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Effects of heat input on microstructure and fracture toughness of simulated coarse-grained heat affected zone for HSLA steels
Materials Characterization 155 (2019) 109818
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Review
Effects of heat input on microstructure and fracture toughness of simulated coarse-grained heat affected zone for HSLA steels ⁎
Xiaocong Yanga,b, Xinjie Dia,b, , Xiuguo Liua,b, Dongpo Wanga,b,c, Chengning Lia,b,
T
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a
School of Materials Science and Engineering, Tianjin University, Tianjin 300350, China Tianjin Key Laboratory of Advanced Joining Technology, Tianjin 300350, China c State Key Laboratory of Metal Material for Marine Equipment and Application, Anshan 114009, China b
A R T I C LE I N FO
A B S T R A C T
Keywords: Coarse-grained heat affected zone HSLA steel Heat input Crack tip opening displacement M-A constituent Fracture toughness
In order to clarify the relationships among welding heat input, microstructural evolution and fracture toughness of coarse-grained heat affected zone (CGHAZ) for HSLA steels, the microstructure of simulated CGHAZ with heat input of 10–100 kJ/cm was quantified and the crack tip opening displacement (CTOD) at −20 °C was evaluated With the increase in heat input, the bainite transformation temperature and duration transformation time increase, which lead to the higher contents of granular bainite and M-A constituents. The fracture toughness of simulated CGHAZ is highly related with the morphologies of M-A constitute. During CTOD test, the island M-A constituents can arrest microcracks and increase the CTOD value, and the fine-stringer ones can deform without excessive damage to fracture toughness. However, the coarse-stringer and massive M-A constituent would crack and debond from the matrix, which seriously decreases the fracture toughness. When the heat input increases from 10 kJ/cm to 100 kJ/cm, the proportion of island M-A constituent decreases from 86% to 35%, while that of coarse-stringer and massive M-A constituent increases from 1% and 6% to 7% and 57%, respectively, which result in the significant deterioration of fracture toughness. The fracture toughness is optimal at heat input of 30 kJ/cm, because the packets of lath bainite are segmented smaller by granular bainite and effectively improved the cracking resistance.
1. Introduction In order to improve the efficiency of transportation and reduce costs, the container ships are required to have larger capacity. The application of high strength low alloy (HSLA) steels, such as EH47, can significant improve the carrying capacity and safety performance of container ships, which have been used in deck and hatch coaming structures [1,2]. The good combination of high strength and toughness for the HSLA steels is well controlled through thermal mechanical controlled processing (TMCP) [3]. However, this excellent balance of strength and toughness will be broken after the high temperature thermal cycling process in welding, such as the reduction of toughness in heat affected zone (HAZ). In generally, coarse-grained HAZ (CGHAZ) is always an attention-attracting local brittle region due to the coarsening prior austenite grain and the formation of M-A constituents, which is highly related to the welding heat input [4–10]. Therefore, the CGHAZ is usually the risk region of fracture failure, especially when the CGHAZ has microcracks produced by cyclic loads of sea wave or welding defect [11], which limits the use of HSLA steels at low
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temperatures. It is known that the hard and brittle M-A constituent can seriously deteriorate fracture toughness. Mohseni et al. concluded that massive M-A constituent decreased the low temperature toughness of CGHAZ seriously, and the fracture initiation occurred preferentially at M-A constituent by a debonding mechanism [12]. Luo et al. pointed out that the slender M-A constituent was more harmful to toughness than massive M-A, and indicated that the massive M-A particles were usually debonded from the matrix while the slender M-A particles were more likely to crack [8]. Chen et al. reported that the micromechanism of cleavage cracking in HAZ was controlled by the size of M-A constituent, and high stress concentrates on the boundary of MA/matrix made it crack or debond due to the heavy deformation of ferrite matrix [13]. However, studies by Lan et al. shows that the small M-A constituent was helpful in improving toughness [10]. Therefore, the effect of M-A constituent on the micromechanism of fracture and fracture toughness is confused, which would be related to the size, content and distribution morphology. In this study, the simulated CGHAZs with different morphology of
Corresponding authors at: School of Materials Science and Engineering, Tianjin University, Tianjin 300350, China. E-mail addresses: [email protected] (X. Di), [email protected] (C. Li).
https://doi.org/10.1016/j.matchar.2019.109818 Received 5 May 2019; Received in revised form 11 July 2019; Accepted 12 July 2019 Available online 12 July 2019 1044-5803/ © 2019 Elsevier Inc. All rights reserved.
Materials Characterization 155 (2019) 109818
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M-A were obtained by welding thermal simulated with different heat input, and the microstructure and crystallographic characteristics of simulated CGHAZs were investigated. At present, the toughness of CGHAZ is mainly evaluated by the Charpy V-notch impact test [4,5,10,14−16]，but it cannot reflect the fracture mechanism of CGHAZ with fatigue cracks or microcracks. Compared with impact toughness, the fracture toughness can effectively characterize the actual service performance of ship hull with microcracks [17]. Therefore, in this paper, the fracture toughness of CGHAZ was measured by three point bend test of crack tip opening displacement (CTOD), which can reflect the plastic deformation performance near the crack tip after cyclic loading. Moreover, the relationship between fracture toughness and microstructure with different heat input was discussed, and the effect of M-A constituent morphologies on fracture micromechanism was discussed in detail.
Table 1 Chemical composition of experimental steel (wt%). C
Si
Mn
P
S
Ni
0.067
0.193
1.34
0.011
0.043
0.831
Cu + Cr + Mo
V + Ti + Nb
≤0.70
≤0.06
Table 2 Mechanical properties of experimental steel. YS Rs/MPa 472
TS Rm/MPa
Yield ratio Rs/Rm
Elongation/%
Hardness/HV
600
0.79
32
200
CTOD value/ mm 0.533
2. Experimental procedure 2.1. Microstructure and properties of the experimental steel The experimental steel is EH47 steel, which is a typical crack arrestability HSLA steels manufactured by TMCP technology. The microstructure of experimental steel is composed of polygonal ferrite (PF), quasi-polygonal ferrite (QPF), acicular ferrite (AF) and bainite (B), presented in Fig. 1. The chemistry and mechanical properties of the experimental steel are listed in Tables 1 and 2, respectively. The yield strength at room temperature is about 472 MPa, and the average CTOD value is 0.533 mm at −20 °C, which indicates that the steel has excellent balance of strength and toughness. In this study, the CTOD test temperature was selected to be −20 °C according to the standard of DNVGL-OS-C401 and DNVGL-OS-C101, which can more effectively reflect the actual service condition and assess the safety of the steel in service.
Fig. 2. Simulated thermal cycle curves with different heat input.
2.2. Welding thermal simulation experiment
50 kJ/cm and 100 kJ/cm, respectively.
Gleeble 3800 simulation machine was used to simulate the CGHAZ in the actual welding. The simulated specimens were cut from a quarter of the 80 mm thick steel plate along the transverse direction and machined into the dimensions 11 mm × 11 mm × 80 mm. The simulated specimens were subjected to different welding thermal cycles according to different actual welding processes. The heat input was 10 kJ/cm, 30 kJ/cm, 50 kJ/cm and 100 kJ/cm, respectively. Simulated thermal cycles are shown in Fig. 2. The three dimensional Rykalin thermodynamic model was used to determine the thermal cycle curve of ultrathick steel plates. The simulated specimens were heated to peak temperature of 1350 °C at a heating rate of 200 °C/s, holding for 1 s at the peak temperature. The cooling time of t8/5 was 4.5 s, 13.4 s, 22.4 s and 44.7 s corresponding to welding heat input of 10 kJ/cm, 30 kJ/cm,
2.3. Fracture toughness test After thermal simulations, the specimens were machined into CTOD specimens with dimension of 10 × 10 × 80 mm according to the standard method of BS7448, and the mechanical notch was located in the thickness direction with a depth of 2.5 mm, shown in Fig. 3. Then the fatigue cracks were prefabricated at room temperature (the maximum value of alternating load is 3.5 kN) until the cracks extend to 1/2 thickness of the specimens. Thereafter, the three point bend test of CTOD was carried out at −20 °C to measure the fracture toughness of simulated CGHAZ. After CTOD test, the specimen was crushed at −80 °C and the fracture morphology was observed. Finally, the validity
Fig. 1. Microstructure of experimental steel: (a) Optical micrograph and (b) SEM micrograph. 2
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laths or within GB ferrite, and displays many different morphologies, including island, stringer and massive. With the increase of heat input, the M-A constituent coarsen obviously and the proportion of island M-A constituent decreased gradually, while the proportion of massive M-A constituent increased continuously. Fig. 5 shows the change of average grain size of prior austenite and Vickers hardness of simulated CGHAZ. It indicates that the average grain size of prior austenite increases with the increase in heat input. The coarsening of prior austenite grain at high heat input is attributed to the increase of duration time at high temperature that the grain can grow easily. In addition, the Vickers hardness decreases with the increase of heat input, which mainly depends on the microstructure and grain size. When the heat input increases, the increase of GB that composed of massive ferrite is the key factor leading to the decrease of hardness. At the same time, the increase of grain size reduces the strengthening effect of effective grain boundary, which results in the decrease of Vickers hardness [14].
Fig. 3. Dimensions of CTOD specimens.
of the test is judged by checking the profiles of all CTOD specimens and ensuring that the crack tip is located in the designated area.
3.2. The crystallographic characteristics of simulated CGHAZ with different heat input
2.4. Microstructure observation of simulated CGHAZ
Fig. 6a–d show the orientation image maps of simulated CGHAZ subjected to different heat input. When the heat input is 10 kJ/cm, the microstructures of simulated CGHAZ exhibits typical LB with various variants, shown in Fig. 6a. Several LB packets appear within a single prior austenite grain. The block is aggregations of the laths with the same crystallographic orientation, Fig. 6a–d shows that LB ferrite or GB ferrite with the same color are bainite blocks. Each packet is aggregations of one or several blocks with the similar crystallographic orientations [18,19]. Fig. 6 indicates that there are many variants with different orientations in LB microstructure when the heat input is low. As the heat input increases to 30 kJ/cm, the LB is divided into several smaller packets due to the presence of more GB blocks (Fig. 6b and f). When the heat input increases to 50 kJ/cm and 100 kJ/cm, the typical GB dominates the microstructure of simulated CGHAZ and the size of GB packet increases (Fig. 6c and d). There are only a few variants with different orientations in GB microstructure when the heat input is high. Fig. 6e–h present the image quality maps with grain boundary misorientation distribution of the simulated CGHAZ. The red lines stand for the low misorientation boundaries of 2–15°and the blue ones shows the high misorientation boundaries of 15–180°. It is seen that the percent of high misorientation boundaries at 10 kJ/cm is high due to the large content of LB blocks. Fig. 7 reveals that the LB block is composed of parallel laths with similar orientation having high misorientation angles in the high range of 55–60°, which can effectively arrest or deflect the microcracks propagation [10,20,21]. However, the percent of low misorientation boundaries decreases with the increase in heat input owning to the rise of GB content. Fig. 7 indicates the low misorientation intragranular GB that has the weak effect on suppressing cracks propagation. Fig. 6i–l shows the distribution of ferrite and retained austenite. The results show that the retained austenite is not detected at low heat input of 10 kJ/cm and 30 kJ/cm. The possible reason is that the volume fraction of retained austenite is too low and their size is too small at these conditions. When the heat input is increased to 50 kJ/cm, a very small amount of retained austenite (~ 0.2%) is detected at the grain boundaries and subgrain boundaries (shown in the white dotted frame in Fig. 6k). With the heat input increase to 100 kJ/cm, the partition fraction of retained austenite increases to about 1.1%, and it also distributes at the grain boundaries and subgrain boundaries (Fig. 6i).
The metallographic specimens of simulated CGHAZs were cut near the monitoring thermocouple, polished and etched with 4vol. % Nital. The microstructure was observed by optical microscope (OM) and JSM7800F scanning electron microscopy (SEM). The hardness of simulated CGHAZ was measured by HV-1000A Vickers hardness tester at 1000 gf load, and the average prior austenite grain size was measured by the line intercept method. In order to reveal the morphological characteristics of M-A constituent, the simulated CGHAZ were etched with Lepera etchant (consisting of 1% aqueous solution of sodium metabisulfite and 4% picric acid in alcohol) for 150 s. The morphology of M-A constituent was investigated by optical microscope and utilized ImagePro Plus to calculate the size and morphology characteristic parameters of M-A constituent. The crystallographic characteristics of simulated CGHAZ were investigated using electron backscattered diffraction (EBSD), and EBSD specimens were electrolytically polished in 5% perchloric acid alcohol solution for 20s. The EBSD maps were obtained at step size of 0.2 μm, and the scanning area was about 80 × 80 μm. 3. Results 3.1. Microstructure of the weld simulated CGHAZ The microstructures of simulated CGHAZ with different heat input are show in Fig. 4. The microstructures of simulated CGHAZ are composed of lath bainite (LB) and granular bainite (GB). The LB ferrite is mostly lath-shaped and M-A constituent (the black slender island indicated by the arrow in Fig. 4a) can be observed between the ferrite lathes. The GB ferrite is greatly massive, and a large number of M-A constituents (the small black island indicated by the arrow in Fig. 4b) are scattered in the massive GB ferrite. To quantitatively analyze the content of those two phases, the GB and LB regions in optical micrograph were labeled using Photoshop software, then the area fraction of LB and GB was calculated by Image-Pro Plus. For accurately statistics of the content of phases, ten different optical micrographs were calculated for each sample. When the heat input is 10 kJ/cm, the dominant microstructure is LB with volume fraction of ~91% (Fig. 4a). With the increase in heat input, the amount of LB decreases and that of GB increases. The microstructure of simulated CGHAZ at 30 kJ/cm consists of ~72% LB and ~28% GB (Fig. 4b), which makes the microstructure more intricate. When the heat input increases to 50 kJ/cm, the amount of LB decreases dramatically to ~22% (Fig. 4c). When the heat input is 100 kJ/cm, the microstructure of simulated CGHAZ is almost composed of GB (~93%) with only a very small amount of LB (Fig. 4d). Fig. 4 also shows that the M-A constituent mainly distributes between LB ferrite
3.3. Effect of heat input on fracture toughness of simulated CGHAZ Fig. 8 shows the variation of load-displacement curves with different heat input. It is evident that the load-displacement curves can be divided into three stages during fracture, which includes linear elastic 3
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Fig. 4. Optical micrograph of simulated CGHAZ with heat input of (a) 10 kJ/cm, (b) 30 kJ/cm, (c) 50 kJ/cm and(d) 100 kJ/cm.
deformation. When the load reaches the maximum, the brittle fracture occurs with dramatically decrease of load. Fig. 8 indicates that the displacement of elastic-plastic deformation stage and peak load for 30 kJ/cm are maximum. This implies that the crack tip region has a strong resistance to brittle fracture, which makes it difficult for microcracks to develop and has good fracture toughness. When the heat input is 50 kJ/cm and 100 kJ/cm the elastic-plastic deformation stage is very short, and the plastic deformation is not easy to occur in the crack tip region, resulting in the low resistance to brittle fracture and the low fracture toughness. Table 3 lists the average fracture toughness parameters of simulated CGHAZ with different heat input. Fip is the load corresponding to the initiation-point. Table 3 shows that Fip is relatively large when the heat input is 50 kJ/cm and 100 kJ/cm, which are 3.69 kN and 3.57 kN, respectively. The high Fip implies that plastic deformation is not easy to occur in the crack tip region and tends to brittle fracture. Fip was relatively low at 10 kJ/cm and 30 kJ/cm, with values of 2.39 kN and 3.30 kN, respectively, which implies low plastic deformation resistance of CGHAZ crack tip region. In addition, the displacement (Dip) at the initiation-point increases continuously from 0.060 to 0.097 with the increase of heat input, further indicating the increasing resistance of plastic deformation. Fmax is the maximum force before cracks unstable propagation that reflects the ability to resist brittle fracture, and Vp is the corresponding plastic opening displacement that represents the accumulation of plastic deformation at the crack tip. According to the BS7448 standard, Vp can be determined by graphical method as shown in Fig. 8. In this study, Fmax increases first and then decreases with the increase of heat input. When the heat input is 30 kJ/cm, the Fmax reaches a maximum of 7.51 kN. With the increase of heat input, Vp increases first and then decreases dramatically. When the heat input is 30 kJ/cm, the Vp reaches its maximum value of 0.513 mm, and the plastic deformation in the crack tip region is the largest before brittle fracture occurs at this moment. Fmax and Vp are the main parameters affecting the CTOD value, which is in direct proportion to the value. When the heat input from 10 kJ/cm to 100 kJ/cm, the CTOD value increases first and then decreases dramatically. The CTOD value
Fig. 5. Average grain size of prior austenite and Vickers hardness of the simulated CGHAZ.
deformation stage (LEDS), elastic-plastic deformation stage (EPDS) and brittle fracture stage (BFS). At the linear elastic deformation stage, the fatigue crack is still in a fully elastic state and does not expand. Once the load reaches the initiation-point, significant plastic deformation occurs in the crack tip region and the elastic-plastic deformation stage starts. When plastic deformation accumulates to the critical value, the micro-cracks begin to propagate in the tip region. The load at the initiation-point can reflect the resistance of CGHAZ to plastic deformation. The low load value implies that plastic deformation is easy to occur in the crack tip region. The plastic deformation at the crack tip can consume energy and limit local stress, which inhibits the brittle fracture [22,23]. In addition, the displacement at the initiation point is also an important parameter associated with plastic deformation initiation. Large displacement indicates the strong resistance to plastic
4
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Fig. 6. (a–d) EBSD inverse pole figure (IPF), (e–h) grain boundary maps and (i–l) phase maps of simulated CGHAZ with heat input of (a,e and i) 10 kJ/cm, (b,f and j) 30 kJ/cm, (c,g and k) 50 kJ/cm and (d,h and l) 100 kJ/cm.
that the width of crack propagation area is the largest about 115 μm at the heat input of 30 kJ/cm. The long elastic-plastic deformation stage leads to the wider crack propagation area, indicating the significant plastic deformation at the crack tip and good fracture toughness. The small dimples indicated by black arrows in Fig. 9b greatly contribute to dissipated energy during crack propagation and arrest the crack [5]. In addition, when the heat input is10 kJ/cm (Fig. 9a), it can be observed that the size of cleavage facet on the fracture surface is relatively small. And the size of cleavage facet is larger at 50 kJ/cm and 100 kJ/cm (Fig. 9c and d), which is prone to brittle fracture. 4. Discussion 4.1. Microstructural transformation of bainite with different heat input The microstructure of the investigated steel is composed of QPF, AF and B, shown in Fig. 2. However, Fig. 3 indicates that the microstructure transforms to bainite when it suffers the welding thermal cycle with heat input of 10–100 kJ/cm. The content of LB and GB in CGHAZ microstructures varies greatly with different heat input, implying the different bainite transformation behavior. In order to further investigate the phase transformation behavior of the CGHAZ, the dilatometric curves of the simulated CGHAZ specimens with different heat input were measured during the welding thermal cycling, presented in Fig. 10. Fig. 10 shows that the lowest starting temperature of phase transformation is 555 °C (10 kJ/cm). According to Andrews' empirical formula of Ms temperature (Ms = 539-423C-30.4Mn-17.7Ni-12.1Cr11.0Si-7.0Mo) for low alloy steel [24], the calculated Ms temperature of EH47 steel is 449 °C, which is much lower than the predicted Ms
Fig. 7. Point-to-point misorientation angle profile in LB and GB packet.
reaches the maximum at the heat input of 30 kJ/cm, but it decreases to only 0.01 mm at the heat input of 50 kJ/cm. The fracture surface with different heat input is presented in Fig. 9. The fracture surface is divided into three regions, including prefabricated fatigue crack area, crack propagation area and brittle fracture area. The crack propagation area is related to the elastic-plastic deformation stage, which reflects the fracture toughness. It is observed 5
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Fig. 8. Load-displacement curves with heat input of (a) 10 kJ/cm, (b) 30 kJ/cm, (c) 50 kJ/cm and (d) 100 kJ/cm.
transform into small island M-A constituent after rapid cooling to ambient temperature. Thus, many small island M-A constituents appear in the CGHAZ with low heat input of 10 kJ/cm, shown in Fig. 4a. However, Fig. 10a indicates that the Bs and Bf rise with increase in heat input. When the heat input is 100 kJ/cm, the Bs and Bf increase to 613 °C and 489 °C, respectively. Fig. 10b shows that the bainite transformation rate decreases with the increase of heat input. The bainite transformation rate of 100 kJ/cm is the lowest with the maximum rate of 12.6%/s. At this high heat input of 100 kJ/cm, the high bainite transformation temperature (489–613 °C) and the long phase transformation time (26 s) provide condition for the sufficient diffusion of carbon, which is beneficial to the transformation of GB. At the same time, the sufficient partitioning and long-range diffusion during bainite transformation are conducive to the formation of large sized carbonrich austenite. These carbon-rich austenites can be transformed into MA constituents at low temperatures, resulting in the appearance of many massive M-A constituents in CGHAZ, shown in Fig. 4d. In view of the above discussion, the microstructural evolution of CGHAZ for HSLA steels during welding thermal cycles can be summarized as Fig. 11. When the steel is rapidly heated to extremely high
temperature. This further confirms that the transformation is mainly bainite transformation in simulated CGHAZ. Previous studies [25−27] have shown that the bainite ferrite nucleates preferentially at the carbon-poor zones of the original austenite grain boundaries and accompanies the portion of carbon from bainite ferrite into the surrounding austenite matrix to form carbon enrichment of austenite. The carbon-rich austenite will transform into martensite or retain as austenitic state during the subsequent cooling stage. When the heat input is 10 kJ/cm, the bainite transformation start temperature (Bs) and transformation finish temperature (Bf) are low (555 °C and 359 °C, respectively) due to the high cooling rate, where the cooling rate from 800 °C to 500 °C is about 67 °C/s. Compared to other heat input conditions, the bainite transformation rate at heat input of 10 kJ/ cm is the highest with the maximum rate of 33.5%/s, presented in Fig. 10b. In this case, the low bainite transformation temperature (359–555 °C) and the short phase transformation time (6 s) lead to the insufficient diffusion of carbon, which will promote the transformation of LB. On the other hand, the small size carbon-rich austenite is formed during austenite-bainite transformation due to the insufficient partitioning and diffusion of carbon. This small carbon-rich austenite can
Table 3 Average fracture toughness parameters of simulated CGHAZ with different heat input. Heat input/kJ/cm
Fip/kN
10 30 50 100
2.39 3.30 3.69 3.57
± ± ± ±
Dip/mm 0.14 0.13 0.18 0.02
0.060 0.078 0.085 0.097
± ± ± ±
Fmax/kN 0.004 0.003 0.004 0.001
7.35 7.51 5.56 5.22
6
± ± ± ±
Vp/mm 0.50 0.22 0.12 0.48
0.340 0.513 0.038 0.104
± ± ± ±
CTOD value/mm 0.028 0.110 0.001 0.020
0.079 0.120 0.010 0.024
± ± ± ±
0.007 0.019 0.001 0.005
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Fig. 9. SEM micrographs of fracture surfaces of the simulated CGHAZ with heat input of (a) 10 kJ/cm, (b) 30 kJ/cm, (c) 50 kJ/cm and (d) 100 kJ/cm.
such as 100 kJ/cm, the nucleation of GB ferrite is predominant because of the high transformation temperature and long duration transformation time. Due to the increases of duration transformation time at high temperature, the bainitic ferrite gradually forms massive ferrite matrix, and conducive to form large sized carbon-rich austenites. During the subsequent cooling process, almost all CGHAZ microstructures were transformed into larger GB with many massive M-A constituents.
temperature (1350 °C) during welding, the microstructure completely transforms into austenite. Then the austenite transforms to bainite during continuous cooling stage of welding due to the relatively high cooling rate. When the heat input is low, such as 10 kJ/cm, a large number of LB ferrite can be formed preferentially due to the low transformation temperature and short duration transformation time. Meanwhile, carbon diffuses into the regions between the ferrite blocks of GB or LB during transformation. After further cooling, the carbonrich austenite regions are formed island or stringer M-A constituent. Ultimately, the microstructures of CGHAZ mainly consist of sufficient LB, accompanied by a large number of island M-A constituent. For the medium heat input, such as 30 kJ/cm, the nucleation of GB ferrite increases and interlaces with LB to form disordered microstructures owing to the higher transformation temperature and longer duration transformation time. The LB packets are segmented smaller due to the increase of GB packets. Therefore, many smaller GB and LB packets can be observed in CGHAZ microstructures. In the case of high heat input,
4.2. Effect of microstructure on the fracture toughness in simulated CGHAZ Considerable studies indicate that the fracture toughness depends on a combination of many factors, such as grain size, microstructural constitution, volume fraction of retained austenite and characteristics of M-A constituent etc. [12,28−34]. Especially, the existence of M-A constituent and retained austenite is the key factor affecting the fracture toughness. Previous studies have shown that retained austenite can effectively improve the low temperature toughness of steel [31−33]. It
Fig. 10. (a) Dilatometric curves and (b) transformation rate curves of bainite with different heat input. 7
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Fig. 13. Area fraction of M-A constituent with different heat input. Fig. 11. Bainite transformation of CGHAZ with different heat input.
increase of M-A area fraction led to the deterioration of toughness [35,36]. The M-A constituent is a hard and brittle phase, which is easy to cause stress concentration [10,37−40]. When the stress concentration exceeds the critical stress of M-A constituent, the microcrack began to forms on the M-A constituent and gradually extends to the matrix. In addition, the average area of M-A constituent with different heat input of 10 kJ/cm, 30 kJ/cm, 50 kJ/cm and 100 kJ/cm is 0.510 ± 0.857 um2, 0.989 ± 1.113 um2, 1.627 ± 3.516 um2 and 2.247 ± 5.227 um2, respectively. For those small M-A constituents, the damage to fracture toughness is not significant. However, the large M-A constituent distinctly worsens the fracture toughness. These large M-A constituents tend to cause stress concentration at the interface with the matrix and causes the poor interface strength. Therefore, it is easy to debond at the interface when subjected to external loads, which leads to the formation of microcracks. The connection of these adjacent
was reported that the toughness of steel containing 15% retained austenite was approximately ten times higher than that of steel without retained austenite because the retained austenite has high resistance to crack propagation [33]. In this study, the volume fraction of retained austenite in CGHAZ specimens is too small (0–1.1%) so that the impact on fracture toughness is not as good as expected. The Lepera etchant was used to reveal the morphology and distribution of M-A constituent, presented in Fig. 12. It is shown that the size of M-A constituent increases with increases in heat input, which is unfavorable to fracture toughness. Furthermore, statistical analysis shows that the area fraction of M-A constituent increases from 15.2% to 28.0% with the increase of heat input, shown in Fig. 13. This can lead to deterioration of fracture toughness. Previous studies also found the similar results that the
Fig. 12. Optical micrograph of M-A constituent (etched by Lepera) with heat input of (a) 10 kJ/cm, (b) 30 kJ/cm, (c) 50 kJ/cm and (d) 100 kJ/cm. 8
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Fig. 14. Effect of M-A constituent morphology on crack propagation: (a) and (b)Island M-A constituents arrest the cracks at 10 kJ/cm and 30 kJ/cm, respectively; (c) Deformation of fine-stringer M-A constituents at 30 kJ/cm; (d) Coarse-stringer M-A constituents cut off by a crack at 50 kJ/cm; (e) Undeformed coarse-stringer M-A constituents near the crack tip at 50 kJ/cm and (f) Interface decohesion of coarse-stringer M-A constituents at 100 kJ/cm.
Fig. 15. Distribution characteristics of M-A constituent with different heat input.
input is 10 kJ/cm and 30 kJ/cm, the small size island M-A constituents can inhibit the propagation of secondary cracks, shown in Fig. 14a and b. According to the classical Griffith theory [41], the critical stress of the island M-A constituent cracking is large due to the small size, which
debonding zones will cause crack propagation and promote brittle fracture. In fact, different morphologies of M-A constituent may have discrepant effects on fracture toughness. In this investigate, when the heat 9
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the increase of GB aggravates the deterioration of fracture toughness at 50 kJ/cm and 100 kJ/cm. However, it is noted that the fracture toughness with heat input of 30 kJ/cm is higher than that of 10 kJ/cm, despite the higher content of GB, higher volume fraction and larger size of M-A constituent. The existence of retained austenite can effectively improve the fracture toughness, but no retained austenite can be detected at 30 kJ/cm according to EBSD phase map (Fig. 6j). Therefore, the abnormal phenomenon is considered to be attributed to the influence of the effect of packet size on fracture toughness. When the heat input is 30 kJ/cm, a large number of GB is generated in CGHAZ microstructures and interlaced with LB. The packets of LB are segmented smaller by GB. The small packet size can shorten the length of microcracks and improve toughness [42]. Therefore, when the heat input is 30 kJ/cm, the fracture toughness of CGHAZ is better than that of 10 kJ/ cm.
makes it difficult to form microcracks. Therefore, these island M-A constituents can arrest the secondary cracks and dissipate more energy of crack propagation, which is beneficial to fracture toughness. Fig. 14c shows that the fine-stringer M-A constituents has obvious plastic deformation in the tip region of the main crack. The local stress concentration is decreased and the energy of crack propagation can be consumed by plastic deformation of fine-stringer M-A constituent during the fracture. Therefore, the fine-stringer M-A constituent does not excessively deteriorate the fracture toughness. However, the coarsestringer M-A constituents are easily cut off when subjected to transverse shear stress, which is helpful for microcracks propagation, shown in Fig. 12d. It is also observed that there is no obvious deformation of the coarse-stringer M-A constituents at the crack tip region (Fig. 12e), which implies that these coarse-stringer M-A constituents are unfavorable to fracture toughness. In addition, Fig. 12f shows that the coarsestringer M-A constituents deboned from the matrix and linked other debonding regions to accelerate crack propagation. Therefore, the fracture toughness is significantly decreased when there are lots of coarse-stringer M-A constituents. In addition, previous studies [8,10,35,37] have confirmed that the massive M-A constituents are prone to initiate microcracks due to the low critical stress of microcracks formation. At the same time, the massive M-A constituents can also debond from the matrix, which leads to serious deterioration of fracture toughness. Based on the above discussion, in order to further confirm the effect of different morphologies of M-A constituent on fracture toughness, the M-A constituents are classified into four types: island, fine-stringer, coarse-stringer and massive, respectively. The single M-A constituent is characterized by its area and aspect ratio (maximal Feret diameter Lmax/minimal Feret diameter Lmin). Previous studies have shown that the average area of many island M-A constituents is considered < 5 um2, while the aspect ratio of stringer M-A constituent is usually > 4 [8,10,38−40]. Fig. 15 shows the detailed morphological distribution of M-A constituent in simulated CGHAZ with different heat input. When the heat input is 10 kJ/cm and 30 kJ/cm, the island M-A constituent (area < 5 um2 and aspect ratio < 4) can arrest the secondary cracks and the area fractions is 86% and 85%, respectively, which is beneficial to fracture toughness. When the heat input rises to 50 kJ/cm and 100 kJ/cm, the island M-A constituent decreases significantly, and the area fraction is only 51% and 35%, respectively, which is detrimental to fracture toughness. In addition, with the increase of heat input, the area fraction of fine-stringer M-A constituent (area < 5 um2 and aspect ratio > 4) decreases from 7% to 1%, whereas that of coarse-stringer MA constituent (area > 5 um2 and aspect ratio > 4) increases from 1% to 7%. The fine-stringer M-A constituent has obvious plastic deformation without excessive damage to fracture toughness, while the coarsestringer M-A constituent exhibits no obvious plastic deformation and is easy to crack and debond from matrix, which seriously damages the fracture toughness. The massive M-A constituent (area > 5 um2 and aspect ratio < 4) is also prone to cracking and debonding from matrix. The area fraction of massive M-A constituent increases dramatically from only 6% to 57% with the increase of heat input, which is one of the reasons for the greatly deterioration of fracture toughness at high heat input (50–100 kJ/cm). Furthermore, the content of LB and GB also significantly influences the fracture toughness of CGHAZ. With the increase of heat input from 10 kJ/cm to 100 kJ/cm, the proportion of LB in CGHAZ decreases dramatically from 91% to 7%. Previous EBSD analyses have concluded that there are a large number of high misorientation boundaries in the LB, as shown in Figs. 6e and 7. High misorientation boundaries can arrest the microcracks during crack propagation, which effectively improves the fracture toughness [10,20,21]. On the contrary, GB increases from 9% to 93% with the increase of heat input from 10 kJ/cm to 100 kJ/cm. The GB is ineffective in preventing the propagation of microcracks due to the existence of many low misorientation boundaries (Figs. 6h and 7). Therefore, the reduction of LB accompanied with
5. Conclusions (1) The microstructure of simulated CGHAZ changes from lath bainite to granular bainite with the increase of heat input, accompanied by the increase of prior austenite grain. The lath bainite has faster transformation rate and more abundant high misorientation boundaries than granular bainite. Excessive coarsening of granular bainite occurs when heat input increases to higher than 50 kJ/cm. (2) The welding heat input significantly influences the content and morphology of M-A constitute. With the increase of heat input, the area fraction of M-A constituent increases from 15.2% to 28.0%, the area proportion of island and fine stringer M-A constituent decrease from 86% and 7% to 35% and 1% respectively, and that of coarsestringer and massive M-A constituent increase from 1% and 6% to 7% and 57% respectively. (3) The fracture toughness of simulated CGHAZ is proved to be sensitive to the morphology of M-A constitute. The island M-A constituent can arrest microcracks, which is beneficial to fracture toughness. The fine stringer M-A constituent does not greatly damage to fracture toughness due to the deformation. However, the coarse stringer and massive M-A constituent can crack and debond from the matrix, which seriously deteriorates the fracture toughness. (4) When the heat input increases from 10 kJ/cm to 30 kJ/cm, the CTOD value at −20 °C increases from 0.079 mm to 0.12 mm, and dramatically decreases as the heat input increases to higher than 50 kJ/cm. The packet of lath bainite is segmented smaller by granular bainite, resulting in higher CTOD values at 30 kJ/cm. and the significant decrease of CTOD value at the heat input of higher than 50 kJ/cm is due to the large content of GB, coarsening of bainite packet, and increase in content of coarse-stringer and massive M-A constituent. Acknowledgements The author is sincerely grateful to the financial support from the National Natural Science Foundation of China (grant No. 51804217 and 51774213), the State Key Laboratory of Metal Material for Marine Equipment and Application (grant No. SKLMEA-K201602) and the Regional Demonstration Project of Marine Economic Innovation and Development (grant No. BHSF2017-10). References [1] Y. Sumi, H. Yajima, M. Toyosada, T. Yoshikawa, S. Aihara, K. Gotoh, Y. Ogawa, T. Matsumoto, K. Hirota, H. Hirasawa, Fracture control of extremely thick welded steel plates applied to the deck structure of large container ships, J. Mar. Sci. Technol. 18 (2013) 497–514. [2] T. Handa, S. Suzuki, N. Kiji, M. Toyoda, T. Miyata, Effect of unwelded length on behaviour of brittle crack arrest in T-joint structure, Weld. Int. 23 (2009) 640–647. [3] J.Y. Wu, B. Wang, B.X. Wang, R.D.K. Misra, Z.D. Wang, Toughness and ductility
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Review
Effects of heat input on microstructure and fracture toughness of simulated coarse-grained heat affected zone for HSLA steels ⁎
Xiaocong Yanga,b, Xinjie Dia,b, , Xiuguo Liua,b, Dongpo Wanga,b,c, Chengning Lia,b,
T
⁎
a
School of Materials Science and Engineering, Tianjin University, Tianjin 300350, China Tianjin Key Laboratory of Advanced Joining Technology, Tianjin 300350, China c State Key Laboratory of Metal Material for Marine Equipment and Application, Anshan 114009, China b
A R T I C LE I N FO
A B S T R A C T
Keywords: Coarse-grained heat affected zone HSLA steel Heat input Crack tip opening displacement M-A constituent Fracture toughness
In order to clarify the relationships among welding heat input, microstructural evolution and fracture toughness of coarse-grained heat affected zone (CGHAZ) for HSLA steels, the microstructure of simulated CGHAZ with heat input of 10–100 kJ/cm was quantified and the crack tip opening displacement (CTOD) at −20 °C was evaluated With the increase in heat input, the bainite transformation temperature and duration transformation time increase, which lead to the higher contents of granular bainite and M-A constituents. The fracture toughness of simulated CGHAZ is highly related with the morphologies of M-A constitute. During CTOD test, the island M-A constituents can arrest microcracks and increase the CTOD value, and the fine-stringer ones can deform without excessive damage to fracture toughness. However, the coarse-stringer and massive M-A constituent would crack and debond from the matrix, which seriously decreases the fracture toughness. When the heat input increases from 10 kJ/cm to 100 kJ/cm, the proportion of island M-A constituent decreases from 86% to 35%, while that of coarse-stringer and massive M-A constituent increases from 1% and 6% to 7% and 57%, respectively, which result in the significant deterioration of fracture toughness. The fracture toughness is optimal at heat input of 30 kJ/cm, because the packets of lath bainite are segmented smaller by granular bainite and effectively improved the cracking resistance.
1. Introduction In order to improve the efficiency of transportation and reduce costs, the container ships are required to have larger capacity. The application of high strength low alloy (HSLA) steels, such as EH47, can significant improve the carrying capacity and safety performance of container ships, which have been used in deck and hatch coaming structures [1,2]. The good combination of high strength and toughness for the HSLA steels is well controlled through thermal mechanical controlled processing (TMCP) [3]. However, this excellent balance of strength and toughness will be broken after the high temperature thermal cycling process in welding, such as the reduction of toughness in heat affected zone (HAZ). In generally, coarse-grained HAZ (CGHAZ) is always an attention-attracting local brittle region due to the coarsening prior austenite grain and the formation of M-A constituents, which is highly related to the welding heat input [4–10]. Therefore, the CGHAZ is usually the risk region of fracture failure, especially when the CGHAZ has microcracks produced by cyclic loads of sea wave or welding defect [11], which limits the use of HSLA steels at low
⁎
temperatures. It is known that the hard and brittle M-A constituent can seriously deteriorate fracture toughness. Mohseni et al. concluded that massive M-A constituent decreased the low temperature toughness of CGHAZ seriously, and the fracture initiation occurred preferentially at M-A constituent by a debonding mechanism [12]. Luo et al. pointed out that the slender M-A constituent was more harmful to toughness than massive M-A, and indicated that the massive M-A particles were usually debonded from the matrix while the slender M-A particles were more likely to crack [8]. Chen et al. reported that the micromechanism of cleavage cracking in HAZ was controlled by the size of M-A constituent, and high stress concentrates on the boundary of MA/matrix made it crack or debond due to the heavy deformation of ferrite matrix [13]. However, studies by Lan et al. shows that the small M-A constituent was helpful in improving toughness [10]. Therefore, the effect of M-A constituent on the micromechanism of fracture and fracture toughness is confused, which would be related to the size, content and distribution morphology. In this study, the simulated CGHAZs with different morphology of
Corresponding authors at: School of Materials Science and Engineering, Tianjin University, Tianjin 300350, China. E-mail addresses: [email protected] (X. Di), [email protected] (C. Li).
https://doi.org/10.1016/j.matchar.2019.109818 Received 5 May 2019; Received in revised form 11 July 2019; Accepted 12 July 2019 Available online 12 July 2019 1044-5803/ © 2019 Elsevier Inc. All rights reserved.
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M-A were obtained by welding thermal simulated with different heat input, and the microstructure and crystallographic characteristics of simulated CGHAZs were investigated. At present, the toughness of CGHAZ is mainly evaluated by the Charpy V-notch impact test [4,5,10,14−16]，but it cannot reflect the fracture mechanism of CGHAZ with fatigue cracks or microcracks. Compared with impact toughness, the fracture toughness can effectively characterize the actual service performance of ship hull with microcracks [17]. Therefore, in this paper, the fracture toughness of CGHAZ was measured by three point bend test of crack tip opening displacement (CTOD), which can reflect the plastic deformation performance near the crack tip after cyclic loading. Moreover, the relationship between fracture toughness and microstructure with different heat input was discussed, and the effect of M-A constituent morphologies on fracture micromechanism was discussed in detail.
Table 1 Chemical composition of experimental steel (wt%). C
Si
Mn
P
S
Ni
0.067
0.193
1.34
0.011
0.043
0.831
Cu + Cr + Mo
V + Ti + Nb
≤0.70
≤0.06
Table 2 Mechanical properties of experimental steel. YS Rs/MPa 472
TS Rm/MPa
Yield ratio Rs/Rm
Elongation/%
Hardness/HV
600
0.79
32
200
CTOD value/ mm 0.533
2. Experimental procedure 2.1. Microstructure and properties of the experimental steel The experimental steel is EH47 steel, which is a typical crack arrestability HSLA steels manufactured by TMCP technology. The microstructure of experimental steel is composed of polygonal ferrite (PF), quasi-polygonal ferrite (QPF), acicular ferrite (AF) and bainite (B), presented in Fig. 1. The chemistry and mechanical properties of the experimental steel are listed in Tables 1 and 2, respectively. The yield strength at room temperature is about 472 MPa, and the average CTOD value is 0.533 mm at −20 °C, which indicates that the steel has excellent balance of strength and toughness. In this study, the CTOD test temperature was selected to be −20 °C according to the standard of DNVGL-OS-C401 and DNVGL-OS-C101, which can more effectively reflect the actual service condition and assess the safety of the steel in service.
Fig. 2. Simulated thermal cycle curves with different heat input.
2.2. Welding thermal simulation experiment
50 kJ/cm and 100 kJ/cm, respectively.
Gleeble 3800 simulation machine was used to simulate the CGHAZ in the actual welding. The simulated specimens were cut from a quarter of the 80 mm thick steel plate along the transverse direction and machined into the dimensions 11 mm × 11 mm × 80 mm. The simulated specimens were subjected to different welding thermal cycles according to different actual welding processes. The heat input was 10 kJ/cm, 30 kJ/cm, 50 kJ/cm and 100 kJ/cm, respectively. Simulated thermal cycles are shown in Fig. 2. The three dimensional Rykalin thermodynamic model was used to determine the thermal cycle curve of ultrathick steel plates. The simulated specimens were heated to peak temperature of 1350 °C at a heating rate of 200 °C/s, holding for 1 s at the peak temperature. The cooling time of t8/5 was 4.5 s, 13.4 s, 22.4 s and 44.7 s corresponding to welding heat input of 10 kJ/cm, 30 kJ/cm,
2.3. Fracture toughness test After thermal simulations, the specimens were machined into CTOD specimens with dimension of 10 × 10 × 80 mm according to the standard method of BS7448, and the mechanical notch was located in the thickness direction with a depth of 2.5 mm, shown in Fig. 3. Then the fatigue cracks were prefabricated at room temperature (the maximum value of alternating load is 3.5 kN) until the cracks extend to 1/2 thickness of the specimens. Thereafter, the three point bend test of CTOD was carried out at −20 °C to measure the fracture toughness of simulated CGHAZ. After CTOD test, the specimen was crushed at −80 °C and the fracture morphology was observed. Finally, the validity
Fig. 1. Microstructure of experimental steel: (a) Optical micrograph and (b) SEM micrograph. 2
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laths or within GB ferrite, and displays many different morphologies, including island, stringer and massive. With the increase of heat input, the M-A constituent coarsen obviously and the proportion of island M-A constituent decreased gradually, while the proportion of massive M-A constituent increased continuously. Fig. 5 shows the change of average grain size of prior austenite and Vickers hardness of simulated CGHAZ. It indicates that the average grain size of prior austenite increases with the increase in heat input. The coarsening of prior austenite grain at high heat input is attributed to the increase of duration time at high temperature that the grain can grow easily. In addition, the Vickers hardness decreases with the increase of heat input, which mainly depends on the microstructure and grain size. When the heat input increases, the increase of GB that composed of massive ferrite is the key factor leading to the decrease of hardness. At the same time, the increase of grain size reduces the strengthening effect of effective grain boundary, which results in the decrease of Vickers hardness [14].
Fig. 3. Dimensions of CTOD specimens.
of the test is judged by checking the profiles of all CTOD specimens and ensuring that the crack tip is located in the designated area.
3.2. The crystallographic characteristics of simulated CGHAZ with different heat input
2.4. Microstructure observation of simulated CGHAZ
Fig. 6a–d show the orientation image maps of simulated CGHAZ subjected to different heat input. When the heat input is 10 kJ/cm, the microstructures of simulated CGHAZ exhibits typical LB with various variants, shown in Fig. 6a. Several LB packets appear within a single prior austenite grain. The block is aggregations of the laths with the same crystallographic orientation, Fig. 6a–d shows that LB ferrite or GB ferrite with the same color are bainite blocks. Each packet is aggregations of one or several blocks with the similar crystallographic orientations [18,19]. Fig. 6 indicates that there are many variants with different orientations in LB microstructure when the heat input is low. As the heat input increases to 30 kJ/cm, the LB is divided into several smaller packets due to the presence of more GB blocks (Fig. 6b and f). When the heat input increases to 50 kJ/cm and 100 kJ/cm, the typical GB dominates the microstructure of simulated CGHAZ and the size of GB packet increases (Fig. 6c and d). There are only a few variants with different orientations in GB microstructure when the heat input is high. Fig. 6e–h present the image quality maps with grain boundary misorientation distribution of the simulated CGHAZ. The red lines stand for the low misorientation boundaries of 2–15°and the blue ones shows the high misorientation boundaries of 15–180°. It is seen that the percent of high misorientation boundaries at 10 kJ/cm is high due to the large content of LB blocks. Fig. 7 reveals that the LB block is composed of parallel laths with similar orientation having high misorientation angles in the high range of 55–60°, which can effectively arrest or deflect the microcracks propagation [10,20,21]. However, the percent of low misorientation boundaries decreases with the increase in heat input owning to the rise of GB content. Fig. 7 indicates the low misorientation intragranular GB that has the weak effect on suppressing cracks propagation. Fig. 6i–l shows the distribution of ferrite and retained austenite. The results show that the retained austenite is not detected at low heat input of 10 kJ/cm and 30 kJ/cm. The possible reason is that the volume fraction of retained austenite is too low and their size is too small at these conditions. When the heat input is increased to 50 kJ/cm, a very small amount of retained austenite (~ 0.2%) is detected at the grain boundaries and subgrain boundaries (shown in the white dotted frame in Fig. 6k). With the heat input increase to 100 kJ/cm, the partition fraction of retained austenite increases to about 1.1%, and it also distributes at the grain boundaries and subgrain boundaries (Fig. 6i).
The metallographic specimens of simulated CGHAZs were cut near the monitoring thermocouple, polished and etched with 4vol. % Nital. The microstructure was observed by optical microscope (OM) and JSM7800F scanning electron microscopy (SEM). The hardness of simulated CGHAZ was measured by HV-1000A Vickers hardness tester at 1000 gf load, and the average prior austenite grain size was measured by the line intercept method. In order to reveal the morphological characteristics of M-A constituent, the simulated CGHAZ were etched with Lepera etchant (consisting of 1% aqueous solution of sodium metabisulfite and 4% picric acid in alcohol) for 150 s. The morphology of M-A constituent was investigated by optical microscope and utilized ImagePro Plus to calculate the size and morphology characteristic parameters of M-A constituent. The crystallographic characteristics of simulated CGHAZ were investigated using electron backscattered diffraction (EBSD), and EBSD specimens were electrolytically polished in 5% perchloric acid alcohol solution for 20s. The EBSD maps were obtained at step size of 0.2 μm, and the scanning area was about 80 × 80 μm. 3. Results 3.1. Microstructure of the weld simulated CGHAZ The microstructures of simulated CGHAZ with different heat input are show in Fig. 4. The microstructures of simulated CGHAZ are composed of lath bainite (LB) and granular bainite (GB). The LB ferrite is mostly lath-shaped and M-A constituent (the black slender island indicated by the arrow in Fig. 4a) can be observed between the ferrite lathes. The GB ferrite is greatly massive, and a large number of M-A constituents (the small black island indicated by the arrow in Fig. 4b) are scattered in the massive GB ferrite. To quantitatively analyze the content of those two phases, the GB and LB regions in optical micrograph were labeled using Photoshop software, then the area fraction of LB and GB was calculated by Image-Pro Plus. For accurately statistics of the content of phases, ten different optical micrographs were calculated for each sample. When the heat input is 10 kJ/cm, the dominant microstructure is LB with volume fraction of ~91% (Fig. 4a). With the increase in heat input, the amount of LB decreases and that of GB increases. The microstructure of simulated CGHAZ at 30 kJ/cm consists of ~72% LB and ~28% GB (Fig. 4b), which makes the microstructure more intricate. When the heat input increases to 50 kJ/cm, the amount of LB decreases dramatically to ~22% (Fig. 4c). When the heat input is 100 kJ/cm, the microstructure of simulated CGHAZ is almost composed of GB (~93%) with only a very small amount of LB (Fig. 4d). Fig. 4 also shows that the M-A constituent mainly distributes between LB ferrite
3.3. Effect of heat input on fracture toughness of simulated CGHAZ Fig. 8 shows the variation of load-displacement curves with different heat input. It is evident that the load-displacement curves can be divided into three stages during fracture, which includes linear elastic 3
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Fig. 4. Optical micrograph of simulated CGHAZ with heat input of (a) 10 kJ/cm, (b) 30 kJ/cm, (c) 50 kJ/cm and(d) 100 kJ/cm.
deformation. When the load reaches the maximum, the brittle fracture occurs with dramatically decrease of load. Fig. 8 indicates that the displacement of elastic-plastic deformation stage and peak load for 30 kJ/cm are maximum. This implies that the crack tip region has a strong resistance to brittle fracture, which makes it difficult for microcracks to develop and has good fracture toughness. When the heat input is 50 kJ/cm and 100 kJ/cm the elastic-plastic deformation stage is very short, and the plastic deformation is not easy to occur in the crack tip region, resulting in the low resistance to brittle fracture and the low fracture toughness. Table 3 lists the average fracture toughness parameters of simulated CGHAZ with different heat input. Fip is the load corresponding to the initiation-point. Table 3 shows that Fip is relatively large when the heat input is 50 kJ/cm and 100 kJ/cm, which are 3.69 kN and 3.57 kN, respectively. The high Fip implies that plastic deformation is not easy to occur in the crack tip region and tends to brittle fracture. Fip was relatively low at 10 kJ/cm and 30 kJ/cm, with values of 2.39 kN and 3.30 kN, respectively, which implies low plastic deformation resistance of CGHAZ crack tip region. In addition, the displacement (Dip) at the initiation-point increases continuously from 0.060 to 0.097 with the increase of heat input, further indicating the increasing resistance of plastic deformation. Fmax is the maximum force before cracks unstable propagation that reflects the ability to resist brittle fracture, and Vp is the corresponding plastic opening displacement that represents the accumulation of plastic deformation at the crack tip. According to the BS7448 standard, Vp can be determined by graphical method as shown in Fig. 8. In this study, Fmax increases first and then decreases with the increase of heat input. When the heat input is 30 kJ/cm, the Fmax reaches a maximum of 7.51 kN. With the increase of heat input, Vp increases first and then decreases dramatically. When the heat input is 30 kJ/cm, the Vp reaches its maximum value of 0.513 mm, and the plastic deformation in the crack tip region is the largest before brittle fracture occurs at this moment. Fmax and Vp are the main parameters affecting the CTOD value, which is in direct proportion to the value. When the heat input from 10 kJ/cm to 100 kJ/cm, the CTOD value increases first and then decreases dramatically. The CTOD value
Fig. 5. Average grain size of prior austenite and Vickers hardness of the simulated CGHAZ.
deformation stage (LEDS), elastic-plastic deformation stage (EPDS) and brittle fracture stage (BFS). At the linear elastic deformation stage, the fatigue crack is still in a fully elastic state and does not expand. Once the load reaches the initiation-point, significant plastic deformation occurs in the crack tip region and the elastic-plastic deformation stage starts. When plastic deformation accumulates to the critical value, the micro-cracks begin to propagate in the tip region. The load at the initiation-point can reflect the resistance of CGHAZ to plastic deformation. The low load value implies that plastic deformation is easy to occur in the crack tip region. The plastic deformation at the crack tip can consume energy and limit local stress, which inhibits the brittle fracture [22,23]. In addition, the displacement at the initiation point is also an important parameter associated with plastic deformation initiation. Large displacement indicates the strong resistance to plastic
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Fig. 6. (a–d) EBSD inverse pole figure (IPF), (e–h) grain boundary maps and (i–l) phase maps of simulated CGHAZ with heat input of (a,e and i) 10 kJ/cm, (b,f and j) 30 kJ/cm, (c,g and k) 50 kJ/cm and (d,h and l) 100 kJ/cm.
that the width of crack propagation area is the largest about 115 μm at the heat input of 30 kJ/cm. The long elastic-plastic deformation stage leads to the wider crack propagation area, indicating the significant plastic deformation at the crack tip and good fracture toughness. The small dimples indicated by black arrows in Fig. 9b greatly contribute to dissipated energy during crack propagation and arrest the crack [5]. In addition, when the heat input is10 kJ/cm (Fig. 9a), it can be observed that the size of cleavage facet on the fracture surface is relatively small. And the size of cleavage facet is larger at 50 kJ/cm and 100 kJ/cm (Fig. 9c and d), which is prone to brittle fracture. 4. Discussion 4.1. Microstructural transformation of bainite with different heat input The microstructure of the investigated steel is composed of QPF, AF and B, shown in Fig. 2. However, Fig. 3 indicates that the microstructure transforms to bainite when it suffers the welding thermal cycle with heat input of 10–100 kJ/cm. The content of LB and GB in CGHAZ microstructures varies greatly with different heat input, implying the different bainite transformation behavior. In order to further investigate the phase transformation behavior of the CGHAZ, the dilatometric curves of the simulated CGHAZ specimens with different heat input were measured during the welding thermal cycling, presented in Fig. 10. Fig. 10 shows that the lowest starting temperature of phase transformation is 555 °C (10 kJ/cm). According to Andrews' empirical formula of Ms temperature (Ms = 539-423C-30.4Mn-17.7Ni-12.1Cr11.0Si-7.0Mo) for low alloy steel [24], the calculated Ms temperature of EH47 steel is 449 °C, which is much lower than the predicted Ms
Fig. 7. Point-to-point misorientation angle profile in LB and GB packet.
reaches the maximum at the heat input of 30 kJ/cm, but it decreases to only 0.01 mm at the heat input of 50 kJ/cm. The fracture surface with different heat input is presented in Fig. 9. The fracture surface is divided into three regions, including prefabricated fatigue crack area, crack propagation area and brittle fracture area. The crack propagation area is related to the elastic-plastic deformation stage, which reflects the fracture toughness. It is observed 5
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Fig. 8. Load-displacement curves with heat input of (a) 10 kJ/cm, (b) 30 kJ/cm, (c) 50 kJ/cm and (d) 100 kJ/cm.
transform into small island M-A constituent after rapid cooling to ambient temperature. Thus, many small island M-A constituents appear in the CGHAZ with low heat input of 10 kJ/cm, shown in Fig. 4a. However, Fig. 10a indicates that the Bs and Bf rise with increase in heat input. When the heat input is 100 kJ/cm, the Bs and Bf increase to 613 °C and 489 °C, respectively. Fig. 10b shows that the bainite transformation rate decreases with the increase of heat input. The bainite transformation rate of 100 kJ/cm is the lowest with the maximum rate of 12.6%/s. At this high heat input of 100 kJ/cm, the high bainite transformation temperature (489–613 °C) and the long phase transformation time (26 s) provide condition for the sufficient diffusion of carbon, which is beneficial to the transformation of GB. At the same time, the sufficient partitioning and long-range diffusion during bainite transformation are conducive to the formation of large sized carbonrich austenite. These carbon-rich austenites can be transformed into MA constituents at low temperatures, resulting in the appearance of many massive M-A constituents in CGHAZ, shown in Fig. 4d. In view of the above discussion, the microstructural evolution of CGHAZ for HSLA steels during welding thermal cycles can be summarized as Fig. 11. When the steel is rapidly heated to extremely high
temperature. This further confirms that the transformation is mainly bainite transformation in simulated CGHAZ. Previous studies [25−27] have shown that the bainite ferrite nucleates preferentially at the carbon-poor zones of the original austenite grain boundaries and accompanies the portion of carbon from bainite ferrite into the surrounding austenite matrix to form carbon enrichment of austenite. The carbon-rich austenite will transform into martensite or retain as austenitic state during the subsequent cooling stage. When the heat input is 10 kJ/cm, the bainite transformation start temperature (Bs) and transformation finish temperature (Bf) are low (555 °C and 359 °C, respectively) due to the high cooling rate, where the cooling rate from 800 °C to 500 °C is about 67 °C/s. Compared to other heat input conditions, the bainite transformation rate at heat input of 10 kJ/ cm is the highest with the maximum rate of 33.5%/s, presented in Fig. 10b. In this case, the low bainite transformation temperature (359–555 °C) and the short phase transformation time (6 s) lead to the insufficient diffusion of carbon, which will promote the transformation of LB. On the other hand, the small size carbon-rich austenite is formed during austenite-bainite transformation due to the insufficient partitioning and diffusion of carbon. This small carbon-rich austenite can
Table 3 Average fracture toughness parameters of simulated CGHAZ with different heat input. Heat input/kJ/cm
Fip/kN
10 30 50 100
2.39 3.30 3.69 3.57
± ± ± ±
Dip/mm 0.14 0.13 0.18 0.02
0.060 0.078 0.085 0.097
± ± ± ±
Fmax/kN 0.004 0.003 0.004 0.001
7.35 7.51 5.56 5.22
6
± ± ± ±
Vp/mm 0.50 0.22 0.12 0.48
0.340 0.513 0.038 0.104
± ± ± ±
CTOD value/mm 0.028 0.110 0.001 0.020
0.079 0.120 0.010 0.024
± ± ± ±
0.007 0.019 0.001 0.005
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Fig. 9. SEM micrographs of fracture surfaces of the simulated CGHAZ with heat input of (a) 10 kJ/cm, (b) 30 kJ/cm, (c) 50 kJ/cm and (d) 100 kJ/cm.
such as 100 kJ/cm, the nucleation of GB ferrite is predominant because of the high transformation temperature and long duration transformation time. Due to the increases of duration transformation time at high temperature, the bainitic ferrite gradually forms massive ferrite matrix, and conducive to form large sized carbon-rich austenites. During the subsequent cooling process, almost all CGHAZ microstructures were transformed into larger GB with many massive M-A constituents.
temperature (1350 °C) during welding, the microstructure completely transforms into austenite. Then the austenite transforms to bainite during continuous cooling stage of welding due to the relatively high cooling rate. When the heat input is low, such as 10 kJ/cm, a large number of LB ferrite can be formed preferentially due to the low transformation temperature and short duration transformation time. Meanwhile, carbon diffuses into the regions between the ferrite blocks of GB or LB during transformation. After further cooling, the carbonrich austenite regions are formed island or stringer M-A constituent. Ultimately, the microstructures of CGHAZ mainly consist of sufficient LB, accompanied by a large number of island M-A constituent. For the medium heat input, such as 30 kJ/cm, the nucleation of GB ferrite increases and interlaces with LB to form disordered microstructures owing to the higher transformation temperature and longer duration transformation time. The LB packets are segmented smaller due to the increase of GB packets. Therefore, many smaller GB and LB packets can be observed in CGHAZ microstructures. In the case of high heat input,
4.2. Effect of microstructure on the fracture toughness in simulated CGHAZ Considerable studies indicate that the fracture toughness depends on a combination of many factors, such as grain size, microstructural constitution, volume fraction of retained austenite and characteristics of M-A constituent etc. [12,28−34]. Especially, the existence of M-A constituent and retained austenite is the key factor affecting the fracture toughness. Previous studies have shown that retained austenite can effectively improve the low temperature toughness of steel [31−33]. It
Fig. 10. (a) Dilatometric curves and (b) transformation rate curves of bainite with different heat input. 7
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Fig. 13. Area fraction of M-A constituent with different heat input. Fig. 11. Bainite transformation of CGHAZ with different heat input.
increase of M-A area fraction led to the deterioration of toughness [35,36]. The M-A constituent is a hard and brittle phase, which is easy to cause stress concentration [10,37−40]. When the stress concentration exceeds the critical stress of M-A constituent, the microcrack began to forms on the M-A constituent and gradually extends to the matrix. In addition, the average area of M-A constituent with different heat input of 10 kJ/cm, 30 kJ/cm, 50 kJ/cm and 100 kJ/cm is 0.510 ± 0.857 um2, 0.989 ± 1.113 um2, 1.627 ± 3.516 um2 and 2.247 ± 5.227 um2, respectively. For those small M-A constituents, the damage to fracture toughness is not significant. However, the large M-A constituent distinctly worsens the fracture toughness. These large M-A constituents tend to cause stress concentration at the interface with the matrix and causes the poor interface strength. Therefore, it is easy to debond at the interface when subjected to external loads, which leads to the formation of microcracks. The connection of these adjacent
was reported that the toughness of steel containing 15% retained austenite was approximately ten times higher than that of steel without retained austenite because the retained austenite has high resistance to crack propagation [33]. In this study, the volume fraction of retained austenite in CGHAZ specimens is too small (0–1.1%) so that the impact on fracture toughness is not as good as expected. The Lepera etchant was used to reveal the morphology and distribution of M-A constituent, presented in Fig. 12. It is shown that the size of M-A constituent increases with increases in heat input, which is unfavorable to fracture toughness. Furthermore, statistical analysis shows that the area fraction of M-A constituent increases from 15.2% to 28.0% with the increase of heat input, shown in Fig. 13. This can lead to deterioration of fracture toughness. Previous studies also found the similar results that the
Fig. 12. Optical micrograph of M-A constituent (etched by Lepera) with heat input of (a) 10 kJ/cm, (b) 30 kJ/cm, (c) 50 kJ/cm and (d) 100 kJ/cm. 8
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Fig. 14. Effect of M-A constituent morphology on crack propagation: (a) and (b)Island M-A constituents arrest the cracks at 10 kJ/cm and 30 kJ/cm, respectively; (c) Deformation of fine-stringer M-A constituents at 30 kJ/cm; (d) Coarse-stringer M-A constituents cut off by a crack at 50 kJ/cm; (e) Undeformed coarse-stringer M-A constituents near the crack tip at 50 kJ/cm and (f) Interface decohesion of coarse-stringer M-A constituents at 100 kJ/cm.
Fig. 15. Distribution characteristics of M-A constituent with different heat input.
input is 10 kJ/cm and 30 kJ/cm, the small size island M-A constituents can inhibit the propagation of secondary cracks, shown in Fig. 14a and b. According to the classical Griffith theory [41], the critical stress of the island M-A constituent cracking is large due to the small size, which
debonding zones will cause crack propagation and promote brittle fracture. In fact, different morphologies of M-A constituent may have discrepant effects on fracture toughness. In this investigate, when the heat 9
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the increase of GB aggravates the deterioration of fracture toughness at 50 kJ/cm and 100 kJ/cm. However, it is noted that the fracture toughness with heat input of 30 kJ/cm is higher than that of 10 kJ/cm, despite the higher content of GB, higher volume fraction and larger size of M-A constituent. The existence of retained austenite can effectively improve the fracture toughness, but no retained austenite can be detected at 30 kJ/cm according to EBSD phase map (Fig. 6j). Therefore, the abnormal phenomenon is considered to be attributed to the influence of the effect of packet size on fracture toughness. When the heat input is 30 kJ/cm, a large number of GB is generated in CGHAZ microstructures and interlaced with LB. The packets of LB are segmented smaller by GB. The small packet size can shorten the length of microcracks and improve toughness [42]. Therefore, when the heat input is 30 kJ/cm, the fracture toughness of CGHAZ is better than that of 10 kJ/ cm.
makes it difficult to form microcracks. Therefore, these island M-A constituents can arrest the secondary cracks and dissipate more energy of crack propagation, which is beneficial to fracture toughness. Fig. 14c shows that the fine-stringer M-A constituents has obvious plastic deformation in the tip region of the main crack. The local stress concentration is decreased and the energy of crack propagation can be consumed by plastic deformation of fine-stringer M-A constituent during the fracture. Therefore, the fine-stringer M-A constituent does not excessively deteriorate the fracture toughness. However, the coarsestringer M-A constituents are easily cut off when subjected to transverse shear stress, which is helpful for microcracks propagation, shown in Fig. 12d. It is also observed that there is no obvious deformation of the coarse-stringer M-A constituents at the crack tip region (Fig. 12e), which implies that these coarse-stringer M-A constituents are unfavorable to fracture toughness. In addition, Fig. 12f shows that the coarsestringer M-A constituents deboned from the matrix and linked other debonding regions to accelerate crack propagation. Therefore, the fracture toughness is significantly decreased when there are lots of coarse-stringer M-A constituents. In addition, previous studies [8,10,35,37] have confirmed that the massive M-A constituents are prone to initiate microcracks due to the low critical stress of microcracks formation. At the same time, the massive M-A constituents can also debond from the matrix, which leads to serious deterioration of fracture toughness. Based on the above discussion, in order to further confirm the effect of different morphologies of M-A constituent on fracture toughness, the M-A constituents are classified into four types: island, fine-stringer, coarse-stringer and massive, respectively. The single M-A constituent is characterized by its area and aspect ratio (maximal Feret diameter Lmax/minimal Feret diameter Lmin). Previous studies have shown that the average area of many island M-A constituents is considered < 5 um2, while the aspect ratio of stringer M-A constituent is usually > 4 [8,10,38−40]. Fig. 15 shows the detailed morphological distribution of M-A constituent in simulated CGHAZ with different heat input. When the heat input is 10 kJ/cm and 30 kJ/cm, the island M-A constituent (area < 5 um2 and aspect ratio < 4) can arrest the secondary cracks and the area fractions is 86% and 85%, respectively, which is beneficial to fracture toughness. When the heat input rises to 50 kJ/cm and 100 kJ/cm, the island M-A constituent decreases significantly, and the area fraction is only 51% and 35%, respectively, which is detrimental to fracture toughness. In addition, with the increase of heat input, the area fraction of fine-stringer M-A constituent (area < 5 um2 and aspect ratio > 4) decreases from 7% to 1%, whereas that of coarse-stringer MA constituent (area > 5 um2 and aspect ratio > 4) increases from 1% to 7%. The fine-stringer M-A constituent has obvious plastic deformation without excessive damage to fracture toughness, while the coarsestringer M-A constituent exhibits no obvious plastic deformation and is easy to crack and debond from matrix, which seriously damages the fracture toughness. The massive M-A constituent (area > 5 um2 and aspect ratio < 4) is also prone to cracking and debonding from matrix. The area fraction of massive M-A constituent increases dramatically from only 6% to 57% with the increase of heat input, which is one of the reasons for the greatly deterioration of fracture toughness at high heat input (50–100 kJ/cm). Furthermore, the content of LB and GB also significantly influences the fracture toughness of CGHAZ. With the increase of heat input from 10 kJ/cm to 100 kJ/cm, the proportion of LB in CGHAZ decreases dramatically from 91% to 7%. Previous EBSD analyses have concluded that there are a large number of high misorientation boundaries in the LB, as shown in Figs. 6e and 7. High misorientation boundaries can arrest the microcracks during crack propagation, which effectively improves the fracture toughness [10,20,21]. On the contrary, GB increases from 9% to 93% with the increase of heat input from 10 kJ/cm to 100 kJ/cm. The GB is ineffective in preventing the propagation of microcracks due to the existence of many low misorientation boundaries (Figs. 6h and 7). Therefore, the reduction of LB accompanied with
5. Conclusions (1) The microstructure of simulated CGHAZ changes from lath bainite to granular bainite with the increase of heat input, accompanied by the increase of prior austenite grain. The lath bainite has faster transformation rate and more abundant high misorientation boundaries than granular bainite. Excessive coarsening of granular bainite occurs when heat input increases to higher than 50 kJ/cm. (2) The welding heat input significantly influences the content and morphology of M-A constitute. With the increase of heat input, the area fraction of M-A constituent increases from 15.2% to 28.0%, the area proportion of island and fine stringer M-A constituent decrease from 86% and 7% to 35% and 1% respectively, and that of coarsestringer and massive M-A constituent increase from 1% and 6% to 7% and 57% respectively. (3) The fracture toughness of simulated CGHAZ is proved to be sensitive to the morphology of M-A constitute. The island M-A constituent can arrest microcracks, which is beneficial to fracture toughness. The fine stringer M-A constituent does not greatly damage to fracture toughness due to the deformation. However, the coarse stringer and massive M-A constituent can crack and debond from the matrix, which seriously deteriorates the fracture toughness. (4) When the heat input increases from 10 kJ/cm to 30 kJ/cm, the CTOD value at −20 °C increases from 0.079 mm to 0.12 mm, and dramatically decreases as the heat input increases to higher than 50 kJ/cm. The packet of lath bainite is segmented smaller by granular bainite, resulting in higher CTOD values at 30 kJ/cm. and the significant decrease of CTOD value at the heat input of higher than 50 kJ/cm is due to the large content of GB, coarsening of bainite packet, and increase in content of coarse-stringer and massive M-A constituent. Acknowledgements The author is sincerely grateful to the financial support from the National Natural Science Foundation of China (grant No. 51804217 and 51774213), the State Key Laboratory of Metal Material for Marine Equipment and Application (grant No. SKLMEA-K201602) and the Regional Demonstration Project of Marine Economic Innovation and Development (grant No. BHSF2017-10). References [1] Y. Sumi, H. Yajima, M. Toyosada, T. Yoshikawa, S. Aihara, K. Gotoh, Y. Ogawa, T. Matsumoto, K. Hirota, H. Hirasawa, Fracture control of extremely thick welded steel plates applied to the deck structure of large container ships, J. Mar. Sci. Technol. 18 (2013) 497–514. [2] T. Handa, S. Suzuki, N. Kiji, M. Toyoda, T. Miyata, Effect of unwelded length on behaviour of brittle crack arrest in T-joint structure, Weld. Int. 23 (2009) 640–647. [3] J.Y. Wu, B. Wang, B.X. Wang, R.D.K. Misra, Z.D. Wang, Toughness and ductility
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