THE ANALYSIS OF MICROSTRUCTURE AND GRAIN ORIENTATION …

17th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors August 9-13, 2015, Ottawa, Ontario, Canada

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THE ANALYSIS OF MICROSTRUCTURE AND GRAIN ORIENTATION OF THERMALLY-AGED ALLOY 600

Seung Chang Yoo, Kyoung Joon Choi, Taeho Kim, Seunghyun Kim and Ji Hyun Kim*.

School of Mechanical and Nuclear Engineering, Ulsan National Institute of Science and Technology, 50, UNIST-gil, Eonyang-eup, Ulju-gun, Ulsan 689-798, Republic of Korea

ABSTRACT

To establish a method to simulate long-term thermal aging in nuclear power plant, a representative thick-walled nickel-based Alloy 600 was prepared and thermally aged at 400 °C in Ar environment to simulate the thermal aging effect of 10 and 20 years in NPP operating conditions (320 °C primary circuit coolant). The basic material properties were investigated by electron microscopy, electron backscatter diffraction, tensile tests and nanoindentation. Formation of abundant semi-continuous precipitates was observed after 10 years of thermal aging that caused hardening of the material. Then, the precipitates were changed to continuous features and induced the softening of the material. These aspects were investigated with various tools in this study. Through this study, reliable data were produced to establish a method for simulating long-term thermal aging, which could contribute to performing precise experiments using thermally-aged materials instead of recently manufactured materials.

Keywords: nickel based alloy, thermal aging, precipitate, microstructure, tensile test

1.0 INTRODUCTION

Components used in nuclear power plant (NPP), and manufactured with Alloy 600, such as steam generator tubes, heater sleeves and head penetration nozzle for control rod drive mechanisms, may suffer from cracking and experiencing leakage. In particular, Alloy 600 has been reported to be susceptible to stress corrosion cracking (SCC) in the primary circuit of NPP, which is caused by a susceptible microstructure, corrosive environment and tensile stress, generally weld residual stress.

Many studies were conducted to determine the mechanism and establish numerical models of SCC. SCC is closely related and influenced by environment, microstructure and tensile stress, so studies were focused on those factors. Scott tested Alloy 600 in various environments to correlate intergranular SCC (IGSCC) with oxidation phenomenon. He suggested that IGSCC is caused by internal oxidation, induced by diffused oxygen through the grain boundary of nickel-based alloys in redox potential [1]. Furthermore, SCC tests evaluated the resistance of nickel-based alloys to crack initiation and crack growth. Alloys 600 and 690 were tested in a primary water coolant environment to determine the influence of the material`s properties on susceptibility to SCC initiation [2]. However, several reports pointed out that there is still a lack of understanding of SCC. Additional work is necessary to propose a more predictive model [3]. It was also reported that the aging of structural materials in a NPP has been accompanied by many SCC cases.

As mentioned before, the effect of thermal aging on a material property is quite important and can be accompanied many problems. Many studies investigated the effect of heat treatment on the properties. Kai et al. investigated Alloy 690 after various heat treatment conditions and found several significant changes [4]. Celin et al. summarized the degradation process of Ni–Cr–Fe alloys in a pressurized water reactor, and they concluded that Alloy 600 seems to pose a problem after being used in long-term operations [5]. Although these valuable studies emphasized the influence of heat treatment on the material properties, most of them involved SCC tests conducted on as-received materials, which never undergoes any thermal aging, except short-term heat treatment generally in high temperature above 600 °C [6]. However, a huge difference in phases and the mechanism of degradation process


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between thermal aging at low temperature (320 °C of NPP operation temperature) and high temperature (above 600 °C) may be suspected.

The effect of thermal aging becomes an issue as the operation time of NPPs increases. But studies considering a heat effect to material were only conducted at elevated temperature as mentioned above. Therefore, it is essential to develop a proper method to simulate long-term thermal aging in NPPs, and better extrapolate laboratory results. For this, the method to simulate the thermal aging in operating NPPs should be established in terms of microstructure and mechanical properties.

The objective of this study is to analyze in details; the microstructure and the mechanical properties of Alloy 600 after applying accelerated aging to simulate the actual long-term thermal aging experienced in NPPs

2.0 EXPERIMENT

2.1 Material

Doosan Heavy Industries & Construction provided a representative thick-walled nickel-based Alloy 600 of cylindrical shape fabricated according to ASTM B166 by an annealing and peeling process which is a process that peel off the damaged part of material after heat treatment. The Alloy 600 cylinder was annealed at 1060 °C for 3.5 h and water quenched. Table 1 shows the chemical composition of Alloy 600.

The material was then heat treated in an Ar environment to simulate the effects of actual long-term thermal aging in a NPP. Accelerated aging was performed via a high temperature heat treatment since it is too long to duplicate the actual thermal history of Alloy 600 in a NPP for 10 or 20 years at 320 °C. The temperature used was 400 °C which limits excessive formation of carbides or sigma phases that do not form at 320 °C according to thermodynamic calculation [7].

The heat treatment conditions were determined and performed based on equation (1), shown below, which is a modified form of the Arrhenius diffusion equation, calculating the required heat treatment time at 400 °C to simulate long-term thermal aging at reactor-operation conditions, i.e., 320 °C.

In this equation, R=8.314 J/mol/K, taging is the heat treatment time needed at 400 °C to simulate long-term thermal aging in a NPP, and tref is the service time at 320°C of a NPP or the time that would be simulated. The heat treatment temperature is represented by Taging (400 °C), and Tref is the actual operation temperature of the NPP (320 °C). For activation energy Q, the Cr diffusion energy at the grain boundary of 180 kJ/mol [8] was selected since Cr is known to play an important role in the precipitation and corrosion resistance, influencing the resistance to SCC of material [9].

Specimens, treated for three different heat treatment durations, were prepared for simulating materials aged in NPP for 0 (as-received state), 10 and 20 years. The heat treatment durations at 400 °C (corresponding to 10 and 20 years at 320 °C) were 1142 and 2284 hours, respectively. Table 2 shows the specimen I.D. (HT + ‘heat treatment temperature [°C]’ – Y + ‘simulated aging time [year]’), simulated aging time and temperature of each specimen.

2.2 Experimental Procedure

The microstructure of the thermally aged Alloy 600 was investigated with optical microscopy (OM) and scanning electron microscopy (SEM). Crystallographic orientations were analyzed with electron

exp (1)


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backscatter diffraction (EBSD). Mechanical properties were analyzed with nanoindentation and by tensile testing.

For the microstructural analysis, the specimens were cut by electrical discharge machining to the size of 10 × 10 × 3 mm3 and they were polished with emery papers, diamond paste, and colloidal silica up to 0.05 μm, and then polished with colloidal silica for about 1 hours to remove additional mechanical deformation caused by polishing process. The specimens were finally polished using a vibratory polisher to achieve an even smoother surface and to remove the residual stress which may be formed during the polishing process. After the polishing, HNO3-based 20%-HCl etching solution was used following ASTM E407-07.

The microstructural characterization was conducted using the Quanta 3D Dual-Beam Focused Ion Beam (FIB) attached to the field-emission gun SEM. EBSD analysis was done with the same instrument with the acceleration voltage of 10 kV, current of 0.64 nA, tilt angle of 70°, and step size of 2 μm. The scanned area used for the EBSD analysis was 298 μm × 881 μm, thus covering a large number of grains to achieve more accurate results.

To analyze the mechanical properties, the nanoindentation test was conducted with the Agilent Technologies Nano Indenter G200. The continuous stiffness measurement (CSM) technique of the Nano Indenter XP head with a three-sided pyramidal Berkovich tip was used for the whole nanoindentation test. The experimental conditions used were a total penetration depth of 2000 nm, a strain rate of 0.05 s-1 and a Poisson’s ratio of 0.3. 32 points were examined by 4 x 4 array matrix in two different regions of the same specimen. The spacing between indentations was about 40 μm and whole size of array is about 200 μm x 200 μm. Tensile tests were done in the Instron 8801 universal fatigue testing machine at the strain rate of 0.4 /s with a proportionally reduced specimen from ASTM standard ASTM E8-E8m. The modified scale specimen is shown in figure 1. Multiple tensile tests were performed to reduce the experimental error.

3.0 RESULTS AND DISCUSSIONS

3.1 Mechanical Properties

The representative stress–strain curves, obtained by tensile test, are shown in figure 2. The yield strengths, ultimate tensile strengths, and elongations are estimated from the measured curves and summarized in Table 3. Evidently, the mechanical properties changed with thermal aging time; this was most pronounced in the case of HT400_Y10 for which the yield strength and ultimate tensile strength were maximized and the elongation was minimized, showing loss of ductility. Meanwhile, it softened following the 20 years of thermal aging as shown in the result of HT400_Y20. This result could be explained by precipitate hardening and overaging due to morphology changes of precipitates during heat treatment [10, 11]. When a size of precipitate exceeds a critical point and forms continuous feature, a hardening of materials, by precipitation effects, is suppressed (i.e., overaging). It will be discussed in section 3.2 with SEM images. The error ranges of tensile properties seem large, however, it could be supported by showing same trends from nanoindentation.

From the fracture surface images of each specimen (as shown in figure 3), a typical ductile failure mode, accompanied many dimples, was observed. It is hard to quantify the size of dimples, since there size is not unique, and small dimples were observed in large dimples. However, the mean dimple size of the specimen HT400_Y10 is relatively large compared to the other two specimens. It may be correlated to the lower ductility found for HT400_Y10 compared to other specimens [12, 13].

The same tendency, observed after tensile tests, was also shown after nanoindentation: hardening in the specimen HT400_Y10 and softening in specimen HT400_Y20 as shown in figure 4. Very small differences were observed between the hardness values at the grain boundary and inside the grains (differences are within the error range) as shown in table 4. This indicates that precipitates are uniformly distributed at the grain boundaries and the plastic deformation during indentation is always influenced by those precipitates [14]. The plastic deformation zone of indent mark is typically 5 times larger than the indent mark (about 20~30 μm in this experiment), which is similar with grain size of thermally aged Alloy 600 which will be discussed in following section.


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3.2 Microstructure

The microstructures of each specimens were observed with OM (figure 5). The average grain sizes were about 26.7 μm for as-received, 24.6 μm for HT400_Y10 and 25.4 μm for HT400_Y20, in agreement with EBSD analysis. As shown in secondary electron images in figure 6, original precipitates grew due to thermal aging. The area fraction and average length of the precipitates were analyzed with the image-analyzer software, I-solution. The as-received specimen contains 0.36 % area fraction of precipitates, HT400_Y10 has 1.20 % and HT400_Y20 contains 1.43 %. The area fraction increased significantly after 10 years of thermal aging, while the change is negligible on further thermal aging time from 10 to 20 years. Meanwhile, the average lengths of precipitates were 0.801 μm for as-received, 0.923 for HT400_Y10 and 1.864 μm for HT400_Y20. The length of precipitates shows little growth after simulated 10 years of thermal aging, while a considerable growth was observed after simulated 20 years of thermal aging.

This result is in good agreement with the results of previous studies, which reported that precipitates are formed during the early stage of thermal aging, and they only are coarsened or migrate to more stable places with time and finally, formed a continuous features of precipitates [15-17]. From these previous studies using identical material as this study, the precipitates are assumed to be Cr7C3 or Cr23C6 [5, 15]. Abundant semi-continuous precipitates are known to increase the tensile properties of materials by providing a role as dislocation sources, which can blunt crack tips [16, 18]. Therefore, thermal aging could increase the resistance against crack growth and the susceptibility to crack initiation in early stage of thermal aging. However, when the size of precipitates reaches a limit and form continuous feature, the strength was decreased again. From the results of this study, it is found that the peak of tensile properties may be reached after 10 years thermal aging in NPP conditions, and softening will be followed due to formation of precipitates with continuous feature.

The hardening and softening mechanism, which is associated with precipitates, has been extensively investigated in nickel-based super alloys for extreme conditions of high temperature. The basic idea of precipitation hardening is that the resistance to the precipitate being cut by dislocation movement, and when precipitates are sufficiently grown, dislocations are forced to make a loop around the precipitate. Therefore, strengthening is weaker than the first phase [19-21]. The transition point is not well defined from hardening to softening, which is often suggested as a critical particle size, where dislocations begin to loop around the particle [19]. This could be one reason for hardening and softening behavior in long-term thermal aging in NPP.

Meanwhile, the grain boundary precipitates are known to have a mechanical effect on IGSCC. Several researchers propose that grain boundary precipitates were found to be a major dislocation sources during deformation at high temperature [22, 23]. They suggest that material with continuous feature of precipitates have lower susceptibility to SCC than materials with both semi-continuous and discontinuous precipitates. In this research, HT400_Y10 have semi-continuous precipitates, HT400_Y20 have continuous precipitates and As-received specimen have discontinuous precipitates. With consideration of chromium depletion at grain boundary due to formation of precipitates, it is expected that HT400_Y10, which have large amount of semi-continuous precipitates, have higher susceptibility to SCC initiation. And this will be verified in further experiment with PWSCC initiation test at primary water environment.

According to the Orowan mechanism, the material strength (τ) depends on the shear modulus (G), magnitude of burgers vector (b), distance between pinning particles (L) and second phase particle radius (r), as shown in equation (2).

(2)


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All the materials, used in this study, have the same chemical composition because the samples were prepared from the same heat of Alloy 600, but heat treatment conditions were different. So, it could be assumed that shear modulus and magnitude of burgers vector is constant. Then, the material strength depends on 1/(L-2r). Through the calculation with the parameters in Table 5, the upper graph, shown in figure 7, could be achieved from the relationship. The feature is same with that of mechanical test which was conducted in this study: hardening by 10 years thermal aging, and softening by 20 years thermal aging. Therefore, it is thought that the aspect of mechanical properties is closely related to the morphology of precipitates [19, 24].

On the other hand, according to the dislocation model by Ashby [16], many voids and overlaps are generated at the grain boundaries during a tensile test. Then, to correct those voids and overlaps, a number of dislocations are formed near grain boundaries, especially near precipitates in a grain boundary. According to this model, disconnected precipitates result in the concentration of stress and dislocations and therefore induce hardening of the material, as shown in the stress–strain curve of HT400_Y10 (in figure 2). In HT400_Y20, the network of connected precipitates disperses the stress, and the role as dislocation barrier will be weaken because of the decreased number of precipitates, at last, resulting in softening of the material [13, 15].

Grain boundary misorientation, with coincidence site lattice (CSL) theory, and kernel average misorientations (KAM) were analyzed by means of EBSD. CSL theory generally categorizes the grain boundaries into three groups: low-angle grain boundaries (2°–15°), high-angle grain boundaries (over 15°) and CSL boundaries which are classified based on Brandon`s criterion Δθ = 15° Σ-0.5, where Δθ is the angular deviation from the exact CSL boundary and Σ is the reciprocal of CSL [25-28]. Only CSL boundaries having Σ numbers below 29 were considered since those boundaries are known to have special characteristics in material degradation such as high corrosion and precipitation resistance [27]. KAM can estimate the residual strain distribution associated with inhomogeneous dislocation distribution [29]. The data cleaning was applied by TSL EBSD software including grain confidence index standardization and grain dilation process. These processes compare neighbor data points around the measured point and reduce abnormal data points.

In figure 8 (a–c), grain-boundary misorientation analysis results are shown. The colors of the figure, including namely green, black, and red, represent low-angle, high-angle, and CSL grain boundaries, respectively. The fraction of CSL boundaries was significantly decreased by thermal aging, as shown in figure 8 (a–c), with comparably large error range. The as-received specimen has 25.5 %, HT400_Y10 has 15.3 %, and HT400_Y20 has 6.8 % of CSL boundaries. No severe changes were observed in the KAM values, indicating that the local strain was not changed much by thermal aging. The microstructural characteristics of specimens are summarized in table 5. It could be carefully considered that drop out of solute near CSL boundaries during formation of precipitates could induce decrease of CSL boundaries by inducing disordering and dislocation pile up near CSL boundary. The decreasing of CSL boundary is remarkable, however, considering the large error range of the result, more studies will be needed for this finding with caution.

It has been observed that CSL boundaries have a higher resistance to SCC and precipitation than other types of grain boundaries [27, 28]. And local strain, which can be measured by KAM, also contributes to local deformation of material, so this parameters can influence micro-level deformation (i.e. SCC initiation or growth), rather than bulk material properties at room temperature. For this reason, these values will be considered as an important parameter in further SCC initiation testing experiments in a primary water environment.

CONCLUSIONS

Long-term thermal aging of nuclear power plant materials were simulated with high temperature heat treatment. Changes of microstructural and mechanical features of thermally aged Alloy 600 were analyzed. Area fraction and length of precipitates were significantly changed by thermal aging, leading to a significant change of the area fraction of precipitates. After simulated 10 year-thermal aging, the area fraction of precipitate was increased, while their average length did not change much. These changes could make more obstacles for dislocation movement and decrease the ductility of the


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material. After simulated 20 year-thermal aging, both the area fraction and average length of precipitate increased. These changes could make dislocations loop around precipitates and decrease the strength of alloy.

Several changes were observed in EBSD analysis such as kernel average misorientation and grain boundary misorientation, which are known to influence to stress corrosion cracking resistance of material. However, those properties are not related to the bulk material`s properties at room temperature. Therefore, these properties will be considered in future experiments; SCC initiation tests in primary water environment with prior thermal aging.

A process to simulate long-term thermal aging was established. 10 years thermally aged material showed the highest mechanical properties and softening occurred after longer durations. Further studies will be conducted to evaluate the resistance of stress corrosion cracking initiation of Alloy 600 with consideration of simulated thermal aging.

ACKNOWLEDGMENTS

This work was financially supported by the Nuclear Power Core Technology Development Program (No.20131520000140) and the Nuclear Safety Research Program through the Korea Radiation Safety Foundation (KORSAFe) and the Nuclear Safety and Security Commission (NSSC), Republic of Korea (Grant No. 1403006)

REFERENCES

[1] P. Scott, "An Overview of Internal Oxidation as a Possible Explanation of Intergranular Stress Corrosion Cracking of Alloy 600 in PWRs", Proceedings of the 9th International Symposium on Environmental Degradation of Materials in Nuclear Power Systems-Water Reactors pp.3-14 (1999)

[2] M. J. O. M. B. Toloczko, D. K. Schreiber, S. M. Bruemmer, "Corrosion and Stress Corrosion Crack Initiation of Cold-Worked Alloy 690 in PWR Primary Water", Technical Milestone Report (2013)

[3] P. Scott, "An Overview of Materials Degradation by Stress Corrosion in PWRs", (2004)

[4] J. Kai, G. Yu, C. Tsai, M. Liu and S. Yao, "The Effects of Heat Treatment on the Chromium Depletion, Precipitate Evolution, and Corrosion Resistance of Inconel Alloy 690", Metallurgical Transactions A 20, pp.2057-2067 (1989)

[5] R. Celin and F. Tehovnik, "Degradation of a Ni-Cr-Fe Alloy in a Pressurised-Water Nuclear Power Plant", Materiali in Tehnologije 45, pp.151-157 (2011)

[6] Y. S. Lim, H. P. Kim and S. S. Hwang, "Microstructural Characterization on Intergranular Stress Corrosion Cracking of Alloy 600 in PWR Primary Water Environment", Journal of Nuclear Materials 440, pp.46-54 (2013)

[7] K. J. Choi, J. J. Kim, B. H. Lee, C. B. Bahn and J. H. Kim, "Effects of Thermal Aging on Microstructures of Low Alloy Steel–Ni Base Alloy Dissimilar Metal Weld Interfaces", Journal of Nuclear Materials 441, pp.493-502 (2013)

[8] J. Boursier, F. Vaillant and B. Yrieix, "Weldability, Thermal Aging and PWSCC Behavior of Nickel Weld Metals Containing 15 to 30% Chromium", ASME/JSME 2004 Pressure Vessels and Piping Conference pp.109-121 (2004)

[9] T. Shoji, Z. Lu and H. Murakami, "Formulating Stress Corrosion Cracking Growth Rates by Combination of Crack Tip Mechanics and Crack Tip Oxidation Kinetics", Corrosion Science 52, pp.769-779 (2010)

[10] V. Mohles, "Simulations of Dislocation Glide in Overaged Precipitation-Hardened Crystals", Philosophical Magazine A 81, pp.971-990 (2001)


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[11] O. A. Ojo and M. C. Chaturvedi, "Liquation Microfissuring in the Weld Heat-Affected Zone of an Overaged Precipitation-Hardened Nickel-Base Superalloy", Metallurgical and Materials Transactions A 38, pp.356-369 (2007)

[12] S. Sajjadi and S. Zebarjad, "Effect of Temperature on Tensile Fracture Mechanisms of a Ni-Base Superalloy", Archives of Materials Science and Engineering 28, pp.34-40 (2007)

[13] A. Shyam and W. W. Milligan, "Effects of Deformation Behavior on Fatigue Fracture Surface Morphology in a Nickel-Base Superalloy", Acta Materialia 52, pp.1503-1513 (2004)

[14] Y. Zhao, I.-C. Choi, Y.-J. Kim and J.-i. Jang, "On the Nanomechanical Characteristics of Thermally-Treated Alloy 690: Grain Boundaries Versus Grain Interior", Journal of Alloys and Compounds 582, pp.141-145 (2014)

[15] Q. Peng, J. Hou, K. Sakaguchi, Y. Takeda and T. Shoji, "Effect of Dissolved Hydrogen on Corrosion of Inconel Alloy 600 in High Temperature Hydrogenated Water", Electrochimica Acta 56, pp.8375-8386 (2011)

[16] K. Mo, G. Lovicu, X. Chen, H.-M. Tung, J. B. Hansen and J. F. Stubbins, "Mechanism of Plastic Deformation of a Ni-Based Superalloy for VHTR Applications", Journal of Nuclear Materials 441, pp.695-703 (2013)

[17] J. Pardal, S. Tavares, V. Terra, M. Da Silva and D. Dos Santos, "Modeling of Precipitation Hardening During the Aging and Overaging of 18Ni–Co–Mo–Ti Maraging 300 Steel", Journal of alloys and compounds 393, pp.109-113 (2005)

[18] B. Alexandreanu, B. Capell and G. S. Was, "Combined Effect of Special Grain Boundaries and Grain Boundary Carbides on IGSCC of Ni-16Cr-9Fe-Xc Alloys", Materials Science and Engineering a-Structural Materials Properties Microstructure and Processing 300, pp.94-104 (2001)

[19] Z. Guo and W. Sha, "Quantification of Precipitation Hardening and Evolution of Precipitates", Materials Transactions 43, pp.1273-1282 (2002)

[20] V. Mohles, D. Rönnpagel and E. Nembach, "Simulation of Dislocation Glide in Precipitation Hardened Materials", Computational materials science 16, pp.144-150 (1999)

[21] D. Kirkwood, "Precipitate Number Density in a Ni Al Alloy at Early Stages of Ageing", Acta Metallurgica 18, pp.563-570 (1970)

[22] G. S. Was, “Grain-Boundary Chemistry and Intergranular Fracture in Austenitic Nickel-Base Alloys- A Review”, Corrosion 46, pp. 319-330, (1990)

[23] J. R. crum, K. A. Heck and T. M. Angeliu, “Effect of different thermal treatments on the corrosion resistance of alloy 690 tubing”, 4th Environmental Degradation of Materials in Nuclear Power Systems-Water Reactors, pp. 293 (1990)

[24] R. Hayes and W. Hayes, "On the Mechanism of Delayed Discontinuous Plastic Flow in an Age-Hardened Nickel Alloy", Acta Metallurgica 30, pp.1295-1301 (1982)

[25] B. S. Kumar, B. S. Prasad, V. Kain and J. Reddy, "Methods for Making Alloy 600 Resistant to Sensitization and Intergranular Corrosion", Corrosion Science 70, pp.55-61 (2013)

[26] D. G. Brandon, "The Structure of High-Angle Grain Boundaries", Acta Metallurgica 14, pp.1479-1484 (1966)

[27] L. Tan, K. Sridharan and T. R. Allen, "Effect of Thermomechanical Processing on Grain Boundary Character Distribution of a Ni-Based Superalloy", Journal of Nuclear Materials 371, pp.171-175 (2007)

[28] H. Grimmer, W. Bollmann and Warringt.Dh, "Coincidence-Site Lattices and Complete Pattern-Shift Lattices in Cubic-Crystals", Acta Crystallographica Section A A 30, pp.197-207 (1974)


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[29] J. Hou, Q. Peng, Y. Takeda, J. Kuniya and T. Shoji, "Microstructure and Stress Corrosion Cracking of the Fusion Boundary Region in an Alloy 182-A533b Low Alloy Steel Dissimilar Weld Joint", Corrosion Science 52, pp.3949-3954 (2010)


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Table 1. Chemical composition of Alloy 600

Element C Si Mn Cr Cu Ni S Fe

wt.% 0.07 0.33 0.56 15.83 0.02 74.79 <0.001 8.40


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Table 2. Specimen name, heat treatment and aging conditions of each specimen

Specimen name Simulated aging time and temperature Heat treatment time and temperature

As-received - -

HT400_Y10 10 years at 320 °C 1142 hours at 400 °C

HT400_Y20 20 years at 320 °C 2284 hours at 400 °C


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Table 3. Summary of mechanical properties achieved from tensile tests.

Specimen As-received HT400_Y10 HT400_Y20

Young`s modulus [GPa] 211.3 ± 12.0 290.7 ± 11.9 256.7 ± 46.4

0.2 % offset yield strength [MPa]

409.5 ± 10.9 587.3 ± 5.9 497.3 ± 2.2

Ultimate tensile strength [MPa] 701.8 ± 24.5 775.7 ± 24.7 767.7 ± 2.5

Elongation [%] 50.0 ± 1.6 36.0 ± 1.6 42.0 ± 3.1


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Table 4. Summary of nanoindentation hardness achieved from nanoindentation test.

Investigated location As-received HT400_Y10 HT400_Y20

Inside grain [GPa] 2.44 ± 0.03 2.75 ± 0.08 2.55 ± 0.09

Cover grain boundary [GPa] 2.5 ± 0.02 2.76 ± 0.05 2.6 ± 0.1

Average hardness [GPa] 2.47 ± 0.14 2.81 ± 0.24 2.59 ± 0.31


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Table 5. Summary of microstructural and orientation characteristics of specimens

Specimen As-received HT400_Y10 HT400_Y20

Area fraction of precipitates [%] 0.36 ± 0.04 1.2 ± 0.1 1.43 ± 0.14

Precipitate length [μm] 0.80 ± 0.30 0.92 ± 0.31 1.86 ± 1.19

Spacing between precipitates [μm] 2.78 ± 0.22 1.92 ± 0.14 3.28 ± 0.31

Grain size [μm] 26.67 ± 0.15 24.61 ± 0.21 25.39 ± 0.25

Kernel average misorientation [°] 0.653 ± 0.083 0.536 ± 0.074 0.701 ± 0.081

Fraction of CSL boundary [%] 25.5 ± 3.8 15.3 ± 4.88 6.8 ± 4.41


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unit [mm]

Figure 1. Proportionally reduced specimen from ASTM standard, used in tensile test.


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Figure 2. Representative stress-strain curves of each specimen.


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Figure 3. Fracture surfaces of (a, d) as-received, (b, e) HT400_Y10 and (c, f) HT400_Y20. (d), (e) and (f) are magnified view of (a), (b) and (c), respectively.


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Figure 4. Nanoindentation results of (a) as-received, (b) HT400_Y10 and (c) HT400_Y20.


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Figure 5. Optical microscope images of (a) as-received, (b) HT400_Y10 and (c) HT400_Y20.


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Figure 6. Secondary electron images of (a, d) as-received, (b, e) HT400_Y10 and (c, f) HT400_Y20. (d), (e) and (f) are magnified view of (a), (b) and (c), respectively.


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Figure 7. Material hardening and softening behavior depending on the precipitates.


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Figure 8. Grain boundary misorientation analysis results of (a) as-received, (b) HT400_Y10, and (c) HT400_Y20. Kernel average misorientation analysis results of (d) as-received, (e) HT400_Y10 and (f) HT400_Y20

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