optimise the welding process parameters by optimising the number of samples, thermo-mechanical simulator like Gleeble is an attractive choice. The objective of this work is to understand the non-uniformity raised up in both microstructure and hardness among different Gleeble simulated subzones of HAZ for P91B creep enhanced ferritic/martensitic steel as well as its minimization by post weld heat treatment.

1.2: Objective of this study 1.2.1: To study different zones of HAZ due to a small region (ranging from a few hundred microns to a few millimetres) formed during actual welding. 1.2.2: To evaluate the effect of different peak temperatures with the same cooling rate on microstructure, precipitates and hardness. 1.1.3: Minimization of hardness variation and microstructural gradient among different simulated HAZ zones by post-weld heat treatment (PWHT).

1.3: Thesis Scope 1.3.1: Understanding the reproduction of weld HAZ subzones during single-pass and multipass Gleeble simulations for creep resistant steels.


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1.3.2: Checking if cooling rate between 800 oC to 500 oC is affecting weld HAZ simulations among different weld HAZ samples and is high enough to produce a fully martensitic structure in P91B steel. 1.3.3: To confirm minimization of hardness variations and microstructural gradient by PWHT for single- pass weld HAZ simulation. 1.3.4: To test out if P91B steel weld HAZ simulation samples underwent uniform grain growth and microstructural stability due to memory effect. 1.3.5: To understand if steel grips are not ideal cooling grips to simulate weld HAZ subzones due to their both poor thermal conductivity and coefficient of thermal expansion. 1.3.6: To check if Gleeble simulated samples are non-homogeneously cooled during cooling part of the Gleeble thermal cycle.

1.4: Outline of this thesis

Chapter 1 and 2

INTRODUCTION AND LITERATURE REVIEW These sections focus on motivation and theoretical background of existing study. It also highlights main problems being addressed so far in literature for P91 steel during service conditions. Furthermore, objectives have been defined.


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Chapter 3 and 4

EXPERIMENTATION, RESULTS AND DISCUSSIONS These sections deal with selection and execution of characterization methods to process the material in order to accomplish productive outcomes. Techniques like hardness testing, optical microscopy (OM), scanning electron microscopy (SEM) with energy dispersive analysis of X-rays spectroscopy (EDAX), X-ray diffraction (XRD),

electron backscattered diffraction (EBSD) and thermodynamic calculation by JMatPro and Thermocalc are briefly given. It also provides systematic assessment and reasonable explanation for the current and future work. Here, hardness and microstructure are correlated with different peak temperatures. The validation of newly formed precipitates and other phases has been done by Thermocalc.

Chapter 5

CONCLUSIONS, REFERENCES AND APPENDIX This chapter represents important conclusions drawn from discussions of this study on P91B steel. Appendices cover the basic understanding of Gleeble and metallographic investigations for P91B.


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Chapter 2: Literature review 2.1: History and codes

Figure 2.1: Plot of (a) CO2 reduction with an increase in thermal efficiency and (b) comparison of thickness reduced of P91 steel with other steels [3].

Figure 2.2: Evaluation of ferritic steel with 105 h creep rupture strength at 600 oC [8]. National Institute of Technology Warangal – 506004

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Figure 2.1 (a) shows the CO2 reduction with increased thermal efficiency for supercritical to ultra-supercritical (USC) power plants. USA has been designing USC power plants for the service temperature of 760 oC and Japan is already converted their supercritical plants to USC plants for an operating temperature of 700 oC. Hot tensile strength and creep strength of P91 steel reduce the thickness of pipes compared to P22 and X20 as shown in Figure 2.1 (b) [3 cf 94]. The growth of Cr - containing steels for use up to 600 oC temperature for a rupture life of hundred thousand hours is shown in Figure 2.2. First high chromium ferritic steel was developed in Europe in the mid-1960s [4]. Minor changes in the composition of 9Cr steels were made to fulfil radiation resistance [19]. The international codes of 9Cr steel for different applications are shown in Table 2.1. Table 2.1: International specifications for 9Cr-1Mo ferritic/martensitic steel [20]. Country

Standard

Grade

USA

ASTM/ASME A 387

P91 (plate)

USA

ASTM/ASME A 213

T91 (seamless tube)

USA

ASTM/ASME A 335

P91 (seamless pipe)

USA

ASTM/ASME A 182

F91 (forged fittings)

USA

ASTM/ASME A 217

C12A (casting variants)

USA

ASTM/ASME A 234

WP91 (wrought piping fittings)

USA

ASTM/ASME A 336

F91 (forging)

USA

ASTM/ASME A 369

FP91 for forged and bored pipes

UK

BS 1503

Grade 91 (forging)

Germany

DIN 17175

X10CrMoVNb 9 1

Japan

----

HCM-9S

France

NF A 49213

T U Z 10CDVNb 09-01

2.2: Physical Metallurgy 2.2.1: Martensite Martensite is a supersaturated carbon in ferrite, the as-quenched structure is body-centered tetragonal [15, 22], with lattice parameters a and c, have a relation as follows [23]. c/a =1.00 + 0.045 *Wt.%C … Equation 2.1 [23]


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Relation in lattice constant and microstrain are represented by linear Equation 2.2 [23]. 𝛽 𝑐𝑜𝑠 𝜃 1 21/2 ×𝑠𝑖𝑛 𝜃 𝛫𝜆

=𝐿 +

𝛫𝜆

… Equation 2.2 [23]

Y=C+MX Where < 𝜖 2 >1/2 is, root mean square strain. The dislocation density is the function of crystallite size and root mean square strain and is given by Equation 2.3 [24]. 𝜌 = 2√3

1⁄2 𝐷×𝑏

…Equation 2.3 [24]

Ms = 635 - 474[C+0.86{N – 0.15(Nb + Zr) – 0.066(Ta + Hf)}]-[17Cr + 33Mn + 21Mo + 17Ni + 39V + 11W] Applicable for 8-14% Cr steels …Equation 2.4 [15] 2.2.1.1: Kinetics of martensite formation in high chromium steels PAGB of martensite is made up of the packets, blocks and laths. Crystallographic packet: This is a collection of laths with nearly the same habit plane. The number of packets formed in a prior austenite grain (PAG) depends on the PAG size, and it is generally greater than 1. For a PAG size of around 10 µm, a single main packet tends to form which occupy practically all the existing space [25]. Block: It is a group of laths of roughly the same habit plane and slightly misoriented (within 10◦) parallel directions. The 6 variants in a crystallographic packet can group into pairs with low misorientation. Experimentally, such low-misoriented variants have been found close to each other for carbon contents up to 0.41 wt. %. In 0.6 wt. % of carbon contents, around low misoriented variants are generally not gathered together [25]. Sub-block: It is a stack of laths with the same habit plane and same parallel direction (i.e., all belonging to the same crystallographic variant). Most lath misorientations within the same subblock are around 3◦ with a maximum at 5◦ [25].

2.2.2: Heat treatments These steels are heated to a normalising temperature of 1040 oC - 1080oC for half to one hour followed by tempering treatment at 750 oC to 780 oC for 1 to 2 h. After normalising, air quenching is enough to produce martensite as high alloying elements shift CCT (continuous cooling transformations) curve to higher transformation time. Martensite produced from normalising has high hardness; this hardness is reduced by tempering operation to enable precipitation of finely distributed V/Nb(C,N). They give strengthening to P91 steel and National Institute of Technology Warangal – 506004

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stabilises matrix during an operating temperature [2]. Dislocation density decreases during tempering, PWHT and service conditions. Lath martensite splits into soft equiaxed subgrains due to tempering then to ferrite at service temperature, as an operating temperature realised further tempering [3]. The hardness of martensite from normalising treatment is 405 - 410 HV whereas tempering decrease it to 240 HV of hardness or less [26]. When the normalising temperature is decreased from 1080 to 1050 oC, the PAG size decreases with a decrease in yield strength and tensile strength that suggests a limitation of grain boundary (GB) strengthening [9]. The Ac1 (lower critical temperature) and Ac3 (upper critical temperature) are 800 and 890 o

C taken from reference respectively [27].

2.2.2.1: Normalizing

Figure 2.3: CCT curve for P91 steel [5]. The purpose of normalising treatment is to homogenise austenite grains and to dissolve coarse carbo-nitrides formed from the hot working of the P91 steel plate. Normal air cooling is required but if the plate/pipe section is thicker then oil or water mist quenching is required for cooling [20]. Normalising treatment dissolves completely all the precipitate but still, NbC left in small quantities that hinders grain growth [22]. As quenched martensitic steel contains RA as a thin film in nm of thickness [25]. After normalising, boron segregates toward boundary [28]. The CCT diagram for this steel is shown in Figure 2.3. 2.2.2.2: Tempering Rejected carbon from martensite has a key role in the nucleation of precipitates during tempering. Iron carbides form by minimising the chemical potential of carbon which diffuses


8

near dislocations during tempering. Entrapped carbon is a linear function of dislocation density [29]. Tempering causes precipitation at GBs, lath, block boundaries and packet boundaries [28]. Matrix

(Fe, Cr)3C

Cr7C3

Cr23C6 …Equation 2.5 [30]

Recovery of ductility is due to the reduction of dislocations [31] and subgrains formed in this treatment from lath martensite grains [6, 31]. Although lower recovery is marked in the case of chromium and molybdenum containing steels. Elements like Ni + Mn bring down the Ac1 temperature during tempering and austenite formation occurs that caused brittle untempered martensite on cooling [22, 32]. Boron enrichment after tempering in M23C6 is observed in boron containing 9Cr steels but more significant during ageing [16]. Normalising treatment is not much affecting mechanical properties at room temperature but tempering has a strong effect. High tempering temperature (especially when it has increased from 650 to 750 °C) produces large amounts of nitride precipitation in the matrix leading to enhanced toughness [33]. 2.2.2.3: Post weld heat treatment PWHT at 740 - 780 oC for 2 to 3 h is recommended for P91 steel, but in actual conditions before PWHT post weld heating at 300 - 350 oC should be given to avoid hydrogen assisted cracking (HAC). Low PWHT temperature (normally 650 – 730°C) is suggested for applications like turbine rotors to enhance tensile strength while high PWHT (750 - 780°C) is for boiler headers and tubes to improve toughness [6, 8]. It has been resolved that time for tempering is based on the application of this steel like for thin tube is limited to 30 min but in the case of turbine rotors a minimum 10 h is employed [14]. Different tempering temperature caused variation in the distribution of precipitates [6]. It produces tempered martensite which microstructurally varies along CGHAZ to over tempered HAZ and decreases residual stress [3] with coarsening of precipitates along with recovery and dynamic recrystallization of martensite laths in HAZ especially ICHAZ occurring during PWHT [34]. Redistribution of carbides and carbonitrides happens at PAGBs, SGBs [5, 34] and at dislocations with a uniform distribution of fine laths of martensite after PWHT [5] which causes the formation of carbon depleted zone (CDZ) and carbon enriched zone (CEZ). CDZ made up of ferrite grains without carbides [35]. Overtempering caused a loss in creep strength whereas, under-tempering resulted in minimised hardness between WM and HAZ [5]. Hardness must be in the range of 200 to 280 HV after PWHT treatment which may be accepted till 303 HV. After that, it represents inadequate PWHT treatment. Though hardness less than 178 HV is considered as overheated weld joint, it is recommended that selection of load to measure hardness should not produce a large depth of


9

the indentation on the surface of thin tubes, minimum tempering time is 2 h. Also, PWHT time is 1 hour per inch thickness is recommended [32].

2.3: Mechanical properties.

Figure 2.4: Micro-hardness variation of welded sample reported by Wang et al. [34]. Mechanical properties of ASTM grade 9Cr steels are yield strength = 415 MPa (min.), tensile strength = 585 to 760 MPa (min.), hardness = 265 VHN (max.) and longitudinal elongation = 20% (min.) and creep rupture strength at 625 oC is 68 MPa [7]. Hardness is the mechanical testing, that can successfully, quantitively and directly apply to the weld samples. See Figure 2.4 which is taken transverse to the sample and shows three thermal histories i.e. as welded conditions, PWHT and crept that tells PWHT minimises hardness variations and creep testing gives minimum hardness at the transition region between FGHAZ and ICHAZ [34]. Another testing requires miniature samples of specified sizes [3]. The hardness value for martensite, ferrite, austenite, TiC, NbC, VC, Fe3C Cr7C3 and Cr23C6 is 500-100, 70-190, 190-350, 28503200, 2000-2400, 2250-2950, 1200, 1400-2150 and 1000-1650 HV respectively [15].

2.3.1: Factors affecting Creep strength Composition, filler materials, joint design, welding procedure and heat treatment [34].

2.4: Key elements for Alloy design of P91B creep resistant steel Higher contents of ferritic stabilisers may increase the risk of δ-ferrite formation [19]. 2.4.1: Carbon Carbon stabilises austenite, helps in suppressing delta ferrite, increases tensile strength, hardness, hardenability, promotes M23C6 and MX precipitation, coarsening of such precipitates


10

depends on its contents [20] and decreases Ms (martensite start temperature) [36]. With the increase in carbon contents, both toughness and weldability decrease [22]. It also delays Zphase precipitation [14]. 2.4.2: Boron In small contents, it enhances creep strength [6, 14, 15, 20], improves ductility [14, 15], increases quench hardenability [6, 20], retains fine subgrains, decreases self-diffusion coefficient, decreases impact toughness, stabilizes matrix, delays M23C6 coarsening, occupies vacancies, promotes M23(C,B)6 precipitation, retards ferrite-austenite transformation, stabilizes tempered martensite against re-austenisation, causes hot shortness [20], refines MX precipitates [13, 37], retards recovery [14, 15] and delays early start of tertiary creep [13]. The presence of enough boron can cause intergranular embrittlement [14, 15]. Generally, segregation to GBs of insoluble elements such as phosphorus, boron and sulphur from solid state, caused precipitation when exceeded the certain limit by interaction through rapidly diffusing species like nitrogen (e.g. BN formation) [15]. Boron is also believed to enhance GB cohesion i.e. providing GB, SGB strengthening [16]. As reported by Abe and by Kondo et al., [17], found high boron (90 130 ppm) containing steels does not suffer from type IV cracking with controlled nitrogen as normalised at 1150 oC to avoid boro-nitride formation, boron has the special feature of memory effect that retains microstructure similar before and after welding. 2.4.3: Chromium Chromium is a carbide former [20, 37], ferrite stabiliser. It decreases Ms, increases Ac1 temperature [37], provides oxidation resistance, creep strength [6, 9, 20, 22], promotes M23C6 precipitation [9] and improves hardenability [15]. The diffusion coefficient of Cr is higher than Mo [40] thus increased diffusion coefficient caused M23C6 coarsening [6]. 2.4.4: Molybdenum It is a carbide former [20, 37] and ferrite stabiliser. Its addition into 9Cr steels decreases Ms, increases Ac1 temperature, increases M23C6 coarsening [37], improves toughness, enhances creep rupture strength, promotes lave phase precipitation [20], provides resistance to pitting corrosion, improves hardenability and secondary hardening due to carbide precipitation during tempering [15]. It provides solid solution strengthening [15, 20, 22], and high-temperature strength by raising the binding force [9]. It prevents graphitization and suppresses temper embrittlement [20].


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2.4.5: Vanadium It is a ferrite stabiliser and carbide former [20, 37] and helps in retaining fine subgrains [6]. 2.4.6: Niobium It is a carbide former [20], z-phase precipitation Nb(C,N) promoter and subgrains refiner [6]. 2.4.7: Titanium It is carbide former [37] and improves creep strength of about 60 MPa when 0.06 % Ti added with 0.035 % Nb [3]. 2.4.8: Nitrogen It is an austenite former and which helps in suppressing delta ferrite [20], promotes VN precipitation [6], increases hardenability. High nitrogen contents may cause porosities in the mould so its fraction is limited to 0.2 % [14]. It also gives solid solution strengthening in ferrite with low solubility and migrates to dislocations at both high and ambient temperature. Nitrogen may cause strain ageing with deleterious consequences for the toughness and its diffusion can be prevented with small additions of aluminium or vanadium, which form AlN or VN respectively. It also lowers Ms temperature and gives resistance to pitting corrosion [15]. 2.4.9: Manganese It is cheapest carbide former, weak strengthening element and austenite stabilisers [15, 20]. Which delays the separation of carbides [15, 37], promotes M23C6 coarsening, M6C precipitation, lowers Ac1 temperature [6], promotes deoxidation and assures sound weld deposits [5]. It also prevents sulphur segregation by forming MnS which is known for hot shortness [15, 20]. With increasing, Mn contents, resulting in a loss in creep strength [20], high hardenability

but

may

suffer

to

temper

embrittlement

[20].

In small amounts,

Mn has also been known to partition to M23C6 [37]. 2.4.9: Silicon It is a ferrite stabiliser [37] and lave phase former [39] and added in small quantiles otherwise it can degrade mechanical properties [37] like creep strength [6, 20], weldability [15], toughness [14] but improves oxidation resistance [6]. It also graphitises carbides [37] and accelerates carbon diffusivity [40], increases diffusion coefficient which resulted in M23C6 coarsening [6].


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2.4.10: Aluminium It graphitises the carbide and added in small quantities otherwise mechanical properties will degrade [37] but it improves high-temperature oxidation, corrosion resistance (for contents >2%) [15] and enhanced toughness due to deoxidation caused clean steel [5]. 2.4.11: Nickel It

graphitises

the

carbide

and

known

as

austenite

stabilisers,

which

delays the separation of carbides and it is added in small quantities otherwise poor mechanical properties [36]. It increases diffusion coefficient i.e. M23C6 coarsening, loses in creep strength, decreases Ac1 [6], gives corrosion resistance in sulphuric acid and lowers the ductile/brittle transition temperature [15]. 2.4.12: Phosphorous It segregates in the interdendritic regions, leading to low incipient melting points and its addition increases the risk of hot-shortness during processing [15]. 2.4.13: Sulphur It segregates in the interdendritic regions, leading to low incipient melting points and its addition increases the risk of hot-shortness during processing. However, it improves machinability and impairs resistance to pitting corrosion [15]. 2.4.14: Cobalt Cobalt is expensive and known as austenite stabiliser [22]. Austenite stabilisers delay the separation of carbides and graphitise the carbides it should be added in small quantities otherwise mechanical properties may degrade [37]. It decreases diffusion coefficient i.e. M23C6 coarsening delayed, delays delta ferrite [6, 9], retards recovery [15]. Coarsening of lave precipitation may increase due to the addition of Co, which combine to form brittle intergranular type fracture. It also increases Curie point which then decreases the diffusivities of atoms [9]. 2.4.15: Copper It is austenite stabiliser which delays the separation of carbides [37], inhibits delta ferrite formation and accelerates lave precipitation [6]. However, it has limited solubility and gives age hardening [15].


13

2.4.16: Tungsten It provides solid solution strengthening [6, 20], dislocation hardening, retarding M23C6 coarsening, improving creep strength and produces fine sub-grain structure. High contents of tungsten cause high Ac1 and low Ms temperature [6]. Tungsten is the carbide former [37], delta ferrite former [28] and Fe2W particle formers [6]. 2.4.17: Rhenium It lowers Ac1 [20] and gives solid solution strengthening [6] including delay in recovery of the dislocation substructure in rhenium-containing steels [20].

2.5: HAZ

Figure 2.5: Schematic representations of (a) single pass and (b) multipass welded bead [41]. Fusion welding produces heterogeneous microstructure in weld HAZ [11] depending upon welding techniques employed. Figure 2.5 represents different zones produced during single and multipass welding. The four characteristic HAZ regions recognised in this steel are given briefly below: 2.5.1: CGHAZ Peak temperature is considerably higher than Ac3 temperature. At this point, all the precipitates completely dissolve and austenite grains grow during welding. Hardness and dislocation are quite higher than any other zones in HAZ [8]. Reprecipitation is noted in this region during PWHT [4]. 2.5.2: FGHAZ Peak temperature is just above Ac3 that is enough to produce fine austenite grains but the temperature is not enough to completely dissolve precipitate mainly M23C6 which limits the National Institute of Technology Warangal – 506004

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growth of austenite grains. This zone is weakest during creep operation at elevated temperature [10]. Hardness is lower than CGHAZ. Very fine laths of martensite with high dislocation density including high distribution of M23C6 type and MX-type precipitates are found in this zone [17, 34] and basically without lath distribution [17]. 2.5.3: ICHAZ Peak temperature is in between Ac1 and Ac3 that produces an incomplete transformation of ferrite into austenite with partially dissolved precipitates including their coarsening that gives the lowest hardness. Further coarsening and hardness caused decrement due to a fast reduction in dislocation density contributed by PWHT. Its microstructure made up of fresh martensite with a tempering of existing martensite also coarsened M23C6 [10, 34, 17]. Recovery of martensite lath structure with high dislocations into subgrains with low dislocations during tempering, softening reported at 900 oC due to unavailability of NbC& VN precipitates [8]. Fresh austenite formed at PAGB and lath boundaries which transform into martensite and existing martensite tempered during tempering [4]. 2.5.4: Over tempered region Peak temperature is lower than Ac1 and microstructure is like candidate material that undergoes tempering [4, 41].

2.6: Cracking

Figure 2.6: Several types of cracking in a weld joint [3, 42]. Four type of cracking are recommended in the weld joint during creep test and working conditions. Type I- propagation from WM itself propagating both transversely and parallelly to National Institute of Technology Warangal – 506004

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WM. Type II- propagation from WM that crosses HAZ transversely and parallelly. Type IIIoccurred in the CGHAZ which diffuses parallel and close to fusion line and eliminated by grain refinement. Type IIIa- cracking located near the fusion line which arose from carbon diffusion effects, changes in the concentration of chromium causes C diffusion which promote the narrow hardened band in WM and finally suffered a loss in creep strength generally found in lower Cr contents steels like 2.25Cr steels. Type IV- occurred and propagated from FGHAZ or ICHAZ, it can never be removed but minimised as shown in Figure 2.6 [3, 42].

2.6.1: Recent failures In UK power plants constructed during the 1960s and 1970s with 0.5%Cr–0.5%Mo–0.25% V steel (BS 3604 grade 660) for service temperature from 540 oC to 565 °C) they had a problem during welding cracking and failure i.e. type IV cracking occurred after 6 to 10 years. In late 1990s attention had paid on type IV cracking after that this was common in the UK, but early cracking observed for steel containing N/Al ratios 1.5. That time Europe was using 12%Cr steel X20CrMoV at moderate temperature with lower thickness sections Grade 91 steel is inexpensive and alternative to X20. Type IV mainly comes due to microstructural variations, heat treatment process given, service temperature and state of stress. But at a higher temperature and low-stress value mechanism of failures seen during transverse tests of WM & BM and vice versa [3]. In 2001, in Danish power plant T122(12Cr-2W-VNbN) steel is failed due to the formation of Z-phase [22]. In the UK, Early type IV cracking observed in high Cr header steels due to coarse aluminium nitride (AlN) precipitates, high aluminium contents, low N/Al ratio and MX free precipitates zone, while high N/Al ratio can postpone early type IV cracking. Brett observed that formation of AlN decrease the matrix strength by reducing the formation of VN precipitation, also they gave 3 possible mechanisms of type IV cracking (1) smaller grain size caused GB cavitation mechanism (2) small grains resists the lath martensite which results in loss in strengthening (3) welding cycle is not enough to reprecipitate different precipitates instead of this coarsening happening [43]. Type IV cracking in ICHAZ is due to deficiency of fine distribution of Nb and V precipitates within the matrix and PAGBs including the formation of fine grains due to polygonization of the tempered martensite [5].

2.6.2: Type IV cracking Over 40 years, this problem in HAZs produces detrimental effects on creep strength. ICHAZ is inserted between two highly creep resistant regions that introduce localised deformation which markedly creates creep cavitation known as type IV cracking [11] at lower strain compared to


16

BM voids nucleation start on precipitate and at PAGB they unite to form crack [3]. Cracking due to the microstructural gradient in the HAZ region is called type IV cracking [44]. ICHAZ, FGHAZ and over tampered HAZ regions introduce metallurgical notch results high triaxiality that finally causes cavitation and creep voids growth this failure is type IV. The stable microstructure of HAZ minimises triaxiality and increases creep life, this failure occurred at the region of minimum creep strength zone [8]. Boron and controlled amount of nitrogen with a lower amount of M23C6 and higher amount of fine MX precipitates mitigate type IV cracking by stabilising microstructure hence enhanced creep strength. But joint design, welding process joint restraint and PWHT also affect type IV cracking. It also forms same grain sizes & lath orientations between BM and HAZ this is called memory effect which means inherently retarded type IV cracking. It also reduces GB energy and delayed austenite formation during heating operations resulted in fine grain formation [8, 13]. Excess boron and nitrogen form intermetallic boro-nitrides (BN) phase during normalising [8, 13, 16, 45]. Enhanced creep strength of P91B steels due to increase in coincident lattice sites [8, 13]. Sakuraya et al., seen many coarse BN under dimples of ductile fracture usually 2 to 5 µm and did not dissolve during annealing at 1150 oC further it starts dissolving with time at 1200 oC whereas at 1250 oC it suddenly dissolves. BN nucleates easily as compared to Fe2B, also 80% of B was forming BN and rest going to the solution of the matrix [13]. Wang et al. [34], reported wedge type cavities at triple points that coalesce to form crack intergranularly with high local deformation of HAGBs these boundaries equipped with coarse M23C6, MoC/Mo2C and M7C3 carbides in FGHAZ leading to type IV cracking although these cavities mainly nucleate between fully recrystallized and un-recrystallized grains. Type IV cracking is a brittle fracture, occurred intergranularly at the transition zone between FGHAZ and ICHAZ due to nucleation of creep cavities on coarse precipitates like M23C6 found at PAGB that are multiplied by applied stress, coarsening and high rate of subgrains formation are accelerated by the recovery of dislocations. Yaguchi et al. [3 cf 19] observed non-uniformity in the number density of cavities throughout the HAZ region this value is maximum at the zone of higher triaxialities. 2.6.2.1: Mitigation of Type IV cracking There are five ways to avoid these failures i.e. normalising and tempering treatment, appropriate welding procedure, modification in the tempering process, alloy development [3], and enhanced design to minimise stress risers [5, 46].


17

2.6.3.2: Repair of Type IV cracking and effect of PWHT After nucleation longer heat treatment reduces interparticle spacing hence degradation in creep strength [3]. Swedish power plants repairs followed up by Storesund and Samuelson as in [47], they suggested following to be considered while repairing. 1.

Welding procedure that produces creep soft HAZ but it cannot be avoided so the region as much as narrow that requires low heat input with interpass temperature

2.

Design of weld groove that reduces state of stress

3.

Selection of filler matching with BM composition with a considerable amount of nickel and manganese contents.

2.7: Strengthening mechanism Creep resistant steels can be strengthened by solid solution, precipitation, dislocation and boundary strengthening [11, 37], which significantly improve strength at RT and elevated temperatures [4].

2.7.1: Solid solution strengthening Hume - Rothery law is suitable for high solid solution strengthening by bigger size of substantial solute atoms like Mo and W than Fe. A cloud of solute atoms accumulates in the solvent lattice mainly at dislocations, stacking faults, LAGBs and other GBs. Very little amount of Ti imparts higher tensile strength followed by Si, W, Mo, Mn and Ni etc. [48]. Mo and W are the prime contributors of solid solution strengthening in 9-12%Cr steels [3, 37]. Substitutional alloying elements whose atomic size is different to that of the solvent metal locally distort the crystal lattice. The resultant elastic stress fields interact with those around dislocations, requiring a higher applied stress for glide to continue. In the case of interstitial solutes, the local lattice distortion depends on the size and shape of the interstices and the type of atom concerned, in case for interstitial elements such as carbon and nitrogen, whose mobility allows them to diffuse to dislocations, where they form a so-called Cottrell atmosphere, which tends to pin the dislocation, impeding its movement, since if it breaks away. The overall energy of the system is increased. A higher stress is required to move the dislocations, but once they have torn free from their atmosphere they can glide under a lower stress, leading to a yield drop [15]. As VN and NbN provide precipitation strengthening but they also subject to get in solid solution [49].


18

2.7.2: Dislocation strengthening Dispersed fine precipitates not only stabilise the grains and subgrains but also contribute dislocation strengthening [28, 45]. Dislocations can only pass precipitates by climb mechanism and when the stress value reaches the value of threshold stress then dislocations bypass Orowan mechanism. This mechanism is faster than climb and the change in this mechanism accelerates Norton’s creep exponent hence creep deformation increases. Dislocation strengthening can be expressed in terms of dislocation density mathematically as following Equation 2.6 [45]. 𝜏𝑑𝑖𝑠𝑙𝑜𝑐𝑎𝑡𝑖𝑜𝑛𝑠 = 𝛼𝐺𝑏√𝜌 …Equation 2.6 Where G = shear modulus, b is Burgers vector, ρ is free dislocation density in the matrix and α is constant.

2.7.3: Grain/subgrain boundary strengthening Grain boundaries are the regions that contain an excess local concentration of defects like vacancies and dislocations, which caused atomic transport facilitates more easily so diffusion occurred in this zone, as they have high defects which imply nucleation sites of stable and metastable precipitates by decreasing interfacial energy and growth is depending on the rate of diffusion. GBs always inhibit dislocation motion due to the difference in orientation of the two crystals, higher the obstacles larger the degree of strengthening. Grain size refinement is achieved by recrystallization with a high nucleation density by controlled thermomechanical processing [15]. Smaller widths of laths and blocks promote SGB strengthening during creep conditions with fine distribution MX and M23C6 nucleated at/around GB [28]. Early cracking is one of the serious issues in the tempered martensite steel due to the high recovery of GBs and SGBs during high-temperature tempering treatment [6].

2.7.4: Precipitation strengthening In the precipitation hardening, strengthening is related to the distribution of particles in the matrix, volume fraction, mean particle size and average interparticle spacing whereas particle shape does not have much influence on strength but for notched impact, it strongly affects. As their solubility in the matrix is little which limit the grain growth at elevated temperature & during ageing. Carbon and carbides cause strengthening that enhance hardness and strength; in the case of lamellar carbides are much stronger than spheroidised carbides but in contrast for notched impact [48]. Precipitations pin the dislocation and SGBs during creep [31]. Group IV


19

to VII elements of periodic table are carbide formers like in group IV Ti, group V, V and Nb form MC type carbides with crystal structure similar to NaCl type, whereas transition metals of group VI (Cr, Mo and W) and VII (Mn) form M23C6 and M7C3 types (M = Cr or Mn). Although, Mo and W form hexagonal type MC and M2C carbides, group VIII (Fe, Co, Ni) all form M3C type carbides, but only Fe3C is stable under the nominal conditions of temperature and pressure. Nickel has little tendency to interact with carbon [15]. Precipitate like M23C6, MX, lave phase and Z-phase nucleate in martensite matrix below the Ac1 temperature till Ms temperature reached after that rate of nucleation is extremely slow they are semi-coherent and creep strength is mainly enhanced by fine MX and M23C6 [50]. Range and size of precipitates depend on working temperature and succession of their reactions during service that is affected coarsening and kinetics cause substantial change on rupture strength [3]. Figure 2.7 represents different precipitates in high Cr steels [22]. Huijun et al. [7], reported from TEM investigations of BM that mean diameter of lath approx. 0.4 µm with Cr-rich M23C6 carbide of 200-400 nm in length, stretched parallel to lath axes with MX type precipitates of approx. 50nm. During service, gigantically recovery of lath and dislocations occurs and the restoring of 9 years exposed to service steel by renormalizing. 2.7.4.1: Nucleation Nucleation is controlled driving force and interfacial energies of nuclei and matrix elements [22, 31]. If precipitation rate is high then nucleation of the small particle at defects like GBs, dislocations or other precipitates, where the distribution of precipitation is the function of nucleation sites and nucleation time [22]. 2.7.4.2: Growth After reaching a critical size (r*), growth begins, these precipitates are stable thermodynamically and grow till reaching to the equal potential of elements that precipitated. Growth results in an increase both volume fraction and precipitate size with respect to time and the process are governed by thermodynamic driving force, diffusion rate and concentration of elements in matrix [22]. 2.7.4.3: Coarsening (Ostwald ripening) Coarsening process consumes small particles and enlarges continuously precipitate that required lower thermodynamic force (interfacial energies) this is the last process where no nucleation occurs with no change in precipitates volume fraction, at this stage creep strength strongly affected. Ostwald ripening is controlled by a total decrease in interfacial energy and


20

its rate is given by following Equation 2.7 [6, 51]. The change in morphology is due solely to extra energy to the total surface area which is small coarsened the precipitates [51]. (𝑟)𝑛 − (𝑟𝑜 )𝑛 = 𝐾(𝜑)𝑡

…Equation 2.7 [6, 28, 51]

Where r and ro are the radii at time t and zero respectively, 𝜑 is precipitate volume fraction, K (𝜑) is coarsening constant including lattice diffusivities and is monotonically increase with 𝜑, t is time [28, 51] and n is exponent and its value depends on coarsening model e.g. for coarsening due to interface diffusion, volume diffusion, GB diffusion and pipe diffusion are 2, 3, 4 and 5 respectively, K must be minimum to delay coarsening, for M23C6 precipitates n=3, forMX rate is 10% of M23C6 due to n = 5 [6]. During creep exposure, Ostwald ripening of M23C6 is decreased due to boron [16]. Pulling of the solute atmosphere through a moving dislocation after it. This phenomenon is the Scavenging effect this is primary concerned in both creep testing and coarsening of precipitates because coarsening model of P91 steel precipitates itself contains 2 or 3 different models including solute transport by lattice diffusion. M23C6 precipitates and high dislocations are chiefly responsible for the boundary and core diffusion respectively causes Ostwald ripening [52]. Hald et al. [53], calculated coarsening rate constants forM23C6, MX and lave phases which evidently support that MX is highly stable than lave phase with a factor of 20 during service conditions and high coarsening of M23C6, at 600 oC, M23C6 coarsen from 50nm to 125 nm whereas when temperature increase to 650 oC size of particle reached to around 225 nm after 10000 h of exposure. 2.7.4.5: Metastable precipitate


21

Figure 2.7: Microstructure of different precipitates which chiefly observed in 9-12%Cr containing steels [22]. Many different types of precipitates nucleate in 9-12%Cr steels with shorter life which will be either dissolved in the matrix during heat treatment like tempering or grown by expensing other precipitates, are called metastable precipitates, one of such condition is shown below in the form of following Equation 2.8 and Figure 2.7 [22]. Primary M7C3is unstable and further precipitated and form M23C6, microhardness of Cr7C3 is 1400-2150 HV [15]. M7C3 has trigonal type structure nucleated at lath boundaries and dislocations [54]. M3 C

M7C3 + M2X

M23C6 + MX

M23C6 + Z-phase …Equation 2.8

2.7.4.6: M23C6 [(Cr, Fe, Mo, W)23(C,B)6] M23C6 carbides are inter and intragranular precipitates having morphologies like globular and cylindrical to lenticular at lath boundaries [39]. During tempering process, large amounts of chromium carbide of M23C6 precipitates nucleated around PAGB, SGB, packets, blocks and lath boundaries [8] with the expense of other carbides except for NbC [22]. High fractions of M23C6 are the main obstacles of climb motion, with tungsten, they also resist recovery of substructure [6]. M23C6 particles stabilise lath martensitic structure. Fine precipitates observed in CGHAZ but in the case of ICHAZ, these are coarse [8]. Coarse M23C6 particles are observed after and before welding at PAGB and packet/block boundaries [34]. M23C6 has face cubic crystal structure of chiefly Cr and C with minor amounts of Fe, Mo, W and B. B of around 100 ppm mitigates coarsening of M23C6 [37]. It precipitates in range 600-9000C and their nucleation decreases chromium in the matrix, microhardness is 1000-1650HV [15]. 2.7.4.7: M2X (M= Cr, Mo and X= C, N) They have hexagonal crystal structure [22, 54], mainly Cr2N type is observed of the faceted plate-like particles. M2X is thermodynamically more stable, even after several hours of exposure, these are found in both high and low nitrogen-containing steels and nucleate during low temperature tempering (1st stage tempering) on PAGB and dislocations (i.e. preferential sites of MX-type) they interfered with fine MX precipitates and are in big size than MX which can further effect strengthening, little amount of V is found in these precipitates as shown in Figure 2.7 [22]. Nucleated during tempering below 700 oC in subgrains and within laths and dissolved during high-temperature creep [37]. M2X also contains in turbine rotors at subgrains


22

[6]. Nitrogen has strong affinity to combine with chromium then after any other elements but manganese can increase the solubility of nitrogen in austenite and retards the nucleation of Cr2N for high chromium-containing steels and found within laths. they are not stable and transformed by MX. Microhardness of Mo2C is 1460-1800 HV [15]. 2.7.4.8: MX [M=Nb and V, X= C and N] NbC precipitation starts from normalising whereas VN during tempering [8, 22, 49], they nucleate very easily at PAGB, SGB, within martensite laths, their boundaries and in dislocations on cooling [8, 22] i.e. intergranular precipitates and stable at high temperature [2, 11]. At working conditions, they oppose movement of free dislocation and substructure growth, which pin the sliding of grains and cavity initiation [6, 8] by striking down driving force aroused from the recovery of densely populated dislocations which suggests MX are the strong obstructers of dislocation climb motions including grains and subgrains [6]. But in service conditions, they are very stable against coarsening and diffusing till Z-phase formation begins after that unstable and beginning to dissolve at long exposure time [2]. Also, the M(C,N) are grain refiners [41]. MX type nitrides coarsening increases with increase in nitrogen amount that creates high interparticle spacing of nitride MX type precipitates during creep test [8]. NbC prevents grain recovery and growth hence increase in creep rupture strength at high temperature [3]. Carbonitrides of aluminium and niobium are distributed at HAGBs [34]. Both Precipitation and dispersion strengthening particles are an obstacle for dislocations, when the precipitates are big enough and dislocation cannot penetrate precipitates then Orowan mechanism begins [31, 53], Orowan stress is given by Equations 2.9 and 2.10. 𝜎𝑂𝑟𝑜𝑤𝑎𝑛 = 3.32

𝐺𝑏 √𝑓𝑝 𝑑𝑝

… Equation2. 9 [31, 53].

Where, G = shear modulus (MPa), b = Burgers vector (nm), dp = precipitate diameter (nm), fp = precipitate volume fraction, it shows strengthening is inversely proportional to precipitate size and directly proportional to the volume fraction of the precipitate. 𝜎𝑂𝑟𝑜𝑤𝑎𝑛 =

𝐶𝐺𝑏 𝜆

…Equation 2.10 [45]

Where C is constant and λ is mean planar spacing of precipitates, Equation 2.10 determines that strengthening of such steel directly depends upon effective numbers of particle intersecting the dislocation glide plane and effective inter-particle spacing between precipitates whereas its absolute precipitate strengthening can be found by taking consideration nature of the dislocation-particle interaction [45]. The nature of the interaction depends on the mechanical


23

properties of the precipitate phase, together with the crystal structure and orientation. Matrix dislocations may shear precipitates that are coherent if their size and shear stress are sufficiently small. In the case of incoherent particles, since the slip planes are not continuous, dislocations must loop around the precipitates, by the classical Orowan mechanism which depends on temperature and high loads or climbs over them on strain rate at higher temperatures with lower stresses. The stress necessary for looping is inversely proportional to the particle spacing. For coherent precipitates that are stronger than the matrix, the stress necessary for shear increases with particle size, so that above a critical dimension looping becomes easier [15, 45]. Cumulative of the different strengthening mechanism is not simply accommodate for combined effects.The value of Orowan stress for M23C6, MX and lave phases are 150 MPa, 120 MPa and 95 MPa with percent fractions 2, 0.2 and 1.2 respectively but in case of tempered martensite 𝜎𝑂𝑟𝑜𝑤𝑎𝑛 , 400-500 MPa [6]. Effect of dislocation interaction can be expressed in terms of resisting force (F) shown pictorially and mathematically as follows [55]. F =2Τ sinθ …Equation 2.11 [55]. Where T = line tension of the dislocation segment as the equation determines, F increases when θ increases i.e. bowing of the dislocation and this suggests two different types of particles viz hard and weak particles.The particle is called hard particle when resisting force higher than 2T then bypassing of dislocation is through Orowan mechanism or cross slip such particles are a high obstacle of dislocation during service exposure and contributing highest strengthening. Similarly, the particle is soft If resisting forces attends the value greater than 2T before reaching θ equal to 90o caused shearing of the particle by the dislocation which may marginally increase work hardening due to slight improve in dislocation interaction [55]. If NbC is a spherical particle with relatively pure composition also it improves toughness by pinning grain growth although remain undissolved completely during normalising hence called primary MX (with minor contents of V, Cr and C) whereas VN (with minor contents of Nb, Cr and C). MX-type precipitates have a face centred cubic crystal structure similar to the crystal structure of NaCl type (but lattice parameter is different) [22]. Fraction and stability of V(C, N) are more than Nb(C,N) in the matrix during coarsening [53]. During normalising nitrogen remained in solid solution and start nucleating as carbonitride throughout tempering, MX precipitates also found on stacking fault, VC crystal structure similar to NaCl [37]. VN is fully miscible with the nitrides of the group IV elements, microhardness of NbC and VC are 2000-2400 HV and 22502950 HV [15].


24

2.7.4.9: M6X M6X nucleated during long-term creep exposure when Mo composition is higher than 1.6 wt.% with the dissolution of Nb(C, N) and M2X in 9-12%Cr steels at PAGBs and SGBs on M23C6 resulted in loss of solid solution due to depletion of Mo and W and precipitation strengthening, also they coarsen rapidly and deteriorates plasticity [15, 37]. 2.7.10: AlN Al is traces in 9Cr steels that forms AlN during tempering at PAGB, SGB [28, 37] and at dislocations [54, 56] their formation not only loses solid solution but also poor plasticity during service conditions which coarsened fast [28, 37]. It has HCP structure [55]. 2.7.11: Z-Phases [Cr (V, Nb, Ta) N] Chromium diffuses into MX phases, mainly VN (i.e. nucleation site) transformed to Z-phase that are thermodynamically more stable & ordered nature compared to any other precipitate with extremely slow nucleation rate markedly after many years [2, 3, 11], After dissolving MX precipitates markedly when VN dissolve, coarsening of Z-phase started and the precipitation rate is accelerated by high Cr contents, formation of this phase significantly reduces creep strength of the steel [3, 45] so it must be minimized [22]. To compensate reduced creep strength which occurred due to the formation of Z-phase, oxide nanoparticles like Y2O3 with unparalleled thermal stability are dispersed into the martensitic steels, thus simultaneously improvement in the upper temperature limit that evolves the development of oxide dispersion strengthened (ODS) martensitic/ferritic steel [19]. M23C6 and Nb(C, N) are preferential nucleation sites of Z-phase [37]. 2.7.11.1: 9–12%Cr steels strengthened by Z-phase In the austenitic steels, fine Z-phase (nitride type) precipitate and contribute to creep strength similarly precipitation of Z-phase that mainly come into picture after longer time and precipitate on fine MX type precipitates and decreases strength of creep resistant steels but if their precipitation occurred during tempering that creates smaller interparticle spacing, prolong the life of such steel with improved serviced temperature in power stations, as they are highly stable than MX and only the growth coarsening kinetics of Z-phase will retard its creep property after longer time period also if they precipitate and distributed finely during service exposure, instead of retarding strength it enhances creep strength [14].


25

Chapter 3: Experimental Set-up and Methodologies 3.1: As Received Material (AR) of P91B steel The material investigated in this study is boron containing modified P91 steel, used in boiler components such as boiler header and boiler tubes. Strengthening of these steels come from the solid solution of solutes like molybdenum and fine carbides and carbonitride of vanadium and columbium which significantly improve strength at RT and elevated temperature [4]. Higher contents of ferritic stabilisers may increase the risk of δ-ferrite formation [19]. Grains of tempered martensite contained carbides along PAGB, SGBs and dislocations. PAGB of martensite was made of many packets, in addition, packets made of several blocks [37]. A plate of this steel had been received in normalized and tempered conditions of 1333 K (1060 oC) for 30 minutes and 1033 K (760 oC) for 2 hours respectively, having dimensions 500 mm × 220 mm × 12 from Indira Gandhi Centre for Atomic Research, Kalpakkam, Tamil Nadu as shown in Figure 3.1. The chemical composition of this steel is given in Table 3.1. Table 3.1: Chemical composition of P91B steel in wt.%. Elements C

Cr

Mo

Mn

Si

V

Nb

Al

B

Wt. %

0.103

8.26

0.88

0.33

0.3

0.186

0.06

0.03

0.01

Ni

Cu

P

O

Ti

N

S

W

Co

Fe

0.01

0.008

0.006

0.006

0.0041 0.004

0.003

0.0019 0.0012 Balance

3.2: Gleeble Simulation and PWHT 3.2.1: HAZ Simulation by Gleeble 3800 Gleeble 3800 is such a device to simulate HAZ which is capable of both rapid heating and cooling as shown in Figure 3.2. A 10-mm square or 11-mm square bar is usually 60mm long (or a little longer). Depending on the cooling rate desired, the free span can be as short as 10mm or less. The shorter the free span, the higher the cooling rate. The principle of heating in Gleeble is direct resistance heating system which can heat specimens at a rate of up to 10,000 oC/s or can hold steady-state equilibrium temperatures. High thermal conductivity grips hold the specimen that makes capable of high cooling rates.


26

Figure 3.1: Plate of P91B steel received in normalised and tempered conditions.

Figure 3.2: Gleeble® 3800 simulator, left side is digital control console, center is Gleeble load unit and the right one mobile conversion unit.

Figure 3.3: Test specimen of length 78mm with 11X11 mm2 cross section after simulation and mid portion showing thermocouple weld.


27

In this Gleeble weld HAZ simulation of single pass and multi-pass welding, samples having dimensions 78 mm X 11 mm X11 mm was cut from Plate of P91B steel. The plate was first cut by wire cut electrical discharge machining (EDM) followed by precision surface grinding. This sample cut is used to replicate sub-zones of HAZ as shown in Figure 3.3. This sample was initially fine polished to ensure free from dirt and oil, so that bead free joint can be obtained by welding of a sample with the K-type thermocouple at 32 V as shown in Figure A.1 (see appendix I). As in this simulation, the peak temperature is 1240 oC so high care is observed during joining of the thermocouple. The heat input was 1.1 kJ/mm with 10 mm of the free span and the temperature profiles were recorded by using the K-type thermocouple. Cooling was accommodated by using full contact 304L austenitic stainless steel grips. Heating process was always linear profile while during the cooling process, nature of curves was linear and exponential as given in Equation 3.1. 𝑻 = 𝑻𝒎𝒂𝒙 ×𝒆(−𝟎.𝟒𝟕𝒕⁄∆𝒕) …Equation 3.1 Where, Tmax= peak temperature (oC). And, ∆t = cooling time from 800 to 500 oC, sec.

3A.2: Post Weld Heat Treatment (PWHT) PWHT has been carried out in the muffle furnace. Four samples i.e. CGHAZ D05, FGHAZ D02, ICHAZ D01 and CG + FGHAZ (Coarse grain and fine grain HAZ) D01 of simulated HAZ are introduced in the furnace when the temperature was reached to the temperature of 760 oC then held for 3 hours after that normally air cooled.

3.3: Hardness testing Hardness is a measure of resistance to localized plastic deformation. The testing machine consists of specimen stage that allows gradual contact with indenter with a preset load in kgf applied for a certain time called dwell time as shown in Figure 3.4. During operation, both application and removal of load care should be taken like no movement of specimen or indenter. A measuring microscope is attached to the testing machine generally objective of 10X and 20X. Due to the used diamond pyramid, Vickers hardness number is also known as diamond pyramid hardness number (DPH). Vickers hardness principle is similar to Brinell except for the indenter, is the square-based pyramidal diamond with face angles of 136° ± 30 min, polished and pointed nose. Figure 3.5 represents the principle of Vickers indentation and line diagram for the macroindentation test done on the surface of Gleeble weld HAZ simulated sample for P91B steel. National Institute of Technology Warangal – 506004

28

Figure 3.4: (a) Vickers macro-hardness testing machine, Economet VH-50MD and (b) Vickers micro-hardness testing machine, Leica VMHT Auto.

Figure 3.5: Schematic diagrams of (a) Principle diagram of Vickers hardness and (b) macroindentation tested surfaces. Formula to calculate is as follows 𝐿𝑜𝑎𝑑 𝑎𝑝𝑝𝑙𝑖𝑒𝑑

𝐻𝑉 = 𝑆𝑢𝑟𝑓𝑎𝑐𝑒 𝑎𝑟𝑒𝑎 𝑜𝑓 𝑖𝑛𝑑𝑒𝑛𝑡𝑎𝑡𝑖𝑜𝑛 =

2𝑃 𝑠𝑖𝑛 𝛼 𝑑2

… Equation 3. 2

Where: P = Force applied in kgf, d = mean diagonal of impression in mm, And α = face angle of diamond = 136°.


29

=

𝟏.𝟖𝟓𝟒𝟒𝑷 𝒅𝟐

… Equation 3. 3

The macro-hardness measurements are carried out using 30 kgf load with 15 seconds’ dwell time by following ASTM E92 procedure at 100X of magnification. These measurements are made on surfaces polished up to 600 grit of emery paper. Whereas, for micro-hardness testing, samples have been polished up to 1000 grit followed by disc polishing with 0.05-micron size of alumina powder and finally etched. Micro-indentation load is 50 g with 15 s’ dwell time by following ASTM E 384 procedure at 500X of magnification which again magnified to 1990X by the preinstalled software in Leica VMHT Auto hardness tester. Generally, the spacing of two indentations should be more than two and one-half times of diagonal of indentations. Testing variable CGHAZ is kept constant throughout the process for easy and quick comparison purpose. Both diagonals of indentations must be measured with appropriate magnification and mean is to be out in Equation 3.2. The accuracy of test results is a function of the force applied, indenter, testing machine and dwell time. The purpose of macro-hardness was only to demonstrate hardness variations on the surfaces of the simulated sample; whereas, microhardness tests was done to study hardness variations before and after PWHT on the cross section of the predefined simulated volume. For that purpose, contour mapping of hardness was first found to be appropriate before and after PWHT.

3.4: X-ray diffraction (XRD) XRD is like a fingerprint of materials, used for crystalline materials to determine crystal structure and lattice parameter etc. In this crystal structure, lattice parameter which is subsequently used to find crystallite size and microstrain with the help of Williamson Hall (WH) method was determined. Results from WH plot are used to find out dislocation density. Samples for XRD characterization were etched with 5 – 7% nital and Vilella reagent and the scanning was done on X’Pert Powder of PANalytical. Scan range is 30 to 120o with Cu Kα source with a step size of 0.0167113 and time/step was 81.915. Whereas, the analysis was done with the help of X’Pert HighScore Plus software. To find out RA, the samples were electropolished in the solution of 200 ml glycerol, 40 ml perchloric acid and 700 ml ethanol at a voltage of 19.2 V with current 1-3 mA under sub-zero temperature for 10 minutes. The Current selected for electro-polishing of the cross section for Gleeble simulated HAZ samples acts an oxidizer for electrolytic etching. Although, scanning range was 400 to 460 with time/step of 107 seconds. Equation 3.3 represents classic Koistenen and Marburger equation [36]. f = exp[β(MS – T)] …Equation 3. 4 [36] National Institute of Technology Warangal – 506004

30

Where f is the austenite fraction, Ms represents martensite start temperature, T represents the lowest temperature reached during quenching, β is a constant which is approximately equal to -0.011/K (oC), that does not give any description of the progress of transformation under different cooling rates [36].

3.5: Metallography 3.5.1: Optical Microscopy (OM)

Figure 3.6: Leica DM 2500M Optical Microscope. An optical microscope is widely used by metallurgists which is based on the principle of reflection of light with a fixed wavelength of light as shown in Figure 3.6 which limits the resolution of the microstructure. In this study, it was used to find grain size distribution of the AR material and Gleeble simulated weld HAZ samples. Samples were initially polished up to 1000 grits of emery papers followed by disc polishing with 0.05 µm of alumina powder and then cleaned with ethanol/methanol. All the samples were simulated from different peak temperature so extensively different types of etchants or similar etchants with optimized compositions and etching time were used to reveal PAGB to measure PAG.


31

3.6.2: Scanning electron microscopy and Energy Dispersive Analysis of Xray spectroscopy Electron microscopes are the powerful tools to improve resolution as the wavelength of an electron is inversely proportional to applied voltage. But in the case of OM, there is fixed wavelength of light that limits both resolution and magnification. In the 9Cr steels, precipitates were smaller typically in the order of hundreds of nanometers which makes such tools essential for characterization of precipitation. There are generally two modes in SEM first is secondary electron mode (due to inelastic scattering and coming from surface) which is suitable topographical imaging) and later is backscattered electrons (produced due to elastic scattering from the surface or near the surface) this is used for compositional contrast thus external morphology is revealed. Sample preparation was same as OM except deep etching is required for SEM due to high magnification. The characterization is done on both VEGA3 TESCAN and NOVA NANO SEM 430. NOVA NANO SEM 430 is equipped with SDD (silicon drift detector) EDAX detector.

3.6.3: Electron backscatter diffraction (EBSD) This technique tells about microstructural-crystallographic features of crystalline materials such that understanding of crystal orientation, structure, deformation, grain morphology and defects etc. will be easy. EBSD is generally directed by using SEM having EBSD detector which comprises at least a phosphor screen, low light CCD camera and compact lens. Sample preparation for EBSD characterization was similar to OM except at the last polishing with silica colloidal suspension in order to reveal microstructural features and to get high IQ value.

3.6: Thermodynamic and kinetic modelling Nowadays thermodynamic and kinetic modelling are used for alloy design, predicting precipitation behavior and phase stability which affect strengthening and stability of matrix.

3.6.1: JMatPro (Java-based Materials Property) JMatPro version 6.1 is Java-based simulation software to evaluate a variety of materials properties. Here, CCT curve is plotted by feeding input data as chemical composition and average grain size of the base metal (see PAGB of AR material in Figure 4.4 (a)). Other than phase transformations like TTT/CCT, JMatPro is also capable of computing many other properties such as temperature dependent thermo-physical properties of each phase formed,


32

including density, Young’s modulus, coefficient of thermal expansion (CTE), and thermal conductivity and temperature dependent mechanical properties of each phase formed, like tensile strength, yield strength and hardness. The Java-based simulation is adopted to validate phase transformations during Gleeble simulation of weld HAZ for P91B steel [57].

3.6.2: Thermocalc (Thermodynamic Calculations) Thermocalc and TC-Prisma are the thermokinetic software used for thermodynamic equilibria to design and process advanced and new materials [40]. Although their thermodynamic functions are termed in polynomials that itself generated from extrapolation of lower order systems. CALPHAD (calculation of phase diagrams) is based on minimizing Gibb’s free energy which contributes every phase in a multicomponent system to reach equilibrium. Thermocalc uses chemical composition and temperature to simulate precipitation behavior except for crystal structure, interfacial energies, mechanical stress and kinetics. 𝐺 = ∑𝑛𝑖=1 𝑛𝑖 𝐺𝑖𝑚 = 𝑚𝑖𝑛𝑖𝑚𝑢𝑚 … …Equation 3. 5 Where ni is the moles of phase i and Gim is the molar Gibbs energy of phase i. A thermodynamic model of P91B steel is calculated from Thermodynamic database, TCFE7 and mobility database MobFe2. Bulk nucleation sites were assumed to calculate the mean radius, critical size, driving force, number density and phase fraction of various precipitates at different PWHT conditions for AR material, CGHAZ D05, FGHAZ D02, ICHAZ D01 and CG + FGHAZ D01.


33

Chapter 4. Experimental Results and Discussions From microstructures of weld HAZ simulation, we observed large PAGBs for P91B steel samples. But from the SEM analysis, we have confirmed that substructures found in the weld HAZ samples were close to microstructures of actual/simulated weld HAZ subzones as reported by [10, 12, 17, 18, 26, 34, 58, 59, 60, 61]. So, further studies have been conducted on these samples to achieve objectives of this research work.

4.1. Thermal profiles of Gleeble simulated HAZ 8 samples of P91B steel were simulated by Gleeble 3800, which are shown in Figures 4.1 (a) to (h). Whereas, their heating rates, cooling rates, exact holding time etc. are given in the following Table 4.1. Whereas, D is the Gleeble profile name followed by a numeric number to distinguish one from another profile. Samples were simulated in Gleeble 3800 by providing inputs like different peak temperatures (Tmax) with time to reach Tmax, holding time, then cooling temperature with time. During cooling from 800oC to 500oC, more points with corresponding time were given. Nature of cooling curve was concerned with the particular simulated sample, generally linear or exponential or in a combination of both linear and exponential were adopted.


34

Figure 4.1: Gleeble thermal profiles of single and multipass welding HAZ simulation of P91B steel.


35

Table 4.1: Thermal histories of Gleeble Simulated HAZs. Where, EHT= Exact holding time. Sample ID

Temperature

Actual

Heating/cooling

range (oC)

Heating/cooling

rate (oC/s)

Nature of curve

rate(oC/s) (a) CGHAZ D01

0 – 1240

100

101.6318

Linear heating

(EHT = 3.19 s)

800 - 500

20

19.86312

Linear cooling

(b) CGHAZ D04

0 – 1240

100

100.7132

Linear heating

(EHT = 3.18 s)

800 - 500

20

19.81067

Linear cooling

(c) CGHAZ D05

0 – 1240

100

101.1918

Linear heating

(EHT = 3.17 s)

800 - 500

200

19.74017

Linear cooling

(d) FGHAZ D02

0 - 1040

100

103.8

Linear heating

(EHT = 4.16 s)

800 - 500

20

19.96

Linear cooling

(e) FGHAZ D03

0 - 1040

100

101.5

Linear heating

(EHT = 4.2 s)

800 - 500

20

19.97

Linear cooling

(f) ICHAZ D01

0 - 865

40

39.73

Linear heating

(EHT = 8.07 s)

800 - 500

20

20.44

Exponential cooling

(g) ICHAZ D03

0 - 865

40

40.2757

Linear heating

(EHT = 8.06 s)

800 - 500

20

20.644

Exponential cooling

(h) CGFGHAZ

0 - 1240

100

101.33

Linear heating

D01 (EHT = 3.1 s)

800 - 500

20

20.79

Exponential cooling

(EHT = 4.13 s)

186 - 1040

100

97.82

Linear heating

800 - 500

20

20.04

Exponential cooling

4.2: Hardness evaluation 4.2.1: Bulk hardness on surface of the sample Different weld HAZ simulated samples were tested before PWHT along the length of the simulated surface of the samples except for AR material. For AR material, arbitrary points were chosen to measure hardness at that point and it has an average hardness of 225 HV/30 with the higher hardness at cross section is due to hot rolling of the P91B steel plate as shown in Figure 4.2 (a). Indentations were longitudinally done leaving 3 mm from both edges of the sample surface by considering about 12 mm of entire length to observe bulk hardness variation as shown in Figure 4.2 (b). Mean, standard deviation, maximum and minimum hardness have been


36

shown in The Figures 4.2 (a) to (j). In these figures, the x-axis represents distance taken from the both sides of thermocouple weld spot in mm. It has clearly shown that CGHAZ samples have higher hardness followed FGHAZ and ICHAZ. Standard deviations are added to understand variations of hardness on the surfaces of the sample. CG + FGHAZ D01 and CGHAZ D04 are the two samples for which significant changes in hardness are estimated. For coarse-grained simulated samples, tooth shaped profiles were observed in Figures 4.2 (b) to (d). The Higher hardness of 405 HV/30 was observed for CGHAZ D05 profile followed by CGHAZ D01. But, the higher standard deviation was measured for CGHAZ D01 & CGHAZ D04, because next surface was suffered large indentation by size i.e. much lower hardness related to other surface areas as shown Figures 4.2 (b) and (c). CGHAZ D05 is showing excellent hardness variation comparing one surface to another surface. The simulation volume is similar in all CG’s samples. For fine-grained simulated samples, dome shape plotted are plotted as shown in Figures 4.2 (e) and (f), 412.5 HV/30 hardness had been measured for FGHAZ D02 sample which was second higher hardness among all the tested samples. The hardness of both the surfaces of FGHAZ D03 sample showing almost no change in high hardness in comparison to FGHAZ D02. In the inter-critical samples, zig-zag hardness variations are observed as shown in Figures 4.2 (g) and (h). Lower hardness than CGHAZ, FGHAZ and AR material are recorded. These regions are generally prime focused zones due to low hardness among all the sub-zones of HAZ including AR Material. ICHAZ D03 profile has a substantial change in hardness between thermocouple welded surface and its adjacent surface. Also, it was difficult to distinguish simulated zone of IC’s samples by observing hardness profile. So, each 6 mm from point of thermocouple welded is assumed for purpose of hardness studies of simulated samples and the centre region was chosen for other different characterization and testing. But by the microstructural characterization, it was concluded that ICHAZ samples have a comparatively low simulated volume of simulated HAZ. CG + FGHAZ D01 simulated sample was an amalgamation of tooth like shape and domed curvature. 93.2 HV/30 was largest standard deviation noted for this surface of sample among all the tested samples. It may be due to two different cycles in which cooling from 1240 oC to 160 oC and in other cycle cooling from 1040 oC to RT. Which was one thing. Another view is that variation in cooling rates between different regions of the sample in both the cycles. Mean hardness was 322.34 ± 111 HV/30. Figure 4.2 (j) shows a summary of all samples with mean and standard deviation as an error bar. It is noticeable that ICHAZs are primarily susceptible to type IV cracking due to least hardness during creep exposure. However, this thesis work is not


37

concerned to creep testing. That means post weld heat treatment is essential to minimise both microstructural and hardness variations before intended to service exposure. Also, maximum macro-hardness values of FGHAZ like D02 and D03 were larger than maximum hardness values of CGHAZs. So, micro-indentation was employed. In the macro-hardness, it has been observed that there was no transient zone between simulated and non-simulated regions. Instead of this low hardness has been remarked for BM due to unavoidable high/low tempering promoted by 304L austenitic stainless steel grips which cause coarse precipitates. Though it is not part of this study but its effects have further observed in the selected simulated sample during further characterization.

4.2.2: Micro-hardness at cross section of the sample Figure 4.3 (a) is showing a summary of micro-hardness testing for both before and after PWHT. The average hardness of the samples is shown by cap and their standard deviation as an error bar. All samples history which is already discussed earlier in their corresponding outline headings. Figures 4.3 (b) to (n) are representing contour hardness mapping where maximum and minimum are divided into ten equal parts as a colour bar indicator of hardness. AR material hardness mapping is shown in Figure 4.3 (b). Before PWHT, the micro-hardness at the cross section of Gleeble simulated weld HAZ samples was successfully done and consistent values are noted. These values are taken to map its hardness. CGHAZs samples had average hardness 530, 444 and 475 HV/0.05 for D01, D04 and D05 profiles respectively. Out of which, the standard deviation was maximum for CGHAZ D04 profiles but CGHAZ D01 and CGHAZ D05 profiles had a similar standard deviation of 24 HV/0.05. Whereas, maximum hardness is same for both CGHAZ D04 and CGHAZ D05. FGHAZ samples have a lower hardness than CGHAZ. However, variation in standard deviations was similar to CGHAZs. It is clearly noted from figures that every hardness reading for a sample differs. Also, maximum and minimum hardness values are not overlapping with another as happened in the case of macro-hardness. FGHAZ D02 has high standard deviation compared to FGHAZ D03 sample.


38


39

Figure 4.2: Hardness plots of various simulated samples including AR material before PWHT. Whereas SL11= hardness profile of the first surface with the first line, SL12= hardness profile of the first surface with line 2nd, SL21= hardness profile of the second surface with the first line and SL22=hardness profile of the second surface with the second line.


40


41

Figure 4.3: (a) Summary of micro-hardness by taking standard deviation as an error bar and (b) to (n) contour mapping of micro-indentation testing with average, standard deviation and hardness colour bar indicator. ICHAZ D01 sample showing lower hardness than ICHAZ D03 which contrasts with macrohardness. It may be due to non-homogeneously and non-uniformly cooling rather than heating National Institute of Technology Warangal – 506004

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during Gleeble simulation. Which causes its effect is opposite at cross-section in comparison to the surface. CG+FGHAZ D01 sample having to mean hardness 544.33 HV. However, the gripping material is not promoting such troubles. Low hardness in ICHAZ may susceptible to type IV cracking. After PWHT of CGHAZ D05, FGHAZ D02, ICHAZ D01 and CG + FGHAZ D01, the percent change in hardness variations from standard deviation are 34.95, 51.37, 54.55 and 2.74 respectively. Which indicates a substantial reduction in hardness after employing PWHT of 760 o

C/3h. But for CG + FGHAZ D01, this PWHT condition is insufficient to minimise hardness

variation. The increase in hardness after PWHT is due to precipitation of large numbers of MX type carbonitrides. But average hardness value is under permissible range so further tempering may cause over tempering of CG + FGHAZ D01.

4.3: Calculation of dislocation density by XRD The XRD pattern of simulated samples before and after PWHT including AR material have been shown in Figures 4.4 (a) to (c). From Equation 2.1 it has been calculated that c/a ratio for this 9Cr steel is 1.004635. Also, XRD pattern in Figures 4.4 (a) to (c) clearly represents this steel has BCC type matrix. Though, in Figure 4.4 (b) it has been shown that AR material has both tempered martensite and a small amount of brittle martensite. But for ICHAZ, XRD peaks of both overtempering martensite tempered martensite and fresh martensite are dominant as shown in Figure 4.4 (b). It has been observed that dislocation densities are decreased after PWHT treatment as shown in Table 4.2. Although, the effort has done to find out RA for HAZ simulated samples but XRD spectra do not show any RA because of high ferritic stabilisers, Ms and Mf (martensite fine temperature). It does not mean that RA is not present in HAZ simulated sample because 100 % phase transformation cannot possible so in this case numerical modelling has been adopted. First Ms temperature is calculated from Equation 2.4 which is 411 oC. Then Koistenen Marburger equation is used to find out RA as shown in Table 4.3. Table 4.2: Dislocation densities of both AR material and simulated HAZ samples before and after PWHT. S.no

1

Sample

AR material

Dislocation density before

Dislocation density after

PWHT (m-2)

PWHT (m-2)

5.80772E+13

--------


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2

CGHAZ D01

1.22021E+15

--------

3

CGHAZ D05

1.02701E+15

3.64178E+13

4

FGHAZ D02

3.0462E+14

1.55066E+14

5

FGHAZ D03

7.90361E+14

--------

6

ICHAZ D01

8.67E+13

6.48E+13

7

ICHAZ D03

2.64192E+13

--------

8

CG + FGHAZ D01

9.44414E+14

8.35826E+13

Figure 4.4: XRD pattern of (a) AR material and HAZ simulated samples before PWHT, (b) high-resolution scanning of AR material, CGHAZ D05, FGHAZ D02, ICHAZ D01 and CG + FGHAZ D01 and (c) HAZ simulated samples after PWHT.


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XRD patterns of both before and after PWHT treated are used to by estimating crystallite size and microstrain through Williamson Hall (WH) as shown in Equation 3.4. The values of crystallite size and microstrain are feed in Equation 2.3 to calculate dislocation densities which are given in Table 4.2. Table 4.3: Amount of RA calculated from classic Koistenen and Marburger equation. Sample

Quenching temperature [QT] (oC)

RA = exp[-0.011(411– QT)]

CGHAZ D05

84

2.74

FGHAZ D02

104

3.41

ICHAZ D01

75

2.48

CGFGHAZ D01

118

3.98

4.4: Microstructural assessment 4.4.1: By OM before and after PWHT treatment Before PWHT process, optical microscopy had been done to reveal PAGBs in the simulated HAZ samples by using different etchants or similar etchants of balanced compositions. Time of etching was different because of different peak temperatures. Only by viewing the surface, it was controlled. AR material was etched with 28% nital to reveal PAGB and lath morphology as shown in Figure 4.5. To reveal PAGBs inCGHAZ D01, CGHAZ D04 and CGHAZ D05, 57% nital + Vilella, 28% nital and 6% picral (acid) respectively are used as shown in Figures 4.6 (a), (c) and (e). FGHAZ D02 and FGHAZ D03 are etched with 6% picral and 6% picral (acid) to reveal PAGBs as shown in Figures 4.6 (g) and (i). FGHAZ D03 are also etched with Curran’s reagent and found revealed and unrevealed grains these grains are packets, not PAGB as shown in Figure 4.6 (j). ICHAZ samples are etched with 5-7% nital + Vilella’s reagent to reveal PAGB as shown in Figures 4.6 (k) and (m). At last CG + FGHAZ D01 sample is etched with 6 % picral to reveal PAGB as shown in Figure 4.6 (o). This method stated that PAGBs are larger for CGHAZs samples. By OM, it was confirmed that BM and ICHAZ samples contain tempered martensite with large numbers of M23C6 at PAGBs, SGBs and within the matrix, as shown in Figure 4.5 and Figures 4.6 (l) and (n). In the CGHAZ samples predominantly lath type martensite was present as shown in Figures 4.6 (b), (d) and (f) whereas the fine and uniform distribution of laths were observed for FGHAZ samples as shown in Figures 4.6 (h) and (j). In the CG + FGHAZ D01, three distinguished features like thin, wide and split laths were observed as shown in Figure 4.6 (p). National Institute of Technology Warangal – 506004

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Figure 4.5: AR material (BM) etched with 10% aqua regia in distilled water. After PWHT, original PAG of the weld HAZ simulated samples are revealed by etching with 5 – 7% nital followed by Vilella etchant as shown in Figures 4.7 (a), (c), (e) and (g). In the CGHAZ D05, FGHAZ D02 and CG + FGHAZ D01 samples, martensite was transformed to tempered martensite with a large amount of primary precipitates which is agreed to Thermocalc property diagram as shown in Figures 4.7 (b), (d) and (h). However, large recovery and recrystallization were observed for CGHAZ D05 and ICHAZ D01 but coarse precipitates were observed for ICHAZ D01 as shown in Figures 4.7 (b) and (f).

CGHAZ D01 etched with 5-7% nital + modified picric acid (a) at 50X and (b) at 1000X.


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CGHAZ D04 etched with (c) 28% nital at 50X and (d) Vilella’s etchant at 1000X.

CGHAZ D05 etched with picral (acid) (e) at 50X and (f) at 1000X.

FGHAZ D02 etched with 6% picral (g) at 50X and (h) at 1000X.


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FGHAZ D03 etched with (i) 6% picral (acid) at 50X and (j) Curran’s reagent at 1000X.

ICHAZ D01 etched with mixture of 5-7% nital + Vilella (k) at 50X and (l) at 1000X.

ICHAZ D03 etched with 5% nital + Vilella (m) at 50X and (n) at 1000X.


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CG + FGHAZ D01 etched with 6% picral (o) at 50X and (p) at 1000X. Figure 4.6: Microstructures from the optical microscope of various samples prior to PWHT with their corresponding etchants.


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Figure 4.7: Light microscope micrographs after PWHT, etched with 5–7 % nital + Vilella’s reagent of 4 g picric acid to CGHAZ D05 at (a) 50X and (b) 1000X, FGHAZ D02 at (c) 50X and (d) 1000X, ICHAZ D01 at (e) 50X and (f) 1000X and CG + FGHAZ D01 at (g) 50X and (h) 1000X.

4.4.2: PAG size estimation through OM microphotographs Grain distributions of various simulated samples and AR material have been shown in Figure 4.8. It was evaluated that BM had a very small grain size of 17.65 µm with minimum standard deviation as shown in Figure 4.8 (a). Whereas, approximately 400 to 550 µm size of grains are evaluated for CGHAZ simulated samples as shown in Figures 4.8 (b) to (d). Summary of PAGB sizes for different HAZ simulated samples is shown in Figure 4.8 (j). There is no doubt that microstructures of FGHAZ and ICHAZ samples contained coarse grains. Excess normal grain growth is due to the additional driving force resulting via, reduction in GB area by 304L austenitic stainless steel grips, which are not only responsible for the growth of grains but also


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promoted tempering when the temperature was less than the Ac1 temperature in both simulated and non-simulated zone. This was also confirmed by viewing microstructures of non-simulated regions of the simulated samples which contained large grains compared to the as-received material that means weld HAZ simulated regions were also suffered growth as shown in Figures A2.1 (a) to (j) (see appendix II).


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Figure 4.8: Grain size distributions in the form of the histogram with normal distribution for different samples. 304L austenitic stainless steel grips have low thermal conductivity (14 – 16.3 W/m-K as compared to the thermal conductivity of copper grips (385 W/m-K, 357 W/m-K @ 727 oC, 398 W/m-K @ 27 oC) and low CTE which causes poor thermal gradient such gradients caused extensive normal grain growth. Also, microstructurally growth took place by either dissolution or annihilation of pinning particles, leading to a decrease in pinning effect which promotes grain growth particularly for CGHAZ and ICHAZ. But the substructures were nearly similar to HAZ subzones that is why the study is continued.

4.4.3: SEM microstructures analysis before and after PWHT treatment Before PWHT, SEM analysis was done by etching with 28% nital to AR material and 5 -7 % nital + Vilella reagent to HAZ simulated samples. In the AR material, mostly laths were still present in the matrix with a cluster of small precipitates.


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53


54

Figure 4.9: SEM microstructures of various samples prior to PWHT, AR material (a) and (b), simulated samples like CGHAZ D01 (c)and (d), CGHAZ D05 (e) and (f), FGHAZ D02 (g) and (h), ICHAZ D01 (i) and (j), ICHAZ D03 (k) and (l) and CG + FGHAZ D01 (m) and (n).


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Figure 4.10: SEM microstructures of various samples after post weld heat treatment at 760 o

C, 3 h, CGHAZ D05 (a) and (b), FGHAZ D02 (c) and (d), ICHAZ D01 (e) and (f) and CG + FGHAZ D01 (g) and (h).

By seeing microstructures of AR material, it has been concluded that in this material recovery and recrystallization of martensite is sluggish as shown in Figures 4.9 (a) and (b). CGHAZ D01 and CGHAZ D05 were not fully etched by this reagent but lath martensite is primarily present in the form of packets in the PAG as shown in Figures 4.9 (c) to (f). In the FGHAZ D02, fine and uniform distribution of lath martensite observed in the packets of PAG with a unrevealed portion of lath martensite. This unrevealed portion has been confirmed by EDAX analysis that these are martensite packets as shown in Figures 4.9 (g) and (h). However, 5–7 % nital and Vilella reagent have not revealed PAGBs in this material. ICHAZ samples contain tempered martensite and coarse M23C6 type of precipitates. It is also observed that some regions of ICHAZ samples fully recovered and transformed to ferrite as shown in Figures 4.9 (i) to (l).


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Whereas in CG + FGHAZ D01 sample wide laths, long and thin laths and thin laths were present as shown in Figures 4.9 (m) and (n). After PWHT, martensite of CGHAZ D05 and FGHAZ D02 is transformed to tempered martensite as shown in Figures 4.10 (a) and (c) respectively. However, large recovery and coarsening of precipitates were found in CGHAZ D05 sample compared to FGHAZ D02 where still fine laths are present as shown in Figures 4.10 (b) and (d). whereas, in the ICHAZ D01 sample large area recovered with extensive coarsening of existing carbides from the simulated sample and new carbides are also nucleated during PWHT as shown in Figures 4.10 (e) and (f). uniform distribution of precipitates was observed for CG + FGHAZ D01 with lath martensite as shown in Figures 4.10 (g) and (h).

4.4.4: Precipitates finding and its analysis before and after PWHT

Figure 4.11: EDAX plots of (a) Fe-rich M23C6 carbide (b) Cr-rich M23C6 carbide and (c) Mo6C carbide.


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Figure 4.12: SEM microphotographs showing several types of precipitates before PWHT of samples (a) AR material, (b) CGHAZ D05 (c) FGHAZ D02 (d) ICHAZ D1 (e) ICHAZ D03 and (f) CG + FGHAZ D01.


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Figure 4.13: SEM microphotographs displaying several types of precipitates after PWHT of at 760 0C/3h (a) CGHAZ D05, (b) FGHAZ D02 (c) ICHAZ D01 and (d) CG + FGHAZ D01. By EDAX analysis, normally two type of precipitates were found and its plots are shown in Figures 4.11 (a) to (c). These precipitates are nothing but Fe and Cr-rich M23C6 carbides and Mo6C carbide. In the AR material, agglomerated Fe/Cr/Mo rich complex carbides are found as shown in Figure 4.12 (a). Before PWHT treatment, 0.2 µm to 1 µm size of precipitates are observed in CGHAZ D05 sample as shown in Figure 4.12 (b). Around on average 0.25 µm size of precipitates are observed in FGHAZ D02 sample as shown in Figure 4.12 (c). In the ICHAZ samples, coarse precipitates of Fe/Cr/Mo rich carbides are generally found as shown in Figures 4.12 (d) and (e). However, in the CG + FGHAZ D01 sample contains Mo rich complex carbides as shown in Figure 4.12 (e). CGHAZ D05 may be undergone tempering for a short period of time during Gleeble simulation because steel grips below lower critical temperature will provide tempering instead of promoting growth that is why these samples contain coarse precipitates. After PWHT, there is no new type of precipitates were found except that nucleation of fresh precipitates of an existing type of carbides with growth and coarsening of existing precipitates as shown in Figures 4.13 (a) to (d).

4.4.5: EBSD of As received material 4.4.5.1: Grain Boundaries The inverse pole figure (IPF) in Z [001] direction (perpendicular to the screen) shows grain orientation as shown in Figure 4.14. Grain size calculation is done by EBSD and shown in Table 4.4 and Figure 4.15 having an average grain of 21.31 µm with a standard deviation of 3.51 µm.


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Figure 4.14: Inverse pole figure (IPF) in Z [001] direction (perpendicular to the screen) shows the grain orientation distribution for AR material.

Figure 4.15: Grain size distribution of AR material in terms of area fraction. Table 4.4: (a) PAG distribution in terms of area fraction and (b) misorientation angle in terms of number fraction. (a) Grain size estimation

(b) Misorientation angle estimation

Grain diameter (µm)

Area fraction [%]

Angle [degrees]

Number fraction

0.830348

1.41734

3.575

30.7466

1.03841

0.437207

6.725

8.66329

1.29861

1.03613

9.875

1.99006

1.62401

1.218

13.025

1.41784

2.03094

1.6897

16.175

1.36493


60

2.53984

1.98356

19.325

0.964338

3.17626

2.64788

22.475

0.594433

3.97214

3.34535

25.625

0.658641

4.96746

4.42045

28.775

3.48704

6.21217

5.34101

31.925

2.84375

7.76878

6.8448

35.075

1.99046

9.71543

5.85571

38.225

2.44598

12.1499

8.60752

41.375

3.04728

15.1943

9.71084

44.525

3.04122

19.0016

8.31142

47.675

2.71412

23.7629

8.22004

50.825

5.63056

29.7172

8.34681

53.975

9.04289

37.1636

2.62817

57.125

8.59828

46.4758

11.0731

60.275

10.6832

58.1214

6.86496

63.425

0.0751118

21.3102

Average

Average number

3.50994

Standard Deviation

29.1203

4.4.5.2: Misorientation Angle Low-angle grain boundaries (LAGB) having a misorientation of 2 to 15 degrees and high-angle grain boundaries (HAGB) having a misorientation greater than 15 degrees of the matrix grains in the AR material. These are divided by determining the misorientation angle between adjacent grains of AR material. Martensite lath and block/packet boundaries are SGBs within PAGs of AR material have a low angle of misorientation than that of PAGBs. PAGBs are always known as HAGBs due to their higher misorientation angle and the distribution of misorientation angle is shown Figure 4.16. The fraction of HAGB and LAGB are tabulated in Table 4.4 where an average of 29.12 comes out.


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Figure 4.16: Distribution of high angle and low angle grain misorientation for AR material. 4.4.5.3: Kernel Average Misorientation (KAM) Kernel average misorientation (KAM) is the arithmetic mean of the scalar misorientation angle between groups of pixels or kernels, that deals with local strain levels of the grains [34]. Local misorientation distribution is shown in Figure 4.17. The blue colour shows zero degrees misorientation while red colour show five degrees of misorientation as in KAM map and local strain level was high in the particle containing grains.

Figure 4.17: Distribution of KAM for AR material.


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4.5: Computational diagrams by thermodynamic and kinetic calculations 4.5.1: CCT and other diagrams by JMatPro It is always kept in mind while developing martensitic steels that their Mf temperature must be above ambient temperature to ensure minimum/no RA. This RA is known to lose its strength before tempering while after tempering production of untempered brittle martensite is generally observed. In the low carbon martensitic steels, carbide precipitation is recorded during tempering. Elements like C, Mn, Cr and V etc. decrease Ms and Mf temperature except for Co, small grain size can also lower Mf [22]. Critical cooling rate to form martensite for P91 steel is about 0.2 0C/s [5]. JMatPro developed continuous cooling plots and suggested there will be martensite formation for all the simulated samples, even in the dual simulation of CGHAZ and FGHAZ has undergone similar profiles. Cooling cycle of Gleeble thermal profiles are reproduced on CCT curves which clearly suggests in this simulation there is only martensite formation no matter if kinetics slow after 800 oC to 500 oC as shown in Figures 4.18 (a) and (b). Whereas, in Figure 4.18 (c) advanced CCT diagram is shown. By CCT diagram it is also confirmed that cooling rate of 20 oC/s was in the zone of fully martensite formation. Figures 4.18 (d), (e) and (f) represent phase fraction when material was heated and hold at 1240 oC, 1040 oC and 865 oC respectively. One of key point is to note that in this case amount of carbon is low which forms BCC martensite structure rather than BCT which is further proof by XRD (see chapter XRD).

4.5.2: Calculation of thermodynamic diagrams by Thermocalc Figures 4.19 (a) to (f) depict phase diagrams of P91B steel where Figure 4.19 (a) is iron carbon diagram where one key point is to follow that delta ferrite formation started nearly at 1225 oC so it can be said that CGHAZ sample was undergone delta ferrite region but by the time of cooling all delta ferrite transformed into austenite which further transformed to martensite. As also from hardness testing there was no indication for the formation of any such phase. It also shows MX and little amount of Z-phase may form in this steel but as this steel was received in normalised and tempered conditions which are only favourable for mostly secondary MX precipitates and Z-phase does not have a driving force to precipitates as it dissolved at given tempering temperature. However, we do not want Z-phase during creep exposure but thermodynamically this phase is highly stable even comparing to MX so fine and uniform National Institute of Technology Warangal – 506004

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distribution of such phase may contribute sufficient strengthening to P91B steel but till now nobody has reported for P91B steel and such study is out of scope for this dissertation. Figure 4.19 (b) is the Fe–Cr diagram for P91B steel which illustrates M7C3 and other chromium carbides should not form as they are not thermodynamically stable and by the time they must be engulfed by M23C6 which is agreed with equation 2.5 [30]. In the Figure 4.19 (c), it is understood that molybdenum contents are not enough to precipitate lave phase and it starts forming when molybdenum contents are raised to 1.1 wt. % Mo. Figures 4.19 (d) to (f) show primary and secondary precipitates in this steel. Although the high amount of large Cr carbides has been detected and which agreed with Thermocalc phase diagrams. Also, Zhang, WenFeng, et al. [33], has reported 760 oC temperature produces a large amount of MX precipitates and its effect in the form of increased hardness after PWHT for ICHAZ sample was observed. Property diagrams from Thermocalc for P91B steel is shown in Figure 4.20, from which it is now known that boron rich M23C6 and MX precipitates are present in this steel especially after PWHT treatment where a large amount of primary and secondary are nucleated. It is also observed from property diagram that dissolution temperatures of M23C6, Ti(N,C,B), Nb(N,C,B), V(N,C,B) and Z-phase are 875 oC, 1230oC, 1150 oC, 860 oC and around 730 oC respectively. Also, delta ferrite formation starts around 1250 oC. Figures 4.21 (b) to (g) are representing the component in various phases shown in property diagram with respect to temperature and the green line is marked for PWHT temperature. Figure 4.20 is self-sufficient to illustrate that there will be no precipitation of boron nitride and other boron related phases for this steel instead of Ti/Nb/V rich MX precipitates. It is always worried for metallurgists and materialists that mean size of precipitates should be finer so that during service conditions it must provide enough strength by pinning the different GBs against diffusion and migration which resulted in smaller voids. So, Figures 4.22 (a) to (e) are plotted for AR material, CGHAZ D05, FGHAZ D02, ICHAZ D01 and CG + FGHAZ D01 at different PWHT conditions. Which represents the size of the precipitates during nucleation and growth of new precipitates during PWHT treatment by assuming bulk nucleation sites. In these graphs M23C6 having very large nucleation barrier at 760 oC and 800 oC due to taking assumption as bulk nucleation sites. However, it is confirmed that 760 oC/3h of PWHT produces the optimum size of precipitates. At 650 oC and 800 oC of temperature, very small and large size of precipitates are nucleating which may fail to provide pinning effect during serving exposure. Even though critical size for such precipitates at 760 oC is also optimum for such precipitates at different simulation conditions as shown in Figures 4.21 (a) to (e) which agrees


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with mean radius plots. Driving force to nucleate these phases into this material at different PWHT conditions are shown in Figures 4.22 (a) to (e). These plots represent that at fixed temperature driving force for nucleation of a precipitate decreases with time in contrast to mean radius of the same precipitates. That means the probability of nucleation of fresh precipitates will decrease and by the time it will be growing with existing precipitates that finally begin to coarsen. Now, the time has come to discuss number densities for these precipitates. However, the size of precipitates matters but they must be uniformly and largely distributed over various GBs to avoid trivial to major undesirable phenomena during service conditions. Figures 4.25 (a) to (e) represent number densities where it has also confirmed that number densities are maximum and minimum for both 650 oC and 800 oC of PWHT conditions among all the samples respectively. Thus, 760 oC/3 h of PWHT treatment is an ideal choice which shows considerable number densities of such precipitates with the desired size of precipitates but the thickness of component may only affect the time for PWHT treatment. At last volume fraction of primary and secondary precipitates are shown in Figures 4.26 (a) to (e). Here, on an average volume fraction of secondary precipitates are 0.5 %. But in the ICHAZ D01 sample, Nb(N,C,B) fraction is raised to 0.9 % which agrees with hardness result as shown in Figures 4.26 (d) and 4.3 (i). Also for CGHAZ D05 sample, Nb(N,C,B) fraction is 0.9 % but due to tempered martensite formation, its effect on hardness is compensated. From the numerical simulation, it was also observed that Mn partition to M23C6 as shown in Figures 4.21 (b) and (c).


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Figure 4.18: JMatPro diagrams.


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Figure 4.19: Thermocalc diagrams for P91B steel, (a) Fe-C diagram, (b) Fe-Cr diagram, (c) Fe-Mo diagram, (d) Fe-V diagram, (e) Fe-N diagram and (f) Fe-N diagram.

Figure 4.20: Binary diagram of P91B steel where 1* region represents no boron nitride.


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68

Figure 4.21: Property diagram and amount of components in a phase diagrams for P91B steel by Thermocalc.


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Figure 4.22: Mean radius (bulk nucleation sites) of various precipitates at different PWHT conditions.


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Figure 4.23: Critical radius of various precipitates at different PWHT conditions.


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Figure 4.24: Driving force of various precipitates at different PWHT conditions.


72

Figure 4.25: Number density of various precipitates at different PWHT conditions.


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Figure 4.26: Volume fraction of various precipitates at different PWHT conditions.

4.6: Other tests 4.6.1: Impact energy for CGHAZ D04 The Charpy test is done on the standard size of the sample for CGHAZ D04 at room temperature which has given impact energy of 46.36 J. For such material is very low so PWHT is required to improve impact energy and the failure was brittle.

4.6.2: Dilatometry for Ac1, Ac3, Ms and Mf The AC1 and AC3 are found for this material at different conditions by using copper grips and taking a span of 30 mm. The different transformation range are shown in Figures 4.27 (a), (c) and (e) and given in Table 4.5. Whereas, their thermal profiles are shown in Figure 4.27 (b), (d) and (f). Ac1 and Ac3 are always a function of heating rate and sometimes also affected by minor changes in the composition of the material. It is also good to see that Mf temperature in all the cases are considerably higher than RT which ensures minimum RA after welding simulation as 100 % phase transformation is difficult to get. After observing all the dilation curve look at ICHAZ simulated sample it shows both Ac1 and Ac3 temperature are above 900 oC. But in earlier simulation of ICHAZ samples as shown in Figures 4.1 (f) and (g) which are done by using 304L austenitic stainless steel cooling grips with 10 mm of span at peak temperature of 865 oC and their microstructure clearly represents these samples were undergone transformation range between Ac1 and Ac3 which resulted in fresh and tempered martensite as shown in Figures 4.6 (l) and (n) and 4.9 (i) to (l). That means for the same material, equal heating rate and identical geometry of sample, Ac1 and Ac3 are depending on span and type of gripping material. National Institute of Technology Warangal – 506004

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Table 4.5: Dilatometry results for P91B steel. S.

Simulation Thermal history

no.

type

01

CGHAZ

100 oC/s-1240 oC -3s-30s

Ac1 [oC]

Ac3 [oC]

Ms [oC]

Mf [oC]

933

1040

392

261

926

1013

416

275

922

924

434

320

(t8/5: Exponential cooling) 02

FGHAZ

100 oC/s-1020 oC -4s-30s (t8/5: Exponential cooling)

03

ICHAZ

40 oC/s-920 oC -8s-30s (t8/5: Exponential cooling)


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Figure 4.27: Plots of dilatometry measurements and thermal histories. Figures 4.28 (a) and (b) represent Gleeble thermal profiles of FGHAZ D04 and ICHAZ D04 respectively although, their microstructures are shown in Figures 4.29 (a) and (b) respectively. These microstructures are free from coarse grains and also in the non-simulated zone of these samples containing nearly similar microstructure which is found in AR material that means simulation is perfect and desirable.

Figure 4.28: Thermal profiles of (a) FGHAZ D04 and (b) ICHAZ D04.


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Figure 4.29: Microstructure of (a) FGHAZ D04 etched with a mixture of 28% nital and Vilella’s reagent and (b) ICHAZ D04 etched with 28% nital.

4.6.3: Furnace heat treatment of P91B steel

Figure 4.30: Thermal profile given to P91B steel during furnace treatment. Furnace treatment of P91B steel is done by the direct introduction of the sample at two different temperatures i.e. 865 oC and 1250 oC for one hour of holding time to observe grain growth during peak temperature which was normally air cooled. Its thermal profile is shown in Figure 4.30. The microstructures of 865 oC sample are shown in Figures 4.31 (a) and (b) which depicts no grain growth except large recovery, transformed region and carbides coarsening in this condition. Figures 4.31 (c) and (d) represent microstructure of furnace heat treated sample at 1240 oC contain martensite and ferrite. The two phases are also identified by microhardness testing which shows dark and light grey is martensite of hardness 449 HV/0.05 and white is ferrite as shown in Figures 4.31 (e) and (f) but it is difficult to say all white phases are alpha


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ferrite as the temperature was in the range of delta ferrite which is calculated from Thermocalc diagrams so it was delta ferrite. Also, the critical cooling rate for this steel 0.2oC/s that means after cooling there will be no alpha ferrite unless it already has undergone hogh temperature ferrite zone and boron delays ferrite formation.


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Figure 4.31: Microphotographs of furnace treated samples (a) and (b) 865 oC/1h, (c) and (d) at 1240 oC/1h, micro-indentation of (e) soft phase and (f) hard phase.

4.7: Effect of peak temperature on microstructure 4.7.1: AR microstructure interpretation P91B steel has been designed to undergo elevated temperatures service of the order 6500C. This steel generally experiences three stages i.e. as welded condition; post weld heat treated conditions and finally elevated temperature exposure with stress levels quite below the yield point of the material. But such low-level stresses and elevated temperatures invoke creep voids in weld HAZ of P91B steel thus reducing the targeted design life of the component. Creep life estimations and investigations are outside the scope of this work for P91B steel. Discussions and inferences drawn are based on the use of different etchants for revealing microstructure, experiments performed for hardness evaluation, and evolution of microstructure before and after PWHT of simulated weld HAZ subzones for P91B steel. High alloying elements already delayed in transformation thus air cooling is sufficient to form martensite after normalising treatment this martensite phase is transformed to the thermodynamically stable phases, like ferrite and carbides (cementite in plain carbon steels) during tempering [15]. As the microstructure also represented some untransformed laths of martensite. The microstructure of P91B steel was tempered martensite with densely populated dislocations by XRD as shown in Table 4.2. These dislocations are present within subgrain of lath martensite and finely dispersed laths of martensite that are restricted by M23C6 carbides. However, thermocalc also predicts the presence of boron enriched Ti/Nb/V carbonitrides which are MX type carbonitrides in AR material as given in Figure 21 (a). Carbide/carbonitrides precipitates of MX-type are found at PAGBs, SGBs, packet, block and lath martensite boundaries etc. [6, 34]. This steel does not have boron nitride as the nitrogen 40 ppm and 100 ppm of boron as shown in Figure 20 and boron is known to give memory effect [8]. The triple point is the point where higher triaxiality encountered and this point promotes nucleation of cavities during service temperature.

4.7.2: CGHAZ samples The simulation temperature for 1240 oC (1513 K) simulated sample is much higher than Ac3, hence this condition is referred as CGHAZ. The simulation was done by taking linear heating up to 1513 K with 3 seconds of holding time as shown in Figure 4.1(a) to (c) which not only


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promoted grain growth of austenite grains but also dissolved all the precipitates. Which were observed in AR material and as revealed by thermodynamic calculations from Thermocalc software (see property diagram). Large grain growth could be held responsible for the use of austenitic steel grips during Gleeble simulation of CGHAZ samples. Since austenitic stainless steel grips of 304L grade have low thermal conductivity than copper which produces low thermal gradient which results in large driving force for grain growth. The defined temperature for CGHAZ simulation was increased to 1250 oC during holding time, but this has less contribution to extensive grain growth compared to low thermal gradient created by steel grips since CGHAZ simulation temperature is already higher than Ac3 temperature. On cooling from this temperature which is above normalising temperature, new Cr and Fe-rich carbides precipitated out of the matrix which could be M23C6 particles. The hardness (see Figure 4.3 (c) to (e)) and dislocation density ( see Table 4.2) of CGHAZ samples are considerably higher than all other simulated subzones of HAZ.

4.7.3: FGHAZ samples The peak temperature for 1040 oC (1313 K) simulated sample is just above the Ac3 temperature i.e. in the range of normalising temperature which infers this sample to be FGHAZ. This sample was linearly heated to 1040 oC and held for 4 seconds at this temperature. As the temperature was not enough to completely austenize and dissolve all precipitates except M23C6 and V(N,C,B), hence the excepted grains must have been finer. But again, due to the use of steel grips which are producing low thermal gradient contributing to either complete dissolution of precipitates (boron enriched Ti/Nb carbonitrides) providing pinning effects or reduction of pinning effect or both. Their presence was assured from property diagram calculation of Thermocalc and reported by Lee, Hongyeob, et al. [21]. However, the programmed temperature of 1040 oC was increased to 1050 oC as revealed in the thermal output profile of FGHAZ as shown in Figures 4.1 (d) and (e) having no effect on grain growth because the normalizing temperature range for 9Cr steels is 1313 K to 1353 K. Also, AR material was received in the normalized condition of 1333 K/30 mins resulting in 17.65 µm average grain size, so major contribution for extensive normal grain growth in FGHAZ goes to steel grips. After cooling, Fe/Cr/Mo rich triple point carbides were formed. Also, very fine laths of martensite with high dislocation density and large distribution of M23C6 type in comparison to as received material were produced.


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4.7.4: ICHAZ samples For 865 oC (1138 K) simulated samples were linearly heated to peak temperature which lies in between Ac1 and Ac3temperature zones so the transformation range corresponds to ICHAZ. In this sample, the holding time kept was 8 seconds that produced incomplete transformation of tempered martensite into austenite with partial dissolution of primary and V(N,C,B) precipitates but secondary Ti/Nb carbonitrides were still present since they dissolve above 1323 K. Rapid coarsening of existing Fe/Cr-rich carbides which are M23C6precipitates (give paper reference)along with high recovery of martensite lath structure with expected high dislocations into subgrains/substructures with low dislocations resulted in lowest hardness among all tested samples but the dislocation density after PWHT increase and the cause of increasing is unclear. Although the grains should have been finer, but the average grains are coarse but finer than FGHAZ samples. Therefore, it is difficult to conclude that only thermal conductivity of 304L austenitic steel grips is alone responsible for delaying in extraction of heat input during holding and cooling time such cause negligible thermal gradient across the sample surface than copper grips, the increased temperature upto 1143 K for the thermocouple spot welded surface have been noted and later the material was exponentially cooled to room temperature as shown in Figure 4.1. The microstructure evolved after cooling consisted of fresh martensite with a tempering of existing martensite, newly formed precipitates and coarse precipitates [4, 23, 28]. However, fresh austenite formed at PAGB and lath boundaries in contrast to as received material [53]. 1138 K simulated sample was primarily characterised by lower hardness than AR material, 1513 K simulated sample and 1313 K simulated sample. Softening is reported at 900 o

C due to unavailability of NbC & VN precipitates [8]. ICHAZ samples have a mixture of fine

grains and coarse grains [3]. Fresh austenite formed at PAGB and lath boundaries that transformed into martensite and normally tempering of existing martensite [4].

4.7.5: CG + FGHAZ D01 CG + FGHAZ is multipass weld HAZ simulation process so it has mixed grains of CGHAZ and FGHAZ. In this sample was treated similar to CGHAZ but during second when the sample was heated to 1040 oC (1313 K) and held for 4 seconds. Which caused three different lath morphologies like thin, wide and split laths have observed as shown in Figures 4.6 (p) and 4.9 (n). Though, martensite laths are finely distributed in a matrix which is surrounded by very fine precipitates nucleated during FGHAZ cycle. In this case, during heating, linear profile is acceptable, while cooling, exponential profile adopted. Both the cycles got separated around National Institute of Technology Warangal – 506004

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702 K. During the2nd cycle, instead of cooling to 415 K, the sample was cooled to 459 K, as shown in Figure 4.1. As the first cycle corresponded to CGHAZ causing complete dissolution of all the precipitates. As all the precipitates dissolved in the matrix and cooling it from the above normalising region, the hardness and dislocation density are considerably high [39]. It is also observed from thermal profiles of the same sample that both the surfaces are non-uniformly heating and cooling.

4.9: Boron effect on microstructure Anyhow, we observed grain growth in the simulated samples but point is to be marked that PAGBs are still present after simulation and growth were uniform such effects may be due to the presence of boron which suppresses non-uniformity in grain size. B inhibits M23C6 coarsening by replacing chromium by itself and reducing its surface energy, M 23C6 particles stabilises lath martensitic structure. Fine precipitates observed in CGHAZ but in thecase of ICHAZ, these are coarse [8]. Coarse M23C6 particles are observed after and before welding at PAGB and packet/block boundaries [34]. As it has been confirmed from the property diagram of P91B steel obtained from Thermocalc that boron is not only present in M23C6 carbides but also is present in MX type precipitate. Thus, boron replaces carbon atom of M23C6 by itself that reduces its coarsening & recovery of martensite matrix along colonies i.e. better creep properties compared to boron free steels, it also forms same grain sizes & lath orientations between BM and HAZ this is called memory effect which means inherently retarded type IV cracking, including reduced GB energy and delayed austenite formation during heating operations resulted in fine grain formation [8, 13]. Excess boron and nitrogen form intermetallic boro-nitrides (BN) phase during normalising [8, 13, 16, 45].


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Chapter 5: Summary and Conclusions 1. The increased peak temperature during HAZ simulation of P91B steel samples was no significant effect on microstructural grain growth by using both copper and steel grips. Output profiles deviated through defined profiles during cooling were not restored for steel grips but for copper grips it restored. At a cross-section, both the surfaces were experienced different thermal profiles and exponential cooling was found better than linear cooling. 2. The high hardness in CGHAZ sample infers the necessity of preheating before welding to avoid hydrogen induced cracking. CG + FGHAZ D01 sample showed large variations in evaluated hardness. 3. All macro-indentation profiles were free from the transient zone between simulated and non-simulated zones which resulted in difficulty to identify simulation zone especially in ICHAZ. 4. Macro-indentation test was not sufficient to understand hardness variations due to the large depth of indentation and limited simulated volume so micro-hardness was employed for further investigation. 5. P91B steel having lath martensite structure with a large number of carbides. PAGB of martensite known as HAGBs due to high misorientation consists of packets, blocks and laths and which are also called subgrains having low angle misorientation. 6. For CGHAZ samples, the wide spacing in laths and a small number of coarse precipitates were observed. Whereas in FGHAZ samples uniform distributions of lath with fine precipitates were observed however ICHAZ samples contained tempered and fresh martensite with coarse precipitates which furthermore coarsen during PWHT. And at last in CG + FGHAZ sample, long and thin, wide and split laths were observed due to second pass welding simulation. 7. From hardness profiles, disc polishing and etching, it was also confirmed that simulated volume of weld HAZ was directly depended upon peak temperature. Sample simulated by using copper grips are a true replica of HAZ in terms of the microstructure. 8. Steel grips caused low thermal gradient along the length of the sample resulting either reduction or annihilation of pinning effects. So, great normal grain growth at above Ac1 temperature. Whereas, at below Ac1 temperature it promoted high/low temperature tempering for comparatively less time period than time for PWHT thus resulting in coarsening of


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precipitates as mainly observed in CGHAZ and loss of hardness in both simulated and nonsimulated zones. 9. CGHAZ D04 sample before PWHT has poor impact toughness so PWHT is necessary. For same material, similar heating rate and identical geometry of specimens, Ac1 and Ac3 are depending on span and type of gripping material. 10. After PWHT, slow recovery and recrystallization took place which caused tempered martensite structure with substantial amounts of carbides/nitrides at different GBs. However, simultaneously nucleation, growth and coarsening happened. 11. Adequate percentage drop in hardness gradient from standard deviation was achieved. This suggested that PWHT of 760 0C/3 h was sufficient for single pass welding whereas the same is not recommended for multi-pass welding simulation. So, further study is required to for this sample. Low hardness in ICHAZ, indicates type IV cracking susceptibility. 12. In these samples, dislocation densities were decreased after PWHT. The numerical calculation was done and found less than 4 % RA. 13. Anyhow, we observed grain growth in the simulated samples but point is to be marked that PAGBs are still present after simulation and both substructure grain and growth were uniform such effects may be due to the presence of boron which suppresses non-uniformity in grain size. So, further studies are required to understand such mechanism. 14. From the CCT diagram, it was confirmed that simulated samples were transformed into martensite on cooling. CGHAZ samples may be undergone delta ferrite region. Thin film of delta ferrite found at PAGB. 15. Thermocalc predicted only chromium carbide of M23C6 was stable without laves, BN and other boron related phases in this steel. Also, it predicted for PWHT condition that no Zphase nucleation rather than nucleation of fine and substantial amounts infraction for both boron rich M23C6 and boron rich MX precipitates with considerable number densities which stabilise matrix. 16. The dissolution temperatures of M23C6, Ti(N,C,B), Nb(N,C,B), V(N,C,B) and Z-phase are 875 oC, 1230 oC, 1150 oC, 860 oC and around 730 oC respectively. 17. The average volume fraction of secondary precipitates was 0.5 % but in the ICHAZ sample Nb(N,C,B) fraction was raised to 0.9 % this could be the reason of increase hardness. It was also concluded that at fixed temperature driving force for nucleation of fresh precipitates decreases with time. It also was observed that Mn partition to M23C6. 18. Furnace heat treated sample at 865 oC was not showing any growth. While growth was observed for 1250 oC sample and produced ferrite and martensite by air cooling. National Institute of Technology Warangal – 506004

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19. Aqua regia and Vilella’s reagent is not always recommended for etching of Gleeble simulated weld HAZ samples. However, HNO3 and concentrated corrosive/oxidising reagents accelerated the effectiveness of etching. Picric acid-based reagents uniformly etch and reveal martensite features like PAGBs and SGBs. 20. Samples with identical heat treatment were etched with same etchants irrespective of past thermal history. Nital does not always reveal PAGB for P91B steel as it reacts with segregated tramp elements and other precipitates at PAGB. Furnace sample of 865 oC where the large transformation of ferrite occurs their GB are not clearly revealing. 21. Non-homogeneously etching may be due to non-homogeneity in microstructures. Also, microstructures and hardness plots suggested that Gleeble simulated samples were nonhomogeneous.

Future work and scope For this research work 1. Effect of Cr2O3 layer on creep strength at 650 oC for P91B steel. 2. Effect of boron, titanium and rhenium on the creep strength of P91B steel . 3. The propensity of memory effect by addition of boron generally 90 to 130 ppm on MX type carbonitrides and M23C6 type carbides in 9Cr steels. 4. PWHT condition to minimize hardness and microstructure gradient for multipass welding simulation. For this material 1. The scope of monitoring, detecting and repairing of weld joints for type IV cracking investigation in countries like India and China where newly built power plants are made of creep resistant steels. 2. Laser, electron beam and magnetic pulse weldings are strong fabrication methods to minimise type IV cracking by the formation of narrow HAZ in the weldments that suppresses microstructural gradients. 3. Study on electron beam welding process that eliminates the selection of filler wire, coupled with normalising and tempering treatments. 4. Plants operating at temperature more than 700 oC there is fabrication of joints between martensitic steels and nickel based alloys, theses joints have difficulty on treatments like cross weld testing, non-destructive testing, filler selection, PWHT (during PWHT carbon diffuses to high alloys from low allow steels caused carbon depleted zone, EPRI


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P87 is nickel based consumable alloyed with 9%Cr that minimizes diffusion of carbon at the time welding. 5. Precipitation by fine Z-phase including welding of such steels followed by PWHT. 6. Avoiding/mitigation of Z phase precipitation. 7. Development of boron containing consumables for these steels. 8. Oxidation resistance of these steels is not systematically studied. 9. Re-containing steels show high creep rupture strength. 10. Effect of W on M23C6 composition change during ageing and solid solution strengthening of the 9Cr matrix.


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Bibliography and references [1] Hyde, C. J., et al. "Thermo-mechanical fatigue testing and simulation using a viscoplasticity model for a P91 steel." Computational Materials Science 56 (2012): 29-33. [2] Prasad, Srinivas, V. B. Rajkumar, and Hari Kumar KC. "Numerical simulation of precipitate evolution in ferritic–martensitic power plant steels." Calphad 36 (2012): 1-7. [3] Abson, D. J., and J. S. Rothwell. "Review of Type IV Cracking of Weldments in 9–12% Cr Creep Strength Enhanced Ferritic Steels." International Materials Reviews 58.8 (2013): 437-473. [4] Francis, J. A., W. Mazur, and H. K. D. H. Bhadeshia. "Review type IV cracking in ferritic power plant steels." Materials Science and Technology 22.12 (2006): 1387-1395. [5] Goswami, P. "P (T) 91 steel—a review of current code and fabrication practices." Proceedings of the Sixth International Conference on Advances in Materials Technology for Fossil Power Plants, La Fonda, USA. 2010. [6] Maruyama, Kouichi, Kota Sawada, and Jun-ichi Koike. "Strengthening mechanisms of creep resistant tempered martensitic steel." ISIJ international 41.6 (2001): 641-653. [7] Li, Huijun, and David Mitchell. "Microstructural Characterization of P91 Steel in the Virgin, Service Exposed and Post‐Service Re‐Normalized Conditions." steel research international 84.12 (2013): 1302-1308. [8] Bhaduri A.K., and Das C.R. “Type IV cracking and its mitigation in 9-12% chromium steels.” National Welding Seminar (2014-15): 1-10. [9] Yan, Wei, et al. 9-12Cr Heat-Resistant Steels. Switzerland: Springer, 2015. [10] Vuherer, Tomaž, et al. "Microstructural investigation of the heat-affected zone of simulated welded joint of P91 steel." Metalurgija 52.3 (2013): 317-320. [11] Laha, K., et al. "Characterization of microstructures across the heat-affected zone of the modified

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[39] Paul, V. Thomas, S. Saroja, and M. Vijayalakshmi. "Microstructural stability of modified 9Cr–1Mo steel during long term exposures at elevated temperatures." Journal of Nuclear Materials 378.3 (2008): 273-281. [40] Shi, Pinfang, et al. "Thermo-Calc and DICTRA enhance materials design and processing." Materials Science Forum. Vol. 475. Trans Tech Publications, 2005. [41] Laitinen, Risto Olavi. Improvement of weld HAZ toughness at low heat input by controlling the distribution of M-A constituents. Vol. 68. No. 01. 2006. [42] Brear, J. M., and A. Fleming. "Prediction of P91 life under plant operating conditions." ETD Int. Conf. on High Temperature Plant Integrity and Life Extension. 2004. [43] Brett, S. J. "Service experience with a retrofit modified 9Cr (Grade 91) steel header." Proceedings of the EPRI 5th International Conference on Advances in materials technology for fossil power plants, Marco Island, Florida. 2007. [44] Coleman, Kent K., and W. F. Newell. "P91 and Beyond." Welding Journal-New York- 86.8 (2007): 29. [45] Holzer, Ivan. Modelling and simulation of strengthening in complex martensitic 9-12% Cr steel and a binary Fe-Cu alloy. Verlag der Techn. Univ. Graz, 2010. [46] Albert, S. K., et al. "Effect of welding process and groove angle on type IV cracking behaviour of weld joints of a ferritic steel." Science and Technology of Welding and Joining 10.2 (2005): 149-157. [47] Storesund, Jan, L. E. Samuelson, and B. Klasen. "Creep life assessment of pipe girth weld repairs with recommendations." 3rd int. HIDA and INTEGRITY conf. on integrity of high temperature repair welds, Oeiras-Lisbon, Portugal. 2002. [48] Dieter, George Ellwood, and David J. Bacon. Mechanical metallurgy. Vol. 3. New York: McGraw-Hill, 1986. [49] Morimoto, H., et al. "Effects of W on mechanical properties at elevated temperature of 9Cr ferritic heat‐resistant steel weld metal." Welding international 14.1 (2000): 35-47. [50] Venugopal, S., G. Sasikala, and Yatindra Kumar. "Creep Crack Growth Behavior of a P91 Steel Weld." Procedia Engineering 86 (2014): 662-668.


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Appendices Appendix. I

Brief introduction to HAZ simulation Guide

A.1.1: Introduction Gleeble has developed to study weld HAZ phenomena founded by Dynamic System Inc. (DSI) in 1957, that is why called pioneers of physical simulation, first export to Society de France in June 1962. A.1.1.1: What is Physical Simulation? a. Physical simulation is the replication of the mechanical and thermal parameters of material processes that materials are actually undergoing a real fabrication process. b. It requires accurate measurement and dynamic control of variables such as strain and strain rate, temperature, heating rate and cooling rate etc. so that material could follow similar thermal and mechanical profile, relatively smaller test piece is used in the physical simulation. c. When the simulation is done accurately, the results can be readily transferred from the laboratory to mass scale productions. A.1.1.2. Why Physical Simulation? a. Let us take an example of welding, is there possible to test FGHAZ to predict type IV cracking, no it is very difficult, but with physical simulation reproduction of actual process is easy and extraction of results to understand metallurgical and mechanical properties is now possible. b. Physical simulators are not only reproduced materials but also save lots of money and time, nowadays they are frequently adopted at industrial and R & D levels for a wide variety of new and existing materials at various thermal and mechanical variables. c. This simulator has the ability to control input parameters and generates a road map of given set of parameters that affects characteristics of simulated sample. In the last 3 decades’ engineers, scientists and operation researchers throughout the world continuously using the simulator to map new and current materials for better understanding. d. Gleeble has a unique capability of maintaining uniform temperature with the selected cross section, known as isothermal planes at mid-portion of a test piece, no matter materials are subjected to heating or cooling rapidly which produces confident to materialist and metallurgist that tests are done correctly.


2

e. Physical simulation of welding and powder metal processes can be performed. In welding and joining, a system allows precise control of energy input which generates accurate profiles. Researchers can simulate heat-affected zones on limited amounts of materials. A.1.1.3: Fundamentals of Physical Simulation Furnace Heating, induction heating and bulk heating have been compared and shown in the Table A1.1 with respect to variables like heating mode, heating rate and sample size etc. Table A1.1: Comparison of heating methods based on different testing variables. Parameters/media

Furnace

Induction heating

Heating mode

Radiation & convection

Eddy current heating Joule Heating

Heating Rate

Slow at 1°C/s

Medium at 200°C/s

Temperature

Uniform after soaking

Uniform

uniformity

Bulk heating

Fast 10,000 °C/s

after Uniform

soaking

heating

Sample size

Large

Medium

Small

Testing Efficiency

Low

Medium

High

during

A.1.2: Gleeble 3800 Gleeble is fully integrated digital closed loop control thermal and mechanical testing system, easy-to-use with Windows based computer software, provides a user-friendly interface to create, run and analyse data from thermal-mechanical tests and physical simulation programs, these systems characterize, optimize and simulate materials. The multitasking graphical user interface is to provide simulation program and analyse the output to create reports and presentations, a number of programming, including QuikSim Software, a spreadsheet-like, fillin-the-blanks software that describes each action in a test sequence in order and duration, allows arbitrary programming of waveforms for both thermal and mechanical systems. Predefined test programs can be run without any change or with change by Virtual Panel Meters (VPMs) to modify the programs when the process is in progress after the end of test results automatically transferred to Origin software that equipped with mathematical tools. 7.1.2.1: Equipment Gleeble Load Unit –it contains hydraulic system, vacuum tank containing the specimen being tested. (Gleeble 3800 models accept Mobile Conversion Units (MCUs) which contain a vacuum tank) as shown in Figure 3.2.


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Digital Control Console – This free-standing control console contains the embedded computer to monitor real-time activity as shown in Figure 3.2. Thermocouple Welder – It is used in many Gleeble applications to record temperature, principle of this welder is the resistance spot welding which is done at 32V and in this simulation type K has been used, and 0.2 mm diameter of thermocouple is welded to the specimen at middle portion and push pots used to record temperature as shown in Figure 3.2. Figure A1.1 shows thermocouple welder and a pipe containing two types of joints out of which beads must be avoided to minimize errors. Two thermocouple wires are to be joint about 1mm apart in the same cross section to avoid temperature measurement errors. Vacuum System – it is used to evacuate air from the Vacuum Tank to control the system around the test piece. This System has a roughing pump (capable of creating a vacuum to 10-2 torr.) Additionally, it can accommodate the backfill of inert gas into the vacuum tank.

Figure A1.1: (a) Thermocouple welder (b) Thermocouple Push Posts that used to couple welded thermocouple to generate thermal/mechanical profiles and (c) weld beads that permit correct simulation. Following points should be kept in mind to avoid above errors. National Institute of Technology Warangal – 506004

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a. Sample surface must be clean with no grease, no scale to do so emery papers can be used. b. The specimen must be in good electrical contact with the specimen holder in the welder. c. The thermocouple wire must be in good electrical contact with the drop arm of the welder. Up to 1250oC, the process can be done using Type K thermocouples. Although, the holding time at peak temperature should decrease when the peak temperature is close to 1250oC. 7.1.3.2: Grips Copper grips are often used for HAZ simulation, Different specimens use different grips. When loading a specimen, the thermocouple must be located in the middle of the free span, and the grips must be clean, otherwise, the thermal profile of the specimen will not be symmetrical, in this simulation 304L austenitic stainless steel square grips. Figure A1.2 depicts the thermal profile of AISI 1018 that concluded copper grips are the best grips for physical simulation followed by full contact that decreases thermal gradient along the length of the specimen.

1200

Temperature (°C)

1000

800

Copper Jaw Full-contact

600

Half-contact Hot Jaw

400

AISI 1018 Steel 10 m m Diam eter

200

0 -20

-15

-10

-5

0

5

10

15

20

Free Span (mm)

Figure A1.2: Thermal profile of AISI 1018 steel with respect to different grips. This profile purpose is to demonstrate how cooling grips may alter your simulation that causes misinterpretations like non-simulated zone also taking part in simulation or this is the actual candidate material, in our case steel grips try to do so but we have observed them.


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Appendix. II

Metallographic preparations and microstructural analysis by Light and Electron Microscopy Guide

Etching response of different etchants on microstructures of Gleeble simulated samples It is well understood that for revealing the microstructures which are heat treated at different peak temperatures require different type and composition of etchants including different time for etching. Therefore, it becomes difficult to prescribe one specific type of etchant which would reveal PAGB for Gleeble simulated HAZ samples particularly in the case of creep resistant steels which have martensitic/ferritic microstructures. It is noticed that 9Cr steels have high oxidation resistance; especially for weld HAZ simulated samples of P91B steel, thus highly concentrated corrosive/oxidation reagents are required to reveal PAGB in this steel. Many etchants were tried to reveal PAGB in simulated HAZ subzones of P91B steel in this work. Table A2.1 shows the effect of different etchants attempted to reveal microstructure on Gleeble simulated weld HAZ samples of P91B steel. The microstructures of the non-simulated zone of different samples have been shown in figures A2.1 (a) to (j).

National Institute of Technology Warangal - 506004

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Figure A2.1: Microstructures of non-simulated region (a) CGHAZ D01 etched with 4% picral, (b) CGHAZ D04 etched with 28% nital, (c) CGHAZ D05 etched with Vilella’s etchant, (d) FGHAZ D02 etched with 28% nital, (e) FGHAZ D02 etched with Marshall reagent, (f) FGHAZ D02 etched with Vilella’s reagent, (g) FGHAZ D02 etched with 6% picral, (h) FGHAZ D03 etched with 6% picral, (i) ICHAZ D01 etched with 6% picral and (j) ICHAZ D03 etched with Vilella’s reagent. Methanol was selected to improve image contrast rather than using ethanol. Alkaline sodium picrate, oxalic acid in distilled water, Marshall’s reagent, glycerine based etchants, a solution National Institute of Technology Warangal - 506004

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of HCl, HNO3 and HF in distilled water and glycerol associated reagents do not reveal PAGB for Gleeble simulated weld HAZ samples of P91B steel as recommended in the literature for 9Cr steel. As in this steel, Sulphur and phosphorous like elements are under ppm level which segregates to GBs and reacts according to the type of etchants thus helping to reveal PAGB. This could be one of the reasons which cause etching little difficult for Gleeble simulated samples. However, modified picric acid and Curran’s etchant have revealed PAGB; although, pitting was a primary problem after etching which suggests distilled water based etchants require skilled personnel to avoid such circumstances. Extron related etchant requires heating and continuous hit and trail approach to etching and polishing. In such cases, Aqua regia is preferable due to strong corrosive nature but aqua regia etched sample requires grinding or emery paper polishing depending upon provided a degree of etching. Table A2.1: Effects of different etchants tried to reveal microstructural features on Gleeble simulated weld HAZ samples of P91B steel Sample

Composition of etchant 28 drops of HNO3 in 10 ml of ethanol (28% Nital) 5g FeCl3, 5mLHClin 100 mL water (Curran’s reagent)

As received material

1g picric acid, 5ml HCl in 95ml methanol (Vilella etchant)

Aqua regia (3:1 = HCl:HNO3)

Time of etching

Etching response

Remarks

15-60 seconds

Reveals PAGB, laths, precipitates, but matrix remains unaffected

20-45 seconds

Reveals sharp PAGB, laths, but matrix unaffected

For safety, use of hand gloves is recommended Pitting is a primary problem and sticking of FeCl3 powder on surface

30-45 seconds

Reveals two regions i.e. white and black

Unidentified microstructure

Reveals sharp and clear PAGB, laths, matrix remains unaffected

Grinding is essential depending upon degree of etching, due to limited simulated HAZ volume this reagent is

10-20 seconds


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6g picric acid, 5ml HNO3 in 100 ml methanol [6%picral (acid)]

60-240 seconds

Reveals sharp and dark PAGB, lath martensite and precipitates

1-4 minutes

Reveals PAGB, packet, lath and precipitates

30 – 40 minutes

Clearly visible unaffected PAGB, lath martensite and precipitates

2 to 3 minutes

Reveal packets, laths and precipitates

CGHAZ Pre-etch with 57% nital and Vilella’s reagent of 2 g picric acid

6 g picric acid in 100 ml methanol (6% picral) FGHAZ and CG + FGHAZ

Pre-etch with 57% nital and Vilella’s reagent of 2 g picric acid

ICHAZ/ All PWHT treated samples

Pre-etch with 57% nital and Vilella’s reagent of 2g picric acid

1 to 2 minutes


Reveals sharp PAGB and martensite lath

not recommended HNO3 provide dark PAGB and increase degree of etching or preetch Confusing between PAGB and packet, and non-uniform etching Poor contrast due to unaffected PAGB and long etching time It is misunderstood that using only this reagent reveals GBs which are actually packets. Similar situation was faced for FGHAZ and CG+ FGHAZ Recommended for heat treated samples either between Ac1 and Ac3 or below Ac1 temperature

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