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Thermal simulation on double-pass welding of a high Cr ferritic steel
Journal of Manufacturing Processes 43 (2019) 9–16
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Thermal simulation on double-pass welding of a high Cr ferritic steel a
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b
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T a
Yi Shao , Biyu Yan , Yihuan Liu , Chunliang Mao , Chen Wei , Yongchang Liu , Zesheng Yan , ⁎ Huijun Lia, Chenxi Liua, a b
State Key Lab of Hydraulic Engineering Simulation and Safety, School of Materials Science & Engineering, Tianjin University, Tianjin, 300354, PR China Tianjin Special Equipment Inspection Institute, Tianjin, 300192, PR China
A R T I C LE I N FO
A B S T R A C T
Keywords: Welding thermal simulation Multi-Pass welding High Cr ferritic steel Creep life
The heat affected zone (HAZ) after double-pass welding in a high Cr ferritic steel was simulated by using a thermomechanical simulator. The microstructural observation and mechanical properties tests were conducted to investigate the effect of double-pass thermal cycle on the microstructure evolution and mechanical properties. The results show that the double-pass thermal cycle leads to the decrease of martensitic fraction, and the refinement of prior austenite. M23C6 precipitates is coarsened, and dissolved subsequently, by increasing the peak temperature during the second-pass thermal cycle, while MX with the higher thermal stability is undissolved until at 1200 °C. Thermal cycle at low temperature (as 900 °C) deteriorates the impact toughness, due to the high δ-ferrite fraction. Short-term creep tests suggest that, only the second-pass thermal cycle temperate of 1200 °C could eliminate the unfavorable effect on creep life resulted from the first-pass thermal cycle. Coarse M23C6 particles and the heterogeneous alloy elements distribution maybe the main reason.
1. Introduction High Cr ferritic steels have been the important structural materials for the main steam pipe and heater with larger diameter and thick wall in ultra-super critical (USC) fossil power plants, due to their high creep rupture strength, low thermal expansion coefficient, good corrosion resistance, and high performance–cost ratio [1–4]. In addition, considering their outstanding resistance to radiation induced void swelling and irradiation embrittlement, the high Cr ferritic steels have become the candidate materials used for nuclear power plants and test blanket modules (TBMs) of International Thermonuclear Experimental Reactor (ITER) [5–7]. Nevertheless, it was found that the creep life of the welding joints in the high Cr ferritic steels are significantly lower than that of the base metal [8]. It has been recognized that the creep cracks usually occur in the fine-grained heat affected zone (FGHAZ) and the intercritical heat affected zone (ICHAZ) of the welds [9]. During welding thermal cycle, FGHAZ is heated just above Ac3, and ICHAZ is heated between Ac1 and Ac3. These creep damage behaviors are called as type IV cracking [10]. Type IV cracking is likely to occur under lower stress and higher temperature conditions [11]. Type IV cracking predominates when the rupture stress is less than about 100MPa [12]. In FGHAZ, the recovery of martensite during creep is heterogeneous so that some coarse subgrains form near the fine subgrains [8]. Moreover, the coarsening of
⁎
M23C6 and the recovery of dislocations are significant in FGHAZ and ICHAZ. The dislocation density in FGHAZ and ICHAZ is significantly lower than the weld metal and base metal [11]. The Creep cavities nucleate preferentially around the coarse M23C6 particles [13]. Besides, the coarse Laves phase forming during creep would also contribute to the formation of the creep cavities [14]. The consumption of solutionstrengthening elements as W, Cr during the formation of Laves phase would also be the key factor for the occurrence of type IV cracking [15]. The post weld heat treatment can eliminate the type IV cracking, though it rarely is a practical option in the actual production [16]. The weld process and parameters also affect the type IV cracking behaviors. It was found that the creep rupture lives of laser weldment are higher than that of shielded metal arc welding (SMAW), since the width of HAZ in the former is narrower than that in the latter [17]. The joint geometry and preheat temperature can ameliorate the type IV failures, while the effect of heat input is less significant [18]. Generally speaking, the width of the heat affected zone (HAZ) in welding joints is only serval millimeters, so that the mechanical tests are difficult to be carried out for each region in HAZ (such as FGHAZ and ICHAZ). Hence, the thermal simulation method is used to “magnify” HAZ, and thus facilitates the mechanical tests, and the further microstructural analysis (as TEM), for each region in HAZ [13]. This method is to simulate the HAZ microstructure by set a thermal cycle that is similar to a certain region of the HAZ which is undergone during
Corresponding author. E-mail address: [email protected] (C. Liu).
https://doi.org/10.1016/j.jmapro.2019.05.012 Received 19 January 2019; Received in revised form 14 April 2019; Accepted 9 May 2019 Available online 15 May 2019 1526-6125/ © 2019 The Society of Manufacturing Engineers. Published by Elsevier Ltd. All rights reserved.
Journal of Manufacturing Processes 43 (2019) 9–16
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up to 900 °C with a rate of 50 °C/s and held for 1 s, followed by cooled to 200 °C with a rate of 25 °C/s. For the second-pass thermal cycle, the samples after the first-pass thermal cycle were reheated to 900, 1000, 1100, 1200 °C with a rate of 50 °C/s, respectively, and held for 1 s, followed by cooled to room temperature with a rate of 25 °C/s. The samples undergoing double-pass thermal cycle are denoted as “900& 900 °C”, “900&1000 °C”, “900&1100 °C” and “900&1200 °C” in terms of their different peak temperatures during the second-pass thermal cycle. The schematic diagram and the summary table of the single- or doublepass thermal cycle procedures is presented in Fig. 2 and Table 4, respectively. After the thermal simulation experiments, the samples were mounted, polished, and etched in solution of nitric acid and ethyl alcohol for metallographic observation. Extraction carbon replicas specimen were prepared to conduct the transmission electron microscopy. For preparation of the replicas, the samples after metallographic observation were etched in a solution of ethanol (90 ml), hydrochloric acid (10 ml) and picric acid (1 g). Then the etched surfaces of the samples were sprayed by carbon nanoparticles in vacuum. After spraying carbon, the scratched carbon films were floated to the surface using a solution of ethanol (70 ml), hydrochloric acid (20 ml) and nitric acid (10 ml), and fished out by the copper grids. The microstructural observation was carried out on Olympus C-35A Optical Microscope (OM), and FEI Tecnai G2 Transmission Electron Microscope (TEM). Due to the size limit of the thermal simulation samples (length for 70 mm and diameter for 6 mm), the subsized specimens for creep and impact tests were used. The creep tests were conducted on RPL100 electron creep testing machine with the subsized samples (φ6 × 63 mm). The creep test parameters are that the creep temperature of 550 °C and the loading stress of 180 MPa. The Charpy impact tests were carried out at room temperature on CBD 500 Charpy impact test machine with the subsized V-notch samples (55 × 2.5 × 2.5 mm). The impact tests were conducted three times for each sample.
the actual welding process. By using this method, it has been found that the high density of the grain boundaries, coarse precipitates, and the homogeneous grain and compositions maybe the key factors for type IV cracking [19–21]. The above research results only focus on the single-pass welding. However, in the actual welding, multi-pass welding is often used to welding the thick plate or pipe, due to the limitation of weld penetration for fusion welding. The multi-pass welding results in the course of thermal cycle in HAZ being more complicated. Depending on factors such as heat input, electrode size and angle and degree of overlap, the different regions in HAZ are exposed to the various multiple thermal cycles [15]. Heterogeneous structure of WM, due to reheating during multi-pass and multi-layer welding, has complex effects on the property of welded joint [22]. So far, the welding thermal simulation for the double-pass (or multi-pass) welding in the high Cr ferritic steels have not been reported. The effect of double-pass (or multi-pass) welding on type IV cracking was not clarified. Therefore, the aim of this work is to investigate the effect of double-pass welding on microstructure evolution and mechanical properties of the high Cr ferritic steel by using the thermal simulation method. According to the previous studies, thermal simulation mainly focuses on FGHAZ and ICHAZ in the welded joint of the high Cr ferritic steels, since these two regions is the weakness of components during their service. Thus, the first-pass thermal cycle temperature was set to be close to the peak temperature for FGHAZ and ICHAZ in this work. The range of the second-pass thermal cycle temperatures were selected to wider. After the welding thermal simulation, the microstructures of the samples were observed, and analyzed quantitatively. The impact toughness and short-term creep time was tested to investigate the effect of microstructure morphology, especially as distribution of the precipitates, on the mechanical properties. 2. Experimental details The employed steel in this project is a high Cr ferritic steel used for nuclear fusion reactor components, whose chemical composition is given by Table 1. Before the thermal simulation experiments, the raw material (as-rolled state) was normalized at 1050 °C for 0.5 h and then tempered at 750 °C for 1.5 h. The microstructure of the initial samples before and after normalizing & tempering is presented in Fig. 1. Martensitic laths and δ-ferrite were found in the initial sample, since low carbon content (0.4 in wt.%, as seen in Table 1) expands the phase region of ferrite in the equilibrium phase diagram [23]. The size of the prior austenitic grain of the sample after N&T is measured as 32.4 μm. Although normalizing & tempering treatment results in decrease of δferrite amount, considerable δ-ferrite distributed along the direction of rolling is still retained. After heat treatments, the samples were cut into the shape of cylinder with a length of 70 mm and a diameter of 6 mm by the wire cutting machine. The thermal simulation experiments were conducted on Gleeble3500 thermomechanical simulator with the Rykalin-3D programme. The detailed information about the Rykalin-3D programme on Gleeble3500 thermomechanical simulator can be seen in Ref. [24]. The required physical parameters and welding process parameters used for the Rykalin-3D procedure are shown in Tables 2 and 3, respectively. The peak temperature of the first thermal cycle was set as 900 °C, and that of the second one was set as 900, 1000, 1100, 1200 °C, to simulate the fine-grained region upon the double-pass welding process. According to the Rykalin-3D procedure, for the first-pass thermal cycle, the samples are preheated to 250 °C with a rate of 78 °C/s, then heated
3. Results 3.1. Microstructural observation Fig. 3 shows the OM images of the samples after single-pass thermal cycle with the peak temperature of 900 °C, and after double-pass thermal cycle with the peak temperature of 900&900 °C, 900&1000 °C, 900&1100 °C and 900&1200 °C. It is found that the microstructures of all samples, whether upon single-pass or double-pass thermal cycle, are composed of lath martensite and ferrite, which is similar to that of the initial sample (see Fig. 1). However, the phase fraction of martensite and ferrite, and the grain size of prior austenite and ferrite is strongly affected by the thermal cycle process. The TEM images of the extraction replicas samples upon singe-pass or double-pass thermal cycle are shown in Fig. 4. The extraction replicas only reserve the second-phase particles, losing the information of the matrix as grain boundary and dislocation. It can be found that almost all samples have the coarse and rod-liked M23C6 particles and the fine MX precipitates, although the number density and average size of the particles may different with the thermal cycle temperature changing. However, M23C6 particles are absent in the sample for 900& 1100 °C, and no precipitate is found in the sample for 900&1200 °C. This suggests that MX precipitate have the higher thermal stability than M23C6. The latter is dissolved at 1100 °C, whereas the former is stable until heating up to 1200 °C. 3.2. Mechanical properties
Table 1 The Chemical compositions of experimental steel (wt.%). C
Cr
W
Mn
Si
V
Ta
Fe
0.05
8.9
1.7
0.45
< 0.1%
0.20
0.15
Bal
To evaluate the effect of single-pass and double-pass welding thermal cycle on the mechanical properties in the employed steel, Charpy impact and short-term creep testing were conducted. Fig. 5 gives the impact energy at room temperature in the samples after 10
Journal of Manufacturing Processes 43 (2019) 9–16
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Fig. 1. Microstructure of the initial samples: (a) before and (b) after normalizing & tempering treatment. The size of the prior austenitic grain of the sample after N&T is measured as 32.4 μm. Table 2 The physical parameters of the experimental steel. Density 7.8 g/cm
Specific heat 3
o
0.58 J/(g· C)
Thermal conductivity 0.353 J/(cm·s·oC)
single-pass or double-pass (with the different second-pass peak temperatures) thermal cycle. It is found that thermal cycle at 900 °C deteriorates the impact toughness. Furthermore, double-pass thermal cycle for 900&900 °C would aggravate this degradation. The sample after double-pass thermal cycle for 900&1000 °C have the highest impact energy. By further increasing the second-pass thermal cycle temperature, the impact energy is decreased. For the high Cr ferritic steels used in nuclear and coal power stations, the creep properties are the most important performance indicators. The creep rupture time, under the condition of 550 °C / 180 MPa, of the samples after single-pass or double-pass (with the different second-pass peak temperatures) thermal cycle is presented in Fig. 6. It is obviously found that thermal cycle at 900 & 1200 °C leads to the longest creep rupture time (up to 1614 h), whereas other samples have the very short rupture time (less than 350 h).
Fig. 2. The schematic diagram of the double-pass thermal cycle procedures. Table 4 The thermal cycles parameters for the single-pass or double-pass thermal cycle samples. Thermal cycle times
Single-pass
Double-pass
Thermal cycle procedures
900 °C
900& 900 °C
900& 1000 °C
900& 1100 °C
900& 1200 °C
4. Discussion Nevertheless, by comparing Figs.1b with 3 a, it is found that the banded distribution feature of ferrite in the initial sample after normalizing & tempering is not obvious in the sample after single-pass thermal cycle. Therefore, it can be inferred that recrystallization of martensite distributes randomly and thus disorders apparently the banded characteristic of the already existed δ-ferrite before thermal cycle. As a result, the amount of ferrite is increased by thermal cycle at 900 °C. Likewise, double-pass thermal cycle for 900&900 °C would promote recrystallization of martensite further, and increase the amount of ferrite accordingly. As the second-pass thermal cycle temperature continues to increase to 1000 °C, which is higher than Ac3 point, martensite transforms to austenite completely. Moreover, the phase fraction of ferrite in the sample for 900&1000 °C is lower than that in the initial sample after normalizing & tempering (see the red dashed line in Fig. 7), since the partial prior δ-ferrite transforms to austenite during the second-pass thermal cycle at 1000 °C. As the second-pass thermal cycle temperature increases from 1000 °C to 1100 and 1200 °C, the fraction of ferrite is increased. This is due to the transformation from
4.1. Effect of thermal cycle on microstructural evolution Fig. 7 gives the phase fraction of ferrite in the samples after singlepass or double-pass (with the different second-pass peak temperatures) thermal cycle. The single-pass thermal cycle with the peak temperature of 900 °C increases the amount of ferrite, comparing to the initial sample after normalizing & tempering (see the red dashed line in Fig. 7). Ac1 and Ac3 points of the employed steel in this work were determined, by using the dilatometer, as 863.4 °C and 966.1 °C, respectively. Obviously, the peak temperature for the sample upon singlepass thermal cycle is between Ac1 and Ac3, resulting in the incomplete austenization. During single-pass thermal cycle at 900 °C, the untransformed martensite would occur recrystallization and thus transform to ferrite. In other words, ferrite in the sample undergoing singlepass thermal cycle includes δ-ferrite formed during hot-rolling, and αferrite due to recrystallization of martensite during thermal cycle at 900 °C. These two types of ferrite are difficult to distinguish under the OM examination, since they have the similar constituent and structure. Table 3 The welding process parameters of the experimental steel. Preheating temperature
Heat input
Heating rate
Holding time
Cooling rate
Interpass temperature
250 °C
16000 J/cm
50 °C/s
1s
25 °C/s
200 °C
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Fig. 3. OM images of the samples after (a) single-pass thermal cycle at 900 °C, and double-pass thermal cycle at (b) 900&900 °C, (c) 900&1000 °C, (d) 900&1100 °C and (e) 900&1200 °C, respectively.
austenite to δ-ferrite at 1100 and 1200 °C. C is an intense austenite stabilization element, which can expand the phase region of austenite in the equilibrium phase diagram, while Cr and W are the ferrite stabilization elements. The employed steel in this study have the low C content, and relatively high Cr and W contents. Hence, the phase region of austenite is very small. As mentioned above, thermal cycle temperature affects not only the phase fraction of ferrite and martensite, but also the grain size of prior austenite and ferrite. Fig. 8 presents the grain size of prior austenite and ferrite in the samples after single-pass or double-pass (with the different second-pass peak temperatures) thermal cycle. Thermal cycle results in the refinement of prior austenite (no larger than 20 μm, comparing to the initial sample after N&T with the grain size of 32.4 μm). The grain size of prior austenite and ferrite with the change of thermal cycle temperature has the same tendency. Double-pass thermal cycle at 900& 900 °C results in the refinement of prior austenite and ferrite. Increase of second-pass thermal cycle temperature leads to grain growth of prior austenite and ferrite. It is easy to understand that high temperature promotes the movement of grain boundary, and thus accelerates the grain growth. Fig. 9 presents the average size of M23C6 and MX particles in the samples after single-pass or double-pass (with the different second-pass peak temperatures) thermal cycle. The average size of M23C6 and MX in
the initial sample after normalizing & tempering is denoted as black and red dashed lines, respectively. The relevant data for the initial sample after normalizing & tempering is derived from our previous works [25]. Comparing with the initial sample after normalizing & tempering, the average size of M23C6 particles is increased remarkably in the samples upon single-pass or double-pass thermal cycle. It has been recognized that the dissolution temperature of M23C6 precipitates is below 1000 °C [26,27]. As a result, single-pass thermal cycle at 900 °C leads to coarsening behavior of M23C6 precipitates merely (not dissolution of them). Nevertheless, double-pass thermal cycle would give rise to partial dissolution of M23C6, appearing as decrease of M23C6 particle size. When the second-pass thermal cycle temperature increases to 1100 and 1200 °C, M23C6 precipitates are totally dissolved into matrix. On the other hand, MX precipitates have the better thermal stability. Thermal cycle temperature affects little on the size of MX particles, and they are not dissolved until heating up to 1200 °C. The volume fractions of M23C6 and MX particles in the samples after single-pass or double-pass (with the different second-pass peak temperatures) thermal cycle are shown in Fig. 10. Due to their high thermal stability, the fraction of MX precipitates remains unchanged basically, despite of the thermal cycle temperature. On the contrary, the fraction of M23C6 is decreased visibly by increase of thermal cycle temperature, reflecting dissolution of M23C6 particles. 12
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Fig. 4. TEM images of the extraction replicas samples after (a) single-pass thermal cycle at 900 °C, and double-pass thermal cycle at (a) 900&900 °C, (b) 900&1000 °C, (c) 900&1100 °C and (d) 900&1200 °C, respectively.
Fig. 6. Creep rupture time of the samples after single-pass or double-pass thermal cycle at 550 °C and 180 MPa stress.
Fig. 5. Impact energy of the samples after single-pass or double-pass thermal cycle.
[28]. Also, it was found that martensitic packets can act as the effective microstructure unit for cleavage during impact tests [29]. Hence, the higher fraction of δ-ferrite leads to the lower impact energy. Furthermore, the coarse precipitates [30] or inclusions [31], and coarse grains [29,32] would decrease the impact energy. According to the microstructure results above, thermal cycle at 900 °C leads to coarsening of M23C6 particles. Besides, the grain size of the samples after thermal cycle at 900 °C, or 900 & 1200 °C, is relatively large. Therefore, it can be
4.2. Effect of microstructural evolution on mechanical properties It can be notable that the trends in impact toughness with changes of thermal cycle parameters are in opposite directions for that in the ferrite fractions, by comparing Figs. 5 to 7 . It has been recognized that δ-ferrite is detrimental to impact toughness, since it can lowered both the crack initiation and propagation energy during the impact progress 13
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Fig. 10. The volume fractions of M23C6 and MX particles in the samples after single-pass or double-pass thermal cycle.
Fig. 7. Phase fraction of ferrite in the samples after single-pass or double-pass thermal cycle (the red dashed line denotes the phase fraction of ferrite in the initial sample after normalizing & tempering). (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article).
concluded that the δ-ferrite fraction, coarse precipitates and coarse grains would decrease the impact energy for the thermal cycle samples, though the most critical factor is the δ-ferrite fraction. S.K. Albert et al. [13] simulated the single-pass welding in a high Cr ferritic steel, and found that the shortest creep rupture time occurs in the sample after singe thermal cycle at 900 °C (between AC1 and AC3), and the longest creep rupture time in the sample at 1200 °C. This is basically in agreement with our results. However, in the present experiment, thermal cycle below 1200 °C consistently manifests the short creep rupture time, while Albert’s results show a continuous increase in rupture time with the increase of thermal cycle temperature. Thus, it is evident that the first-pass thermal cycle at 900 °C have a remarkable effect on the creep properties of the samples upon double-pass thermal cycle. As discussed above, thermal cycle, regardless of single-pass or double-pass thermal cycle, refines the prior austenitic grain, even if the thermal cycle temperature is up to 1200 °C. Due to the short creep time, Laves phase and Z phase does not form. Obviously, the evolution of the precipitates especially as M23C6 would be the main reason of the short creep life for the samples after thermal cycle below 1200 °C. According to the TEM observation and analysis, it has been recognized that thermal cycle results in coarsening and dissolution of the precipitates. As well known, creep failure in the ferritic steel weld joints at high temperature tends to occur in the fine grain heat affected zone (FGHAZ) and intercritical heat affected zone (ICHAZ) [33]. The peak temperature during welding thermal cycle for FGHAZ is just above AC3, and that for ICHAZ is between AC1 and AC3 [15]. As mentioned above, Ac1 and Ac3 points of the samples were determined as 863.4 °C and 966.1 °C, respectively. Thus, it can be considered that the samples after thermal cycle at 900 °C and 1000 °C would correspond to FGHAZ and ICHAZ, respectively. It should be pointed out, that the creep failure occurred in FGHAZ and ICHAZ in the welded joints is generally attributed to so-called type IV cracking. However, type IV cracking is considered to be happened during the long-term creep (more than 10000 h). In this project, the longest creep time of the samples is only as long as 1614 h. Thus, the rupture behaviors in the present work would not be regarded as type IV cracking. Although, these should have something in common, since the sites of fracture are both located in FGHAZ and ICHAZ. Generally speaking, type IV cracking which occurs in FGHAZ and ICHAZ may be attributed to too high density of grain boundaries due to the small size of the prior austenitic grain [21], coarsening or insufficient dissolution of the second phases [34], and the formation of the coarse precipitates (as Laves phase and Z phase) [14]. According to the measurement of the prior austenitic grain size (see Fig. 8), it can be recognized that the prior austenitic grain size has little effect on short-term creep life. Due to the
Fig. 8. Grain size of prior austenite and ferrite in the samples after single-pass or double-pass thermal cycle.
Fig. 9. Average size of M23C6 and MX in the samples after single-pass or doublepass thermal cycle (The black and red dashed lines represent the average size of M23C6 and MX in the initial sample after normalizing & tempering, respectively [25]). (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article).
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Data availability
short creep time in this project, the formation of the coarse Laves phase and Z phase particles is seldom possible. Thus, coarsening or insufficient dissolution of the second phases maybe the main reason for the very poor creep property of the samples after double-pass thermal cycle below 1200 °C. Y. Liu et al. [34] indicated that dissolution and re-formation of M23C6 results in lack of sufficient boundary strengthening effect by precipitates, since M23C6 particles would scarcely precipitated at the newly formed grain boundaries. The alloying elements’ redistribution after dissolution and precipitation of the second-phase particles during thermal cycle would also cause the early creep failure [35,36]. The precipitates formed at the prior austenitic boundary play an important role to stabilize the microstructure during creep [21]. Coarsening of M23C6 particles no longer give enough pinning effect to the grain boundary, and thus leads to degradation of creep properties. The above results and discussion only focus on the single-pass welding process, and there is no report about the creep properties of the high Cr ferritic steels after double-pass welding. In this work, it is found double-pass thermal cycle results in coarsening of M23C6 particles at the relatively low second cycle temperature. When the second cycle temperature is higher than 1100 °C, M23C6 particles will be dissolved totally. Hence, the mechanism for the low creep life in the sample after double-pass thermal cycle below 1200 °C is similar to that for type IV cracking. The first thermal cycle at 900 °C has a significant influence on the distribution of M23C6, and the short-term creep life. Until the second thermal cycle temperature increases to 1200 °C, M23C6 is dissolved into matrix, and the alloy elements are homogenized. Under this condition, the effect of the first thermal cycle at 900 °C is removed, giving rise to the longest creep rupture time for the sample after 900&1200 °C double-pass thermal cycle.
The raw/processed data required to reproduce these findings cannot be shared at this time as the data also forms part of an ongoing study. Acknowledgments The authors are grateful to the National Magnetic Confinement Fusion Energy Research Project (granted No. 2015GB119001), the National Natural Science Foundation of China (granted No. U1660201), and the Project of Natural Science Foundation of Tianjin (granted Nos. 18JCQNJC03300 and 18YFZCGX00050) for grant and financial support. References [1] Zhou X, Liu C, Yu L, Liu Y, Li H. Phase transformation behavior and microstructural control of High-Cr Martensitic/Ferritic heat-resistant steels for power and nuclear plants: a review. J Mater Sci Technol 2015;31:235–42. [2] Fujita T. Current progress in advanced high Cr ferritic steels for high-temperature applications. Isij Int 1992;32(2):175–81. [3] Abe F. Progress in Creep-Resistant steels for high efficiency coal-fired power plants. J Press Vessel Technol 2016;138(4):040804. [4] Lu Q, van der Zwaag S, Xu W. Charting the ‘composition–strength’space for novel austenitic, martensitic and ferritic creep resistant steels. J Mater Sci Technol 2017;33(12):1577–81. [5] Klueh RL. Elevated temperature ferritic and martensitic steels and their application to future nuclear reactors. Int Mater Rev 2005;50(5):287–310. [6] Zhang C, Cui L, Liu Y, Liu C, Li H. Microstructures and mechanical properties of friction stir welds on 9% Cr reduced activation ferritic/martensitic steel. J Mater Sci Technol 2018;34(5):756–66. [7] Huang Q, Baluc N, Dai Y, Jitsukawa S, Kimura A, Konys J, et al. Recent progress of R &D activities on reduced activation ferritic/martensitic steels. J Nucl Mater 2013;442(1-3):S2–8. [8] Abe F, Tabuchi M. Microstructure and creep strength of welds in advanced ferritic power plant steels. Sci Technol Weld Join 2004;9(1):22–30. [9] Francis JA, Mazur W, Bhadeshia HKDH. Review Type IV cracking in ferritic power plant steels. Mater Sci Technol 2006;22(12):1387–95. [10] Tabuchi M, Watanabe T, Kubo K, Matsui M, Kinugawa J, Abe F. Creep crack growth behavior in the HAZ of weldments of W containing high Cr steel. Int J Press Vessel Pip 2001;78(11–12):779–84. [11] Watanabe T, Tabuchi M, Yamazaki M, Hongo H, Tanabe T. Creep damage evaluation of 9Cr–1Mo–V–Nb steel welded joints showing Type IV fracture. Int J Press Vessel Pip 2006;83(1):63–71. [12] Francis J, Mazur W, Bhadeshia H. Estimation of type IV cracking tendency in power plant steels. Isij Int 2004;44(11):1966–8. [13] Albert SK, Matsui M, Hongo H, Watanabe T, Kubo K, Tabuchi M. Creep rupture properties of HAZs of a high Cr ferritic steel simulated by a weld simulator. Int J Press Vessel Pip 2004;81(3):221–34. [14] Wang X, Xu Q, Yu S-m, Liu H, Hu L, Ren Y-y. Laves-phase evolution during aging in fine grained heat-affected zone of a tungsten-strengthened 9% Cr steel weldment. J Mater Process Technol 2015;219:60–9. [15] Parker J. In-service behaviour of creep strength enhanced ferritic steels Grade 91 and Grade 92 – Part 2 weld issues. Int J Press Vessel Pip 2014;114–115:76–87. [16] Pandey C, Mahapatra MM, Kumar P, Kumar S, Sirohi S. Effect of post weld heat treatments on microstructure evolution and type IV cracking behavior of the P91 steel welds joint. J Mater Process Technol 2019;266:140–54. [17] Divya M, Das CR, Albert SK, Goyal S, Ganesh P, Kaul R, et al. Influence of welding process on Type IV cracking behavior of P91 steel. Mater Sci Eng A 2014;613:148–58. [18] Francis JA, Cantin GMD, Mazur W, Bhadeshia HKDH. Effects of weld preheat temperature and heat input on type IV failure. Sci Technol Weld Join 2009;14(5):436–42. [19] Mayr P, Martín FM, Albu M, Cerjak H. Correlation of creep strength and microstructural evolution of a boron alloyed 9Cr3W3CoVNb steel in as-received and welded condition. Mater High Temp 2010;27(1):67–72. [20] Das CR, Albert SK, Swaminathan J, Raju S, Bhaduri AK, Murty BS. Transition of crack from type IV to type II resulting from improved utilization of boron in the modified 9Cr-1Mo steel weldment. Metall Mater Trans A 2012;43(10):3724–41. [21] Liu Y, Tsukamoto S, Sawada K, Abe F. Role of boundary strengthening on prevention of type IV failure in high Cr ferritic heat-resistant steels. Metall Mater Trans A 2013;45(3):1306–14. [22] Liu W, Lu F, Wei Y, Ding Y, Wang P, Tang X. Special zone in multi-layer and multipass welded metal and its role in the creep behavior of 9Cr 1Mo welded joint. Mater Des 2016;108:195–206. [23] Chen J-g, Liu Y-c, Liu C-x, Yan B-y, Li H-j. Effects of tantalum on austenitic transformation kinetics of RAFM steel. J Iron Steel Res Int 2017;24(7):705–10. [24] Feng YY, Luo ZA, Zhang DH, Wei J, Wang LJ. The application of welding heat cycle computer software. Mater Sci Forum 2008;575–578:821–6. [25] Chen J, Liu C, Liu Y, Yan B, Li H. Effects of tantalum content on the microstructure and mechanical properties of low-carbon RAFM steel. J Nucl Mater
5. Conclusions Thermal simulation on double-pass welding of a high Cr ferritic steel was conducted. The microstructures and mechanical properties of the samples after double-pass thermal cycle were investigated. The main conclusions are as follows: Thermal cycle almost leads to the increase of ferrite faction and the decrease of martensite fraction, except for the sample after thermal cycle at 900&1000 °C. Thermal cycle results in the refinement of prior austenite, and increase of second-pass thermal cycle temperature leads to grain growth of prior austenite and ferrite. The average size of M23C6 precipitates with the low thermal stability is increased remarkably by thermal cycle. Due to its high thermal stability, the average size of nano-sized MX precipitates remains unchanged basically. M23C6 precipitates are dissolved totally when the second thermal cycle increases to 1100 °C, while MX particles are undissolved until at 1200 °C. Thermal cycle at low temperature (as 900 °C) deteriorates the impact toughness. The sample after double-pass thermal cycle for 900& 1000 °C have the highest impact energy. By further increasing the second-pass thermal cycle temperature, the impact energy is decreased. This is mainly due to the increase of δ-ferrite fraction. Short-term creep tests indicate that, thermal cycle at 900 & 1200 °C leads to the longest creep rupture time, while other samples have the very short rupture time. Coarse M23C6 particles and the heterogeneous alloy elements distribution maybe the main reason. The second thermal cycle at 1200 °C dissolves all the second-phase particles, and homogenize the compositions, promoting the creep properties. The mechanism for that is similar to the type IV cracking. Declaration of interests The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. 15
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carbon microalloyed steels. Metall Mater Trans A 2006;37(7):2147–58. [32] Hyzak JM, Bernstein IM. The role of microstructure on the strength and toughness of fully pearlitic steels. Metall Trans A 1976;7(8):1217–24. [33] Sakthivel T, Vasudevan M, Laha K, Parameswaran P, Chandravathi KS, Panneer Selvi S, et al. Creep rupture behavior of 9Cr–1.8W–0.5Mo–VNb (ASME grade 92) ferritic steel weld joint. Mater Sci Eng A 2014;591:111–20. [34] Liu Y, Tsukamoto S, Shirane T, Abe F. Formation mechanism of type IV failure in high Cr ferritic heat-resistant steel-welded joint. Metall Mater Trans A 2013;44(10):4626–33. [35] Falat L, Čiripová L, Kepič J, Buršík J, Podstranská I. Correlation between microstructure and creep performance of martensitic/austenitic transition weldment in dependence of its post-weld heat treatment. Eng Fail Anal 2014;40:141–52. [36] Zhang W, Wang X, Gong J, Jiang Y, Huang X. Experimental and simulated characterization of creep behavior of P92 steel with prior cyclic loading damage. J Mater Sci Technol 2017;33(12):1540–8.
2016;479:295–301. [26] Schneider A, Inden G. Simulation of the kinetics of precipitation reactions in ferritic steels. Acta Mater 2005;53(2):519–31. [27] Xiao X, Liu G, Hu B, Wang J, Ma W. Microstructure stability of V and Ta microalloyed 12% Cr reduced activation ferrite/martensite steel during long-term aging at 650 C. J Mater Sci Technol 2015;31(3):311–9. [28] Wang P, Lu SP, Xiao NM, Li DZ, Li YY. Effect of delta ferrite on impact properties of low carbon 13Cr–4Ni martensitic stainless steel. Mater Sci Eng A 2010;527(13–14):3210–6. [29] Wang C, Wang M, Shi J, Hui W, Dong H. Effect of microstructural refinement on the toughness of low carbon martensitic steel. Scr Mater 2008;58(6):492–5. [30] Wei Z, Laizhu J, Jincheng H, Hongmei S. Effect of ageing on precipitation and impact energy of 2101 economical duplex stainless steel. Mater Charact 2009;60(1):50–5. [31] Yan W, Shan YY, Yang K. Effect of TiN inclusions on the impact toughness of low-
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Contents lists available at ScienceDirect
Journal of Manufacturing Processes journal homepage: www.elsevier.com/locate/manpro
Thermal simulation on double-pass welding of a high Cr ferritic steel a
a
b
a
b
a
T a
Yi Shao , Biyu Yan , Yihuan Liu , Chunliang Mao , Chen Wei , Yongchang Liu , Zesheng Yan , ⁎ Huijun Lia, Chenxi Liua, a b
State Key Lab of Hydraulic Engineering Simulation and Safety, School of Materials Science & Engineering, Tianjin University, Tianjin, 300354, PR China Tianjin Special Equipment Inspection Institute, Tianjin, 300192, PR China
A R T I C LE I N FO
A B S T R A C T
Keywords: Welding thermal simulation Multi-Pass welding High Cr ferritic steel Creep life
The heat affected zone (HAZ) after double-pass welding in a high Cr ferritic steel was simulated by using a thermomechanical simulator. The microstructural observation and mechanical properties tests were conducted to investigate the effect of double-pass thermal cycle on the microstructure evolution and mechanical properties. The results show that the double-pass thermal cycle leads to the decrease of martensitic fraction, and the refinement of prior austenite. M23C6 precipitates is coarsened, and dissolved subsequently, by increasing the peak temperature during the second-pass thermal cycle, while MX with the higher thermal stability is undissolved until at 1200 °C. Thermal cycle at low temperature (as 900 °C) deteriorates the impact toughness, due to the high δ-ferrite fraction. Short-term creep tests suggest that, only the second-pass thermal cycle temperate of 1200 °C could eliminate the unfavorable effect on creep life resulted from the first-pass thermal cycle. Coarse M23C6 particles and the heterogeneous alloy elements distribution maybe the main reason.
1. Introduction High Cr ferritic steels have been the important structural materials for the main steam pipe and heater with larger diameter and thick wall in ultra-super critical (USC) fossil power plants, due to their high creep rupture strength, low thermal expansion coefficient, good corrosion resistance, and high performance–cost ratio [1–4]. In addition, considering their outstanding resistance to radiation induced void swelling and irradiation embrittlement, the high Cr ferritic steels have become the candidate materials used for nuclear power plants and test blanket modules (TBMs) of International Thermonuclear Experimental Reactor (ITER) [5–7]. Nevertheless, it was found that the creep life of the welding joints in the high Cr ferritic steels are significantly lower than that of the base metal [8]. It has been recognized that the creep cracks usually occur in the fine-grained heat affected zone (FGHAZ) and the intercritical heat affected zone (ICHAZ) of the welds [9]. During welding thermal cycle, FGHAZ is heated just above Ac3, and ICHAZ is heated between Ac1 and Ac3. These creep damage behaviors are called as type IV cracking [10]. Type IV cracking is likely to occur under lower stress and higher temperature conditions [11]. Type IV cracking predominates when the rupture stress is less than about 100MPa [12]. In FGHAZ, the recovery of martensite during creep is heterogeneous so that some coarse subgrains form near the fine subgrains [8]. Moreover, the coarsening of
⁎
M23C6 and the recovery of dislocations are significant in FGHAZ and ICHAZ. The dislocation density in FGHAZ and ICHAZ is significantly lower than the weld metal and base metal [11]. The Creep cavities nucleate preferentially around the coarse M23C6 particles [13]. Besides, the coarse Laves phase forming during creep would also contribute to the formation of the creep cavities [14]. The consumption of solutionstrengthening elements as W, Cr during the formation of Laves phase would also be the key factor for the occurrence of type IV cracking [15]. The post weld heat treatment can eliminate the type IV cracking, though it rarely is a practical option in the actual production [16]. The weld process and parameters also affect the type IV cracking behaviors. It was found that the creep rupture lives of laser weldment are higher than that of shielded metal arc welding (SMAW), since the width of HAZ in the former is narrower than that in the latter [17]. The joint geometry and preheat temperature can ameliorate the type IV failures, while the effect of heat input is less significant [18]. Generally speaking, the width of the heat affected zone (HAZ) in welding joints is only serval millimeters, so that the mechanical tests are difficult to be carried out for each region in HAZ (such as FGHAZ and ICHAZ). Hence, the thermal simulation method is used to “magnify” HAZ, and thus facilitates the mechanical tests, and the further microstructural analysis (as TEM), for each region in HAZ [13]. This method is to simulate the HAZ microstructure by set a thermal cycle that is similar to a certain region of the HAZ which is undergone during
Corresponding author. E-mail address: [email protected] (C. Liu).
https://doi.org/10.1016/j.jmapro.2019.05.012 Received 19 January 2019; Received in revised form 14 April 2019; Accepted 9 May 2019 Available online 15 May 2019 1526-6125/ © 2019 The Society of Manufacturing Engineers. Published by Elsevier Ltd. All rights reserved.
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up to 900 °C with a rate of 50 °C/s and held for 1 s, followed by cooled to 200 °C with a rate of 25 °C/s. For the second-pass thermal cycle, the samples after the first-pass thermal cycle were reheated to 900, 1000, 1100, 1200 °C with a rate of 50 °C/s, respectively, and held for 1 s, followed by cooled to room temperature with a rate of 25 °C/s. The samples undergoing double-pass thermal cycle are denoted as “900& 900 °C”, “900&1000 °C”, “900&1100 °C” and “900&1200 °C” in terms of their different peak temperatures during the second-pass thermal cycle. The schematic diagram and the summary table of the single- or doublepass thermal cycle procedures is presented in Fig. 2 and Table 4, respectively. After the thermal simulation experiments, the samples were mounted, polished, and etched in solution of nitric acid and ethyl alcohol for metallographic observation. Extraction carbon replicas specimen were prepared to conduct the transmission electron microscopy. For preparation of the replicas, the samples after metallographic observation were etched in a solution of ethanol (90 ml), hydrochloric acid (10 ml) and picric acid (1 g). Then the etched surfaces of the samples were sprayed by carbon nanoparticles in vacuum. After spraying carbon, the scratched carbon films were floated to the surface using a solution of ethanol (70 ml), hydrochloric acid (20 ml) and nitric acid (10 ml), and fished out by the copper grids. The microstructural observation was carried out on Olympus C-35A Optical Microscope (OM), and FEI Tecnai G2 Transmission Electron Microscope (TEM). Due to the size limit of the thermal simulation samples (length for 70 mm and diameter for 6 mm), the subsized specimens for creep and impact tests were used. The creep tests were conducted on RPL100 electron creep testing machine with the subsized samples (φ6 × 63 mm). The creep test parameters are that the creep temperature of 550 °C and the loading stress of 180 MPa. The Charpy impact tests were carried out at room temperature on CBD 500 Charpy impact test machine with the subsized V-notch samples (55 × 2.5 × 2.5 mm). The impact tests were conducted three times for each sample.
the actual welding process. By using this method, it has been found that the high density of the grain boundaries, coarse precipitates, and the homogeneous grain and compositions maybe the key factors for type IV cracking [19–21]. The above research results only focus on the single-pass welding. However, in the actual welding, multi-pass welding is often used to welding the thick plate or pipe, due to the limitation of weld penetration for fusion welding. The multi-pass welding results in the course of thermal cycle in HAZ being more complicated. Depending on factors such as heat input, electrode size and angle and degree of overlap, the different regions in HAZ are exposed to the various multiple thermal cycles [15]. Heterogeneous structure of WM, due to reheating during multi-pass and multi-layer welding, has complex effects on the property of welded joint [22]. So far, the welding thermal simulation for the double-pass (or multi-pass) welding in the high Cr ferritic steels have not been reported. The effect of double-pass (or multi-pass) welding on type IV cracking was not clarified. Therefore, the aim of this work is to investigate the effect of double-pass welding on microstructure evolution and mechanical properties of the high Cr ferritic steel by using the thermal simulation method. According to the previous studies, thermal simulation mainly focuses on FGHAZ and ICHAZ in the welded joint of the high Cr ferritic steels, since these two regions is the weakness of components during their service. Thus, the first-pass thermal cycle temperature was set to be close to the peak temperature for FGHAZ and ICHAZ in this work. The range of the second-pass thermal cycle temperatures were selected to wider. After the welding thermal simulation, the microstructures of the samples were observed, and analyzed quantitatively. The impact toughness and short-term creep time was tested to investigate the effect of microstructure morphology, especially as distribution of the precipitates, on the mechanical properties. 2. Experimental details The employed steel in this project is a high Cr ferritic steel used for nuclear fusion reactor components, whose chemical composition is given by Table 1. Before the thermal simulation experiments, the raw material (as-rolled state) was normalized at 1050 °C for 0.5 h and then tempered at 750 °C for 1.5 h. The microstructure of the initial samples before and after normalizing & tempering is presented in Fig. 1. Martensitic laths and δ-ferrite were found in the initial sample, since low carbon content (0.4 in wt.%, as seen in Table 1) expands the phase region of ferrite in the equilibrium phase diagram [23]. The size of the prior austenitic grain of the sample after N&T is measured as 32.4 μm. Although normalizing & tempering treatment results in decrease of δferrite amount, considerable δ-ferrite distributed along the direction of rolling is still retained. After heat treatments, the samples were cut into the shape of cylinder with a length of 70 mm and a diameter of 6 mm by the wire cutting machine. The thermal simulation experiments were conducted on Gleeble3500 thermomechanical simulator with the Rykalin-3D programme. The detailed information about the Rykalin-3D programme on Gleeble3500 thermomechanical simulator can be seen in Ref. [24]. The required physical parameters and welding process parameters used for the Rykalin-3D procedure are shown in Tables 2 and 3, respectively. The peak temperature of the first thermal cycle was set as 900 °C, and that of the second one was set as 900, 1000, 1100, 1200 °C, to simulate the fine-grained region upon the double-pass welding process. According to the Rykalin-3D procedure, for the first-pass thermal cycle, the samples are preheated to 250 °C with a rate of 78 °C/s, then heated
3. Results 3.1. Microstructural observation Fig. 3 shows the OM images of the samples after single-pass thermal cycle with the peak temperature of 900 °C, and after double-pass thermal cycle with the peak temperature of 900&900 °C, 900&1000 °C, 900&1100 °C and 900&1200 °C. It is found that the microstructures of all samples, whether upon single-pass or double-pass thermal cycle, are composed of lath martensite and ferrite, which is similar to that of the initial sample (see Fig. 1). However, the phase fraction of martensite and ferrite, and the grain size of prior austenite and ferrite is strongly affected by the thermal cycle process. The TEM images of the extraction replicas samples upon singe-pass or double-pass thermal cycle are shown in Fig. 4. The extraction replicas only reserve the second-phase particles, losing the information of the matrix as grain boundary and dislocation. It can be found that almost all samples have the coarse and rod-liked M23C6 particles and the fine MX precipitates, although the number density and average size of the particles may different with the thermal cycle temperature changing. However, M23C6 particles are absent in the sample for 900& 1100 °C, and no precipitate is found in the sample for 900&1200 °C. This suggests that MX precipitate have the higher thermal stability than M23C6. The latter is dissolved at 1100 °C, whereas the former is stable until heating up to 1200 °C. 3.2. Mechanical properties
Table 1 The Chemical compositions of experimental steel (wt.%). C
Cr
W
Mn
Si
V
Ta
Fe
0.05
8.9
1.7
0.45
< 0.1%
0.20
0.15
Bal
To evaluate the effect of single-pass and double-pass welding thermal cycle on the mechanical properties in the employed steel, Charpy impact and short-term creep testing were conducted. Fig. 5 gives the impact energy at room temperature in the samples after 10
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Fig. 1. Microstructure of the initial samples: (a) before and (b) after normalizing & tempering treatment. The size of the prior austenitic grain of the sample after N&T is measured as 32.4 μm. Table 2 The physical parameters of the experimental steel. Density 7.8 g/cm
Specific heat 3
o
0.58 J/(g· C)
Thermal conductivity 0.353 J/(cm·s·oC)
single-pass or double-pass (with the different second-pass peak temperatures) thermal cycle. It is found that thermal cycle at 900 °C deteriorates the impact toughness. Furthermore, double-pass thermal cycle for 900&900 °C would aggravate this degradation. The sample after double-pass thermal cycle for 900&1000 °C have the highest impact energy. By further increasing the second-pass thermal cycle temperature, the impact energy is decreased. For the high Cr ferritic steels used in nuclear and coal power stations, the creep properties are the most important performance indicators. The creep rupture time, under the condition of 550 °C / 180 MPa, of the samples after single-pass or double-pass (with the different second-pass peak temperatures) thermal cycle is presented in Fig. 6. It is obviously found that thermal cycle at 900 & 1200 °C leads to the longest creep rupture time (up to 1614 h), whereas other samples have the very short rupture time (less than 350 h).
Fig. 2. The schematic diagram of the double-pass thermal cycle procedures. Table 4 The thermal cycles parameters for the single-pass or double-pass thermal cycle samples. Thermal cycle times
Single-pass
Double-pass
Thermal cycle procedures
900 °C
900& 900 °C
900& 1000 °C
900& 1100 °C
900& 1200 °C
4. Discussion Nevertheless, by comparing Figs.1b with 3 a, it is found that the banded distribution feature of ferrite in the initial sample after normalizing & tempering is not obvious in the sample after single-pass thermal cycle. Therefore, it can be inferred that recrystallization of martensite distributes randomly and thus disorders apparently the banded characteristic of the already existed δ-ferrite before thermal cycle. As a result, the amount of ferrite is increased by thermal cycle at 900 °C. Likewise, double-pass thermal cycle for 900&900 °C would promote recrystallization of martensite further, and increase the amount of ferrite accordingly. As the second-pass thermal cycle temperature continues to increase to 1000 °C, which is higher than Ac3 point, martensite transforms to austenite completely. Moreover, the phase fraction of ferrite in the sample for 900&1000 °C is lower than that in the initial sample after normalizing & tempering (see the red dashed line in Fig. 7), since the partial prior δ-ferrite transforms to austenite during the second-pass thermal cycle at 1000 °C. As the second-pass thermal cycle temperature increases from 1000 °C to 1100 and 1200 °C, the fraction of ferrite is increased. This is due to the transformation from
4.1. Effect of thermal cycle on microstructural evolution Fig. 7 gives the phase fraction of ferrite in the samples after singlepass or double-pass (with the different second-pass peak temperatures) thermal cycle. The single-pass thermal cycle with the peak temperature of 900 °C increases the amount of ferrite, comparing to the initial sample after normalizing & tempering (see the red dashed line in Fig. 7). Ac1 and Ac3 points of the employed steel in this work were determined, by using the dilatometer, as 863.4 °C and 966.1 °C, respectively. Obviously, the peak temperature for the sample upon singlepass thermal cycle is between Ac1 and Ac3, resulting in the incomplete austenization. During single-pass thermal cycle at 900 °C, the untransformed martensite would occur recrystallization and thus transform to ferrite. In other words, ferrite in the sample undergoing singlepass thermal cycle includes δ-ferrite formed during hot-rolling, and αferrite due to recrystallization of martensite during thermal cycle at 900 °C. These two types of ferrite are difficult to distinguish under the OM examination, since they have the similar constituent and structure. Table 3 The welding process parameters of the experimental steel. Preheating temperature
Heat input
Heating rate
Holding time
Cooling rate
Interpass temperature
250 °C
16000 J/cm
50 °C/s
1s
25 °C/s
200 °C
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Fig. 3. OM images of the samples after (a) single-pass thermal cycle at 900 °C, and double-pass thermal cycle at (b) 900&900 °C, (c) 900&1000 °C, (d) 900&1100 °C and (e) 900&1200 °C, respectively.
austenite to δ-ferrite at 1100 and 1200 °C. C is an intense austenite stabilization element, which can expand the phase region of austenite in the equilibrium phase diagram, while Cr and W are the ferrite stabilization elements. The employed steel in this study have the low C content, and relatively high Cr and W contents. Hence, the phase region of austenite is very small. As mentioned above, thermal cycle temperature affects not only the phase fraction of ferrite and martensite, but also the grain size of prior austenite and ferrite. Fig. 8 presents the grain size of prior austenite and ferrite in the samples after single-pass or double-pass (with the different second-pass peak temperatures) thermal cycle. Thermal cycle results in the refinement of prior austenite (no larger than 20 μm, comparing to the initial sample after N&T with the grain size of 32.4 μm). The grain size of prior austenite and ferrite with the change of thermal cycle temperature has the same tendency. Double-pass thermal cycle at 900& 900 °C results in the refinement of prior austenite and ferrite. Increase of second-pass thermal cycle temperature leads to grain growth of prior austenite and ferrite. It is easy to understand that high temperature promotes the movement of grain boundary, and thus accelerates the grain growth. Fig. 9 presents the average size of M23C6 and MX particles in the samples after single-pass or double-pass (with the different second-pass peak temperatures) thermal cycle. The average size of M23C6 and MX in
the initial sample after normalizing & tempering is denoted as black and red dashed lines, respectively. The relevant data for the initial sample after normalizing & tempering is derived from our previous works [25]. Comparing with the initial sample after normalizing & tempering, the average size of M23C6 particles is increased remarkably in the samples upon single-pass or double-pass thermal cycle. It has been recognized that the dissolution temperature of M23C6 precipitates is below 1000 °C [26,27]. As a result, single-pass thermal cycle at 900 °C leads to coarsening behavior of M23C6 precipitates merely (not dissolution of them). Nevertheless, double-pass thermal cycle would give rise to partial dissolution of M23C6, appearing as decrease of M23C6 particle size. When the second-pass thermal cycle temperature increases to 1100 and 1200 °C, M23C6 precipitates are totally dissolved into matrix. On the other hand, MX precipitates have the better thermal stability. Thermal cycle temperature affects little on the size of MX particles, and they are not dissolved until heating up to 1200 °C. The volume fractions of M23C6 and MX particles in the samples after single-pass or double-pass (with the different second-pass peak temperatures) thermal cycle are shown in Fig. 10. Due to their high thermal stability, the fraction of MX precipitates remains unchanged basically, despite of the thermal cycle temperature. On the contrary, the fraction of M23C6 is decreased visibly by increase of thermal cycle temperature, reflecting dissolution of M23C6 particles. 12
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Fig. 4. TEM images of the extraction replicas samples after (a) single-pass thermal cycle at 900 °C, and double-pass thermal cycle at (a) 900&900 °C, (b) 900&1000 °C, (c) 900&1100 °C and (d) 900&1200 °C, respectively.
Fig. 6. Creep rupture time of the samples after single-pass or double-pass thermal cycle at 550 °C and 180 MPa stress.
Fig. 5. Impact energy of the samples after single-pass or double-pass thermal cycle.
[28]. Also, it was found that martensitic packets can act as the effective microstructure unit for cleavage during impact tests [29]. Hence, the higher fraction of δ-ferrite leads to the lower impact energy. Furthermore, the coarse precipitates [30] or inclusions [31], and coarse grains [29,32] would decrease the impact energy. According to the microstructure results above, thermal cycle at 900 °C leads to coarsening of M23C6 particles. Besides, the grain size of the samples after thermal cycle at 900 °C, or 900 & 1200 °C, is relatively large. Therefore, it can be
4.2. Effect of microstructural evolution on mechanical properties It can be notable that the trends in impact toughness with changes of thermal cycle parameters are in opposite directions for that in the ferrite fractions, by comparing Figs. 5 to 7 . It has been recognized that δ-ferrite is detrimental to impact toughness, since it can lowered both the crack initiation and propagation energy during the impact progress 13
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Fig. 10. The volume fractions of M23C6 and MX particles in the samples after single-pass or double-pass thermal cycle.
Fig. 7. Phase fraction of ferrite in the samples after single-pass or double-pass thermal cycle (the red dashed line denotes the phase fraction of ferrite in the initial sample after normalizing & tempering). (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article).
concluded that the δ-ferrite fraction, coarse precipitates and coarse grains would decrease the impact energy for the thermal cycle samples, though the most critical factor is the δ-ferrite fraction. S.K. Albert et al. [13] simulated the single-pass welding in a high Cr ferritic steel, and found that the shortest creep rupture time occurs in the sample after singe thermal cycle at 900 °C (between AC1 and AC3), and the longest creep rupture time in the sample at 1200 °C. This is basically in agreement with our results. However, in the present experiment, thermal cycle below 1200 °C consistently manifests the short creep rupture time, while Albert’s results show a continuous increase in rupture time with the increase of thermal cycle temperature. Thus, it is evident that the first-pass thermal cycle at 900 °C have a remarkable effect on the creep properties of the samples upon double-pass thermal cycle. As discussed above, thermal cycle, regardless of single-pass or double-pass thermal cycle, refines the prior austenitic grain, even if the thermal cycle temperature is up to 1200 °C. Due to the short creep time, Laves phase and Z phase does not form. Obviously, the evolution of the precipitates especially as M23C6 would be the main reason of the short creep life for the samples after thermal cycle below 1200 °C. According to the TEM observation and analysis, it has been recognized that thermal cycle results in coarsening and dissolution of the precipitates. As well known, creep failure in the ferritic steel weld joints at high temperature tends to occur in the fine grain heat affected zone (FGHAZ) and intercritical heat affected zone (ICHAZ) [33]. The peak temperature during welding thermal cycle for FGHAZ is just above AC3, and that for ICHAZ is between AC1 and AC3 [15]. As mentioned above, Ac1 and Ac3 points of the samples were determined as 863.4 °C and 966.1 °C, respectively. Thus, it can be considered that the samples after thermal cycle at 900 °C and 1000 °C would correspond to FGHAZ and ICHAZ, respectively. It should be pointed out, that the creep failure occurred in FGHAZ and ICHAZ in the welded joints is generally attributed to so-called type IV cracking. However, type IV cracking is considered to be happened during the long-term creep (more than 10000 h). In this project, the longest creep time of the samples is only as long as 1614 h. Thus, the rupture behaviors in the present work would not be regarded as type IV cracking. Although, these should have something in common, since the sites of fracture are both located in FGHAZ and ICHAZ. Generally speaking, type IV cracking which occurs in FGHAZ and ICHAZ may be attributed to too high density of grain boundaries due to the small size of the prior austenitic grain [21], coarsening or insufficient dissolution of the second phases [34], and the formation of the coarse precipitates (as Laves phase and Z phase) [14]. According to the measurement of the prior austenitic grain size (see Fig. 8), it can be recognized that the prior austenitic grain size has little effect on short-term creep life. Due to the
Fig. 8. Grain size of prior austenite and ferrite in the samples after single-pass or double-pass thermal cycle.
Fig. 9. Average size of M23C6 and MX in the samples after single-pass or doublepass thermal cycle (The black and red dashed lines represent the average size of M23C6 and MX in the initial sample after normalizing & tempering, respectively [25]). (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article).
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Data availability
short creep time in this project, the formation of the coarse Laves phase and Z phase particles is seldom possible. Thus, coarsening or insufficient dissolution of the second phases maybe the main reason for the very poor creep property of the samples after double-pass thermal cycle below 1200 °C. Y. Liu et al. [34] indicated that dissolution and re-formation of M23C6 results in lack of sufficient boundary strengthening effect by precipitates, since M23C6 particles would scarcely precipitated at the newly formed grain boundaries. The alloying elements’ redistribution after dissolution and precipitation of the second-phase particles during thermal cycle would also cause the early creep failure [35,36]. The precipitates formed at the prior austenitic boundary play an important role to stabilize the microstructure during creep [21]. Coarsening of M23C6 particles no longer give enough pinning effect to the grain boundary, and thus leads to degradation of creep properties. The above results and discussion only focus on the single-pass welding process, and there is no report about the creep properties of the high Cr ferritic steels after double-pass welding. In this work, it is found double-pass thermal cycle results in coarsening of M23C6 particles at the relatively low second cycle temperature. When the second cycle temperature is higher than 1100 °C, M23C6 particles will be dissolved totally. Hence, the mechanism for the low creep life in the sample after double-pass thermal cycle below 1200 °C is similar to that for type IV cracking. The first thermal cycle at 900 °C has a significant influence on the distribution of M23C6, and the short-term creep life. Until the second thermal cycle temperature increases to 1200 °C, M23C6 is dissolved into matrix, and the alloy elements are homogenized. Under this condition, the effect of the first thermal cycle at 900 °C is removed, giving rise to the longest creep rupture time for the sample after 900&1200 °C double-pass thermal cycle.
The raw/processed data required to reproduce these findings cannot be shared at this time as the data also forms part of an ongoing study. Acknowledgments The authors are grateful to the National Magnetic Confinement Fusion Energy Research Project (granted No. 2015GB119001), the National Natural Science Foundation of China (granted No. U1660201), and the Project of Natural Science Foundation of Tianjin (granted Nos. 18JCQNJC03300 and 18YFZCGX00050) for grant and financial support. References [1] Zhou X, Liu C, Yu L, Liu Y, Li H. Phase transformation behavior and microstructural control of High-Cr Martensitic/Ferritic heat-resistant steels for power and nuclear plants: a review. J Mater Sci Technol 2015;31:235–42. [2] Fujita T. Current progress in advanced high Cr ferritic steels for high-temperature applications. Isij Int 1992;32(2):175–81. [3] Abe F. Progress in Creep-Resistant steels for high efficiency coal-fired power plants. J Press Vessel Technol 2016;138(4):040804. [4] Lu Q, van der Zwaag S, Xu W. Charting the ‘composition–strength’space for novel austenitic, martensitic and ferritic creep resistant steels. J Mater Sci Technol 2017;33(12):1577–81. [5] Klueh RL. Elevated temperature ferritic and martensitic steels and their application to future nuclear reactors. Int Mater Rev 2005;50(5):287–310. [6] Zhang C, Cui L, Liu Y, Liu C, Li H. Microstructures and mechanical properties of friction stir welds on 9% Cr reduced activation ferritic/martensitic steel. J Mater Sci Technol 2018;34(5):756–66. [7] Huang Q, Baluc N, Dai Y, Jitsukawa S, Kimura A, Konys J, et al. Recent progress of R &D activities on reduced activation ferritic/martensitic steels. J Nucl Mater 2013;442(1-3):S2–8. [8] Abe F, Tabuchi M. Microstructure and creep strength of welds in advanced ferritic power plant steels. Sci Technol Weld Join 2004;9(1):22–30. [9] Francis JA, Mazur W, Bhadeshia HKDH. Review Type IV cracking in ferritic power plant steels. Mater Sci Technol 2006;22(12):1387–95. [10] Tabuchi M, Watanabe T, Kubo K, Matsui M, Kinugawa J, Abe F. Creep crack growth behavior in the HAZ of weldments of W containing high Cr steel. Int J Press Vessel Pip 2001;78(11–12):779–84. [11] Watanabe T, Tabuchi M, Yamazaki M, Hongo H, Tanabe T. Creep damage evaluation of 9Cr–1Mo–V–Nb steel welded joints showing Type IV fracture. Int J Press Vessel Pip 2006;83(1):63–71. [12] Francis J, Mazur W, Bhadeshia H. Estimation of type IV cracking tendency in power plant steels. Isij Int 2004;44(11):1966–8. [13] Albert SK, Matsui M, Hongo H, Watanabe T, Kubo K, Tabuchi M. Creep rupture properties of HAZs of a high Cr ferritic steel simulated by a weld simulator. Int J Press Vessel Pip 2004;81(3):221–34. [14] Wang X, Xu Q, Yu S-m, Liu H, Hu L, Ren Y-y. Laves-phase evolution during aging in fine grained heat-affected zone of a tungsten-strengthened 9% Cr steel weldment. J Mater Process Technol 2015;219:60–9. [15] Parker J. In-service behaviour of creep strength enhanced ferritic steels Grade 91 and Grade 92 – Part 2 weld issues. Int J Press Vessel Pip 2014;114–115:76–87. [16] Pandey C, Mahapatra MM, Kumar P, Kumar S, Sirohi S. Effect of post weld heat treatments on microstructure evolution and type IV cracking behavior of the P91 steel welds joint. J Mater Process Technol 2019;266:140–54. [17] Divya M, Das CR, Albert SK, Goyal S, Ganesh P, Kaul R, et al. Influence of welding process on Type IV cracking behavior of P91 steel. Mater Sci Eng A 2014;613:148–58. [18] Francis JA, Cantin GMD, Mazur W, Bhadeshia HKDH. Effects of weld preheat temperature and heat input on type IV failure. Sci Technol Weld Join 2009;14(5):436–42. [19] Mayr P, Martín FM, Albu M, Cerjak H. Correlation of creep strength and microstructural evolution of a boron alloyed 9Cr3W3CoVNb steel in as-received and welded condition. Mater High Temp 2010;27(1):67–72. [20] Das CR, Albert SK, Swaminathan J, Raju S, Bhaduri AK, Murty BS. Transition of crack from type IV to type II resulting from improved utilization of boron in the modified 9Cr-1Mo steel weldment. Metall Mater Trans A 2012;43(10):3724–41. [21] Liu Y, Tsukamoto S, Sawada K, Abe F. Role of boundary strengthening on prevention of type IV failure in high Cr ferritic heat-resistant steels. Metall Mater Trans A 2013;45(3):1306–14. [22] Liu W, Lu F, Wei Y, Ding Y, Wang P, Tang X. Special zone in multi-layer and multipass welded metal and its role in the creep behavior of 9Cr 1Mo welded joint. Mater Des 2016;108:195–206. [23] Chen J-g, Liu Y-c, Liu C-x, Yan B-y, Li H-j. Effects of tantalum on austenitic transformation kinetics of RAFM steel. J Iron Steel Res Int 2017;24(7):705–10. [24] Feng YY, Luo ZA, Zhang DH, Wei J, Wang LJ. The application of welding heat cycle computer software. Mater Sci Forum 2008;575–578:821–6. [25] Chen J, Liu C, Liu Y, Yan B, Li H. Effects of tantalum content on the microstructure and mechanical properties of low-carbon RAFM steel. J Nucl Mater
5. Conclusions Thermal simulation on double-pass welding of a high Cr ferritic steel was conducted. The microstructures and mechanical properties of the samples after double-pass thermal cycle were investigated. The main conclusions are as follows: Thermal cycle almost leads to the increase of ferrite faction and the decrease of martensite fraction, except for the sample after thermal cycle at 900&1000 °C. Thermal cycle results in the refinement of prior austenite, and increase of second-pass thermal cycle temperature leads to grain growth of prior austenite and ferrite. The average size of M23C6 precipitates with the low thermal stability is increased remarkably by thermal cycle. Due to its high thermal stability, the average size of nano-sized MX precipitates remains unchanged basically. M23C6 precipitates are dissolved totally when the second thermal cycle increases to 1100 °C, while MX particles are undissolved until at 1200 °C. Thermal cycle at low temperature (as 900 °C) deteriorates the impact toughness. The sample after double-pass thermal cycle for 900& 1000 °C have the highest impact energy. By further increasing the second-pass thermal cycle temperature, the impact energy is decreased. This is mainly due to the increase of δ-ferrite fraction. Short-term creep tests indicate that, thermal cycle at 900 & 1200 °C leads to the longest creep rupture time, while other samples have the very short rupture time. Coarse M23C6 particles and the heterogeneous alloy elements distribution maybe the main reason. The second thermal cycle at 1200 °C dissolves all the second-phase particles, and homogenize the compositions, promoting the creep properties. The mechanism for that is similar to the type IV cracking. Declaration of interests The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. 15
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carbon microalloyed steels. Metall Mater Trans A 2006;37(7):2147–58. [32] Hyzak JM, Bernstein IM. The role of microstructure on the strength and toughness of fully pearlitic steels. Metall Trans A 1976;7(8):1217–24. [33] Sakthivel T, Vasudevan M, Laha K, Parameswaran P, Chandravathi KS, Panneer Selvi S, et al. Creep rupture behavior of 9Cr–1.8W–0.5Mo–VNb (ASME grade 92) ferritic steel weld joint. Mater Sci Eng A 2014;591:111–20. [34] Liu Y, Tsukamoto S, Shirane T, Abe F. Formation mechanism of type IV failure in high Cr ferritic heat-resistant steel-welded joint. Metall Mater Trans A 2013;44(10):4626–33. [35] Falat L, Čiripová L, Kepič J, Buršík J, Podstranská I. Correlation between microstructure and creep performance of martensitic/austenitic transition weldment in dependence of its post-weld heat treatment. Eng Fail Anal 2014;40:141–52. [36] Zhang W, Wang X, Gong J, Jiang Y, Huang X. Experimental and simulated characterization of creep behavior of P92 steel with prior cyclic loading damage. J Mater Sci Technol 2017;33(12):1540–8.
2016;479:295–301. [26] Schneider A, Inden G. Simulation of the kinetics of precipitation reactions in ferritic steels. Acta Mater 2005;53(2):519–31. [27] Xiao X, Liu G, Hu B, Wang J, Ma W. Microstructure stability of V and Ta microalloyed 12% Cr reduced activation ferrite/martensite steel during long-term aging at 650 C. J Mater Sci Technol 2015;31(3):311–9. [28] Wang P, Lu SP, Xiao NM, Li DZ, Li YY. Effect of delta ferrite on impact properties of low carbon 13Cr–4Ni martensitic stainless steel. Mater Sci Eng A 2010;527(13–14):3210–6. [29] Wang C, Wang M, Shi J, Hui W, Dong H. Effect of microstructural refinement on the toughness of low carbon martensitic steel. Scr Mater 2008;58(6):492–5. [30] Wei Z, Laizhu J, Jincheng H, Hongmei S. Effect of ageing on precipitation and impact energy of 2101 economical duplex stainless steel. Mater Charact 2009;60(1):50–5. [31] Yan W, Shan YY, Yang K. Effect of TiN inclusions on the impact toughness of low-
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