Effects of Ni and Mn addition on critical crack tip opening displacement (CTOD) of weld-simulated heat-affected zones of three high-strength low-alloy (HSLA) steels

Materials Science & Engineering A 697 (2017) 55–65

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Effects of Ni and Mn addition on critical crack tip opening displacement (CTOD) of weld-simulated heat-affected zones of three high-strength lowalloy (HSLA) steels

MARK

⁎

Seok Gyu Leea, Dong Ho Leea, Seok Su Sohna, , Woo Gyeom Kimb, Kyung-Keun Umb, Ki-Seok Kimc, Sunghak Leea a b c

Center for Advanced Aerospace Materials, Pohang University of Science and Technology, Pohang 790-784, Republic of Korea Steel Products Research Group 1, Technical Research Laboratories, POSCO, Pohang 790-785, Republic of Korea Structural Research Group, Steel Solution Marketing Department, POSCO, Incheon 406-840, Republic of Korea

A R T I C L E I N F O

A B S T R A C T

Keywords: Heat affected zone (HAZ) High-strength low alloy (HSLA) steel Crack tip opening displacement (CTOD) Coarse-grained HAZ Inter-critically heated HAZ

In order to understand and improve fracture toughness of heat affected zones (HAZs) of high-strength low alloy (HSLA) steels, complex microstructures including quasi-polygonal ferrite (QPF), acicular ferrite (AF), granular bainite (GB), bainitic ferrite (BF), and martensite-austenite (MA) constituent should be identified, quantified, and then correlated with critical crack tip opening displacement (CTOD). In this study, microscopic analysis methods were achieved for identification and quantitation of microstructures in the HAZs of three HSLA steels. The coarse-grained HAZ (CGHAZ) consisted of AF, GB, and BF together with a small amount of MA, while the inter-critically heated HAZ (ICHAZ) consisted of QPF, GB, and MA. In the CGHAZ, Ni promoted the formation of AF, while it prevented the formation of GB, and the addition of Ni resulted in very high critical CTOD. In the CGHAZ, both Ni and Mn promoted the formation of AF and prevented the formation of GB, while Ni was more effective than Mn. Thus, the addition of Ni resulted in very high critical CTOD. In the ICHAZ, both Ni and Mn promoted the formation MA. However, in the high-Ni-containing steel, a number of MAs were densified along Ni-segregated bands, and thus readily provided void initiation sites. This played an important role in reducing the mean free path for coalescence of voids and crack propagation, which easily led to the serious deterioration of critical CTOD.

1. Introduction Very low fracture toughness values have been revealed in heat affected zones (HAZs) of high-strength low alloy (HSLA) steels used for offshore platform constructions [1,2]. Metallographic analyses of HSLA steel welds have noticed different regions in HAZ microstructures [3–5]. In single-pass welds, for example, there are four characteristic regions in the HAZ determined by the peak temperature to which the region was exposed during the weld thermal cycles, i.e., coarse-grained HAZ, fine-grained HAZ, inter-critically heated HAZ, and sub-critically heated HAZ. In multi-pass welds, these regions are affected by multiple thermal cycles, thereby resulting in more inhomogeneous and complicated microstructures. It is generally accepted that the coarse-grained HAZ (CGHAZ) has the lowest fracture toughness because of its undesirable microstructural features such as large prior austenite grain size, low-temperature transformed bainitic microstructures, and martensite-austenite (MA) constituents [6,7]. The inter-critically heated

⁎

HAZ (ICHAZ) is also a low-toughness region because a number of hard MAs are formed during the cooling from the (ferrite+austenite) region. Carbon equivalent (Ceq) and alloying elements such as Ni and Mn seriously affect the fracture toughness of the CGHAZ and ICHAZ [8–10]. Mn and Ni are often added to improve strength and fracture toughness of HSLA steels, particularly in thermo-mechanically controlled processed (TMCP) steels whose main microstructures are bainitic ones such as quasi-polygonal ferrite (QPF), acicular ferrite (AF), granular bainite (GB), and bainitic ferrite (BF), together with MA. Under welding conditions, mechanical and fracture properties are significantly varied with more complicated HAZ microstructures. Here in HSLA steel welds, Mn and Ni stabilize the austenite at high temperatures [11–14], and raise the volume fraction of fine AF which is known to have a good combination of strength and toughness [15–18]. However, since the addition of Mn or Ni may promote the formation of crack-susceptible MA, the optimization of alloying elements is required to understand details of fracture toughness degradation. In order to systematically

Corresponding author. E-mail address: [email protected] (S.S. Sohn).

http://dx.doi.org/10.1016/j.msea.2017.04.115 Received 22 December 2016; Received in revised form 23 April 2017; Accepted 25 April 2017 Available online 04 May 2017 0921-5093/ © 2017 Elsevier B.V. All rights reserved.

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understand and improve fracture toughness of these HAZs, thus, various complex microstructures should be verified and then correlated with macroscopic fracture toughness [19,20]. A simple way widely used to investigate complex microstructures is a phenomenal observation by optical or scanning electron microscope, the results of which might not be credible because identification and classification methods of complex HAZ microstructures have not been sufficiently developed yet [5,21–24]. Recently, the microscopic understanding is extended as electron microscopes are more widely utilized [25–29]. In spite of these efforts, many difficulties still remain to be addressed to quantitatively define microstructures and to systematically analyze them. Particularly in the HSLA steel HAZs where very inhomogeneous and complicated microstructures are mixed together, microscopic features are occasionally varied with observation locations and microscopic classification methods even in the same HAZ, but only limited information on microscopic identification and quantitation is available. In the present study, therefore, reliable microscopic analysis methods were achieved for identification and quantitation of microstructures existed in the HAZs of three HSLA steels. Effects of various microstructural features on critical crack tip opening displacement (CTOD), which has been generally used for evaluation of fracture toughness in offshore-application HSLA steel HAZs, were investigated by conducting weld thermal cycle simulation tests and by evaluating CTOD values of thermally simulated HAZ specimens. These simulated HAZ microstructures were so complicated that the electron back-scatter diffraction (EBSD) analysis coupled with optical and SEM micrographs were required for identification of each microstructure. The resultant quantitative analysis data of simulated HAZ microstructures were then correlated with the measured CTOD test results, and effects of Mn or Ni addition on HAZ microstructures and CTOD values were verified.

Table 2 Room-temperature tensile test results of the three HSLA steels. Steel

Yield Strength (MPa)

Tensile Strength (MPa)

Elongation (%)

T HM HN

524 532 563

609 617 652

24 23 22

Fig. 1. Schematic diagrams showing thermal simulation cycles of the coarse-grained HAZ (CGHAZ) and inter-critically heated HAZ (ICHAZ).

characterized by the peak temperature and the cooling time from the peak temperature as shown in Fig. 1. After reaching the peak temperature of 1350 °C or 720 °C, the specimens were cooled to 500 °C at a rate of 15 °C/s, and were cooled further down to 200 °C at a rate of 7.5 °C/s. The weld-simulated specimens were polished and electro-etched in a solution of 8% perchloric acid and 92% acetic acid, and microstructures of longitudinal-short transverse (L-S) plane were observed by an optical microscope and a scanning electron microscope (SEM, model; S4300SE, Hitachi, Tokyo, Japan). Electron back-scatter diffraction (EBSD) analysis (step size; 0.5 or 0.1 µm) was conducted by a field emission scanning electron microscope (FE-SEM, model; Quanta 3D FEG, FEI Company, USA), and the data were interpreted by an orientation imaging microscopy (OIM) analysis software provided by TexSEM Laboratories, Inc. Electron probe micro analysis (EPMA) measurements employing wavelength-dispersive spectrometry (WDS) were also performed by an EPMA microprobe (model; JXA 8530 F microprobe, JEOL, Japan) at an electron beam voltage of 15 keV. The critical CTOD was evaluated in accordance with the BS 7448 standard test method [30]. Three-point bending tests were conducted on pre-cracked rectangular bar specimens (size: 10×10×60 mm, orientation: transverse-longitudinal (T-L)), whose shape and dimensions are shown in Fig. 2. The initial machined-notch depth was 3 mm. The fatigued pre-crack was introduced into the specimen within a stress ratio (R) range of 0.01 < R < 0.1, and the pre-crack length was controlled in a range of 0.45 < a0/W < 0.70, where a0 and W were total crack length (mm) and width (mm) of the specimen, respectively. The CTOD tests were conducted at a cross-head speed of 0.02 mm/s by a universal testing machine (model: 810, MTS Systems Corporation, Minneapolis, USA) with a 100 kN capacity. The CTOD specimen was

2. Experimental The HSLA steels used in this study are 80-mm-thick commercial S500~S550 grade (yield strength; 500–550 MPa) steels produced by TMCP, whose chemical compositions are shown in Table 1. The ‘T’ steel has a reference composition, and a small amount of Mn or Ni was added to the T steel to fabricate the ‘HM’ or ‘HN’ steel, respectively, in order to investigate effects of Mn or Ni addition. The TMCP treatment was similar to commercial one consisted of reheating at 1100–1200 °C and followed by controlled rolling and accelerated cooling [19]. An overall grain refinement effect was expected by rolling with a high-rolling reduction ratio of 75% in the non-recrystallized region of austenite after the austenitization. Rolling was finished at 760–800 °C in the austenite region above Ar3, and the final thickness was 80 mm. After the finish rolling, the steels were rapidly cooled from 740 °C to finish cooling temperature below 300 °C at a cooling rate of 4–10 °C/s, cooled in the air to room temperature, and then tempered at 500–600 °C for 160 min in order to improve a toughness. Basic room-temperature tensile properties are listed in Table 2. Weld thermal cycle simulation tests were conducted on rectangular bar specimens (size: 10×10×60 mm, orientation: transverse) by a metal thermal cycle simulator (Phase Transformation Simulator, POSCO, Korea). The specimen temperature was controlled by a K-type thermocouple, which was wire-percussion-welded at the midsection of the specimen. The thermal cycle of the weld simulation in the coarsegrained HAZ (CGHAZ) and inter-critically heated HAZ (ICHAZ) was Table 1 Chemical compositions of the three commercial S500~S550 grade HSLA steels. (wt.%). Steel

C

Mn

Ni

Si

Cr+Al+Ca

B+N

P+S

Mo+Cu

Ti+Nb+V

T HM HN

< 0.07 < 0.07 < 0.07

0.8~1.6 1.2~2.0 0.8~1.6

1.0~1.8 1.0~1.8 1.4~2.2

0.1 0.1 0.1

< 0.04 < 0.03 < 0.04

< 0.004 < 0.004 < 0.003

< 0.006 < 0.005 < 0.006

< 0.7 < 0.7 < 0.7

< 0.03 < 0.03 < 0.03

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laths. It shows a good combination of strength and toughness because of its high interior dislocation density and fine grain. GB has relatively large packets with island-type MA constituents. BF is well formed at faster cooling rates and lower finish-cooling temperatures than AF or GB. Since secondary phases are finely distributed along lath boundaries, BF shows high strength but low toughness. MA is a secondary phase formed at fast cooling rates. According to these microstructural features, QPF, AF, GB, BF, and MA are differentiated in this study by using optical and SEM micrographs coupled with EBSD analyses. Fig. 2. Shape and dimensions of the CTOD test specimen. (unit: mm) The critical CTOD was evaluated in accordance with the BS 7448 standard test method [30].

3.2. Simulated CGHAZ microstructure placed in a low-temperature chamber, in which the test temperature (−20 °C) was controlled by liquid nitrogen and ethyl alcohol. Critical CTOD values were determined by using the following equation in accordance with the CTOD test standard of BS 7448 [30]:

Fig. 3a-c shows optical micrographs of the simulated CGHAZ of the three steels. These CGHAZs are generally composed of AF, BF, and GB as marked by arrows. They also have coarse packets whose sizes are larger than 50 µm, and prior austenite grain boundaries are observable. Since the differentiation between MA and other microstructures was difficult in optical micrographs, the CGHAZs were etched in a LePera solution [36], and LePera-etched micrographs are shown in Fig. 3d-f. MA and other microstructures (AF, GB, and BF) are colored in brightwhite and brown, respectively. The volume fractions of MA are similar in the range of 2.7–3.5% in the three steels. In order to elucidate and quantify the complex microstructures, more detailed analyses including SEM and EBSD are essentially needed. Fig. 4a shows an EBSD inverse pole figure (IPF) map of the CGHAZ of the T steel. Here, AF, BF, and GB can be identified by fine irregularshaped, parallel lath-shaped, and coarse equi-axed microstructures, respectively. Fig. 4b-d shows misorientation profiles of black arrows in black-dashed-circle-marked AF, BF, and GB areas in Fig. 4a. The spacing between AF boundaries is small, and most of AF boundaries are high-angle boundaries having misorientations of 50–60° (Fig. 4b). BF packets defined by boundaries having misorientations of 15° are coarse, and are composed of parallel substructures of low-angle (1–10°) boundaries (Fig. 4c). Inside GB packets, substructures are also developed (Fig. 4d). These GBs are coarse (about 25 µm) and irregularshaped. The same areas of black-dashed-circle-marked AF, BF, and GB areas in Fig. 4a are observed by an SEM, as shown in Fig. 4e-g. The SEM micrographs are well matched with the microstructural definition by the EBSD. These microstructural definition methods indicate that the

CTOD,δ = (K2 (1 − n2))/(2σys E ) + (0.4(w − a 0)Vpl)/(0.4W + 0.6a 0 + z ) (1) where K (MPa m1/2) is stress intensity factor for the critical load, E (MPa) is Young's modulus, ν is Poisson's ratio, and σys (MPa) is yield strength of the material. W, a0, and z are specimen dimensions (mm), and Vpl (mm) is plastic component of crack mouth opening displacement. In order to investigate fracture modes and stretched zones, fracture surfaces were observed by SEM. 3. Results 3.1. Microstructural classification All the steels and their simulated HAZ specimens basically consist of quasi-polygonal ferrite (QPF), acicular ferrite (AF), granular bainite (GB), bainitic ferrite (BF), and martensite-austenite constituent (MA). These microstructures are classified by their morphologies and characteristics defined from many researches on previously developed HALS steels [23,31–35]. For example, QPF is transformed at lower temperatures and higher cooling rates than polygonal ferrite, and has irregular grain boundaries. AF is a very fine, needle-shaped ferrite, and can be grouped into packets according to orientations between neighboring

Fig. 3. Optical micrographs of the simulated CGHAZ of the (a, d) T, (b,e) HM, and (c, f) HN steels. The CGHAZs are generally composed of acicular ferrite (AF), bainitic ferrite (BF), and granular bainite (GB) as marked by arrows. (d-f) optical micrographs of the CGHAZs etched in a LePera solution [36]. Martensite-austenite constituent (MA) and other microstructures (AF, GB, and BF) are colored in bright-white and brown, respectively.

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Fig. 4. (a) EBSD inverse pole figure (IPF) maps, and (b-d) EBSD misorientation profiles, and (e-g) SEM micrographs of the CGHAZ of the T steel. AF, BF, and GB can be identified by fine irregular-shaped, parallel lath-shaped, and coarse equi-axed microstructures, respectively.

200 °C at a rate of 7.5 °C/s, a small amount of MA is formed from the retained austenite. Both Mn and Ni play a role in promoting the formation of AF and in preventing the formation of GB, while Ni promotes the formation of BF, but Ni is more effective than Mn [43,44]. Fig. 5a-f shows EBSD IPF and grain boundary maps, and SEM micrographs of the CGHAZ of the HM and HN steels, from which microstructures such as AF, BF, and GB are defined by the aforementioned microstructure analysis methods and their volume fractions are summarized in Table 4. These CGHAZs consist of AF, GB, and BF together with a small amount of MA, like in the CGHAZ of the T steel, and the volume fractions of AF, GB, and BF are highest in the HN, T, and HN steels, respectively. The volume fraction of AF tends to be higher in the HM and HN steels than in the T steel, whereas the volume fraction of GB is lower. Particularly in the CGHAZ of the HN steel, AF is very populated (Fig. 5d-f). These microstructural analysis results indicate that the volume fraction of AF increases as the Mn and Ni addition increases, while that of GB decreases, and that the Ni addition is more effective than the Mn addition (Table 4).

combination of EBSD maps and misorientation profiles, optical micrographs, and SEM micrographs is very useful for the detailed classification of various microstructures such as AF, BF, and GB existed in the CGHAZ in consideration of their morphologies and substructures. In order to estimate HAZ microstructures of the present HSLA steels, Ar3, Ar1, bainite start temperature (Bs), and martensite start temperature (Ms) were calculated by using the following equations [37–42], and the results are summarized in Table 3. Ar3 (°C) = 910 – 230C – 21Mn – 15Ni + 32Mo + 45Si

(2)

Ar1 (°C) = 723 – 10.7Mn – 16.9Ni + 29.1Si + 16.9Cr

(3)

Bs (°C) = 830 – 270C – 90Mn – 37Ni – 70Cr – 88Mo

(4)

Ms (°C) = 539 – 423C – 30.4Mn – 17.7Ni – 12.1Cr – 7.5Si

(5)

The Ar3, Ar1, Bs, and Ms temperatures are highest in the T steel, and tend to decrease as Mn or Ni is added. The Bs and Ms temperatures are lowest in the HM steel. The Ar1, Bs, and Ms temperature ranges are also shown in Fig. 1. In the simulated CGHAZ, the peak temperature of 1350 °C lies in the austenite single phase region (Fig. 1). When cooling from 1350 °C to 500 °C at a fast rate of 15 °C/s, low-temperature transformed microstructures such as AF, GB, and BF except martensite are formed because 500 °C lies between Bs and Ms temperatures, while polygonal ferrites are hardly formed as they need the sufficient diffusion. According to the rapid cooling rate, these microstructures are mostly nucleated and grown at grain boundaries or inside grains rather than along grain boundaries, which results in the clear observation of prior austenite grain boundaries as shown in Fig. 3a-c. When further cooling down to

3.3. Simulated ICHAZ microstructure Fig. 6a-c shows optical micrographs of the ICHAZ of the three steels. In the ICHAZs, fine grains are mostly aligned along the rolling direction, unlike in the large-packet-shaped CGHAZs, and GB microstructures are populated, together with some QPF microstructures whose substructures are not sufficiently developed. Island-shaped MAs are present, but are not well defined in the optical micrographs. Fig. 6d-f shows optical micrographs of the ICHAZs etched by a LePera solution [36]. The volume fractions of MA are much higher in the ICHAZs than in the CGHAZs, and tend to increase in the order of the T, HM, and HN steels. MAs tend to be densified in a band shape along the rolling direction. In the T steel, most of MAs are discontinuously distributed in the bands, and only a few fine MAs are found between the bands (Fig. 6d). In the HN steel whose MA volume fraction is highest, coarse MAs are continuously formed in the bands, while a considerable number of MAs are also observed between the bands (Fig. 6f). In order to investigate causes of MA-densified band formation, chemical composition maps of C, Ni, and Mn and composition profiles of Ni and Mn as well as SEM images were obtained by the EPMA analysis, as shown in Fig. 6a-c. Island-shaped particles in the SEM images are identified to be

Table 3 Ar3, Ar1, bainite start temperature (Bs), and martensite start temperature (Ms) of the three HSLA steels. (unit: °C). Steel

Ar3 temperature

Ar1 temperature

Bs temperature

Ms temperature

T HM HN

849 843 843

682 678 676

607 573 592

442 431 435

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Fig. 5. EBSD inverse pole figure (IPF) and grain boundary maps, and SEM micrographs of the CGHAZ of the (a-c) HM and (d-f) HN steels. These CGHAZs are composed of AF, BF, and GB together with a small amount of MA.

MA ones because carbon is concentrated at them. MAs are relatively homogeneously distributed in grain boundary areas in the T steel, and the segregation of Ni and Mn is not clearly visible (Fig. 6a). In the HM steel having high Mn content, Ni and Mn contents are higher by 0.1–0.2 wt% in coarse-MA-rich band areas than in MA-depleted matrix areas (Fig. 6b). In the HN steel having high Ni content, however, Ni and Mn are distinctly segregated in coarse-MA-rich band areas (Fig. 6c). Ni and Mn contents are higher by 0.58 wt% and 0.77 wt%, respectively, in coarse-MA-rich band areas than in MA-depleted matrix areas. This indicates that the high Ni content causes Ni- and Mn-segregated bands in the HN steel, while the segregation tendency of Ni and Mn is small in the HM steel (Fig. 7). Fig. 8a shows an EBSD IPF map of the ICHAZ of the T steel.

Table 4 Volume fractions of quasi polygonal ferrite (QPF), acicular ferrite (AF), granular bainite (GB), bainitic ferrite (BF), and martensite-austenite constituent (MA) in the simulated HSLA steel HAZs. Steel

T HM HN T HM HN

HAZ

CGHAZ

ICHAZ

Volume Fraction (%) QPF

AF

GB

BF

MA

– – – 20.7 ± 2.8 19.5 ± 2.6 27.3 ± 2.5

52.0 ± 3.4 58.0 ± 2.8 62.2 ± 4.1 – – –

16.2 ± 0.8 16.3 ± 1.7 23.0 ± 1.7 77.6 ± 1.5 75.2 ± 2.3 62.4 ± 1.1

29.1 ± 3.2 22.2 ± 2.0 11.9 ± 1.5 – – –

2.7 ± 0.2 3.5 ± 0.2 2.9 ± 0.3 1.7 ± 0.4 5.3 ± 0.4 10.3 ± 0.6

Fig. 6. Optical micrographs of the simulated ICHAZ of the (a, d) T, (b,e) HM, and (c, f) HN steels. The ICHAZs are generally composed of granular bainite (GB), quasi-polygonal ferrite (QPF), and MA as marked by arrows. (d-f) are optical micrographs of the ICHAZs etched in a LePera solution [36]. MA and other microstructures (GB and QPF) are colored in bright-white and brown, respectively.

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Fig. 7. SEM micrographs, EPMA composition maps of C, Ni, and Mn, and composition profiles of Ni and Mn of the (a) T, (b) HM, and (c) HN steels. Ni is segregated along MA-densified bands in the HN steel.

Fig. 8. EBSD inverse pole figure (IPF) and grain orientation spread (GOS) maps, and SEM micrographs of the ICHAZ of the (a-c) T, (d-f) HM steels, and (g-i) HN steels. These ICHAZs are composed of GB, QPF, and MA.

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and their volume fractions are summarized in Table 4. These ICHAZs are also composed of QPF, GB, and MA, like in the ICHAZ of the T steel. The volume fractions of QPF and GB are similar in the T and HM steels within error ranges, while that of MA is somewhat higher in the HM steel. In the HN steel, the volume fraction of QPF increases while that of GB decreases, in comparison with the T or HM steel, which indicates the increasing QPF and decreasing GB effects of Ni. It is noted that MAs are most populated in the HN steel. In the simulated ICHAZ, the peak temperature of 720 °C lies in the (ferrite+austenite) two phase region because it is higher than the Ar1. In this two phase region, fine austenites are formed near carbides or grain boundaries whose C contents are high. As-received (non-HAZsimulated) microstructures such as AF and BF are recovered, and internal substructures are changed by the reduction in overall dislocation density. When cooling from 720 °C to 500 °C at a fast rate of 15 °C/ s, low-temperature transformed microstructures are hardly formed because fine austenite grains can keep the C content. During the further cooling to 200 °C, a number of MAs are formed as the temperature decreases down below the Ms. As-received fine AF having high-angle boundaries is changed to a kind of recovered ferrite, i.e., QPF. Asreceived BF having coarse packet boundaries is readily changed into GB containing MAs formed from carbides located at BF while internal substructures are released. Mn and Ni, which are typical austenite stabilizers, decrease the Ar1, and thus the addition of Mn or Ni raises the volume fraction of austenite in the (ferrite+austenite) region. Since this austenite is finally transformed to MA, MA is more populated in the HM and HN steels than in the T steel. Particularly in the HN steel, the volume fraction of MA reaches about 10%. This can be explained by effects of Ni-segregated bands of Fig. 7c. Transformation temperatures of Ar3 and Ar1 of the HN steel are similar to those of the HM steel because contents of Mn and Ni are higher by about 0.4 wt% than those in the T steel (Table 1). When the Ni content is high, like in the HN steel, Ni becomes to be segregated along the rolling direction, thereby forming Ni-segregated bands whose Ni content is high (about 1.8 wt%). In the inter-critical temperature region, thus, the reverse transformation to austenite and its grain growth occur preferentially, which results in the large size and high

Table 5 Critical CTOD values measured at −20 °C and stretched zone widths of the three HSLA steels. (unit: mm). Steel

HAZ

Maximum CTOD

Minimum CTOD

Average CTOD

Stretched Zone Width (μm)

T HM HN T HM HN

CGHAZ

0.28 0.32 0.41 0.37 0.25 0.25

0.08 0.04 0.22 0.24 0.24 0.16

0.16 ± 0.06 0.20 ± 0.10 0.30 ± 0.08 0.31 ± 0.05 0.24 ± 0.01 0.22 ± 0.04

11.42 ± 4.47 18.96 ± 4.48 53.15 ± 20.70 38.98 ± 8.72 33.33 ± 6.37 21.19 ± 7.16

ICHAZ

According to morphological categories and existence of substructures, overall ICHAZ microstructure consists of GB, together with some QPF microstructures. In order to classify and quantify microstructural differences by the extent of substructure development, detailed substructures were analyzed by a grain orientation spread (GOS) map, as shown in Fig. 8b. The GOS color map provides the degree of average misorientation between certain datum points inside a grain. In the present simulated ICHAZ, the angle of 15° is determined to be a misorientation for differentiating grain boundaries, and grains having misorientations of 5° or smaller are classified into QPFs, as colored by yellow areas in Fig. 8b. By using this GOS technique, QPFs can be differentiated from GBs because substructures are hardly developed in QPFs, although both QPFs and GBs have irregular-shaped high-angle boundaries. According to the GOS analysis data, the volume fraction of QPF is about 21% in the ICHAZ of the T steel. Fig. 8c shows SEM micrographs of GB and QPF regions marked in Fig. 8b. Simple morphological differentiation between QPF and GB is not easy in these micrographs. GB can contain island-type MAs, but relatively coarse MAs can be occasionally formed in the ICHAZ, which makes the easy differentiation between QPF and GB. These results imply that the GOS technique or substructure defining method would be a more quantitatively clarified accurate method. Fig. 8d-i shows EBSD IPF and GOS grain boundary maps, and SEM micrographs of the ICHAZ of the HM and HN steels, from which microstructures are defined by the analysis methods shown in Fig. 8a-c

Fig. 9. SEM fractographs of the CTOD specimens of the (a-c) CGHAZ and (d-f) ICHAZ. Stretched zones are observed in front of the fatigued pre-crack tip. After the crack-tip stretching, the crack propagates mostly in a cleavage fracture mode.

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the fractured CTOD specimens of the CGHAZ of the HN steel were examined by EBSD IPF, grain boundary, and kernel average misorientation (KAM) maps, and the results are shown in Fig. 11a,b. Here, the KAM was calculated up to the first neighbor shell with a maximum misorientation angle of 5° [45]. The KAM maps serve as a measure of the deformation-induced local orientation gradients inside grains, and KAM values are high in grains whose dislocation densities are high. Since the CGHAZ specimens have large deviations in critical CTOD values, two specimens having minimum and maximum critical CTOD values (0.22 vs. 0.41 mm) were selected from the fractured CTOD specimens, and cross-sectional areas were examined. In the minimumCTOD-valued (0.22 mm) specimen, coarse packets of GB exist in front of the fatigued pre-crack tip, and the crack propagation path is relatively linear because the crack initiated at the pre-crack tip propagates rapidly across GBs (Fig. 11a). In this case, the stretched zone located in front of the pre-crack tip is very small, and the plastic deformation hardly occurs. KAM values of the crack tip region are relatively low (average; 0.95 ± 0.59°), and confirm the small plastic deformation and the rapid and linear crack propagation with a little resistance to fracture. In the maximum-CTOD-valued (0.41 mm) specimen, on the other hand, fine AF grains whose boundaries are highangled ones are populated, as shown in Fig. 11b. Here, the crack propagation path is tortuous in a zig-zag pattern as the crack frequently deviates at fine and highly deformable AF grains. The stretched zone is quite large, and KAM values of the crack tip region are high (average; 2.04 ± 0.97°) because the large plastic deformation occurs at the crack tip region. According to the aforementioned results in the CGHAZ, the decrease in GB fraction and increase in AF fraction improve critical CTOD values. Thus, the HN steel shows the best critical CTOD value because the GB fraction is reduced by the addition of Ni. It is interesting to note that these EBSD results of the two specimens having minimum and maximum critical CTOD values indicate that critical CTOD values can be largely deviated, depending on what microstructure exists mostly in the crack tip region. In order to accurately measure critical CTOD values, thus, the distribution of microstructures near the crack tip region should be carefully examined, and critical CTOD values need to be averaged after testing of many specimens. In the present EBSD analysis data (Fig. 11), two major microstructures, i.e., GB and AF, are considered for estimating CTOD values because the volume fractions of MA are similar in the three steels. In the ICHAZ, however, brittle MAs and their size, volume fraction, and distribution should also be correlated with CTOD values because they seriously deteriorate the resistance to fracture [2,46–48]. It is noted that the HN steel shows the lowest CTOD value as the MA fraction reaches about 10%, although it has the most favorable structure for plastic deformation because of the lowest GB fraction and the highest QPF fraction. This indicates that only a small increase in MA fraction can greatly influence critical CTOD values. When hard and brittle MA and relatively soft QPF and GB are mixed together, like in the ICHAZ of the HN steel, the negative effect of MA on critical CTOD readily

volume fraction of austenite in the HN steel. This also indicates that Ni activates the formation of austenite more readily than Mn according to the formation of Ni-segregated band. 3.4. Critical CTOD Critical CTOD data are summarized in Table 5. Deviations of critical CTOD values are large in the CGHAZs, while they are relatively small in the ICHAZs. The average critical CTOD values are higher in the ICHAZ than in the CGHAZ for the T and HM steels. The critical CTOD in the CGHAZ increases in the order of the T, HM, and HN steels, but this trend is reversed in the ICHAZ. Fig. 9a-f shows SEM fractographs of the CTOD specimens. In all the specimens, stretched zones are observed in front of the fatigued precrack tip. After the crack-tip stretching, the crack propagates mostly in a cleavage fracture mode. The width of the stretched zone was measured, and the results are shown in Table 5. The stretched zone width ranges in 11~55 µm, and shows the same trends of the critical CTOD data. 4. Discussion HSLA steels produced by TMCP generally contain low-temperature transformed microstructures such as AF and BF having a good combination of strength and toughness. In the case of welding, however, their toughness is seriously deteriorated by forming lowtoughness microstructures at HAZs. The simulated HAZ microstructures mainly consist of AF, GB, BF, QPF, and MA, as mentioned in the Sections 3.3 and 3.4. Among them, critical microstructural factors deteriorating critical CTOD in each HAZ should be examined in detail to improve the fracture toughness of the HAZs by appropriately controlling HAZ microstructures. The addition of Mn or Ni influences microstructures of the simulated HAZs and consequent critical CTOD values. Fig. 10a,b shows the relation between microstructural fractions and critical CTOD values in the CGHAZ and ICHAZ for the three steels. Critical CTOD values of the CGHAZ increase as the volume fraction of AF increases or the volume fraction of GB decreases (Fig. 10a). In the T steel, the critical CTOD is lowest because of the highest volume fraction of GB. In order to improve the critical CTOD in the CGHAZ, the activation of AF formation as well as the prevention of GB formation are needed, and can be effectively achieved by adding Ni. In the ICHAZ, on the other hand, the Ni-added HN steel shows the lowest critical CTOD (Fig. 10b). This is because the volume fraction of MA is highest (about 10%). The T steel has the highest critical CTOD because of the lowest volume fraction of MA (about 2%), although the volume fraction of GB is high. The increased volume fraction of QPF whose fracture resistance and plastic deformability are excellent also favorably influences the critical CTOD. In order to confirm effects of GB and AF on critical CTOD in the CGHAZ, the cross-sectional areas in front of the fatigued pre-crack tip of

Fig. 10. Relation between microstructural fractions and critical CTOD values in the (a) CGHAZ and (b) ICHAZ.

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Fig. 11. EBSD IPF, grain boundary, and kernel average misorientation (KAM) maps of the cross-sectional areas of the fractured CTOD specimens having (a) minimum and (b) maximum critical CTOD values (0.22 vs. 0.41 mm) for the CGHAZ of the HN steel. The fatigued pre-crack tip is marked by arrows.

Fig. 12. SEM micrograph of the cross-sectional area in front of the fatigued pre-crack tip of the fractured CTOD specimens of the ICHAZ of the HN steel. A considerable number of voids or microcracks are found mainly at MA/matrix interfaces beneath the fracture surface, while some voids or microcracks are formed by the cracking of MAs themselves, which indicates that the major fracture mechanism is a void or microcrack initiation at MA/matrix interfaces. The number of voided MAs is graphed as a function of MA size range in the right side of the SEM micrograph.

can actively work as void initiation sites when its size is not too small or large, e.g., intermediate range of 1.5–3 µm. Another noteworthy characteristic of the ICHAZ of the HN steel is a populated distribution of MA along Ni-segregated bands (Fig. 7c). Thus, a number of MAs densified along the bands readily provide void initiation sites. This plays an important role in reducing the mean free path for coalescence of voids and crack propagation, which easily leads to the serious deterioration of critical CTOD, as shown in Table 5. Since MAs are densified along these bands, they readily provide void initiation sites, and play a role in reducing the mean free path for coalescence of voids and crack propagation, thereby leading to the serious deterioration of critical CTOD. When considering the addition of Mn or Ni with critical CTOD values of the HAZs simultaneously, critical CTOD values are appropriate at about 0.2 mm in both CGHAZ and ICHAZ of the HM steel. In the HN steel, Ni promotes the formation of AF in the CGHAZ, while it

overrides the positive effect of QPF and GB. When MAs are populated, it can be simply expected that the CTOD values are deteriorated by the increase number of MA, but there exists an effective MA size range for the void or microcrack initiation [2,46–49]. Fig. 12 shows an SEM micrograph of the cross-sectional area in front of the fatigued pre-crack tip of the fractured CTOD specimens of the ICHAZ of the HN steel. A considerable number of voids or microcracks are found mainly at MA/matrix interfaces beneath the fracture surface, while some voids or microcracks are formed by the cracking of MAs themselves. This indicates that the major fracture mechanism is a void or microcrack initiation at MA/matrix interfaces. The number of voided MAs was measured, and the results are shown in the right side of Fig. 12 as a function of MA size range. Average size of voided MA is 2 ± 1.1 µm, and the void occurrence frequency is highest in the MA size range of 1.5–3 µm, although a few voids are formed at very fine (0.5 µm or finer) or coarse (3.5 µm or coarser) MAs. Thus, MA 63

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provided void initiation sites. This played an important role in reducing the mean free path for coalescence of voids and crack propagation, which easily led to the serious deterioration of critical CTOD.

prevents the formation of GB, and thus the addition of Ni results in very high critical CTOD value (0.30 mm). However, it reduces the critical CTOD to 0.22 mm in the ICHAZ because Ni promotes the formation of MA. With respect to critical CTOD in both HAZs, therefore, the HN steel are better overall than the HM steel, which indicates the beneficial effect of Ni addition on improving critical CTOD, but the careful control of Ni addition is needed for restraining the populated formation of MA. The present study on microstructural definition and quantitation of microstructures existed in the HAZs of offshore-application HSLA steels would prove a good way to investigate effects of Mn or Ni addition on critical CTOD. It is also useful to understand fracture characteristics in relation with complicated microstructures of the simulated CGHAZ and ICHAZ having low fracture toughness. Critical CTOD data can be plausibly interpreted by volume fractions of various complicated HAZ microstructures which are successfully defined by the detailed EBSD microscopic analyses. In fact, the detailed microstructural analyses of the HSLA steel HAZs, whose microstructural definition and quantification are quite difficult because of their microstructural complexity, have been hardly performed. The present EBSD microscopic definition and quantification methods are outstanding ones for the detailed analyses of HSLA steel HAZ microstructures. They can also lead to the reliable evaluation of critical CTOD, and provide an important idea for the alloy designs of HSLA steels demanded for offshore platform constructions.

Acknowledgements This work was supported by POSCO [grant number; 2016Y001], and Korea Agency for Infrastructure Technology Advancement [grant number; 17IFIP-B067108-05]. References [1] J. Liao, K. Ikeuchi, F. Matsuda, Effect of cooling time on HAZ toughness and microstructure: simulated HAZ toughness of low-alloy SQV-2A pressure vessel steel (1st report), Weld. Int. 10 (1996) 552–558. [2] C.L. Davis, J.E. King, Cleavage initiation in the intercritically reheated coarsegrained heat affected zone: Part I. Fractographic evidence, Metall. Mater. Trans. A 25A (1994) 563–573. [3] T. Haze, S. Aihara, Influence of toughness and size of local brittle zone on HAZ toughness of HSLA steels, in: Proceedings of the 7th International Conference on Offshore Mechanics and Arctic Engineering, Houston, TX, ASME, Golden, CO, 1988, pp. 515–523. [4] K. Uchino, Y. Ohno, Proceedings of the 7th International Conference on Offshore Mechanics and Arctic Engineering, Houston, TX, ASME, Golden, CO, 1988, pp. 159–165. [5] B.C. Kim, S. Lee, N.J. Kim, D.Y. Lee, Microstructure and local brittle zone phenomena in high-strength low-alloy welds, Metall. Trans. A 22A (1991) 139–149. [6] K. Easterling, Introduction to the Physical Metallurgy of Welding, ButterworthHeinemann, Oxford, 1992. [7] Y. Li, D.N. Crowther, M.J.W. Green, P.S. Mitchell, T.N. Baker, The effect of vanadium and niobium on the properties and microstructure of the intercritically reheated coarse grained heat affected zone in low carbon microalloyed steels, ISIJ Int. 41 (2001) 46–55. [8] E. Keehan, H.O. Andren, L. Karlsson, M. Murugananth, H.K.D.H. Bhadeshia, Microstructural and mechanical effects of nickel and manganese on high strength steel weld metals, Trends Weld. Res. (2002) 695–700. [9] M. Lord, Design and modelling of ultra-high strength steel weld deposits (Ph. D. thesis), 1999. [10] B.Y. Kang, H.J. Kim, S.K. Hwang, Effect of Mn and Ni on the variation of the microstructure and mechanical properties of low-carbon weld metals, ISIJ Int. 40 (2000) 1237. [11] M. Shome, O.P. Gupta, O.N. Mohanty, Effect of simulated thermal cycles on the microstructure of the heat-affected zone in HSLA-80 and HSLA-100 steel plates, Metall. Trans. A 35A (2004) 985–996. [12] Q.Y. Long, D. Tseng, K. Tangri, Retained austenite in intercritically annealed HSLA steel, Metallography 20 (1987) 61–73. [13] K.E. Easterling, H.M. Miekk-Oja, The martensitic transformation of iron precipitates in a copper matrix, Acta Metall. 11 (1967) 1133–1141. [14] I.W. Chen, Y.H. Chiao, Martensitic nucleation in ZrO2, Acta Metall. 31 (1983) 1627–1638. [15] M.C. Zhao, K. Yang, Y.Y. Shan, The effects of thermos-mechanical control process on microstructures and mechanical properties of a commercial pipeline steel, Mater. Sci. Eng. A 335 (2002) 14–20. [16] J.M. Gregg, H.K.D.H. Bhadeshia, Solid-state nucleation of acicular ferrite on minerals added to molten steel, Acta Mater. 45 (1997) 739–748. [17] Y.E. Smith, A.P. Coldren, R.L. Cryderman, Toward Improved Ductility and Toughness, Climax Molybdenum Company (Japan) Ltd, Tokyo, 1972, pp. 119–142. [18] D. Jeong, W. Jung, Y. Kim, M. Goto, S. Kim, Stress corrosion cracking behavior of X80 steel in artificial seawater under controlled strain rate and applied potentials, Met. Mater. Int. 21 (2015) 785–792. [19] S. Lee, B.C. Kim, D. Kwon, Correlation of microstructure and fracture properties in weld heat-affected zones of thermomechanically controlled processed steels, Metall. Trans. A 25 (1992) 2803–2816. [20] M. Kang, H. Kim, S. Lee, S.Y. Shin, Correlation of microstructure with tensile and crack tip opening displacement properties at low temperatures in API linepipe steels, Met. Mater. Int. 21 (2015) 628–638. [21] B.L. Bramfitt, J.G. Speer, A perspective on the morphology of bainite, Metall. Trans. A 21A (1990) 817–829. [22] H. Ohtani, S. Okaguchi, Y. Fujishiro, Y. Ohmori, Morphology and properties of lowcarbon bainite, Metall. Trans. A 21A (1990) 877–888. [23] G. Krauss, S.W. Thompson, Ferritic microstructures in continuously cooled low-and ultralow-carbon steels, ISIJ Int. 35 (1995) 937–945. [24] L.C. Chang, Microstructures and reaction kinetics of bainite transformation in Sirich steels, Mater. Sci. Eng. A 368 (2004) 175–182. [25] J.S. Kang, S.S. Ahn, C.Y. Yoo, C.G. Park, FIB and TEM studies on the bainitic microstructure in low carbon HSLA steels, Adv. Mater. Res. 26–28 (2007) 73–76. [26] J.S. Kang, J.B. Seol, C.G. Park, Three-dimensional characterization of bainitic microstructures in low-carbon high-strength low-alloy steel studied by electron backscatter diffraction, Mater. Charact. 79 (2013) 110–121. [27] J.H. Sung, Y.K. Kim, J.G. Moon, K.W. Kim, K.B. Kang, K.M. Cho, Effect of plastic deformation on hydrogen induced crack resistance of API X65 linepipe steel, Korean

5. Conclusions In the present study, effects of Ni and Mn addition on weldsimulated HAZ microstructures and CTOD properties of three HSLA steels were analyzed, and their correlation was explained by verifying fracture characteristics. (1) The CGHAZs consisted of AF, GB, and BF together with a small amount of MA. The volume fraction of AF tended to be higher in the HM and HN steels than in the T steel, whereas the volume fraction of GB was lower. Both Mn and Ni played a role in promoting the formation of AF and in preventing the formation of GB, while Ni was more effective than Mn. (2) The ICHAZ microstructures were composed of QPF, GB, and MA. The volume fractions of QPF and GB were highest in the HN and T steels, respectively. The volume fraction of MA was much higher in the ICHAZ than in the CGHAZ, and tended to increase in the order of the T, HM, and HN steels. When the Ni content was high, it became to be segregated along the rolling direction, thereby forming Ni-segregated bands where coarse MAs were continuously formed. (3) When considering the addition of Mn or Ni with critical CTOD values of the HAZs simultaneously, Ni promoted the formation of AF in the CGHAZ, while it prevented the formation of GB, and thus the addition of Ni resulted in very high critical CTOD value (0.30 mm). However, Ni reduced the critical CTOD to 0.22 mm in the ICHAZ because it promoted the formation of MA. Therefore, the careful control of Ni addition was needed for restraining the populated formation of MA. (4) In the CGHAZ where the volume fraction of MA was similar in the three steels, coarse packets of GB in front of the fatigued pre-crack tip induced the small plastic deformation and the rapid and linear crack propagation with a little resistance to fracture. However, fine AF grains whose boundaries were high-angled ones led to the large plastic deformation at the crack tip region. Thus, the crack propagation path was tortuous in a zig-zag pattern as the crack frequently deviated at fine and highly deformable AF grains. (5) In the HN steel, the high Ni content caused Ni- and Mn-segregated bands, while the segregation tendency of Ni and Mn was small in the HM steel. Particularly in the ICHAZ of the HN steel, a number of MAs were densified along Ni-segregated bands, and thus readily 64

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