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The effect of alloy compositions on the joint strength of carbon steels welded by the extended hot-rolling process
Materials Science & Engineering A 558 (2012) 319–325
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
The effect of alloy compositions on the joint strength of carbon steels welded by the extended hot-rolling process Jin-Woo Park a,n, Chan-Woo Yang a, Youn-Hee Kang b, Jong-Sub Lee b a b
Department of Materials Science and Engineering, Yonsei University, Seoul 120-749, Korea Steel Solution Group, POSCO R&D Center, POSCO Co. Ltd., Pohang, Korea
a r t i c l e i n f o
abstract
Article history: Received 4 June 2012 Received in revised form 30 July 2012 Accepted 3 August 2012 Available online 8 August 2012
In this study, we investigated the factors that determine the joint strength of carbon steels welded by an extended joining process in air. Our main focus was on the effects of alloying elements (e.g., C, Si, Cr, and Mn, which are more reactive with O than with Fe) on scale adhesion, joint microstructure, and, ultimately, joint strength. Five carbon steels with different compositions of the same set of alloying elements are selected for this investigation. The model steels are hot-rolled, and surface descaling is performed several times during processing. The hot-rolled plates are welded with 1000% deformation at approximately 1000 1C. The joint strengths are evaluated by uniaxial tension testing. The joint microstructures and fractured surfaces are investigated with field emission scanning electron microscopy (FE-SEM). In addition, chemical analysis for phase identification is performed by electron probe X-ray micro-analysis (EPMA) and transmission electron microscopy (TEM). The results of our analyses reveal that Si increases the scale content at the joint by forming Fe2SiO4 along the interfaces between the Fe oxides and alloys and that this phase anchors the scale to the substrate. When both the C and Si contents are increased, Fe2SiO4 constitutes a larger portion of the scale–substrate interface, and the density of the internal oxide particles in the substrate near the joint is markedly increased. The tension test results demonstrate that the joint strength is decreased with increasing contents of C and Si because the joint failure is initiated by brittle fracture of the scale at the joint and the void nucleation sites around the internal oxides. & 2012 Elsevier B.V. All rights reserved.
Keywords: Extended hot rolling Scales Carbon steels Joint strength Adhesion Si
1. Introduction Endless hot rolling is an innovative process that combines thin slab casting, hot rolling, and solid-state joining to produce hotrolled strips [1]. In the extended hot-rolling process, metal slabs are joined before they are inserted between the finishing mills, such that the rolling step can be performed without pausing during the introduction of separate slabs, which is a feature of conventional hot rolling. Thus, extended hot rolling is significantly more efficient than conventional hot rolling because of the reduced lead time [2]. The mechanism of solid-state joining achieved during hot rolling is the result of the shear deformation of the overlapping regions of the two metal slabs; the shear deformation can be carried out to more than 1000% at an approximate processing temperature of 1000 1C [2]. Because the joining temperature typically exceeds 60% of the melting temperatures of the metal alloys used, the metal slabs tend to be significantly deformed during this process; this result is
n
Corresponding author. Fax: þ 82 2 312 5375. E-mail address: [email protected] (J.-W. Park).
0921-5093/$ - see front matter & 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.08.006
especially true for the mating surfaces, which are compressed to achieve an atomic-scale level of separation, allowing chemical bonding to occur at the interface. However, the duration of the compressive load applied to induce shear in the two mating surfaces is often less than a few seconds [2]. The strength of the joint made during extended hot rolling is highly affected by the processing conditions, such as the joining pressure and atmosphere, the temperature, and the area of the mating surfaces; these dependencies are frequently observed in joining processes. Nevertheless, in reality, the processing conditions are highly limited in variability [2,3]. For a fixed joining Table 1 Alloy compositions (in wt%) and fracture stresses (at 25 1C) of the model alloys. Sample alloy names
C
Si
Cr
Mn
Ni
Cu
Fe
Fracture stress (MPa)
LCS1 LCS2 MCS HCS1 HCS2
0.07 0.15 0.47 0.82 1.05
0.45 0.01 0.19 0.17 0.2
0.44 0.05 0.02 0.163 0.15
0.408 0.481 0.708 0.397 0.4
0.122 0.012 0.01 0.008 0.0001
0.284 bal. 500 0.011 450 0.006 650 0.008 900 0.0011 950
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condition, it has been reported that the ratio of the joint strength to the fracture strength of the alloy varies widely and is strongly dependent on the compositions of the alloys used [2].
Fig. 1. (a) Schematic description of the joining process used in the extended hotrolling process and (b) a description of the tensile test specimens of the model steel joints.
Additionally, fracture tends to occur along the interfaces of weaker joints [4,5]. The debonding of joints at low stresses is frequently observed in joints with poor interfacial adhesion, which is often the result of contamination at the mating surfaces. When a thick reaction layer forms along the bonding interface, brittle fracture can also occur at the joint because most intermetallic compounds are very brittle [5]. The extended hot-rolling process, which is carried out before the metal slabs are joined, comprises several major steps, including slab reheating, hot rolling and coiling. Both during and between each of these steps, the steel slabs are heated to the desired temperature by the combustion of natural gas in a reheat furnace. Because the atmosphere is oxidizing in nature, layers of oxide scales with thickness of up to several millimeters tend to form on the slab surfaces [6,7]. The oxide scale must be removed from the surface before the slabs are stacked and joined, particularly at the joining interface. Oxide scale removal, i.e., descaling, is typically performed with a hydraulic descaler using high-pressure water at each pass or at selective passes in the hot-rolling process [2]. In practice, however, the complete removal of the oxide scale is impossible. A thin scale layer or fragments of the original thick scale remain on the substrate surface, even after the surface has been subjected to multiple descaling passes, and the remaining scale is then included at the
Fig. 2. Optical microscopy (OM) images of the joints of (a) LCS1, (b) LCS2, (c) MCS, (d) HCS1 and (e) HCS2 with inset images at low magnification to show the joint interfaces.
J.-W. Park et al. / Materials Science & Engineering A 558 (2012) 319–325
joint interface [8]. Such scale inclusions can serve as initiation sites for brittle fractures. The thick reaction layer that occurs at a joint is one example of these scale inclusions. Additionally, the presence of altered surface compositions, which may result from either the formation of scales or the continued phase transformation of scales, has a tendency to further degrade an alloy’s weldability [2,9]. According to previous studies [10], a bandtype depletion zone of a specific element can occur along the interfaces between the scale and substrate surfaces. The respective microstructures of the joint interface and the region adjacent to the interface are both highly dependent on the microstructures and compositions of the alloying metals used [7]. However, there is little understanding of the joint microstructures that are formed during the extended hot-rolling process, especially with respect to the identities of the substrate–scale interfaces before joining. In this study, we investigated the fracture mechanisms of the two steel slab joints that are produced in the extended hot-rolling process; in particular, we focused on the interfacial microstructures of the joint regions in light of the major alloying elements added to balance the strength, ductility and high-temperature oxidation resistance of the steels. We selected five steels with different relative compositions of several key elements. The model steels were hot rolled, joined, and coiled in the extended hot-rolling pilot plant of the POSCO research laboratory. The joints of the model steels were cut into standard tensile test specimens, and uniaxial tension tests were carried out at room temperature to evaluate the joint strengths.
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The joint microstructures were analyzed by FE-SEM, EPMA, and TEM. The results of our analyses were then used to determine the effects of the alloying elements on the inclusion of scale, the microstructure and the strength of the joints.
2. Experimental procedures The compositions of the five model alloys are summarized in Table 1. The alloys labeled LCS1 and LCS2 in Table 1 are low carbon steels with C concentrations of 0.05–0.15 wt%. MCS is a medium carbon steel (o0.59 wt%), and HCS1 and HCS2 are both high carbon steels (40.6 wt%). LCS1 has the highest Si concentration of the five model alloys, and LCS2 has the lowest. The MCS and HCS alloys have the same Si content. The model alloys are first hot-rolled into thin slabs and then welded in the extended hot-rolling pilot plant of the POSCO research laboratory. The joining process induces a superdeformation of the two slabs and is schematically described in Fig. 1(a). The joint microstructures were analyzed using optical microscopy (OM) and FE-SEM. For microstructural analysis, the specimens are first etched in Nital. In particular, this analysis is focused on the amount of scale that remains on the slab surface following the descaling process, as well as the amount included at the joint of each sample. The interfaces between the scale inclusions and substrates are analyzed by EPMA and TEM. The phases of the
Fig. 3. FE-SEM images of the etched joints of (a) MCS, (b) HCS1 and (c) HCS2.
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scales are also identified with TEM. The specimens for TEM analysis are prepared by focused ion beam (FIB). For TEM analysis, selected area diffraction patterns (SADP) and energy-dispersive X-ray spectroscopy (EDS) are used to identify the oxide phases, which consist of the scale layers and other internal oxides within the substrates. The joint strengths are evaluated with a uniaxial tension test at 25 1C. The test specimens are described in Fig. 1(b). The strain rate is 1 mm/min, and 7–10 specimens are tested for each alloy, as shown in Table 1. Following the strain rate tests, the fracture surfaces are analyzed by FE-SEM, and an additional chemical analysis is performed with EDS.
3. Results and discussion Fig. 2 presents the OM images of the joint microstructures. As the C content is increased, the colors of the substrates under OM vary from white to dark gray, as illustrated in Fig. 2. The inset images in Fig. 2(a–e) shows the joints at low magnification ( 50), and the main images present typical scale inclusions observed along the joint interfaces of each alloy at higher magnifications (from 200 to 1000). Except for LCS2 in Fig. 2(b), all the joint regions are clearly visible. In the LCS1 joint, either a discrete thin layer of scale or small particle-type scale inclusions are visible. As shown in Fig. 2(b), the joint of LCS2 is almost invisible, except for a few scale inclusions that are sparsely distributed along the joint seam. According to Fig. 2(b), the grain dimensions (2.5 mm) at the scale interfaces are smaller than those of the grains in the substrate interior (12 mm). Recrystallization often occurs as a result of diffusion during and after the joining process [11].
The joints of the MCS, HCS1, and HCS2 specimens are clearly visible, as shown in the inset images of Fig. 2(c–e). Similar to the LCS1 joint, layer-type scale inclusions can be identified in discrete distributions along the joints of MCS, HCS1, and HCS2. However, these layers are significantly thicker than those in LCS1 and LCS2, based on the comparison of Fig. 2(a–e). Additionally, there are significantly more scale inclusions present in MCS, HCS1, and HCS2 than in LCS1 and LCS2. According to Fig. 2(c–e), the regions of the alloys that interface with the scale inclusions are whiter than the substrate bulk regions, indicating that compositional changes occurred across the joint–substrate interfaces [10,12]. According to previous publications by others [10,13] and by the current authors [12], the C content in the regions close to the surfaces of medium and high C steels are highly modified upon alloy oxidation. The surfaces of the alloys are often decarburized several microns deep as the C diffuses toward the metal surfaces and oxidizes into gaseous CO and CO2, which escape from the metal surface into the surrounding atmosphere. The decarburization continues until the alloy surface has been fully covered with an oxide layer. Because CO and CO2 are not soluble in iron oxides, the oxide forms at a higher rate than the C can diffuse and evaporate through the thick, continuous oxide surface layer, which forces the CO and CO2 to decompose and causes the decarburization rate to decay [13]. FE-SEM micrographs of the etched joints are presented in Fig. 3(a–c). These micrographs clearly present the carbon depletion zones, which measure several to tens of microns in thickness; additionally, these zones are pearlite-free near the interfaces. The thickness of the depletion zone increases with the alloy’s C content (Table 1). According to Fig. 3(b and c), the thickness of the depletion zone in one joint interface is greater than that in the other.
Fig. 4. EPMA element maps and results of the joints of (a) LCS1, (b) LCS2, (c) MCS, (d) HCS1 and (e) HCS2.
J.-W. Park et al. / Materials Science & Engineering A 558 (2012) 319–325
This observation implies that one of the two mating surfaces was better descaled than the other before joining was carried out. Based on previous analyses [12,13], for a fixed C content, the decarburization rate increases with increasing content of reactive elements such as Si and Cr, which are more readily oxidized than Fe is. These elements react with the O that diffuses into the substrate at the initial stage of surface oxidation, which is called internal oxidation. This process results in a steep gradient of the oxygen partial pressure from the surface into the substrate. Thus, the rate of oxygen dissolution is increased, and more C reacts with
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the O to form the gaseous phases (CO and CO2) at the initial oxidation stage of the alloys. The gaseous phases are not soluble in the alloys or the Fe oxides; hence, they must evaporate into the air. As the alloys continue to oxidize, the alloy surfaces become covered with layers of oxide. As mentioned above, the rate of decarburization decreases as the path to evaporation becomes increasingly blocked with the thickening oxide layers [10]. However, it has also been reported that defects and cracks are likely to form in the oxide layers of alloys with high contents of reactive
Fig. 5. TEM SADP and EDS results: (a) STEM images of MCS and (b) STEM and TEM images of HCS1.
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elements [7,12,13]. The gaseous phases diffuse through the cracks and defects into the atmosphere; otherwise, they decompose within the defects, which allows the decarburization to continue even though the reaction rate is decreased. According to the EPMA analysis, significant Si contents can be found in the scale or along the scale–substrate interfaces. Si concentration maps across the joints of the model alloys are Table 2 TEM EDS results of MCS (in wt%). Sample alloy name
Region number
O
C
Si
Cr
Mn
Fe
MCS
A 1 2 B 3 4 C
24.74 29.85 29.86 28.43 6.39 1.26 27.40
0.33 1.03 0.63 – 0.51 0.14 0.02
– 3.11 6.15 0.15 3.07 0.11 13.32
– 0.04 0.57 – 0.22 0.00 0.06
2.24 8.37 11.31 1.06 14.90 0.68 5.02
bal.
Table 3 TEM EDS results of HCS1 (in wt%). Sample alloy names
Site number
O
C
Si
Cr
Mn
Fe
HCS1
D E F 5 6 H 7 8 9
35.91 26.69 36.55 28.19 36.90 26.69 23.41 20.67 21.47
0.53 0.48 0.16 0.66 0.65 – – 0.20 0.07
15.54 0.25 14.08 9.99 14.68 13.15 0.32 0.47 5.0
– 0.30 – – 0.04 – 34.82 0.36 5.01
1.02 0.71 23.43 16.26 1.66 2.65 0.33 0.70 16.31
bal.
summarized in Fig. 4(a–e). All the model alloys in Table 1 have high Si contents, up to 2 wt%, except LCS2. In LCS1, the Sicontaining oxide particles are distributed throughout the scale inclusion layer at the joint, as demonstrated in Fig. 4(a). LCS2 has the lowest Si content of the model alloys examined (Table 1), and the Si-containing oxide cannot be found in the substrate or scale layers (Fig. 4(b)). Contrary to LCS1, the Si-containing oxides congregate along the interfaces between the MCS, HCS1, and HCS2 substrates and the scale layers, as shown in Fig. 4(c–e). Based on Fig. 4(c–e) and Table 1, the amount of Si oxide present along the interface appears to increase with the C content for an approximately fixed Si content (0.2 wt%). Fig. 5(a and b) presents scanning TEM (STEM) and regular TEM images of the scale inclusions at the joints of MCS and HCS1, respectively. The EDS results of the scales and particle-type internal oxides are summarized in Tables 2 and 3, respectively. According to the SADP analyses in Fig. 5(a and b), the scale inclusions are composed of Fe and Fe–Si oxides, such as FeO, Fe3O4, and Fe2SiO4. The particles in the substrate, i.e., the regions numbered and indicated by arrows in Fig. 5(a and b), are internal oxides whose major elements are Si and Mn, as presented in Tables 2 and 3. Based on the TEM results, it should be noted that Fe2SiO4 constitutes a larger portion of the scale inclusions in HCS1 than in MCS and that the density of internal oxide in HCS1 is significantly greater than that in MCS. Additionally, some of the internal oxides were found to be embedded in the scale–substrate interface and appear to anchor the scale to the substrate (Regions 1 and 2 in Fig. 5(a) and Regions 6 and 7 in Fig. 5(b)). According to the TEM analysis, more of these anchoring oxides can be found in HCS1 than in MCS. In light of the alloy compositions in Table 1, HCS2 will have a scale phase composition and a density and distribution of internal oxides that are similar to those in HCS1. Scales containing Si and/or Cr that form along interfaces are known to be highly adherent to steel substrates [14,15]. Hence, the large fractions of Fe2SiO4 in the scales and anchoring internal
Fig. 6. FE-SEM images of typical fractured surfaces (in MCS).
J.-W. Park et al. / Materials Science & Engineering A 558 (2012) 319–325
Table 4 EDS analysis of the selected fractured surface areas of S45C in Fig. 6 (in wt%). Area number
Fe
O
Si
Mn
I II III IV
89.9 100 65.3 65.62
11.1 – 20.81 15.21
– – 11.24 19.17
– – 2.67 –
325
Based on the fracture mechanism and microstructural analysis, the joints of HCS1 and HCS2 are weaker than those of the other alloys owing to the larger amounts of scale inclusions and internal oxides.
4. Conclusions In this study, we experimentally demonstrated that the strength of a carbon steel joint produced by extended hot rolling is highly dependent on the compositions of the alloys used. In weak joints, which contain high fractions of C and Si, joint failure can be observed along the joint interfaces. Interfacial fracture is initiated by the brittle fracture of scale inclusions and by the interconnection of the voids that have nucleated around the internal oxides. Si is known to be a more readily oxidized element than Fe, and Si oxides are highly adherent to Fe alloys. According to our analysis, Fe–Si oxides are increasingly formed along the scale–substrate interfaces that possess higher Si contents, which includes more Fe-oxide scales at the joint regions. The presence of brittle inclusions results in joint fracture at reduced stress. However, the joint strength is drastically decreased when the substrates include a large number of internal oxides, as observed in alloys with high contents of both C and Si. This finding is due to the rate of O dissolution throughout the substrate near the joint, which is increased with C and reacts with the O to form gaseous phases. The extended decarburization regions present near the C-containing interfaces confirm this fact.
Acknowledgments Fig. 7. Schematic representation of the fracture mechanism of the joints that failed along the interface.
oxides in HCS1 and HCS2 explain why the amounts of scale inclusions are greater in HCS1 and HCS2 than in the other model alloys in Table 1 despite the fact that identical descaling processes were undertaken before joining. In the uniaxial tension tests, fracturing of the LCS1 and LCS2 joints mostly occurs in the substrate. Hence, the fracture strengths of the joints are as high as those of the alloys (Table 1). However, fracturing of the MCS, HCS1, and HCS2 specimens occurs along the joint, and the fracture strengths are only 57%, 24%, and 36% of the failure strengths of the alloys (Table 1), respectively. The tension test results reveal that the joint strength is highly dependent on the C and Si contents, and the fracture surfaces presented in Fig. 6 depict the factors that contributed to the brittle joint fractures. Fig. 6(a–c) presents the typical fractured surfaces found in MCS, HCS1, and HCS2 (the images are the surfaces of MCS) as analyzed by FE-SEM. The fracture surfaces are either faceted or dimpled, as presented in Fig. 6(a). The results of the chemical analysis by EDS (Table 4) reveal that the faceted areas (Region I in Fig. 6(a and c)) are scale inclusions and that the dimpled area (Region II in Fig. 6(a and b)) is the substrate. As shown in Fig. 6(b) and Table 4, particles of Si-containing oxides can be identified inside the dimples. These particles must be the internal oxides observed by the EPMA and TEM results in Figs. 4 and 5. Based on the surface analysis presented in Fig. 6, joint fracture behavior is initiated by the brittle fracture of scale inclusions and by void nucleation and growth around the internal oxides. The fracture mechanism is schematically described in Fig. 7.
We acknowledge the POSCO R&D center for their financial support of this work. This work was also supported by a grant (M-2009-01-0014) from the Fundamental R&D Program for the Core Technology of Materials, which is funded by the Ministry of Knowledge Economy, South Korea. References [1] G. Arvedi, F. Mazzolari, J. Siegl, G. Hohenbichler, G. Holleis, Ironmak. Steelmak. 37 (2010) 271–275. [2] Y.-H. Kang, S.-H. Um, J.-S. Lee, J.-B. Lee, Y.-B. Gong, K. Horii, in: Korean Welding Society, Busan, Korea, 2006, pp. 117–119. [3] W.H. Lee, S.R. Lee, J. Korean Soc. Mech. Eng. A 23 (1999) 839–852. [4] T. Vigraman, D. Ravindran, R. Narayanasamy, Mater. Des. 34 (2012) 594–602. [5] W.H. Zhang, D.Q. Sun, L.J. Han, W.Q. Gao, X.M. Qiu, ISIJ Int. 51 (2011) 1870–1877. [6] R.Y. Chen, W.Y.D. Yuen, Oxid. Metals 53 (2000) 539–560. [7] Y.L. Yang, C.H. Yang, S.N. Lin, C.H. Chen, W.T. Tsai, Mater. Chem. Phys. 112 (2008) 566–571. [8] S. Taniguchi, K. Yamamoto, D. Megumi, T. Shibata, Mater. Sci. Eng. A 308 (2001) 250–257. [9] B.A. Baker, V.W. Hartmann, L.E. Shoemaker, S.A. McCoy, S. Rajendran, Trans. Indian Inst. Metals 56 (2003) 327–333. [10] D. Geneve, D. Rouxel, B. Weber, M. Confente, Mater. Sci. Eng. A 435–436 (2006) 1–11. [11] Y. Nakao, K. Shinozaki, M. Hamada, ISIJ Int. 31 (1991) 1260–1266. [12] C.W. Yang, S.M. Cho, Y.H. Kang, J.S. Lee, J.-W. Park, Mater. Sci. Eng. A, http://dx.doi.org/10.1016/j.msea.2012.06.082, in press. [13] D. Geneve, D. Rouxel, P. Pigeat, B. Weber, M. Confente, Appl. Surf. Sci. 254 (2008) 5348–5358. [14] M. Takeda, H. Kushida, T. Onishi, M. Toyama, F. Koizumi, S. Fujimoto, Oxid. Metals 73 (2010) 1–13. [15] T. Amano, K. Yamada, M. Okazaki, M. Takeda, T. Onishi, Mater. Sci. Forum 522–523 (2006) 451–460.
Contents lists available at SciVerse ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
The effect of alloy compositions on the joint strength of carbon steels welded by the extended hot-rolling process Jin-Woo Park a,n, Chan-Woo Yang a, Youn-Hee Kang b, Jong-Sub Lee b a b
Department of Materials Science and Engineering, Yonsei University, Seoul 120-749, Korea Steel Solution Group, POSCO R&D Center, POSCO Co. Ltd., Pohang, Korea
a r t i c l e i n f o
abstract
Article history: Received 4 June 2012 Received in revised form 30 July 2012 Accepted 3 August 2012 Available online 8 August 2012
In this study, we investigated the factors that determine the joint strength of carbon steels welded by an extended joining process in air. Our main focus was on the effects of alloying elements (e.g., C, Si, Cr, and Mn, which are more reactive with O than with Fe) on scale adhesion, joint microstructure, and, ultimately, joint strength. Five carbon steels with different compositions of the same set of alloying elements are selected for this investigation. The model steels are hot-rolled, and surface descaling is performed several times during processing. The hot-rolled plates are welded with 1000% deformation at approximately 1000 1C. The joint strengths are evaluated by uniaxial tension testing. The joint microstructures and fractured surfaces are investigated with field emission scanning electron microscopy (FE-SEM). In addition, chemical analysis for phase identification is performed by electron probe X-ray micro-analysis (EPMA) and transmission electron microscopy (TEM). The results of our analyses reveal that Si increases the scale content at the joint by forming Fe2SiO4 along the interfaces between the Fe oxides and alloys and that this phase anchors the scale to the substrate. When both the C and Si contents are increased, Fe2SiO4 constitutes a larger portion of the scale–substrate interface, and the density of the internal oxide particles in the substrate near the joint is markedly increased. The tension test results demonstrate that the joint strength is decreased with increasing contents of C and Si because the joint failure is initiated by brittle fracture of the scale at the joint and the void nucleation sites around the internal oxides. & 2012 Elsevier B.V. All rights reserved.
Keywords: Extended hot rolling Scales Carbon steels Joint strength Adhesion Si
1. Introduction Endless hot rolling is an innovative process that combines thin slab casting, hot rolling, and solid-state joining to produce hotrolled strips [1]. In the extended hot-rolling process, metal slabs are joined before they are inserted between the finishing mills, such that the rolling step can be performed without pausing during the introduction of separate slabs, which is a feature of conventional hot rolling. Thus, extended hot rolling is significantly more efficient than conventional hot rolling because of the reduced lead time [2]. The mechanism of solid-state joining achieved during hot rolling is the result of the shear deformation of the overlapping regions of the two metal slabs; the shear deformation can be carried out to more than 1000% at an approximate processing temperature of 1000 1C [2]. Because the joining temperature typically exceeds 60% of the melting temperatures of the metal alloys used, the metal slabs tend to be significantly deformed during this process; this result is
n
Corresponding author. Fax: þ 82 2 312 5375. E-mail address: [email protected] (J.-W. Park).
0921-5093/$ - see front matter & 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.08.006
especially true for the mating surfaces, which are compressed to achieve an atomic-scale level of separation, allowing chemical bonding to occur at the interface. However, the duration of the compressive load applied to induce shear in the two mating surfaces is often less than a few seconds [2]. The strength of the joint made during extended hot rolling is highly affected by the processing conditions, such as the joining pressure and atmosphere, the temperature, and the area of the mating surfaces; these dependencies are frequently observed in joining processes. Nevertheless, in reality, the processing conditions are highly limited in variability [2,3]. For a fixed joining Table 1 Alloy compositions (in wt%) and fracture stresses (at 25 1C) of the model alloys. Sample alloy names
C
Si
Cr
Mn
Ni
Cu
Fe
Fracture stress (MPa)
LCS1 LCS2 MCS HCS1 HCS2
0.07 0.15 0.47 0.82 1.05
0.45 0.01 0.19 0.17 0.2
0.44 0.05 0.02 0.163 0.15
0.408 0.481 0.708 0.397 0.4
0.122 0.012 0.01 0.008 0.0001
0.284 bal. 500 0.011 450 0.006 650 0.008 900 0.0011 950
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condition, it has been reported that the ratio of the joint strength to the fracture strength of the alloy varies widely and is strongly dependent on the compositions of the alloys used [2].
Fig. 1. (a) Schematic description of the joining process used in the extended hotrolling process and (b) a description of the tensile test specimens of the model steel joints.
Additionally, fracture tends to occur along the interfaces of weaker joints [4,5]. The debonding of joints at low stresses is frequently observed in joints with poor interfacial adhesion, which is often the result of contamination at the mating surfaces. When a thick reaction layer forms along the bonding interface, brittle fracture can also occur at the joint because most intermetallic compounds are very brittle [5]. The extended hot-rolling process, which is carried out before the metal slabs are joined, comprises several major steps, including slab reheating, hot rolling and coiling. Both during and between each of these steps, the steel slabs are heated to the desired temperature by the combustion of natural gas in a reheat furnace. Because the atmosphere is oxidizing in nature, layers of oxide scales with thickness of up to several millimeters tend to form on the slab surfaces [6,7]. The oxide scale must be removed from the surface before the slabs are stacked and joined, particularly at the joining interface. Oxide scale removal, i.e., descaling, is typically performed with a hydraulic descaler using high-pressure water at each pass or at selective passes in the hot-rolling process [2]. In practice, however, the complete removal of the oxide scale is impossible. A thin scale layer or fragments of the original thick scale remain on the substrate surface, even after the surface has been subjected to multiple descaling passes, and the remaining scale is then included at the
Fig. 2. Optical microscopy (OM) images of the joints of (a) LCS1, (b) LCS2, (c) MCS, (d) HCS1 and (e) HCS2 with inset images at low magnification to show the joint interfaces.
J.-W. Park et al. / Materials Science & Engineering A 558 (2012) 319–325
joint interface [8]. Such scale inclusions can serve as initiation sites for brittle fractures. The thick reaction layer that occurs at a joint is one example of these scale inclusions. Additionally, the presence of altered surface compositions, which may result from either the formation of scales or the continued phase transformation of scales, has a tendency to further degrade an alloy’s weldability [2,9]. According to previous studies [10], a bandtype depletion zone of a specific element can occur along the interfaces between the scale and substrate surfaces. The respective microstructures of the joint interface and the region adjacent to the interface are both highly dependent on the microstructures and compositions of the alloying metals used [7]. However, there is little understanding of the joint microstructures that are formed during the extended hot-rolling process, especially with respect to the identities of the substrate–scale interfaces before joining. In this study, we investigated the fracture mechanisms of the two steel slab joints that are produced in the extended hot-rolling process; in particular, we focused on the interfacial microstructures of the joint regions in light of the major alloying elements added to balance the strength, ductility and high-temperature oxidation resistance of the steels. We selected five steels with different relative compositions of several key elements. The model steels were hot rolled, joined, and coiled in the extended hot-rolling pilot plant of the POSCO research laboratory. The joints of the model steels were cut into standard tensile test specimens, and uniaxial tension tests were carried out at room temperature to evaluate the joint strengths.
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The joint microstructures were analyzed by FE-SEM, EPMA, and TEM. The results of our analyses were then used to determine the effects of the alloying elements on the inclusion of scale, the microstructure and the strength of the joints.
2. Experimental procedures The compositions of the five model alloys are summarized in Table 1. The alloys labeled LCS1 and LCS2 in Table 1 are low carbon steels with C concentrations of 0.05–0.15 wt%. MCS is a medium carbon steel (o0.59 wt%), and HCS1 and HCS2 are both high carbon steels (40.6 wt%). LCS1 has the highest Si concentration of the five model alloys, and LCS2 has the lowest. The MCS and HCS alloys have the same Si content. The model alloys are first hot-rolled into thin slabs and then welded in the extended hot-rolling pilot plant of the POSCO research laboratory. The joining process induces a superdeformation of the two slabs and is schematically described in Fig. 1(a). The joint microstructures were analyzed using optical microscopy (OM) and FE-SEM. For microstructural analysis, the specimens are first etched in Nital. In particular, this analysis is focused on the amount of scale that remains on the slab surface following the descaling process, as well as the amount included at the joint of each sample. The interfaces between the scale inclusions and substrates are analyzed by EPMA and TEM. The phases of the
Fig. 3. FE-SEM images of the etched joints of (a) MCS, (b) HCS1 and (c) HCS2.
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scales are also identified with TEM. The specimens for TEM analysis are prepared by focused ion beam (FIB). For TEM analysis, selected area diffraction patterns (SADP) and energy-dispersive X-ray spectroscopy (EDS) are used to identify the oxide phases, which consist of the scale layers and other internal oxides within the substrates. The joint strengths are evaluated with a uniaxial tension test at 25 1C. The test specimens are described in Fig. 1(b). The strain rate is 1 mm/min, and 7–10 specimens are tested for each alloy, as shown in Table 1. Following the strain rate tests, the fracture surfaces are analyzed by FE-SEM, and an additional chemical analysis is performed with EDS.
3. Results and discussion Fig. 2 presents the OM images of the joint microstructures. As the C content is increased, the colors of the substrates under OM vary from white to dark gray, as illustrated in Fig. 2. The inset images in Fig. 2(a–e) shows the joints at low magnification ( 50), and the main images present typical scale inclusions observed along the joint interfaces of each alloy at higher magnifications (from 200 to 1000). Except for LCS2 in Fig. 2(b), all the joint regions are clearly visible. In the LCS1 joint, either a discrete thin layer of scale or small particle-type scale inclusions are visible. As shown in Fig. 2(b), the joint of LCS2 is almost invisible, except for a few scale inclusions that are sparsely distributed along the joint seam. According to Fig. 2(b), the grain dimensions (2.5 mm) at the scale interfaces are smaller than those of the grains in the substrate interior (12 mm). Recrystallization often occurs as a result of diffusion during and after the joining process [11].
The joints of the MCS, HCS1, and HCS2 specimens are clearly visible, as shown in the inset images of Fig. 2(c–e). Similar to the LCS1 joint, layer-type scale inclusions can be identified in discrete distributions along the joints of MCS, HCS1, and HCS2. However, these layers are significantly thicker than those in LCS1 and LCS2, based on the comparison of Fig. 2(a–e). Additionally, there are significantly more scale inclusions present in MCS, HCS1, and HCS2 than in LCS1 and LCS2. According to Fig. 2(c–e), the regions of the alloys that interface with the scale inclusions are whiter than the substrate bulk regions, indicating that compositional changes occurred across the joint–substrate interfaces [10,12]. According to previous publications by others [10,13] and by the current authors [12], the C content in the regions close to the surfaces of medium and high C steels are highly modified upon alloy oxidation. The surfaces of the alloys are often decarburized several microns deep as the C diffuses toward the metal surfaces and oxidizes into gaseous CO and CO2, which escape from the metal surface into the surrounding atmosphere. The decarburization continues until the alloy surface has been fully covered with an oxide layer. Because CO and CO2 are not soluble in iron oxides, the oxide forms at a higher rate than the C can diffuse and evaporate through the thick, continuous oxide surface layer, which forces the CO and CO2 to decompose and causes the decarburization rate to decay [13]. FE-SEM micrographs of the etched joints are presented in Fig. 3(a–c). These micrographs clearly present the carbon depletion zones, which measure several to tens of microns in thickness; additionally, these zones are pearlite-free near the interfaces. The thickness of the depletion zone increases with the alloy’s C content (Table 1). According to Fig. 3(b and c), the thickness of the depletion zone in one joint interface is greater than that in the other.
Fig. 4. EPMA element maps and results of the joints of (a) LCS1, (b) LCS2, (c) MCS, (d) HCS1 and (e) HCS2.
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This observation implies that one of the two mating surfaces was better descaled than the other before joining was carried out. Based on previous analyses [12,13], for a fixed C content, the decarburization rate increases with increasing content of reactive elements such as Si and Cr, which are more readily oxidized than Fe is. These elements react with the O that diffuses into the substrate at the initial stage of surface oxidation, which is called internal oxidation. This process results in a steep gradient of the oxygen partial pressure from the surface into the substrate. Thus, the rate of oxygen dissolution is increased, and more C reacts with
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the O to form the gaseous phases (CO and CO2) at the initial oxidation stage of the alloys. The gaseous phases are not soluble in the alloys or the Fe oxides; hence, they must evaporate into the air. As the alloys continue to oxidize, the alloy surfaces become covered with layers of oxide. As mentioned above, the rate of decarburization decreases as the path to evaporation becomes increasingly blocked with the thickening oxide layers [10]. However, it has also been reported that defects and cracks are likely to form in the oxide layers of alloys with high contents of reactive
Fig. 5. TEM SADP and EDS results: (a) STEM images of MCS and (b) STEM and TEM images of HCS1.
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elements [7,12,13]. The gaseous phases diffuse through the cracks and defects into the atmosphere; otherwise, they decompose within the defects, which allows the decarburization to continue even though the reaction rate is decreased. According to the EPMA analysis, significant Si contents can be found in the scale or along the scale–substrate interfaces. Si concentration maps across the joints of the model alloys are Table 2 TEM EDS results of MCS (in wt%). Sample alloy name
Region number
O
C
Si
Cr
Mn
Fe
MCS
A 1 2 B 3 4 C
24.74 29.85 29.86 28.43 6.39 1.26 27.40
0.33 1.03 0.63 – 0.51 0.14 0.02
– 3.11 6.15 0.15 3.07 0.11 13.32
– 0.04 0.57 – 0.22 0.00 0.06
2.24 8.37 11.31 1.06 14.90 0.68 5.02
bal.
Table 3 TEM EDS results of HCS1 (in wt%). Sample alloy names
Site number
O
C
Si
Cr
Mn
Fe
HCS1
D E F 5 6 H 7 8 9
35.91 26.69 36.55 28.19 36.90 26.69 23.41 20.67 21.47
0.53 0.48 0.16 0.66 0.65 – – 0.20 0.07
15.54 0.25 14.08 9.99 14.68 13.15 0.32 0.47 5.0
– 0.30 – – 0.04 – 34.82 0.36 5.01
1.02 0.71 23.43 16.26 1.66 2.65 0.33 0.70 16.31
bal.
summarized in Fig. 4(a–e). All the model alloys in Table 1 have high Si contents, up to 2 wt%, except LCS2. In LCS1, the Sicontaining oxide particles are distributed throughout the scale inclusion layer at the joint, as demonstrated in Fig. 4(a). LCS2 has the lowest Si content of the model alloys examined (Table 1), and the Si-containing oxide cannot be found in the substrate or scale layers (Fig. 4(b)). Contrary to LCS1, the Si-containing oxides congregate along the interfaces between the MCS, HCS1, and HCS2 substrates and the scale layers, as shown in Fig. 4(c–e). Based on Fig. 4(c–e) and Table 1, the amount of Si oxide present along the interface appears to increase with the C content for an approximately fixed Si content (0.2 wt%). Fig. 5(a and b) presents scanning TEM (STEM) and regular TEM images of the scale inclusions at the joints of MCS and HCS1, respectively. The EDS results of the scales and particle-type internal oxides are summarized in Tables 2 and 3, respectively. According to the SADP analyses in Fig. 5(a and b), the scale inclusions are composed of Fe and Fe–Si oxides, such as FeO, Fe3O4, and Fe2SiO4. The particles in the substrate, i.e., the regions numbered and indicated by arrows in Fig. 5(a and b), are internal oxides whose major elements are Si and Mn, as presented in Tables 2 and 3. Based on the TEM results, it should be noted that Fe2SiO4 constitutes a larger portion of the scale inclusions in HCS1 than in MCS and that the density of internal oxide in HCS1 is significantly greater than that in MCS. Additionally, some of the internal oxides were found to be embedded in the scale–substrate interface and appear to anchor the scale to the substrate (Regions 1 and 2 in Fig. 5(a) and Regions 6 and 7 in Fig. 5(b)). According to the TEM analysis, more of these anchoring oxides can be found in HCS1 than in MCS. In light of the alloy compositions in Table 1, HCS2 will have a scale phase composition and a density and distribution of internal oxides that are similar to those in HCS1. Scales containing Si and/or Cr that form along interfaces are known to be highly adherent to steel substrates [14,15]. Hence, the large fractions of Fe2SiO4 in the scales and anchoring internal
Fig. 6. FE-SEM images of typical fractured surfaces (in MCS).
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Table 4 EDS analysis of the selected fractured surface areas of S45C in Fig. 6 (in wt%). Area number
Fe
O
Si
Mn
I II III IV
89.9 100 65.3 65.62
11.1 – 20.81 15.21
– – 11.24 19.17
– – 2.67 –
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Based on the fracture mechanism and microstructural analysis, the joints of HCS1 and HCS2 are weaker than those of the other alloys owing to the larger amounts of scale inclusions and internal oxides.
4. Conclusions In this study, we experimentally demonstrated that the strength of a carbon steel joint produced by extended hot rolling is highly dependent on the compositions of the alloys used. In weak joints, which contain high fractions of C and Si, joint failure can be observed along the joint interfaces. Interfacial fracture is initiated by the brittle fracture of scale inclusions and by the interconnection of the voids that have nucleated around the internal oxides. Si is known to be a more readily oxidized element than Fe, and Si oxides are highly adherent to Fe alloys. According to our analysis, Fe–Si oxides are increasingly formed along the scale–substrate interfaces that possess higher Si contents, which includes more Fe-oxide scales at the joint regions. The presence of brittle inclusions results in joint fracture at reduced stress. However, the joint strength is drastically decreased when the substrates include a large number of internal oxides, as observed in alloys with high contents of both C and Si. This finding is due to the rate of O dissolution throughout the substrate near the joint, which is increased with C and reacts with the O to form gaseous phases. The extended decarburization regions present near the C-containing interfaces confirm this fact.
Acknowledgments Fig. 7. Schematic representation of the fracture mechanism of the joints that failed along the interface.
oxides in HCS1 and HCS2 explain why the amounts of scale inclusions are greater in HCS1 and HCS2 than in the other model alloys in Table 1 despite the fact that identical descaling processes were undertaken before joining. In the uniaxial tension tests, fracturing of the LCS1 and LCS2 joints mostly occurs in the substrate. Hence, the fracture strengths of the joints are as high as those of the alloys (Table 1). However, fracturing of the MCS, HCS1, and HCS2 specimens occurs along the joint, and the fracture strengths are only 57%, 24%, and 36% of the failure strengths of the alloys (Table 1), respectively. The tension test results reveal that the joint strength is highly dependent on the C and Si contents, and the fracture surfaces presented in Fig. 6 depict the factors that contributed to the brittle joint fractures. Fig. 6(a–c) presents the typical fractured surfaces found in MCS, HCS1, and HCS2 (the images are the surfaces of MCS) as analyzed by FE-SEM. The fracture surfaces are either faceted or dimpled, as presented in Fig. 6(a). The results of the chemical analysis by EDS (Table 4) reveal that the faceted areas (Region I in Fig. 6(a and c)) are scale inclusions and that the dimpled area (Region II in Fig. 6(a and b)) is the substrate. As shown in Fig. 6(b) and Table 4, particles of Si-containing oxides can be identified inside the dimples. These particles must be the internal oxides observed by the EPMA and TEM results in Figs. 4 and 5. Based on the surface analysis presented in Fig. 6, joint fracture behavior is initiated by the brittle fracture of scale inclusions and by void nucleation and growth around the internal oxides. The fracture mechanism is schematically described in Fig. 7.
We acknowledge the POSCO R&D center for their financial support of this work. This work was also supported by a grant (M-2009-01-0014) from the Fundamental R&D Program for the Core Technology of Materials, which is funded by the Ministry of Knowledge Economy, South Korea. References [1] G. Arvedi, F. Mazzolari, J. Siegl, G. Hohenbichler, G. Holleis, Ironmak. Steelmak. 37 (2010) 271–275. [2] Y.-H. Kang, S.-H. Um, J.-S. Lee, J.-B. Lee, Y.-B. Gong, K. Horii, in: Korean Welding Society, Busan, Korea, 2006, pp. 117–119. [3] W.H. Lee, S.R. Lee, J. Korean Soc. Mech. Eng. A 23 (1999) 839–852. [4] T. Vigraman, D. Ravindran, R. Narayanasamy, Mater. Des. 34 (2012) 594–602. [5] W.H. Zhang, D.Q. Sun, L.J. Han, W.Q. Gao, X.M. Qiu, ISIJ Int. 51 (2011) 1870–1877. [6] R.Y. Chen, W.Y.D. Yuen, Oxid. Metals 53 (2000) 539–560. [7] Y.L. Yang, C.H. Yang, S.N. Lin, C.H. Chen, W.T. Tsai, Mater. Chem. Phys. 112 (2008) 566–571. [8] S. Taniguchi, K. Yamamoto, D. Megumi, T. Shibata, Mater. Sci. Eng. A 308 (2001) 250–257. [9] B.A. Baker, V.W. Hartmann, L.E. Shoemaker, S.A. McCoy, S. Rajendran, Trans. Indian Inst. Metals 56 (2003) 327–333. [10] D. Geneve, D. Rouxel, B. Weber, M. Confente, Mater. Sci. Eng. A 435–436 (2006) 1–11. [11] Y. Nakao, K. Shinozaki, M. Hamada, ISIJ Int. 31 (1991) 1260–1266. [12] C.W. Yang, S.M. Cho, Y.H. Kang, J.S. Lee, J.-W. Park, Mater. Sci. Eng. A, http://dx.doi.org/10.1016/j.msea.2012.06.082, in press. [13] D. Geneve, D. Rouxel, P. Pigeat, B. Weber, M. Confente, Appl. Surf. Sci. 254 (2008) 5348–5358. [14] M. Takeda, H. Kushida, T. Onishi, M. Toyama, F. Koizumi, S. Fujimoto, Oxid. Metals 73 (2010) 1–13. [15] T. Amano, K. Yamada, M. Okazaki, M. Takeda, T. Onishi, Mater. Sci. Forum 522–523 (2006) 451–460.