Study on the Microstructural Degradation of the Boiler Tubes for Coal Fired Power Plants

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1 Article KEPCO Journal on Electric Power and Energy Vol. 4, No. 1, June 2018 ISSN (Print), (Online) DOI : /KEPCO Study on the Microstructural Degradation of the Boiler Tubes for Coal Fired Power Plants Keun Bong Yoo 1, Yinsheng He 1, Han Sang Lee 1, Si Yeon Bae 1, Doo Soo Kim 1 Abstract A boiler system transforms water to pressured supercritical steam which drives the running of the turbine to rotate in the generator to produce electricity in power plants. Materials for building the tube system face challenges from high temperature creep damage, thermal fatigue/expansion, fireside and steam corrosion, etc. A database on the creep resistance strength and steam oxidation of the materials is important to the long term reliable operation of the boiler system. Generally, the ferritic steels, i.e., grade 1, grade 2, grade 9, and X20, are extensively used as the superheater (SH) and reheater (RH) in supercritical (SC) and ultra supercritcal (USC) power plants. Currently, advanced austenitic steel, such as TP347H (FG), Super304H and HR3C, are beginning to replace the traditional ferritic steels as they allow an increase in steam temperature to meet the demands for increased plant efficiency. The purpose of this paper is to provide the state of the art knowledge on boiler tube materials, including the strengthening, metallurgy, property/microstructural degradation, oxidation, and oxidation property improvement and then describe the modern microstructural characterization methods to assess and control the properties of these alloys. The paper covers the limited experience and experiment results with the alloys and presents important information on microstructural strengthening, degradation, and oxidation mechanisms. Keywords: Boiler Tubes, Strengthening, Degradation, Assessment, Microstructural Analysis I. GENERATION INTRODUCTION Materials selection is a key factor in the construction of cost effective and reliable boiler systems in power plants [1]. Fig. 1 schematically summarizes code approved boiler tube materials with the evolution of the steam conditions. The most effective way to increase the thermal efficiency of plants is to increase the steam temperature. Four types of power plant are classified as the subcritical (SC) and ultra supercritical (USC), and recently widely under construction of advanced USC (A USC), depending the steam condition. Increasing the steam parameter from 540 C/22.1 MPa to as high as 725 C/35 MPa can increase the thermal efficiency from 35% to 52%. Increasing the steam temperature mostly depends on developing materials with better creep strength and oxidation resistance. For subcritical power plants, the property requirement can be easily achieved by the low alloy steels, e.g., A210, T11, and more recently, the developed Grade 2 (T/P22 and 23). In the next generation of plants operating with SC steam parameters, 9 12%Cr martensitic steels, particularly Grade 9 (T/P91, 92) steels, are extensively used. USC plants adopt the improved 9 12%Cr ferritic steels and high alloyed austenitic TP347H steels. The 12%Cr X20CrMoV12.1 (X20) alloys and TP347H have been partially used for the final SH/RH and header. The A USC needs the steam temperature to reach 700 C, and currently most of the candidate materials are under field testing worldwide [2][3] : MARBN, Super304H, HR3C, and the serial 2, 6, 7 of Ni based alloys. In a modern boiler, for example the one shown in inset (a) of Fig. 1, selection of the right materials in consideration of the properties/tasks listed in 2~7 in inset b, is the critical mission for its design and construction [4][5]. Among the properties, the creep life is a key factor because the long term safe operation of the boiler must be guaranteed. An important task is the steam oxidation. Scale formed on the steam side is serious because it s decreases the heat transfer efficiency and may bring damage to the steam blade. Fireside corrosion is another challenge of the materials, and coatings have been used in some of the highest temperature parts. Another factor that must be taken into account is welding dissimilar welding components because failures occur in the heat affected zone (HAZ) of the welded parts. The fabricability, such as the cold working of bent tubes, is important because there are Manuscript received September 8, 2017, Accepted June 5, KEPCO Research Institute, Korea Electric Power Corporation, 105 Munji Ro Yusung Gu, Daejeon 34056, Korea yinsheng.he@kepco.co.kr, heyinsheng@gmail.com This paper is licensed under a Creative Commons Attribution NonCommercial NoDerivatives 4.0 International Public License. To view a copy of this license, visit nc nd/4.0. This paper and/or Supplementary information is available at 25

2 Keun Bong Yoo, et al.: Study on the Microstructural Degradation of the Boiler Tubes for Coal Fired Power Plants Fig. 3. Morphological, compositional and structural analysis of the microstructure using XRD, OM, SEM/EDS/EBSD and TEM. The general analysis area decreased while resolution increased from XRD to TEM. Fig. 1. Commonly used boiler tube materials as evolution of steam parameters in power plants. Insets (a) and (b) are the schematic images of the typical boiler and tasks for its construction. Fig. 2. Schematic diagrams showing the strengthening mechanisms of the 9 12 %Cr ferritic (a) and 18Cr8Ni austenitic (b) boiler tube materials. frequent failures in this area. II. BASIC STRENGTHENING MECHANISM The basic concept of strengthening of materials involves a solid solution, fine grain, precipitation and dislocation hardening as listed in Fig. 2. For ferritic materials, i.e., grade 11, 22 and grade 91, 92 as well as the X20, the composition of the alloy and the heat treatment process give the materials create a microstructure of a tempered martensite structure as showing in Fig. 2(a). Normally, the microstructure possesses a ferritic (α Fe) matrix with dispersion of substructures, including the prior austenite grain (PAG) with a size of ~50 μm and block boundaries dividing the PAG into several parts (block) with a size of ~10 μm. A larger number of lath structures are formed within the block with a size (lath width) of ~1 μm. Amounts of subgrains (~500 nm) are formed within the lath structure. The dislocation structure is densely dispersed within the subgrains and lath structure. Two kinds of the precipitates (PPTs) are dispersed in the microstructure. One is the Cr riched M23C6 with a size of ~200 nm, which mainly forms on the PAGB, block boundaries and lath boundaries. The other is the Nb riched MX with a size of ~50 nm, which mostly disperses in the grain interior. Therefore, the ferritic boiler tube materials are strengthened by the above four aspects, which the fine grain, PPTs and dislocations are the mainly strengthening elements. The MX PPTs blocking the motion of dislocation during creep is considered the most important strengthening mechanism [6]. Austenitic boiler tubes of TP347, Super304H and HR3C, basically contain a high alloy concentration of Cr and Ni, which give the solute solution strengthening. After heat treatment, as shown in Fig. 2(b), the PAG always has a size of ~60 μm. Only some annealing twin and NbC (~1 μm) PPTs as well as the dislocation structure were formed within the austenite matrix (γ Fe). Therefore, the strengthening mechanism is simpler than the ferritic materials, which are dominated by the solute solution. We monitored the properties involved to investigate the evolution of the microstructural parameters during aging, creep and service testing [7]. III. ASSESSMENT METHODS OF BOILER TUBES Sampling of the boiler tubes is the starting point of the assessment work. Laboratory testing and on site service are the two basic ways for the tube sampling and mechanical properties/structural characterization. Mechanical testing always includes creep rupture life, aging, and steam tests in the laboratory. The creep rupture test is a simulation of the accelerated environment of the tubes during service. It is possible to hold the same creep temperature as the on site service one, but the creep stress is much higher than the service one. For example, for the T/P91 boiler tube, when serviced at 550 C with a steam pressure of 20 MPa, the rupture life is more than 100 years, which is almost impossible with the laboratory test. It is hard to simulate the steam pressure in the laboratory, so most of the reported data is a kind of oxidation behavior upon steam at elevated temperature (500~800 C) with no pressure. A short term (~1,000 hours) oxidation test was aimed at the understanding of the oxidation mechanism, while a long term test is mostly to collect data on the thickness of the oxidation 26

3 KEPCO Journal on Electric Power and Energy, Vol. 4, No. 1, June 2018 scale with different steam temperatures and test durations. Another way is to collect the on site serviced tubes with known service history, such as different temperature/stress (SH, RH) with different durations to obtain the oxidation scale thickness during the services. All of the testing is to establish a database on the mechanical properties of these boiler tubes. Analysis of the boiler tubes includes non destructive tests, such as the replicas, and ultrasonic and destructive methods taken tubes from the power plant. Normally, to reveal the above microstructural strengthening aspects in terms of morphological, compositional and structural analysis, the investigation includes XRD (X ray diffractometry), OM (optical microscope), SEM (scanning electron microscopy), EBSD (electron backscatter diffraction) and TEM (transmission electron microscope). Fig. 3 schematically shows these modern analytical techniques. Briefly, the equipment is composed of a beam source, e.g., x ray for XRD, light for OM, and accelerated electrons for SEM and TEM, etc., and several detectors to detect the interaction of the beam with the prepared samples surfaces in XRD, OM, SEM/EBSD, while throughout the thin sample in TEM. The analysis concepts and sequence are from macro to nano, that is from XRD to TEM, during which the spatial resolution decreases from 500 mm 2 by XRD to ~100 μm 2 by TEM. In most cases, several analysis methods are combined to well understand the details of morphologies, composition and structures of the interested microstructures. In addition, Fig. 3 shows the specific analysis objectives and focus using each technique [8] [11]. The following sections describe a limited collection of the microstructural evolution of some typical boiler tubes alloys in terms of ferritic and austenitic alloy during service and creep test. IV. MICROSTRUCTURAL DEGRADATION A. Ferritic Steels Grade 2 materials are the bainitic alloys generally used in low temperature parts (<540 C) of the boiler, such as the waterwall or the inlet of SH/RH, which are not frequently reported to fail during long term operations. 1) Grade 9 Alloys Grade 9 alloys, including the T/P91 and T/P92, are the mostly widely used boiler tubes worldwide in power plants. T/P92 is a modified version of the 91 with the addition of 2 wt.% of W and reduction of 0.5 wt.% of Mo to increase the high temperature creep resistance. The microstructure consists of the tempered martensite matrix (α Fe) with the distribution of M23C6 and MX precipitates. Fig. 4(a1)~(a4) present a set of representative microstructures of T92 boiler tubes in the as received (AR) condition (before service). Fig. 4(a1), the EBSD image, well reveals the mentioned PAG, lath, block and subgrain structures. Fig. 4(a2) shows the SEM image of the etched surface, and various precipitates can be observed. Classification of the types of PPTs is carried out by the combined use of structural and compositional analysis in TEM observations of the carbon extraction replica samples, Fig. 4. EBSD, SEM, TEM micrographs of T92 tubes in the states of (a1, a4) AR, (b1~b4) crept ruptured at 600 C/130 MPa/4,970 hours and (c1~c3) operated as RH at a condition of 596 C/25 MPa/6,630 hours (c1~c3). where the PPTs are extracted on the carbon films. Fig. 4(a3) is a typical TEM image of the carbon extraction replication samples of the T92 AR specimen showing the domain M23C6 PPT with a size of ~150 nm and is mostly distributed in various grain boundaries. Within the grain interior, the tiny size (~30 nm) of MX (NbC) can be easily observed. Fig. 4(a4) is a TEM thin foil image of the T92 AR specimen, which also shows the M23C6 distribution in grain boundaries; however, it is difficult to observe the MX PPTs because they were overlapped by the dense dislocation structure in the matrix. With the combined use of the EBSD, SEM and TEM, all of the microstructural features are well characterized. In particular, we can see the grain (boundaries) by using the EBSD, while we see the overall PPT features by using SEM, and finally we see the details of the PPTs with the dislocation structures by using the TEM. The microstructural changes of T92 upon the laboratory creep test, for example, that ruptured at the condition of 600 C/130 MPa for 4,970 hours are shown in Fig. 4(b1)~(b4). It is clear that the lath and subgrain of the tempered martensite became equiaxed due to the dislocation rehabilitation/recovery as shown in the EBSD and TEM micrographs of Fig. 4(b1) and (b4). The domain PPTs M23C6 shows a slight coarsening after creep rupture [Fig. 4(b2), (b4)]. One of the most significant changes is the newly formed W riched Laves phase, which can be easily imaged by the backscattered image (BSI) of the SEM in Fig. 4(b3). The Laves phase always has a size of over 200 nm, in which the formation absorbed the solute W and Cr from the matrix. Additionally, the Laves phase has a high coarsening rate during creep, which was considered as the main reason for the microstructure and properties deteriorating. In Fig. 4(b4), there are tiny MX PPTs within the matrix because the matrix became clean due to the dislocation annihilation during the creep test. In the T92 alloy upon on site service as the RH tube with the conditions of 600 C/25 MPa for 6,630 hours, the microstructural evolution can be characterized as the formation of the Laves phase and even slight coarsening of the M23C6 as shown in Fig. 4(c1)~(c3). Quantitative analysis 27

4 Keun Bong Yoo, et al.: Study on the Microstructural Degradation of the Boiler Tubes for Coal Fired Power Plants Fig. 6. (a) OM image of the cracks formed in the HAZ of TP347H tube after operation at 596 C/25 MPa for about 2 years in a power plant. (b) SEM image of the AR TP347H. (c) SEM image and (d, e) TEM images near the crack part of (a). Fig. 5. EBSD, SEM, TEM micrographs of X20 tubes in the states of (a1, a4) AR, (b1~b4) crept ruptured at 550 C/180 MPa/6,886 hours and (c1~c4) operated as R/H at a condition of 541 C/25 MPa/112, 567 hours (c1~c3). of the PPTs for the average size and area fraction was completed with tens of BSE (back scatter electron) images for Laves and TEM thin foil images for M23C6 and MX and shows the highest coarsening rate of the Laves phase but extremely stable features of the MX. 2) X20 Alloy The X20 alloy is a 12Cr tempered martensitic ferritic steel and is used for the typical boiler tubes in Korea 500 MW standard coal fired power plants. The microstructure is similar to the Grade 9 boiler tubes with a tempered martensite matrix [Fig. 5(a1)] and the distribution of M23C6 and MX precipitates [Fig. 5(a2)~(a4)]. With short term creep in the conditions of 550 C/180 MPa for 6,886 hours, the tempered martensite [Fig. 5(b1)] and PPTs of M23C6 and MX [Fig. 5(b2) and (b4)] show no significant change, while formation of the newly Mo riched Laves phase [Fig. 5(b3)] is the mainly due to the microstructure degradation mechanism. When the X20 tubes are used as the SH with the steam parameters of 541 C/25 MPa for 112,567 hours, the microstructural evolution is characterized as the recovery of the tempered martensite [Fig. 5(c1)] and formation of the Laves phase [Fig. 5(c3)]. Even serviced for 12.8 years, the M23C6 is still very fine and only shows slight clustering along the grain boundary [Fig. 5(c)]. These experimental results indicate that the formation of the coarse Laves phase with a high growth rate is the main degradation mechanism of the ferritic steels upon creep and services [8][10]. B. Austenitic Steels The austenitic stainless steels are increasingly being used in new USC plant construction and also as replacement materials for traditional ferritic steels in boilers. The high 28 alloy concentration gives those materials a favorable property, which make them applicable in complex elevated temperature boiler environments involving fireside/ steamside corrosion and creep resistance. Indeed, TP347H (FG), Super304H and HR3C nowadays continue to be used as the SH and RH for boilers. 1) The SIPH in TP347H TP347H is an 18Cr 10NiNb austenitic steel, designed with the dispersion of primary NbC to strengthen the creep resistance [Fig. 6(b)]. After service, failures are often found in the HAZ of the welding components. For example, in a TP347H tube, operating at the steam conditions of 596 C/25 MPa for about two years, a crack in the HAZ began from the steamside and permeated to the fireside as shown in Fig. 6(a) [12]. The microstructural investigation of the failed part revealed the formation of coarse sized Cr riched PPTs along the GB [Fig. 6(c)] but only scattered dispersion of the NbC in the GB in AR as shown in Fig. 6(b). These GB Cr riched PPTs were confirmed to be M23C6 as shown in Fig. 6(d), which is a TEM image of the HAZ of the failed tube. In the grain interior, a larger fraction of the NbC with a size of ~30 nm was formed along the dislocations [Fig. 6(e)], which is called strain induced precipitation hardening (SIPH). Though it could be due to reasons like stress accumulation in the HAZ, we supposed that the reason for the failure was the formation of coarse M23C6 along the GB by consuming Cr to form a Cr depletion zone and finally weaken the boundaries, while the formation of nano sized NbC within the grain strengthened the grain interior. The inhomogeneous microstructure eventually resulted in brittle tubes. The methods for extending the service life of the TP347H HAZ aim to delay the formation and retard the coarsening of the M23C6 on GB by heat treatment, including the solution treatment (ST) and post weld heat treatment (PWHT). The idea is to stabilize the C by the ST before the welding process, which will reduce its content to form M23C6 during afterward

5 KEPCO Journal on Electric Power and Energy, Vol. 4, No. 1, June 2018 Fig. 7. TEM micrographs of the TP347 in the state of (a) AR, (b) ST at 1120 C/30 min and (c) ST+PHWT at 900 C/2 hours. SEM micrographs of the (d1) AR, (e1) ST, (f1) PWHT, (g1) ST+PWHT and aging tested at 600 C/1,000 hours of (d2) AR1000, (e2) ST1000, (f2) PWHT1000, (g2) ST+PWHT1000. high temperature exposure. In addition, the PWHT process was designed to release the dislocations and further stabilize the C after the welding process. The AR TP347H contained the primary NbC scattered dispersed within the grain interior as shown in Fig. 7(a). After the ST at a temperature of 1,120 C for 30 min, a larger number of NbC particles with a size of ~30 nm was formed within the grain interior [Fig. 7(b)], indicating that the C atoms were partially stabilized. Then, after two pass welding, PWHT was carried out at a temperature of 900 C for 2 hours, and the microstructure is shown in Fig. 7(c), and the clear result was that the dislocation density decreased with a slight increase in the fraction of NbC, just as expected. Finally, the welded tubes with the heat treatment of as received (AR), ST, PWHT of the AR, and PWHT of the ST (ST+PWHT) were aged at 600 C up to 1,000 hours, and the microstructures are shown in the SEM images of Fig. 7(d1)~(d2), (e1)~(e2), (f1)~(f2) and (g1)~(g2), respectively. It can be seen that the GBs were clean in the specimens of AR, ST, PWHT and ST+PWHT in Fig. 7(d1), (e1), (f1) and (g1), respectively, suggesting the free M23C6 before the aging test. After the aging test, one may find that the fraction and size of the M23C6 decreased in the order of AR1000 > ST1000 > PWHT > ST+PWHT1000, indicating that the ST and PWHT retarded the formation of M23C6 during aging, and PWHT had a better effect, while ST with PWHT gave impressive results. Therefore, we suggest that the ST and PWHT of the austenitic boiler tubes are essential processes for improving the service life of weld components. 2) The Sigma Phase in Super304H Alloys Super304H is a 18Cr9NiCuNbN alloy that is an enhancement of TP347HFG with the addition of Cu and N for the improvement of creep resistance and beneficial oxidation properties. The microstructure of the AR S304H is similar to the TP347H and it has dispersion of coarse primary NbC within an austenite matrix (γ Fe) as shown in Fig. 8(a1) and Fig. 8(a2). After aging at a temperature above 600 C, for example, at 700 C for 20,000 hours, as shown in Fig. 8(b1) and Fig. 8(b2), the microstructure changes are the formation of a Cr riched σ phase (~10 μm) and M23C6 (~5 μm) in the Fig. 8. SEM and TEM micrographs of the S304H tube in the AR state (a1, a2), and after aging tested at 700 C/20,000 hours (b1, b2). The inset table in b1 is the typical composition of the newly formed M23C6 and phases. Fig. 9. (a1~a4) SEM/EDS compositional maps and (a5) EBSD phase map micrographs of the oxide scale formed on the T91 after steam test at 650 C/10,000 hours. (b) EBSD micrograph of scale formed on the T91 after steam test at 650 C/15,000 hours. (c1~c4) and (d1~d4) SEM/EDS compositional maps of the oxide formed on S304H after steam test at 650 C/10,000 hours without and with inner surface shot peening treatment, respectively. (e) Schematic images of the oxide phase distribution in the scale formed on the S304H tubes. GBs [Fig. 8(b1)]. The nucleation and coarsening of these PPTs consumed the solute elements like Cr and Ni from the adjacent matrix. In addition, within the grain interior, instead of SIP in TP347H, amounts of Cu particles with a size of ~20 nm were nucleated [Fig. 8(b2)]. The Cu was in the solute state in the AR condition, with clusters that formed Cu particles when aged at a high temperature, which enhanced the hardness and creep resistance. An aging test of the S304H steels at a temperature range of 600~700 C with different durations revealed that the σ phase and M23C6 had a high coarsening rate, while the Cu was relatively stable as the time elapsed. Currently, the short term creep behavior of the S304H is under testing, and the properties and microstructural degradation will be reported later. V. STEAM OXIDATION BEHAVIORS AND ENHANCEMENT Oxidation of tubes is a significant challenge to the safe operation of boilers. Issues involved in the oxidation of boiler tubes include the following: 1) collecting data on the oxide scale thickness to determine the overhaul time needed to 29

6 Keun Bong Yoo, et al.: Study on the Microstructural Degradation of the Boiler Tubes for Coal Fired Power Plants remove of the scale, 2) study the oxide scale features including the oxide layer distribution underlying the total scale to understand the oxidation resistant layer, 3) selection cost effective ways to increase the steam oxidation resistance. [13] [15] Fig. 9 presents a set of SEM and EBSD micrographs showing the morphology, composition and structure of the typical oxide scale formed on the T91 tube after a steam test at 650 C for 10,000 hours. Generally, the cracks parallel to the scale/metal interface are presented inside the scale [Fig. 9(a1)]. An Fe O enriched [Fig. 9(a2) and (a3)] outer layer and a O Cr enriched [Fig. 9(a3) and (a4)] inner layer are observed above and beyond the cracks, respectively. The oxide phase distribution was identified to be in the sequence of Fe2O3/Fe3O4/Cracks/FeCr2O4/metal from the steam/scale interface to the scale/metal interface as clearly shown in the EBSD phase map of Fig. 9(a5). The Fe2O3 and Fe3O4 grains in the outer layer were much coarser than the FeCr2O4 in the inner layer. The FeCr2O4 was at nano size and densely dispersed. With the elapse of time, the outer layer was subject to flaking off along the interface cracks as shown in Fig. 9(b), which is an EBSD phase map of the scale on T91 after a test at 650 C for 15,000 hours. In addition, with the growth of the inner layer, pores and cracks formed within the inner FeCr2O4 layer, which may also the inner layer to partially flake off. With a further increase in the test duration, the dropped oxide layer will re form, that is, recovery, and cause the scale owned a same structure as in Fig. 9(a5). Indeed, the evolution of the scale is the repeat process of the scale formation flake off recovery. The nano sized and dense FeCr2O4 inner layer isolated the metal from the steam oxidation environment, which is supposed as the oxidation protection layer. In addition, in case of the T91 tested at 650 C, α grains near the scale/metal interface with a thickness of ~20 μm are observed to be grown from their original size of ~400 nm to ~10 μm, which is known as grain recrystallization. The oxidation resistance decreases dramatically due to this grain recovered layer. Fig. 9(c1)~(c4) is a set of SEM/EDS micrographs showing the morphology and composition of the scale formed on a S304H boiler tube after it was steam tested at 650 C for 10,000 hours. Clearly, the scale thickness is much thinner and more compact (free of inside cracks) than the T91 one. The scale consisted of an outer layer of enriched Fe O [Fig. 9(c2), (c3)] and an inner O Cr enriched layer. Detailed analysis using TEM of the inner O enriched layer suggested that it was composed of nano sized Cr2O3 and FeCr2O4 oxide. Closer to the scale/metal interface, a discontinuous amorphous SiO2 layer with a thickness ~30 nm was formed as schematically shown in Fig. 9(e). The investigation of the evolution of the oxide layer thickness according to test duration found that the inner layer, particularly the Cr2O3 and SiO2, was stable, suggesting the critical role of these oxide layers on the oxidation resistances. The idea to enhance the oxidation resistance is to develop more Cr2O3 and SiO2 [16]. Fig. 9(d1)~(d4) are a set of SEM/EDS micrographs showing scale formed on the inner surface shot peened S304H tube after a steam oxidation test at 650 C for 10,000 hours. Obviously, the thickness is around ~4 μm which is much thinner than the un peened one, indicating that the SP enhanced the oxidation properties effectively. The scale also contained an outer Fe O enriched layer [Fig. 9(c2), (c3)] and an inner O Cr enriched layer, which is mostly composed of nano sized Cr2O3 and a continuously amorphous SiO2 layer. Our even longer steam test of the steamside shot peened tubes for 25,000 hours found that the oxide layer was still stable. Therefore, we concluded that once the austenitic boiler tubes have been shot peened, the oxide scale problem will not have an effect on the service of the tubes throughout the whole designed service life. VI. SUMMARY AND CONCLUSIONS Long term safe operation of a boiler highly depends on the microstructural stabilities of the tube materials. The domain degradation of boiler tubes is considered as the formation of a coarse new phase and easily flaked off oxide layers. In the ferritic steel of T91, T92 and X20, formation of a high growth rate (Mo, W) riched Laves phase along the PAG, block and lath boundaries are characterized as the main degradation mechanism. The coarsening of the main strengthening precipitates of M23C6 is not significant upon creep test and service below the 600 C. The oxidation scale contained the outer (Fe2O3 and Fe3O4) and inner (FeCr2O4) layers, where the outer layer is easy to flake off along the scale interface cracks. The degradation of the scale is the flake off of the outer layer and its re formation, which consumed the metal. In the austenitic steel of TP347, formation of coarse Crriched M23C6 on the GB and nano sized NbC (SIP) within the grain interior were found to be the main degradation mechanism. This phenomenon was even significant in the HAZ of the welding components, which can be effectively modified by ST and PWHT processes. In S304H steel, formation of coarse Cr riched M23C6 and a σ phase on GB were the main degradation mechanisms. Steam oxidation properties of the austenitic steel is much better than the ferritic steels due to the high Cr and Ni alloy concentrations, which developed dense and nano sized Cr2O3 oxides that covered the metal surface and prevented the fast oxidation of the metal with high temperature steam. Inner surface shot peening treatment was an effective way to enhance the oxidation resistance because it generated a larger fraction of GB that guaranteed the formation of dense Cr2O3 and a continuous amorphous SiO2 protective layer. This paper covered limited experiences and experimental results on the microstructural degradation of the boiler tubes. Indeed, studies are being carried out to determine the evolution of the detailed microstructure from the quantitative data on the grains, boundaries, PPTs, scale and oxide phase layer thickness. It is essential to establish a method to predict the evolution of GB, PPTs and oxide scale at elevated temperatures using modern microstructural characterization. These efforts will contribute to developing materials with the best combination of cost effectiveness and long term reliability for modern power plants. 30

7 KEPCO Journal on Electric Power and Energy, Vol. 4, No. 1, June 2018 ACKNOWLEDGEMENT This research was supported by Energy Technology Development Program of Korea Institute of Energy Technology Evaluation and Planning (KETEP) grant funded by the Korea government s Ministry of Trade, Industry and Energy ( ). REFERENCES [1] Abe, F., Progress in Creep Resistant Steels for High Efficiency Coal Fired Power Plants, Journal of Pressure Vessel Technology, Vol. 138, 2016, pp [2] Shingledecker, J., The Advanced Austenitic Stainless Steels Handbook: 347HFG, SUPER304H, HR3C., EPRI, [3] Abe, F., Research and Development of Heat Resistant Materials for Advanced USC Power Plants with Steam Temperatures of 700 C and Above. Engineering, Vol. 1, pp [4] Eggler, G, The effect of long term creep on particle coarsening in tempered martensitic ferritic steels, Acta Metallurgica, vol, 37, 1989, pp [5] Milović, L., et al., Microstructures and mechanical properties of creep resistant steel for application at elevated temperatures, Materials & Design, Vol, 46, 2013, pp [6] Hald, J., Microstructure and long term creep properties of 9 12% Cr steels, International Journal of Pressure Vessels and Piping, Vol. 85, 2008, pp [7] Klotz, U.E., C. Solenthaler, and P.J. Uggowitzer, Martensitic austenitic 9 12% Cr steels Alloy design, microstructural stability and mechanical properties. Materials Science and Engineering: A, Vol , pp [8] Y.S. He, J.C. Chang, J.L. Dong and K. Shin, Microstructural Evolution of X20CrMoV12.1 Steel Upon Long Term On Site Exposure in Power Plants, Advanced Science Letters, Vol. 4, 2011, pp [9] Y.S. He, J.L. Dong, J.S. Juang and K. Shin, An Improved nondestructive replication metallography method for investigation of the precipitates in Cr Mo V turbine steel, Surface and Interface Analysis, Vol. 44, 2012, pp [10] Hino, M, Y.S. He, K.J. Li, J.C. Chang and K. Shin, Microstructural Evolution of X20CrMoV12.1 Steel upon Short term Creep Rupture Test, Applied Microscopy, Vol. 43, 2013, pp [11] Y.S. He, J.C. Chang, J.H. Lee and K. Shin, Microstructural Evolution of Grade 91 Steel upon Heating at 760~1000 C. Korean Journal of Materials Research, Vol. 25, Pp [12] H.S. Lee, J.S. Jung, D.S. Kim, K.B. Yoo, Failure Analysis on welded joints of 347H austenitic boiler tubes, Engineering Failuer Analysis, Vol.57, 2015, pp [13] Fry, A., S. Osgerby, and M. Wright, Oxidation of Alloys in Steam Environments A Review. NPL Report. MATC (A) 90, [14] Tan, L., X. Ren, and T.R. Allen, Corrosion behavior of 9 12% Cr ferritic martensitic steels in supercritical water. Corrosion Science, Vol. 52, 2010, pp [15] Xia, Z.X., et al., Improve oxidation resistance at high temperature by nanocrystalline surface layer. Scientific Reports, 13027, [16] Hoelzer, D.T., B.A. Pint, and I.G. Wright, A microstructural study of the oxide scale formation on ODS Fe 13Cr steel. Journal of Nuclear Materials, Vol , 2000, pp