Effect of Si content on microstructure and cryogenic toughness of heat affected zone of low nickel steel

Materials Science & Engineering A 771 (2020) 138621

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Effect of Si content on microstructure and cryogenic toughness of heat affected zone of low nickel steel Qi-Yuan Chen, Jun Chen, Jia-Kuan Ren, Zeng-Hui Wang, Zhen-Yu Liu * State Key Laboratory of Rolling and Automation, Northeastern University, Shenyang, 110819, China

A R T I C L E I N F O

A B S T R A C T

Keywords: Low nickel steel Si content Heat affected zone Cryogenic toughness

Compared with 9%Ni steel, low nickel steel has great advantages in cost and weldability. However, cryogenic toughness of the heat affected zone (HAZ) of conventional low nickel steel deteriorates, which limits its appli­ cation in the construction of liquefied natural gas tanks. In this work, two low nickel steels with different Si contents were designed, and the relationship between Si content and HAZ cryogenic toughness of these steel samples was investigated. After reducing Si content, the required Charpy impact absorbed energy was obtained in the HAZ at fusion line and 1 mm from fusion line. Microstructural characterization revealed that the reduction in Si content promoted auto-tempering of coarse-grained HAZ (CGHAZ), leading to improved plastic deformation ability of the matrix. In intercritically reheated CGHAZ (ICCGHAZ), martensite-austenite (M-A) constituents were formed, which consisted of twin martensite and austenite. The decrease in Si content reduced the amount and size of M-A constituents and changed their elongated morphology to blocky morphology, which is also believed to be beneficial for ICCGHAZ cryogenic toughness. This work provides a novel way to soften the HAZ micro­ structure for improving its cryogenic toughness.

1. Introduction

The presence of local hard-brittle zones in the HAZ leads to poor toughness, and their embrittlement is generally considered to be related to grain coarsening and formation of martensite-austenite (M-A) con­ stituents [8,9]. Extensive studies have shown that chemical composition of steel has a significant effect on the microstructure and toughness of HAZ. For example, the addition of Nb and V can refine grains, alter the amount and size of M-A constituents in HAZ, and thereby, improve the HAZ toughness [9–12]. Al is also reported to play an important role in modifying the morphologies of M-A constituents [13]. However, these second phase forming elements are detrimental to cryogenic toughness of base metal (BM), and it is not feasible to add them to improve HAZ cryogenic toughness of low nickel steels. It is well known that Si inhibits cementite precipitation and facilitates austenite stabilization and M-A formation [14–19]. Taillard et al. found that the decrease in Si content greatly increased the CGHAZ toughness for microalloyed steels by lowering the amount of M-A constituents [20,21]. Bonnevie et al. studied the effect of Si content on morphologies of M-A constituents in CGHAZ and ICCGHAZ of structural steels [22]. They found that Si increased the proportion of large and massive M-A constituents, which were primarily responsible for the decrease in HAZ toughness. There­ fore, it is feasible to soften local hard-brittle zones by reducing the Si

There is a huge demand for liquefied natural gas (LNG) due to ecological and environmental requirements, and consequently, a large number of giant storage tanks need to be built for storage and trans­ portation of LNG. The most important structural material used for LNG storage tanks is 9%Ni steel, which possesses excellent cryogenic toughness and strength [1–4]. However, because of a relatively high Ni content, 9%Ni steels are susceptible to magnetic fields, leading to poor weldability [5]. Also, the price of Ni is typically high in the international market, causing mounting pressure on the cost of 9%Ni steels. There­ fore, the development of new LNG tank steels, which contain low Ni content and possess cryogenic properties as good as those of 9%Ni steels, is essential for the safety and economic production of the giant LNG tanks. A suitable combination of cryogenic toughness and strength of low nickel steel plate can be achieved via a novel UFC-LT treatment consisting of ultra-fast cooling (UFC), intercritical quenching (i.e., lamellarizing, L), and tempering (T) [6,7]. However, after simply reducing the Ni content, cryogenic toughness of the heat affected zone (HAZ) deteriorates, especially for coarse-grained HAZ (CGHAZ) and intercritically reheated CGHAZ (ICCGHAZ).

* Corresponding author. E-mail address: [email protected] (Z.-Y. Liu). https://doi.org/10.1016/j.msea.2019.138621 Received 16 January 2019; Received in revised form 30 October 2019; Accepted 31 October 2019 Available online 1 November 2019 0921-5093/© 2019 Elsevier B.V. All rights reserved.

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content, thereby improving HAZ toughness. Based on the above background, this research focuses on improving HAZ cryogenic toughness of low nickel steels by reducing the Si content. The effect of Si content on HAZ microstructure was investigated using multi-pass shielded metal arc welding (SMAW). This work can provide an understanding of the microstructural evolution of the HAZ with Si content, which would be helpful for the design of novel LNG tank steels.

Average prior austenite grain size was determined from SEM micro­ graphs using linear intercept method. TEM specimens with 3 mm diameter were cut from HAZ, mechanically thinned to 50 μm thickness, and then twin-jet electropolished using a solution of 10.0% perchloric acid and 90.0% alcohol at 25 � C and 32 V. The holes of TEM specimens were located near FLþ1 mm (brown and dark blue solid dots in Fig. 1). Fractographs of Charpy impact specimens were also observed using SEM.

2. Experimental

3. Results

The experimental samples were of two types of low nickel steels with different Si contents (0.05%Si and 0.29%Si), which also contained 6.5% Ni, 0.05%C, 0.80%Mn, 0.58% (Cr þ Mo), and balance Fe. The experi­ mental steels were melted in a vacuum induction furnace and cast as 100 kg ingots. The as-cast ingots were homogenized at 1200 � C for 2 h, hot rolled into 15 mm thick plates, and then cooled to room temperature at a rate of 50 � C/s. The rolling reduction at non-recrystallized region was ~68%, and the finishing rolling temperature was ~770 � C. The rolled plates were further quenched after intercritical annealing at 700 � C for 40 min and tempered at 580 � C for 60 min. A pair of plates with dimensions of 15 mm � 90 mm � 300 mm (longitudinal direction) were cut from each UFC-LT treated plate for butt welding. Asymmetrical double-V grooves were machined in the plates. SMAW was performed using LINCOLN NiCrMo-6 electrodes with 4 mm diameter. The chemical composition of electrodes was 69.0%Ni, 0.05%C, 0.40%Si, 2.90%Mn, 12.9%Cr, 6.0%Mo, 1.40%Nb, 0.10%Cu, 1.50%W, and balance Fe. The parameters used during the SMAW process were: voltage of 22 V, cur­ rent of 150 A, welding speed of 19.8 cm min 1, inter-pass temperature below 120 � C, number of welding passes: 9, and the corresponding heat input of ~10 kJ cm 1. Charpy impact specimens with dimensions of 10 mm � 10 mm � 55 mm (transverse direction) were machined from the thickness center of welded plates. In order to ensure the accuracy of Charpy V-notch locations, specimens were macro-etched in a 4% nital solution to visualize the HAZ outline. According to IGF code 16.2.2.3 [23], Charpy V-notch was prepared at weld metal (WM), fusion line (FL), and 1 mm from FL (FLþ1 mm) to test cryogenic toughness of these locations, as shown in yellow, red, and dark blue solid lines in Fig. 1. FL is defined as the location where half of the V-notch is in WM [24]. Charpy impact tests were conducted at 196 � C in a liquid nitrogen environment. Vickers microhardness distribution of welded joints before and after intercritical reheating was determined using FM-700 hardness tester with 200 g load. The selected locations for indentations were marked to facilitate microhardness measurements (purple and orange solid lines in Fig. 1). Microstructure of the HAZ was characterized by a Zeiss Ultra 55 scanning electron microscopy (SEM) and a Tecnai G2 F20 transmission electron microscopy (TEM). SEM specimens were cut from the welded plates along transverse section perpendicular to welding direction, polished, and etched in a 4% nital solution. The locations represented by SEM micrographs are shown by the three colored solid dots in Fig. 1.

3.1. Microstructural characterization Fig. 2(a) through 2(e) present the SEM micrographs at FL and FLþ1 mm locations, as indicated by the green and brown solid dots in Fig. 1. The WM containing high Ni content was not etched by 4% nital solution, so its microstructure was not revealed. For both steels, microstructures at FL and FLþ1 mm were characterized by coarse lath martensite, indicating that the region between FL and FLþ1 mm corresponded to CGHAZ. Average prior austenite grain size in CGHAZ of 0.05%Si steel and 0.29%Si steel reached 78.4 μm and 76.6 μm, respectively. As shown in Fig. 2(e) and (f), a certain amount of fine carbides were dispersed in CGHAZ martensite matrix of 0.05%Si steel. These carbides showed the characteristics of cementite formed in auto-tempered martensite [25–28]. However, no cementite precipitation was found in the matrix of 0.29%Si steel. Fig. 3 presents the SEM micrographs and TEM micrographs at FLþ1 mm, represented by the dark blue solid dot in Fig. 1, which characterize the ICCGHAZ microstructure. After intercritical reheating, the martensite lath boundary became unclear and M-A constituents were formed. For 0.05%Si steel, the morphologies of M-A constituents were predominantly blocky, and a small amount of fine elongated M-A con­ stituents were distributed between laths (Fig. 3(a) and (b)). Selected area diffraction pattern confirmed the presence of austenite in the M-A constituents (see inset of Fig. 3(b)). TEM observations confirmed that MA constituents also contained twin martensite. Generally, the austenite and twin martensite are associated with high C content and result from C partitioning during intercritical reheating. In addition, carbides in CGHAZ were retained, which were identified as cementite by selected area diffraction pattern (Fig. 3(e)). When Si content was increased to 0.29%, most of the M-A constituents exhibited elongated morphology (Fig. 3(c) and (d)), and no cementite precipitation was observed. Area fraction of M-A constituents was obtained by counting at least 20 SEM micrographs. With the increase in Si content from 0.05% to 0.29%, the area fraction of M-A constituents increased from 9.8% to 24.7%. For morphology characterization, maximal size Lmax and mini­ mal size Lmin of each M-A constituent were measured, which are indi­ cated by the maximal and minimal “Feret diameter” in the image analysis shown in Fig. 4. Aspect ratio Lmax/Lmin characterizes the elon­ gation of a given M-A constituent. For each steel sample, at least 500 particles were included to avoid bias. The statistical analysis in Fig. 5

Fig. 1. Schematic diagram of Charpy V-notch locations (yellow, red, and dark blue solid lines), microhardness measurement locations (purple and orange solid lines) and microstructure characterization locations (green, brown and dark blue solid dots). (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.) 2

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Fig. 2. SEM micrographs of welded joints without intercritical reheating ((a) FL of 0.05%Si steel, (b) FL of 0.29%Si steel, (c) (e) FLþ1 mm of 0.05%Si steel and (d) FLþ1 mm of 0.29%Si steel), TEM micrograph of welded joints without intercritical reheating ((f) FLþ1 mm of 0.05%Si steel).

0.05%Si steel. On the other hand, microhardness remained high for 0.29%Si steel initially, and then started to decrease, which suggests that the ICCGHAZ microhardness of 0.29%Si steel was higher than that of 0.05%Si steel. This can be ascribed to the formation of a large amount of hard-brittle M-A constituents in 0.29%Si steel and the precipitation of cementite in 0.05%Si steel.

shows that Si significantly increased the proportion of M-A constituents with large Lmax and large Lmax/Lmin, indicating coarsening and morphological changes in M-A constituents. 3.2. Vickers microhardness distribution Vickers microhardness distribution of welded joints without inter­ critical reheating (at the locations marked by purple solid line in Fig. 1) is shown in Fig. 6(a). It is evident that the microhardness increased significantly when crossing from WM to CGHAZ. Average CGHAZ microhardness was 308 HV for 0.05%Si steel and 342 HV for 0.29%Si steel, indicating that CGHAZ microstructure was clearly softened after reducing the Si content. With further increase in the distance from FL, microhardness decreased gradually, and then tended to be stable after reaching BM. Fig. 6(b) shows Vickers microhardness distribution of welded joints after intercritical reheating (represented by orange solid line in Fig. 1). Similar to Fig. 6(a), microhardness increased sharply near FL. As the distance from FL increased, microhardness decreased gradually for

3.3. Charpy impact toughness Fig. 7 presents the Charpy impact absorbed energy at different lo­ cations. Absorbed energy at each location was the average of three measurements. Cryogenic toughness of both steels at WM was similar, due to the same electrodes and welding parameters employed for both steels. However, absorbed energy at FL and FLþ1 mm locations decreased sharply when Si content was increased to 0.29%. Since the Vnotch at FL and FLþ1 mm partially crosses ICCGHAZ, absorbed energy at FL and FLþ1 mm can be used as a reference for measuring cryogenic toughness of ICCGHAZ. In other words, higher Si content deteriorated cryogenic toughness of ICCGHAZ for low nickel steel. 3

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Fig. 3. SEM micrographs of welded joints after intercritical reheating ((a) FLþ1 mm of 0.05%Si steel and (c) FLþ1 mm of 0.29%Si steel), TEM micrographs of welded joints after intercritical reheating ((b) (e) FLþ1 mm of 0.05%Si steel and (d) FLþ1 mm of 0.29%Si steel). Selected area diffraction patterns from austenite in (b) and cementite in (e) are displayed in the insets of (b) and (e), respectively.

4. Discussion

steel was 34 HV lower than that of 0.29%Si steel, which is possibly because the lower Si content promoted auto-tempering of CGHAZ. During the auto-tempering phenomenon, martensite formed during cooling is simultaneously tempered, leading to the formation of cementite particles [28]. The solid solubility of Si in cementite is negligible. Cementite can nucleate and grow only in a region where Si concentration is less than a critical value [16]. Martensite auto-tempering occurs at a relatively low temperature at which the diffusion of substitutional atoms is limited. In this case, if the initial Si content of steel is too high, the precipitation kinetics of cementite would be dramatically reduced. Supersaturated C atoms in 0.05%Si steel

4.1. Microstructural evolution Fig. 8 shows the microstructural evolution of HAZ in 0.05%Si steel and 0.29%Si steel. During the welding process, BM adjacent to FL was rapidly heated to above 1300 � C, leading to prior austenite grain coarsening. Coarse lath martensite was obtained after fast weldingcooling due to the addition of small amounts of Cr and Mo to improve hardenability. There was no obvious distinction in prior austenite grain size in both steels. However, average CGHAZ microhardness of 0.05%Si 4

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fully recovered [29]. On the other hand, part of the matrix reverted to austenite and became enriched in the C atoms [11,30]. The austenite, which was not sufficiently stable, then partly transformed to twin martensite, and the remaining austenite was retained at room temper­ ature. In 0.05%Si steel, blocky reversed austenite, i.e. blocky M-A

Fig. 4. Definition of maximal size Lmax and minimal size Lmin of each M-A constituent.

preferred to form cementite, thereby softening the matrix. In ICCGHAZ, two transformation processes occurred because sub­ sequent welding passes reheated CGHAZ to the temperatures between Ac1 and Ac3. On the one hand, original martensite microstructure was

Fig. 7. Charpy impact absorbed energy at WM, FL, and FLþ1 mm.

Fig. 5. Statistics on Lmax/Lmin and Lmax of at least 500 M-A constituents in ICCGHAZ of (a) 0.05%Si steel and (b) 0.29%Si steel.

Fig. 6. Vickers microhardness distribution of welded joints (a) before and (b) after intercritical reheating. 5

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initiation. The existence of blocky M-A constituents did not cause a significant decrease in cryogenic toughness of 0.05%Si steel. For 0.29%Si steel, a quasi-cleavage fracture was evident from the fractograph. Dotted circles in Fig. 9(b) represent examples of cleavage initiation sites. The particles marked by dotted circles 1 show typical characteristics of elongated M-A constituents. The holes marked by dotted circles 2 correspond to the locations where elongated M-A con­ stituents were located and pulled out. Cleavage facets initiated at these sites propagated along a straight crack path and exhibited a river pattern. Mohseni et al. concluded that the fracture mechanism was debonding between the M-A constituent and surrounding matrix based on a similar fractographic observation [31]. Higher hardness of elon­ gated M-A constituents can introduce stress concentration during loading. When the stress concentration exceeds a critical value, elon­ gated M-A constituent is prone to be debonded from the surrounding matrix, and cleavage crack is easily initiated. Besides, several nearby holes were observed at the initiation site of river pattern, indicated by dotted circles 3. These holes could be related to the initiation of cleavage crack from the region between several adjacent M-A constituents, where there are overlapping stress concentrations due to debonding of M-A constituent from the matrix [31,32]. Li et al. and Luo et al. reported that elongated M-A constituents were more likely to be debonded from matrix compared to blocky M-A con­ stituents [33,34]. This is consistent with the phenomenon observed herein. The work done by Li et al. showed that elemental segregation at the interface between the elongated M-A constituent and matrix may promote debonding of M-A constituent from the matrix and assist nucleation of cleavage crack [35]. According to the classical Griffith theory, as the size of M-A constituent corresponding to critical micro­ crack size increases, critical cleavage stress decreases, which facilitates cleavage crack initiation [36]. Considering the presence of a large amount of coarse elongated M-A constituents in ICCGHAZ of 0.29%Si steel, cracks immediately coalesced, and fracture of impact specimens took place. All the above results suggest that low nickel steels should be designed with reduced Si content to obtain a softened microstructure and the required cryogenic toughness in HAZ. The Charpy impact absorbed energy at 3 mm and 5 mm from fusion line was tested to be 175 J and 162 J, respectively. Cryogenic toughness at all five locations meets the requirements of 9%Ni steels reported in EN 10028-4 [37]. Therefore, the newly developed low nickel steels can potentially replace 9%Ni steels for the construction of giant LNG tanks.

Fig. 8. Microstructural evolution of HAZ.

constituents, were formed preferentially at high-angle grain boundaries such as prior austenite grain boundaries, and only a few fine elongated M-A constituents were formed at low energy lath boundaries. This may be related to the cementite precipitation, which can reduce C content in the surrounding matrix and inhibit the formation of reversed austenite during intercritical reheating. In contrast, in 0.29%Si steel, higher C content in the matrix caused the formation of a large amount of coarse elongated M-A constituents at lath boundaries. 4.2. Fracture analysis Fractographs of Charpy impact specimens at FLþ1 mm are presented in Fig. 9. Several ductile dimples were observed in the fractograph of 0.05%Si steel, indicating that the fracture occurred through micro-void growth. A clear plastic deformation before fracture consumes a lot of energy; hence, a higher absorbed energy can be obtained. This effect is related to the softening of matrix by martensite auto-tempering, which improves plastic deformation ability of the matrix and delays crack

5. Conclusions The effect of Si content in low nickel steels on the microstructure and resultant cryogenic toughness of HAZ was studied through multi-pass SMAW. The major conclusions are summarized as follows:

Fig. 9. Fractographs of Charpy impact specimens at FLþ1 mm of (a) 0.05%Si steel and (b) 0.29%Si steel. 6

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(1) CGHAZ microstructures of 0.05%Si steel and 0.29%Si steel were composed of coarse lath martensite. The decrease in Si content promoted auto-tempering of CGHAZ, which softened the matrix and improved plastic deformation ability of the matrix. (2) In ICCGHAZ, apparent recovery and formation of M-A constitu­ ents occurred. With decrease in Si content, the amount and size of M-A constituents significantly decreased, and their elongated morphologies were transformed to blocky morphology. Blocky M-A constituents in 0.05%Si steel were less detrimental to toughness than elongated M-A constituents in 0.29%Si steel, which were easily debonded from the matrix. (3) Charpy impact absorbed energy in the HAZ at FL and FLþ1 mm were significantly improved by reducing the Si content. Low nickel steels with reduced Si content would be suitable for the preparation of LNG tanks.

[11] [12] [13] [14] [15] [16] [17] [18]

Declaration of competing interest

[19]

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

[20] [21] [22]

Acknowledgements The authors gratefully appreciate the financial support by the Na­ tional Key R&D Program of China (Grant No. 2017YFB0305000), the Fundamental Research Funds for Central Universities (Grant No. N170708018), and the State Natural Science Foundation of China (Grant No. U1660117).

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Appendix A. Supplementary data

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Supplementary data to this article can be found online at https://doi. org/10.1016/j.msea.2019.138621.

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