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ISSN : 1598-6721(Print)
ISSN : 2288-0771(Online)
The Korean Society of Manufacturing Process Engineers Vol.18 No.1 pp.31-37
DOI : https://doi.org/10.14775/ksmpe.2019.18.1.031

Effects of Long-term Artificial-Aging on the Hardness Variation of Dissimilar Metal Weldments

Chung-Seok Kim*#
*Dep. of Materials Science and Engineering, Chosun UNIV.
Corresponding Author : chs2865@chosun.ac.kr Tel: +82-62-230-7197, Fax: +82-62-230-7197
23/10/2018 09/11/2018 14/11/2018

Abstract


This study investigates the effects of long-term artificial-aging on hardness variation in the dissimilar metal weldments for nuclear power plant facilities. These dissimilar welds are inevitably required to join the components in nozzle parts of pressurized vessels, such as austenitic stainless steels and ferritic steels. A artificial thermal aging was conducted in an electrical furnace to simulate material degradation at high temperatures. The test materialswere held at the temperature of 600°C for 10000 hours and interrupted at various levels of degraded specimens. The degradation of hardness is a well-known phenomenon resulting from long-term aging or high-temperature degradation of structural materials. In this study, the variation of hardness at each position was different, and complicated in relation to microstructures such as twins, grains, precipitates, phase transformations, and residual stresses in dissimilar weldments. We discussed the variation of hardness in terms of microstructural changes during long-term aging.



이종금속 용접부의 경도변화에 대한 장시간 인공열화의 영향

김 정석*#
*조선대학교 재료공학과

초록


    © The Korean Society of Manufacturing Process Engineers. All rights reserved.

    This is an Open-Access article distributed under the terms of the Creative Commons Attribution Non-Commercial License (http://creativecommons.org/licenses/by-nc/3.0) which permits unrestricted non-commercial use, distribution, and reproduction in any medium, provided the original work is properly cited.

    1. Introduction

    Dissimilar metal welds (DMW) is the joining of two different metals which is not ordinarily weldments as they have different thermal, chemical, and physical characteristics. DMW is welded together in order to maximize on the benefits that each metal produces while minimizing the drawbacks[1,2]. In general, austenitic stainless steel and ferritic carbon steel are a popular combination for nozzle components of the reactor pressurized vessels in nuclear power plants. The stainless steel should be connected to the austenitic steels or Ni-base alloy weldments. Also, the carbon steel is joined to the buttering materials firstly that is deposited a layer of weld on one surface to enhance a metallurgical compatibility for weld performance[3,4].

    Problems of dissimilar welds between the two materials (stainless steel and ferritic carbon steel) have been a common occurrence in the boiler industry. The primary causes of failures in dissimilar welds are carbon migration from the heat-affected zone (HAZ) of the ferritic carbon steel into the weld metal, expansion differences between the two varieties of steel, and the differences in corrosion resistance[5].

    Also, when these steels are exposed for a long period of time, they are generally subjected to softening of materials because of the coarsening of the intermetallic phases and the phase transformation[6]. In recent, there are many reports on the primary water stress corrosion cracking (PWSCC) at dissimilar metal weld to the nozzle component in pressurized vessels[7,8]. The major physical causes of this PWSCC failure is tensile residual stress on the heat affected zone (HAZ) generated by the welding processing. In addition, these dissimilar welding components are used under high-temperature and high-pressure within corrosion atmosphere. The mechanical and microstructural characterization is very important for the materials integrity and reliability in fossil and nuclear power facilities.

    However, there are not enough reports about the effect of the long-term artificial-aging on hardness variation in the dissimilar metal weldments. In this study, we investigate the long-term artificial aging of dissimilar metal weldments at elevated temperature and examine the thermal degradation of them.

    2. Experimental method

    2.1 Test materials

    The dissimilar metal weld joint was prepared by the gas tungsten arc welding (GTAW) process with the base materials, ferritic carbon steel (SA508) and austenitic stainless steel (AISI316L). The Ni-base alloy (alloy82) was used for the weld metal and the buttering material that is deposited a layer of weld on the surface of ferritic carbon steel to enhance a metallurgical compatibility for weld performance. Fig. 1 depicts the schematic diagram and macro-structure of dissimilar metal weld joint showing each part of weldments.

    The artificial thermal aging was achieved in an electrical furnace to simulate the material degradation at high temperature. The test materials are holding at the temperature of 600 for 100, 1000, 5000, and 10000 hours and interrupted at various level of degraded specimens. The test samples are prepared with a bar shape of the size of 20 mm width, 20 mm thickness, and 150 mm long. The residual stress distribution was reported previously through the continuous indentation technique[9]. The HAZ of SA508 and AISI316L prior to heat treatment showed a compressive residual stress as shown in Fig. 2. The HAZ of SA508 showed a tensile residual stress after heat treatment. But, the HAZ of AISI316L showed a compressive residual stress due to its high level of initial stress[10].

    2.2 Microstructure test

    The small-sized samples were prepared by a band saw and a low-speed diamond saw, and then mounted with cold-mounting resins. The microstructural observations were performed using an optical microscope (OM) and scanning electron microscope (SEM). The test specimen for microstructure was mechanically polished with SiC grinding papers up to grit 2400, then fine polished with Al2O3 powder solution of 1 mm. The Buehler's final include colloidal silica solution was used through the vibrating polishing machine. The chemical etching with Vilella’s reagent was conducted for the SA508 carbon steel and we use Aqua regia (30 ml of distilled water, 20 ml HCl and 15 ml HNO3) at room temperature during 120 seconds for austenitic AISI316L.

    2.3 Hardness test

    The hardness variation in the dissimilar metal weldments was measured to study the effects of the long-term artificial-aging on material degradation. The micro Vicker’s hardness was measured at each position through the HAZ and welds metals in line on the cross-section of weldments with a load of 1 kg and hold time of 10 seconds following the ASTM E 348[11]. The test was achieved ten times at each welding, and then, the average value was taken as the final hardness result.

    3. Experimental results and discussion

    3.1 Microstructure

    The surface microstructure of SA508 base metals was depicted in Fig. 3. The typical macrostructures of ferritic carbon steel are obviously different between before heat treatment and long-term aging for 10000 hours. The initial microstructure shows an upper bainite with well-developed laths. The fine precipitates are on the lath boundaries. These precipitates play a role in obstacles to move of dislocations and hardening mechanism of low carbon steel as shown in Fig. 3(a).

    However, as the aging time increased, these lath substructures are getting collapsed and also the precipitates on the lath boundaries are dissolved in the matrix as shown in Fig. 3(b). The initial bainite microstructure is finally changed to ferritic microstructure during long-term thermal aging. These microstructural changes possibly cause the softening of a matrix of SA508 base metals during long-term thermal aging

    The polygonal shape of grains and well-developed twins are the typical characteristics of the austenitic stainless steel (AISI316L).

    The surface microstructure of AISI316 base metals was shown in Fig. 4. The delta ferrite formed along the intergrain boundary of austenite stainless steels, which mainly depends on the chemical composition and cooling rate[12]. The amount of delta ferrite of AISI316L base metals can be estimated through the Schaeffler diagram that shows the limits of the austenitic, ferritic and martensitic phases in relation to the chromium and nickel equivalent[13]. The Fig. 4(a) shows the OM micrograph of the initial microstructure of AISI316L base metals.A plenty of annealing twins can be clearly shown in grain interior and delta ferrite along the grain boundaries. As the aging time increases, the annealing twins are disappeared and recovered. The polygonal-shape grains were changed to round-shape grains as shown in Fig. 4(b). The delta ferrite is not a stable phase in stainless steel AISI316L base metals, and thus it may transform to a more stable phase during long-term thermal aging. It is well-known from the literature that the delta ferrite might decompose into sigma phase after long-term exposure at high temperatures in Cr-Ni-steel[14]. The sigma phase forms at ferrite/austenite interfaces. A sigma phase is an intermolecular stage that causes metals to lose ductility, toughness, stability and corrosion resistance. Therefore, it is undesirable in physical and chemical properties for structural materials due to embrittled effect near ferrite/austenite interfaces and chromium depletion effect in the matrix[15]. The coarse phases of sigma are clearly observed along grain boundaries in Fig. 4(b).

    For more precise observation, we observed the surface microstructure through an electron scanning microscope (SEM). The fine precipitates along grain boundaries and lath subgrain boundaries of SA508 are clearly investigated as shown in Fig. 5(a). The bainite structure having well-developed lath substructures and boundary precipitates along lath subgrain boundaries are recovered and changed to a ferritic structure that depicts the recovery of lath subgrains and dissolution of boundary precipitates after the long-term aging time of 10000 hours as shown in Fig. 5(b).

    For the austenite stainless steel AISI316L base metal, the initial well developed twins and delta ferrite along grain boundaries were shown in Fig. 5(c). The annealing twins are nearly recovered and the delta ferrite is decomposed with sigma phase and chromium carbides (Cr23C6). Also, the boundary precipitation was developed as shown in Fig. 5(d).

    We measured the Vickers hardness at each welding region in dissimilar welds metals to investigate the variation of hardness with long-term aging time. The hardness of austenite AISI316L base metal region decreased a little, and also, for the HAZ of AISI316, hardness was dropped which caused by the release of residual stress after long-term thermal aging. The hardness of ferritic SA508 base metal decreased a lot with long-term aging because of the recovery of lath subgrain and dissolution of boundary precipitations as already shown in Fig. 5. In addition, the hardness of HAZ of SA508 region decreased because of the recovery of residual stress generated on the welding process as previously shown in Fig. 2. The microstructure in HAZ of SA508 region is complex and mixed microstructure such as fine martensite, bainite, coarse martensite, and carbides in the initial state after welding processing. This complex microstructure is recovered and transformed to ferrite phase having coarse carbides with long-term thermal aging. Therefore, the hardness decreased very fast in HAZ of SA508 region. On the contrary, the hardness of the buttering alloy82 region increased a lot continuously with long-term thermal aging. In this study, we used the alloy82 as a buttering metal that is deposited a layer of weld on the surface of the ferritic SA508 steel. The strength mechanism of this nickel-base superalloy (alloy82) is a precipitation hardening such as NbC and Cr23C6 intermetallic phases along grain boundaries and interior of grains[16]. Also, adiffusion of carbon atoms of SA508 HAZ part may happen and thus the hardness increased very fast and a lot with long-term thermal aging. It needs to study further for this physical phenomenon.

    4. Conclusion

    We investigated the effects of the long-term artificial-aging on hardness variation in the dissimilar metal weldments for nuclear power plant facilities. The microstructural characteristics in all welding regions have been investigated by the optical and electron scanning microscope. In this study, the variation of hardness at each position was different and complicate in relation with microstructures such as twins, grains, precipitates, phase transformation, and residual stress in dissimilar weldments. As the aging time increased, the lath substructures of ferritic SA508 are getting collapsed and also the precipitates on the lath boundaries are dissolved in the matrix. The initial bainite microstructure is finally changed to ferritic microstructure during long-term thermal aging. The annealing twins are disappeared and recovered. The polygonal-shape grains were changed to round-shape grains in austenitic AISI316L. The hardness of austenite AISI316L base metal and the HAZ of AISI316L decreased a little caused by the release of residual stress after long-term thermal aging. But, the hardness of ferritic SA508 base metal decreased a lot because of the recovery of lath subgrain and dissolution of boundary precipitations. The hardness of buttering alloy82 region increased a lot continuously due to precipitation hardening such as NbC and Cr23C6 intermetallic phases along grain boundaries and interior of grains. Consequently, we successfully investigated the long-term artificial-aging of dissimilar metal weldments by measuring Vickers hardness.

    Acknowledgements

    This research was supported by Basic Science Research Program through the National Research Foundation of Korea (NRF) funded by the Ministry of Education (2017R1D1A3B03028681).

    Figure

    KSMPE-18-31_F1.gif

    Macro micrograph and geometry of dissimilar metal weld joint; A: SA508, B: Buttering (Alloy82), C: Weldment (alloy82), D: AISI 316L

    KSMPE-18-31_F2.gif

    Residual stress distribution of dissimilar metal weld joint[6]

    KSMPE-18-31_F3.gif

    Optical micrographs of ferritic SA508 base metal microstructure; 0 h and (b) 10000 h

    KSMPE-18-31_F4.gif

    Optical micrographs of AISI316L base metal; 0 h and (b) 10000 h

    KSMPE-18-31_F5.gif

    Scanning electron micrographs of SA508 and AISI316L base metal microstructures; (a-b) SA508 0 h and (c-d) AISI316L 10000 h

    KSMPE-18-31_F6.gif

    Hardness variation with long-term aging time at each welding regions

    Table

    Reference

    1. Ogawa, T., Itatani, M., Saito, T., Hayashi, T., Narazaki, C. and Tsuchihashi, K., “Fracture assessment for a dissimilar metal weld of low alloy steel and Ni-base alloy,” International Journal of Pressure Vessels and Piping, Vol. 90–91, pp. 61-68, 2012.
    2. Mortezaie, A. and Shamanian, M., “An assessment of microstructure, mechanical properties and corrosion resistance of dissimilar welds between Inconel 718 and 310S austenitic stainless steel,” International Journal of Pressure Vessels and Piping, Vol. 116, pp. 37-46, 2014.
    3. Rathod, D. W., Singh, P. K., Pandey, S. and Aravindan, S., “Effect of buffer-layered buttering on microstructure and mechanical properties of dissimilar metal weld joints for nuclear plant application,” Materials Science and Engineering: A, Vol. 666, No. 1, pp. 100-113, 2016.
    4. Mvola, B., Kah, P. and Martikainen, J., “Dissimilar ferrous metal welding using advanced gas metal arc welding processes,” Reviews on Advanced Materials Science, Vol. 38, pp. 125-137, 2014.
    5. Al-Hajri, M., Malik, A. U., Meroufel, A. and Al-Muaili, F., “Premature failure of dissimilar weld joint at intermediate temperature superheater tube,” Case Studies in Engineering Failure Analysis, Vol. 3, pp. 96-103, 2015.
    6. Kim, C. S., “A Study of the Heat Treatment Effect on the Fatigue Crack Growth Behavior in Dissimilar Weld Metal Joints of SA508 Low-Carbon Steel and AISI316 Austenitic Stainless Steel,” Journal of the Korean Society of Manufacturing Process Engineers, Vol. 17, No. 3, pp. 16-21, 2018.
    7. Yeh, T. K., Huang, G. R., Wang, M. Y. and Tsai, C. H., “Stress corrosion cracking in dissimilar metal welds with 304L stainless steel and Alloy 82 in high temperature water,” Progress in Nuclear Energy, Vol. 63, pp. 7-11, 2013.
    8. Xu, H., Mahmoud, S., Nana, A. and Killian, D., “A new modeling method for natural PWSCC cracking simulation in a dissimilar metal weld,” Progress in Nuclear Energy, Vol. 116, pp. 20-26, 2014.
    9. Choi, H. M., Jung, Y. G. and Cho, Y. T., “Comparison of Hardness and Elastic Modulus of Polymer Thin Film using Nano Indentation Test,” Fall Conference of the Korean Society of Manufacturing Process Engineers, Vol. 1, pp. 84-84, 2017.
    10. Kim, C. S., “Optimal double heat treatment process to improve the mechancial properties of lightweight AlSiCu alloy,” Journal of the Korean Society of Manufacturing Process Engineers, Vol. 17, No. 3, pp. 102-108, 2018.
    11. Sasikala, G., Ray, S. K. and Mannan, S. L., “Kinetics of transformation of delta ferrite during creep in a type 316(N) stainless steel weld metal,” Materials. Science and Engineering A, Vol. 359, pp. 86-90, 2003.
    12. Chun, E. J., Baba, H., Nishimoto, K. and Saida, K., “Effect of Phosphorus on Solidification Characteristics in Austenitic Stainless Steels,” Material Characterization, Vol. 86, pp. 152-166, 2013.
    13. Guiraldenq, P. and Duparc, O. H., “The genesis of the Schaeffler diagram in the history of stainless steel,” Metallurgical Research Technology, Vol. 51, No. 4, pp. 613, 2017.
    14. Neidel, A. Fischer, B. Riesenbeck, S. and Cagliyan, E., “Transformation of Delta Ferrite Into Sigma Phase in Metastable Austenitic Stainless Steels After Long-Term High-Temperature Service Exposure,” Practical Metallography, Vol. 114, pp. 259-279, 2014.
    15. Chun, E. J., and Saida, K., “Prediction of sigma-phase embrittlement and its influence on repair weldability for type 316FR stainless steel weld metals with different solidification modes,” Journal of Nuclear Materials, Vol. 505, pp. 212-226, 2018.
    16. Sennour, M., Chaumun, E., Crépin, J., Duhamel, C., Gaslain, F., Guerre, C. and Curières, I., “TEM investigations on the effect of chromium content and of stress relief treatment on precipitation in Alloy 82,” Journal of Nuclear Materials, Vol. 442, pp. 262-269, 2013.