Creep Behavior of High Temperature Alloys for Generation IV Nuclear Energy Systems

Home , Dislocation creep, Yield strength anomaly

A dissertation submitted to the Division of Research and Advanced Studies of the University of Cincinnati in partial fulfillment of the requirements for the degree of

Doctor of Philosophy

in Department of Mechanical and Materials Engineering of College of Engineering and Applied Science University of Cincinnati by

Xingshuo Wen

M.S., University of Cincinnati, Ohio, USA, 2009 B.S., Shaanxi University of Science and Technology, Shaanxi, China, 2005 March, 2014

Committee: Dr. Vijay K. Vasudevan (Chair) Dr. Relva C. Buchanan Dr. Rodney D. Roseman Dr. Vesselin N. Shanov Abstract

The Very High Temperature Reactor (VHTR) is one of the leading concepts of the

Generation IV nuclear reactor development, which is the core component of Next Generation

Nuclear Plant (NGNP). The major challenge in the research and development of NGNP is the performance and reliability of structure materials at high temperature.

Alloy 617, with an exceptional combination of high temperature strength and oxidation resistance, has been selected as a primary candidate material for structural use, particularly in

Intermediate Heat Exchanger (IHX) which has an outlet temperature in the range of 850 to

950°C and an inner pressure from 5 to 20MPa. In order to qualify the material to be used at the operation condition for a designed service life of 60 years, a comprehensive scientific understanding of creep behavior at high temperature and low stress regime is necessary. In addition, the creep mechanism and the impact factors such as precipitates, grain size, and grain boundary characters need to be evaluated for the purpose of alloy design and development.

In this study, thermomechanically processed specimens of alloy 617 with different grain sizes were fabricated, and creep tests with a systematic test matrix covering the temperatures of

850 to 1050°C and stress levels from 5 to 100MPa were conducted. Creep data was analyzed, and the creep curves were found to be unconventional without a well-defined steady-state creep.

Very good linear relationships were determined for minimum creep rate versus stress levels with the stress exponents determined around 3-5 depending on the grain size and test condition.

Activation energies were also calculated for different stress levels, and the values are close to

400kJ/mol, which is higher than that for self-diffusion in nickel. Power law dislocation climb- glide mechanism was proposed as the dominant creep mechanism in the test condition regime.

Dynamic recrystallization happening at high strain range enhanced dislocation climb and are

believed to be responsible for the monotonically increasing creep rates. Apart from dislocation

creep, diffusional creep in existence at low stress level in fine-grained (ASTM 8) material also

contributed partly to the creep rates. A reasonable prediction on the long term performance of

alloy 617 was also made by extrapolation method using optimized parameters based on creep test

data.

Furthermore, microstructure characterization was performed utilizing Optical Microscopy

(OM), Scanning Electron Microscopy (SEM), Electron Backscattered Diffraction (EBSD),

Transmission Electron Microscopy (TEM) and related analytical techniques on samples from both before and after creep, with special attention given to grain size effects, grain boundary type, and dislocation substructures. Evidences for dislocation climb and dislocation glide were found through detailed dislocation analysis by TEM, proving the dislocation climb-glide mechanism.

The formation of subgrain boundary, the changes in boundary characters and grain sizes was confirmed by EBSD analysis for dynamic recrystallization. The effects of initial grain size and grain boundary character distribution on the creep behavior and mechanism were also evaluated.

Through the results obtained from this experimental study, new insights were provided into how changes in microstructure take place during high temperature creep of alloy 617, creep mechanism at different conditions was identified, and the creep deformation model was discussed. The results will also serve to technological and code case development and design of materials for NGNP.

Acknowledgements

The completion of this Ph.D. would not be possible without the encouragement and

support from many individuals. First and foremost, I would like to express my sincere gratitude

to my advisor Prof. Vijay K. Vasudevan, who provides his guidance, inspiration, patience, and

support throughout my study in University of Cincinnati and, more importantly, in life. Without

him, it would be impossible for me to accomplish this long journey.

I would also like to thank the committee members Prof. Relva C. Buchanan, Prof.

Rodney D. Roseman, and Prof. Vesselin N. Shanov for their precious time and valuable

suggestions in completing my dissertation.

I would like to thank DOE-NEUP (Program# 09-013, Sub-contract# DE-FX07-00088635) for funding this work. Part of the creep tests were conducted in Idaho National Laboratory, and

Metcut Research Inc. also contributed to some of the low stress tests. The microstructural analysis was carried out in Advanced Materials Characterization Center, University of Cincinnati.

I would like to acknowledge Dr. Laura Carroll and Dr. Richard Wright from Idaho

National Laboratory and Dr. T.-L. (Sam) Sham from Oak Ridge National Laboratory for providing the test specimens and creep data. Their insightful discussions and valuable inputs in this research project are very helpful to my research. Special thanks also go to Randy Lloyd from

INL for coaching me to setup creep tests.

I would also like to thank my friends and colleagues, Dr. Hyojin Song, Dr. Amrinder Gill,

Dr. Yixiang Zhao, Dr. Chaswal Vibhor, Dr. Chang Ye, Dr. Behrang Poorganji, Gokul

Ramakrishnan, Sethuraman Subramanian, Abhishek Telang, and Deepthi Tammana, who shared much of their experience in research, helped me with experiments, and had numerous meaningful discussions.

I am grateful to University of Cincinnati and Department of Mechanical and Materials

Engineering for providing an excellent education and research environment, and for financial support. My thanks are due to the faculty and staff, who contributed to my study and life.

Last but not least, I would like to deeply thank my father for his encouragement and patience. To my wife Yuanyuan Niu, who is always by my side with constant support and endless love, I am so lucky to have you in my life, and I could not go thus far without you.

Table of Contents

1 Introduction ...... 1 1.1 Background ...... 1 1.2 Material Selection for Intermediate Heat Exchanger ...... 4 1.3 Motivation of Research ...... 5 1.4 Objective and Scope ...... 6 2 Literature Review ...... 8 2.1 Chemistry and Microstructure of Nickel Base Superalloys ...... 8 2.2 Strengthening Mechanisms of Nickel Base Superalloys ...... 13 2.2.1 Solid Solution Strengthening ...... 13 2.2.2 Precipitation Strengthening ...... 13 2.2.3 Grain Boundary Strengthening ...... 17 2.3 Background of Alloy 617 ...... 18 2.3.1 Chemistry and Microstructure of Alloy 617 ...... 18 2.3.2 Mechanical Properties of Alloy 617 ...... 20 2.3.2.1 Tensile Property ...... 21 2.3.2.2 Fatigue Strength ...... 24 2.3.2.3 Creep-fatigue Property ...... 25 2.3.2.4 Creep Behavior ...... 27 3 Experimental Details ...... 39 3.1 Specimens ...... 39 3.2 Creep Testing ...... 39 3.2.1 Creep Testing Equipment ...... 39 3.2.2 Creep Testing Condition ...... 42 3.3 Microstructure Characterization...... 44 4 Results and Discussion ...... 47 4.1 Creep Test Data Analysis ...... 47 4.1.1 Specimens with Grain Size of ASTM 3.5 ...... 47 4.1.2 Specimens with Grain Size of ASTM 2 ...... 64 4.1.3 Specimens with Grain Size of ASTM 8 ...... 70 4.2 Microstructure Characterization...... 75 vi

4.2.1 Average Grain Size Determination by Optical Microscopy ...... 75 4.2.2 SEM Analysis ...... 78 4.2.3 EBSD Analysis and Grain Boundary Characterization ...... 80 4.2.4 TEM Observation and Analysis ...... 98 4.3 Static Recrystallization Study ...... 111 4.4 Discussion ...... 117 4.4.1 Dislocation Creep ...... 117 4.4.2 Dynamic Recrystallization ...... 118 4.4.3 Diffusional Creep ...... 127 5 Conclusions and Future Works ...... 132 References ...... 135

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List of Figures

Figure 1.1 General diagram of VHTR design [9]...... 3 Figure 2.1 Alloying elements present in Ni-based superalloys [12]...... 9 Figure 2.2 The unit cell of face-centered cubic (FCC) crystal structure...... 11 Figure 2.3 Unit cell of L12 crystal structure showing arrangements of Ni and Al, Ti atoms in the ordered Ni3(Al,Ti) phase...... 11 Figure 2.4 (a) Yield stress as a function of temperature in Ni3Al, (b) effect of alloying condition on flow stress of γ’ [21]...... 14 Figure 2.5 Formation of a Kear-Wilsdorf lock by cross-slip in the L12 structure [25]...... 15 Figure 2.6 Schematic diagram illustrating the development of γ’ morphologies during aging [31]...... 16 Figure 2.7 TTT diagram for Alloy 617 long-term aged samples [38]...... 19 Figure 2.8 Creep rupture strength of high temperature alloys at 962°C in air [7]...... 21 Figure 2.9 High temperature tensile properties of solution-annealed, hot rolled Alloy 617 rod [33]...... 22 Figure 2.10 Tensile properties of as-received Alloy 617 and 230, (a) Yield Stress vs. Temperature and (b) UTS vs. Temperature [43]...... 22 Figure 2.11 High temperature tensile properties of the alloy 617 weldments, (a) 0.2% yield stress, (b) UTS [44]...... 23 Figure 2.12 (a) Low-cycle fatigue results of Alloy 617 at 950°C in air [46]; (b) Low-cycle fatigue behavior of Alloy 617 at three different temperatures in air environment [47]...... 24 Figure 2.13 The life reduction factor R against the cycle period for different load conditions. R is defined as the ratio of NF/NF0 where NF is the fatigue life for a given strain rate or wave shape and hold time and NF0 is the reference life value for continuous cycling with a strain rate of 4x10-3s-1. Test temperature: 950°C [49]...... 26 Figure 2.14 Stress relaxation for alloy 617 during 60 and 1800 seconds tensile hold in air at (a) 0.3% and (b) 1.0% strain [50]...... 26 Figure 2.15 Comparison of creep and fatigue damage endured by alloy 617 and Haynes 230 based on linear damage summation [51]...... 27 Figure 2.16 Creep strength of solution-annealed alloy 617 [33]...... 30 Figure 2.17 Rupture strength of solution-annealed alloy 617. Arrows denote tests discontinued before fracture [33]...... 30 Figure 2.18 Creep curves of alloy 617 at 850 and 950°C [52]...... 31 Figure 2.19 The effect of fine intragranular carbides on the creep rate. The ratio of creep rates, / are plotted against the creep times for each pair of the same grain size samples. is the creep rate for the solution treated sample, and is that for aged sample휖퐵 휖퐴 [36]...... 32 Figure 2.20 The grain휖퐴 size dependence of (a) the steady state creep rate and (b)휖퐵 the creep rupture life at 1000°C, 24.5MPa [36]...... 32 Figure 2.21 Heat-to-heat variation of creep behavior of Alloy 617 from past US HTGR programs; (a) stress vs. rupture life, (b) stress vs. time to 1% strain [42]...... 33 Figure 2.22 Variation in creep rupture time as a function of applied stress for alloy 617 at 950°C in air and in helium of low steam content [57]...... 34

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Figure 2.23 Effect of oxygen content on (a) creep strain-time curves and (b) creep rupture time and decarburization of alloy 617 in helium [58]...... 35 Figure 3.1 Schematic drawing shows geometry of creep test specimen...... 39 Figure 3.2 Lever Arm Creep/Stress Rupture Testing Systems with lever arm ratio of (a) 3:1 and (b) 20:1...... 40 Figure 3.3 Temperature curve of creep test...... 41 Figure 4.1 Creep strain vs. time plots for alloy 617 specimens tested at (a) 850°C, (b) 950°C, (c) 1000°C, and (d) 1050°C for various stresses...... 50 Figure 4.2 Creep strain vs. time plot for alloy 617 specimens tested at 20MPa for various temperatures...... 50 Figure 4.3 Normalized creep curves for typical high-temperature materials [70]...... 51 Figure 4.4 Creep rate vs. time and creep strain plots for alloy 617 specimens tested at (a) 850°C 50MPa, (b) 850°C 100MPa, (c) 950°C 50MPa, and (d) 950°C 100MPa...... 54 Figure 4.5 Creep rate vs. creep strain plots for alloy 617 specimens tested at (a) 850°C, (b) 950°C, and (c) 1050°C for various stress levels...... 55 Figure 4.6 Plot of minimum creep rate vs. stress for determining stress exponent n of alloy 617 ASTM 3.5 specimens...... 58 Figure 4.7 Minimum creep rate vs. stress plot for alloy 617 specimens compared with historical data...... 58 Figure 4.8 Plot of minimum creep rate vs. 1/T for determining the activation energy Q...... 60 Figure 4.9 Normalized minimum creep rate vs. normalized stress for alloy 617 ASTM 3.5 specimens...... 61 Figure 4.10 Plot for Larson-Miller parameter relationship and projected long term performance of alloy 617...... 63 Figure 4.11 Stress for (a) 1% creep strain and (b) rupture vs. Larson-Miller parameter, for alloy 617 at temperatures between 800 to 1000°C [52]...... 64 Figure 4.12 Plots of creep strain vs. time for alloy 617 ASTM 2 specimens at (a) 850°C, (b) 950°C, and (c) 1050°C...... 66 Figure 4.13 Creep rate vs. creep strain plots for alloy 617 specimens with grain size of ASTM 2 tested at (a) 850°C, (b) 950°C, and (c) 1050°C for various stress levels...... 68 Figure 4.14 Creep strain vs. time plots for alloy 617 specimens with ASTM 3.5 and 2 grain sizes tested at (a) 850°C, (b) 950°C, and (c) 1050°C...... 69 Figure 4.15 Plot of minimum creep rate vs. stress for determining stress exponent n of ASTM 2 specimens...... 70 Figure 4.16 Plots of creep strain vs. time for alloy 617 ASTM 8 specimens at (a) 950°C, and (b) 1050°C...... 72 Figure 4.17 Plot of minimum creep rate vs. stress for determining stress exponent n of ASTM 8 specimens...... 72 Figure 4.18 Creep strain vs. time plots for alloy 617 specimens with ASTM 3.5 and 2 grain sizes tested at (a) 950°C and (b) 1050°C...... 73 Figure 4.19 Comparison among three sets of specimens of alloy 617 at (a) 950°C 5MPa and (b) 1050°C 5MPa ...... 74 Figure 4.20 Optical images taken from (a) rolling direction, (b) transverse direction, and (c) cross-sectional direction of as-received alloy 617 material showing average grain size determination by circular intercept method...... 77

Figure 4.21 Optical images taken from (a) rolling direction, (b) transverse direction, and (c) cross-sectional direction of hot-rolled alloy 617 material showing average grain size determination by circular intercept method...... 78 Figure 4.22 Phase composition in as received alloy 617 by (a) SEM image, and EDS spectrum for (b) Ni matrix and (c) Ti(C,N)...... 80 Figure 4.23 IPF maps for as received alloy 617 samples in (a) cross-section direction and (b) rolling direction...... 82 Figure 4.24 Pole figures for as-received alloy 617 sample from (a) cross section direction and (b) rolling direction...... 83 Figure 4.25 IPF maps for alloy 617 samples underwent creep tests at 850°C 50MPa in (a) cross section direction and (b) rolling direction...... 83 Figure 4.26 IPF maps for alloy 617 samples underwent creep tests at 850°C 100MPa in (a) cross section direction and (b) rolling direction...... 84 Figure 4.27 IPF maps for alloy 617 samples underwent creep tests at 950°C 50MPa in (a) cross section direction and (b) rolling direction...... 84 Figure 4.28 IPF maps for alloy 617 samples underwent creep tests at 950°C 100MPa in (a) cross section direction and (b) rolling direction...... 85 Figure 4.29 Pole figures for creep tested alloy 617 samples at different test conditions, (a) 850°C 50MPa, (b) 850°C 100MPa, (c) 950°C 50MPa, and (d) 950°C 100MPa...... 86 Figure 4.30 Grain size distribution for (a) as-received, and creep-tested at (b) 850°C 50MPa, (c) 850°C 100MPa, (d) 950°C 50MPa, and (e) 950°C 100MPa samples of alloy 617. .. 88 Figure 4.31 IPF maps for shoulder section of alloy 617 specimen underwent 950°C 100MPa creep test in (a) cross-section direction and (b) rolling direction...... 89 Figure 4.32 Grain Boundary Character Distribution (GCBD) charts for (a) as-received sample, (b) creep tested at 850°C 50MPa sample, (c) 950°C creep tested shoulder section (0MPa) sample, and (d) creep tested at 950°C 100MPa sample...... 92 Figure 4.33 CLS boundary fraction vs. stress at 850 and 950°C...... 93 Figure 4.34 CSL boundary fraction vs. creep strain at 950°C 50MPa...... 93 Figure 4.35 Grain boundary maps of (a) as-received specimen, (b) creep tested at 850°C 50MPa specimen, (c) 950°C creep tested shoulder section (0MPa) specimen and (d) creep tested at 950°C 100MPa specimen...... 96 Figure 4.36 Twin boundary fraction in HAGB vs. stress for specimens creep tested at 850 and 950 °C...... 97 Figure 4.37 TEM micrographs of specimen D-61, creep test condition 950°C 50MPa...... 100 Figure 4.38 Stereographic projections for (a) dislocation trace analysis and (b) dislocation slip plane determination for specimen D-61 tested at 950°C 50MPa...... 101 Figure 4.39 TEM images of specimen D-62 area 1, creep test condition 850°C 50MPa...... 102 Figure 4.40 TEM images of specimen 617-14 area 1, creep test condition 850°C 30MPa...... 105 Figure 4.41 TEM images of specimen 617-14 area 2, creep test condition 850°C 30MPa...... 106 Figure 4.42 TEM images of specimen D-62 area 2, creep test condition 850°C 50MPa...... 107 Figure 4.43 TEM images of specimen D-62 area 3, creep test condition 850°C 50MPa...... 108 Figure 4.44 TEM images of specimen 617-15, creep test condition 950°C 50MPa, terminated at 10% creep strain...... 109 Figure 4.45 TEM images of specimen D-61, creep test condition 950°C 50MPa...... 109 Figure 4.46 TEM images of specimen 617-13 area 1, creep test condition 1050°C 30MPa...... 110 Figure 4.47 TEM images of specimen 617-13 area 2, creep test condition 1050°C 30MPa...... 111

Figure 4.48 Hardness for cold rolled alloy 617 samples...... 112 Figure 4.49 Optical microscopy images for alloy 617 in (a) as-received, (b) 10% deformed, (c) 30% deformed, and (d) 60% deformed conditions...... 113 Figure 4.50 IPF maps for alloy 617 in (a) as-received, (b) 10% deformed, (c) 30% deformed, and (d) 60% deformed conditions...... 114 Figure 4.51 IPF maps for alloy 617 samples for (a) 10% deformation annealed at 900°C for 15 minutes, (b) 30% deformation annealed at 900°C for 15 minutes, (c) 60% deformation annealed at 900°C for 15 minutes and (d) 180 minutes...... 115 Figure 4.52 Hardness of samples annealed at 900°C for different time...... 115 Figure 4.53 Grain size vs. cold deformation percentage for alloy 617 samples after annealing at 900 and 1000°...... 116 Figure 4.54 Schematic diagram showing dislocation climb-glide mechanism...... 118 Figure 4.55 Schematic representation of the process of new grain formation during DRX [82]...... 119 Figure 4.56 DRX in Ni and Ni-Co alloys in torsion [87]...... 121 Figure 4.57 EBSD image quality maps with highlighted small grains for (a) as-received, (b) creep tested at 950°C 100MPa, (c) creep tested at 850°C 50MPa, and (d) grain size distribution showing highlighted grain sizes...... 122 Figure 4.58 DRX in tensile test at 1000°C, (a) 2mm from neacking zone, (b) near necking zone [83]...... 123 Figure 4.59 Relationship between grain size and Zener-Hollomon parameter [90]...... 126 Figure 4.60 Relative contribution from diffusional creep and dislocation creep...... 129 Figure 4.61 Stress exponent calculation showing both dislocation and diffusional creep contribute to total creep strain...... 129 Figure 4.62 Deformation mechanism map for pure nickel with a grain size of 0.1mm [93]. .... 130

List of Tables

Table 2.1 Alloying Effects on Nickel Base Superalloys [13]...... 10 Table 2.2 Limiting chemical composition (wt%) of alloy 617 [33]...... 18 Table 2.3 Observation of Second Phase Precipitates in Alloy 617 [7]...... 19 Table 2.4 Summary of the Huntington Alloy Inc. data set of alloy 617 ...... 29 Table 2.5 Summary of ORNL data set of alloy 617 ...... 29 Table 3.1 Creep test conditions for alloy 617 with grain size of ASTM 3.5 ...... 43 Table 3.2 Creep test conditions for Alloy 617 with grain size of ASTM 2 ...... 43 Table 3.3 Creep test conditions for Alloy 617 with grain size of ASTM 8 ...... 44 Table 4.1 Creep rate curve type summary...... 56 Table 4.2 Stress exponent for alloy 617 historical data ...... 57 Table 4.3 Predicted time for alloy 617 to reach 1% and 20% creep...... 63 Table 4.4 Average grain size determination of as-received alloy 617 material...... 76 Table 4.5 Average grain size determination of hot-rolled alloy 617 material...... 77 Table 4.6 Average grain size of alloy 617 from EBSD analysis ...... 87 Table 4.7 Data for Grain Boundary Character Distribution (GCBD) charts ...... 92 Table 4.8 Boundary length data for specimens tested at different conditions ...... 96 Table 4.9 Dislocation analysis results for specimens creep tested at different conditions ...... 103

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1 Introduction

1.1 Background

Nuclear energy is an important energy source on earth, providing 5.1% of the world’s

energy and 11.7% of the world’s electricity production today [1]. International Atomic Energy

Agency (IAEA) reports that there are 437 operational nuclear power reactors and 67 under

construction in 31 countries by the end of 2012 [2]. In the United States, according to U.S.

Energy Information Administration data [3], there are 104 nuclear energy reactors in 65 nuclear

power plants, producing a total of 769.3 TWh of electricity as of 2012, consisting 19% of

national total electricity generation.

There are many debates over the pros and cons of nuclear energy. The biggest advantages

of nuclear energy are its sustainability and low emission. Nuclear energy generally utilizes the

heat produced during the atomic fission process to generate electricity. Uranium as the major

fuel of nuclear energy can be found within all rock, soil and water. It is a naturally occurring

element with an average concentration of 2.8 parts per million in the Earth’s crust, more

abundant than gold, silver, or mercury, about the same as tin, and slightly less than cobalt, lead

or molybdenum [4]. It is estimated that 5.5 million tonnes of uranium exist in ore reserves and more than 35 million tones is available for exploitation [5,6]. The abundance of fuel together with its unparalleled energy density makes nuclear energy a sustainable and secure source of energy. In addition, nuclear energy is a clean method of power generation. Water vapor is the major waste released during the electricity generation process, and there are no air pollutants released, such as nitrogen oxides, sulfur dioxide, volatile organic compounds and heavy metals, which are produced by fossil fuels. It does not even release greenhouse gas like carbon dioxide.

Other major benefits of nuclear energy include high efficiency, reliability, low cost, to name a

1 few. On the other hand, it suffers from some drawbacks, such as radioactivity, storage of nuclear wastes, and safety. Amongst these, the reliability of nuclear reactor is of most concern in industry. Because the consequences of the damage to the reactor, if it happens, are disastrous to the environment, living creatures, and human society, and would have worldwide effect over an extended period of time. Therefore, the structural materials to build nuclear reactor and power plant will have to be carefully tested and proven safe and sound over the designed life time.

Overall, nuclear energy serves as an important and promising energy source to people all over the world. It is necessary to study and develop the technology to make it more efficient and reliable, and thereby to generate more power that benefits the human society.

Next Generation Nuclear Plant (NGNP) is a Department of Energy (DOE) sponsored research program that focuses on the validation of reactor physics, mechanical design, materials research, and safety and risk analysis of the Generation IV nuclear reactor. Among a few types of design, Very High Temperature Reactor (VHTR) or High Temperature Gas-cooled Reactor

(HTGR) has been selected as the basis of the project. The VHTR is a graphite-moderated, helium-cooled reactor, which generates electricity and heat with high efficiency at the same time.

The basic design of VHTR is shown in Figure 1.1. The helium gas turbine system can be set in the primary coolant loop for electricity generation, and nuclear heat is transferred through an

Intermediate Heat Exchanger (IHX) for applications such as refineries, metallurgy, hydrogen production, etc. The VHTR incorporates substantive safety and operational enhancements over existing nuclear reactors. The envisioned NGNP would be able to produce process heat that can be used to power a variety of high energy consuming industries, including electricity generation, hydrogen production, oil refinery, chemical plants, etc., and has a designed service life of 60 years [7].

One of the major advances of NGNP to the existing water-cooled reactor technologies is

the elevated operating temperature, which can be as high as 950°C for the outlet temperature.

Although there are numbers of materials and alloys for high temperature applications having

been in use in the aerospace or chemical engineering industries, a very limited number of

materials have been qualified for use in nuclear reactor systems at such a high temperature.

American Society of Mechanical Engineers (ASME) Codes for reactor applications have only

approved materials for use as high as 760°C, but higher temperature is required in some primary

components for NGNP system [8]. From the commercial point of view, cost is also significant factor in material selection and design. The goal of NGNP is to establish a sound foundation for future nuclear power plant both technologically and economically viable.

Figure 1.1 General diagram of VHTR design [9].

1.2 Material Selection for Intermediate Heat Exchanger

Intermediate Heat Exchanger (IHX), via which heat is transferred from the reactor, is the

fundamental part for VHTR. It separates the primary and secondary power generation circuits and also separates the hydrogen generation plant from the major reactor systems, thus minimizes the potential for accidents, and assures the operation of the reactor will not be impacted by any harmful condition from hydrogen production facility. Based on three VHTR designs, the IHX may be arranged in parallel or in series with the power conversion systems, and designs for IHX include countercurrent shell and tube, involutes, plate with fin, etched plate, microchannel heat exchanger [1,7]. The designed outlet gas temperature is 850°C, with the future target as high as

1000°C for higher efficiency. The pressure differential between primary and secondary loops is relatively low, with 5-10MPa being recommended. The ideal service life for IHX would be 60

years, same as that of the plant.

When it comes to the material selection of IHX, regardless of the design, the material

must be easy to manufacture with availability in both plate and sheet, and good weldability.

More importantly, the candidate materials need to perform well at 850°C or above for long

period of time. The properties brought into consideration include strength, toughness, creep,

fatigue, creep-fatigue, thermal expansion, and thermal conductivity. Irradiation, corrosion, and

thermal aging effects on mechanical properties have to be considered as well. Murty [10]

summarized the desirable characteristics for VHTR structural materials:

(1) Excellent dimensional stability against thermal and irradiation creep, void swelling,

etc.

(2) Favorable mechanical properties such as strength, ductility, creep rupture, fatigue,

creep–fatigue interactions, etc.

(3) Acceptable resistance to radiation damage under high neutron doses, helium

embrittlement, etc.

(4) High degree of chemical compatibility between the structural materials and the

coolant as well as with the fuel.

Currently Alloy 617 by Special Metals and Alloy 230 by Haynes are the two leading

candidates for NGNP IHX application. Both are nickel-based alloys with chromium, cobalt, molybdenum etc. as alloying elements, and have attractive properties for high temperature applications. There have been more studies on Alloy 617 due to its earlier introduction, while

Alloy 230 may have valuable potential for NGNP application. Iron-based Alloy 800H has also been brought into consideration because it is already in the code, but the high temperature properties are not as good. In addition, consideration has also been given to Hastelloy X because

of its excellent oxidation resistance.

1.3 Motivation of Research

Research on structural materials is one of the most critical aspects of the development for

NGNP, because VHTR requires materials performing beyond the limitations of current reactors in high temperature exposure, environmental contaminants, and long-term irradiation. This means materials used in VHTR are expected to have better chemical and mechanical properties in stability, strength, ductility, creep resistance, corrosion resistance, etc. A number of ongoing

NGNP materials research programs are specifically related to the selection, development, and qualification of metallic alloys for high temperature applications such as IHX, core barrel, and core internals such as control rod sleeves.

Among all the aforementioned properties, creep is one of the most important for IHX structural materials, because of the challenge for long term high temperature low stress operation condition. With an exceptional combination of high temperature strength, oxidation resistance, and corrosion resistance to a wide range of environments, Alloy 617, therefore, is considered as the primary candidate material. Introduced in early 1970s, Alloy 617 is largely a solid solution strengthened, nickel-based superalloy, with a nominal chemistry (in wt%) of Ni-20Cr-10Co-

8Mo-1Al-0.6Ti. Although studies on the creep property of Alloy 617 were conducted throughout

1970s to early 1980s by institutions in the US, as well as in Germany, France and Japan, some data are not accessible, some may even be missing or irretrievable. In addition, the data and models are still not completely adequate and the ASME codes and standards need to be expanded to cover the low stresses expected in the application. It is clear that creep behavior varies a lot under different conditions, such as heat, temperature, stress, oxidation and corrosion environment. Therefore, a systematic study on the creep behavior of Alloy 617, particularly used in IHX in the NGNP, is an essential need and of great importance.

1.4 Objective and Scope

This research focuses on Alloy 617, the candidate material for applications as intermediate heat exchangers, with an emphasis on the effects of grain size, grain boundaries and second phases on the creep properties; the mechanisms of dislocation creep, diffusional creep and cavitation; the onset of tertiary creep; and theoretical modeling for long-term predictions of materials behavior and for high temperature alloy design. The research program includes the following key elements:

(1) Thermomechanical processing of Alloys 617 to obtain samples with different grain

sizes;

(2) Determination of the creep properties and phenomenology over a range of temperatures and stresses from low to high covering the diffusional and dislocation creep regimes, respectively;

(3) Characterization of the microstructure of the samples before and after creep using

modern techniques to understand the microstructural mechanisms associated with creep

deformation and damage;

(4) Determination of the parameters in the creep equations and modeling to allow

predictions of the long term creep behavior for design and ASME code case purposes.

2 Literature Review

2.1 Chemistry and Microstructure of Nickel Base Superalloys

Superalloys are a class of alloys that exhibit excellent mechanical strength and creep

resistance at high temperatures, with good corrosion and oxidation resistance. Since the first

development in 1940s for jet turbine engines, superalloys have been widely used in high

temperature applications including gas turbines, rocket engines, chemical and petroleum

processing plants. They are particularly well suited for these demanding applications because of

their ability to retain most of their strength even after long exposure times above 650°C (1200°F).

Their versatility stems from the fact that they combine this high strength with good low-

temperature ductility and excellent surface stability. Depending on the base alloying element,

superalloys are usually categorized into nickel-, iron-, and cobalt-based alloys, and common

addition elements include W, Mo, Nb, Ta, Ti, and Al [11].

Nickel base superalloys develop high temperature properties through solid solution

strengthening and/or precipitation strengthening. For applications requiring modest strength,

solid solution strengthened alloys are usually used, while precipitation strengthened alloys are

used in the most demanding applications such as hot sections of gas turbine engines. The

corrosion and oxidation resistance property is provided by a protective layer of oxidized coat

formed by Al or Cr. Most nickel base superalloys contain 10-20% Cr, 5-10% Co, up to 8% Al

and Ti, and small amount of B, Zr, and C. Other common additions are Mo, W, Ta, Hf, and Nb.

The typical elements alloyed with nickel to form a superalloy are highlighted in Figure 2.1.

Figure 2.1 Alloying elements present in Ni-based superalloys [12].

In general, these alloying elements in nickel base alloys can be categorized by their function into four types. (1) γ formers, elements that tend to partition to the γ matrix. These

elements are Group V, VI, VII, and VIII elements, having an atomic diameter very close to that of the matrix element Ni, such as Co, Fe, Mo, W, Cr. They provide solid solution strengthen to

both austenitic matrix and γ’ precipitate at elevated temperatures. (2) γ’ formers, elements that

partition to the γ’ precipitate. They are from Group III, IV, and V, including Al, Ti, Nb, Ta. They tend to help the formation of stable intermetallic precipitates (Ni3Al, Ti, Ta, Nb), which results in

precipitation strengthening on top of solid solution strengthening. (3) Carbide formers, such as

Cr, Mo, W, Nb. Strengthening is also derived from M23C6, M6C, Ti(C, N) and other precipitates formed by these elements when in appropriate sizes, distributions, and volume fractions. (4)

Grain boundary elements strengthen grain boundaries. They have atomic diameters 21-27%

different to that of Ni, and are from Group III or IV, including B, C, and Zr. They are usually

added in small quantities for control of grain structure and mechanical properties which are

highly related to grain boundaries. The alloying effects of major component elements are

summarized by Wu [13] in Table 2.1.

Table 2.1 Alloying Effects on Nickel Base Superalloys [13].

Alloying Element Effect Co, Cr, Fe, Mo, W, Ta solid solution strengthening carbide formers W, Ta, Ti, Mo, Nb, Hf MC Cr M7C3 Cr, Mo, W M23C6 Mo, W M6 Al, Ti, Ta, Nb γ’ formers, Ni3(Al, Ti) Co reduce stacking fault energy Al, Cr provide oxidation resistance B improve creep properties C, B, Zr strengthen grain boundaries

The microstructure of a typical nickel base superalloy consists of three phases, which

have been reviewed by Sabol and Stickler [14], Decker and Sims [15], and Bridges [16].

Gamma (γ) phase forms a continuous matrix in which other phases reside. It is a face- centered cubic (FCC) structure as shown in Figure 2.2. It contains a significant amount of solid

solution elements such as Co, Cr, Mo, and W. Depending on the solute atom diameters relative

to that of Ni, they will form substitutional solution or insterstitial solution. With the introduction

of these solute atoms, the movement of dislocations will be impeded by the local stress field

created by lattice distortion or modulus change. Therefore, the yield stress of the alloy is

increased, meaning an increase in the strength.

Figure 2.2 The unit cell of face-centered cubic (FCC) crystal structure.

Gamma Prime (γ’) is the primary strengthening phase in nickel base superalloys. This phase has a nominal composition of Ni3(Al,Ti), which formed by the precipitates as the level of

Al or Ti added to Ni matrix increases. It has an ordered intermetallic L12 crystal structure with

Al atoms at the cube corners and Ni atoms at the centers of each face, as shown in Figure 2.3.

Figure 2.3 Unit cell of L12 crystal structure showing arrangements of Ni and Al, Ti atoms in the ordered Ni3(Al,Ti) phase.

The properties of the superalloys are found to be dependent on the coherency of the γ/γ’ interface,

defined by the value of lattice misfit, δ,

′ = 2 × ′+ 푎훾 −푎훾 훿 � 훾 훾� 푎 푎 (2.1)

The closely matched lattice parameters of Ni and Al, Ti, together with chemical compatibility

allows γ’ to precipitate homogeneously throughout the matrix. Therefore, the presence of γ’ is the key to strengthening.

Carbides are usually formed by minor addition of Ti, Ta, Hf, Cr, etc. in the form of MC.

These carbides are in FCC structure with carbon atoms sit in the octahedral sites. They tend to be

distributed on the grain boundaries because diffusion is faster along grain boundaries. They may

be decomposed to form M6C, M23C6, and M7C3 during heat treatment, as described in the

following equations.

+ + (2.2)

23 6 푀퐶 + 훾 → 푀 퐶+ 훾 ′ (2.3)

6 The presence of carbides푀퐶 inhibits훾 → 푀 sliding퐶 훾 ′ by increasing rupture strength at high

temperatures [17]. On the other hand, however, they may be crack-initiation sites which degrade

mechanical properties [18]. The carbides are rigid, and cannot accommodate plastic deformation,

they stress concentration caused between carbides and matrix will lead to microcracking. In

addition, clustering of carbides can result in microcrack joining, which causes intergranular and

transgranular failure [19].

Topologically Close-Packed (TCP) phases have closed-packed atoms in layers separated

by relatively large distance. They usually form as plates or sheets that aligned with octahedral

planes in FCC nickel matrix. TCP phases are detrimental to mechanical properties by depleting

12 strengthening elements and serving as crack initiation sites. This is a major concern when it comes to new alloy design and development.

2.2 Strengthening Mechanisms of Nickel Base Superalloys

2.2.1 Solid Solution Strengthening

When solute elements present in the solid solution, the crystal lattice are distorted by substitutional or interstitial atoms. These solute elements have atomic diameters different from nickel by only 3-13%. Local stress fields are created by introduced strain in the lattice, and with the interaction between local stress field and that of dislocation, the movement of dislocation will be impeded. In another scenario, the modulus of solute atoms differs from the host atom, which results in an increase of the necessary force to propagate the dislocation. At high temperatures

(T>0.6Tm), where diffusion process dominate creep, elements with large atomic diameters having smaller diffusivities in the solid solution become more important in providing strengthening. The governing equation is

= 3 2 Δ휏 퐺푏휖 √푐 (2.4) where c is the concentration of the solute atoms, G is the shear modulus, b is the magnitude of

Burger’s vector, and is the lattice strain due to the solute, it is composed of two terms respectively describing휖 lattice distortion and local modulus change.

2.2.2 Precipitation Strengthening

The γ’ precipitates provide the major parts of the strengthening in Ni-based superalloys.

Al and Ti are the major elements to precipitate a high volume fraction (around 70%) in the Ni

matrix. These precipitates have ordered structure and are stable at high temperatures. In creep,

dislocations must bypass or shear through the precipitates in order to accommodate plastic strain.

As a result, a phenomenon named yield strength anomaly happens. Over a considerable

temperature range, unlike most alloys, yield strength increases with temperature. Westbrook [20]

and Thornton [21] described the phenomenon and discussed the effect of alloying elements on

the flow stress, illustrated in Figure 2.4. Many theories have been discussed on yield strength anomaly in literature [22,23,24], Kear-Wilsdorf lock mechanism is one of them and has been

observed by weak beam TEM [25]. Dislocations with FCC Burgers vectors produce antiphase

boundaries (APB) in L12 structure. They move in pairs and connected by an APB, but dissociate

with stacking fault. Because APB energy on {100} is smaller, dislocations cross-slip to {100},

but leave stacking fault on {111}. Hirsch validated the Kear-Wilsdorf lock at high temperature

and stress conditions [26]. The strengthening effects from the presence of γ’ precipitates are a

related to their size, volume fraction, distribution, fault energy, and the precipitate interface.

(a) (b)

Figure 2.4 (a) Yield stress as a function of temperature in Ni3Al, (b) effect of alloying condition on flow stress of γ’ [21].

Figure 2.5 Formation of a Kear-Wilsdorf lock by cross-slip in the L12 structure [25].

The morphology of γ’ precipitates also plays an important role in strengthening. These

precipitates can grow and change morphoology during heat treatment. Therefore, it is important

to understand the kinetics and its effects. The morphology of γ’ precipitates evolves due to two

different mechanisms: (a) competitive coarsening in order to reduce γ-γ’ interface area (Ostwald ripening), and (b) shape changes in order to minimize interfacial and elastic interaction energies.

Kim [27] and Fisher [28] have examined the coarsening of γ’ precipitates in superalloys at high temperatures, and find out the growth rate fits Lifshitz-Slyozov-Wagner (LSW) theory [29,30], expressed in equation as,

= (2.5) 3 3 0 where d is average precipitate diameter, 푑d0 is− initial푑 푘푡precipitate diameter, k is a constant related to

interfacial energy, diffusivity, temperature, etc. and t is time.

Figure 2.6 Schematic diagram illustrating the development of γ’ morphologies during aging [31].

The morphology of γ’ precipitates in nickel base superalloys generally occurs in a sequence of sphere, cube, array of cubes, and dendrite as coarsening developed by aging, as illustrated in Figure 2.6. The precipitates are firstly nucleate as spheres to minimize surface area.

As the particles grow, the misfit strain energy continues to increase to that a cuboidal and then a cubic morphology that minimizes the interfacial energy forms. This is because crystal planes of the cubic matrix and precipitates remain continuous. However, coherency can be lost by overaging, resulting in directional coarsening and rounding of cube edges.

2.2.3 Grain Boundary Strengthening

Grain boundaries act as strengtheners in superalloys by impeding dislocation movement.

Because of the large misorientation between adjacent grains, it requires more energy for a

dislocation to slip. Dislocation pileup occurs as an array of dislocation travel towards grain

boundary but unable to move across it. This causes a stress against the grain boundary and act as

a driving force for dislocations. Larger grains have more dislocations pileup leading to bigger driving force than smaller grains, so that less applied stress is needed. Thus materials with smaller grain size exhibit higher yield stress. Hall-Petch relationship shows this inverse relationship between grain size and yield strength as,

= + 푘푦 휎푦 휎0 √푑 (2.6)

where is the yield stress, is a material constant for starting stress for dislocation movement,

푦 0 is the휎 strengthening coefficient,휎 and d is the average grain size.

푦 푘 Subgrains within the grain have a small misorientation between each other. It has been

shown that the higher density of the subgrains, the higher yield stress of the materials due to the

increase subgrain boundaries.

Carbides, as mentioned in the previous section, present along the grain boundaries act to

stabilize them and increase the resistance to dislocations across grains. They also act to restrict

grain boundary sliding. Creep resistance is increased thereby.

2.3 Background of Alloy 617

2.3.1 Chemistry and Microstructure of Alloy 617

Alloy 617, also designated as Inconel 617 or UNS N06617, was initially developed for

high temperature applications above 800°C. The combination of high strength and oxidation

resistance at high temperature makes alloy 617 an attractive material for use in aircraft and land

based gas turbines, chemical manufacturing components, metallurgical processing facilities, and

power generating plants [32]. The chemical composition of alloy 617 is listed in Table 2.2 as

follows.

Table 2.2 Limiting chemical composition (wt%) of alloy 617 [33].

Ni Cr Co Mo Fe Mn Al Min 44.5 20.0 10.0 8.0 - - 0.8 Max - 24.0 15.0 10.0 3.0 1.0 1.5 C Cu Si S Ti B Min 0.05 - - - - - Max 0.15 0.5 1.0 0.015 0.6 0.006

Although there are various alloying additions contained in Ni-based superalloys, the typical microstructure of these alloys only consists of γ matrix and γ’ intermetallic precipitates.

However, for Alloy 617, the observations by Mankins[34], Kimball [35], Kihara [36],

Kirchhofer [37], and Wu [38], showed that precipitates formed in the alloy at given temperature ranges have not been consistent, and other phases including CrMo(C, N), TiN, M12C, and

Ni2(MoCr) were also identified. These observations are summarized in Table 2.3. Wu [38]

studied the microstructure of long-term aged alloy 617 up to 65,000 hours and modified the TTT

diagram with identified phases on it, as shown in Figure 2.7. According to his observation, the

curve represents γ’ precipitates formation was extended below 550°C, and the γ’ completion

curve was constructed.

Table 2.3 Observation of Second Phase Precipitates in Alloy 617 [7].

Investigator M23C6 M6C γ’ Ti(C, N) Observation in material aged for 10,000 hours and less Mankins Small wt% persist to T = T≤1093°C Not observed Not observed Kimball 760°C Kihara 1000°C 1000°C Not reported Not reported Kirchhofer Observed Observed 550-1000°C 400-1000°C Observation in material aged for much longer than 10,000 hours at 482-871°C Observed at 482, 538, and Observed, 593°C, not at 704°C for Wu Observed also observed TiN observed 43,000h and longer, nor eta-MC 870°C after long time

Figure 2.7 TTT diagram for Alloy 617 long-term aged samples [38].

Efforts are still needed to investigate whether any of these phases influence the mechanical behavior under aging effects. Comprehensive reviews of the precipitates in Alloy

617 have been conducted by Ren and Swinderman [39,40], and Natesan et al. [41]. It is clear that the kinetics of the precipitation and coarsening processes are important in determining the effects

of aging on properties. Moreover, in the temperature range of 760 to about 1000°C, which is of

interest to the NGNP IHX application, the alloy may be stronger at initial exposure because of

precipitates, but most of the precipitates will be dissolved after long-term exposure and the alloy will depend on solid solution strengthening in the long run.

The grain size also plays an important role in the strength of the alloy. For general applications, a grain size of ASTM 6 (~45µm) or coarser is typically preferred, but it has been shown that creep strength increases with increasing grain size so microstructures of 100-200 µm grain size are often produced. A tradeoff exists, however, when fatigue is an issue, since finer grain sizes are preferred for fatigue resistance. In addition, for compact IHX, the thin sheet form restricts large grain size. Whether the grains will significantly coarsen after the dissolution of certain grain boundary precipitates at long-term exposure is not clear.

2.3.2 Mechanical Properties of Alloy 617

As 60 years is the designed service life for NGNP, such as strength, creep, and fatigue are crucial aspects in the selection of candidate materials. Compared with other candidate materials, alloy 617 is known for its good oxidation and corrosion resistance at temperatures up to 1093°C, and higher creep-rupture strength at temperatures from 649 to 1093°C, as shown in Figure 2.8. It also has good weldability and lower thermal expansion than most austenitic stainless steels and high thermal conductivity relative to other candidate materials. After being considered for the

VHTR, efforts have been made to study the mechanical properties of alloy 617 by research institutions all around the world. Some of the results will be reviewed in the following sections.

Figure 2.8 Creep rupture strength of high temperature alloys at 962°C in air [7].

2.3.2.1 Tensile Property

Tensile properties at high temperatures of solution-annealed, hot-rolled rods were

reported in Technical Bulletin [33] by Special Metals and are shown in Figure 2.9. The yield

strength decreases from room temperature up to about 400°C, roughly stays constant between

400-800°C, beyond which it decreases again. McCoy and King [42] from ORNL reported similar

trend in 1985. They found out that the tensile property varies with different heat, and fracture

strain decreased with aging. Recently, Burlet et al. [43] from the French VHTR program also studied the tensile property of Alloy 617. The variation of yield and ultimate tensile strengths with temperature up to 950°C of alloy 617 is given in Figure 2.10. Their data is consistent with

Special Metals. Kim et al. [44] compared high temperature tensile property of alloy 617 base and

weld metals, and their results show that base metal has lower yield stress but higher ultimate

tensile strength than weld metal below 800°C, however they are almost same at above 800°C, as

shown in Figure 2.11.

Figure 2.9 High temperature tensile properties of solution-annealed, hot rolled Alloy 617 rod [33].

(a) (b)

Figure 2.10 Tensile properties of as-received Alloy 617 and 230, (a) Yield Stress vs. Temperature and (b) UTS vs. Temperature [43].

(a)

(b)

Figure 2.11 High temperature tensile properties of the alloy 617 weldments, (a) 0.2% yield stress, (b) UTS [44].

2.3.2.2 Fatigue Strength

Several studies have been conducted to evaluate low-cycle fatigue, thermal-fatigue, creep-fatigue, and crack growth behavior on Alloy 617. Figure 2.12(a) shows the results of low- cycle fatigue in air environment by Meurer et al. [45] in 1984 and McGreevy et al. [46] in 2005, both of which are comparable. Low-cycle fatigue tests were also performed at different temperatures of 760, 871 and 982°C by Srivastava and Klarstrom [47], and results of which is shown in Figure 2.12(b). From the figures, it is evident that the fatigue life decreases at higher temperatures, and decreases dramatically up to about 104 cycles. At higher cyclic strain ranges, failure predominantly occurs by mechanical loading whereas at lower cyclic strain ranges, the effects of materials microstructure and environment on failure mechanisms and resulting fatigue life are significant [41].

(a) (b)

Figure 2.12 (a) Low-cycle fatigue results of Alloy 617 at 950°C in air [46]; (b) Low-cycle fatigue behavior of Alloy 617 at three different temperatures in air environment [47].

2.3.2.3 Creep-fatigue Property

Creep-fatigue damage is of concern for alloy 617 in VHTR applications, as the damage

arises on startup, shutdown, or power transient during operation. It was found that the allowable

number of cycles of the materials was significantly reduced when hold times at peak tensile

strain was introduced. Efforts have been put to obtain more creep-fatigue data, understand the

damage mechanism, and establish predicting models.

Lu et al. [48] pointed out that the combined creep-fatigue damage resulted in a much

shorter component life than the sum of creep and fatigue damages incurred separately. Both the

length of holding period and its position with respect to the cyclic strain wave are important

factors in creep-fatigue process. Holding only in the tension portion of the cycle led to extensive

damage and drastically reduced fatigue life of alloy 617, as shown in Figure 2.13. Fatigue life

showed a monotonic decrease with increasing holding time. Stress relaxation is usually to be

found during holding time, as shown in Figure 2.14, and results in the development of creep

strain and its accumulation in the crack-tip region. This would cause intergranular crack growth

and propagation, leading to shorter fatigue life.

Since both creep and fatigue component of the cycle cause damage, more complicated

material behavior is displayed. There are many theories try to explain and predict the creep- fatigue behavior, and linear damage summation is one of the most widely accepted. It assumes the damage from creep and fatigue is cumulative in a linear fashion, and is expressed as

+ = (2.7)

퐷푓 퐷푐 푡표푡푎푙 푑푎푚푎푔푒 = (2.8) 푁 퐷푓 푁푓 = = (2.9) ( ) ( ) Δ푡 푁 푡0 푑푡 퐷푐 Σ 푡푟 휎 ∫0 ∫0 푡푟 휎 푑푁0

Figure 2.13 The life reduction factor R against the cycle period for different load conditions. R is defined as the ratio of NF/NF0 where NF is the fatigue life for a given strain rate or wave shape - and hold time and NF0 is the reference life value for continuous cycling with a strain rate of 4x10 3s-1. Test temperature: 950°C [49].

(a) (b)

Figure 2.14 Stress relaxation for alloy 617 during 60 and 1800 seconds tensile hold in air at (a) 0.3% and (b) 1.0% strain [50].

where is fatigue damage, is creep damage, is cycles to failure in creep-fatigue test, is

푓 푐 푓 cycles 퐷to failure in low cycle퐷 fatigue test at same푁 test condition, is creep time, ( ) is creep푁

푟 rupture time at stress , is incremental time interval at instantaneousΔ푡 stress푡 ,휎 and is

푡 0 instantaneous cycle number.휎 푑 Chen et al. [51] compared creep and fatigue damage휎 endured푁 by

Alloy 617 and Haynes 230 based on linear damage summation, as shown in Figure 2.15. The black dot line outlines the failure criterion of linear damage summation, and the region below the

line is safe while the region above means material would fail. Here alloy 617 showed much less

damage at same creep-fatigue condition compared with Haynes 230.

Figure 2.15 Comparison of creep and fatigue damage endured by alloy 617 and Haynes 230 based on linear damage summation [51].

2.3.2.4 Creep Behavior

The components made from alloy 617 are usually subjected to high temperatures in

services, such as piping, liners of turbines, and heat exchangers. Therefore, the understanding of

creep behavior of the material is of great importance when predicting in-service behavior. 27

Research on creep and rupture properties of alloy 617 at high temperature have been brought into interest by scientists since 1980s. Ren [8] summarized the currently available creep data for alloy

617 from previous investigations. These data include times to 1%, 2%, and 5% total strains, time to tertiary creep, time to creep rupture, creep-rupture stress, minimum creep rate, loading strain, creep-rupture stress, minimum creep rate, loading strain, creep-rupture strain, and etc. First data set was produced by the manufacturer Huntington Alloys, Inc., which was obtained from 249 creep tests at temperatures from 593 to 1093°C, as summarized in Table 2.4. Oak Ridge National

Laboratory (ORNL) and General Electric Co. (GE) studied alloy 617 when it was considered for

HTGR programs. Their data set include 51 creep tests at temperatures from 593 to 871°C, as summarized in Table 2.5. Although these data sets are available, some raw data were missing or irretrievable due to different reasons. Scientists in Germany also investigated the alloy and produced data sets including 1947 tests at temperatures ranging from 500 to over 1000°C.

However, the access to the data set is restricted and creep test curves were not in fine resolution for further analysis. In order to thoroughly study the creep behavior of alloy 617 for NGNP applications, a systematic database is needed to be constructed with reasonable choice of test conditions, consistent treatment of specimens, well controlled test environments, and accurate record of data, as well as secured storage.

Table 2.4 Summary of the Huntington Alloy Inc. data set of alloy 617

Test temperature (°C) Base metal Weld metal 593 7 649 22 704 8 760 24 816 28 3 871 65 3 927 6 982 29 2 1000 16 1038 5 1093 31 Total 241 8

Table 2.5 Summary of ORNL data set of alloy 617

Test temperature (°C) Specimen Treatment Test environment 593 649 704 760 871 Base metal He aging He 3 2 1 Inert aging Air 2 He 1 1 1 Solution anneal Air 2 1 1 2 4 He 1 4 3 3 3 Weld metal As-welded He 3 3 2 1 1 He aging He 1 2

Alloy 617 generally displays exceptionally high levels of creep-rupture strength, even at temperatures of 980°C and above [33]. Figure 2.16 and Figure 2.17 show creep and rupture strength tested on bar, tubing, and sheet specimens by Special Metals. Creep strain vs. Time curves obtained under a series of stresses from the study by Schubert [52] are shown in Figure

2.18. The curves obtained from 850°C tests displayed classical three-stage creep behavior with an obvious secondary creep. However, curves for 950°C tests showed a different trend, a monotonically increasing creep rate from the start of the tests and virtually no secondary creep region, and the onset of tertiary creep is also not clear. Schubert [52] argued that the onset of tertiary creep as an indication of local instability is not applicable at high temperatures.

Figure 2.16 Creep strength of solution-annealed alloy 617 [33].

Figure 2.17 Rupture strength of solution-annealed alloy 617. Arrows denote tests discontinued before fracture [33].

Figure 2.18 Creep curves of alloy 617 at 850 and 950°C [52].

There are a number of factors that have great effects on the creep behavior of alloy 617, including aging, environment condition, grain size and grain boundary characteristics, and so on.

Kihara [36] found fine intragranular carbides precipitated on grain boundaries after only 1 hour aging at 1000°C, and they are effective in lowering creep rate during the early stage, as shown in

Figure 2.19. They also correlated the grain size to steady state creep rate and creep rupture life.

The results showed that steady state creep rate depends more on grain size than on the heat treatment, but heat treatment condition has greater effects on the creep rupture lives of small grain size samples, as shown in Figure 2.20.

Figure 2.19 The effect of fine intragranular carbides on the creep rate. The ratio of creep rates, / are plotted against the creep times for each pair of the same grain size samples. is the creep rate for the solution treated sample, and is that for aged sample [36]. 퐵 퐴 퐴 휖̇ 휖̇ 휖̇ 휖퐵̇

(a) (b) Figure 2.20 The grain size dependence of (a) the steady state creep rate and (b) the creep rupture life at 1000°C, 24.5MPa [36].

Creep data from studies by Baldwin [53], McCoy and King [42] show much different

creep rupture lives and time to 1% strain as in Figure 2.21, because the samples are from

different heats. From McCoy and King’s data, the grain size of Heat A sample is ≈130µm and

that of B is ≈210µm. Surprisingly the creep strength of Heat A is better than Heat B, which has not been explained. In this set of comparison, it is clear that creep behavior varies among

different heats, which indicates the importance of composition and microstructure control.

(a) (b) Figure 2.21 Heat-to-heat variation of creep behavior of Alloy 617 from past US HTGR programs; (a) stress vs. rupture life, (b) stress vs. time to 1% strain [42].

Research has shown that grain size and grain boundary characteristics have great impact

on crack propagation and void formation during creep of alloy 617 [54]. Larger grain size

provides lower chance for diffusion along grain boundaries, but higher chance of crack

propagation. Low angle boundary and special boundary have better resistance to creep damage

than random boundary, because with high atomic fit they are more effective in impeding crack

propagation. In addition, voids are more easily generated on random boundary rather than

coherent twin boundary, which have lower grain boundary energy [55]. Precipitation of

intergranular carbides is suppressed on special boundaries as compared to random boundaries,

and the applied stress plays an important role in precipitate evolution and distribution [56].

Therefore, creep resistance can be improved by manufacturing materials containing higher

percentage of preferred boundary type by means of heat treatment at proper condition.

Helium is the medium of heat transformation in VHTR, and it typically contains impurities such as H2, H2O, CH4, CO, and CO2 at different levels. First, it is clear that alloy 617

is adversely affected by helium environment as a significant decrease in creep life in helium

compared with that in air is observed (Figure 2.22). The reason is that in helium environment, aluminum will oxidize internally instead of forming a protective oxide scale. A loss of carbon is found in the samples. In addition, the impurity concentration also has impact on creep behavior.

Hosoi and Abe [58] investigated the effect of helium environment on creep properties of alloy

617. It is clearly illustrated in Figure 2.23(a) that alloy 617 creeps faster as oxygen concentrations in helium increase, though the creep rates decrease in a similar fashion. Oxygen concentration is also closely related to creep rupture time and carbon content, which in turn affects the creep property because of carbide precipitation as mentioned above.

Figure 2.22 Variation in creep rupture time as a function of applied stress for alloy 617 at 950°C in air and in helium of low steam content [57].

(a) (b) Figure 2.23 Effect of oxygen content on (a) creep strain-time curves and (b) creep rupture time and decarburization of alloy 617 in helium [58].

The IHX components are typically expected to experience relatively low stresses (5-

20MPa) and high temperatures (750-950°C), but may also encounter higher stresses during service excursions. Thus, both diffusional (Nabarro-Herring, Coble) and dislocation creep processes are expected to be quite important. The minimum creep rates for Nabarro-Herring

(NH), Coble (C) and dislocation (D) creep mechanisms are given by:

o DL σ ⋅ Ω ε NH = ANH ⋅ 2 ⋅  d kT (2.10)

o DGB ⋅δ ′ σ ⋅ Ω ε C = AC ⋅ 3 ⋅  d kT (2.11)

n o  σ    QD ε D = AD ⋅ DL ⋅  DL = Do exp −    G  RT  (2.12)

where σ is stress, Ω is atomic volume, d is the grain size, δ’ is the grain boundary width, k is the

Boltzman constant, T is the absolute temperature and DL and DGB are the lattice and grain

boundary diffusivity, respectively. The pre-exponential constants ANH and AC are typically 10-14

and 50, respectively, whereas AD can vary a lot depending on the alloy system and stacking fault

energy. However, more accurate values for these constants have to be determined through

experiment. The NH and Coble mechanisms show a similar linear stress dependence of the creep

rate, but differ in the activation energy and grain size dependence, with diffusion through the

lattice controlling NH creep and that along grain boundaries controlling Coble creep. Grain size

has a larger effect on Coble creep (1/d3) compared with NH creep (1/d2). Grain boundary sliding

is also an important mechanism that operates concomitantly during NH and Coble creep and the

difficulty of this process can lead to early void formation at grain boundaries and contribute to an

increase in the creep rates past the minimum creep rate, i.e. in the tertiary creep stage. The stress

exponents for dislocation creep are usually higher (4-8) and lattice diffusion is expected to

control the creep rates.

Efforts have been directed towards modeling and life prediction of the creep behavior of

alloy 617, including the well known parameter-based approaches, such as Larson-Miller, Sherby-

Dorn, Manson-Haferd, etc. [59]. Various models developed up to date for nickel alloys were

recently reviewed by Swindeman [60], including the Norton [61] secondary creep equation:

= 푄 푛 푒̇ 푎1푒푥푝 �− � 휎 푡 푅푇 (2.13)

the Garofalo [62] transient-secondary creep equation:

= [1 ( )] + (2.14)

푡 1 푚푖푛 the theta [63] transient tertiary creep푒 푒 equation:− 푒푥푝 −푏 푡 푒̇ 푡

= [1 ( )] + [ ( 1)] (2.15)

1 2 3 4 and the Dyson and McLean푒 [64]휃 constitutive− 푒푥푝 −휃 model:푡 휃 푒푥푝 휃 푡 −

(1 ) = (1 + ) × sinh 1 (1 ) 0 푑 푄 휎 − 퐻 푒̇ 푒̇ 퐷 푒푥푝 �− � � 0 푝 � 푅푇 휎 � − 퐷 � − 휔 (2.16)

where is creep strain, is creep rate, is time, is stress, is absolute temperature, is activation푒 energy, is the푒 ̇ gas constant, 푡is the stress휎 exponent, 푇the s are constants in the theta푄 equation, , , and푅 are damage parameters푛 in the DM model whose휃 values range from 0 to

푑 푝 1, is a hardening퐷 퐷 parameter,휔 is a rate constant, and is a constant. The DM model may be

1 1 considered퐻 as a strong candidate푏 in unified models which푎 represent both plasticity and creep.

To describe the plastic displacements and strains, they usually apply some phenomenological tensor form of the Norton law relating the plastic strain (or strain rate) to the power function of the local stress as in equation (3), including as necessary an internal or threshold stress [65]. Viscoplastic constitutive models [66] extend these approaches and there is work being conducted by ORNL to develop a unified viscoplastic constitutive model for GEN-

IV alloys to cover the high temperature, low stress regime [67]. Preliminary numerical and simulation results show that the developed approach is promising. However, experiments and determination of material constants are needed. Continuum creep damage mechanics (CDM) models that have also been developed employ internal state variables to quantify the strain rate response to an applied stress and/or stressing rate, up to and including fracture [68]. These various CDM-based modeling approaches do not explicitly take into account microstructural features like lattice grains, sub-grains and dislocations. Consequently, these models do not provide mechanistic insight into microstructural processes that control creep. On the other hand,

physically-based CDM models [69] have been developed to fill this gap by explicitly accounting

37 for microstructural mechanisms of creep and degradation in the form of damage parameters and kinetics of damage evolution into the unified creep strain rate equations, to predict creep curves and failures in structures accumulating large strains. The physically based CDM approach has been argued to provide a unifying framework for some of the seemingly diverse methods of predicting creep lifetimes, such as the Larson-Miller parameter.

3 Experimental Details

3.1 Specimens

Two sets of alloy 617 test specimens with different grain sizes were received from Idaho

National Laboratory. One set of specimens were machined from 1.5-inch-thick alloy 617 plates

that were hot rolled and solution annealed at 1150°C for 2 hours, with an average grain size of

~95µm (ASTM 3.5). The other set as machined from hot-rolled 0.5 inch thick plate with an

average grain size of ~180µm (ASTM 2). A third set of specimens with an average grain size of

~20µm (ASTM 8) were manufactured from an alloy 617 plate provided by INL, by first cold

rolling the plate to 50% thickness, and then annealing at 1000°C for 15 hours. All specimens

were machined with the long axis along the rolling direction of the raw plates.

All test specimens conform to ASTM Standard E8/8M-09, and the design is shown in

Figure 3.1. The diameter of the gage section is ~6.35mm, and the effective gage length is

~34mm. Detailed dimensions are shown in the attached drawing.

Figure 3.1 Schematic drawing shows geometry of creep test specimen.

3.2 Creep Testing

3.2.1 Creep Testing Equipment

Creep tests were mostly performed on Lever Arm Creep/Stress Rupture Testing Systems

Model 2330 manufactured by Applied Test Systems Inc., as shown in Figure 3.2, and installed in the lab of University of Cincinnati; some of the tests were performed on the same systems at

Idaho National Laboratory, and other creep data also obtained from tests conducted by Metcut

Research Inc. in Cincinnati, Ohio. Lever arm creep testing system uses sharp knife edge as the pivot, and specimen is gripped by pull rods on one side and dead weights providing constant force are hung on the other side. In order to ensure the sensitivity and accuracy of applied stress,

high stress tests (≥50MPa) were performed on two creep testing machines with a lever arm ratio of 20:1, and another creep testing machine with a lever arm ratio of 3:1 was used for low stress

tests (<50MPa).

(a) (b) Figure 3.2 Lever Arm Creep/Stress Rupture Testing Systems with lever arm ratio of (a) 3:1 and (b) 20:1.

A three-zone furnace with a capability of 1100°C is equipped on the creep testing systems. The temperature was monitored by two type K thermocouples, which were separately spot-welded to the necking part of gage section on top and bottom of the specimens. It was precisely controlled to within ±1.60°C of the test temperatures. The heating process was programmed as first heat to target temperature at a ramp rate of 950°C per hour, and a soak time of 2 hours was applied in which the specimen was held at test temperature with 0 load in order to ensure a stabilized and uniformed temperature. The dead weights were then automatically loaded after the soak time finished. A typical temperature curve is shown in Figure 3.3.

Figure 3.3 Temperature curve of creep test.

The creep displacement was measured by dual channel digital displacement LVDT transducers with an accuracy of ±1.0µm attached on the averaging extensometer. Using two transducers to measure the displacement on both sides of the specimen and then averaging the

values could minimize any error caused by bending. Data of time, temperature, displacement,

and creep were recorded 5 times per minute automatically by using Windows Computer Creep

System (WinCCS) software contained in the PC interfaced to the creep system.

Most tests were terminated automatically at 20% creep strain of specimen by the end-of-

test criterion. However, some of low stress tests were terminated after the minimum creep rate

was reached depending on the length of the tests.

For the purpose of yielding high quality results, major components of the test systems,

such as the load cell, extensometer, and LVDT sensors, were all calibrated to satisfy the relevant

ASTM standards. The alignment of the frames was also checked with strain-gauged specimens and the alignment fixture was adjusted to insure the minimum bending.

3.2.2 Creep Testing Condition

Since this research project aims to systematically study the creep behavior of alloy 617, a test matrix covering both high and low stress regimes were proposed. Meanwhile, more attention was given to the low stress conditions, since IHX for NGNP application would normally undergo high temperature low stress conditions in service at most times, and similar condition for creep is desirable. Furthermore, the test matrix is designed to generate enough data for modeling which may serve for a reasonable prediction of long term performance of alloy 617. These data also contribute to the database for ASME code case development.

In summary, creep tests were performed at a series of conditions with temperatures in the range of 850 to 1050°C and stresses in the range of 5 to 100MPa. The test condition for each specimen by different grain sizes are listed in Table 3.1, Table 3.2, and Table 3.3.

Table 3.1 Creep test conditions for alloy 617 with grain size of ASTM 3.5

Specimen Temperature (°C) Stress (MPa) 617-22 850 20 617-14 850 30 617-D62 850 50 617-8 850 75 617-D64 850 100 61735-1 950 5 617-20 950 10 617-21 950 20 617-12 950 30 61735-3 950 40 617-D61 950 50 617-5 950 75 617-D63 950 100 617-26 1000 10 617-D16 1000 20 617-D17 1000 50 61735-2 1050 5 617-19 1050 10 617-17 1050 20 617-13 1050 30 617-7 1050 50 617-11 1050 75

Table 3.2 Creep test conditions for Alloy 617 with grain size of ASTM 2

Specimen Temperature (°C) Stress (MPa) HR617-12 850 50 HR617-11 850 75 HR617-5 850 100 6172-1 950 5 HR617-2 950 30 6172-3 950 40 HR617-4 950 50 HR617-10 950 75 HR617-6 950 100 6172-2 1050 5 6172-4 1050 30 HR617-8 1050 50 HR617-9 1050 75

Table 3.3 Creep test conditions for Alloy 617 with grain size of ASTM 8

Specimen Temperature (°C) Stress (MPa) 6178-1-INL 950 5 6178-3-INL 950 10 6178-11 950 15 6178-1 950 20 6178-2-INL 1050 5 6178-4-INL 1050 10 6178-12 1050 15 6178-3 1050 20

3.3 Microstructure Characterization

In order to obtain insights into the governing creep mechanism of alloy 617, it is important to gain a firm understanding of the microstructure of the material and its change from before to after creep testing. Detailed information of interest includes grain size and distribution, character distribution of grain boundaries, presence, distribution and chemistry of carbides, dislocation characters, etc. Therefore, a number of modern techniques were utilized to characterize the microstructure of the materials before and after creep testing, including Optical

Microscopy (OM), Scanning Electron Microscopy (SEM), Electron Back Scattered Diffraction

(EBSD) and Orientation Imaging Microscopy (OIM), and Transmission Electron Microscopy

(TEM).

Metallographic study and average grain size measurements were carried out using a

Keyence Digital Microscope VHX-600, equipped with a 54 million pixel 3CCD Camera with

1000x magnification capability. Sample preparation procedure includes first sectioning the sample with low speed precision saw, followed by hot or cold mounting the sample in an epoxy, and then polishing using successively finer silicon carbide paper, and last diamond solution for a mirror finish. To reveal grain structure under microscope, samples were immersed into the

etchant composed of 100mL hydrochloric acid and 0.5mL hydrogen peroxide for 10 to 15

seconds and then rinsed by water.

An FEI XL30 Environmental SEM with secondary electron (SE) and backscattered

electron (BSE) imaging was employed to provide high resolution images. Energy Dispersive X-

ray Spectroscopy (EDS) technique was utilized for chemistry analysis and phase identification.

EDS analysis was performed by using EDAX Genesis 4000 thin-window system attached to FEI

XL30 SEM.

Quantitative information of grain size, grain size distribution, texture, grain boundary character distribution was obtained by utilizing Electron Backscatter Diffraction (EBSD)

technique. Kikuchi bands are formed by backscattered electrons which satisfy Bragg’s condition

related to the spacing of lattice planes. The width of Kikuchi band is related to interplanar spacing and the angle between the bands is related to the angle between diffracting planes. Thus,

a unique solution to the crystal orientation can be obtained based on the indexing of each band.

Orientation Imaging Microscopy (OIM) map is thereby constructed by scanning the electron

beam to obtain crystallographic information point by point. EBSD requires a highly tilted angle

(~70°C) from horizontal of the sample to enhance the scattered electron and increase the contrast

in Kikuchi pattern. FEI XL30 SEM with an EDAX/TSL 4000 EBSD system controlled by TSL

OIM Data Collection software was employed to collect EBSD data, and TSL OIM Data Analysis

software was used to analyze collected data. Samples from both before and after creep tests

materials were cut in cross-sectional and longitudinal directions by linear precision saw with a

silicon carbide abrasive cut-off blade, followed by hot mounting into conductive epoxy, and then progressively polished with silicon carbide abrasive paper with grit sizes of 320, 640, 800, and

1200, in sequence. Diluted alumina paste with particle sizes of 1 and 0.3µm were used for finer

polishing. Colloidal silica solution was applied on a vibratory polisher as the final step to obtain a very smooth sample surface. The size of EBSD scan was 600 by 600µm, and the step size of

2µm was used.

An FEI CM20 TEM equipped with Gatan CCD based image acquisition system was also utilized for microstructure characterization and dislocation analysis. An accelerating voltage of

200kV was used. Cross-sectional samples were first cut off from the gage section of creep tested specimens by Electrical Discharge Machine (EDM) with a thickness of ~100-120µm, and then they were ground to a thickness of ~50-70µm. Thin foils with 3mm diameter were then punched out from the ground samples. Small holes which give thin areas that can be transmitted by electron beam were created in the center of the thin foils by Fischione Model 120 Twin Jet

Electro-polisher. The electrolytic solution was composed of 10% perchloric acid and 90% methanol. The solution was maintained at the temperature of -30°C during polishing.

4 Results and Discussion

4.1 Creep Test Data Analysis

The recorded data from creep tests include time, temperature, stress, displacement, and creep. Plots are constructed based on these data in order to reveal creep behavior of alloy 617 at different conditions and the underlying creep mechanism. Parameters of the deformation model are also calculated based on the analysis of the obtained data. Comparison is also made with historical data. It is noted that the creep behavior not only depends on temperature and stress, but grain size also has a great influence. The detailed data analysis is presented in the following sections.

4.1.1 Specimens with Grain Size of ASTM 3.5

Specimens with an average grain size of ASTM 3.5 were machined from the original plated received from INL. They are the set which underwent most tests, therefore most data was obtained. Since ASTM 3.5 is considered as standard grain size for alloy 617, the data for these specimens is set as a benchmark.

Creep curves for every test condition are plotted as creep strain vs. time and smoothed by averaging method, as shown in Figure 4.1. They are presented in four plots by test temperature as 850, 950, 1000, and 1050°C, which show a clear dependence on the stress level. Creep curves grouped by stress level are also plotted in order to show the temperature dependence of creep. A typical representation at 20MPa for different temperatures is shown in Figure 4.2. Creep process happens much faster at temperatures greater than 850°C, but the difference becomes smaller as temperature increases. Creep curves representing classical creep behavior of metals are usually composed of three stages as primary, secondary, and tertiary stage. These creep curves of alloy

617, however, display a very different trend with little primary creep stage, and a well-defined secondary creep stage is not recognizable as well. The onset of tertiary creep stage is unclear but initiated early. Similar findings were also reported by previous studies [52,70]. A comparison to

Cr-Mo steels and 316 stainless steel, which shows a well-defined secondary creep stage, has been made as shown in Figure 4.3. A steady state creep is the result of a dynamic balance between the competing process of strain hardening and recovery. A much more complicated process is applicable here for creep in alloy 617. However, the minimum creep rate can still be determined by more detailed analysis.

(a)

(b)

(c)

(d) Figure 4.1 Creep strain vs. time plots for alloy 617 specimens tested at (a) 850°C, (b) 950°C, (c) 1000°C, and (d) 1050°C for various stresses.

Figure 4.2 Creep strain vs. time plot for alloy 617 specimens tested at 20MPa for various temperatures.

Figure 4.3 Normalized creep curves for typical high-temperature materials [70].

The initiation of tertiary creep stage is an important criterion for creep property, because

it is an indication of localized necking with increased stress, leading to local instability and thus

material failure. In order to determine tertiary creep stage, creep rate is plotted against time and

creep strain as shown in Figure 4.4. The curves are presented individually to display their trend clearly. Creep rate vs. creep strain is plotted in order to reveal the creep rate changes with respect to creep development stages and compare among various conditions. Generally, the curves can be categorized into two types by their trends. For the first type (Type I), the creep rate increases rapidly at the beginning of the test, followed by a decrease for a short period of time to reach minimum, and then increase monotonically until failure. For the second type (Type II), the creep rate drops quickly at the beginning of the test, followed by an increase for a short period of time, then gradually decreases to reach minimum, and finally increases monotonically until failure.

The creep rate curve type for every test condition is summarized in Table 4.1, and comparison

among different stresses at same temperature conditions is illustrated in Figure 4.5. It is easy to

tell that the type of creep rate curve is only determined by stress level, while temperature has no

impact. All the high stress tests that have stress levels greater than 50MPa, specifically 75 and

100MPa, exhibit Type I creep rate curve; and all the low stress tests which have stress levels

equal or lower than 50MPa (5 to 50MPa) exhibit Type II creep rate curve. This observation indicates the effects of stress on creep behavior of alloy 617. Since there is no steady state creep stage in alloy 617, we may define the point when minimum creep rate is reached as the initiation of tertiary creep stage. For high stress condition, tertiary creep stage starts at about 2% creep strain, while for low stress condition, tertiary creep begins much later at about 10% creep strain.

(a)

(b)

(c)

(d) Figure 4.4 Creep rate vs. time and creep strain plots for alloy 617 specimens tested at (a) 850°C 50MPa, (b) 850°C 100MPa, (c) 950°C 50MPa, and (d) 950°C 100MPa.

(a) 54

(b)

(c) Figure 4.5 Creep rate vs. creep strain plots for alloy 617 specimens tested at (a) 850°C, (b) 950°C, and (c) 1050°C for various stress levels.

Table 4.1 Creep rate curve type summary.

Specimen Temperature (°C) Stress (MPa) Creep Rate Curve Type 617-22 850 20 II 617-14 850 30 II 617-D62 850 50 II 617-8 850 75 I 617-D64 850 100 I 617-20 950 10 II 617-21 950 20 II 617-12 950 30 II 617-D61 950 50 II 617-5 950 75 I 617-D63 950 100 I 617-26 1000 10 II 617-D16 1000 20 II 617-D17 1000 50 I 617-19 1050 10 II 617-17 1050 20 II 617-13 1050 30 II 617-7 1050 50 II 617-11 1050 75 I

The values of minimum creep rate for each test were determined from the creep rate vs. time and creep rate vs. creep strain plots, and used for creep mechanism determination and

parameter calculation. Figure 4.6 shows the plot of minimum creep rate vs. stress in log-log

scales. A linear relationship is revealed by regression fit, and the slopes of fitting lines are

calculated in the range of 4.8-5.7 This suggests a power law dislocation creep mechanism is

dominant in the creep process, where the following equations hold,

= 푛 휎 휀퐷̇ 퐴퐷 ∙ 퐷퐿 ∙ � � 퐺 (4.1)

and

= 푄퐷 퐷퐿 퐷0푒푥푝 �− � 푅푇 (4.2) 56

where is creep rate, A is a constant, D is diffusion coefficient, σ is applied stress, n is stress

exponent,휀̇ Q is activation energy, R is universal gas constant, T is absolute temperature. The calculated stress exponent values are in good agreement with the values reported by Kim et al.

[70], and Shankar and Natesan [71].

However, the stress exponent depends on a number of factors and could vary a lot at different conditions for specimens with different metallurgical history. Figure 4.7 is generated from historical data provided by ORNL showing the stress exponent variation and a comparison with our results. Dark lines are the fitting lines for historical data and red lines are for our data.

The stress exponent values lie in a pretty wide range of 4.3 to 15.6, as summarized in Table 4.2.

It can be seen from the figure that the fitting for historical data is not as good as our data, and no

clear linear relationship is even presented at some test conditions. More importantly, the stress

exponent values are not consistent; neither does it show a clear trend depending on temperature.

The possible reason is that the data was obtained from samples from different supplier, different

heat, and even different product form such as bar, tube, and sheet.

Table 4.2 Stress exponent for alloy 617 historical data

Temperature (°C) Stress exponent 593 15.53 649 12.06 704 4.39 760 10.32 816 4.99 871 6.48 982 6.51 1038 8.52 1093 8.63

Figure 4.6 Plot of minimum creep rate vs. stress for determining stress exponent n of alloy 617 ASTM 3.5 specimens.

Figure 4.7 Minimum creep rate vs. stress plot for alloy 617 specimens compared with historical data.

The values of activation energy Q can be calculated from the plot of minimum creep rate vs. inverse of temperature in a log-log scale, as shown in Figure 4.8. By measuring the slopes of fitted lines, activation energy for different test conditions are calculated as 420, 394, 443, and

389kJ/mol for stresses at 20, 30, 50, and 75MPa, respectively. For stress at 5 and 100MPa, the activation energy are calculated as 319 and 379 kJ/mol by only two data points determining each fitting line. Therefore, these two values are just for reference, more data points are needed for a reliable estimation. Although the values are stress dependents, they are close to the value of

408kJ/mol observed in pure helium environment test [71].

The activation energy is expected to be close to that for self-diffusion of Ni of 275kJ/mol in theory for dislocation creep. However, many observations of higher values of activation energy have been reported on solid solution strengthened alloys, and a few theories have been proposed to try to explain the phenomenon. Tien et al. [72] proposed that vacancy sources may be exhausted in very fine dislocation network produced at the γ/γ’ interface, and thus creep can only occur by processes involving formation and migration of interstitial ions, therefore, the activation energy for this process could be much higher than that for self-diffusion. Leverant and

Kear [73] suggested that the self-diffusion of γ’ precipitates may be involved in the process, and these precipitates have considerably higher activation energy comparing with that of Ni, resulting in a much higher activation energy overall.

Figure 4.8 Plot of minimum creep rate vs. 1/T for determining the activation energy Q.

The obtained data and analysis can be represented in the form of the Mukherjee-Bird-

Dorn equation [74], relating minimum creep rate as a function of applied stress, temperature, and grain size, as following:

= 푛 푝 퐴퐺푏 푄퐶 휎 푑 휀̇ 퐷0푒푥푝 �− � � � � � 푘푇 푅푇 퐺 푏 (4.3) where is minimum creep rate, is a dimensionless constant, is shear modulus, is Burgers vector,휀 ̇ is Boltzmann’s constant,퐴 is absolute temperature, 퐺 is the self-diffusion푏 coefficient,

0 is activation푘 energy for creep, usually푇 considered being equal퐷 to that for self-diffusion, is

퐶 푄universal gas constant, is applied stress, is stress exponent, is grain size, and is 푅the inverse grain-size exponent휎 , equals to 0 for power푛 law dislocation 푑creep. A unified relationship푝

can therefore be established by plotting normalized minimum creep rate as ( / ) against

normalized stress as ( / ) for all conditions as shown in Figure 4.9. Here 휀̇푘푇 = 퐷405kJ/mol퐺푏 is

퐶 obtained by averaging휎 activation퐺 energy values for all stress levels calculated푄 in the previous

section. Data points for different temperatures converge and follow a good linearity, and the

stress exponent by the slope of the fitting line is 4.96, indicating a typical power law dislocation

creep is the dominant mechanism for the range of temperature and stress conditions.

Figure 4.9 Normalized minimum creep rate vs. normalized stress for alloy 617 ASTM 3.5 specimens.

Due to the lengthy test, sometimes it is impossible to obtain the results from tests at conditions close to operation. Hence, a few extrapolation methods have been developed to predict the performance of materials based on data obtained from tests conducted over a relatively short period of time, such as Larson-Miller [75], Manson-Haferd [76], and Sherby-

Dorn [77] methods. Larson-Miller method is used here to predict the long term performance of alloy 617, the equation of which is given as follows.

(log + ) = (4.4)

푟 where T is temperature, tr is the time to푇 rupture푡 at 퐶a constant푚 stress, C is a constant depending

on the alloy, usually being close to 20, m is a parameter depending on stress and rupture time,

also known as Larson-Miller parameter.

In this study, we use Larson-Miller method to predict the time for alloy 617 to reach 1%

and 20% creep strain at high temperature and low stress conditions. First, constant C is

calculated by plotting test data of time to reach 1% and 20% creep strain vs. inverse of

temperature. Fitting lines for different stresses will converge and intercept to one point on the y

axis for time, and the intercept value is C for 1% and 20% creep, respectively. C for 1% creep of

alloy 617 specimens is calculated as 16.7 and C for 20% creep is calculated as 16.1. Then the

Larson-Miller parameter m, can be calculated by equation 4.4. Consequently, a linear

relationship can be obtained by plotting stress in log scale vs. Larson-Miller parameter m, as shown in Figure 4.10. The Larson-Miller parameter m at lower stress levels can be determined by extrapolating the fit lines for test data points. Therefore, the time t to both 1% and 20% creep of alloy 617 at different creep conditions can be reverse calculated according to equation 4.4.

Table 4.3 lists the time to 1% and 20% creep at temperature of 850, 950, and 1050°C, and stress of 5, 10, and 15MPa, which is a good reference for the long term performance of the alloy in service condition. As can be seen from the data, the creep strain varies dramatically with respect to both temperatures and stress levels. Only considering creep damage, alloy 617 with grain size of ASTM 3.5 will definitely last as long as the life of the plant at 850°C and stress lower than

5MPa, may be able to meet the requirement of 60 years service life at 850°C and stress lower

62 than 15MPa, and 950°C stress lower than 5MPa depending on the criteria of failure in creep strain, and cannot meet the requirement at 1050°C. Schubert et al. [52] reported similar prediction for alloy 617 to reach 1% and rupture as shown in Figure 4.11, and he also compared

LMP with other extrapolation methods and proved its validity.

Figure 4.10 Plot for Larson-Miller parameter relationship and projected long term performance of alloy 617.

Table 4.3 Predicted time for alloy 617 to reach 1% and 20% creep.

Temperature (°C) Stress (MPa) Time to 1% Creep Time to 20% Creep 850 5 1849 years 139899 years 850 10 38 years 2155 years 850 15 4 years 188 years 950 5 21 years 1219 years 950 10 7 months 26 years 950 15 27 days 2.8 years 1050 5 5.4 months 22 years 1050 10 6 days 7.6 months 1050 15 21 hours 29 days

(a) (b) Figure 4.11 Stress for (a) 1% creep strain and (b) rupture vs. Larson-Miller parameter, for alloy 617 at temperatures between 800 to 1000°C [52].

4.1.2 Specimens with Grain Size of ASTM 2

Creep tests at the same conditions were also performed on the hot rolled Alloy 617

samples with a larger grain size of ASTM 2 (~180µm). Similar analysis on this set of specimens is conducted and presented as follows.

Creep curve for each test condition is generated by plotting creep strain vs. time, as shown in Figure 4.12. It is obvious from the plots that ASTM 2 specimens display a classical three-stage creep behavior with well-recognized primary, secondary steady-state, and tertiary creep at 850 and 950°C. It is a big departure compared with what ASTM 3.5 specimens show.

This observation means the dynamic equilibrium between work hardening and relaxation is reached in secondary creep. The possible explanation is that the annealing process not only provides a good condition for grain growth, but also causes precipitates ripening and carbides redistribution, which, in turn, provide impedance to dislocation movement. In addition, the larger

64 grain size means less grain boundaries, and less chance of grain boundary sliding would occur.

When the temperature reaches 1050°C, the classical creep behavior is replaced by what is seen in the ASTM 3.5 grain size specimens, namely no steady-state secondary creep occurs.

Creep rate vs. creep strain curves for ASTM 2 specimens at three temperatures are illustrated in Figure 4.13. The creep rates experience a suddenly increase to a very high level after a very slow increase. Interestingly, the corresponding creep strain to the sudden increase does not decrease with the stress increase.

(a)

(b)

(c)

Figure 4.12 Plots of creep strain vs. time for alloy 617 ASTM 2 specimens at (a) 850°C, (b) 950°C, and (c) 1050°C.

(a)

(b)

(c)

Figure 4.13 Creep rate vs. creep strain plots for alloy 617 specimens with grain size of ASTM 2 tested at (a) 850°C, (b) 950°C, and (c) 1050°C for various stress levels.

By comparing the creep behavior between two sets of alloy 617 samples at same test conditions, it was found that grain size has a significant effect on creep behavior, as shown in

Figure 4.14. Specimens with larger grain size show much lower creep rates and extended period of time for minimum creep rate at equivalent temperature of 850 and 950°C and stress conditions.

But after entering the tertiary creep regime, where creep rate suddenly increases in a very short period of time, the creep strain for ASTM 2 specimens exceeds ASTM 3.5 specimens. In addition, at 1050°C, ASTM 2 specimens creep faster than ASTM 3.5 specimens from beginning or after a short period of time. This suggests the large grain size in alloy 617 may not be favorable for long-term creep.

(a)

(b)

(c) Figure 4.14 Creep strain vs. time plots for alloy 617 specimens with ASTM 3.5 and 2 grain sizes tested at (a) 850°C, (b) 950°C, and (c) 1050°C.

The stress exponents n determined by plotting minimum creep rate vs. stress in log-log scale were 2.28 and 3.33 for 850 and 950°C respectively, as shown in Figure 4.15. Shankar and

Natesan [71] reported a stress exponent of 2.38 for alloy 617 tested in He + O2 environment, and suggested a viscous glide mechanism for creep. A stress exponent of 4.86 was found in 1050°C for ASTM 2 grain size specimens. It is close to 5, same as ASTM 3.5 grain size specimens, meaning the power law dislocation creep mechanism is still in dominant at high temperature.

Figure 4.15 Plot of minimum creep rate vs. stress for determining stress exponent n of ASTM 2 specimens.

4.1.3 Specimens with Grain Size of ASTM 8

Creep tests on specimens with ASTM 8 grain sizes were conducted at 950 and 1050°C at low stress level regime from 5 to 20MPa. Similar analysis on obtained data were performed and

70 plotted as shown in Figure 4.16. ASTM 8 grain size specimens exhibit same creep behavior as

ASTM 3.5 specimens.

Stress exponent was calculated by plotting minimum creep rate vs. stress, as shown in

Figure 4.17. Very good linearity was obtained, and the stress exponents were determined as 3.27 for 950°C and 3.23 for 1050°C. The lower values of stress exponents may indicate a mixed creep mechanism with both dislocation creep and diffusional creep involved.

Creep curves of ASTM 8 specimens were compared with those of ASTM 3.5 specimens, as shown in Figure 4.18. Although both sets of specimen showed similar trend in creep curves,

ASTM 3.5 specimens clearly have lower creep rate at same conditions, thus leading to better creep property.

(a)

(b)

Figure 4.16 Plots of creep strain vs. time for alloy 617 ASTM 8 specimens at (a) 950°C, and (b) 1050°C.

Figure 4.17 Plot of minimum creep rate vs. stress for determining stress exponent n of ASTM 8 specimens.

(a)

(b)

Figure 4.18 Creep strain vs. time plots for alloy 617 specimens with ASTM 3.5 and 2 grain sizes tested at (a) 950°C and (b) 1050°C.

The comparison was made among all three sets of the specimens at low stress conditions, as shown in Figure 4.19. The set of specimens with ASTM 3.5 grain size have the best creep strength at both 950 and 1050°C 5MPa conditions, the specimens with smallest grain size of

ASTM 8 have the worst performance, and the largest ASTM 2 grain size specimens are in between. Therefore, small grain size is not favorable for high temperature creep strength.

(a)

(b) Figure 4.19 Comparison among three sets of specimens of alloy 617 at (a) 950°C 5MPa and (b) 1050°C 5MPa 74

4.2 Microstructure Characterization

4.2.1 Average Grain Size Determination by Optical Microscopy

Optical microscopy was utilized for determining grain size for different sets of samples.

Specimens were cut from longitudinal, cross-section, and transverse directions, then mounted, polished, and etched to reveal grain boundaries. Optical images were then recorded from each

specimen. Average grain size was determined by circular intercept Hilliard Sing-Circle

procedure per ASTM E112-96, where a circle with certain circumference length was overlaid on

the optical image, and the number of grain boundaries it intercepted with was counted and

recorded, then the grain size was calculated by the following equations:

= / 푁푖 푁�퐿 퐿 푀 (4.5)

1 = ̅ 푙 퐿 푁� (4.6)

where is mean intercept count per unit length, is the number of intercepts counted in each

퐿 푖 image, 푁�L is the total test line length, M is the magnification,푁 is the mean lineal intercept value.

Figure 4.20 shows the images taken from rolling direction,푙̅ cross-sectional direction and

transverse direction of the as-received alloy 617 material, and each of which was overlaid by a

circle with a circumference length of 5mm in blue, and the intercepts of the circle and grain

boundaries were marked in red cross. Rolling direction was also noted as RD with an arrow in

the images. The number of intercepts were counted, recorded, and calculated by the equations

4.5 and 4.6, as listed in Table 4.4. As a result, the average grain size can be calculated as 94.4µm

which is 3.5 in ASTM micro-grain size number (ASTM 3.5).

Same method was applied to the material that underwent hot rolling and annealing. The optical images are shown in Figure 4.21, and the average grain size was determined as 181µm

(ASTM 2), as the data listed in Table 4.5.

In addition, EBSD was also used to help determining average grain size. The details are presented in the following section.

Table 4.4 Average grain size determination of as-received alloy 617 material.

Rolling Transverse Cross-sectional Direction Direction Direction Direction N (# of intercepts) 54 54 51 L (circumference, mm) 5 5 5 NL (# of intercepts per unit length) 10.8 10.8 10.2

(a) (b)

(c)

Figure 4.20 Optical images taken from (a) rolling direction, (b) transverse direction, and (c) cross-sectional direction of as-received alloy 617 material showing average grain size determination by circular intercept method.

Table 4.5 Average grain size determination of hot-rolled alloy 617 material.

Rolling Transverse Cross-sectional Direction Direction Direction Direction N (# of intercepts) 26 28 29 L (circumference, mm) 5 5 5 NL (# of intercepts per unit length) 5.2 5.6 5.8

(a) (b)

(c)

4.2.2 SEM Analysis

Microstructure analysis of as-received alloy 617 material was performed on the cross- section specimens by using SEM and EDS. A secondary electron (SE) image was obtained with a magnification of 500×, the accelerating voltage of 15kV, and the working distance of 10 mm, as shown in Figure 4.22(a). It can be seen from the image that it generally consists of matrix in grey and some inclusions shown in dark color. With the help of EDS, chemical composition of these two phases can be identified as Ni-Cr-X matrix and Ti(C,N), as shown in Figure 4.22(b) and (c).

(a)

(b)

(c)

Figure 4.22 Phase composition in as received alloy 617 by (a) SEM image, and EDS spectrum for (b) Ni matrix and (c) Ti(C,N).

4.2.3 EBSD Analysis and Grain Boundary Characterization

EBSD analysis has been conducted on samples in both as-received and creep-tested

conditions. The results obtained were able to provide a variety of information which is inaccessible from conventional microstructure characterization techniques, including grain size distribution, texture, recrystallization, grain growth, grain boundary development, and so on.

Since EBSD data was obtained from point-by-point scans, statistical analysis could be performed on volume data to give significant results on the effect of grain boundary on creep property of alloy 617.

Inverse pole figure (IPF) maps were obtained from as-received samples in both rolling

direction and transverse direction as shown in Figure 4.23. The color on each grain represents the

orientation of the crystal lattice as shown in IPF. It can be seen from the maps that grains are

randomly oriented without specific preference. The grains are generally equiaxed, and there is no

80 elongation in any direction. This means that there is no texture found in the sample, which also can be confirmed by the pole figures obtained from same scan shown in Figure 4.24.

Furthermore, each grain is uniformly colored, which means there is no small angle misorientation existing within the grain. It is an indication that the grains are well developed during the annealing.

IPF maps were also generated for creep tested specimens from both cross-section direction and rolling direction. The representative maps of a few test conditions are shown here in Figure 4.25 to Figure 4.28. Compared with as-received samples, creep tested samples show strong color variations within each grain, which means that regions with small angle misorientations exist within the grains. This is an indication of subgrain boundary formation and subgrain development process caused by plastic deformation. The subgrain boundaries are usually composed of arrays of dislocations, and they will pile up against precipitates and grain boundaries as creep proceeds. However, no strong texture can be found in crept samples either, as there is no strong preference in grain orientation shown in the maps. This observation is also confirmed by and the pole figures for crept samples as shown in Figure 4.29. Although less random distribution is displayed than that for as-received samples, no obvious preference is to be found.

(a)

(b)

Figure 4.23 IPF maps for as received alloy 617 samples in (a) cross-section direction and (b) rolling direction.

(a) (b)

Figure 4.24 Pole figures for as-received alloy 617 sample from (a) cross section direction and (b) rolling direction.

(a) (b)

Figure 4.25 IPF maps for alloy 617 samples underwent creep tests at 850°C 50MPa in (a) cross section direction and (b) rolling direction.

(a) (b)

Figure 4.26 IPF maps for alloy 617 samples underwent creep tests at 850°C 100MPa in (a) cross section direction and (b) rolling direction.

(a) (b)

Figure 4.27 IPF maps for alloy 617 samples underwent creep tests at 950°C 50MPa in (a) cross section direction and (b) rolling direction.

(a) (b)

Figure 4.28 IPF maps for alloy 617 samples underwent creep tests at 950°C 100MPa in (a) cross section direction and (b) rolling direction.

(a) (b)

Figure 4.29 Pole figures for creep tested alloy 617 samples at different test conditions, (a) 850°C 50MPa, (b) 850°C 100MPa, (c) 950°C 50MPa, and (d) 950°C 100MPa.

Statistical results of grain size, and grain size distribution were also obtained from EBSD data. Twin boundaries were excluded in calculation for grain size. It also needs to be noticed that because the scan size with respect to the grain size is small, the statistical sample size for included number of grains is not ideal, for grain size statistics. Therefore, a larger scan size covering more grains would be preferred in order to obtain an accurate value on the average grain size. The average grain size data for as-received and some of the creep tested samples are listed in Table 4.6. The grain size distributions for same samples are shown in Figure 4.30. The grain sizes of both before and after creep samples exhibits a bimodal distribution with large grains in ~150-250µm diameter and small grains in ~10µm. This is due to the pinning effect of carbides during solid solution annealing. These carbides segregate along the grain boundaries, impeding the grain boundary migration, so that grain growth becomes much slower in the carbide rich areas. On the contrary, grain growth is faster in the areas free of carbides, leading to

86 the large sizes of grains. Lillo et al. [54] reported similar findings by showing that the bands of small grains were approximately parallel with tensile direction, and high density of carbides pinned grain boundaries in these areas and prevented grain growth. In addition, a slight decrease in average grain size after creep is found. It is found to be due to the formation of new small grains by dynamic recrystallization (DRX), which is favored at high temperature and high strain range. The DRX in alloy 617 was also observed in elevated temperature tensile tests, especially near strain concentrated regions [78].

Table 4.6 Average grain size of alloy 617 from EBSD analysis

Sample Temperature (°C) Stress (MPa) Average Grain Size (µm) As-received n/a n/a 128.4 617-D62 850 50 107.5 617-D64 850 100 84.3 617-D61 950 50 86.2 617-D63 950 100 80.1

Grain Size (diameter) Grain Size (diameter)

0.5 0.6

0.4 0.4 0.3

0.2 0.2 Number Fraction 0.1 Number Fraction

0.0 0.0 20 40 60 80 100 120 140 160 180 100 200 Grain Size (Diameter) [microns] Grain Size (Diameter) [microns]

(a) (b)

Grain Size (diameter) Grain Size (diameter)

0.6 0.5

0.4 0.4 0.3

0.2 0.2

Number Fraction Number Fraction 0.1

0.0 0.0 20 40 60 80 100 120 140 160 20 40 60 80 100 120 140 160 Grain Size (Diameter) [microns] Grain Size (Diameter) [microns]

Grain Size (diameter)

0.5

0.4

0.3

0.2

Number Fraction 0.1

0.0 20 40 60 80 100 120 140 160 Grain Size (Diameter) [microns]

(e) Figure 4.30 Grain size distribution for (a) as-received, and creep-tested at (b) 850°C 50MPa, (c) 850°C 100MPa, (d) 950°C 50MPa, and (e) 950°C 100MPa samples of alloy 617.

Another samples was cut from the shoulder section of the specimen that underwent creep test at 950°C 100MPa, and EBSD was conducted on it. It went through exact same test, but there

is no stress applied on this part of the specimen. Therefore, the effect of stress is eliminated, and

the sample is like being annealed at 950°C for about 1 hour. The IPF maps for both cross-section and rolling directions are shown in Figure 4.31.

(a) (b)

Figure 4.31 IPF maps for shoulder section of alloy 617 specimen underwent 950°C 100MPa creep test in (a) cross-section direction and (b) rolling direction.

Besides orientation, grain size, and texture, EBSD data is also very useful in grain boundary characterization, which provides important information on creep property and mechanism. Grain boundaries can be characterized by coincident site lattice (CSL) model according to their geometry. In CSL model, the degree of fit Σ between the structures of the two grains is described by the reciprocal of the ratio of coincidence sites to the total number of sites.

Generally a boundary with high Σ might be expected to have a higher energy than those with low

Σ. Low Σ boundaries such as coherent twin boundaries (Σ = 3) have special properties as the boundary plane contains a high density of coincident sites. Therefore, grain boundaries can be categorized into three types: (1) low angle grain boundaries (LAGB) with Σ = 1 and

misorientation angles between 2°-15°, LAGB usually are subgrain boundaries; (2) special

boundaries or CSL boundaries with Σ ≤ 27, majority of these boundaries are twin boundaries for

FCC materials; and (3) random high angle grain boundaries (HAGB) with misorientation angles

higher than 15°, these are grain boundaries. As a result, Grain Boundary Character Distribution

chart can be generated based on the length fraction of boundary type. Representative charts for a

few conditions are shown in Figure 4.32, and data for each sample are listed in Table 4.7.

It can be seen from the GBCD for as-received sample and shoulder section sample, both of which had not been subjected to stress, no LAGBs exist, while LAGB fraction dramatically

increased after creep test. It indicates that subgrain formation is introduced by stress or

deformation in creep, and these subgrain boundaries consisting of arrays of dislocations actively

rearrange and interact with each other to accommodate the deformation. For random HAGB,

there is no obvious change in their fraction before or after creep, meaning that grain structures

are still intact and no cracks or voids have been formed.

CSL boundaries are of great interest due to their desirable characters. Although they may have high degree of misorientation, this type of boundaries has lower grain boundary energy than random HAGB. Hence, they are more stable and are more effective in hindering the dislocation movements so as to improve creep property. From the GBCD charts, the CSL boundary fraction is lower for crept samples than as-received sample, which is due to the emerging of LAGB during creep. Putting the CSL boundary fraction to the perspective of creep conditions shows an interesting correlation with stress and creep strain. CSL boundary fraction generally decreases with the increase of stress levels at the same temperature, as shown in Figure 4.33. This observation indicates that higher stress forces more boundary reconfiguration such as grain rotation, so that CSL boundaries break-up and lose their special characters. It also suggests that

LAGB introduced in creep has close correlation with stress. In addition, there is no obvious difference in CSL boundary fraction between different temperatures, meaning there is no

90 correlation between boundary character and temperature in this regime. Moreover, one alloy 617 specimen was tested at 950°C 50MPa but the test was terminated at 10% strain of the specimen corresponding to minimum creep rate range. Together with the as-received specimen and same test condition but 20% creep strain specimen, they form a series with progressive creep strain.

Another important observation is obtained thereafter by plotting CSL boundary fraction vs. creep strain, as shown in Figure 4.34. There is a clear drop in CSL boundary fraction as creep strain increases, which suggests that the transformation of CSL boundaries happens progressively during creep process, and it may because of the interaction with dislocations and breakdown of these boundaries.

GBCD - Grain Boundary Character Distribution GBCD - Grain Boundary Character Distribution

0.8 0.4

0.6 0.3

0.4 0.2

0.2 0.1 Number Fraction Number Fraction

0.0 0.0 2.0<= angle <15.0 CSL angle > 15 2.0<= angle <15.0 CSL angle > 15 Boundary Character Boundary Character

(a) (b)

GBCD - Grain Boundary Character Distribution GBCD - Grain Boundary Character Distribution

0.8 0.4

0.6 0.3

0.4 0.2

0.2 0.1 Number Fraction Number Fraction

0.0 0.0 2.0<= angle <15.0 CSL angle > 15 2.0<= angle <15.0 CSL angle > 15 Boundary Character Boundary Character

Figure 4.32 Grain Boundary Character Distribution (GCBD) charts for (a) as-received sample, (b) creep tested at 850°C 50MPa sample, (c) 950°C creep tested shoulder section (0MPa) sample, and (d) creep tested at 950°C 100MPa sample.

Table 4.7 Data for Grain Boundary Character Distribution (GCBD) charts

Temperature Stress LAGB CSL Random HAGB Specimen (°C) (MPa) fraction fraction fraction As-received n/a n/a 0.0185169 0.666411 0.315072 617-D62 850 50 0.341627 0.381081 0.277292 617-8 850 75 0.341031 0.332469 0.3265 617-D64 850 100 0.486432 0.257986 0.255582 617-12 950 30 0.34263 0.433154 0.224216 617-D61 950 50 0.442648 0.393763 0.163589 617-15 950 50 0.26577 0.517558 0.216673 (10% creep) 617-5 950 75 0.382183 0.336843 0.280974 617-D63 950 100 0.395535 0.32401 0.280455 617-D63 950 n/a 0.037699 0.695595 0.266706 (shoulder) 617-13 1050 30 0.705294 0.153364 0.141343 617-7 1050 50 0.308553 0.406601 0.284846 617-11 1050 75 0.600485 0.164938 0.234577

Figure 4.33 CLS boundary fraction vs. stress at 850 and 950°C.

Figure 4.34 CSL boundary fraction vs. creep strain at 950°C 50MPa.

In FCC materials, twin boundaries consist majority of the special boundaries. Primary twin boundaries are Σ3 boundaries formed by a 60° rotation about <111> direction. Secondary twin boundaries are Σ3 boundaries formed by 38.5° rotation about <110> direction. Unlike random HAGB, twin boundaries have high degree of atomic matching, low grain boundary energy, and thus low mobility. Therefore, twin boundaries are believed to possess desirable properties such as resistance to dislocation movement, cracking and corrosion [79]. In addition, twin boundaries can be categorized into coherent twin boundaries and incoherent twin boundaries. Coherent twin boundary requires the boundary plane to be aligned with the twin plane, while incoherent twin boundary doesn’t require that. In TSL OIM software, boundary plane trace and twin plane trace are examined, if they align to each other, the twin boundary is considered coherent, and vice versa.

Grain boundary maps were generated based on boundary characters for specimens as- received and creep tested at different conditions, as shown in Figure 4.35. Green lines are LAGB, blues lines are HAGB, and red lines are twin boundary. Actually, in the maps, twin boundaries and HAGB overlap with each other as twin boundaries are also considered as HAGB by their misorientation angles. By comparing Figure 4.35 (a) as-received specimen and (b) specimen creep tested at 850°C and 50MPa, it is obvious that a large number of LAGB emerged during creep, as also confirmed by the GBCD data mentioned previously. By close observation of the twin boundaries, we can find that the twin boundaries are smoother in as-received specimen, while in crept specimens, some of them are curved, and some are even discontinuous. This suggests that the break-up of twin boundaries happened during creep due to the interaction with dislocation, the grain rotation, and even carbide redistribution and grain boundary diffusion. On the other hand, comparing as-received specimen with shoulder section of the specimen creep

tested at 950°C 50MPa, it is easy to find that not much change happened during creep, which

also confirms the dislocation and subgrain formation are introduced by stress.

Data of the length of each type of boundary can also be obtained from grain boundary

maps, as listed in Table 4.8. Therefore, correlations between test conditions and boundary lengths or fractions could be established. As twin boundaries are considered part of HAGB, the

length fraction of those in HAGB could reflect the change made by creep, and it, in the mean

time, also excludes the effect of the emerging of LAGB. By plotting the length fraction of twin

boundary in HAGB vs. stress, as shown in Figure 4.36, we can find that twin boundary fraction clearly decreases with the increase of stress, and there is not much variation due to different temperature. It is indicated that twin boundaries are active elements in the creep process in terms

of rotation, deforming and interacting with dislocations, and all these activities are related to

stress levels.

(a) (b)

Figure 4.35 Grain boundary maps of (a) as-received specimen, (b) creep tested at 850°C 50MPa specimen, (c) 950°C creep tested shoulder section (0MPa) specimen and (d) creep tested at 950°C 100MPa specimen.

Table 4.8 Boundary length data for specimens tested at different conditions

Temperature Stress LAGB HAGB Twin Boundary Twin/HAGB Specimen (°C) (MPa) (cm) (cm) (cm) fraction (%) As-received n/a n/a 0.075 3.98 2.54 63.8 617-D62 850 50 1.66 3.21 1.30 40.5 617-D64 850 100 2.80 2.96 0.82 27.7 617-12 950 30 1.04 1.99 1.09 54.8 617-D61 950 50 1.74 2.19 1.01 46.1 617-15 950 50 1.05 2.90 1.60 56.9 (10% strain) 617-5 950 75 2.10 3.40 0.99 29.1 617-D63 950 100 2.80 2.96 0.82 27.7 617-D63 950 n/a 0.096 2.45 1.71 69.8 (shoulder)

Figure 4.36 Twin boundary fraction in HAGB vs. stress for specimens creep tested at 850 and 950 °C.

To sum up, EBSD could provide a lot of very useful information regarding grain size, grain boundary, texture, etc. Detailed analysis is the key to extract the important messages from the data. In this study, grain size and grain size distribution information was obtained by statistical analysis on EBSD data. Grain boundaries were characterized according to their misorientations and properties. Various maps and plots were generated to show the change happened during creep and to ascertain the responsible factors. Correlation between the change in length and fraction for different types of grain boundary and test conditions was also established. Grain size and grain boundary characters have great influence on material properties and can be altered by means of grain boundary engineering such as thermomechanical processing so as to control the material properties. The creep property of alloy 617 would be

benefitted by increasing the fraction of special boundaries, resulting in a reduction of grain

boundary diffusion and precipitation of carbides, thus delaying the onset of tertiary creep, and improve creep rupture life.

4.2.4 TEM Observation and Analysis

TEM imaging and detailed dislocation analysis has been performed on creep tested

specimens in different conditions, with special attention given to subgrain and grain boundary

areas, in order to obtain the information of dislocation properties and evidence of creep

mechanisms. TEM images were taken in different zone axes in order to analyze the Burgers

vector, orientation, and type of the dislocations. Selected Area Diffraction (SAD) patterns were

recorded for zone direction identification. The double tile specimen holder was used to give most

flexibility in orienting the specimen by allowing the specimen to rotate about two orthogonal

axes. Based on the Kikuchi pattern obtained from electron diffraction, the specimen can then be

oriented to the desired zone axes. The images were recorded under “two-beam” conditions

achieved by orienting the specimen slightly off the zone axis for best contrast. Then bright field

condition was achieved by aligning the apertures with the transmitted beam. The tilting angles

were recorded for zone axes and two-beam conditions when the images were taken, for detailed

orientation calculation of features of interests.

As shown in Figure 4.37 for specimen D-61 tested at 950°C 50MPa, five images were

recorded from the exact same area at two different zone directions, namely [011] and [101].

From the images, one array of dislocations can be seen crossing the subgrain boundary formed by dislocations and dark area in the upper right corner is carbide. Red arrow on the images marked the dislocation to be analyzed, which also represents other dislocations in the array since

they are in the same orientation. The dislocations are visible when g = (200), (11-1), and (020),

but invisible when g = (-11-1), and (-111). According to g·b invisibility criterion, the Burgers vector of the dislocation is determined as b = a/2[110]. With the help of the software named

Desktop Microscopist, a stereographic projection can be generated and used for trace analysis to

determine the line direction of dislocation being UT = [011], as shown in Figure 4.38(a). The type of dislocation can then be determined by calculating the angle ρ between Burgers vector and line direction. Pure edge dislocation has a Burgers vector perpendicular to the line direction.

Burgers vector is aligned with the line direction in pure screw dislocation. If the Burgers vector is neither perpendicular nor parallel to line direction, it is determined as mixed dislocation, which has both edge and screw character. The angle ρ for this dislocation is calculated as 54.2° so that

it is a mixed dislocation. By rotating the stereographic projection until both dislocation trace and

Burgers vector on the circumference, the slip plane of dislocation can be determined being (-11-

1), as shown in Figure 4.38(b). To sum up the results, the analyzed set of dislocations are mixed dislocations with Burgers vector of a/2[110], line direction of [011], and slip plane of (-11-1).

(a) (b)

(e) (f)

(g)

Figure 4.37 TEM micrographs of specimen D-61, creep test condition 950°C 50MPa.

100

(a) (b)

Figure 4.38 Stereographic projections for (a) dislocation trace analysis and (b) dislocation slip plane determination for specimen D-61 tested at 950°C 50MPa.

TEM images for specimen D-62 creep tested at 850°C 50MPa show a subgrain boundary composed of an array of dislocations, as shown in Figure 4.39. Dislocation analysis results show the Burgers vector of the dislocation is a/2[110], and the line direction is perpendicular to

Burgers vector. Therefore it is an edge dislocation. The slip plane is determined as (-110), giving the evidence of dislocation climb. This observation indicates that when there is low-angle misorientation, dislocations can move by climbing to overcome and thus form the subgrain boundary.

(a) (b)

101

(e)

Figure 4.39 TEM images of specimen D-62 area 1, creep test condition 850°C 50MPa.

Similar analyses were conducted on the specimens that are tested at different conditions, and results are listed in Table 4.9. From the table, most dislocations being analyzed are mixed type, not dependent on temperature or stress. The major slip system is found to be {111}<110>, typical for FCC material, suggesting that dislocation glide mechanism is in dominant at most conditions.

Some features are of great interest in providing evidence for creep mechanism. As images shown in Figure 4.40 for specimen 617-14 tested at 850°C 30MPa, two arrays of dislocations cross with a third array of dislocation, and pile up against the precipitates. It is clear that the distances between dislocations are lower in the area close to precipitate, indicating the pinning

102

effect of the precipitates. The images also show the precipitates have dislocation pass them, which confirms the power law dislocation climb-glide mechanism is in dominant. Another area for the same specimen shows more precipitates in different shapes, as in Figure 4.41. Generally all the observed precipitates are somewhat cuboidal shape, not spherical, and they provide good effects in impeding dislocation movements.

Table 4.9 Dislocation analysis results for specimens creep tested at different conditions

Temperature Stress Burgers Line Dislocation Specimen Slip Plane (°C) (MPa) Vector Direction Type 617-14-1 850 30 a/2[101] [011] (11-1) mixed 617-14-2 850 30 a/2[101] [011] (11-1) mixed 617-14-2 850 30 a/2[01-1] [001] (-100) mixed D-62-1 850 50 a/2[101] [-110] (11-1) mixed D-62-2 850 50 a/2[110] [001] (-110) edge D-62-3 850 50 a/2[011] [101] (-1-11) mixed D-62-4 850 50 a/2[101] [011] (11-1) mixed D-64-1 850 100 a/2[0-11] [0-11] (111) screw 617-12-1 950 30 a/2[10-1] [001] (-11-1) mixed 617-15-1 950 50 a/2[011] [101] (-1-11) mixed 10% strain 617-15-1 950 50 a/2[0-11] [101] (-111) mixed 10% strain D-61-1 950 50 a/2[110] [011] (-11-1) mixed D-61-2 950 50 a/2[110] [122] (-22-1) mixed D-61-3 950 50 a/2[01-1] [001] (-100) mixed D-61-4 950 50 a/2[01-1] [001] (-100) mixed 617-13-1 1050 30 a/2[01-1] [001] (-100) mixed

TEM images recorded for specimen D-62 creep tested at 850°C 50MPa are shown in

Figure 4.42. These images clearly displayed that subgrain boundaries are composed of arrays of

dislocations. These dislocations move in different directions and interact with one another to

form subgrain boundaries. Some dislocations within the subgrain are also rearranging themselves

103

and moving towards boundaries. Another area from the same specimen showed an array of

dislocation in very high density piling up against precipitate, as shown in Figure 4.43.

Stacking faults were found in the TEM images for specimen 617-15 creep tested at

950°C 50MPa and terminated at 10% creep strain, as shown in Figure 4.44. Dislocation slip

generally occurs by the glissile motion of perfect dislocations on close packed planes and in

close packed directions in one of the twelve {111}<110> slip systems. Stacking fault is

generated in certain conditions when providing favorable lower strain energy by dislocation

reaction. In the reaction, dislocations dissociate into Shockley partials with a Burgers vector of

a/6<112>, as

[101] [211] + [112] 2 6 6 푎 푎 푎 � → �� (4.7)

The region between these two Shockley partials is the stacking fault. As the result of atomic

plane disruption, stacking faults can be categorized into intrinsic or extrinsic type. An intrinsic

stacking fault is formed when an atomic plane is removed, for FCC materials, the original

ABCABCABC sequence is changed to ABCACABC. On the contrary, the extrinsic stacking

fault is formed when an additional atomic plane is inserted, altering sequence to

ABCACBCABC. These stacking faults are possibly formed by the collapse of vacancies,

condensation of interstitials, or by shear displacements of atomic planes. Unfortunately, more

information is needed for a complete stacking fault analysis in this specimen.

Interaction between dislocation and subgrain or grain boundaries were also recorded in

some specimens. Figure 4.45 shows an array of dislocations pile up against grain boundary but

cannot pass through it at 950°C 50MPa test condition. For the specimen tested at a higher

temperature of 1050°C as shown in Figure 4.46, an array of dislocations slip through a subgrain

104 boundary without any difficulty. In Figure 4.47 for another area on the same specimen, totally three arrays slip through subgrain boundaries can be observed at different orientation. However, boundaries still have impedance to the slip of dislocations, as can be seen from the higher dislocation density close to boundaries. Higher temperature leads to higher diffusivity in matrix, which is easier for dislocation movement.

(a) (b)

Figure 4.40 TEM images of specimen 617-14 area 1, creep test condition 850°C 30MPa.

105

(a) (b)

Figure 4.41 TEM images of specimen 617-14 area 2, creep test condition 850°C 30MPa.

(a) (b)

106

(e)

Figure 4.42 TEM images of specimen D-62 area 2, creep test condition 850°C 50MPa.

(a) (b)

107

(e)

Figure 4.43 TEM images of specimen D-62 area 3, creep test condition 850°C 50MPa.

(a) (b)

108

Figure 4.44 TEM images of specimen 617-15, creep test condition 950°C 50MPa, terminated at 10% creep strain.

(a) (b)

Figure 4.45 TEM images of specimen D-61, creep test condition 950°C 50MPa.

109

(a) (b)

Figure 4.46 TEM images of specimen 617-13 area 1, creep test condition 1050°C 30MPa.

(a) (b)

110

Figure 4.47 TEM images of specimen 617-13 area 2, creep test condition 1050°C 30MPa.

4.3 Static Recrystallization Study

In order to study the effects of grain size and grain boundary character on creep mechanisms, alloy 617 with wide range of grain sizes was developed by static recrystallization.

As-received alloy 617 plate was first cold rolled by multiple passed to reach the reduction of 10,

20, 30, 40, 50, and 60% of the original thickness. Then the samples were cut from deformed plates to be annealed at 900 and 1000°C for different length of time. Hardness for each sample was measured by Micro Indenter. Optical Microscope and EBSD were utilized for microstructure and grain boundary character examination.

Evidence for working hardening during cold rolling is displayed as the hardness increases

with thickness reduction, as shown in Figure 4.48. In the optical microscope images, as-received

sample displays equiaxial grains, 30% deformed sample displays somewhat elongated grains,

and 60% deformed sample clearly shows elongated grains in layered structure along rolling

direction, as shown in Figure 4.49.

111

Figure 4.48 Hardness for cold rolled alloy 617 samples.

(a) (b)

112

Figure 4.49 Optical microscopy images for alloy 617 in (a) as-received, (b) 10% deformed, (c) 30% deformed, and (d) 60% deformed conditions.

IPF maps generated from EBSD data for same samples as shown in Figure 4.50. As- received sample shows uniformed color in each grain meaning no misorientation within grains.

10% deformed sample starts showing a little variation in color, and 30% deformed sample shows a lot color variations, indicatating the formation of LAGB, and some evidence of grain boundary breaking are also discovered. In 60% deformed samples, grains are highly deformed and forced to intertwine with each other, large part of grain boundaries are broken. This indicates that the deformation by cold rolling results in extensive work hardening, and it increases with the deformation strain.

IPF maps for annealed samples show active recrystallization process happens during annealing, as shown in Figure 4.51. For 10 and 30% deformed sample annealed at 900°C for only 15 minutes, LAGB were greatly reduced and some small size grains were already formed.

For 60% deformed sample annealed at same temperature for 15 minutes, the recovery happened even faster. Small grains with well-developed grain boundaries were formed, and barely LAGB could be found. After 3 hours annealing at 900°C, the fine grains were fully developed with high fraction of twin boundaries, which may provide excellent resistance to creep. These observations 113

suggest that the process of recovery, nucleation, and grain growth are active from the early stage

of annealing. It is also confirmed by the hardness measurements of annealed samples, as shown

in Figure 4.52. The hardness values dropped dramatically within the first 15 minutes of

annealing, but generally maintained at same level for the rest of time. The rates of recovery are also correlated to the deformation, as it increases with the deformation percentage.

(a) (b)

Figure 4.50 IPF maps for alloy 617 in (a) as-received, (b) 10% deformed, (c) 30% deformed, and (d) 60% deformed conditions.

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(a) (b)

Figure 4.51 IPF maps for alloy 617 samples for (a) 10% deformation annealed at 900°C for 15 minutes, (b) 30% deformation annealed at 900°C for 15 minutes, (c) 60% deformation annealed at 900°C for 15 minutes and (d) 180 minutes.

Figure 4.52 Hardness of samples annealed at 900°C for different time.

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The grain size information was also obtained from EBSD data, and plotted vs.

deformation percentage, as shown in Figure 4.53. It can be seen that a wide range of grain size

was successfully manufactured by cold rolling and annealing process. A very fine grain size at 2-

3µm was produced by 60% cold rolling followed by 15 minutes annealing at 900°C. The grain size decreases with the increase in deformation. In addition, the grain sizes varied a lot among different cold deformation, but are very close to each other at same deformation percentage with different annealing conditions. It suggests the grain size has greater dependent on cold deformation than annealing temperature and time.

Figure 4.53 Grain size vs. cold deformation percentage for alloy 617 samples after annealing at 900 and 1000°.

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4.4 Discussion

4.4.1 Dislocation Creep

Power law dislocation climb-glide mechanism is proposed as the dominant creep mechanism for alloy 617 at 850-1050°C, 5-100MPa, as the stress exponent values are determined close to 5. This mechanism has two components, dislocation climb and dislocation glide, as schematically shown in Figure 4.54, and the creep rate is controlled by the slower of the two rates. Usually at high temperature, dislocation climb happens much slower than glide, and thus is rate controlling.

The TEM images show relatively low dislocation density within grains, and majority of dislocations are observed as straight lines moving in arrays, indicating the dislocation movement and annihilation happen quickly during creep. The analysis give the result of most dislocations in typical FCC {111}<011> slip system. These are the clear evidences to dislocation glide mechanism.

The dislocation climb process is driven and controlled by the diffusion of vacancies to or from the climbing dislocation through lattice, which depends on the temperature and activation energy. At high temperatures, the pinning effect of precipitates is not efficient as at moderate temperatures, and dislocations with enhanced mobility are able to pass them by climb or cross- slip. The low dislocation density within grains suggests that dislocation climb rate is relatively high at high temperatures. Dislocation climb mechanism allows edge dislocation to move perpendicular to its slip plane to overcome the obstacles. Evidence from TEM observations were found as edge dislocations on {110} plane to confirm the dislocation climb. Dislocation analysis determined them being mixed type dislocation which also have the capability to climb over the

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obstacles. Combined the evidences observed by TEM, dislocation climb-glide mechanism is

proved to be the dominant mechanism in alloy 617 creep.

Figure 4.54 Schematic diagram showing dislocation climb-glide mechanism.

4.4.2 Dynamic Recrystallization

According to the study by Weertman [80], a steady-state creep should be achieved for

dislocation climb by the balancing between strain hardening and stress relaxation recovery.

However, a well-defined secondary creep stage was not observed in ASTM 3.5 samples, and the

creep rate curves keep increasing after their minimum values were reached, which is in good

agreement with previous studies. The initial drop in creep rate found in low stress tests (≤50MPa)

is due to M23C6 carbides precipitation which hardens the matrix and pins dislocations during the temperature stabilization period before the stress applied, and this phenomenon is significantly

reduced by aging treatment [81].

This unconventional creep behavior with increasing creep rate after its minimum value is

reached, is the result of dynamic recrystallization (DRX). Unlike static recrystallization, DRX

occurs during straining of materials at high temperature and it only occurs after a critical amount

of strain is reached. Thereafter, the strain-hardening and recovery stop to be the dominant

mechanism, and DRX is more responsible for creep. The representation of DRX process is 118

shown in Figure 4.55, and described by Sakai [82]. First, dislocations are developed near the

grain boundaries and original boundaries start to corrugate. Then as the strain carry on, some

grain boundary sliding takes place, leading to the development of subgrain boundaries or twin

boundaries. Finally, subgrains are formed at the serrated grain boundaries.

(a) (b)

(c)

Figure 4.55 Schematic representation of the process of new grain formation during DRX [82].

DRX has been observed and reported in alloy 617 during creep by Chomette [81], Mo

[83], Sharma [84,85], and etc. It is preferred to happen in the areas with high concentration of carbides, because carbides have pinning effects on HAGB resulting in small grain sizes strengthen this area, and thus the HAGB together with carbides impede the movement of dislocation. So-called “particle deformation zones” are created by high density dislocations and large orientation gradient. In this way, HAGB and carbides provide concentrated nucleation sites, and higher dislocation density provides the source for subgrain boundary formation. During high

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temperature creep, therefore, subgrain structures can easily develop into new grains via the

transformation of subgrain boundaries into LAGB [86], and then new HAGB. In addition to the

carbides concentrated areas, the original HAGB are also preferred regions for DRX as mentioned previously.

DRX could be characterized by stress-strain curves as one or multiple peaks will appear as material hardens to a peak stress depending on deformation condition. Then softening happens attributed to the growth of recrystallized grains accompanied by dislocation annihilation. Typical

stress-strain curves for DRX are illustrated in Figure 4.56. In addition, the hardening rate, , as a

function of stress, usually decreases at a constant or decreasing rate, but DRX, on the other휃 hand,

causes an acceleration of the decreasing rate. Therefore, the overall creep rate increases after

DRX taking over in creep. Furthermore, in DRX, uniform and dislocation-free grain structure will not be attained as in static recrystallization, because the dislocations continue to be generated with straining, and keep providing the sources for DRX. Since DRX only happens in

some specific regions for alloy 617, and the rest areas have much less activities, the

recrystallized area is much smaller than the unrecrystallized area, so that a fully recrystallized

microstructure featuring steady-state creep could not be attained.

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Figure 4.56 DRX in Ni and Ni-Co alloys in torsion [87].

Evidences for DRX were observed in EBSD, as shown in Figure 4.57. Recrystallized grains with size smaller than 40µm were highlighted in the image quality maps, with dark lines being HAGB, and green lines being LAGB. There are only a limited number of small grains in the as-received sample, and they are randomly distributed on the HAGB. These grains, as mentioned previously, are formed during annealing because of the strong pinning effects of the carbides, preventing them from growth. On the contrary, recrystallized small grains can be observed as agglomerated in groups in both crept samples. These recrystallized grains are distributed along grain boundaries and areas with rich carbides, as they are the effective obstacles to dislocation movement, and thus provides sites suitable for DRX. The sample tested at 850°C 50MPa shows more recrystallized grains and they are more grouped together than the sample tested at 950°C 100MPa. It is because the creep happens much slower at lower

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temperature and stress, allowing more time for nucleation and grain growth in DRX. Therefore

DRX explains the slight decrease in the average grain size after creep as the number of small

recrystallized grain increases.

(a) (b)

Grain Size (diameter)

0.6

0.5

0.4

0.3

0.2 Number Fraction Number

0.1

0.0 20 40 60 80 100 120 140 160 180 200 220 240 Grain Size (Diameter) [microns] (c) (d)

Figure 4.57 EBSD image quality maps with highlighted small grains for (a) as-received, (b) creep tested at 950°C 100MPa, (c) creep tested at 850°C 50MPa, and (d) grain size distribution showing highlighted grain sizes.

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The recrystallized grain size is also dependent sensitively on the deformation conditions.

The strain dependence of DRX has also been reported by Mo [83]. Figure 4.58 shows IPF maps for samples cut from necking zone and 2mm away from necking zone after short term tensile test at 1000°C. Large fraction of recrystallized grains indicates DRX is more intense in higher strain area. Seemingly texture is formed in the original large grains, while recrystallized small grains have various orientations. Twin boundaries are also generated during DRX due to the alloy’s low stacking fault energy [79]. On the contrary, original twin boundaries lose their coherency during creep and evolve into high Σ CSL boundaries or HAGB. These two process leads to a small reduction in twin boundary fraction as observed in ASTM 3.5 samples.

(a) (b)

Figure 4.58 DRX in tensile test at 1000°C, (a) 2mm from neacking zone, (b) near necking zone [83].

Statistical data obtained from EBSD show that there is no obvious change in grain size during creep in alloy 617. The slightly decrease in average grain size at some conditions is due to

DRX. Alloy 617 is known for exceptional thermal stability because the carbides and other intergranular particles inhibit the grain boundary migration and grain growth, and thus the grain

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boundary characters stay unchanged during thermal aging. Tan et al. reported well preserved

GBCD after thermal aging of alloy 617 at 850 and 1000°C for 4 and 6 weeks [79,88]. However,

GBCD in alloy 617 during high temperature creep did evolve due to plastic deformation. While

HAGB fraction stayed unchanged, LAGB increased dramatically during creep, indicating the

subgrain boundary formation. More importantly, the fraction of CSL boundary decreases during

creep with the increase of stress levels, but independent on temperature. CSL boundary fraction

also depends on the creep strain, as it decreases with the increase in strain. Moreover, the

fraction of twin boundary over HAGB shows similar dependence on the stress levels. These

observations are on the contrary to Schlegel’s results [56], where there is slight increase or no

obvious change in the special boundary fraction. It is possibly because of the difference in creep

test condition and strain range of the specimens.

TEM observations confirmed the presence of carbides as the effective impedance to

dislocation movements. These carbides are believed to be M23C6 type, as they are mostly coherent with the γ matrix through a cube-on-cube orientation relationship according to previous studies by Wu [13],[38]. Dislocations are concentrated in the areas close to precipitates or grain boundaries, providing sources for DRX to further transform into subgrain boundaries.

The onset of DRX is marked by the reach of critical strain, which depends on the combination of a number of factors. The critical strain decreases with temperature increase and with decrease in strain rate as well as with decrease in the grain size. On the other hand, the precipitation of second phase particles has the opposite effect on DRX. When the pinning force due to the particles exceeds the driving force for boundary bulging, DRX is suppressed and finer

DRX grains are formed at larger strains, if at all [82]. This would well explain the inconsistency in creep behavior of ASTM 2 samples. The conventional creep curves were obtained at 850 and

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950°C, indicating steady-state power law dislocation was in dominant, and no DRX happening.

Because the dislocation density is not high enough to form subgrain boundaries, and thus DRX was suppressed. While the test temperature went up to 1050°C, the mobility of dislocation was higher and a lower critical strain was attained. Therefore DRX happened quickly and led to monotonic increasing creep curves. Similarly at 950°C for ASTM 2 samples, low stress level of

5MPa could not provide enough driving force for DRX to happen, but at 30MPa or higher stresses, DRX happened with ease.

The DRX is also sensitive to initial grain size. The flow curve by stress vs. strain shows

two froms, a single peak or multiple peaks with reducing amplitude leading to an effective

steady-state. It has been suggested that the fluctuation in flow curve indicates that grain coarsening is occurring while a single peak is associated with grain refinement [89]. When it is

fine grained material, grain coarsening takes place until a stable grain size is reached. In contrast,

the single peak form is associated with grain refinement in initial coarse grained materials.

Zener-Holloman parameter could be used to evaluate the initial and recrystallized grain sizes,

expressed by,

= 푄 푍 휀̇푒푥푝 � � 푅푇 (4.8)

= (4.9) 푚 where is Zener-Hollomon parameter, 푍퐷 is strain퐾 rate, is apparent activation energy, is gas

constant,푍 D is initial or recrystallized grain휀̇ size, K and푄 m are constants related to materials.푅

Zener-Hollomon parameter is also seen as temperature compensated strain rate. The value of

activation energy, , is an experimental constant. It is usually higher than that of self-diffusion, indicating that formation푄 under DRX is thermally activated and involves self-diffusion [91]. The

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relationship between grain size and Zener-Hollomon parameter is illustrated in Figure 4.59. In

general, with the increase in grain size, the strain rate by DRX decreases. This also explains why

DRX does not occur in ASTM 2 samples at 850 and 950°C.

Figure 4.59 Relationship between grain size and Zener-Hollomon parameter [90].

In general, DRX is the major reason for the unconventional creep behavior in alloy 617, particularly the increase in the creep rate after the minimum value is reached, and no well- defined steady-state creep. It is considered an effective metallurgical process for grain refinement during hot deformation. Therefore, thermal mechanical processing needs to be properly controlled in order to obtain suitable distribution of large grain size and inhibit grain refinement so as to develop desired mechanical properties for alloy 617.

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4.4.3 Diffusional Creep

As dislocation creep dominates at most test conditions, however, at high temperature and very low stress level conditions, there may be diffusional creep existing, especially in the fine- grained materials. Diffusional creep is caused by the mass transport of vacancies within grain as

Nabarro-Herring creep or along the grain boundary as Coble creep.

The minimum creep rate of diffusional creep for alloy 617 ASTM 8 specimens can be calculated by the equation as following,

Nabarro-Herring creep:

= 퐷퐿 휎Ω 휀푁퐻̇ 퐴푁퐻 ∙ 2 ∙ � � 푑 푘푇 (4.10)

Coble creep:

= 퐷퐺퐵 ∙ 훿′ 휎Ω 휀퐶̇ 퐴퐶 ∙ 3 ∙ � � 푑 푘푇 (4.11) where, and are constants, is grain size, is atomic volume, is Boltzmann constant,

푁퐻 퐶 and 퐴and are퐴 diffusivity in lattice푑 and grain Ωboundary, and can be 푘calculated by

퐷퐿 퐷퐺퐵 = 푄 퐷 퐷0푒푥푝 �− � 푅푇 (4.12)

Given a high temperature low stress condition of 1050°C 2MPa, assuming there is no grain size effect to the value of activation energy for Nabarro-Herring creep, the average activation energy of 405kJ/mol obtained from ASTM 3.5 specimens could be applied. =

0 8.5×10-5 m2/s and = 12 (usually 10~14). The minimum creep rate of Nabarro-Herring creep퐷

푁퐻 is calculated being퐴 1.12×10-9 h-1. In Coble creep, the vacancy diffusion happens along grain

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boundaries, giving a lower value of activation energy and diffusion coefficient. of 1.94×10-5

0 m2/s and Q of 128 kJ/mol calculated by Gust [92], and of 40 (usually 30~50)퐷 are used here.

퐶 The minimum creep rate of Coble creep therefore is calculated퐴 being 1.30×10-3 h-1. Considering

the temperature and stress conditions in this study together with the calculation results, Coble

creep contributes the majority part of the diffusional creep.

For creep conditions in the high temperature low stress regime, the creep mechanism for

alloy 617 is the result of a combined contribution from both dislocation creep and diffusional

creep, and can be expressed as following:

= + (4.13)

푡표푡푎푙 푑푖푠푙표푐푎푡푖표푛 푑푖푓푓푢푠푖표푛푎푙 Extrapolating the test data휀̇ for ASTM휀̇ 8 specimens,휀̇ the minimum creep rate at 1050°C

2MPa will be 1.87×10-3 h-1. This suggests that about 70% the creep is from diffusional creep at this condition, and dislocation creep only contributes about 30%. However, the relative contribution from diffusional and dislocation creep depends on temperature and stress, as shown in Figure 4.60. Since the difference in stress exponent, diffusional creep contributes more at

lower stress levels while dislocation creep contributes more at relatively higher stress levels.

Since either of the diffusional creep has a stress exponent of 1 in creep rate and stress

relationship, a relatively lower stress exponent is expected. Therefore, stress exponents being

around 3 for ASTM 8 specimens is a good evidence to that diffusional creep takes part in the

creep. As shown in Figure 4.61, dislocation creep for ASTM 3.5 specimens is in red lines,

having a stress exponent of ~5; diffusional creep by calculation is shown in green line with a

stress exponent of 1; and the test data from ASTM 8 specimens are shown in blue line. From the

combination of dislocation and diffusional creep, the stress exponent of ~3 is obtained.

128

Figure 4.60 Relative contribution from diffusional creep and dislocation creep.

Figure 4.61 Stress exponent calculation showing both dislocation and diffusional creep contribute to total creep strain.

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Deformation maps are usually used to indicate the dominant mechanism at different

temperatures and stresses. Although a complete deformation mechanism map for alloy 617 has

not been developed yet, the map for pure nickel could be a good reference, as shown in Figure

4.62. Alloy 617 has a melting temperature range of 1332-1380°C, shear modulus of 61-49GPa

between 800-1100°C. Therefore, the homologous temperature is about 0.55-0.65, normalized shear stress is 10-3-10-4. The covered region, as shown in Figure 4.62 in red, is close to the

boundary of power law creep. The boundary line for different mechanisms could be different

depending on grain size. Diffusional creep participates creep process in fine-grained alloy 617, as samples with different grain sizes show different creep behavior and stress expoents. Nabarro-

Herring creep mostly happens at extremely high temperature, above 0.9 of homologous temperature. Therefore the diffusional creep happening here in alloy 617 is most possibly Coble

creep. Detailed evidence from microstructural analysis will be needed to further investigate the

diffusional creep.

Figure 4.62 Deformation mechanism map for pure nickel with a grain size of 0.1mm [93].

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For specimens with ASTM 8 grain size, DRX was also observed and reported in the microcreep tests on foil specimens [94]. A decrease in average grain size is expected by DRX, and as a result, the diffusional creep which depends on the grain size will also be easier.

Therefore, after the critical creep strain is reached, DRX also participates in the creep, together with diffusional and dislocation creep. Microstructure examination on crept samples is needed in the future to provide direct evidences to different creep mechanisms.

To sum up, the alloy 617 materials with different grain sizes generally creep as in dislocation climb-glide mechanism. DRX happens at certain stain after the minimum creep is reached depending on the creep condition and grain size, and is the reason for the unconventional creep behavior. For fine-grained material, diffusional creep also participate in the creep process, at high temperature low stress level regime.

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5 Conclusions and Future Works

In this study, creep behavior of Alloy 617 was investigated by creep tests in a systematic test matrix consisted by various temperatures (850-1050°C) and stress levels (5-100MPa) for samples with three different average grain sizes (ASTM 8, 3.5, 2). Obtained creep data was analyzed to determine creep mechanism and parameters for creep model. Microstructure characterization was performed by OM, SEM, EBSD, and TEM in order to obtain the information of grain size, grain boundary, dislocation, etc, associated with creep behavior at different conditions. The conclusions are summarized as follows.

1. Alloy 617 samples with an average grain size of ASTM 3.5 showed an unconventional creep curve with little primary creep stage and unrecognizable secondary creep stage, while samples with an average grain size of ASTM 2 showed a well-defined secondary creep stage at temperatures of 850 and 950°C, but similar creep curve to ASTM 3.5 samples at 1050°C.

Samples with larger grain size showed a much lower creep rate compared with those having smaller grain size at same test conditions, suggesting grain size is an important factor to creep property. Larger grain size provide less chance for grain boundary diffusion and sliding, resulted from dislocation pile up and cracking. For ASTM 8 samples, the creep curves are also similar to

ASTM 3.5 samples at low stress levels, but they have the worst creep properties among all three grain sizes, suggesting small grain size is not favorable for long term creep applications.

2. Analysis of creep data from ASTM 3.5 samples gave stress exponents of 5-6, indicating power law dislocation climb-glide creep is the dominant mechanism at stress levels between 5-100MPa and temperatures between 850-1050°C. The activation energy was observed between 380-450 kJ/mol, which is higher than that of nickel self-diffusion. It is due to the involvement of precipitates which has much higher activation energies.

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3. For ASTM 2 samples, a lower stress exponent of ~3 was obtained at 850 and 950°C,

suggesting a solute drag creep mechanism. The stress exponent for 1050°C is determined as

close to 5, suggesting a power law dislocation climb-glide mechanism.

4. ASTM 8 sample showed stress exponents around 3, suggesting a combination contribution from both diffusional creep and dislocation creep.

5. EBSD observations of creep tested samples gave evidence for significant subgrain boundary formation within grains and dynamic recrystallization, which is the major reason of the

unconventional creep behavior of alloy 617. No strong texture was observed in both as-received

and crept materials. Grain boundary character distribution showed a dramatic increase in LAGB

and decrease in CSL boundaries during creep, and they depend on stress levels. Twin boundary

fractions also decrease with the increase of stress level, but no impact was found from

temperature.

6. TEM observations gave evidence for the presence of dislocations around precipitates

and grain boundaries. Detailed dislocation analysis confirmed dislocation climb-glide

mechanism. The pinning effects of carbides and the impedance of HAGB to dislocations were

also confirmed.

7. Static recrystallization experiments were carried out by cold rolling and then annealing

process to produce a wide range of grain sizes, as fine as 2-3µm. Hardness measurements and

EBSD analysis confirmed the recovery and twinning happens rapidly in the early stage of

annealing. The grain size was found to be dependent on the deformation but not annealing time.

8. Combining all the information obtained from creep testing and microstructural analysis,

the unconventional creep behavior of alloy 617 is due to enhanced dislocation climb rate at high

temperature, subgrain boundaries not being effective obstacles for dislocation movement, and

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most importantly, dynamic recrystallization continuously providing new volume for dislocation

activities. Diffusional creep also contributed to the creep for fine grain sized material at high

temperature low stress conditions.

Based on the analyses on creep testing and microstructure characterization, with

temperature, stress, grain size, and grain boundary character being taken into consideration, a

new insight is gained into the governing mechanism more of alloy 617 creep behavior. This

information will be helpful to the alloy development and code case establishment for the

applications in Generation IV nuclear energy system.

To fully understand the creep behavior of alloy 617, more studies are needed in the future.

First, more creep data is required in the low stress high temperature regime to establish a

quantitative deformation model for alloy 617. Detailed microstructural analysis is needed to

evaluate diffusional creep condition and effects on creep behavior. In addition, more

investigation is needed in dynamic recrystallization mechanism in alloy 617, including its onset condition, the preferred grain size, and the effects of grain boundary character. Grain boundary engineering techniques could also be employed to control the grain boundary characters and thereby improve the mechanical properties. Furthermore, for nuclear energy applications, environmental effects, such as irradiation and corrosion have to be taken into consideration for the evaluation of the material performance.

For the structure materials in future nuclear energy applications, the Ni-based oxide dispersion strengthened alloys (ODS) could be potential candidates. These alloys employ oxide dispersion as strengthener into the Ni alloying system, and could have better high temperature strength, creep resistance, and corrosion resistance. Therefore, more efforts could be directed towards the research on ODS.

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