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Effect of Si content on the microstructures and the impact properties in the coarse-grained heat-affected zone (CGHAZ) of typical weathering steel
Materials Science & Engineering A 762 (2019) 138082
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Effect of Si content on the microstructures and the impact properties in the coarse-grained heat-affected zone (CGHAZ) of typical weathering steel
T
Yue Zhanga,b, Genhao Shia,b, Rui Suna,b, Kai Guoa,b, Chunling Zhanga,b,*, Qingfeng Wanga,b a b
State Key Laboratory of Metastable Materials Science and Technology, Yanshan University, 066004, Qinhuangdao, China National Engineering Research Center for Equipment and Technology of Cold Strip Rolling, Yanshan University, 066004, Qinhuangdao, China
A R T I C LE I N FO
A B S T R A C T
Keywords: Silicon content Toughness Coarse-grained heat-affected zone Weathering steel M/A constituents
This study aims to make clear the content limit of silicon (Si), which normally favors weathering steel but probably impairs its impact properties in the coarse-grained heat-affected zone (CGHAZ) induced by the conventional arc welding. Thermal simulation with a Gleeble 3500 simulator was thus carried out to reproduce the CGHAZ of this type of steel with differing silicon (Si) content from 0.12% to 0.48%–0.75%. Charpy V-notch impact tests and microstructure observations using optical microscopy, transmission and scanning electron microscopy (TEM and SEM), and electron backscatter diffraction (EBSD) were performed. The results showed that a microstructure composed of granular bainite (GB), acicular ferrite (AF), lath bainite (LB), and martensite/ austenite (M/A) constituents formed in 0.12%Si and 0.48%Si samples, while the LB scarcely appeared in 0.75% Si sample. When the Si content increased from 0.12% to 0.48%–0.75%, the LB decreased and in contrast, the GB increased. In addition, the area fraction (average size) of the M/A constituents increased from 5.8% (0.70 μm) to 9.2% (1.32 μm), while the number fraction (mean size) of effective grains with high-angle grain boundaries (HAGBs) decreased (increased) from 27% (3.77 μm) to 10% (5.71 μm). Accordingly, the Charpy impact energy of the CGHAZ decreased sharply from 208 to 42 J. The increased large M/A constituents significantly enable the microcrack nucleation, while the reduced HAGBs inefficiently impede the microcrack propagation, finally leading to a deteriorated toughness in the CGHAZ of 0.75%Si steel. The silicon content in this steel should not exceed 0.48%.
1. Introduction Typical weathering steel with an enhanced atmospheric corrosion resistance by the addition of Ni, Cr and Cu, has been developed over decades and extensively applicated for fabricating the uncoated building and bridge structures via the conventional arc welding process. This type of low-alloyed steel can normally obtain excellent comprehensive mechanical properties via thermo-mechanical controlled process. However, previous studies [1–4] reported that the welding-induced heat-affected zone can suffer from the inhomogeneity of microstructure and impact toughness. It is generally accepted that the worst toughness may frequently occur in the coarse-grained heat-affected zone (CGHAZ), owing to the martensite/austenite (M/A) constituents acting as a hard and brittle phase, the coarse prior austenite grain (PAG), and the correspondingly coarse granular banaite (GB) [5–7]. Some studies [8,9] have focused on how to enhance the poor impact properties in CGHAZ through optimizing the alloying element
and the corresponding content, including Ni, Cu and Nb, etc. Silicon (Si) is also an alloying element frequently added to the typical weathering steel, for it can enable the formation of a stable and compact rust layer and thereby enhance the atmospheric corrosion resistance [10,11]. Additionally, the Si can notably enhance the tensile strength of any steel through solid solution strengthening [12]. However, an excessive of Si content can also weaken the welding performance by degrading the impact toughness of CGHAZ [13]. Ruan et al. [14] reported that the decreasing Si content from 1.55% to 0.54% improves the toughness of Cr–Mn steel. The carbide particles can be refined by a lowered Si content, leading to an enhanced the low-temperature impact toughness. Taillard et al. [13] proved that the carbon enrichment in retained austenite of industrial E355 structural steel facilitated by the increasing Si content, significantly promotes the formation of M/A constituents, resulting a local brittle zone subjected to twice welding thermal cycles. Furthermore, a high Si content inhibits the decomposition of M/A constituents in the tempering process,
*
Corresponding author. State Key Laboratory of Metastable Materials Science and Technology, Yanshan University, 066004, Qinhuangdao, China. E-mail addresses: [email protected] (Y. Zhang), [email protected] (G. Shi), [email protected] (R. Sun), [email protected] (K. Guo), [email protected] (C. Zhang), [email protected] (Q. Wang). https://doi.org/10.1016/j.msea.2019.138082 Received 11 March 2019; Received in revised form 15 June 2019; Accepted 27 June 2019 Available online 28 June 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.
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leading to a lower toughness. Cao et al. [15] supposed that the increasing Si contents of X120 steel from 0.3% to 2.0% increase the amount fraction of GB, leaving the unaltered density of the high angle grain boundaries (HAGBs) in the CGHAZ [4]. The Si content limits of 50 W (0.30%–0.65% for Type A and 0.15%–0.50% for Type B), HPS 50W/70 W (0.30%–0.50%) and Q420qNH (0.15%–0.50%) weathering steels have been specified by ASTM A709/A709 M [16] and Chinese GB/T 714 [17] standards, respectively. It is necessary to understand the effect of Si content on the microstructures and impact properties in the welding-induced CGHAZs of these typical weathering steels, and the underlying mechanisms. This can help optimize the Si content. However, to our knowledge, the closely relevant research has not been reported yet. For this attempt to make clear the Si content limits due to the possible toughness loss in the CGHAZ of typical weathering steel, three experimental steels with the targeted/actual Si contents of 0.15%/ 0.12%, 0.50%/0.48% and 0.65%/0.75% were particularly designed/ prepared for welding thermal simulation tests. We investigated the correlation between the Si content, the microstructures, and the impact properties in the simulated CGHAZ. The microstructural features for each Si-bearing sample were systematically clarified. In particular, we focused on elucidating the influence of the Si content on the impact behavior and the mechanism of microcrack formation and propagation. Finally, an optimal range of Si content was proposed to improve the welding performance.
Fig. 1. Specimen orientation in the rolled plate and geometry. (RD-rolling direction; ND-normal direction; TD-transverse direction).
2. Experimental procedure The steels with the target Si content of 0.15%, 0.50%, and 0.65% chosen for the study were melted in a 50 kg vacuum furnace and hotrolled to 18-mm thick plates. The chemical compositions of three steels with the actual Si content of 0.12%, 0.48%, and 0.75% are listed in Table 1, which were described as 12Si, 48Si, and 75Si, respectively. Welding thermal cycle simulations were performed on a Gleeble3500 thermal cycle simulator for estimating the microstructure and impact property evolutions in the CGHAZ. The samples for simulation were cut along the rolling direction and processed into the rods of 10.5 × 10.5 × 80 mm3, as shown in Fig. 1. The thermal cyclying was conducted with the welding heat input of 20 kJ/cm, the mean heating rate of 100 °C/s and the peak temperature of 1320 °C for 1s, as shown in Fig. 2. Expansion curves were also measured and the onset/finish temperatures (Ar3/Ar1) for the austenite (γ) → ferrite (α) transformation were determined using the tangent method. To analyze the influence of Si addition on the microstructures and the Charpy impact properties, the simulated sample was cut along the location of the thermocouple. The samples obtained were mechanically polished and etched in a 4% Nital solution in volume. The morphology of the microstructure was characterized with an axiover-200mat metallographic microscope (OM) and an S3400 scanning electron microscope (SEM). To quantify the phase, the mean equivalent grain size, and the grain boundary misorientation, a solution of glycerin: perchloric acid: alcohol at a ratio of 0.5:1:8.5 was used to electropolish the samples. The samples were then observed with a Hitachi s-3400 scanning electron microscope using electron backscatter diffraction (EBSD) for orientation imaging microscopy (OIM) with the following parameters: a working distance of 15 mm, a step size of 0.3 μm, and acceleration voltage of 20 kV, and a tilt angle of 70°. The effective grains are regions
Fig. 2. The welding thermal cycle curves.
enclosed by low and high angle boundaries revealed by EBSD which was calculated as the equivalent circle diameter related to the individual grain area. Tolerance angles ranging from 2° to 30° are applied. Furthermore, the ferrite plates and M/A constituents were characterized using a JEM-2010 transmission electron microscope (TEM). The thin foils used for TEM were prepared by the twin-jet method and electro-polished in a 7% perchloric acid/glacial acetic acid solution at room temperature. The metallographic samples were etched with the Lepera reagent (4% picric acid alcohol solution: 1% sodium metabisulfite aqueous solution at a ratio of 1:1) to observe the M/A constituents. With this technique, the morphological distribution of the M/ A constituents (bright white) in the ferrite matrix (yellow) can be simply identified by OM. Subsequently, the image-pro plus software was used to perform a quantitative analysis of the images of the M/A constituents and at least 5 fields of view were measured for each sample at a magnification of 1000. The size of the M/A constituents was calculated as the equivalent diameter of the individual M/A constituent area. The average size was statistically obtained by measuring at least 1000 M/A constituents for each specimen. To conduct the Charpy impact tests, the thermal simulation samples were machined into normative Charpy v-notch samples with a
Table 1 Chemical composition of the test steels in wt. %. Steel
C
Si
Mn
P
S
Ni
Cr
Mo
Cu
Nb
Ti
12Si 48Si 75Si
0.053 0.058 0.058
0.12 0.48 0.75
1.25 1.27 1.27
0.007 0.007 0.008
0.002 0.005 0.002
0.31 0.32 0.31
0.51 0.52 0.52
0.05 0.05 0.05
0.26 0.25 0.27
0.028 0.028 0.027
0.013 0.014 0.016
2
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Fig. 3. Optical micrographs of the simulated CGHAZ samples of (a) 12Si, (b) 48Si and (c) 75Si steels, together with the M/A constituents observations of (d) 12Si, © 48Si and (f) 75Si samples, where the ferrite matrix is yellow and the M/A constituent is white. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
dimension of 10 × 10 × 55 mm3, as shown in Fig. 1. The Charpy V impact tests were conducted on a JB-300B impact tester at −40 °C and the average value was recorded. SEM was used to observe the fracture surfaces of the specimens and the sections perpendicular to the notch were used to characterize the second crack propagation by SEM and EBSD.
Table 2 Summary of the microstructure observations and quantifications.
3. Results
PAG—prior austenite grain; Dγ—mean size of PAG; M/A—martensite-austenite constituents; fM/A— area fraction of M/A constituents; dM/A— mean size of M/A constituents; MEDθ≥15°—mean equivalent diameter of ferrite grain with boundary misorientation angle θ > 15°; f (θ≥15°) —fraction of boundaries with misorientation angle θ ≥ 15°; pDFZ—area proportion of ductile fracture zone.
3.1. Microstructure Typical optical micrographs of each Si-containing CGHAZ sample are presented in Fig. 3. They show the microstructures on the left and the M/A constituents on the right. The quantitative analysis results are shown in Table 2. The mixed microstructure of granular bainite (GB), lath bainite (LB), acicular ferrite (AF), and martensite/austenite (M/A) constituent formed in 12Si and 48Si samples, while the LB scarcely appeared in 75Si sample. The microstructure observations showed a decrease in LB and an increase in GB when the Si content increased from 0.12% to 0.75%. In addition, the area fraction fM/A and the average size dM/A of the M/A constituents in CGHAZ increased from 5.8% to 9.2% and from 0.70 μm to 1.32 μm, respectively. Fig. 4 shows the size distribution of the M/A constituents varying with the increasing Si content, indicating that the M/A constituents with the size larger
Steel
Dγ/μm
fM/A/%
dM/A/μm
MEDθ≥15°/μm
f
(θ≥15°)/%
12Si 48Si 75Si
57.2 ± 2 62.8 ± 2 65.0 ± 3
5.8 ± 0.2 7.7 ± 0.3 9.2 ± 0.2
0.70 0.89 1.32
3.77 4.45 5.71
27 19 10
PDFZ/% 41.5 25.4 0.0
than 1 μm increased remarkably. Fig. 5 shows the typical microstructure morphology obtained by TEM. The TEM micrographs showed the formation of ferrite plates and island-like phases, as shown in Fig. 5(a) and (b). There are two types of bainitic ferrite distinguished by their morphologies: a plate or a block. The plate/block-like bainitic ferrite could be categorized as LBF/GBF [18]. In addition, typical M/A constituents in stripe and block shapes were observed in the bright field images (Fig. 5(c) and (f)), dark field images (Fig. 5(d) and (g)), and in the selected area diffraction pattern (Fig. 5(e) and (h)). They suggest the selected area was composed of martensite and austenite phases. The M/A constituents were dispersed 3
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between the bainitic ferrites in elongated (Fig. 5(c)) or blocky (Fig. 5(f)) shapes. Extensive TEM observations showed that the increasing Si content led to an increase in GB, a decrease in LB, and an increase of the area fraction and the average size of the M/A constituents, which was consistent with the OM observation results. The mean equivalent diameter (MED) of ferrite grains with boundaries defined by a misorientation tolerance angle (MTA) ranging 2o-60° for each sample was determined via EBSD, and the typical inverse pole figures are given in Fig. 6. When the MTA varied from 2° to 60° as a function of the MTA, the MED had a monotonic increasing trend with the increase of Si content, as shown in Fig. 7(a). In addition, the boundaries of ferrite grains could be divided into two categories: low angle grain boundary (LAGB, 2°≤MTA < 15°) and high angle grain boundary (HAGB, MTA≥15°). The white and black lines in Fig. 6 represented the LAGB and the HAGB, respectively. The HAGB can effectively impede crack propagation in a notched impact sample, according to Refs. [19], and hence the MED of grains with the HAGB can be considered as the effective grain size. Table 2 summarizes the number fraction of high-MTA, fθ≥15° boundary and the corresponding MEDθ≥15°. As shown in Table 2, both of them increased with the increase in Si content. For further clarification, Fig. 7(b) summarizes the
Fig. 4. Number density distribution of the M/A constituents vs. size for different Si contents.
Fig. 5. Typical TEM observations of the microstructure in CGHAZ samples of (a) 12Si and (b) 75Si steels, with the M/A constituents in stripe/block shape indicated particularly by (c/f) bright field, (d/g) dark field, and (e/h) the selected area diffraction pattern. GB- granular bainite; LB- lath bainite; M/A-martensite-austenite. 4
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Fig. 6. EBSD inverse pole figures for CGHAZ samples of (a) 12Si, (b) 48Si and (c) 75Si steels, with the white lines indicating LAGBs and the black lines indicating HAGBs.
Fig. 7. MED varied with MTA (a) and number fraction of grain boundary against misorientation angle for each Si-containing steel CGHAZ obtained by EBSD.
significantly, which might be associated to the evolution of the microstructures. The impact fracture surface of the CGHAZ for each sample with varying Si content was observed by SEM (Fig. 9). The blue line represented the margin of the v-shaped notch, and the yellow line indicated the boundary between the ductile fracture zone and the brittle fracture zone, while the red line showed the boundary between the brittle fracture zone and the hammer-impacted zone. According to the morphological characteristics of the fractures and the corresponding local positions, the complete fracture surfaces were of a mixed type where a ductile fracture zone and a brittle fracture zone coexist for 12Si
misorientation distribution for each sample. The statistics showed that the number fraction of HAGBs increased from 27% to 10% when the Si content increased from 0.12% to 0.75%. 3.2. Charpy impact toughness The Charpy impact energies determined at −40 °C for the simulated CGHAZ with different Si contents are shown in Fig. 8. The Charpy Vnotch (CVN) impact energy decreased sharply from 208 J to 42 J when the Si content increased from 0.12% to 0.75%. When the Si content reached 0.75%, the impact toughness of the test steels deteriorated 5
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constituents increased with the Si content, whereas the LB decreased. The evolution of the microstructure in CGHAZ had a close relationship with the Si content in steel, which can be explained as follows. 4.1.1. Effect of increased Si content on mixed LB + GB + AF microstructure The expansion curve for the CGHAZ of each Si-containing sample, which was subjected to the welding thermal cycle (Fig. 2), was measured and shown in Fig. 11. The temperature (Ar3) at which the phase transformation of γ→α started to take place on cooling, was determined and was also indicated in this figure. As Fig. 11 shows, the CGHAZs of 12Si and 48Si steels have the relatively lower Ar3 (586 °C and 599 °C) than that of 75Si steel, indicating that the phase change of γ→α proceeds at a relatively low onset temperature in these two samples. According to Ref. [22], the γ→GB transition normally precedes the γ→LB transition, and the GB also commonly nucleates and grows from the prior austenite grain boundary (PAGB). This enables the GB to constitute the leading phase of mixed GB + AF + LB microstructure. By contrast, the LB which mainly nucleates intragranularly according to Ref. [23] and grows in the retained space, accounts for a small proportion. This indicates that the GB could form with a higher thermodynamic driving force in the CGHAZ of 12Si and 48Si steels than the LB. These phase changes and products are in partly agreement with those of a low carbon Mo–V–Ti–N steel subjected to the controlled cooling at different rates [24]. Furthermore, with the Si content increasing to 0.75%, the Ar3 increased to 611 °C. The onset temperature for the phase change of γ→α was raised accordingly. This leads to an obviously increasing amount of the GBF at the expanse of LB. Also for this reason, the GB dominates, while the LB scarcely appears in the CGHAZ of 75Si steel. These discussions have clarified that increasing Si content facilitated the formation of GBF, and hence increased its fraction in the CGHAZ of weathering steel. The CGHAZs of 12Si and 48Si steels have the higher supercooling degree for the γ→α transformation than that of 75Si steel, owing to their lower Ar3. This may lead to a higher driving force and a higher nucleation rate of α in these two steels, which eventually result in the smaller equivalent grain size or MED at any MTA. Correspondingly, the effective grain size (MED MTA≥15°) for the CGHAZs of these two steels are smaller than that for 75Si steel, as indicated in Fig. 7(a) and Table 2. Meanwhile, the lowered Ar3 also promoted the formation, and accordingly led to an increasing fraction, of the thinner LB and probably the finer AF in the CGHAZs of 12Si and 48Si steels, as shown in Figs. 3 and Figs. 5. The boundaries of bainitic packet [22] and the AF [25] can additionally constitute the HAGBs in a bainite and ferrite dual-phase steel. This leads to the elevated number fraction of HAGBs for 12Si and 48Si steels in comparison with the 75Si steel, as indicated in Fig. 7(b) and Table 2, owing to the lowered Ar3 with the increasing Si content.
Fig. 8. Charpy impact energy for each sample varied with Si content.
and 48Si samples, but there was no ductile fracture zone on the fracture surface of 75Si sample. The area proportion of the ductile fracture zone of 12Si (41.5%) was much larger than that of 48Si (25.4%). The ductile fracture zone of 12Si and 48Si were comprised of a large number of cup-shaped and cone-shaped dimples, as shown in Fig. 9Ⅰ and Fig. 9Ⅲ. The brittle fracture zones of 12Si and 48Si were dominated by high density river-like patterns and tear ridges with small dimples, as seen in Fig. 9Ⅱ and Fig. 9Ⅳ. The junction of the cleavage facets by ductile fracture bands was considered as the main energy absorbing mechanism in the crack propagation process and kept the impact toughness at a high level [20]. As shown in Fig. 9(c) and Ⅴ, and Fig. 9Ⅵ for 75Si, the fracture was a whole cleavage fracture surface with large cleavage facets. Numerous large cleavage facets propagated in a linear manner and the straight river pattern extended from the cleavage crack to the periphery, forming a fracture step when encountering the highangle grain boundary and following a typical brittle fracture mechanism with low impact toughness. The apparent difference in fracture morphology between 12Si, 48Si and 75Si indicated that the impact toughness of the CGHAZ of 12Si sampe was better than that for 48Si sample, while 75Si sample showed the lowest impact toughness. To further analyze the influence of the microstructure on crack initiation and propagation, the secondary cracks underneath the fracture of the impact specimen were observed by SEM, as shown in Fig. 10. Numerous microvoids were observed in the matrix for the ductile fracture zone of 0.12% Si samples. The failure process of the dimple fracture was by accumulation of microvoids and their initiation. With the increasing Si content, the proportion of cleavage zone increased. The main fracture failure process went from the microvoids accumulation to a linear expansion of the microcracks. For 75Si sample (Fig. 10 (b)), the microcrack propagated through the interior of the prior austenite grain following a linear pattern until they were deflected or stopped at the prior austenite grain boundaries. As seen in Fig. 10, the microvoids and microcracks nucleated frequently at or around the smaller (< 1 μm, typically) and larger (≥1 μm, typically) M/A constituents, respectively, and propagated through the interface between the M/A constituents and the bainitic matrix. The measured impact energies showed a significant toughness degradation when the Si content varied from 0.12% to 0.75%, which was usually related to the distribution of the high angle grain boundaries in the matrix and the size/amount of the M/A constituents [21].
4.1.2. Effect of increased Si content on M/A constituents As described in Section 3.1, a mixed microstructure consisting of GB, AF and a small amount of LB and M/A constituents formed in the CGHAZs of 12Si and 48Si steels, while the GB + AF + M/A constituents in that of 75Si steel. The continuous cooling phase transformation of γ→α (bainitic ferrite and AF) + γ′ (metastable austenite) actually started from the Ar3 and finished at the Ar1, during which the γ/α interface normally moves toward to the γ′ phase. Since the solubility of carbon in the metastable austenite is different from that in ferrite [26], this phase transformation process is supposed to be accompanied by the re-distribution of carbon atoms from α to γ′ via bulk diffusion and dislocation-related pipe diffusion [27], leading to a C-rich γ′ with an accordingly enhanced stability. Fig. 12 schematically illustrates the carbon atom diffusion from bainitic and acicular ferrite to γ′ during the continuous cooling transformation. Since the Ar3 increased, and the Ar1 decreased with the Si addition, the residence time in the two-phase (α + γ) zone is extended under the same thermal cycle condition. The degree of C-enrichment in γ′ increases, leading to the
4. Discussion 4.1. Effect of the increased Si content on the microstructures During the continuous cooling process, a mixed GB + AF + LB + M/A microstructure formed in CGHAZ of each Sicontaining steel. As shown in Figs. 3 and 5, the GB and M/A 6
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Fig. 9. SEM micrographs of the fracture surface of 12Si (a), 48Si (b), and 75Si(c). Fig.9Ⅰ, Fig.9 Ⅱ, Fig.9Ⅲ, Fig.9Ⅳ, Fig.9Ⅴ, and Fig.9Ⅵ are the magnified SEM micrographs of the selected areas marked as Ⅰ, Ⅱ, Ⅲ, Ⅳ, Ⅴ, and Ⅵ.
further stabilization of γ'. In addition, Si can inhibit the precipitation of carbides from the super-cooled austenite [13] during this phase transformation, and keep the carbon solid solution in the γ′. This leads to the increase in residual austenite content at room temperature. Therefore,
the increasing Si content promotes the C-enrichment and accordingly the stabilization of γ′. The C-rich γ′ would further transform to the M/A constituents at the next continuous cooling stage and determine their morphology and inner-structure. The area fraction of the M/A 7
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Fig. 10. The crack path underneath the ductile fracture surface for (a) 0.12% Si sample and the brittle fracture surface for (b) 0.75% Si sample at the testing temperature of −40 °C.
Fig. 11. Expansion curve and Ar3 temperature measured for each CGHAZ sample with a different Si content.
Fig. 13. EBSD inverse pole figures of the secondary cracks underneath the ductile fracture surface for (a) 12Si sample and the brittle fracture surface for (b) 75Si sample.
could directly alter the impact behavior of the test steels. In the ductile fracture zone of the mixed fracture, the microvoids in the matrix tended to nucleate at carbides or at fine M/A constituents by means of particle fracture or stripping of the particle/matrix interface, as shown in Figs. 10(a) and Fig. 13(a). The microvoids grow via the plastic deformation of the surrounding matrix and their accumulation, which requires a large energy [4,28]. This helps to provide a high impact toughness. Compared with the ductile fracture zone, the cleavage fracture zone is more likely to cause the failure of the steel sample. The secondary crack in 75Si was taken as an example (Figs. 10(b) and Fig.13 (b)) to understand the crack propagation path of the microcracks in the cleavage zone in the CGHAZ of the test steel and the underlying mechanism. The influence of Si on microcrack initiation and propagation can be explained as follows.
Fig. 12. Schematic diagram of the diffusion of C atoms from the bainitic ferrite to the carbon-enriched austenite during the continuous cooling transformation.
constituents in the CGHAZ is ultimately increased, as shown in Fig. 3 and Table 2. Simultaneously, the intersecting area of adjacent GBF and AF plates, where the γ′ is located during this phase transformation, is enlarged with coarsening of these plates, owing to the Si addition and accordingly the elevated Ar3. The average size dM/A of M/A constituents in the CGHAZ is also increased, as indicated in Fig. 5 and Table 2.
4.2.1. Effect of the M/A constituents on the impact behavior in the CGHAZ On the basis of the results from Fig. 8, we conclude that the impact toughness of CGHAZ is closely associated with the number and the size of the M/A constituents in the matrix. The M/A constituents is supposed to be the brittle phase embedded in the ferrite matrix, and hence is one of the main factors determining the impact toughness [29,30]. Therefore, it is particularly important to understand the mirocrack nucleation and propagation assisted by the M/A constituents as displayed in Fig. 10(b). Usually due to the high hardness and brittleness of the M/A constituents, there is a significant difference in micro-hardness between the M/A constituents and the bainitic ferrite matrix [31]. Therefore, the
4.2. Effect of increased Si content on the impact behavior of the CGHAZ To further explore the mechanism of the influence of Si on the impact properties of the simulated CGHAZ, it is necessary to comprehensively consider the effect of the microstructural evolution on the impact toughness. When the Si content increased, the microstructures became coarser, the GB increased, and the LB decreased, while a relatively high area fraction and large size of M/A constituents ensues. These are the reasons why the impact energy remained low, which 8
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content results in the elevated Ar3 and hence, the decreased LB as well as the increased GB and M/A constituents. 2. Since the increasing Si content from 0.12% to 0.48%–0.75% leads to the obviously deteriorated impact properties of the CGHAZ, the silicon content in typical weathering steel should not exceed 0.48%. 3. Due to the increasing Si addition, both the increasing amount of larger M/A constituents and the decreasing fraction of HAGBs enables the microcrack propagation.
soft matrix structure in the sample for impact test usually undergoes an initial plastic deformation, resulting in the local stress concentration on or around the hard M/A constituents. The propagation of microcrack nucleating from the M/A constituents could occur if the locally concentrated stress exceeds the critical stress [32]. This critical stress can be described quantitatively by the classical Griffith local cracking model [1–4], given in equation (1):
σc =
2E γs 1/2 αa
(1)
Author contributions
where σc is the critical stress, E is the Young modulus, γs is the effective surface energy of the interface fracture between the M/A constituents and the bainitic matrix, α is a constant related to the crack shape, and a is the size of the microcrack nucleating at or around the M/A constituents. From the model, the smaller M/A constituents at or around which only a smaller microcrack might nucleate frequently, can cause an increase in the critical stress. Accordingly, this makes it difficult for the smaller microcrack to propagate. When the Si content is increased, the increasing amount of the larger M/A constituents can significantly promote the propagation of microcracks, which is primarily responsible for the lowered impact toughness. In the 12Si sample, the number fraction of the large M/A constituents with the dM/A ≥ 1 μm was as low as less than 50%, as indicated in Fig. 4. Then the nucleation and propagation of microcrack can get more difficult. By contract, a Si content as high as 0.75% turned out to be obviously detrimental to the toughness in the CGHAZ. This can be attributed to low difficulty for the microcrack to nucleate and propagate with respect to a high fraction of the large M/A constituents.
Chunling Zhang and Yue Zhang conceived and designed the experiments; Genhao Shi and Rui Sun performed the experiments; Kai Guo contributed analysis tools and analyzed the data; Qingfeng Wang directed the testing and editing; Chunling Zhang and Yue Zhang wrote the paper. Acknowledgments This work was support by the National Key Research and Development Program of China (Grant No. 2017YFB0304800 and Grant No. 2017YFB0304802 for the second sub project). References [1] L. Lan, C. Qiu, D. Zhao, X. Gao, L. Du, Microstructural characteristics and toughness of the simulated coarse grained heat affected zone of high strength low carbon bainitic steel, Mater. Sci. Eng. A 529 (2011) 192–200. [2] S.G. Lee, S.S. Sohn, B. Kim, W.G. Kim, K.-K. Um, S. Lee, Effects of martensite-austenite constituent on crack initiation and propagation in inter-critical heat-affected zone of high-strength low-alloy (HSLA) steel, Mater. Sci. Eng. A 715 (2018) 332–339. [3] L. Rancel, M. Gómez, S.F. Medina, I. Gutierrez, Measurement of bainite packet size and its influence on cleavage fracture in a medium carbon bainitic steel, Mater. Sci. Eng. A 530 (2011) 21–27. [4] J. Chen, S. Tang, Z.-Y. Liu, G.-D. Wang, Microstructural characteristics with various cooling paths and the mechanism of embrittlement and toughening in low-carbon high performance bridge steel, Mater. Sci. Eng. A 559 (2013) 241–249. [5] S. Kim, Y. Kang, C. Lee, Variation in microstructures and mechanical properties in the coarse-grained heat-affected zone of low-alloy steel with boron content, Mater. Sci. Eng. A 559 (2013) 178–186. [6] J. Chen, M.-y. Lv, Z.-y. Liu, G.-d. Wang, Combination of ductility and toughness by the design of fine ferrite/tempered martensite–austenite microstructure in a low carbon medium manganese alloyed steel plate, Mater. Sci. Eng. A 648 (2015) 51–56. [7] X. Zou, J. Sun, H. Matsuura, C. Wang, In situ observation of the nucleation and growth of ferrite laths in the heat-affected zone of EH36-Mg shipbuilding steel subjected to different heat inputs, Metall. Mater. Trans. B 49 (5) (2018) 2168–2173. [8] X.L. Wan, K.M. Wu, G. Huang, K.C. Nune, Y. Li, L. Cheng, Toughness improvement by Cu addition in the simulated coarse-grained heat-affected zone of high-strength low-alloy steels, Sci. Technol. Weld. Join. 21 (4) (2016) 295–302. [9] X. Zou, D. Zhao, J. Sun, C. Wang, H. Matsuura, An integrated study on the evolution of inclusions in EH36 shipbuilding steel with Mg addition: from casting to welding, Metall. Mater. Trans. B 49 (2) (2017) 481–489. [10] J.A. Mejía Gómez, J. Antonissen, C.A. Palacio, E. De Grave, Effects of Si as alloying element on corrosion resistance of weathering steel, Corros. Sci. 59 (2012) 198–203. [11] T. Nishimura, Corrosion behavior of silicon-bearing steel in a wet/dry environment containing chloride ions, Mater. Trans. 48 (6) (2007) 1438–1443. [12] W.M. Garrison, The effect of silicon and nickel additions on the sulfide spacing and fracture toughness of a 0.4 carbon low alloy steel, Metall. Trans. A (Phys. Metall. Mater. Sci.) 17 (4) (1986) 669–678. [13] R. Taillard, P. Verrier, T. Maurickx, J. Foct, Effect of silicon on CGHAZ toughness and microstructure of microalloyed steels, Metall. Mater. Trans. A 26 (2) (1995) 447–457. [14] L.H. Ruan, K.M. Wu, J.A. Qiu, A.A. Shirzadi, I.G. Rodionova, Effect of silicon content on carbide precipitation and low-temperature toughness of pressure vessel steels, Met. Sci. Heat Treat. 59 (1–2) (2017) 97–101. [15] R. Cao, Z.S. Chan, J.J. Yuan, C.Y. Han, Z.G. Xiao, X.B. Zhang, Y.J. Yan, J.H. Chen, The effects of Silicon and Copper on microstructures, tensile and Charpy properties of weld metals by refined X120 wire, Mater. Sci. Eng. A 718 (2018) 350–362. [16] ASTM, Standard Specification for Structural Steel for Bridges, A709/A709M-16a, West Conshohocken, 2016. [17] Structural Steel for Bridge, GB/T 714, (2015), p. 16. [18] H.K.D.H. Bhadeshia, J.W. Christian, Bainite in steels, Metall. Trans. A 21 (3) (1990) 767–797. [19] Z. Xiong, S. Liu, X. Wang, C. Shang, X. Li, R.D.K. Misra, The contribution of intragranular acicular ferrite microstructural constituent on impact toughness and
4.2.2. Effect of grain boundary misorientation on the impact behavior in the CGHAZ Whether or not the microcracks propagation is deviated and/or arrested depends highly on the misorientations between the cracked grains and the adjacent grains. Some grain boundaries can suppress the propagation of microcracks, which is usually accompanied by the formation of a crack contusion or a small plastic deformation zone [4,19,33]. Figs. 13(b) and 10(b) show that the crack could not be deflected or be arrested at the LAGBs, but the HAGBs significantly inhibited the propagation of the secondary crack. Since the HAGB can change the propagation direction of a crack and effectively prevent its rapid growth, a high density of the HAGBs may help enhance the toughness of the CGHAZ. The density of the HAGBs should be, therefore, crucial in determining the impact toughness. According to Fig. 7(b), When the Si content increased from 0.12% to 0.48%–0.75%, the fraction of HAGBs decreased from 27% to 19%–10%. The distribution of grain misorientation in the prior austenite grain (PAG) became more uniform, as displayed in Fig. 13(b), owing to the increasing amount of the GB that normally has a larger equivalent grain size than the LB. The microcracks propagated almost in line towards the interior of the PAG until a deflection or an interception occurred at the PAG boundaries. As these obstacles encountered by crack propagation are reduced, and accordingly the energy consumed for crack propagation is lowered, which encourages crack propagation. This is the second major reason the toughness of CGHAZ was deteriorated obviously by the excessive addition of Si to the weathering steel. 5. Conclusions From our investigations into the effect of the Si content on the microstructures and the impact properties in the CGHAZ of typical weathering steel subjected to welding thermal simulation, the following major conclusions can be drawn: 1. A mixture of GB + AF + LB + M/A constituent forms in the CGHAZ of typical weathering steel with differing Si content. Increasing Si 9
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[20]
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Metall. Trans. A 10 (7) (1979) 895–907. [27] B. Dutta, E.J. Palmiere, C.M. Sellars, Modelling the kinetics of strain induced precipitation in Nb microalloyed steels, Acta Mater. 49 (5) (2001) 785–794. [28] S.H. Goods, L.M. Brown, Overview No. 1: the nucleation of cavities by plastic deformation, Acta Metall. 27 (1) (1979) 1–15. [29] A. Lambert-Perlade, T. Sturel, A.F. Gourgues, Mechanisms and modeling of cleavage fracture in simulated heat-affected zone microstructures of a high-strength low alloy steel, Metall. Mater. Trans. A 35 (3) (2004) 1039–1053. [30] E. Mazancová, K. Mazanec, Physical metallurgy characteristics of the M/A constituent formation in granular bainite, J. Mater. Process. Technol. 64 (1) (1997) 287–292. [31] X. Li, X. Ma, S.V. Subramanian, C. Shang, R.D.K. Misra, Influence of prior austenite grain size on martensite–austenite constituent and toughness in the heat affected zone of 700MPa high strength linepipe steel, Mater. Sci. Eng. A 616 (2014) 141–147. [32] C.L. Davis, J.E. King, Cleavage initiation in the intercritically reheated coarsegrained heat-affected zone: Part I. Fractographic evidence, Metall. Mater. Trans. A 25 (3) (1994) 563–573. [33] H.B. Cui, G.M. Xie, Z.A. Luo, J. Ma, G.D. Wang, R.D.K. Misra, The microstructural evolution and impact toughness of nugget zone in friction stir welded X100 pipeline steel, J. Alloy. Comp. 681 (2016) 426–433.
impeding crack initiation and propagation in the heat-affected zone (HAZ) of lowcarbon steels, Mater. Sci. Eng. A 636 (2015) 117–123. E. Bouyne, H.M. Flower, T.C. Lindley, A. Pineau, Use of EBSD technique to examine microstructure and cracking in a bainitic steel, Scripta Mater. 39 (3) (1998) 295–300. H. Terasaki, Y.-i. Komizo, Morphology and crystallography of bainite transformation in a single prior-austenite grain of low-carbon steel, Metall. Mater. Trans. A 44 (6) (2013) 2683–2689. B. Guo, L. Fan, Q. Wang, Z. Fu, Q. Wang, F. Zhang, The role of the bainitic packet in control of impact toughness in a simulated CGHAZ of X90 pipeline steel, Metals 6 (11) (2016). A. Lambert-Perlade, A.F. Gourgues, A. Pineau, Austenite to bainite phase transformation in the heat-affected zone of a high strength low alloy steel, Acta Mater. 52 (8) (2004) 2337–2348. S. Zhang, K. Liu, H. Chen, X. Xiao, Q. Wang, F. Zhang, Effect of increased N content on microstructure and tensile properties of low-C V-microalloyed steels, Mater. Sci. Eng. A 651 (2016) 951–960. L. Fan, D. Zhou, T. Wang, S. Li, Q. Wang, Tensile properties of an acicular ferrite and martensite/austenite constituent steel with varying cooling rates, Mater. Sci. Eng. A 590 (2014) 224–231. H.K.D.H. Bhadeshia, D.V. Edmonds, The bainite transformation in a silicon steel,
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Effect of Si content on the microstructures and the impact properties in the coarse-grained heat-affected zone (CGHAZ) of typical weathering steel
T
Yue Zhanga,b, Genhao Shia,b, Rui Suna,b, Kai Guoa,b, Chunling Zhanga,b,*, Qingfeng Wanga,b a b
State Key Laboratory of Metastable Materials Science and Technology, Yanshan University, 066004, Qinhuangdao, China National Engineering Research Center for Equipment and Technology of Cold Strip Rolling, Yanshan University, 066004, Qinhuangdao, China
A R T I C LE I N FO
A B S T R A C T
Keywords: Silicon content Toughness Coarse-grained heat-affected zone Weathering steel M/A constituents
This study aims to make clear the content limit of silicon (Si), which normally favors weathering steel but probably impairs its impact properties in the coarse-grained heat-affected zone (CGHAZ) induced by the conventional arc welding. Thermal simulation with a Gleeble 3500 simulator was thus carried out to reproduce the CGHAZ of this type of steel with differing silicon (Si) content from 0.12% to 0.48%–0.75%. Charpy V-notch impact tests and microstructure observations using optical microscopy, transmission and scanning electron microscopy (TEM and SEM), and electron backscatter diffraction (EBSD) were performed. The results showed that a microstructure composed of granular bainite (GB), acicular ferrite (AF), lath bainite (LB), and martensite/ austenite (M/A) constituents formed in 0.12%Si and 0.48%Si samples, while the LB scarcely appeared in 0.75% Si sample. When the Si content increased from 0.12% to 0.48%–0.75%, the LB decreased and in contrast, the GB increased. In addition, the area fraction (average size) of the M/A constituents increased from 5.8% (0.70 μm) to 9.2% (1.32 μm), while the number fraction (mean size) of effective grains with high-angle grain boundaries (HAGBs) decreased (increased) from 27% (3.77 μm) to 10% (5.71 μm). Accordingly, the Charpy impact energy of the CGHAZ decreased sharply from 208 to 42 J. The increased large M/A constituents significantly enable the microcrack nucleation, while the reduced HAGBs inefficiently impede the microcrack propagation, finally leading to a deteriorated toughness in the CGHAZ of 0.75%Si steel. The silicon content in this steel should not exceed 0.48%.
1. Introduction Typical weathering steel with an enhanced atmospheric corrosion resistance by the addition of Ni, Cr and Cu, has been developed over decades and extensively applicated for fabricating the uncoated building and bridge structures via the conventional arc welding process. This type of low-alloyed steel can normally obtain excellent comprehensive mechanical properties via thermo-mechanical controlled process. However, previous studies [1–4] reported that the welding-induced heat-affected zone can suffer from the inhomogeneity of microstructure and impact toughness. It is generally accepted that the worst toughness may frequently occur in the coarse-grained heat-affected zone (CGHAZ), owing to the martensite/austenite (M/A) constituents acting as a hard and brittle phase, the coarse prior austenite grain (PAG), and the correspondingly coarse granular banaite (GB) [5–7]. Some studies [8,9] have focused on how to enhance the poor impact properties in CGHAZ through optimizing the alloying element
and the corresponding content, including Ni, Cu and Nb, etc. Silicon (Si) is also an alloying element frequently added to the typical weathering steel, for it can enable the formation of a stable and compact rust layer and thereby enhance the atmospheric corrosion resistance [10,11]. Additionally, the Si can notably enhance the tensile strength of any steel through solid solution strengthening [12]. However, an excessive of Si content can also weaken the welding performance by degrading the impact toughness of CGHAZ [13]. Ruan et al. [14] reported that the decreasing Si content from 1.55% to 0.54% improves the toughness of Cr–Mn steel. The carbide particles can be refined by a lowered Si content, leading to an enhanced the low-temperature impact toughness. Taillard et al. [13] proved that the carbon enrichment in retained austenite of industrial E355 structural steel facilitated by the increasing Si content, significantly promotes the formation of M/A constituents, resulting a local brittle zone subjected to twice welding thermal cycles. Furthermore, a high Si content inhibits the decomposition of M/A constituents in the tempering process,
*
Corresponding author. State Key Laboratory of Metastable Materials Science and Technology, Yanshan University, 066004, Qinhuangdao, China. E-mail addresses: [email protected] (Y. Zhang), [email protected] (G. Shi), [email protected] (R. Sun), [email protected] (K. Guo), [email protected] (C. Zhang), [email protected] (Q. Wang). https://doi.org/10.1016/j.msea.2019.138082 Received 11 March 2019; Received in revised form 15 June 2019; Accepted 27 June 2019 Available online 28 June 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.
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leading to a lower toughness. Cao et al. [15] supposed that the increasing Si contents of X120 steel from 0.3% to 2.0% increase the amount fraction of GB, leaving the unaltered density of the high angle grain boundaries (HAGBs) in the CGHAZ [4]. The Si content limits of 50 W (0.30%–0.65% for Type A and 0.15%–0.50% for Type B), HPS 50W/70 W (0.30%–0.50%) and Q420qNH (0.15%–0.50%) weathering steels have been specified by ASTM A709/A709 M [16] and Chinese GB/T 714 [17] standards, respectively. It is necessary to understand the effect of Si content on the microstructures and impact properties in the welding-induced CGHAZs of these typical weathering steels, and the underlying mechanisms. This can help optimize the Si content. However, to our knowledge, the closely relevant research has not been reported yet. For this attempt to make clear the Si content limits due to the possible toughness loss in the CGHAZ of typical weathering steel, three experimental steels with the targeted/actual Si contents of 0.15%/ 0.12%, 0.50%/0.48% and 0.65%/0.75% were particularly designed/ prepared for welding thermal simulation tests. We investigated the correlation between the Si content, the microstructures, and the impact properties in the simulated CGHAZ. The microstructural features for each Si-bearing sample were systematically clarified. In particular, we focused on elucidating the influence of the Si content on the impact behavior and the mechanism of microcrack formation and propagation. Finally, an optimal range of Si content was proposed to improve the welding performance.
Fig. 1. Specimen orientation in the rolled plate and geometry. (RD-rolling direction; ND-normal direction; TD-transverse direction).
2. Experimental procedure The steels with the target Si content of 0.15%, 0.50%, and 0.65% chosen for the study were melted in a 50 kg vacuum furnace and hotrolled to 18-mm thick plates. The chemical compositions of three steels with the actual Si content of 0.12%, 0.48%, and 0.75% are listed in Table 1, which were described as 12Si, 48Si, and 75Si, respectively. Welding thermal cycle simulations were performed on a Gleeble3500 thermal cycle simulator for estimating the microstructure and impact property evolutions in the CGHAZ. The samples for simulation were cut along the rolling direction and processed into the rods of 10.5 × 10.5 × 80 mm3, as shown in Fig. 1. The thermal cyclying was conducted with the welding heat input of 20 kJ/cm, the mean heating rate of 100 °C/s and the peak temperature of 1320 °C for 1s, as shown in Fig. 2. Expansion curves were also measured and the onset/finish temperatures (Ar3/Ar1) for the austenite (γ) → ferrite (α) transformation were determined using the tangent method. To analyze the influence of Si addition on the microstructures and the Charpy impact properties, the simulated sample was cut along the location of the thermocouple. The samples obtained were mechanically polished and etched in a 4% Nital solution in volume. The morphology of the microstructure was characterized with an axiover-200mat metallographic microscope (OM) and an S3400 scanning electron microscope (SEM). To quantify the phase, the mean equivalent grain size, and the grain boundary misorientation, a solution of glycerin: perchloric acid: alcohol at a ratio of 0.5:1:8.5 was used to electropolish the samples. The samples were then observed with a Hitachi s-3400 scanning electron microscope using electron backscatter diffraction (EBSD) for orientation imaging microscopy (OIM) with the following parameters: a working distance of 15 mm, a step size of 0.3 μm, and acceleration voltage of 20 kV, and a tilt angle of 70°. The effective grains are regions
Fig. 2. The welding thermal cycle curves.
enclosed by low and high angle boundaries revealed by EBSD which was calculated as the equivalent circle diameter related to the individual grain area. Tolerance angles ranging from 2° to 30° are applied. Furthermore, the ferrite plates and M/A constituents were characterized using a JEM-2010 transmission electron microscope (TEM). The thin foils used for TEM were prepared by the twin-jet method and electro-polished in a 7% perchloric acid/glacial acetic acid solution at room temperature. The metallographic samples were etched with the Lepera reagent (4% picric acid alcohol solution: 1% sodium metabisulfite aqueous solution at a ratio of 1:1) to observe the M/A constituents. With this technique, the morphological distribution of the M/ A constituents (bright white) in the ferrite matrix (yellow) can be simply identified by OM. Subsequently, the image-pro plus software was used to perform a quantitative analysis of the images of the M/A constituents and at least 5 fields of view were measured for each sample at a magnification of 1000. The size of the M/A constituents was calculated as the equivalent diameter of the individual M/A constituent area. The average size was statistically obtained by measuring at least 1000 M/A constituents for each specimen. To conduct the Charpy impact tests, the thermal simulation samples were machined into normative Charpy v-notch samples with a
Table 1 Chemical composition of the test steels in wt. %. Steel
C
Si
Mn
P
S
Ni
Cr
Mo
Cu
Nb
Ti
12Si 48Si 75Si
0.053 0.058 0.058
0.12 0.48 0.75
1.25 1.27 1.27
0.007 0.007 0.008
0.002 0.005 0.002
0.31 0.32 0.31
0.51 0.52 0.52
0.05 0.05 0.05
0.26 0.25 0.27
0.028 0.028 0.027
0.013 0.014 0.016
2
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Fig. 3. Optical micrographs of the simulated CGHAZ samples of (a) 12Si, (b) 48Si and (c) 75Si steels, together with the M/A constituents observations of (d) 12Si, © 48Si and (f) 75Si samples, where the ferrite matrix is yellow and the M/A constituent is white. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
dimension of 10 × 10 × 55 mm3, as shown in Fig. 1. The Charpy V impact tests were conducted on a JB-300B impact tester at −40 °C and the average value was recorded. SEM was used to observe the fracture surfaces of the specimens and the sections perpendicular to the notch were used to characterize the second crack propagation by SEM and EBSD.
Table 2 Summary of the microstructure observations and quantifications.
3. Results
PAG—prior austenite grain; Dγ—mean size of PAG; M/A—martensite-austenite constituents; fM/A— area fraction of M/A constituents; dM/A— mean size of M/A constituents; MEDθ≥15°—mean equivalent diameter of ferrite grain with boundary misorientation angle θ > 15°; f (θ≥15°) —fraction of boundaries with misorientation angle θ ≥ 15°; pDFZ—area proportion of ductile fracture zone.
3.1. Microstructure Typical optical micrographs of each Si-containing CGHAZ sample are presented in Fig. 3. They show the microstructures on the left and the M/A constituents on the right. The quantitative analysis results are shown in Table 2. The mixed microstructure of granular bainite (GB), lath bainite (LB), acicular ferrite (AF), and martensite/austenite (M/A) constituent formed in 12Si and 48Si samples, while the LB scarcely appeared in 75Si sample. The microstructure observations showed a decrease in LB and an increase in GB when the Si content increased from 0.12% to 0.75%. In addition, the area fraction fM/A and the average size dM/A of the M/A constituents in CGHAZ increased from 5.8% to 9.2% and from 0.70 μm to 1.32 μm, respectively. Fig. 4 shows the size distribution of the M/A constituents varying with the increasing Si content, indicating that the M/A constituents with the size larger
Steel
Dγ/μm
fM/A/%
dM/A/μm
MEDθ≥15°/μm
f
(θ≥15°)/%
12Si 48Si 75Si
57.2 ± 2 62.8 ± 2 65.0 ± 3
5.8 ± 0.2 7.7 ± 0.3 9.2 ± 0.2
0.70 0.89 1.32
3.77 4.45 5.71
27 19 10
PDFZ/% 41.5 25.4 0.0
than 1 μm increased remarkably. Fig. 5 shows the typical microstructure morphology obtained by TEM. The TEM micrographs showed the formation of ferrite plates and island-like phases, as shown in Fig. 5(a) and (b). There are two types of bainitic ferrite distinguished by their morphologies: a plate or a block. The plate/block-like bainitic ferrite could be categorized as LBF/GBF [18]. In addition, typical M/A constituents in stripe and block shapes were observed in the bright field images (Fig. 5(c) and (f)), dark field images (Fig. 5(d) and (g)), and in the selected area diffraction pattern (Fig. 5(e) and (h)). They suggest the selected area was composed of martensite and austenite phases. The M/A constituents were dispersed 3
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between the bainitic ferrites in elongated (Fig. 5(c)) or blocky (Fig. 5(f)) shapes. Extensive TEM observations showed that the increasing Si content led to an increase in GB, a decrease in LB, and an increase of the area fraction and the average size of the M/A constituents, which was consistent with the OM observation results. The mean equivalent diameter (MED) of ferrite grains with boundaries defined by a misorientation tolerance angle (MTA) ranging 2o-60° for each sample was determined via EBSD, and the typical inverse pole figures are given in Fig. 6. When the MTA varied from 2° to 60° as a function of the MTA, the MED had a monotonic increasing trend with the increase of Si content, as shown in Fig. 7(a). In addition, the boundaries of ferrite grains could be divided into two categories: low angle grain boundary (LAGB, 2°≤MTA < 15°) and high angle grain boundary (HAGB, MTA≥15°). The white and black lines in Fig. 6 represented the LAGB and the HAGB, respectively. The HAGB can effectively impede crack propagation in a notched impact sample, according to Refs. [19], and hence the MED of grains with the HAGB can be considered as the effective grain size. Table 2 summarizes the number fraction of high-MTA, fθ≥15° boundary and the corresponding MEDθ≥15°. As shown in Table 2, both of them increased with the increase in Si content. For further clarification, Fig. 7(b) summarizes the
Fig. 4. Number density distribution of the M/A constituents vs. size for different Si contents.
Fig. 5. Typical TEM observations of the microstructure in CGHAZ samples of (a) 12Si and (b) 75Si steels, with the M/A constituents in stripe/block shape indicated particularly by (c/f) bright field, (d/g) dark field, and (e/h) the selected area diffraction pattern. GB- granular bainite; LB- lath bainite; M/A-martensite-austenite. 4
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Fig. 6. EBSD inverse pole figures for CGHAZ samples of (a) 12Si, (b) 48Si and (c) 75Si steels, with the white lines indicating LAGBs and the black lines indicating HAGBs.
Fig. 7. MED varied with MTA (a) and number fraction of grain boundary against misorientation angle for each Si-containing steel CGHAZ obtained by EBSD.
significantly, which might be associated to the evolution of the microstructures. The impact fracture surface of the CGHAZ for each sample with varying Si content was observed by SEM (Fig. 9). The blue line represented the margin of the v-shaped notch, and the yellow line indicated the boundary between the ductile fracture zone and the brittle fracture zone, while the red line showed the boundary between the brittle fracture zone and the hammer-impacted zone. According to the morphological characteristics of the fractures and the corresponding local positions, the complete fracture surfaces were of a mixed type where a ductile fracture zone and a brittle fracture zone coexist for 12Si
misorientation distribution for each sample. The statistics showed that the number fraction of HAGBs increased from 27% to 10% when the Si content increased from 0.12% to 0.75%. 3.2. Charpy impact toughness The Charpy impact energies determined at −40 °C for the simulated CGHAZ with different Si contents are shown in Fig. 8. The Charpy Vnotch (CVN) impact energy decreased sharply from 208 J to 42 J when the Si content increased from 0.12% to 0.75%. When the Si content reached 0.75%, the impact toughness of the test steels deteriorated 5
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constituents increased with the Si content, whereas the LB decreased. The evolution of the microstructure in CGHAZ had a close relationship with the Si content in steel, which can be explained as follows. 4.1.1. Effect of increased Si content on mixed LB + GB + AF microstructure The expansion curve for the CGHAZ of each Si-containing sample, which was subjected to the welding thermal cycle (Fig. 2), was measured and shown in Fig. 11. The temperature (Ar3) at which the phase transformation of γ→α started to take place on cooling, was determined and was also indicated in this figure. As Fig. 11 shows, the CGHAZs of 12Si and 48Si steels have the relatively lower Ar3 (586 °C and 599 °C) than that of 75Si steel, indicating that the phase change of γ→α proceeds at a relatively low onset temperature in these two samples. According to Ref. [22], the γ→GB transition normally precedes the γ→LB transition, and the GB also commonly nucleates and grows from the prior austenite grain boundary (PAGB). This enables the GB to constitute the leading phase of mixed GB + AF + LB microstructure. By contrast, the LB which mainly nucleates intragranularly according to Ref. [23] and grows in the retained space, accounts for a small proportion. This indicates that the GB could form with a higher thermodynamic driving force in the CGHAZ of 12Si and 48Si steels than the LB. These phase changes and products are in partly agreement with those of a low carbon Mo–V–Ti–N steel subjected to the controlled cooling at different rates [24]. Furthermore, with the Si content increasing to 0.75%, the Ar3 increased to 611 °C. The onset temperature for the phase change of γ→α was raised accordingly. This leads to an obviously increasing amount of the GBF at the expanse of LB. Also for this reason, the GB dominates, while the LB scarcely appears in the CGHAZ of 75Si steel. These discussions have clarified that increasing Si content facilitated the formation of GBF, and hence increased its fraction in the CGHAZ of weathering steel. The CGHAZs of 12Si and 48Si steels have the higher supercooling degree for the γ→α transformation than that of 75Si steel, owing to their lower Ar3. This may lead to a higher driving force and a higher nucleation rate of α in these two steels, which eventually result in the smaller equivalent grain size or MED at any MTA. Correspondingly, the effective grain size (MED MTA≥15°) for the CGHAZs of these two steels are smaller than that for 75Si steel, as indicated in Fig. 7(a) and Table 2. Meanwhile, the lowered Ar3 also promoted the formation, and accordingly led to an increasing fraction, of the thinner LB and probably the finer AF in the CGHAZs of 12Si and 48Si steels, as shown in Figs. 3 and Figs. 5. The boundaries of bainitic packet [22] and the AF [25] can additionally constitute the HAGBs in a bainite and ferrite dual-phase steel. This leads to the elevated number fraction of HAGBs for 12Si and 48Si steels in comparison with the 75Si steel, as indicated in Fig. 7(b) and Table 2, owing to the lowered Ar3 with the increasing Si content.
Fig. 8. Charpy impact energy for each sample varied with Si content.
and 48Si samples, but there was no ductile fracture zone on the fracture surface of 75Si sample. The area proportion of the ductile fracture zone of 12Si (41.5%) was much larger than that of 48Si (25.4%). The ductile fracture zone of 12Si and 48Si were comprised of a large number of cup-shaped and cone-shaped dimples, as shown in Fig. 9Ⅰ and Fig. 9Ⅲ. The brittle fracture zones of 12Si and 48Si were dominated by high density river-like patterns and tear ridges with small dimples, as seen in Fig. 9Ⅱ and Fig. 9Ⅳ. The junction of the cleavage facets by ductile fracture bands was considered as the main energy absorbing mechanism in the crack propagation process and kept the impact toughness at a high level [20]. As shown in Fig. 9(c) and Ⅴ, and Fig. 9Ⅵ for 75Si, the fracture was a whole cleavage fracture surface with large cleavage facets. Numerous large cleavage facets propagated in a linear manner and the straight river pattern extended from the cleavage crack to the periphery, forming a fracture step when encountering the highangle grain boundary and following a typical brittle fracture mechanism with low impact toughness. The apparent difference in fracture morphology between 12Si, 48Si and 75Si indicated that the impact toughness of the CGHAZ of 12Si sampe was better than that for 48Si sample, while 75Si sample showed the lowest impact toughness. To further analyze the influence of the microstructure on crack initiation and propagation, the secondary cracks underneath the fracture of the impact specimen were observed by SEM, as shown in Fig. 10. Numerous microvoids were observed in the matrix for the ductile fracture zone of 0.12% Si samples. The failure process of the dimple fracture was by accumulation of microvoids and their initiation. With the increasing Si content, the proportion of cleavage zone increased. The main fracture failure process went from the microvoids accumulation to a linear expansion of the microcracks. For 75Si sample (Fig. 10 (b)), the microcrack propagated through the interior of the prior austenite grain following a linear pattern until they were deflected or stopped at the prior austenite grain boundaries. As seen in Fig. 10, the microvoids and microcracks nucleated frequently at or around the smaller (< 1 μm, typically) and larger (≥1 μm, typically) M/A constituents, respectively, and propagated through the interface between the M/A constituents and the bainitic matrix. The measured impact energies showed a significant toughness degradation when the Si content varied from 0.12% to 0.75%, which was usually related to the distribution of the high angle grain boundaries in the matrix and the size/amount of the M/A constituents [21].
4.1.2. Effect of increased Si content on M/A constituents As described in Section 3.1, a mixed microstructure consisting of GB, AF and a small amount of LB and M/A constituents formed in the CGHAZs of 12Si and 48Si steels, while the GB + AF + M/A constituents in that of 75Si steel. The continuous cooling phase transformation of γ→α (bainitic ferrite and AF) + γ′ (metastable austenite) actually started from the Ar3 and finished at the Ar1, during which the γ/α interface normally moves toward to the γ′ phase. Since the solubility of carbon in the metastable austenite is different from that in ferrite [26], this phase transformation process is supposed to be accompanied by the re-distribution of carbon atoms from α to γ′ via bulk diffusion and dislocation-related pipe diffusion [27], leading to a C-rich γ′ with an accordingly enhanced stability. Fig. 12 schematically illustrates the carbon atom diffusion from bainitic and acicular ferrite to γ′ during the continuous cooling transformation. Since the Ar3 increased, and the Ar1 decreased with the Si addition, the residence time in the two-phase (α + γ) zone is extended under the same thermal cycle condition. The degree of C-enrichment in γ′ increases, leading to the
4. Discussion 4.1. Effect of the increased Si content on the microstructures During the continuous cooling process, a mixed GB + AF + LB + M/A microstructure formed in CGHAZ of each Sicontaining steel. As shown in Figs. 3 and 5, the GB and M/A 6
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Fig. 9. SEM micrographs of the fracture surface of 12Si (a), 48Si (b), and 75Si(c). Fig.9Ⅰ, Fig.9 Ⅱ, Fig.9Ⅲ, Fig.9Ⅳ, Fig.9Ⅴ, and Fig.9Ⅵ are the magnified SEM micrographs of the selected areas marked as Ⅰ, Ⅱ, Ⅲ, Ⅳ, Ⅴ, and Ⅵ.
further stabilization of γ'. In addition, Si can inhibit the precipitation of carbides from the super-cooled austenite [13] during this phase transformation, and keep the carbon solid solution in the γ′. This leads to the increase in residual austenite content at room temperature. Therefore,
the increasing Si content promotes the C-enrichment and accordingly the stabilization of γ′. The C-rich γ′ would further transform to the M/A constituents at the next continuous cooling stage and determine their morphology and inner-structure. The area fraction of the M/A 7
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Fig. 10. The crack path underneath the ductile fracture surface for (a) 0.12% Si sample and the brittle fracture surface for (b) 0.75% Si sample at the testing temperature of −40 °C.
Fig. 11. Expansion curve and Ar3 temperature measured for each CGHAZ sample with a different Si content.
Fig. 13. EBSD inverse pole figures of the secondary cracks underneath the ductile fracture surface for (a) 12Si sample and the brittle fracture surface for (b) 75Si sample.
could directly alter the impact behavior of the test steels. In the ductile fracture zone of the mixed fracture, the microvoids in the matrix tended to nucleate at carbides or at fine M/A constituents by means of particle fracture or stripping of the particle/matrix interface, as shown in Figs. 10(a) and Fig. 13(a). The microvoids grow via the plastic deformation of the surrounding matrix and their accumulation, which requires a large energy [4,28]. This helps to provide a high impact toughness. Compared with the ductile fracture zone, the cleavage fracture zone is more likely to cause the failure of the steel sample. The secondary crack in 75Si was taken as an example (Figs. 10(b) and Fig.13 (b)) to understand the crack propagation path of the microcracks in the cleavage zone in the CGHAZ of the test steel and the underlying mechanism. The influence of Si on microcrack initiation and propagation can be explained as follows.
Fig. 12. Schematic diagram of the diffusion of C atoms from the bainitic ferrite to the carbon-enriched austenite during the continuous cooling transformation.
constituents in the CGHAZ is ultimately increased, as shown in Fig. 3 and Table 2. Simultaneously, the intersecting area of adjacent GBF and AF plates, where the γ′ is located during this phase transformation, is enlarged with coarsening of these plates, owing to the Si addition and accordingly the elevated Ar3. The average size dM/A of M/A constituents in the CGHAZ is also increased, as indicated in Fig. 5 and Table 2.
4.2.1. Effect of the M/A constituents on the impact behavior in the CGHAZ On the basis of the results from Fig. 8, we conclude that the impact toughness of CGHAZ is closely associated with the number and the size of the M/A constituents in the matrix. The M/A constituents is supposed to be the brittle phase embedded in the ferrite matrix, and hence is one of the main factors determining the impact toughness [29,30]. Therefore, it is particularly important to understand the mirocrack nucleation and propagation assisted by the M/A constituents as displayed in Fig. 10(b). Usually due to the high hardness and brittleness of the M/A constituents, there is a significant difference in micro-hardness between the M/A constituents and the bainitic ferrite matrix [31]. Therefore, the
4.2. Effect of increased Si content on the impact behavior of the CGHAZ To further explore the mechanism of the influence of Si on the impact properties of the simulated CGHAZ, it is necessary to comprehensively consider the effect of the microstructural evolution on the impact toughness. When the Si content increased, the microstructures became coarser, the GB increased, and the LB decreased, while a relatively high area fraction and large size of M/A constituents ensues. These are the reasons why the impact energy remained low, which 8
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content results in the elevated Ar3 and hence, the decreased LB as well as the increased GB and M/A constituents. 2. Since the increasing Si content from 0.12% to 0.48%–0.75% leads to the obviously deteriorated impact properties of the CGHAZ, the silicon content in typical weathering steel should not exceed 0.48%. 3. Due to the increasing Si addition, both the increasing amount of larger M/A constituents and the decreasing fraction of HAGBs enables the microcrack propagation.
soft matrix structure in the sample for impact test usually undergoes an initial plastic deformation, resulting in the local stress concentration on or around the hard M/A constituents. The propagation of microcrack nucleating from the M/A constituents could occur if the locally concentrated stress exceeds the critical stress [32]. This critical stress can be described quantitatively by the classical Griffith local cracking model [1–4], given in equation (1):
σc =
2E γs 1/2 αa
(1)
Author contributions
where σc is the critical stress, E is the Young modulus, γs is the effective surface energy of the interface fracture between the M/A constituents and the bainitic matrix, α is a constant related to the crack shape, and a is the size of the microcrack nucleating at or around the M/A constituents. From the model, the smaller M/A constituents at or around which only a smaller microcrack might nucleate frequently, can cause an increase in the critical stress. Accordingly, this makes it difficult for the smaller microcrack to propagate. When the Si content is increased, the increasing amount of the larger M/A constituents can significantly promote the propagation of microcracks, which is primarily responsible for the lowered impact toughness. In the 12Si sample, the number fraction of the large M/A constituents with the dM/A ≥ 1 μm was as low as less than 50%, as indicated in Fig. 4. Then the nucleation and propagation of microcrack can get more difficult. By contract, a Si content as high as 0.75% turned out to be obviously detrimental to the toughness in the CGHAZ. This can be attributed to low difficulty for the microcrack to nucleate and propagate with respect to a high fraction of the large M/A constituents.
Chunling Zhang and Yue Zhang conceived and designed the experiments; Genhao Shi and Rui Sun performed the experiments; Kai Guo contributed analysis tools and analyzed the data; Qingfeng Wang directed the testing and editing; Chunling Zhang and Yue Zhang wrote the paper. Acknowledgments This work was support by the National Key Research and Development Program of China (Grant No. 2017YFB0304800 and Grant No. 2017YFB0304802 for the second sub project). References [1] L. Lan, C. Qiu, D. Zhao, X. Gao, L. Du, Microstructural characteristics and toughness of the simulated coarse grained heat affected zone of high strength low carbon bainitic steel, Mater. Sci. Eng. A 529 (2011) 192–200. [2] S.G. Lee, S.S. Sohn, B. Kim, W.G. Kim, K.-K. Um, S. Lee, Effects of martensite-austenite constituent on crack initiation and propagation in inter-critical heat-affected zone of high-strength low-alloy (HSLA) steel, Mater. Sci. Eng. A 715 (2018) 332–339. [3] L. Rancel, M. Gómez, S.F. Medina, I. Gutierrez, Measurement of bainite packet size and its influence on cleavage fracture in a medium carbon bainitic steel, Mater. Sci. Eng. A 530 (2011) 21–27. [4] J. Chen, S. Tang, Z.-Y. Liu, G.-D. Wang, Microstructural characteristics with various cooling paths and the mechanism of embrittlement and toughening in low-carbon high performance bridge steel, Mater. Sci. Eng. A 559 (2013) 241–249. [5] S. Kim, Y. Kang, C. Lee, Variation in microstructures and mechanical properties in the coarse-grained heat-affected zone of low-alloy steel with boron content, Mater. Sci. Eng. A 559 (2013) 178–186. [6] J. Chen, M.-y. Lv, Z.-y. Liu, G.-d. Wang, Combination of ductility and toughness by the design of fine ferrite/tempered martensite–austenite microstructure in a low carbon medium manganese alloyed steel plate, Mater. Sci. Eng. A 648 (2015) 51–56. [7] X. Zou, J. Sun, H. Matsuura, C. Wang, In situ observation of the nucleation and growth of ferrite laths in the heat-affected zone of EH36-Mg shipbuilding steel subjected to different heat inputs, Metall. Mater. Trans. B 49 (5) (2018) 2168–2173. [8] X.L. Wan, K.M. Wu, G. Huang, K.C. Nune, Y. Li, L. Cheng, Toughness improvement by Cu addition in the simulated coarse-grained heat-affected zone of high-strength low-alloy steels, Sci. Technol. Weld. Join. 21 (4) (2016) 295–302. [9] X. Zou, D. Zhao, J. Sun, C. Wang, H. Matsuura, An integrated study on the evolution of inclusions in EH36 shipbuilding steel with Mg addition: from casting to welding, Metall. Mater. Trans. B 49 (2) (2017) 481–489. [10] J.A. Mejía Gómez, J. Antonissen, C.A. Palacio, E. De Grave, Effects of Si as alloying element on corrosion resistance of weathering steel, Corros. Sci. 59 (2012) 198–203. [11] T. Nishimura, Corrosion behavior of silicon-bearing steel in a wet/dry environment containing chloride ions, Mater. Trans. 48 (6) (2007) 1438–1443. [12] W.M. Garrison, The effect of silicon and nickel additions on the sulfide spacing and fracture toughness of a 0.4 carbon low alloy steel, Metall. Trans. A (Phys. Metall. Mater. Sci.) 17 (4) (1986) 669–678. [13] R. Taillard, P. Verrier, T. Maurickx, J. Foct, Effect of silicon on CGHAZ toughness and microstructure of microalloyed steels, Metall. Mater. Trans. A 26 (2) (1995) 447–457. [14] L.H. Ruan, K.M. Wu, J.A. Qiu, A.A. Shirzadi, I.G. Rodionova, Effect of silicon content on carbide precipitation and low-temperature toughness of pressure vessel steels, Met. Sci. Heat Treat. 59 (1–2) (2017) 97–101. [15] R. Cao, Z.S. Chan, J.J. Yuan, C.Y. Han, Z.G. Xiao, X.B. Zhang, Y.J. Yan, J.H. Chen, The effects of Silicon and Copper on microstructures, tensile and Charpy properties of weld metals by refined X120 wire, Mater. Sci. Eng. A 718 (2018) 350–362. [16] ASTM, Standard Specification for Structural Steel for Bridges, A709/A709M-16a, West Conshohocken, 2016. [17] Structural Steel for Bridge, GB/T 714, (2015), p. 16. [18] H.K.D.H. Bhadeshia, J.W. Christian, Bainite in steels, Metall. Trans. A 21 (3) (1990) 767–797. [19] Z. Xiong, S. Liu, X. Wang, C. Shang, X. Li, R.D.K. Misra, The contribution of intragranular acicular ferrite microstructural constituent on impact toughness and
4.2.2. Effect of grain boundary misorientation on the impact behavior in the CGHAZ Whether or not the microcracks propagation is deviated and/or arrested depends highly on the misorientations between the cracked grains and the adjacent grains. Some grain boundaries can suppress the propagation of microcracks, which is usually accompanied by the formation of a crack contusion or a small plastic deformation zone [4,19,33]. Figs. 13(b) and 10(b) show that the crack could not be deflected or be arrested at the LAGBs, but the HAGBs significantly inhibited the propagation of the secondary crack. Since the HAGB can change the propagation direction of a crack and effectively prevent its rapid growth, a high density of the HAGBs may help enhance the toughness of the CGHAZ. The density of the HAGBs should be, therefore, crucial in determining the impact toughness. According to Fig. 7(b), When the Si content increased from 0.12% to 0.48%–0.75%, the fraction of HAGBs decreased from 27% to 19%–10%. The distribution of grain misorientation in the prior austenite grain (PAG) became more uniform, as displayed in Fig. 13(b), owing to the increasing amount of the GB that normally has a larger equivalent grain size than the LB. The microcracks propagated almost in line towards the interior of the PAG until a deflection or an interception occurred at the PAG boundaries. As these obstacles encountered by crack propagation are reduced, and accordingly the energy consumed for crack propagation is lowered, which encourages crack propagation. This is the second major reason the toughness of CGHAZ was deteriorated obviously by the excessive addition of Si to the weathering steel. 5. Conclusions From our investigations into the effect of the Si content on the microstructures and the impact properties in the CGHAZ of typical weathering steel subjected to welding thermal simulation, the following major conclusions can be drawn: 1. A mixture of GB + AF + LB + M/A constituent forms in the CGHAZ of typical weathering steel with differing Si content. Increasing Si 9
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