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Fatigue crack growth behavior of the simulated HAZ of 800 MPa grade high-performance steel
Materials Science and Engineering A 528 (2011) 2331–2338
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Fatigue crack growth behavior of the simulated HAZ of 800 MPa grade high-performance steel Sanghoon Kim a , Donghwan Kang b , Tae-Won Kim b , Jongkwan Lee c , Changhee Lee a,∗ a b c
Division of Materials Science and Engineering, Hanyang University, 17 Haengdang-dong, Seongdong-gu, Seoul 133-791, Republic of Korea School of Mechanical Engineering, Hanyang University, 17 Haengdang-dong, Seongdong-gu, Seoul 133-791, Republic of Korea Research Institute of Industrial Science and Technology, 75-9, Youngcheon, Dongtan, Hwaseong, Gyeonggi-do 445-813, Republic of Korea
a r t i c l e
i n f o
Article history: Received 20 August 2010 Received in revised form 4 November 2010 Accepted 29 November 2010 Available online 5 December 2010 Keywords: Fatigue crack High performance steel for bridges and buildings Local brittle zone M–A constituent Crack propagation test
a b s t r a c t The present study focuses on the fatigue properties in the weld heat-affected zone (HAZ) of 800 MPa grade high-performance steel, which is commonly used in bridges and buildings. Single- and multi-pass HAZs were simulated by the Gleeble system. Fatigue properties were estimated using a crack propagation test under a 0.3 stress ratio and 0.1 load frequencies. The microstructures and fracture surfaces were analyzed by optical microscopy, scanning electron microscopy, and transmission electron microscopy. The results of the crack propagation test showed that the fatigue crack growth rate of coarse-grained HAZ (CGHAZ) was faster than fine-grained HAZ (FGHAZ), although both regions have identical fully martensite microstructures, because FGHAZ has smaller prior austenite grain and martensite packet sizes, which can act as effective barriers to crack propagation. The fatigue crack growth rate of intercritically reheated CGHAZ (ICCGHAZ) was the fastest among local zones in the HAZ, due to rapid crack initiation and propagation via the massive martensite–austenite (M–A) constituent. © 2010 Elsevier B.V. All rights reserved.
1. Introduction Recently, demands for balanced development between regions and the effective use of space are growing as by-products of increasing population and population density. Therefore, new structures such as bridges and buildings are being built bigger and higher, and high performance steels, which have excellent strength properties for their weights, are in great demand as materials for construction. Considering these demands, high-performance steel for bridges and buildings (HSB) has been developed to meet modern construction needs. To obtain the necessary mechanical properties, various alloying elements are added to HSB to increase hardenability, such as manganese, boron, and niobium, instead of carbon, and a thermo-mechanical controlled process (TMCP) is applied during fabrication. Considering the production cost, the addition of expensive alloying elements, such as Ni and Mo, is decreased. However, the welding process is still required to construct structures, and various heat-affected zones (HAZs) with different microstructures and mechanical properties from base steel are formed according to peak temperature (the distance from the heat source) after the welding process [1]. Hence, these HAZs have lower impact toughness and fatigue resistance than base steel and act as local brittle zones (LBZs) in the welding part.
∗ Corresponding author. Tel.: +82 2 2220 0388; fax: +82 2 2299 0389. E-mail address: [email protected] (C. Lee). 0921-5093/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2010.11.089
Thus, many investigations of the mechanical properties of HAZs have been carried out [2–6]. In most of these previous studies, impact toughness and fatigue tests were carried out to estimate the mechanical properties of HAZs according to the distance from the fusion line of the weld metal. However, each HAZ possesses different microstructures and mechanical properties, and distinguishing each HAZ is difficult due to their continuity and narrowness. Hence, studies of the mechanical properties of each HAZ have rarely been performed [7,8]. In particular, studies of the fatigue properties of each HAZ have not been carried out, and therefore fatigue properties remain poorly understood. In this study, the fatigue properties of weld HAZs in 800 MPa grade HSB were investigated to consider the relationships between the microstructures of each HAZ and fatigue using simulated HAZ specimens. 2. Experimental procedures 2.1. Materials In this study, 800 MPa grade HSB (HSB800, KS D 3868:2009) was used for HAZ fatigue property tests. HSB800 is a recently developed steel that is currently being used in the construction of long span bridges. The chemical composition and mechanical properties of HSB800 are shown in Table 1. A thermo-mechanical controlled process (TMCP) was applied for the fabrication of base steel. Therefore, HSB800 has a bainite microstructure, and a pancaked prior austen-
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S. Kim et al. / Materials Science and Engineering A 528 (2011) 2331–2338
Table 1 Chemical composition (wt.%) and mechanical properties of base steels. Base steel HSB800
C ≤0.10
Mechanical properties
YS ≥690 MPa
Si ≤0.65
Mn ≤2.20
P ≤0.015
S ≤0.006
Cu 0.10–0.50
TS ≥800 MPa
Cr 0.45–0.75
Ni 0.05–0.80
Nb + V + Ti ≤0.12
Ceq ≤0.55
Pcm ≤0.25
CVN(−20 ◦ C) ≥47 J
EL ≥22%
Ceq = C + Mn/6 + Si/24 + Ni/40 + Cr/5 + Mo/4 + V/14. Pcm = C + Si/30 + Mn/20 + Cu/20 + Ni/60 + Cr/20 + Mo/15 + V/10 + 5B.
Fig. 1. Optical micrograph and SEM image of base steel.
ite grain boundary can be also observed in the matrix, as shown in Fig. 1. As a result, HSB800 has a good balance between strength and toughness despite its decreased carbon and micro alloying elements, as shown in Table 1. HSB800 is also suitable for welding because it was designed to possess low carbon equivalents (Ceq ) and welding crack susceptibility index (Pcm ) values. 2.2. Simulations The Gleeble system was used to perform simulations of various weld HAZs before the crack propagation test. Table 2 shows the welding cycles used in simulations, with the temperature cycle calculated using the Rosenthal equation for thick plate [9]. The Rosenthal equation is as follows: T − T0 =
q/v exp 2t
−
r2 4at
(1)
where T is the temperature (K), T0 is the pre-heat temperature, q/v is the heat input (kJ/cm), is the thermal conductivity of carbon steel (J/msK), t is the time, a is the thermal diffusivity of carbon steel (m2 /s), and r is the radial distance from the heat source. In this study, coarse-grained (CG) and fine-grained (FG) HAZs were simulated for single-pass welding and intercritically reheated (IC) CGHAZ was simulated for multi-pass welding, because CG and ICCGHAZ are known to act as local brittle zones in welds. Each thermal cycle of HAZs was approximately equivalent to that of welding with 30 kJ/cm heat input. The peak temperatures of CG, FG, and ICHAZ are 1350 ◦ C, 1050 ◦ C, and 850 ◦ C, respectively. To confirm the peak temperature of ICHAZ, Ac1 and Ac3 temperatures were esti-
mated using a dilatometer, and they were confirmed as 780 and 880 ◦ C, respectively. 2.3. Crack propagation tests To investigate the relationships between microstructures and fatigue behaviors of various HAZs, crack propagation tests were carried out at room temperature using base steel and simulated HAZ specimens. A schematic diagram of the crack propagation test setup and specimen geometry is shown in Fig. 2 [10]. The cyclic load for fatigue crack propagation was applied by three-point bending under a 0.3 stress ratio and 0.1 load frequencies. The detailed conditions of the crack propagation tests on base steel and each HAZ are outlined in Table 3. 2.4. Microstructural observations The analysis of microstructures was performed using optical microscopy (OM), and transmission electron microscopy (TEM). For TEM observation, the sample was mechanically polished below 100 m and then thinned by jet-polishing (Tenupol-3) using a solution of 5 vol.% perchloric acid and 95 vol.% methanol at −40 ◦ C. TEM specimens were examined by a FEI Tecnai F20 (∼200 kV). For observation of the fracture surfaces and cross sections of the fractured specimens after crack propagation tests, scanning electron microscopy (SEM) was used. To observe martensite–austenite (M–A) constituents, specimens were etched by a two-step etching process using chemical and electrochemical methods, as summarized in Table 4. The sizes of the prior austenite grains, martensite packets, and M–A constituents were
Table 2 Thermal cycles for HAZ simulations. HAZ
Peak temp. (◦ C)
Heating rate (◦ C/s)
Cooling rate (PT ∼ 800 ◦ C/s)
Cooling rate (800 ∼ RT ◦ C/s)
Cycle and heat input
CGHAZ FGHAZ ICCGHAZ
1350 1050 850 1350
260 90 30 260
70 50 30 70
35 30
1 cycle (30 kJ/cm) 1 cycle (30 kJ/cm) 2 cycle (30 kJ/cm)
35
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Table 3 Conditions of crack propagation tests. R (stress ratio) Base steel CGHAZ FGHAZ ICCGHAZ
0.3
f (frequency) 0.1
A0 (initial crack length, mm)
Pmax (maximum load, kgf)
Pmin (minimum load, kgf)
6.809 7.319 6.625 6.339
6354 6447 6438 6487
1957 1942 1931 1901
Temp. (temperature, ◦ C) 26
Table 4 Etchant used in two-step etching for observation of M–A constituents. 1st step 2nd step
Solution A: Na2 SO4 1 g + distilled water Solution B: picric acid 4 g + 100 ml ethanol → 1:1 portion of mixing NaOH 25 g + picric acid 5 g + distilled water → 5 V, 140 s electrolytic etching
measured by an image analyzer (Image-Pro Plus, Media Cybernetics). 3. Results 3.1. Fatigue behavior Crack propagation tests were performed using base steel and simulated CG, FG, and ICCGHAZ specimens to analyze the fatigue properties of various HAZs. CG and ICCGHAZ were selected for crack propagation tests because they have relatively higher ductile brittle transition temperatures (DBTTs) than other HAZs, although HSB800 base steel and various HAZs have excellent mechanical properties compared to other high-strength low-alloy (HSLA) steels. The crack propagation tests that were conducted on base steel and FGHAZ were also carried out for comparisons with CG and ICCGHAZ. The fatigue crack growth rates (da/dN) are shown in Fig. 3 as a function of the stress intensity factor range (K) for base steel and each HAZ. The fatigue crack growth rates increase significantly with increases in the stress intensity factor range in all specimens. Base steel has the highest resistance to crack growth compared with other HAZs under identical load conditions. In the case of HAZ specimens, ICCGHAZ possesses the highest crack propagation rate, and the crack propagation rate of FGHAZ is lower than that of CGHAZ. These results confirmed that sensitivity to fatigue cracks for ICCG and CGHAZ was high in welding part.
Fig. 2. Schematic diagram of the crack propagation test setup and specimen geometry.
3.2. Microstructures of each HAZ The results of microstructural observation of each HAZ by OM and TEM are shown in Fig. 4. In CG and FGHAZ, martensite is the main microstructural component due to the rapid cooling rate during the welding thermal cycle despite very low carbon content. Thus, a carbide-free lath structure was observed in the CG and FGHAZ specimens by TEM analysis, and consequently these specimens have high hardness and strength. However, CGHAZ withstood a higher peak temperature than FGHAZ, and the prior austenite grain and packet size of CGHAZ were greater than that of FGHAZ. The lath width of FGHAZ was also narrower. Meanwhile, as shown in Fig. 4, ICCGHAZ was composed of lath type ferrite (or tempered martensite) and M–A constituents. M–A constituents can be formed by rapid heating and cooling due to carbon enrichment during the welding cycle, and therefore M–A constituents possess high carbon contents and hardness. Therefore, M–A constituents are one of the main factors involved in toughness deterioration [7,8,11–15]. M–A constituents can be morphologically classified into two types: elongated and massive. Changes in type are governed by cooling rates and formation sites. In this study,
Fig. 3. Fatigue crack growth behaviors of base steel and HAZs.
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Fig. 4. Optical micrographs and TEM images of CG, FG, and ICCGHAZ.
the two types of M–A constituent were observed simultaneously. The TEM micrograph in Fig. 4 shows an elongated M–A constituent that formed at the lath boundary.
a very complicated and long crack propagation path. Thus, many secondary cracks are observed. In contrast, CG and ICCGHAZ have relatively smooth crack propagation paths, and secondary cracks are rarely detected.
3.3. Fractographs and crack propagation paths 4. Discussion SEM fractography was carried out after the crack propagation tests to verify the results of the crack propagation tests. Fig. 5 shows the results of SEM fractography. In these figures, cracks propagated from the bottom-right corner to the top-left corner. As shown in Fig. 5, transgranular type fatigue fracture occurred, and fatigue striation, which was not clear but distinguishable, was observed in every specimen. A correlation was found between the distance between each fatigue striation and the fatigue crack growth rate. In other words, when fatigue crack growth is faster, the distance between fatigue striations is broader. Thus, in this study, ICCGHAZs have the largest distances between fatigue striations. Meanwhile, Fig. 6 shows cross-sectional SEM micrographs beneath the fracture surface of base steel and the HAZ specimens. Base steel, which has the slowest fatigue crack growth rate, has
4.1. Crack propagation behavior in CG and FGHAZ As mentioned above, CG and FGHAZ were composed of identical lath martensite structures. However, as shown in Figs. 3 and 6, they have significantly different fatigue crack resistances and crack propagation paths. To observe the crack propagation paths in detail, Fig. 7 shows a high magnification SEM image of a cross section. Fatigue cracks mainly propagated into prior austenite grain or martensite packets, and lath boundaries may not have affected the crack propagation in either specimen. Almost all secondary crack propagation was stopped by the prior austenite grain or martensite packet boundary. Thus, the prior austenite grain and martensite packet boundary may significantly affect fatigue crack propagation.
S. Kim et al. / Materials Science and Engineering A 528 (2011) 2331–2338
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Fig. 5. Fractography of base steel and HAZs specimen.
According to previous studies, crack propagation is mainly affected by the microstructure of the matrix, and obstacles such as the secondary phase in the matrix and grain or subgrain boundaries [16]. Especially, grain and subgrain boundaries such as prior austenite grain boundaries, martensite packet boundaries, block boundaries, and lath boundaries act as effective barriers to crack propagation in lath martensite [17]. Thus, various grain and sub-
grain sizes of CG and FGHAZ were measured (Fig. 8). CGHAZ exhibits large values of prior austenite grain, martensite packet, and lath width compared to FGHAZ. The difference in prior austenite grain size between CG and FGHAZ is especially great, and the prior austenite grain boundary is the most effective barrier because it has a relatively high angle boundary compared to other boundaries. Therefore, FGHAZ incorporates many effective barriers to
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Fig. 6. SEM images of cross sections beneath fracture surfaces of base steel and HAZ specimens.
Fig. 7. Crack propagation path of CG and FGHAZ.
Fig. 8. Martensite packet and prior austenite grain size and lath width of CG and FGHAZ.
Fig. 9. Crack propagation path of ICCGHAZ.
S. Kim et al. / Materials Science and Engineering A 528 (2011) 2331–2338
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Fig. 10. SEM images and size distribution of M–A constituent in base steel and HAZs.
crack propagation and excellent fatigue crack resistance compared to CGHAZ, despite their identical microstructures, lath martensites. 4.2. Crack propagation behavior in ICCGHAZ ICCGHAZ has a relatively large amount of critical microstructure that is composed of M–A constituent because of its unique thermal cycle. Large amounts of M–A constituent are detrimental to impact toughness [11,14,17]. However, the morphological effects of M–A constituents on crack propagation are not fully understood [11]. In this study, we assumed that the M–A constituent was the main factor causing the deterioration of fatigue crack resistance (Figs. 3 and 5) although some prior studies reported that the M–A constituent, which has a film structure, can be an effective barrier to crack propagation [17]. To estimate the effects of M–A constituents on crack propagation, high magnification cross sectional SEM images were analyzed. As shown in Fig. 9, M–A constituents were observed on the crack propagation path and identified as massive M–A constituents. M–A constituents act at crack initiation sites, and cracks can therefore propagate along the M–A constituent [7,11,15]. The observed massive M–A constituent exists mainly at prior austenite grain or martensite packet boundaries because, during the thermal cycle, nucleation and growth of austenite at prior austenite grain and martensite packet boundaries easily occurs. Elongated M–A constituents are mainly observed at the lath boundary in the matrix, and the effect of elongated M–A constituents on crack propagation may therefore be nonexistent. Meanwhile, the size distribution of M–A constituents was measured in ICCGHAZ and the size distribution of M–A constituents in base steel and ICHAZ were also measured for comparison. As a result, a considerable amount of M–A constituent was found in ICCGHAZ compared with base steel and ICHAZ, as shown in Fig. 10. Interestingly, M–A constituents larger than 1 m2 were observed only in ICCGHAZ. We concluded that the deterioration of fatigue crack resistance in ICCGHAZ is caused by crack initiation and propagation by massive M–A constituents that were greater than 1 m2 in size.
5. Conclusions In this study, the fatigue properties in various weld HAZs of HSB800 were investigated to consider the relationships between the microstructures of each HAZ and fatigue. Our conclusions were as follows: (1) Crack propagation tests were carried out on base steel, CG, FG, and ICCGHAZ. Our results showed that the fatigue crack growth rate of ICCGHAZ was the fastest, while base steel had the best fatigue crack resistance. (2) The fatigue crack growth rate of CGHAZ was faster than that of FGHAZ although both regions have identical fully martensite microstructures, because FGHAZ has smaller prior austenite grain and martensite packet sizes that act as effective barriers to crack propagation. (3) The fatigue crack growth rate of ICCGHAZ was the fastest among local zones because fatigue cracks were rapidly initiated and propagated through massive M–A constituents that were larger than 1 m2 in size.
Acknowledgement The authors would like to acknowledge the financial support of the Research Institute of Industrial Science & Technology (RIST), Korea. References [1] K. Easterling, Introduction to the Physical Metallurgy of Welding, Butterworth, 1983, pp. 104–155. [2] L.W. Tsay, T.S. Chern, C.Y. Gau, J.R. Yang, Int. J. Fatigue 21 (1999) 857. [3] L.W. Tsay, C.C. Liu, Y.-H. Chao, Y.H. Shieh, Mater. Sci. Eng. A 299 (2001) 16. [4] J.H. Kim, Y.J. Oh, I.S. Hwang, D.J. Kim, J.T. Kim, J. Nucl. Mater. 2999 (2001) 132. [5] J. Jang, B. Lee, J. Ju, D. Kwon, W. Kim, Eng. Fract. Mech. 70 (2003) 1245. [6] B.J. Kim, S.H. Ryu, B.S. Lim, Met. Mater. Int. 10 (2004) 19. [7] J. Jang, J. Ju, B. Lee, D. Kwon, W. Kim, Mater. Sci. Eng. A 340 (2003) 68. [8] S. Moeinifar, A.H. Kokabi, H.R. Madaah Hosseini, Mater. Des. 31 (2010) 2948. [9] K. Easterling, Introduction to the Physical Metallurgy of Welding, Butterworth, 1983, pp. 17–33.
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[10] ISO 11782-2, International Organization for Standardization, 1998. [11] F. Matsuda, K. Ikeuchi, Y. Fukada, Y. Horii, H. Okada, T. Shiwaku, C. Shiga, S. Suzuki, Trans. JWRI 24 (1995) 1. [12] K. Ikeuchi, J. Liao, H. Tanabe, F. Matsuda, ISIJ Int. 35 (1995) 1203. [13] J. Janovec, M. Takahashi, T. Kuroda, K. Ikeuchi, ISIJ Int. 40 (2000) S44.
[14] Y. Li, D.N. Crowther, M.J.W. Green, P.S. Mitchell, T.N. Baker, ISIJ Int. 41 (2001) 46. [15] Y. Shi, Z. Han, J. Mater. Process. Technol. 207 (2008) 30. [16] S.Y. Shin, S.Y. Han, B. Hwang, C.G. Lee, S. Lee, Mater. Sci. Eng. A 517 (2009) 212. [17] Y. Zhong, F. Xiao, J. Zhang, Y. Shan, W. Wang, K. Yang, Acta Mater. 54 (2006) 435.
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Fatigue crack growth behavior of the simulated HAZ of 800 MPa grade high-performance steel Sanghoon Kim a , Donghwan Kang b , Tae-Won Kim b , Jongkwan Lee c , Changhee Lee a,∗ a b c
Division of Materials Science and Engineering, Hanyang University, 17 Haengdang-dong, Seongdong-gu, Seoul 133-791, Republic of Korea School of Mechanical Engineering, Hanyang University, 17 Haengdang-dong, Seongdong-gu, Seoul 133-791, Republic of Korea Research Institute of Industrial Science and Technology, 75-9, Youngcheon, Dongtan, Hwaseong, Gyeonggi-do 445-813, Republic of Korea
a r t i c l e
i n f o
Article history: Received 20 August 2010 Received in revised form 4 November 2010 Accepted 29 November 2010 Available online 5 December 2010 Keywords: Fatigue crack High performance steel for bridges and buildings Local brittle zone M–A constituent Crack propagation test
a b s t r a c t The present study focuses on the fatigue properties in the weld heat-affected zone (HAZ) of 800 MPa grade high-performance steel, which is commonly used in bridges and buildings. Single- and multi-pass HAZs were simulated by the Gleeble system. Fatigue properties were estimated using a crack propagation test under a 0.3 stress ratio and 0.1 load frequencies. The microstructures and fracture surfaces were analyzed by optical microscopy, scanning electron microscopy, and transmission electron microscopy. The results of the crack propagation test showed that the fatigue crack growth rate of coarse-grained HAZ (CGHAZ) was faster than fine-grained HAZ (FGHAZ), although both regions have identical fully martensite microstructures, because FGHAZ has smaller prior austenite grain and martensite packet sizes, which can act as effective barriers to crack propagation. The fatigue crack growth rate of intercritically reheated CGHAZ (ICCGHAZ) was the fastest among local zones in the HAZ, due to rapid crack initiation and propagation via the massive martensite–austenite (M–A) constituent. © 2010 Elsevier B.V. All rights reserved.
1. Introduction Recently, demands for balanced development between regions and the effective use of space are growing as by-products of increasing population and population density. Therefore, new structures such as bridges and buildings are being built bigger and higher, and high performance steels, which have excellent strength properties for their weights, are in great demand as materials for construction. Considering these demands, high-performance steel for bridges and buildings (HSB) has been developed to meet modern construction needs. To obtain the necessary mechanical properties, various alloying elements are added to HSB to increase hardenability, such as manganese, boron, and niobium, instead of carbon, and a thermo-mechanical controlled process (TMCP) is applied during fabrication. Considering the production cost, the addition of expensive alloying elements, such as Ni and Mo, is decreased. However, the welding process is still required to construct structures, and various heat-affected zones (HAZs) with different microstructures and mechanical properties from base steel are formed according to peak temperature (the distance from the heat source) after the welding process [1]. Hence, these HAZs have lower impact toughness and fatigue resistance than base steel and act as local brittle zones (LBZs) in the welding part.
∗ Corresponding author. Tel.: +82 2 2220 0388; fax: +82 2 2299 0389. E-mail address: [email protected] (C. Lee). 0921-5093/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2010.11.089
Thus, many investigations of the mechanical properties of HAZs have been carried out [2–6]. In most of these previous studies, impact toughness and fatigue tests were carried out to estimate the mechanical properties of HAZs according to the distance from the fusion line of the weld metal. However, each HAZ possesses different microstructures and mechanical properties, and distinguishing each HAZ is difficult due to their continuity and narrowness. Hence, studies of the mechanical properties of each HAZ have rarely been performed [7,8]. In particular, studies of the fatigue properties of each HAZ have not been carried out, and therefore fatigue properties remain poorly understood. In this study, the fatigue properties of weld HAZs in 800 MPa grade HSB were investigated to consider the relationships between the microstructures of each HAZ and fatigue using simulated HAZ specimens. 2. Experimental procedures 2.1. Materials In this study, 800 MPa grade HSB (HSB800, KS D 3868:2009) was used for HAZ fatigue property tests. HSB800 is a recently developed steel that is currently being used in the construction of long span bridges. The chemical composition and mechanical properties of HSB800 are shown in Table 1. A thermo-mechanical controlled process (TMCP) was applied for the fabrication of base steel. Therefore, HSB800 has a bainite microstructure, and a pancaked prior austen-
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S. Kim et al. / Materials Science and Engineering A 528 (2011) 2331–2338
Table 1 Chemical composition (wt.%) and mechanical properties of base steels. Base steel HSB800
C ≤0.10
Mechanical properties
YS ≥690 MPa
Si ≤0.65
Mn ≤2.20
P ≤0.015
S ≤0.006
Cu 0.10–0.50
TS ≥800 MPa
Cr 0.45–0.75
Ni 0.05–0.80
Nb + V + Ti ≤0.12
Ceq ≤0.55
Pcm ≤0.25
CVN(−20 ◦ C) ≥47 J
EL ≥22%
Ceq = C + Mn/6 + Si/24 + Ni/40 + Cr/5 + Mo/4 + V/14. Pcm = C + Si/30 + Mn/20 + Cu/20 + Ni/60 + Cr/20 + Mo/15 + V/10 + 5B.
Fig. 1. Optical micrograph and SEM image of base steel.
ite grain boundary can be also observed in the matrix, as shown in Fig. 1. As a result, HSB800 has a good balance between strength and toughness despite its decreased carbon and micro alloying elements, as shown in Table 1. HSB800 is also suitable for welding because it was designed to possess low carbon equivalents (Ceq ) and welding crack susceptibility index (Pcm ) values. 2.2. Simulations The Gleeble system was used to perform simulations of various weld HAZs before the crack propagation test. Table 2 shows the welding cycles used in simulations, with the temperature cycle calculated using the Rosenthal equation for thick plate [9]. The Rosenthal equation is as follows: T − T0 =
q/v exp 2t
−
r2 4at
(1)
where T is the temperature (K), T0 is the pre-heat temperature, q/v is the heat input (kJ/cm), is the thermal conductivity of carbon steel (J/msK), t is the time, a is the thermal diffusivity of carbon steel (m2 /s), and r is the radial distance from the heat source. In this study, coarse-grained (CG) and fine-grained (FG) HAZs were simulated for single-pass welding and intercritically reheated (IC) CGHAZ was simulated for multi-pass welding, because CG and ICCGHAZ are known to act as local brittle zones in welds. Each thermal cycle of HAZs was approximately equivalent to that of welding with 30 kJ/cm heat input. The peak temperatures of CG, FG, and ICHAZ are 1350 ◦ C, 1050 ◦ C, and 850 ◦ C, respectively. To confirm the peak temperature of ICHAZ, Ac1 and Ac3 temperatures were esti-
mated using a dilatometer, and they were confirmed as 780 and 880 ◦ C, respectively. 2.3. Crack propagation tests To investigate the relationships between microstructures and fatigue behaviors of various HAZs, crack propagation tests were carried out at room temperature using base steel and simulated HAZ specimens. A schematic diagram of the crack propagation test setup and specimen geometry is shown in Fig. 2 [10]. The cyclic load for fatigue crack propagation was applied by three-point bending under a 0.3 stress ratio and 0.1 load frequencies. The detailed conditions of the crack propagation tests on base steel and each HAZ are outlined in Table 3. 2.4. Microstructural observations The analysis of microstructures was performed using optical microscopy (OM), and transmission electron microscopy (TEM). For TEM observation, the sample was mechanically polished below 100 m and then thinned by jet-polishing (Tenupol-3) using a solution of 5 vol.% perchloric acid and 95 vol.% methanol at −40 ◦ C. TEM specimens were examined by a FEI Tecnai F20 (∼200 kV). For observation of the fracture surfaces and cross sections of the fractured specimens after crack propagation tests, scanning electron microscopy (SEM) was used. To observe martensite–austenite (M–A) constituents, specimens were etched by a two-step etching process using chemical and electrochemical methods, as summarized in Table 4. The sizes of the prior austenite grains, martensite packets, and M–A constituents were
Table 2 Thermal cycles for HAZ simulations. HAZ
Peak temp. (◦ C)
Heating rate (◦ C/s)
Cooling rate (PT ∼ 800 ◦ C/s)
Cooling rate (800 ∼ RT ◦ C/s)
Cycle and heat input
CGHAZ FGHAZ ICCGHAZ
1350 1050 850 1350
260 90 30 260
70 50 30 70
35 30
1 cycle (30 kJ/cm) 1 cycle (30 kJ/cm) 2 cycle (30 kJ/cm)
35
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Table 3 Conditions of crack propagation tests. R (stress ratio) Base steel CGHAZ FGHAZ ICCGHAZ
0.3
f (frequency) 0.1
A0 (initial crack length, mm)
Pmax (maximum load, kgf)
Pmin (minimum load, kgf)
6.809 7.319 6.625 6.339
6354 6447 6438 6487
1957 1942 1931 1901
Temp. (temperature, ◦ C) 26
Table 4 Etchant used in two-step etching for observation of M–A constituents. 1st step 2nd step
Solution A: Na2 SO4 1 g + distilled water Solution B: picric acid 4 g + 100 ml ethanol → 1:1 portion of mixing NaOH 25 g + picric acid 5 g + distilled water → 5 V, 140 s electrolytic etching
measured by an image analyzer (Image-Pro Plus, Media Cybernetics). 3. Results 3.1. Fatigue behavior Crack propagation tests were performed using base steel and simulated CG, FG, and ICCGHAZ specimens to analyze the fatigue properties of various HAZs. CG and ICCGHAZ were selected for crack propagation tests because they have relatively higher ductile brittle transition temperatures (DBTTs) than other HAZs, although HSB800 base steel and various HAZs have excellent mechanical properties compared to other high-strength low-alloy (HSLA) steels. The crack propagation tests that were conducted on base steel and FGHAZ were also carried out for comparisons with CG and ICCGHAZ. The fatigue crack growth rates (da/dN) are shown in Fig. 3 as a function of the stress intensity factor range (K) for base steel and each HAZ. The fatigue crack growth rates increase significantly with increases in the stress intensity factor range in all specimens. Base steel has the highest resistance to crack growth compared with other HAZs under identical load conditions. In the case of HAZ specimens, ICCGHAZ possesses the highest crack propagation rate, and the crack propagation rate of FGHAZ is lower than that of CGHAZ. These results confirmed that sensitivity to fatigue cracks for ICCG and CGHAZ was high in welding part.
Fig. 2. Schematic diagram of the crack propagation test setup and specimen geometry.
3.2. Microstructures of each HAZ The results of microstructural observation of each HAZ by OM and TEM are shown in Fig. 4. In CG and FGHAZ, martensite is the main microstructural component due to the rapid cooling rate during the welding thermal cycle despite very low carbon content. Thus, a carbide-free lath structure was observed in the CG and FGHAZ specimens by TEM analysis, and consequently these specimens have high hardness and strength. However, CGHAZ withstood a higher peak temperature than FGHAZ, and the prior austenite grain and packet size of CGHAZ were greater than that of FGHAZ. The lath width of FGHAZ was also narrower. Meanwhile, as shown in Fig. 4, ICCGHAZ was composed of lath type ferrite (or tempered martensite) and M–A constituents. M–A constituents can be formed by rapid heating and cooling due to carbon enrichment during the welding cycle, and therefore M–A constituents possess high carbon contents and hardness. Therefore, M–A constituents are one of the main factors involved in toughness deterioration [7,8,11–15]. M–A constituents can be morphologically classified into two types: elongated and massive. Changes in type are governed by cooling rates and formation sites. In this study,
Fig. 3. Fatigue crack growth behaviors of base steel and HAZs.
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Fig. 4. Optical micrographs and TEM images of CG, FG, and ICCGHAZ.
the two types of M–A constituent were observed simultaneously. The TEM micrograph in Fig. 4 shows an elongated M–A constituent that formed at the lath boundary.
a very complicated and long crack propagation path. Thus, many secondary cracks are observed. In contrast, CG and ICCGHAZ have relatively smooth crack propagation paths, and secondary cracks are rarely detected.
3.3. Fractographs and crack propagation paths 4. Discussion SEM fractography was carried out after the crack propagation tests to verify the results of the crack propagation tests. Fig. 5 shows the results of SEM fractography. In these figures, cracks propagated from the bottom-right corner to the top-left corner. As shown in Fig. 5, transgranular type fatigue fracture occurred, and fatigue striation, which was not clear but distinguishable, was observed in every specimen. A correlation was found between the distance between each fatigue striation and the fatigue crack growth rate. In other words, when fatigue crack growth is faster, the distance between fatigue striations is broader. Thus, in this study, ICCGHAZs have the largest distances between fatigue striations. Meanwhile, Fig. 6 shows cross-sectional SEM micrographs beneath the fracture surface of base steel and the HAZ specimens. Base steel, which has the slowest fatigue crack growth rate, has
4.1. Crack propagation behavior in CG and FGHAZ As mentioned above, CG and FGHAZ were composed of identical lath martensite structures. However, as shown in Figs. 3 and 6, they have significantly different fatigue crack resistances and crack propagation paths. To observe the crack propagation paths in detail, Fig. 7 shows a high magnification SEM image of a cross section. Fatigue cracks mainly propagated into prior austenite grain or martensite packets, and lath boundaries may not have affected the crack propagation in either specimen. Almost all secondary crack propagation was stopped by the prior austenite grain or martensite packet boundary. Thus, the prior austenite grain and martensite packet boundary may significantly affect fatigue crack propagation.
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Fig. 5. Fractography of base steel and HAZs specimen.
According to previous studies, crack propagation is mainly affected by the microstructure of the matrix, and obstacles such as the secondary phase in the matrix and grain or subgrain boundaries [16]. Especially, grain and subgrain boundaries such as prior austenite grain boundaries, martensite packet boundaries, block boundaries, and lath boundaries act as effective barriers to crack propagation in lath martensite [17]. Thus, various grain and sub-
grain sizes of CG and FGHAZ were measured (Fig. 8). CGHAZ exhibits large values of prior austenite grain, martensite packet, and lath width compared to FGHAZ. The difference in prior austenite grain size between CG and FGHAZ is especially great, and the prior austenite grain boundary is the most effective barrier because it has a relatively high angle boundary compared to other boundaries. Therefore, FGHAZ incorporates many effective barriers to
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Fig. 6. SEM images of cross sections beneath fracture surfaces of base steel and HAZ specimens.
Fig. 7. Crack propagation path of CG and FGHAZ.
Fig. 8. Martensite packet and prior austenite grain size and lath width of CG and FGHAZ.
Fig. 9. Crack propagation path of ICCGHAZ.
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Fig. 10. SEM images and size distribution of M–A constituent in base steel and HAZs.
crack propagation and excellent fatigue crack resistance compared to CGHAZ, despite their identical microstructures, lath martensites. 4.2. Crack propagation behavior in ICCGHAZ ICCGHAZ has a relatively large amount of critical microstructure that is composed of M–A constituent because of its unique thermal cycle. Large amounts of M–A constituent are detrimental to impact toughness [11,14,17]. However, the morphological effects of M–A constituents on crack propagation are not fully understood [11]. In this study, we assumed that the M–A constituent was the main factor causing the deterioration of fatigue crack resistance (Figs. 3 and 5) although some prior studies reported that the M–A constituent, which has a film structure, can be an effective barrier to crack propagation [17]. To estimate the effects of M–A constituents on crack propagation, high magnification cross sectional SEM images were analyzed. As shown in Fig. 9, M–A constituents were observed on the crack propagation path and identified as massive M–A constituents. M–A constituents act at crack initiation sites, and cracks can therefore propagate along the M–A constituent [7,11,15]. The observed massive M–A constituent exists mainly at prior austenite grain or martensite packet boundaries because, during the thermal cycle, nucleation and growth of austenite at prior austenite grain and martensite packet boundaries easily occurs. Elongated M–A constituents are mainly observed at the lath boundary in the matrix, and the effect of elongated M–A constituents on crack propagation may therefore be nonexistent. Meanwhile, the size distribution of M–A constituents was measured in ICCGHAZ and the size distribution of M–A constituents in base steel and ICHAZ were also measured for comparison. As a result, a considerable amount of M–A constituent was found in ICCGHAZ compared with base steel and ICHAZ, as shown in Fig. 10. Interestingly, M–A constituents larger than 1 m2 were observed only in ICCGHAZ. We concluded that the deterioration of fatigue crack resistance in ICCGHAZ is caused by crack initiation and propagation by massive M–A constituents that were greater than 1 m2 in size.
5. Conclusions In this study, the fatigue properties in various weld HAZs of HSB800 were investigated to consider the relationships between the microstructures of each HAZ and fatigue. Our conclusions were as follows: (1) Crack propagation tests were carried out on base steel, CG, FG, and ICCGHAZ. Our results showed that the fatigue crack growth rate of ICCGHAZ was the fastest, while base steel had the best fatigue crack resistance. (2) The fatigue crack growth rate of CGHAZ was faster than that of FGHAZ although both regions have identical fully martensite microstructures, because FGHAZ has smaller prior austenite grain and martensite packet sizes that act as effective barriers to crack propagation. (3) The fatigue crack growth rate of ICCGHAZ was the fastest among local zones because fatigue cracks were rapidly initiated and propagated through massive M–A constituents that were larger than 1 m2 in size.
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