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Influence of welding pass on microstructure and toughness in the reheated zone of multi-pass weld metal of 550 MPa offshore engineering steel
Author’s Accepted Manuscript Influence of welding pass on microstructure and toughness in the reheated zone of multi-pass weld metal of 550MPa offshore engineering steel X.L. Wang, Y.R. Nan, Z.J. Xie, Y.T. Tsai, J.R. Yang, C.J. Shang www.elsevier.com/locate/msea
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S0921-5093(17)30848-1 http://dx.doi.org/10.1016/j.msea.2017.06.081 MSA35219
To appear in: Materials Science & Engineering A Received date: 28 April 2017 Revised date: 19 June 2017 Accepted date: 20 June 2017 Cite this article as: X.L. Wang, Y.R. Nan, Z.J. Xie, Y.T. Tsai, J.R. Yang and C.J. Shang, Influence of welding pass on microstructure and toughness in the reheated zone of multi-pass weld metal of 550MPa offshore engineering steel, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2017.06.081 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Influence of welding pass on microstructure and toughness in the reheated zone of multi-pass weld metal of 550 MPa offshore engineering steel X.L. Wang1, Y.R. Nan1, Z.J. Xie1, Y.T. Tsai2, J.R. Yang2*, C.J. Shang3* 1
School of Materials Science and Engineering, University of Science and Technology Beijing, Beijing, China 2 Department of Materials Science and Engineering, National Taiwan University, Taipei, Taiwan 3 Collaborative Innovation Center of Steel Technology, University of Science and Technology Beijing, Beijing, China
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Corresponding author: Tel: +886 2 33661314, Fax: +886 2 23634562.
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Corresponding author: Tel: +86 010 62322428, Fax: +86 10 62332428.
Abstract The objective of this paper is to study the influence of thermal cycles produced by the latter welding pass on properties of the sub-regions in reheated zone of multi-pass weld metal for a 550 MPa grade offshore engineering steel. A Gleeble-3500 simulator was applied to simulate microstructural evolution in sub-regions of the reheated zone and its effect on the properties. The results indicated that the reheated process changed the prior austenitic morphology from columnar structure to equiaxed structure and similar columnar structure with quasi-polygonal ferrite (QPF) or blocky M-A (martensite/austenite) constituent distributed on the grain boundaries while the matrix microstructure (acicular ferrite) changed slightly. Charpy impact results indicated that WM region (as-deposited) had the highest impact energy. However, the actual impact sample showed lower impact energy because of the machined notch which contained one or more brittle reheated zones. In these brittle reheated zones, the necklace-type M-A constituent as hard phase, along prior columnar or equiaxed austenite grain boundaries, yielded stress concentration obviously and was mainly responsible for lower toughness of the entire weld metal. But fortunately, this influence could be reduced by multiple thermal cycles during multi-pass welding process, due to the degree of decomposition of necklace-type M-A constituent.
Keywords: offshore engineering steel; multi-pass weld metal; reheated zone; acicular ferrite; M-A constituent; low temperature toughness
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1. Introduction The microstructures of multi-pass submerged arc welds generally consist not only of as-deposited weld metal but also of metal which has been reheated, during subsequent passes, to temperatures below Ac1, between Ac1 and Ac3, or above Ac3. This reheated process usually destroys the prior columnar morphology and divides the reheated zone into several sub-regions which resulted in inhomogeneous and complicated microstructures. In previous work [1-3], the welding parameters (i.e., interpass temperature and heat input) have significant effect on microstructure and properties of weld metal. However, the determinant factor that controls the low temperature toughness of the entire weld metal was more like the reheated zone [4]. Because the low-temperature toughness is severely reduced due to the presence of the unfavorable microstructural feature, i.e., necklace-type M-A islands. Previous studies [5-7] show that necklace-type M-A constituents are always formed in intercriticallly reheated coarse grained heat affected zone (ICCGHAZ) in high-strength low-alloyed (HSLA) steels. Moreover, it has been reported [8-10] that the toughness of ICCGHAZ is predominantly controlled by the size and volume fraction of the high carbon M-A products which can act as nucleation sites for crack initiation. Furthermore, the formation of necklace-type M-A (M-A presents as a continuous constituent along the prior austenite grain boundary) is more easy to crack propagation and then forms large cleavage fracture [6], and causes a dramatic loss of the toughness. However, in regard to the multi-pass weld, only few studies focused on the microstructure and impact toughness of the real weld metal because of its complexity. Additionally, the precise definition on the correlation between necklace-type M-A in reheated zone of weld metal and impact toughness is uncertain. Thus, it is very difficult to analyze the effect of a characteristic microstructure on fracture toughness using the true weldment. Moreover, the microstructural evolution of multi-pass weld metal in reheated process is not very clear, especially the transformation process and the effect on the properties. In the present study, the quantitative analysis of the influence of welding thermal cycles produced by reheated processing on mechanical properties is discussed using the method of thermal simulation. In addition, the microstructural evolution in reheated zone of multi-pass
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weld metal is illustrated in the form of a schematic diagram, in terms of the actual and simulated microstructures characterized in present work and elsewhere [11].
2. Experimental 2.1. Material and submerged arc welding conditions The base plate was a quenching and tempering offshore engineering steel with yield strength of 550 MPa and thickness of 36 mm. Welding was performed using gas metal arc welding (GMAW) and submerged arc welding (SAW) according to the welding parameters listed in Table 1. GMAW as back welding, was only used for the first pass, and the welding wire with diameter of 1.2 mm was a conventional wire for commercial use. After back welding, SAW with interpass temperature of 130 °C was designed to the welding experiment. The filler material was Mn-Ni-Mo alloy welding wire with diameter of 4.0 mm. After welding, the compositions of the weld metal and base metal are listed in Table 2. 2.2. Simulated thermal cycles The macro morphology of weld joint is shown in Fig. 1a. In our study, the multi-pass weld metal is firstly divided into two regions, i.e., as-deposited weld metal and reheated weld metal. The as-deposited region (referred as WM) does not undergo any reheated process, while the reheated region is WM region that will subject to the same thermal cycles as heat affected zone (HAZ). In this reheated region, it also can be divided into several sub-regions, i.e., WM-CG, WM-FG, WM-IC, WM-ICCG and WM-ICCG' (WM subject to coarse grained (CG), fine grained (FG), intercritical (IC), and intercritical coarse grained (ICCG) thermal cycles) according to the distance form the fusion line, as shown in Fig. 1b and c. WM-ICCG' region is WM-ICCG region but is tempered one or more times by subsequent welding pass. As these actual reheated zones in multi-pass weld metal are so narrow and the testing position of the impact specimen (referred as WM-C in Fig. 1c) contains several sub-regions mentioned above, it is impossible to correlate the microstructures and mechanical properties quantitatively. To overcome this problem, thermal cycles were simulated by a Gleeble-3500 simulator. Although the simulated method may not fully reflect the real microstructure, it can enlarge the narrow reheated zones and be helpful to study the microstructural evolution and the variation tendency of mechanical properties in real weld joint. The location of the Gleeble
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sample is shown in Fig. 1d, and is simply the weld metal of the second last weld where no real thermal cycles yield, according to the microstructure. To simulate the heat treatment of reheated zone, the time-temperature parameters were deduced from actual welding process, and were given in Fig. 2a. Actually, in this work the peak temperatures using to simulate possible heating histories of WM-CG, WM-FG and WM-IC regions, are from HAZ regions superimposed on the iron-carbon phase diagram [12]. Moreover, a middle temperature (780 °C) between Ac1 and Ac3 (Fig. 2b) was slected as the peak temperature of WM-IC and the second peak temperature of WM-ICCG. Meanwhile, to simplify the experimental process, the thermal cycle of simulated WM-ICCG' region was only designed as three times, and the third peak value as tempering temperature was designed as 600 °C. Additionally, all the cooling rates were calculated from the analytical formulation of Rykalin [13], which is based on conductive heat transfer. 2.3. Mechanical properties Impact toughness was measured at -40 °C according to ASTM: E2298 (length × width × thickness = 55 mm ×10 mm × 10 mm) using standard Charpy v-notch (CVN) test. The samples were from the real weld joint of the second last weld with or without simulated thermal cycles, as shown in Fig. 1c and d. Vickers microhardness across the cross-section (from cap to root) of polished weld joint and of the simulated samples were measured with 1 kgf and a dwell time of 15 s, according to ASTM: E384. 2.4. Microstructure Real welded samples and simulated samples for microstructure observation were cut from the marked area in Fig. 1c and the heat treatment regions respectively, and then mounted and mechanically polished using standard metallographic procedures. The polished specimens for optical microscopy (OM) and scanning electron microscopy (SEM) analysis were etched with 4% nital for 10-15 s. To reveal the morphology and distribution of M-A constituent, the specimens were observed by OM using the method of LePera etchant [14]. Electron back-scattered diffraction (EBSD) was used to analyze the distribution of high angle boundaries (above 15° and 45°) and to explore stress distribution induced by M-A constituent. The samples were electropolished using the solution (glycerol: perchloric acid: alcohol=0.5: 1: 8.5). EBSD scanning was performed at an acceleration voltage of 20 kV, working distance of 4
16mm, title angle of 70°, and step size of 0.15 μm. The fine-scale microstructures were further examined using a transmission electron microscope (TEM, FEI Tecnai G2 F20 FEG-TEM) equipped with an energy-dispersive X-ray spectrometer (EDS) at 200 kV. TEM samples were carried out using 3 mm disks and ground to a thickness of 50 nm, and then twin-jet electropolished with electrolyte of 10% perchloric acid and 90% ethanol.
3. Results and discussion 3.1. Real and simulated microstructure in reheated zone of weld metal The microstructure of as-deposited region (WM) is shown in Fig. 3. It can be clearly seen that, from Fig. 3a, the WM region was dominated by fine acicular ferrite (AF), distinguished by the characteristic ferrite grains of different spatial orientations inside a large prior austenite columnar grain. Large amount of inclusions can be observed, and were possibly the nucleation site of AF. Fig. 3b shows the micrograph which was LePera etched to reveal the distribution of the second phase, i.e., M-A constituent; the M-A constituents appeared as white etched particles, and the volume fraction was measured to be ~10.4%, with an average size of ~0.7 μm. Shown in Figs. 4 are respectively the optical micrographs of the real reheated zone of weld metal (Figs. 4a-e), the correspondingly simulated reheated zone (Figs. 4a1-e1) and the distribution of M-A constituent in simulated reheated zone (Figs. 4a2-e2). It can be seen that the simulated microstructure could more or less reflect the microstructure in corresponding real reheated zone. The subtle difference could be attributed to the accuracy of simulated time-temperatuer parameters and the variation of stress during transformation (greater in real reheated zone [15]). Nevertheless, Gleeble simulations can be helpful in correlating microstructure and mechanical properties, especially low temperature toughness. After thermal (real or simulated) cycles, the prior austenite morphology changed from columnar structure (Fig. 3) to equiaxed structure (Figs. 4a-a2, 4d-d2 and 4e-e2), or to similar columnar structure but with columnar boundaries decorated with quasi-polygonal ferrite (QPF) and blocky M-A constituent (Figs. 4b-b2 and 4c-c2). SEM images of the simulated microstructure are shown in Fig. 5 to reveal the microstructural details more clearly. As shown in Fig. 5b, the microstructures of the sample
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WM-CG were mainly AF and some fine M-A constituents within the large prior austenite grain, and comparing to Fig. 5c the sample WM-FG , the main difference is the increase in the fraction of QPF. In sample WM-IC, as shown in Fig. 5d, there was a large amount of blocky M-A, which could also be observed in sample WM-ICCG (Fig. 5e), with the difference being the types of prior austenite grain boundaries; the sample WM-IC is columnar while the sample WM-ICCG is equiaxed. This distinct difference in austenite morphology is due to the reheating process. For these two samples, the previous microstructures (WM and WM-CG) were reheated to an intercritical temperature between Ac1 and Ac3 by subsequent weld bead, which prompted part of the matrix to revert to austenite with inhomogeneity in carbon-content and residual carbides [16]. These reversed austenite islands preferentially nucleate at the prior austenite grain boundary and become enriched in carbon. On cooling, part of the reverted austenite transforms first into bainitic ferrite at relatively high temperature [7] and when cooled to Ms temperature, martensite is obtained (Fig. 5d). Only a small part of austenite is retained at room temperature. Thus, the high-carbon austenite islands transformed to second-phase particles as a "necklace" around the prior austenite grain. Based on LePera etched images, the size and volume fraction of M-A of all samples were measured, as shown in Fig. 6. For WM-IC and WM-ICCG samples, the size and especially the volume fraction of M-A are larger than the other samples; the average size is ~1.1 (~1.4 μm) and the volume fraction is ~18.8% (~19.3%) in WM-IC (WM-ICCG) sample. In addition, it is noteworthy that the size of necklace-type M-A (up to ~2 μm is possible) can be significantly greater than average value. As a result, WM-IC and WM-ICCG should be the weakest region for impact toughness. For sample WM-ICCG', the M-A constituents can decompose during multiple thermal cycles, as shown in Figs. 4e2 and 5f. The overall volume fraction and average size are reduced by tempering (Fig. 6), and we can therefore conclude that the third thermal cycle could be helpful to improve the toughness. TEM is used to further reveal the detailed microstructure of M-A and tempered M-A/martensite, and the results on samples WM-ICCG and WM-ICCG' are shown in Fig. 7. Typical AF nucleated on the Al-Si-Ti-Mn-O complex inclusion can be clearly observed in the interior of a large prior austenite grain, and refines the large austenite grain by the interlocking structures (Fig. 7a and c). While large blocky M-A constituents with the size of 6
~2 μm are mainly distributed on the prior austenite grain boundary, and some of them have the twin sub-structure, i.e., twin martensite (Fig. 7b), which should be transformed at lower temperature. For sample WM-ICCG', some of the M-A constituents decompose into ferrite and carbides during the 600 °C tempering, and as shown in Fig. 7d and e, the precipitates mainly distributed on the interfaces between AF grains and inside the tempered blocky M-A (also shown in Fig. 5f). All of them tended to be spheroidal or rod-like. The size of the spheroidal carbides is smaller than 100 nm, but the size of rod-like carbides can be up to 200 nm in the long axis direction. Selected area (marked circle in Fig. 7g) electron diffraction pattern analysis confirmed that the carbides were M3C-type cementite carbides. Previous studies [17] on the effect of cementite carbides on properties in low carbon steel suggested that the spheroidal nano-sized cementite carbide (˂100 nm) is good for low temperature toughness. Regarding large rod-like carbides, their role on impact toughness is still unclear. However, as the volume fraction is extremely small, it is plausible that toughness deterioration by rod-like carbides are negligible. In this work, the degree of decomposition of M-A constituents should have a decisive effect on the improvement of toughness. In addition, the film-like M-A (thickness of ~70 nm), which can be obtained during the third thermal cycle, can deflect brittle cracks and is good for impact toughness [18]. 3.2. Hardness and stress distribution Fig. 8 shows the from-cap-to-root hardness distribution in the center line of the real weld joint, and the indents are marked in Fig. 1c. The average hardness is 260 HV, but the hardness shows significant deviation, which is associated with the microstructural variations in the real weld joint. By comparing the microstructure of real weld metal (Figs. 3a and 4a-e), the four peaks in hardness profile corresponds to the reheated zone. For further verification, the hardness of all simulated samples are presented in Fig. 9 together with the impact energy obtained at the test temperature of -40 °C. The results of hardness test revealed that the maximum hardness can be obtained in samples WM-IC, WM-ICCG and WM-ICCG', which corresponded to the peaks in Fig. 8. That is to say the hardness of these three regions in reheated zone of real weld metal is higher than the matrix (WM). To further analyze the stress induced by M-A constituent, Kernel average misorientation (KAM) map was used to study the stress distribution [19-21], as shown in Fig. 10c and d, 7
with the distribution of high angle boundary illustrated in Fig. 10a (WM) and b (WM-ICCG). The results indicated that there was no significant decrease in the density of high angle boundaries in WM-ICCG region of multi-pass weld metal. However, the yellow and red colors represent stress concentration due to the high KAM value, and the KAM maps suggested that the highest values are present in the locations of M-A constituents presence, indicating that stress concentration is higher in M-A constituents than in acicular ferrite. Thus, the interactions between these hard regions and the surrounding as-deposited regions (WM) possibly lead to the formation of local brittle zone (LBZ) in weld metal, which, in turn, cause local unstable fracture behavior [22-24], and seriously deteriorate the low temperature toughness. 3.3. Low temperature toughness From the results of impact test (Fig. 9), the as-deposited region (WM) had excellent toughness, and the impact energy at -40 °C could be up to ~120 J while it was only ~62 J for the test sample located at the center of the real weld metal (WM-C). The simulated results showed that samples WM-CG and WM-FG also have a better impact toughness, though they are slightly lower than WM. Because the volume fraction of AF obtained in sample WM is higher than samples WM-CG and WM-FG, which is correlated to the size of prior austenite grains [25,26]. WM region shows relatively large columnar grain and it is more easy to occur intragranular nucleation on inclusions, due to a higher activation energy compared with grain boundary nucleation [25]. That means the number of AF nucleation site in WM region is higher than WM-CG and WM-FG. Moreover, the reason why the impact energy of sample WM-FG is lower than WM-CG should be due to the transformation process with insufficient hardenability as the lower austenitizing temperature [27]. For sample WM-FG, the smaller reversed austenite grains formed at the prior columnar grain boundary during reheating and holing processes, and in subsequent cooling process, they transformed to high temperature products which are evidenced by the presence of QPF in Figs. 4b1 and 5c. Therefore, even though both of the size and fraction of M-A particles in these two samples are similar, the toughness of sample WM-FG is lower than WM-CG. In contrast, the low temperature toughness of samples WM-IC and WM-ICCG are not very satisfactory. The impact energy is ~64 and 67 J, respectively, indicating that the weakest region of multi-pass 8
weld metal is WM-IC and WM-ICCG regions, which contained blocky necklace-type M-A constituents (Fig. 5d and e). However, it is worth noting that the impact energy of sample WM-C is ~62 J, which is a little lower than samples WM-IC and WM-ICCG. This reason could be attributed to the peak temperature which was designed as 780 °C. Li et al. [6] suggested that the lowest toughness will be obtained for simulated ICCGHAZ when the second peak temperature is only a little higher than Ac1, while the toughness can be improved by increasing the second peak temperature from Ac1 to Ac3. Because the size of M-A constituent decreases with increasing of the second peak temperature, especially the temperature up to half of the intercritical region. Meanwhile, the distribution of M-A changes from necklace-type to dispersive type. Therefore, in our studies, samples WM-IC and WM-ICCG did not get the worst toughness. Nevertheless, it can also be concluded that the reason for low impact energy of the real weld metal (WM-C) is that the machined notch contains a brittle zone, i.e., WM-IC and/or WM-ICCG region. For multi-pass welding, as shown in Fig. 9, sample WM-ICCG' (~73 J) had increased toughness comparing to WM-IC and WM-ICCG, indicating that the toughness can be improved by later thermal cycles due to the decomposition of brittle M-A constituents, especially the necklace-type M-A constituents. 3.4 Microstructure-mechanical peroperties relationship From microstructure and mechanical property data, it is clear that the microstructure of multi-pass weld metal is complex and diverse, and it has a decisive effect on the mechanical properties. According to the actual and simulated microstructures characterized in present work and elsewhere [11], the schematic diagram of microstructural evolution in the reheated zone of multi-pass weld metal is shown in Fig. 11. The microstructure of WM region (as-deposited) was composed mainly of AF in large columnar prior austenite grain, which has high strength and good Charpy impact toughness (Fig. 9), attributed to its small effective packet size and fine interlocking microstructure with high angle boundaries between ferrite plates (Fig. 10a). Moreover, the uniform distribution of fine M-A particles (~0.7 μm) have no significant effect on toughness [9], but can enhance the strength of matrix. However, in reheated zone of multi-pass weld metal, the microstructural morphology has been changed based on the different welding thermal cycles. The main changes in reheated zone are the prior austenitic morphology and the formation of necklace-type M-A. Although there was 9
slightly change to the matrix microstructure, the reheated process promoted the formation of large amount of necklace-type M-A constituent in WM-IC and WM-ICCG regions. Moreover, the hardness results indicated that local hardness in the reheated zone of weld metal (Fig. 8) is mainly caused by increase of blocky necklace-type M-A constituents. Davis et al. [5] reported that when there is necklace-type M-A constituent, there is “overlap of transformation-induced residual tensile stresses” and enhancement of stress concentration due to strength mismatch between M-A and matrix during deformation. Thus, M-A constituent present as a continuous and interconnected structure is detrimental to toughness. This result can also be concluded from the analysis of high angle boundary and stress distribution (Fig. 10). The reheated process has hardly changed the density of high angle boundary, but lead to stress concentration present in the location of M-A constituents presence, especially necklace-type M-A constituent. That is to say the decrease of low temperature toughness was mainly caused by the necklace-type M-A products even if large amount of high angle boundary could be obtained in the matrix. Figs. 12 and 13 present the fracture surface of all samples. From Fig. 12a-c it can be seen that the fracture surface was characterized by quasi-cleavage and intergranular fracture with large facets in sample WM-C with notch contained reheated zone. Moreover, Fig. 12d shows that the crack propagates along the original columnar grain boundaries, where large amount of M-A constituents formed. Previous study [5] suggested that the crack is most likely to initiate from the brittle M-A constituent. In the case where the initiation sites were found, they were observed to have originated from brittle M-A constituents along prior austenite grain boundaries. For WM-IC or WM-ICCG regions, the necklace-type M-A constituents increase the stress concentration (Fig. 10c and d) and provide more nucleation sites for brittle fracture on grain boundaries. Then many microcracks nucleate in the front of the crack tip through cracking of M-A or debonding of M-A from the matrix [5,9]. Finally, the cleavage or intergranular fracture happen through interconnection of microcracks. However, for simulated samples, only quasi-cleavage fracture can be found in samples WM-IC and WM-ICCG. That is to say the actual regions of WM-IC and WM-ICCG are more brittle than simulated samples. Because the size of britile phase (necklace-type M-A) in these two actual brittle regions increases gradually with the peak temperature dropping to Ac1 [6], indicating that the peak 10
temperature (780 °C) used for simulated samples is not the worst. For other simulated samples, many small ductile dimples with different amount can be obtained. The dimples on the fracture surface in sample WM-ICCG', which mainly formed along the tear ridges, should be attributed to the nano cementite and film-like M-A from the tempered necklace-type M-A/martensite. Thus, the formation of dimples and tear ridges can increase the impact energy due to energy dissipating micromechanisms [28]. From the above analysis, it can be concluded that the thermal cycling process produced by the latter welding pass determines the microstructure and mechanical properties of the reheated zone. The lower impact energy of the real weld metal was attributed to the testing position of the impact specimen contained WM (as-deposited) and reheated zone (Fig. 1c), where WM, WM-CG and WM-FG usually exhibited higher toughness, but WM-IC and WM-ICCG displayed inferior toughness because of the formation of necklace-type M-A constituents. However, from the result of the multiple thermal cycles (sample WM-ICCG') it can be seen that changing the thermal cycling process could effectively improve the impact toughness of the reheated zone, which was due to the decomposition of necklace-type M-A constituent. Additionally, others [29,30] indicated that the necklace-type M-A product can be suppressed through refinement of austenite grain size, or via addition of Ni which can promote complete transformation of austenite during welding in the intercritical temperature between Ac1 and Ac3. Thus, for improving low temperature toughness of multi-pass weld metal, the effective way is to control the size and fraction of necklace-type M-A constituent by changing the chemical composition of welding wire or the welding parameters between different welding pass. Moreover, Wang et al. [4] suggested that the toughness of multi-pass weld metal can also be improved by intercritical heat treatment.
4. Conclusions (1) In reheated zone of multi-pass weld metal, the prior austenitic morphology changed from columnar structure to equiaxed structure and similar columnar structure with quasi-polygonal ferrite or blocky M-A constituent distributed on the grain boundaries while the matrix microstructure (acicular ferrite) changed slightly. Necklace-type M-A constituents formed in WM-IC and WM-ICCG regions irrespective of the real or simulated samples.
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(2) The WM region (without reheated process) showed the highest impact toughness, while lower toughness was recorded for the actual sample WM-C and the simulated samples of WM-IC and WM-ICCG. Results showed that the lower impact energy of the real weld metal was attributed to the machined notch which contained the brittle zones of WM-IC and/or WM-ICCG. The WM-IC and WM-ICCG regions with higher hardness and stress concentration in necklace-type M-A constituent, exhibited a greater loss in absorbed energy than other reheated regions. (3) Multiple welding thermal cycles could improve the impact toughness which was contributed to the degree of decomposition of necklace-type M-A constituent.
Acknowledgement This work is financially supported by the Natural Science Foundation of China (51371001). Thanks for Mr. Min Li from Technology Center, Jinan Iron & Steel Co., Ltd., for the operation of welding experiment.
References [1] S. Dirk, K. Thomas, K. Arne, Adv. Mater. Res. 996 (2014) 475-480. [2] L.Y. Lan, X.W. Kong, C.L. Qiu, D.W. Zhao, Mater. Des. 90 (2016) 488-498. [3] Y. Peng, A.H. Wang, H.J. Xiao, Z.L. Tian, Mater. Sci. Forum. 706-709 (2012) 2246-2252. [4] X.L. Wang, X.M. Wang, C.J. Shang, R.D.K. Misra, Mater. Sci. Eng. A 649 (2016) 282-292. [5] C.L. Davis, J.E. King, Metall. Mater. Trans. A 25 (1994) 563-573. [6] X.D. Li, Y.R. Fan, X.P. Ma, S.V. Subramanian, C.J. Shang, Mater. Des. 67 (2015) 457-463. [7] Y. You, C.J. Shang, L. Chen, S.V. Subramanian, Mater. Des. 43 (2013) 485-491. [8] S. Lee, B.C. Kim, D. Kwon, Metall. Mater. Trans. A 23A (1992) 2803-2816. [9] A. Lambert-Perlade, T. Sturel, A.F. Gourgues, J. Besson, A. Pineau, Metall. Mater. Trans. A 35 (2004) 1039-1053. [10] B.C. Kim, S. Lee, N.J. Kim, D.Y. Lee, Metall. Trans. A 22A (1991) 139-149. [11] S.S. Babu, Curr. Opin. Solid. St. Mater. Sci. 8 (2004) 267-278. [12] D.P. Fairchild, N.V. Bangaru, J.Y. Koo, P.L. Harrison, A. Ozekcin, Weld. J. 70 (1991) 321-329. [13] N.N. Rykalin, Calculation of Heat Processes inWelding, U.S.S.R., Moscow, 1960. [14] F.S. LePera, Metallography. 12 (1979) 263-268. [15] J.G. Nawrocki, J.N. Dupont, C.V. Robino, J.D. Puskar, A.R. Marder, Weld. J., 79 (2000) 355-362. [16] R. Danilo, G. Vladimir, Mater. Manuf. Process. 20 (2005) 833-849. 12
[17] Z.J. Xie, Y.P. Fang, G. Han, H. Guo, R.D.K. Misra, C.J. Shang, Mater. Sci. Eng. A 618 (2014) 112-117. [18] Y. Zhong, F.R. Xiao, J.W. Zhang, Y.Y. Shan, W. Wang, K. Yang, Acta. Mater. 54 (2006) 435-443. [19] L.P. Kubin, A. Mortensen, Scr. Mater. 48 (2003) 119-125. [20] A. Ramazani, S. Bruehl, T. Gerber, W. Bleck, U. Prahl, Mater. Des. 57 (2014) 479-486. [21] K. Radwański, Arch. Civ. Mech. Eng. 16 (2016) 282-293. [22] G.M. Evans, Weld. J. 74 (1995) 249-261. [23] R. Danilo, G. Vladimir, Mater. Manuf. Process. 20 (2005) 833-849. [24] F. Matsuda, K. Ikeuchi, Y. Fukada, Y. Horii, H. Okada, T. Shiwaku, C. Shiga, S. Suzuki, Trans. JWRI. 24 (1995) 1-24. [25] J.R. Yang, H.K.D.H. Bhadeshia, J. Mater. Sci. 26 (1991) 839-845. [26] J.R. Yang, C.C. Yang, C.Y. Huang, J. Mater. Sci. Lett. 11 (1992) 1145-1146. [27] X.C. Li, D.X. Xia, X.L. Wang. X.M. Wang, C.J. Shang, Sci. China: Technol. Sci. 56 (2013) 66-70. [28] L.Y. Lan, C.L. Qiu, D.W. Zhao, X.H. Guo, L.X. Du, Mater. Sci. Eng. A 529 (2011) 192-200. [29] X.D. Li, X.P. Ma, S.V. Subramanian, C.J. Shang, R.D.K. Misra, Mater. Sci. Eng. A 616 (2014) 141-147. [30] Y. You, C.J. Shang, S.V. Subramanian, Met. Mater. Int. 4 (2014) 659-668.
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Fig. 1. (a) Macro morphology of weld joint, (b) magnified view of local zone in (c) (OM observation area), (c) location of the measurement of mechanical properties in the test samples and (d) location of the samples to simulate the reheated zone of weld metal in (b).
Fig. 2. (a) Schematic diagram of thermal cycles to simulate reheated zone of weld metal and (b) dilatometric curve of weld metal to test Ac1 and Ac3 temperatures.
Fig. 3. Microstructure of WM (as-deposited). (a) optical micrograph, (b) LePera etched micrograph. (PCGB- prior columnar grain boundary)
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Fig. 4. (a-e) real microstructure of weld metal in reheated zone, (a1-e1) simulated microstructure of weld metal in reheated zone and (a2-e2) distribution of M-A constituent of all simulated samples. (a, a1, a2) WM-CG, (b, b1, b2) WM-FG, (c, c1, c2) WM-IC, (d, d1, d2) WM-ICCG and (e, e1, e2) WM-ICCG'. (PAGB- prior equiaxed-austenite grain boundary) 15
Fig. 5. SEM images of the simulated microstructure. (a) WM, (b) WM-CG, (c) WM-FG, (d) WM-IC, (e) WM-ICCG and (f) WM-ICCG'.
Fig. 6. Volume fraction and average size of M-A constituent of all simulated samples.
Fig. 7. TEM micrograph of simulated samples WM-ICCG (a-c) and WM-ICCG' (d-g) showing (a) necklace-type M-A, (b) M-A with twin sub-structure, (c) EDS result of the inclusion in (a), (d) and (e) tempered microstructure, (f) and (g) magnified view of M-A film and cementite in (e). (α- ferrite, θ- cementite) 16
Fig. 8. Hardness of the center line from the cap to the root in real weld metal.
Fig. 9. -40 °C impact toughness and hardness of the simulated samples.
Fig. 10. EBSD showing the distribution of high angle boundary and stress distribution of sample (a) WM and (b-d) WM-ICCG. (a,b) boundary distribution (above 15° and 45°) shown by black and yellow lines, (c) BC (Band contrast) + KAM (Kernel average misorientation) map, (d) magnified view of the marked area in (c). 17
Fig. 11. Schematic diagram adapted from Babu [11] to show the microstructural evolution in reheated zone of multi-pass weld metal.
Fig. 12. SEM and OM micrographs showing impact fracture surface morphologies and secondary cracks propagation paths underneath the fracture surface of sample WM-C. (a) low magnification fractograph, (b) quasi-cleavage fracture, (c) intergranular fracture and (d) secondary crack initiation and propagation along the prior columnar grain boundary.
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Fig. 13. SEM micrographs showing impact fracture surface morphologies of simulated samples at the test temperature of -40°C. (a) WM, (b) WM-CG, (c) WM-FG, (d) WM-IC, (e) WM-ICCG and (f) WM-ICCG'.
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Table 1 Welding process parameters. Current (A)
GMAW
1.2
250
28
30
12.6
12
SAW
4.0
700
34
28
45.9
130
Method Backing weld Filling weld
Welding variables
Wire diameter (mm)
Interpass Heat input temperature Voltage Speed (kJ/cm) (°C) (V) (cm/min)
Table 2 Chemical compositions of base metal and weld metal (wt%). C
Si
Base 0.063 0.26 metal Weld 0.077 0.49 metal
Mn
Ni
Mo
Nb
Cu
Cr
Ti
1.38 0.65 0.22 0.044 0.35 0.40 0.017
B
N
O
-
-
-
1.76 1.41 0.32 0.010 0.18 0.12 0.013 0.0054 0.072 0.032
20
PII: DOI: Reference:
S0921-5093(17)30848-1 http://dx.doi.org/10.1016/j.msea.2017.06.081 MSA35219
To appear in: Materials Science & Engineering A Received date: 28 April 2017 Revised date: 19 June 2017 Accepted date: 20 June 2017 Cite this article as: X.L. Wang, Y.R. Nan, Z.J. Xie, Y.T. Tsai, J.R. Yang and C.J. Shang, Influence of welding pass on microstructure and toughness in the reheated zone of multi-pass weld metal of 550MPa offshore engineering steel, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2017.06.081 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Influence of welding pass on microstructure and toughness in the reheated zone of multi-pass weld metal of 550 MPa offshore engineering steel X.L. Wang1, Y.R. Nan1, Z.J. Xie1, Y.T. Tsai2, J.R. Yang2*, C.J. Shang3* 1
School of Materials Science and Engineering, University of Science and Technology Beijing, Beijing, China 2 Department of Materials Science and Engineering, National Taiwan University, Taipei, Taiwan 3 Collaborative Innovation Center of Steel Technology, University of Science and Technology Beijing, Beijing, China
[email protected] [email protected] *
Corresponding author: Tel: +886 2 33661314, Fax: +886 2 23634562.
*
Corresponding author: Tel: +86 010 62322428, Fax: +86 10 62332428.
Abstract The objective of this paper is to study the influence of thermal cycles produced by the latter welding pass on properties of the sub-regions in reheated zone of multi-pass weld metal for a 550 MPa grade offshore engineering steel. A Gleeble-3500 simulator was applied to simulate microstructural evolution in sub-regions of the reheated zone and its effect on the properties. The results indicated that the reheated process changed the prior austenitic morphology from columnar structure to equiaxed structure and similar columnar structure with quasi-polygonal ferrite (QPF) or blocky M-A (martensite/austenite) constituent distributed on the grain boundaries while the matrix microstructure (acicular ferrite) changed slightly. Charpy impact results indicated that WM region (as-deposited) had the highest impact energy. However, the actual impact sample showed lower impact energy because of the machined notch which contained one or more brittle reheated zones. In these brittle reheated zones, the necklace-type M-A constituent as hard phase, along prior columnar or equiaxed austenite grain boundaries, yielded stress concentration obviously and was mainly responsible for lower toughness of the entire weld metal. But fortunately, this influence could be reduced by multiple thermal cycles during multi-pass welding process, due to the degree of decomposition of necklace-type M-A constituent.
Keywords: offshore engineering steel; multi-pass weld metal; reheated zone; acicular ferrite; M-A constituent; low temperature toughness
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1. Introduction The microstructures of multi-pass submerged arc welds generally consist not only of as-deposited weld metal but also of metal which has been reheated, during subsequent passes, to temperatures below Ac1, between Ac1 and Ac3, or above Ac3. This reheated process usually destroys the prior columnar morphology and divides the reheated zone into several sub-regions which resulted in inhomogeneous and complicated microstructures. In previous work [1-3], the welding parameters (i.e., interpass temperature and heat input) have significant effect on microstructure and properties of weld metal. However, the determinant factor that controls the low temperature toughness of the entire weld metal was more like the reheated zone [4]. Because the low-temperature toughness is severely reduced due to the presence of the unfavorable microstructural feature, i.e., necklace-type M-A islands. Previous studies [5-7] show that necklace-type M-A constituents are always formed in intercriticallly reheated coarse grained heat affected zone (ICCGHAZ) in high-strength low-alloyed (HSLA) steels. Moreover, it has been reported [8-10] that the toughness of ICCGHAZ is predominantly controlled by the size and volume fraction of the high carbon M-A products which can act as nucleation sites for crack initiation. Furthermore, the formation of necklace-type M-A (M-A presents as a continuous constituent along the prior austenite grain boundary) is more easy to crack propagation and then forms large cleavage fracture [6], and causes a dramatic loss of the toughness. However, in regard to the multi-pass weld, only few studies focused on the microstructure and impact toughness of the real weld metal because of its complexity. Additionally, the precise definition on the correlation between necklace-type M-A in reheated zone of weld metal and impact toughness is uncertain. Thus, it is very difficult to analyze the effect of a characteristic microstructure on fracture toughness using the true weldment. Moreover, the microstructural evolution of multi-pass weld metal in reheated process is not very clear, especially the transformation process and the effect on the properties. In the present study, the quantitative analysis of the influence of welding thermal cycles produced by reheated processing on mechanical properties is discussed using the method of thermal simulation. In addition, the microstructural evolution in reheated zone of multi-pass
2
weld metal is illustrated in the form of a schematic diagram, in terms of the actual and simulated microstructures characterized in present work and elsewhere [11].
2. Experimental 2.1. Material and submerged arc welding conditions The base plate was a quenching and tempering offshore engineering steel with yield strength of 550 MPa and thickness of 36 mm. Welding was performed using gas metal arc welding (GMAW) and submerged arc welding (SAW) according to the welding parameters listed in Table 1. GMAW as back welding, was only used for the first pass, and the welding wire with diameter of 1.2 mm was a conventional wire for commercial use. After back welding, SAW with interpass temperature of 130 °C was designed to the welding experiment. The filler material was Mn-Ni-Mo alloy welding wire with diameter of 4.0 mm. After welding, the compositions of the weld metal and base metal are listed in Table 2. 2.2. Simulated thermal cycles The macro morphology of weld joint is shown in Fig. 1a. In our study, the multi-pass weld metal is firstly divided into two regions, i.e., as-deposited weld metal and reheated weld metal. The as-deposited region (referred as WM) does not undergo any reheated process, while the reheated region is WM region that will subject to the same thermal cycles as heat affected zone (HAZ). In this reheated region, it also can be divided into several sub-regions, i.e., WM-CG, WM-FG, WM-IC, WM-ICCG and WM-ICCG' (WM subject to coarse grained (CG), fine grained (FG), intercritical (IC), and intercritical coarse grained (ICCG) thermal cycles) according to the distance form the fusion line, as shown in Fig. 1b and c. WM-ICCG' region is WM-ICCG region but is tempered one or more times by subsequent welding pass. As these actual reheated zones in multi-pass weld metal are so narrow and the testing position of the impact specimen (referred as WM-C in Fig. 1c) contains several sub-regions mentioned above, it is impossible to correlate the microstructures and mechanical properties quantitatively. To overcome this problem, thermal cycles were simulated by a Gleeble-3500 simulator. Although the simulated method may not fully reflect the real microstructure, it can enlarge the narrow reheated zones and be helpful to study the microstructural evolution and the variation tendency of mechanical properties in real weld joint. The location of the Gleeble
3
sample is shown in Fig. 1d, and is simply the weld metal of the second last weld where no real thermal cycles yield, according to the microstructure. To simulate the heat treatment of reheated zone, the time-temperature parameters were deduced from actual welding process, and were given in Fig. 2a. Actually, in this work the peak temperatures using to simulate possible heating histories of WM-CG, WM-FG and WM-IC regions, are from HAZ regions superimposed on the iron-carbon phase diagram [12]. Moreover, a middle temperature (780 °C) between Ac1 and Ac3 (Fig. 2b) was slected as the peak temperature of WM-IC and the second peak temperature of WM-ICCG. Meanwhile, to simplify the experimental process, the thermal cycle of simulated WM-ICCG' region was only designed as three times, and the third peak value as tempering temperature was designed as 600 °C. Additionally, all the cooling rates were calculated from the analytical formulation of Rykalin [13], which is based on conductive heat transfer. 2.3. Mechanical properties Impact toughness was measured at -40 °C according to ASTM: E2298 (length × width × thickness = 55 mm ×10 mm × 10 mm) using standard Charpy v-notch (CVN) test. The samples were from the real weld joint of the second last weld with or without simulated thermal cycles, as shown in Fig. 1c and d. Vickers microhardness across the cross-section (from cap to root) of polished weld joint and of the simulated samples were measured with 1 kgf and a dwell time of 15 s, according to ASTM: E384. 2.4. Microstructure Real welded samples and simulated samples for microstructure observation were cut from the marked area in Fig. 1c and the heat treatment regions respectively, and then mounted and mechanically polished using standard metallographic procedures. The polished specimens for optical microscopy (OM) and scanning electron microscopy (SEM) analysis were etched with 4% nital for 10-15 s. To reveal the morphology and distribution of M-A constituent, the specimens were observed by OM using the method of LePera etchant [14]. Electron back-scattered diffraction (EBSD) was used to analyze the distribution of high angle boundaries (above 15° and 45°) and to explore stress distribution induced by M-A constituent. The samples were electropolished using the solution (glycerol: perchloric acid: alcohol=0.5: 1: 8.5). EBSD scanning was performed at an acceleration voltage of 20 kV, working distance of 4
16mm, title angle of 70°, and step size of 0.15 μm. The fine-scale microstructures were further examined using a transmission electron microscope (TEM, FEI Tecnai G2 F20 FEG-TEM) equipped with an energy-dispersive X-ray spectrometer (EDS) at 200 kV. TEM samples were carried out using 3 mm disks and ground to a thickness of 50 nm, and then twin-jet electropolished with electrolyte of 10% perchloric acid and 90% ethanol.
3. Results and discussion 3.1. Real and simulated microstructure in reheated zone of weld metal The microstructure of as-deposited region (WM) is shown in Fig. 3. It can be clearly seen that, from Fig. 3a, the WM region was dominated by fine acicular ferrite (AF), distinguished by the characteristic ferrite grains of different spatial orientations inside a large prior austenite columnar grain. Large amount of inclusions can be observed, and were possibly the nucleation site of AF. Fig. 3b shows the micrograph which was LePera etched to reveal the distribution of the second phase, i.e., M-A constituent; the M-A constituents appeared as white etched particles, and the volume fraction was measured to be ~10.4%, with an average size of ~0.7 μm. Shown in Figs. 4 are respectively the optical micrographs of the real reheated zone of weld metal (Figs. 4a-e), the correspondingly simulated reheated zone (Figs. 4a1-e1) and the distribution of M-A constituent in simulated reheated zone (Figs. 4a2-e2). It can be seen that the simulated microstructure could more or less reflect the microstructure in corresponding real reheated zone. The subtle difference could be attributed to the accuracy of simulated time-temperatuer parameters and the variation of stress during transformation (greater in real reheated zone [15]). Nevertheless, Gleeble simulations can be helpful in correlating microstructure and mechanical properties, especially low temperature toughness. After thermal (real or simulated) cycles, the prior austenite morphology changed from columnar structure (Fig. 3) to equiaxed structure (Figs. 4a-a2, 4d-d2 and 4e-e2), or to similar columnar structure but with columnar boundaries decorated with quasi-polygonal ferrite (QPF) and blocky M-A constituent (Figs. 4b-b2 and 4c-c2). SEM images of the simulated microstructure are shown in Fig. 5 to reveal the microstructural details more clearly. As shown in Fig. 5b, the microstructures of the sample
5
WM-CG were mainly AF and some fine M-A constituents within the large prior austenite grain, and comparing to Fig. 5c the sample WM-FG , the main difference is the increase in the fraction of QPF. In sample WM-IC, as shown in Fig. 5d, there was a large amount of blocky M-A, which could also be observed in sample WM-ICCG (Fig. 5e), with the difference being the types of prior austenite grain boundaries; the sample WM-IC is columnar while the sample WM-ICCG is equiaxed. This distinct difference in austenite morphology is due to the reheating process. For these two samples, the previous microstructures (WM and WM-CG) were reheated to an intercritical temperature between Ac1 and Ac3 by subsequent weld bead, which prompted part of the matrix to revert to austenite with inhomogeneity in carbon-content and residual carbides [16]. These reversed austenite islands preferentially nucleate at the prior austenite grain boundary and become enriched in carbon. On cooling, part of the reverted austenite transforms first into bainitic ferrite at relatively high temperature [7] and when cooled to Ms temperature, martensite is obtained (Fig. 5d). Only a small part of austenite is retained at room temperature. Thus, the high-carbon austenite islands transformed to second-phase particles as a "necklace" around the prior austenite grain. Based on LePera etched images, the size and volume fraction of M-A of all samples were measured, as shown in Fig. 6. For WM-IC and WM-ICCG samples, the size and especially the volume fraction of M-A are larger than the other samples; the average size is ~1.1 (~1.4 μm) and the volume fraction is ~18.8% (~19.3%) in WM-IC (WM-ICCG) sample. In addition, it is noteworthy that the size of necklace-type M-A (up to ~2 μm is possible) can be significantly greater than average value. As a result, WM-IC and WM-ICCG should be the weakest region for impact toughness. For sample WM-ICCG', the M-A constituents can decompose during multiple thermal cycles, as shown in Figs. 4e2 and 5f. The overall volume fraction and average size are reduced by tempering (Fig. 6), and we can therefore conclude that the third thermal cycle could be helpful to improve the toughness. TEM is used to further reveal the detailed microstructure of M-A and tempered M-A/martensite, and the results on samples WM-ICCG and WM-ICCG' are shown in Fig. 7. Typical AF nucleated on the Al-Si-Ti-Mn-O complex inclusion can be clearly observed in the interior of a large prior austenite grain, and refines the large austenite grain by the interlocking structures (Fig. 7a and c). While large blocky M-A constituents with the size of 6
~2 μm are mainly distributed on the prior austenite grain boundary, and some of them have the twin sub-structure, i.e., twin martensite (Fig. 7b), which should be transformed at lower temperature. For sample WM-ICCG', some of the M-A constituents decompose into ferrite and carbides during the 600 °C tempering, and as shown in Fig. 7d and e, the precipitates mainly distributed on the interfaces between AF grains and inside the tempered blocky M-A (also shown in Fig. 5f). All of them tended to be spheroidal or rod-like. The size of the spheroidal carbides is smaller than 100 nm, but the size of rod-like carbides can be up to 200 nm in the long axis direction. Selected area (marked circle in Fig. 7g) electron diffraction pattern analysis confirmed that the carbides were M3C-type cementite carbides. Previous studies [17] on the effect of cementite carbides on properties in low carbon steel suggested that the spheroidal nano-sized cementite carbide (˂100 nm) is good for low temperature toughness. Regarding large rod-like carbides, their role on impact toughness is still unclear. However, as the volume fraction is extremely small, it is plausible that toughness deterioration by rod-like carbides are negligible. In this work, the degree of decomposition of M-A constituents should have a decisive effect on the improvement of toughness. In addition, the film-like M-A (thickness of ~70 nm), which can be obtained during the third thermal cycle, can deflect brittle cracks and is good for impact toughness [18]. 3.2. Hardness and stress distribution Fig. 8 shows the from-cap-to-root hardness distribution in the center line of the real weld joint, and the indents are marked in Fig. 1c. The average hardness is 260 HV, but the hardness shows significant deviation, which is associated with the microstructural variations in the real weld joint. By comparing the microstructure of real weld metal (Figs. 3a and 4a-e), the four peaks in hardness profile corresponds to the reheated zone. For further verification, the hardness of all simulated samples are presented in Fig. 9 together with the impact energy obtained at the test temperature of -40 °C. The results of hardness test revealed that the maximum hardness can be obtained in samples WM-IC, WM-ICCG and WM-ICCG', which corresponded to the peaks in Fig. 8. That is to say the hardness of these three regions in reheated zone of real weld metal is higher than the matrix (WM). To further analyze the stress induced by M-A constituent, Kernel average misorientation (KAM) map was used to study the stress distribution [19-21], as shown in Fig. 10c and d, 7
with the distribution of high angle boundary illustrated in Fig. 10a (WM) and b (WM-ICCG). The results indicated that there was no significant decrease in the density of high angle boundaries in WM-ICCG region of multi-pass weld metal. However, the yellow and red colors represent stress concentration due to the high KAM value, and the KAM maps suggested that the highest values are present in the locations of M-A constituents presence, indicating that stress concentration is higher in M-A constituents than in acicular ferrite. Thus, the interactions between these hard regions and the surrounding as-deposited regions (WM) possibly lead to the formation of local brittle zone (LBZ) in weld metal, which, in turn, cause local unstable fracture behavior [22-24], and seriously deteriorate the low temperature toughness. 3.3. Low temperature toughness From the results of impact test (Fig. 9), the as-deposited region (WM) had excellent toughness, and the impact energy at -40 °C could be up to ~120 J while it was only ~62 J for the test sample located at the center of the real weld metal (WM-C). The simulated results showed that samples WM-CG and WM-FG also have a better impact toughness, though they are slightly lower than WM. Because the volume fraction of AF obtained in sample WM is higher than samples WM-CG and WM-FG, which is correlated to the size of prior austenite grains [25,26]. WM region shows relatively large columnar grain and it is more easy to occur intragranular nucleation on inclusions, due to a higher activation energy compared with grain boundary nucleation [25]. That means the number of AF nucleation site in WM region is higher than WM-CG and WM-FG. Moreover, the reason why the impact energy of sample WM-FG is lower than WM-CG should be due to the transformation process with insufficient hardenability as the lower austenitizing temperature [27]. For sample WM-FG, the smaller reversed austenite grains formed at the prior columnar grain boundary during reheating and holing processes, and in subsequent cooling process, they transformed to high temperature products which are evidenced by the presence of QPF in Figs. 4b1 and 5c. Therefore, even though both of the size and fraction of M-A particles in these two samples are similar, the toughness of sample WM-FG is lower than WM-CG. In contrast, the low temperature toughness of samples WM-IC and WM-ICCG are not very satisfactory. The impact energy is ~64 and 67 J, respectively, indicating that the weakest region of multi-pass 8
weld metal is WM-IC and WM-ICCG regions, which contained blocky necklace-type M-A constituents (Fig. 5d and e). However, it is worth noting that the impact energy of sample WM-C is ~62 J, which is a little lower than samples WM-IC and WM-ICCG. This reason could be attributed to the peak temperature which was designed as 780 °C. Li et al. [6] suggested that the lowest toughness will be obtained for simulated ICCGHAZ when the second peak temperature is only a little higher than Ac1, while the toughness can be improved by increasing the second peak temperature from Ac1 to Ac3. Because the size of M-A constituent decreases with increasing of the second peak temperature, especially the temperature up to half of the intercritical region. Meanwhile, the distribution of M-A changes from necklace-type to dispersive type. Therefore, in our studies, samples WM-IC and WM-ICCG did not get the worst toughness. Nevertheless, it can also be concluded that the reason for low impact energy of the real weld metal (WM-C) is that the machined notch contains a brittle zone, i.e., WM-IC and/or WM-ICCG region. For multi-pass welding, as shown in Fig. 9, sample WM-ICCG' (~73 J) had increased toughness comparing to WM-IC and WM-ICCG, indicating that the toughness can be improved by later thermal cycles due to the decomposition of brittle M-A constituents, especially the necklace-type M-A constituents. 3.4 Microstructure-mechanical peroperties relationship From microstructure and mechanical property data, it is clear that the microstructure of multi-pass weld metal is complex and diverse, and it has a decisive effect on the mechanical properties. According to the actual and simulated microstructures characterized in present work and elsewhere [11], the schematic diagram of microstructural evolution in the reheated zone of multi-pass weld metal is shown in Fig. 11. The microstructure of WM region (as-deposited) was composed mainly of AF in large columnar prior austenite grain, which has high strength and good Charpy impact toughness (Fig. 9), attributed to its small effective packet size and fine interlocking microstructure with high angle boundaries between ferrite plates (Fig. 10a). Moreover, the uniform distribution of fine M-A particles (~0.7 μm) have no significant effect on toughness [9], but can enhance the strength of matrix. However, in reheated zone of multi-pass weld metal, the microstructural morphology has been changed based on the different welding thermal cycles. The main changes in reheated zone are the prior austenitic morphology and the formation of necklace-type M-A. Although there was 9
slightly change to the matrix microstructure, the reheated process promoted the formation of large amount of necklace-type M-A constituent in WM-IC and WM-ICCG regions. Moreover, the hardness results indicated that local hardness in the reheated zone of weld metal (Fig. 8) is mainly caused by increase of blocky necklace-type M-A constituents. Davis et al. [5] reported that when there is necklace-type M-A constituent, there is “overlap of transformation-induced residual tensile stresses” and enhancement of stress concentration due to strength mismatch between M-A and matrix during deformation. Thus, M-A constituent present as a continuous and interconnected structure is detrimental to toughness. This result can also be concluded from the analysis of high angle boundary and stress distribution (Fig. 10). The reheated process has hardly changed the density of high angle boundary, but lead to stress concentration present in the location of M-A constituents presence, especially necklace-type M-A constituent. That is to say the decrease of low temperature toughness was mainly caused by the necklace-type M-A products even if large amount of high angle boundary could be obtained in the matrix. Figs. 12 and 13 present the fracture surface of all samples. From Fig. 12a-c it can be seen that the fracture surface was characterized by quasi-cleavage and intergranular fracture with large facets in sample WM-C with notch contained reheated zone. Moreover, Fig. 12d shows that the crack propagates along the original columnar grain boundaries, where large amount of M-A constituents formed. Previous study [5] suggested that the crack is most likely to initiate from the brittle M-A constituent. In the case where the initiation sites were found, they were observed to have originated from brittle M-A constituents along prior austenite grain boundaries. For WM-IC or WM-ICCG regions, the necklace-type M-A constituents increase the stress concentration (Fig. 10c and d) and provide more nucleation sites for brittle fracture on grain boundaries. Then many microcracks nucleate in the front of the crack tip through cracking of M-A or debonding of M-A from the matrix [5,9]. Finally, the cleavage or intergranular fracture happen through interconnection of microcracks. However, for simulated samples, only quasi-cleavage fracture can be found in samples WM-IC and WM-ICCG. That is to say the actual regions of WM-IC and WM-ICCG are more brittle than simulated samples. Because the size of britile phase (necklace-type M-A) in these two actual brittle regions increases gradually with the peak temperature dropping to Ac1 [6], indicating that the peak 10
temperature (780 °C) used for simulated samples is not the worst. For other simulated samples, many small ductile dimples with different amount can be obtained. The dimples on the fracture surface in sample WM-ICCG', which mainly formed along the tear ridges, should be attributed to the nano cementite and film-like M-A from the tempered necklace-type M-A/martensite. Thus, the formation of dimples and tear ridges can increase the impact energy due to energy dissipating micromechanisms [28]. From the above analysis, it can be concluded that the thermal cycling process produced by the latter welding pass determines the microstructure and mechanical properties of the reheated zone. The lower impact energy of the real weld metal was attributed to the testing position of the impact specimen contained WM (as-deposited) and reheated zone (Fig. 1c), where WM, WM-CG and WM-FG usually exhibited higher toughness, but WM-IC and WM-ICCG displayed inferior toughness because of the formation of necklace-type M-A constituents. However, from the result of the multiple thermal cycles (sample WM-ICCG') it can be seen that changing the thermal cycling process could effectively improve the impact toughness of the reheated zone, which was due to the decomposition of necklace-type M-A constituent. Additionally, others [29,30] indicated that the necklace-type M-A product can be suppressed through refinement of austenite grain size, or via addition of Ni which can promote complete transformation of austenite during welding in the intercritical temperature between Ac1 and Ac3. Thus, for improving low temperature toughness of multi-pass weld metal, the effective way is to control the size and fraction of necklace-type M-A constituent by changing the chemical composition of welding wire or the welding parameters between different welding pass. Moreover, Wang et al. [4] suggested that the toughness of multi-pass weld metal can also be improved by intercritical heat treatment.
4. Conclusions (1) In reheated zone of multi-pass weld metal, the prior austenitic morphology changed from columnar structure to equiaxed structure and similar columnar structure with quasi-polygonal ferrite or blocky M-A constituent distributed on the grain boundaries while the matrix microstructure (acicular ferrite) changed slightly. Necklace-type M-A constituents formed in WM-IC and WM-ICCG regions irrespective of the real or simulated samples.
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(2) The WM region (without reheated process) showed the highest impact toughness, while lower toughness was recorded for the actual sample WM-C and the simulated samples of WM-IC and WM-ICCG. Results showed that the lower impact energy of the real weld metal was attributed to the machined notch which contained the brittle zones of WM-IC and/or WM-ICCG. The WM-IC and WM-ICCG regions with higher hardness and stress concentration in necklace-type M-A constituent, exhibited a greater loss in absorbed energy than other reheated regions. (3) Multiple welding thermal cycles could improve the impact toughness which was contributed to the degree of decomposition of necklace-type M-A constituent.
Acknowledgement This work is financially supported by the Natural Science Foundation of China (51371001). Thanks for Mr. Min Li from Technology Center, Jinan Iron & Steel Co., Ltd., for the operation of welding experiment.
References [1] S. Dirk, K. Thomas, K. Arne, Adv. Mater. Res. 996 (2014) 475-480. [2] L.Y. Lan, X.W. Kong, C.L. Qiu, D.W. Zhao, Mater. Des. 90 (2016) 488-498. [3] Y. Peng, A.H. Wang, H.J. Xiao, Z.L. Tian, Mater. Sci. Forum. 706-709 (2012) 2246-2252. [4] X.L. Wang, X.M. Wang, C.J. Shang, R.D.K. Misra, Mater. Sci. Eng. A 649 (2016) 282-292. [5] C.L. Davis, J.E. King, Metall. Mater. Trans. A 25 (1994) 563-573. [6] X.D. Li, Y.R. Fan, X.P. Ma, S.V. Subramanian, C.J. Shang, Mater. Des. 67 (2015) 457-463. [7] Y. You, C.J. Shang, L. Chen, S.V. Subramanian, Mater. Des. 43 (2013) 485-491. [8] S. Lee, B.C. Kim, D. Kwon, Metall. Mater. Trans. A 23A (1992) 2803-2816. [9] A. Lambert-Perlade, T. Sturel, A.F. Gourgues, J. Besson, A. Pineau, Metall. Mater. Trans. A 35 (2004) 1039-1053. [10] B.C. Kim, S. Lee, N.J. Kim, D.Y. Lee, Metall. Trans. A 22A (1991) 139-149. [11] S.S. Babu, Curr. Opin. Solid. St. Mater. Sci. 8 (2004) 267-278. [12] D.P. Fairchild, N.V. Bangaru, J.Y. Koo, P.L. Harrison, A. Ozekcin, Weld. J. 70 (1991) 321-329. [13] N.N. Rykalin, Calculation of Heat Processes inWelding, U.S.S.R., Moscow, 1960. [14] F.S. LePera, Metallography. 12 (1979) 263-268. [15] J.G. Nawrocki, J.N. Dupont, C.V. Robino, J.D. Puskar, A.R. Marder, Weld. J., 79 (2000) 355-362. [16] R. Danilo, G. Vladimir, Mater. Manuf. Process. 20 (2005) 833-849. 12
[17] Z.J. Xie, Y.P. Fang, G. Han, H. Guo, R.D.K. Misra, C.J. Shang, Mater. Sci. Eng. A 618 (2014) 112-117. [18] Y. Zhong, F.R. Xiao, J.W. Zhang, Y.Y. Shan, W. Wang, K. Yang, Acta. Mater. 54 (2006) 435-443. [19] L.P. Kubin, A. Mortensen, Scr. Mater. 48 (2003) 119-125. [20] A. Ramazani, S. Bruehl, T. Gerber, W. Bleck, U. Prahl, Mater. Des. 57 (2014) 479-486. [21] K. Radwański, Arch. Civ. Mech. Eng. 16 (2016) 282-293. [22] G.M. Evans, Weld. J. 74 (1995) 249-261. [23] R. Danilo, G. Vladimir, Mater. Manuf. Process. 20 (2005) 833-849. [24] F. Matsuda, K. Ikeuchi, Y. Fukada, Y. Horii, H. Okada, T. Shiwaku, C. Shiga, S. Suzuki, Trans. JWRI. 24 (1995) 1-24. [25] J.R. Yang, H.K.D.H. Bhadeshia, J. Mater. Sci. 26 (1991) 839-845. [26] J.R. Yang, C.C. Yang, C.Y. Huang, J. Mater. Sci. Lett. 11 (1992) 1145-1146. [27] X.C. Li, D.X. Xia, X.L. Wang. X.M. Wang, C.J. Shang, Sci. China: Technol. Sci. 56 (2013) 66-70. [28] L.Y. Lan, C.L. Qiu, D.W. Zhao, X.H. Guo, L.X. Du, Mater. Sci. Eng. A 529 (2011) 192-200. [29] X.D. Li, X.P. Ma, S.V. Subramanian, C.J. Shang, R.D.K. Misra, Mater. Sci. Eng. A 616 (2014) 141-147. [30] Y. You, C.J. Shang, S.V. Subramanian, Met. Mater. Int. 4 (2014) 659-668.
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Fig. 1. (a) Macro morphology of weld joint, (b) magnified view of local zone in (c) (OM observation area), (c) location of the measurement of mechanical properties in the test samples and (d) location of the samples to simulate the reheated zone of weld metal in (b).
Fig. 2. (a) Schematic diagram of thermal cycles to simulate reheated zone of weld metal and (b) dilatometric curve of weld metal to test Ac1 and Ac3 temperatures.
Fig. 3. Microstructure of WM (as-deposited). (a) optical micrograph, (b) LePera etched micrograph. (PCGB- prior columnar grain boundary)
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Fig. 4. (a-e) real microstructure of weld metal in reheated zone, (a1-e1) simulated microstructure of weld metal in reheated zone and (a2-e2) distribution of M-A constituent of all simulated samples. (a, a1, a2) WM-CG, (b, b1, b2) WM-FG, (c, c1, c2) WM-IC, (d, d1, d2) WM-ICCG and (e, e1, e2) WM-ICCG'. (PAGB- prior equiaxed-austenite grain boundary) 15
Fig. 5. SEM images of the simulated microstructure. (a) WM, (b) WM-CG, (c) WM-FG, (d) WM-IC, (e) WM-ICCG and (f) WM-ICCG'.
Fig. 6. Volume fraction and average size of M-A constituent of all simulated samples.
Fig. 7. TEM micrograph of simulated samples WM-ICCG (a-c) and WM-ICCG' (d-g) showing (a) necklace-type M-A, (b) M-A with twin sub-structure, (c) EDS result of the inclusion in (a), (d) and (e) tempered microstructure, (f) and (g) magnified view of M-A film and cementite in (e). (α- ferrite, θ- cementite) 16
Fig. 8. Hardness of the center line from the cap to the root in real weld metal.
Fig. 9. -40 °C impact toughness and hardness of the simulated samples.
Fig. 10. EBSD showing the distribution of high angle boundary and stress distribution of sample (a) WM and (b-d) WM-ICCG. (a,b) boundary distribution (above 15° and 45°) shown by black and yellow lines, (c) BC (Band contrast) + KAM (Kernel average misorientation) map, (d) magnified view of the marked area in (c). 17
Fig. 11. Schematic diagram adapted from Babu [11] to show the microstructural evolution in reheated zone of multi-pass weld metal.
Fig. 12. SEM and OM micrographs showing impact fracture surface morphologies and secondary cracks propagation paths underneath the fracture surface of sample WM-C. (a) low magnification fractograph, (b) quasi-cleavage fracture, (c) intergranular fracture and (d) secondary crack initiation and propagation along the prior columnar grain boundary.
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Fig. 13. SEM micrographs showing impact fracture surface morphologies of simulated samples at the test temperature of -40°C. (a) WM, (b) WM-CG, (c) WM-FG, (d) WM-IC, (e) WM-ICCG and (f) WM-ICCG'.
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Table 1 Welding process parameters. Current (A)
GMAW
1.2
250
28
30
12.6
12
SAW
4.0
700
34
28
45.9
130
Method Backing weld Filling weld
Welding variables
Wire diameter (mm)
Interpass Heat input temperature Voltage Speed (kJ/cm) (°C) (V) (cm/min)
Table 2 Chemical compositions of base metal and weld metal (wt%). C
Si
Base 0.063 0.26 metal Weld 0.077 0.49 metal
Mn
Ni
Mo
Nb
Cu
Cr
Ti
1.38 0.65 0.22 0.044 0.35 0.40 0.017
B
N
O
-
-
-
1.76 1.41 0.32 0.010 0.18 0.12 0.013 0.0054 0.072 0.032
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