- Email: [email protected]
Orientation relationships and interfaces of the Fe–Zn intermetallic phases in galvannealed CMnSi-TRIP steels
Materials Characterization 107 (2015) 23–28
Contents lists available at ScienceDirect
Materials Characterization journal homepage: www.elsevier.com/locate/matchar
Orientation relationships and interfaces of the Fe–Zn intermetallic phases in galvannealed CMnSi-TRIP steels Kuang-Kuo Wang a, Guan-Lin You a, Liuwen Chang a,⁎, Dershin Gan a, Lung-Jen Chiang b a b
Research Center for Physical Properties and Microstructure of Metals, Department of Materials and Optoelectronic Science, National Sun Yat-Sen University, Kaohsiung 80424, Taiwan, ROC Steel Research and Development Department, China Steel Corporation, Kaohsiung 81233, Taiwan, ROC
a r t i c l e
i n f o
Article history: Received 19 March 2015 Received in revised form 10 June 2015 Accepted 30 June 2015 Available online 2 July 2015 Keywords: Galvannealed coating Mn–Si TRIP steels Fe–Zn intermetallic Orientation relationship Interface
a b s t r a c t Hot-dip galvannealed coatings of two Mn–Si TRIP steels have been characterized by transmission electron microscopy to study the crystallographic orientation relationships between the α-Fe and Γ1 phases, the Γ1 and Γ2 phases, as well as the Γ2 and δ phases. Both the α-Fe/Γ1 phases and the Γ1/Γ2 phases were found to exhibit a cube-on-cube orientation relationship. The orientation relationship at the Γ2/δ interface was identified to be ½1120 δ//[011]Γ2 and (0002) δ//ð111Þ Γ2. The interfaces of Γ2/δ and Γ1/Γ2 are basically the close-packed plane of each phase. The nucleation of Γ2 phase between the δ and Γ1 phases was discussed by the heterogeneous nucleation theory based on the orientation relationships and the interfaces. © 2015 Elsevier Inc. All rights reserved.
1. Introduction Hot-dip galvannealed (GA) steel sheets have been extensively used in car manufacture. An optimized GA coating consists of two consecutive layers of Γ1 and δ on the steel matrix, though totally four Fe–Zn intermetallic compounds (IMCs), Γ1, Γ2, δ and ζ, exist below 803 K according to the equilibrium phase diagram [1]. The symbols and crystallographic information of the Fe–Zn IMCs is listed in Table 1 for reference [2–4]. It was reported that the ζ phase is the first phase to form at the Fe/Zn interface followed by the δ phase and finally the Γ1 and/or Γ2 phases on annealing at or below 773 K for Zn coating containing 0.1 wt.% Al or less [5–7]. However, the Fe/Zn reactions are inhibited temporarily by the formation of an inhibition layer of Fe2Al5 − xZnx in hotdip galvanizing if some Al is added to the Zn bath. An additional annealing at 773–813 K is thus necessary to accelerate the interdiffusion across the inhibition layer for the occurrence of the Fe/Zn reactions. The δ, instead of the ζ, phase was suggested to be the first IMC to form at the Fe2Al5 − xZnx/α-Fe interface followed by the Γ1 phase at the δ/α-Fe interface [8–10]. The formation kinetics of the Fe–Zn IMCs attracts great attentions due to its importance on optimizing the microstructure and the properties of the GA coating. One of the important issues which have not been clarified yet is whether the IMCs are formed epitaxially to the ferrite grains and the prior-formed IMCs. Adachi and Kamei [11] argued that the Γ1/α-Fe interface should be coherent or partly coherent due to the ⁎ Corresponding author. E-mail address: [email protected] (L. Chang).
http://dx.doi.org/10.1016/j.matchar.2015.06.041 1044-5803/© 2015 Elsevier Inc. All rights reserved.
presence of growth ledge at the interface. Various orientation relationships (ORs) have been proposed between the α-Fe and Γ1 [12,13], αFe and Γ2 [13] and α-Fe and ζ phases [7]. However, the OR between the α-Fe and Γ1 phases derived from the electrodeposited sample is different from that from the hot-dip sample. The crystallographic characteristics of the Fe–Zn reactions are difficult to be revealed because the reactions are so rapid and sometimes locally intense that the inherent ORs do not show up. Accordingly, many cross-sectional TEM works have been conducted recently [14–16], but no results on the ORs and the interfaces of the Fe–Zn intermetallic phases have been reported. In addition, Cho et al. [17] reported that no ORs were found between the α-Fe and Γ1, Γ1 and Γ2, as well as Γ2 and δ phases during the galvannealing of a Mn–Si TRIP steel. In this work, we conducted transmission electron microscopy (TEM) analysis on two galvannealed samples of Mn–Si transformation induced plasticity (TRIP) steels. It is known that the oxides formed by selective oxidation on the surface of the Mn–Si TRIP steels during intercritical annealing can inhibit the Fe–Zn reactions on subsequent annealing. By changing the annealing conditions, various reaction results can therefore be observed. In this work, we report, for the first time, the ORs between the consecutive pairs of α-Fe/Γ1, Γ1/Γ2, and Γ2/δ phases. The crystallographic data can provide valuable information for further discussion of the nucleation and growth kinetics of the Fe–Zn IMCs. 2. Experimental GA coatings on two TRIP steels containing 1.0 wt.% Si-2.0 wt.% Mn (steel A) and 1.8 wt.% Si-1.9 wt.% Mn (steel B), respectively, were
24
K.-K. Wang et al. / Materials Characterization 107 (2015) 23–28
Table 1 Crystallographic characteristics of the Fe–Zn phases [2–4]. Phase
Formula
Crystal structure
Space group
Lattice parameters (nm)
ζ
FeZn13–15
Monoclinic
C2/m
δ (δk/δp) Γ2
FeZn7–10 Fe11Zn40
Hexagonal FCC
P63/mmc
a = 1.3424, b = 0.7608 c = 0.5061, β = 127.3° a = 1.2787, c = 5.7222 a = 1.7963
Γ1
Fe3Zn10
BCC
I43m
a = 0.9018
α-Fe
Fe
BCC
Im3m
a = 0.2866
F43m
TEM (FEI Tecnai G2, at 200 kV). After galvannealing, the central part (50 × 50 mm2) of the samples were cut into 10 × 10 mm2 coupons, cleaned in ethanol, and loaded in an FEI Nova-200 dual beam instrument. The TEM cross-sectional sample was prepared by milling an electron-transparent membrane of approximately 10 μm × 10 μm in size out of the desired region of the bulk GA sample. The crosssectional sample was then transferred to an Omniprobe grid and further thinned to a thickness of 60 nm or less to enable high-resolution imaging in TEM. 3. Results and discussion
studied. The steels were prepared by vacuum induction melting. The ingots were reheated and soaked at 1523 K, and hot rolled to approximately 5 mm thick with a finish-rolling temperature of about 1173 K. The strip was rapidly cooled to about 773 K and then soaked immediately at 823 K for 2 h. Afterwards, the hot band was machined to 4 mm thick to remove the surface scale and cold rolled to 1.0 mm thick. The cold-rolled sheets were cut into samples of 120 mm × 250 mm for simulating the continuous annealing and hot-dip galvannealing process using an Iwatani hot-dip simulator. Prior to annealing, the samples were cleaned by dipping in a 2% alkaline solution for 20 min at 343 K, mechanical brushing, water-spray rinsing and hot-air drying in sequence. The steels were first annealed at 1073 K for 60 s in a 10%H2–N2 atmosphere with a dew point (dp) of 243 K (PO2 = 7.24 × 10− 24 atm) or 273 K (PO2 = 1.87 × 10− 21 atm), cooled to 733 K at a rate of 15 K/s after soaking, held for 60 s, and dipped in a Zn bath for 3 s. The Al content in the bath was 0.16 wt.% with saturated Fe. The high Al content ensured the formation of the Fe–Al layer with a high coverage ratio at the Fe–Zn interface in galvanizing which provides a sufficient wettability to the re-melted Zn in galvannealing. High quality GA coatings were obtained accordingly [18]. The samples were finally galvannealed at 793–833 K for 20 s in an induction furnace of the simulator after hot-dipping. The GA coatings were analyzed by scanning electron microscopy (SEM, Zeiss Supra 55, at 5–15 kV) and
The effects of processing parameters and steel composition on the galvanizing and galvannealing characteristics of the Mn–Si TRIP steels used in this study were reported previously [18,19]. In this work we focus on the ORs at the various interfaces. Fig. 1 shows the SEM secondary electron images of sample A (steel A, annealed at dp = 273 K and galvannealed at 793 K for 20 s) and sample B (steel B, annealed at dp = 243 K and galvannealed at 833 K for 20 s), using two extreme GA conditions. Fig. 1a shows that sample A exhibits a relatively flat surface on which Fe–Zn grains of less than 3 μm in size are observed. The cross-sectional micrograph in Fig. 1b verifies that the GA coating is uniform in thickness with a thin inner layer (1–1.5 μm) and a thick outer layer. The average Fe content of the coating determined by energy dispersive spectroscopy is 10.7 wt.%, which is within the typical GA region of 9–11 wt.%. Coating B, however, exhibits a rough surface scattered with coarse IMC particles as large as 20 μm (Fig. 1c). The crosssectional micrograph in Fig. 1d shows that the coating consists of a thin layer less than 2 μm thick and the large particles on it. Furthermore, the thin layer is not continuous at the interface. Accordingly, sample A represents a well-optimized GA coating but sample B is not. The results agree relatively well with a recent study carried out for a 2.2Mn–1.4Si TRIP steel [17]. Because we will focus on the ORs and interfaces of the various phases, the complex growth mechanisms of the Fe–Zn phases are out of the scope of this paper and will be discussed elsewhere.
Fig. 1. SEM secondary electron images of samples A (a, b) and B (c, d).
K.-K. Wang et al. / Materials Characterization 107 (2015) 23–28
25
Fig. 2. Cross sectional TEM of sample A (a) BFI, (b) SADP taken from grain 2 in (a) (zone axis = [011]Γ1), (c) SADP from position A in (a) (zone axis = ½1120 δ and [011]Γ2), (d) corresponding high resolution image of position A, (e) SADP from position B in (a) (zone axis = [011]Γ2 and [011]Γ1) and (f) corresponding high resolution image of position B. All the SADPs were taken at the same tilting angle.
Fig. 2a is a TEM bright field image (BFI) taken from a cross-sectional FIB sample of the sample A showing the Fe–Zn IMCs formed at and near the interface. The Fe–Zn interface is located slightly below the bottom of the BFI. The Fe–Zn grains are labeled from 1 to 5. The selected area diffraction pattern (SADP) acquired from grain 2 is shown in Fig. 2b, verifying that the grain 2 has a BCC crystal structure with a = 0.90 nm, that is, it is a Γ1 grain. Grain 1 and 3 are also found to be the Γ1 phase. As indicate in the figure, a grain (grain 4) about 100–200 nm in size lies between grain 2 and grain 5 of the outer layer. Fig. 2c shows a SADP acquired from position A at the interface of grains 4 and 5. Diffraction
patterns of a δ grain in the ½1120 δ zone and a Γ2 grain in the [011]Γ2 zone are identified. Dark field images (not shown) show that grain 4 is a Γ2 grain and grain 5 is a δ one. A well-defined OR is determined between them as ½1120 δ//[011]Γ2, (0001)δ//ð111Þ Γ2 and ð1100Þ δ//ð211Þ Γ2. In addition, a large portion of the interface is parallel to the (0001)δ// ð111Þ Γ2 planes with two short segments which are inclined approximately 70 degrees from (0001)δ as revealed in Fig. 2a. These segments are therefore the ð111Þ Γ2 plane. Fig. 2d shows the high-resolution image of position A, which reveals that the interface is a (0001)δ// ð111Þ Γ2 interface. Structural ledges of less than 10 nm high are observed
26
K.-K. Wang et al. / Materials Characterization 107 (2015) 23–28
at the interfaces with a ð111Þ Γ2 side surface, indicating that the {111}Γ2 planes possess a low interfacial energy. Although grain 1 and 3 are also Γ1 grains, no OR was found with the δ grain. The SADP in Fig. 2e taken from the Γ1/Γ2 interface (position B in Fig. 2a) shows that these two phases exhibit a cube-on-cube OR, that is, (100)Γ2//(100)Γ1 and [001]Γ2//[001]Γ1. The interface has two segments parallel to the ð111Þ Γ1/Γ2 and ð111Þ Γ1/Γ2 planes, respectively. The high-resolution image of position B is shown in Fig. 2f. Ledges of one lattice-plane height are also observed. No misfit dislocations are observed at both the δ/Γ2 and the Γ1/Γ2 interfaces. In sample B of poor uniformity the same ORs are observed too. Fig. 3a is a TEM BFI showing that the Γ1, Γ2 and δ phases are formed irregularly after galvannealing. The position of the original Fe–Zn interface can be identified unambiguously by the presence of a thin residual oxide layer. The coating is not uniform and the thickness varies from 4 to 8 μm, which is also evident in Fig. 1d. The Γ2 grain is approximately 3 μm in size with a δ grain lies next to it. Two smaller Γ1 grains are also present at the corners of the Γ2/δ interface. No OR was found between the δ and Γ2 grains here. However, the SADP in Fig. 3b taken
Fig. 3. Cross-sectional TEM (a) BFI and (b) SADP taken from the circle in (a) (zone axis = [011]Γ2 and [011]Γ1) for sample B.
from the circle in Fig. 3a shows the same cube-on-cube OR between Γ1 and Γ2 as that shown in Fig. 2d. Again, the interface is basically zigzagged ð111Þ Γ1/Γ2 and ð111Þ Γ1/Γ2 planes. Fig. 4a shows another BFI taken from an area at and beneath the original Fe–Zn interface (dashed line). Γ1 grains are formed extensively below the interface along the α-Fe grain boundaries to a depth of about 4 μm. This area is similar to the Γ1/Γ2 area in Fig. 1d. Two α-Fe grains in Fig.4a are thus surrounded by the Γ1 grains. The SADP of Fig. 4b taken from position A in Fig. 4a shows both the α-Fe and Γ1 patterns are in the [111] zone. The SADP in Fig. 4c, taken from position B after some tilting, shows that the α-Fe and Γ1 patterns are both in the [001] zone. Accordingly, both α-Fe/Γ1 grains exhibit the same cube-on-cube OR. The interfaces, however, are wavy and have no preferred crystallographic planes. The ORs of δ/Γ2 and Γ2/Γ1 observed in this study and the corresponding lattice mismatches are summarized in Table 2. The lattice mismatches are calculated by the coincidence site lattice model, which results in a good indication of the interfacial strain [20]. Crystals having a cube-on-cube ORs usually possess interfaces of the close-packed {111} planes which are believed to exhibit the lowest interfacial energy among the low index planes. This has been observed in Fig. 2f of the Γ1/Γ2 interface. The {111}/(0001) interface of the close-packed planes were also frequently observed for FCC/HCP crystals [20,21]. The high coincidence of the interfacial atoms across the interface can effectively reduce the interfacial energy. It is therefore evident that the extremely low lattice mismatches of 0.3–0.7% between the Γ1/Γ2 and δ/Γ2 interfaces allow excellent lattice coincidence at the interfaces. In Fig. 2d and f, no misfit dislocations are observed at the δ/Γ2 and Γ2/Γ1 interfaces, further confirming the excellent lattice coherence. Because of the OR and the very low interfacial energy between the Γ1/Γ2 interface, it is argued that the same OR and interface is also applicable to the δ/Γ1 interface which is, however, different from the (0001)δ//ð123Þ Γ1 one reported [12]. Considering the heterogeneous nucleation at the interface of two existing phases, the activation energy can be greatly reduced by the low interfacial energies of the nucleus. As a consequence, the energy barrier for the nucleation of Γ2 at the coherent δ/Γ1 interface should be very low. It thus explains why the Γ2 grain was observed to form between the δ and Γ1 (grain 2) in Fig. 2a. In the figure, the δ grain has no orientation relationship with grain 1 and 3, and the nucleation of the Γ2 grain at the interfaces is impeded due to the lack of OR and therefore the larger activation energy of nucleation. The cube-on-cube OR between cubic/cubic or cubic/tetragonal crystals having a lattice mismatch of 15% or less is predicted theoretically [22–24] and has indeed been frequently observed [25–27]. This is also confirmed by this research on the Γ2/Γ1 and α-Fe/Γ1 ORs. No OR, nevertheless, between Γ1 and Γ2 was reported previously, even though it appears obviously from Table 2 that the two phases have a low lattice misfit. The low lattice mismatch and excellent lattice coherence also allows the Γ1 phase to nucleate epitaxially at the α grain boundaries. This is indeed the case for sample B. Moreover, the formation of the Γ1 phase during the galvannealing of sample A might follow the same mechanism and exhibited a cube-on-cube OR with one of the α grains. Other ORs between the α-Fe and Γ1 phases, such as (110)α-Fe//(420)Γ1 and [001]α-Fe//½123 Γ1 in electrodeposited samples or ð112Þ α-Fe//ð111Þ Γ1 and [330]α-Fe 7° from [110]Γ1 in hot-dipped samples were reported [12,13]. However, these ORs are lack of consistence and reproducibility. Due to the strong reaction between Zn and Fe, the IMCs can grow rapidly and change the microstructure rapidly. A Γ1 grain and a nextneighbored α-Fe grain may be brought together after the original αFe grain was consumed by the Γ1 grain. In addition, no OR should be observed between the δ and Γ1 phases if the latter nucleated at the α grain boundaries instead of at the δ/α−Fe interface. The rare observation of the OR between the δ and Γ1 phases can be attributed to this reason. Although the α-Fe/Γ1 OR was observed, no well-defined interfaces are present between them (Fig. 4). This may be due to the much larger
K.-K. Wang et al. / Materials Characterization 107 (2015) 23–28
27
Table 2 ORs of observed in this study and the corresponding lattice mismatches. Phases δ/Γ2 Γ2/Γ1 α-Fe/Γ1
OR
Lattice mismatch (%)
½1120δ ==½011Γ2
0.7 8.1
ð0001Þδ ==ð111ÞΓ2 (100)Γ2//(100)Γ1 [001]Γ2//[001]Γ1 (100)α ‐ Fe//(100Γ1) [001]α ‐ Fe//[001]Γ1
0.3 4.9
lattice misfit of 4.9%. However, the lattice mismatch of 4.9% between αFe and Γ1 phases is still low for remaining interfacial coherency by forming misfit dislocations. Another possibility is that the growth of the Γ1 phase may be associated with the redistribution of the solute atoms, such as Si, which have a low solubility in Γ1 [28,29]. The accumulation of the solute atoms at the interface therefore changes the interfacial energy. ORs between various Fe–Zn IMCs were reported to be present in some TEM studies of the galvannealed coating [7,12,13] but not in others [11,14–17] of the galvannealed coating. The major reason can be attributed to the inhomogeneous local Fe–Zn reactions. Due to the presence of the inhibition layer and the residual oxides at the interface, the rapid interdiffusion of Fe and Zn occurs solely through such short cuts as grain and phase boundaries, which results in outburst at the positions and removes the overlaid layer(s). The outburst and the fast refill of liquid Zn break the presumed sequential growth of the intermetallic layers [30,31]. The TEM observation shown in Fig. 3a clearly demonstrated that the IMCs do not stack sequentially on substrate as Fe dissolves rapidly into Zn through the holes in the residual oxide film. Moreover, the inward movement of the interface during galvannealing consumed the original α-Fe grains at the surface which have a small size of 1–2 μm, destruct the ORs. In the present study of sample B, the Zn which flowed in through the holes in the oxide film further diffused relatively deep into α-Fe along the grain boundaries at the high GA temperature (833 K) and formed the Γ1 phase shown in Fig. 4a. The surface α-Fe grains were retained possibly because of the high Mn and Si contents inside the grains. The solubility of Si in the Γ1 phase was estimated to be as low as 0.25 at.% (or 0.11 wt.%) at 753 K [32]. The growth of Γ1 is dragged by the redistribution of Si at the interface. The OR between Γ1 and α-Fe was therefore reserved. The fast formation of the Γ1 phase at the α-Fe grain boundaries was also reported by Adachi and Kamei [11]. 4. Conclusions The microstructures of the GA coatings of two Mn–Si TRIP steels were studied by TEM. Crystallographic ORs between the α-Fe and Γ1 phases, the Γ1 and Γ2 phases, as well as the Γ2 and δ phases were characterized. Both the α-Fe/Γ1 phases and the Γ1/Γ2 phases exhibit a cube-oncube orientation relationship. The orientation relationship between Γ2 and δ phases is identified as ½1120 δ//[011]Γ2 and (0002)δ//ð111Þ Γ2. The interfaces of Γ2/δ and Γ1/Γ2 are basically the close-packed (111)Γ1, (111)Γ2 and (0001)δ planes. The transformation of Γ2 phase is facilitated by the reduced activation energy of heterogeneous nucleation by the coherent interface with both δ and Γ1 phases. Acknowledgments The NSYSU team acknowledges the Ministry of Science and Technology, R. O. C. and China Steel Corporation (CSC) for supporting the work Fig. 4. Cross sectional TEM (a) BFI and SADPs taken from (b) position A (zone axis = [111]Γ1 and [111]α-Fe) and (c) position B (zone axis = [001]Γ1 and [001]α-Fe) in (a) for sample B.
28
K.-K. Wang et al. / Materials Characterization 107 (2015) 23–28
under grant numbers MOST 103-2622-E-006-037 and 02T1D-RE030, respectively. The authors gratefully thank Dr. Luke Hsiung of the Lawrence Livermore National Laboratory for discussions as well as Mr. J.-S. Lin and Mr. J.-M. Huang of the CSC for helping with the rolling/ galvanizing processes.
References [1] E. Marquard, ASM Handbook, Vol. 3, Alloy Phase Diagrams, 10th ed. American Society for Metals, Materials Park, OH, 1992. [2] H.E. Swanson, E. Tatge, Standard X-Ray diffraction powder patterns, Natl. Bur. Stand. (U.S.) Circ. 359 (1953) 1. [3] P. Villars, L.D. Calvert, Pearson's Handbook of Crystallographic Data for Intermetallic Phases, Vols. 1–3, 2nd ed. American Society for Metals, Materials Park, OH, 1991. [4] C.H.E. Belin, R.C.H. Belin, Synthesis and crystal structure determinations in the Γ and δ phase domains of the iron–zinc system: electronic and bonding analysis of Fe13Zn39 and FeZn10, a subtle deviation from the Hume–Rothery standard, J. Solid State Chem. 151 (2000) 85–95. [5] C.E. Jordan, A.R. Marder, Fe–Zn phase formation in interstitial-free steels hot-dip galvanized at 450 °C: part I 0.00 wt% Al-Zn baths, J. Mater. Sci. 32 (1997) 5593–5602. [6] J.-I. Inagaka, M. Saturai, T. Watanabe, Alloying reactions in hot dip galvanizing and galvannealing processes, ISIJ Int. 35 (1995) 1388–1393. [7] H. Ohtsubo, T. Yagi, K. Nakai, Y. Ohmori, Crystallography of intermetallic interface layers in hot-dip galvanizing steel, ISIJ Int. 36 (1996) 317–323. [8] C.E. Jordan, A.R. Marder, Fe–Zn phase formation in interstitial-free steels hot-dip galvanized at 450 °C: part II 0.20 wt% Al–Zn baths, J. Mater. Sci. 32 (1997) 5603–5610. [9] M. Saito, Y. Uchida, T. Kittaka, Y. Hirose, Y. Hisamatsu, Formation behavior of alloy layer in initial stages of galvanizing, Tetsu to Hagane 77 (1991) 947–954. [10] H. Yamaguchi, Y. Hisamatsu, Reaction mechanism of the sheet galvanizing, Trans. ISIJ 19 (1979) 649–658. [11] Y. Adachi, K. Kamei, TEM study of diffusional alfa-Fe/Zn interface, Acta Metall. Mater. 43 (1995) 3189–3197. [12] Y. Lin, W.A. Chiou, M. Meshii, Atomic structures of interfaces in annealed electrogalvanized steel sheets, Proc. of 2th Intern. Conf. Zinc and Zinc Alloy Coated Steel Sheets, Verlag Stahleisen, Düsseldorf 1992, pp. 429–435. [13] C.S. Lin, M. Meshii, The effect of steel chemistry on the formation of Fe–Zn intermetallic compounds of galvanneal-coated steel sheets, Metall. Mater. Trans. B 25B (1994) 721–730. [14] F.H. Kao, W.C. Li, C.Y. Chen, C.Y. Huang, J.R. Yang, S.H. Wang, Cross-sectional observation of the intermetallic phase in a galvannealed steel, Mater. Sci. Eng. A A499 (2009) 45–48.
[15] J. Yu, J. Liu, W. Zhou, J. Zhang, J. Wu, Cross-sectional TEM observation of iron–zinc intermetallic Γ and Γ1 phases in commercial galvannealed IF steel sheets, Mater. Des. 28 (2007) 249–253. [16] T. Kato, M.H. Hong, K. Nunome, K. Sasaki, K. Kuroda, H. Saka, Cross-sectional TEM observation of multilayer structure of a galvannealed steel, Thin Solid Films 319 (1998) 132–139. [17] L. Cho, M.S. Kim, Y.H. Kim, B.C. De Cooman, Influence of gas atmosphere dew point on the galvannealing of CMnSi TRIP steel, Metall. Mater. Trans. A 44A (2013) 5081–5095. [18] K.K. Wang, C.W. Hsu, L.W. Chang, D. Gan, Key to successful galvannealing of CMnSi TRIP steels, Proc. of 9th Intern. Conf. Zinc and Zinc Alloy Coated Steel Sheets, The Chinese Society for Metals, Beijing 2013, pp. 68–72. [19] K.K. Wang, C.W. Hsu, L. Chang, D. Gan, K.C. Yang, Role of Al in Zn bath on the formation of the inhibition layer during hot-dip galvanizing for a 1.2Si–1.5Mn transformation-induced plasticity steel, Appl. Surf. Sci. 285P (2013) 458–468. [20] M.X. Zhang, P.M. Kalley, Crystallographic features of phase transformations in solids, Prog. Mater. Sci. 54 (2009) 1101–1170. [21] W.Z. Zhang, On the crystallography of precipitation, Prog. Mater. Sci. 50 (2005) 181–292. [22] A. Seki, K. Kamei, Epitaxial growth of Zn and Zn–Ni electrodeposits on steel sheets, ISIJ Int. 32 (1992) 1306–1310. [23] A. Van de Walle, Q. Hong, L. Miljacic, C. Balaji Gopal, S. Demers, G. Pomrehn, A. Kowalski, P. Tiwary, Ab initio calculation of anisotropic interfacial excess free energies, Phys. Rev. B89 (2014) 184101. [24] Z. Mao, C. Booth-Morrison, E. Plotnikov, D. Seidman, Effects of temperature and ferromagnetism on the γ-Ni/γ′-Ni3Al interfacial free energy from first principles calculations, J. Mater. Sci. 47 (2012) 7653–7659. [25] Y. Gao, K.L. Merkle, High-resolution electron microscopy of metal/metal and metal/ metal-oxide interfaces in the Ag/Ni and Au/Ni systems, J. Mater. Res. 5 (1990) 1995–2003. [26] P. Lu, F. Cosandey, Dislocation structures at Cu–MgO and Pd–MgO interfaces, Acta Metall. Mater. 40 (1992) S259–S266. [27] B.J. Kooi, H.B. Groen, J.T.M. De Hosson, Atomic structure of interfaces between Mn3O4 precipitates and Ag studied with HRTEM, Acta Mater. 45 (1997) 3587–3607. [28] M.S. Kozdras, P. Niessen, Silicon-induced destabilization of galvanized coatings in the sandelin peak region, Metallography 22 (1989) 253–267. [29] C. Sha, S. Liu, Y. Du, H. Xu, L. Zhang, Y. Liu, Experimental investigation and thermodynamic reassessment of the Fe–Si–Zn system, Calphad 34 (2010) 405–414. [30] M. Guttmann, Diffusive phase transformations in hot dip galvanizing, Mater. Sci. Forum 155–156 (1994) 527–548. [31] C.E. Jordan, A.R. Marder, The effect of iron oxide as an inhibition layer on iron–zinc reaction during hot-dip galvanizing, Metall. Mater. Trans. 29B (1998) 479–483. [32] J. Wang, X. Su, F. Yin, Z. Li, M. Zhao, The 480 °C and 405 °C isothermal sections of the phase diagram of Fe–Zn–Si ternary system, J. Alloys Compd. 399 (2005) 214–218.
Contents lists available at ScienceDirect
Materials Characterization journal homepage: www.elsevier.com/locate/matchar
Orientation relationships and interfaces of the Fe–Zn intermetallic phases in galvannealed CMnSi-TRIP steels Kuang-Kuo Wang a, Guan-Lin You a, Liuwen Chang a,⁎, Dershin Gan a, Lung-Jen Chiang b a b
Research Center for Physical Properties and Microstructure of Metals, Department of Materials and Optoelectronic Science, National Sun Yat-Sen University, Kaohsiung 80424, Taiwan, ROC Steel Research and Development Department, China Steel Corporation, Kaohsiung 81233, Taiwan, ROC
a r t i c l e
i n f o
Article history: Received 19 March 2015 Received in revised form 10 June 2015 Accepted 30 June 2015 Available online 2 July 2015 Keywords: Galvannealed coating Mn–Si TRIP steels Fe–Zn intermetallic Orientation relationship Interface
a b s t r a c t Hot-dip galvannealed coatings of two Mn–Si TRIP steels have been characterized by transmission electron microscopy to study the crystallographic orientation relationships between the α-Fe and Γ1 phases, the Γ1 and Γ2 phases, as well as the Γ2 and δ phases. Both the α-Fe/Γ1 phases and the Γ1/Γ2 phases were found to exhibit a cube-on-cube orientation relationship. The orientation relationship at the Γ2/δ interface was identified to be ½1120 δ//[011]Γ2 and (0002) δ//ð111Þ Γ2. The interfaces of Γ2/δ and Γ1/Γ2 are basically the close-packed plane of each phase. The nucleation of Γ2 phase between the δ and Γ1 phases was discussed by the heterogeneous nucleation theory based on the orientation relationships and the interfaces. © 2015 Elsevier Inc. All rights reserved.
1. Introduction Hot-dip galvannealed (GA) steel sheets have been extensively used in car manufacture. An optimized GA coating consists of two consecutive layers of Γ1 and δ on the steel matrix, though totally four Fe–Zn intermetallic compounds (IMCs), Γ1, Γ2, δ and ζ, exist below 803 K according to the equilibrium phase diagram [1]. The symbols and crystallographic information of the Fe–Zn IMCs is listed in Table 1 for reference [2–4]. It was reported that the ζ phase is the first phase to form at the Fe/Zn interface followed by the δ phase and finally the Γ1 and/or Γ2 phases on annealing at or below 773 K for Zn coating containing 0.1 wt.% Al or less [5–7]. However, the Fe/Zn reactions are inhibited temporarily by the formation of an inhibition layer of Fe2Al5 − xZnx in hotdip galvanizing if some Al is added to the Zn bath. An additional annealing at 773–813 K is thus necessary to accelerate the interdiffusion across the inhibition layer for the occurrence of the Fe/Zn reactions. The δ, instead of the ζ, phase was suggested to be the first IMC to form at the Fe2Al5 − xZnx/α-Fe interface followed by the Γ1 phase at the δ/α-Fe interface [8–10]. The formation kinetics of the Fe–Zn IMCs attracts great attentions due to its importance on optimizing the microstructure and the properties of the GA coating. One of the important issues which have not been clarified yet is whether the IMCs are formed epitaxially to the ferrite grains and the prior-formed IMCs. Adachi and Kamei [11] argued that the Γ1/α-Fe interface should be coherent or partly coherent due to the ⁎ Corresponding author. E-mail address: [email protected] (L. Chang).
http://dx.doi.org/10.1016/j.matchar.2015.06.041 1044-5803/© 2015 Elsevier Inc. All rights reserved.
presence of growth ledge at the interface. Various orientation relationships (ORs) have been proposed between the α-Fe and Γ1 [12,13], αFe and Γ2 [13] and α-Fe and ζ phases [7]. However, the OR between the α-Fe and Γ1 phases derived from the electrodeposited sample is different from that from the hot-dip sample. The crystallographic characteristics of the Fe–Zn reactions are difficult to be revealed because the reactions are so rapid and sometimes locally intense that the inherent ORs do not show up. Accordingly, many cross-sectional TEM works have been conducted recently [14–16], but no results on the ORs and the interfaces of the Fe–Zn intermetallic phases have been reported. In addition, Cho et al. [17] reported that no ORs were found between the α-Fe and Γ1, Γ1 and Γ2, as well as Γ2 and δ phases during the galvannealing of a Mn–Si TRIP steel. In this work, we conducted transmission electron microscopy (TEM) analysis on two galvannealed samples of Mn–Si transformation induced plasticity (TRIP) steels. It is known that the oxides formed by selective oxidation on the surface of the Mn–Si TRIP steels during intercritical annealing can inhibit the Fe–Zn reactions on subsequent annealing. By changing the annealing conditions, various reaction results can therefore be observed. In this work, we report, for the first time, the ORs between the consecutive pairs of α-Fe/Γ1, Γ1/Γ2, and Γ2/δ phases. The crystallographic data can provide valuable information for further discussion of the nucleation and growth kinetics of the Fe–Zn IMCs. 2. Experimental GA coatings on two TRIP steels containing 1.0 wt.% Si-2.0 wt.% Mn (steel A) and 1.8 wt.% Si-1.9 wt.% Mn (steel B), respectively, were
24
K.-K. Wang et al. / Materials Characterization 107 (2015) 23–28
Table 1 Crystallographic characteristics of the Fe–Zn phases [2–4]. Phase
Formula
Crystal structure
Space group
Lattice parameters (nm)
ζ
FeZn13–15
Monoclinic
C2/m
δ (δk/δp) Γ2
FeZn7–10 Fe11Zn40
Hexagonal FCC
P63/mmc
a = 1.3424, b = 0.7608 c = 0.5061, β = 127.3° a = 1.2787, c = 5.7222 a = 1.7963
Γ1
Fe3Zn10
BCC
I43m
a = 0.9018
α-Fe
Fe
BCC
Im3m
a = 0.2866
F43m
TEM (FEI Tecnai G2, at 200 kV). After galvannealing, the central part (50 × 50 mm2) of the samples were cut into 10 × 10 mm2 coupons, cleaned in ethanol, and loaded in an FEI Nova-200 dual beam instrument. The TEM cross-sectional sample was prepared by milling an electron-transparent membrane of approximately 10 μm × 10 μm in size out of the desired region of the bulk GA sample. The crosssectional sample was then transferred to an Omniprobe grid and further thinned to a thickness of 60 nm or less to enable high-resolution imaging in TEM. 3. Results and discussion
studied. The steels were prepared by vacuum induction melting. The ingots were reheated and soaked at 1523 K, and hot rolled to approximately 5 mm thick with a finish-rolling temperature of about 1173 K. The strip was rapidly cooled to about 773 K and then soaked immediately at 823 K for 2 h. Afterwards, the hot band was machined to 4 mm thick to remove the surface scale and cold rolled to 1.0 mm thick. The cold-rolled sheets were cut into samples of 120 mm × 250 mm for simulating the continuous annealing and hot-dip galvannealing process using an Iwatani hot-dip simulator. Prior to annealing, the samples were cleaned by dipping in a 2% alkaline solution for 20 min at 343 K, mechanical brushing, water-spray rinsing and hot-air drying in sequence. The steels were first annealed at 1073 K for 60 s in a 10%H2–N2 atmosphere with a dew point (dp) of 243 K (PO2 = 7.24 × 10− 24 atm) or 273 K (PO2 = 1.87 × 10− 21 atm), cooled to 733 K at a rate of 15 K/s after soaking, held for 60 s, and dipped in a Zn bath for 3 s. The Al content in the bath was 0.16 wt.% with saturated Fe. The high Al content ensured the formation of the Fe–Al layer with a high coverage ratio at the Fe–Zn interface in galvanizing which provides a sufficient wettability to the re-melted Zn in galvannealing. High quality GA coatings were obtained accordingly [18]. The samples were finally galvannealed at 793–833 K for 20 s in an induction furnace of the simulator after hot-dipping. The GA coatings were analyzed by scanning electron microscopy (SEM, Zeiss Supra 55, at 5–15 kV) and
The effects of processing parameters and steel composition on the galvanizing and galvannealing characteristics of the Mn–Si TRIP steels used in this study were reported previously [18,19]. In this work we focus on the ORs at the various interfaces. Fig. 1 shows the SEM secondary electron images of sample A (steel A, annealed at dp = 273 K and galvannealed at 793 K for 20 s) and sample B (steel B, annealed at dp = 243 K and galvannealed at 833 K for 20 s), using two extreme GA conditions. Fig. 1a shows that sample A exhibits a relatively flat surface on which Fe–Zn grains of less than 3 μm in size are observed. The cross-sectional micrograph in Fig. 1b verifies that the GA coating is uniform in thickness with a thin inner layer (1–1.5 μm) and a thick outer layer. The average Fe content of the coating determined by energy dispersive spectroscopy is 10.7 wt.%, which is within the typical GA region of 9–11 wt.%. Coating B, however, exhibits a rough surface scattered with coarse IMC particles as large as 20 μm (Fig. 1c). The crosssectional micrograph in Fig. 1d shows that the coating consists of a thin layer less than 2 μm thick and the large particles on it. Furthermore, the thin layer is not continuous at the interface. Accordingly, sample A represents a well-optimized GA coating but sample B is not. The results agree relatively well with a recent study carried out for a 2.2Mn–1.4Si TRIP steel [17]. Because we will focus on the ORs and interfaces of the various phases, the complex growth mechanisms of the Fe–Zn phases are out of the scope of this paper and will be discussed elsewhere.
Fig. 1. SEM secondary electron images of samples A (a, b) and B (c, d).
K.-K. Wang et al. / Materials Characterization 107 (2015) 23–28
25
Fig. 2. Cross sectional TEM of sample A (a) BFI, (b) SADP taken from grain 2 in (a) (zone axis = [011]Γ1), (c) SADP from position A in (a) (zone axis = ½1120 δ and [011]Γ2), (d) corresponding high resolution image of position A, (e) SADP from position B in (a) (zone axis = [011]Γ2 and [011]Γ1) and (f) corresponding high resolution image of position B. All the SADPs were taken at the same tilting angle.
Fig. 2a is a TEM bright field image (BFI) taken from a cross-sectional FIB sample of the sample A showing the Fe–Zn IMCs formed at and near the interface. The Fe–Zn interface is located slightly below the bottom of the BFI. The Fe–Zn grains are labeled from 1 to 5. The selected area diffraction pattern (SADP) acquired from grain 2 is shown in Fig. 2b, verifying that the grain 2 has a BCC crystal structure with a = 0.90 nm, that is, it is a Γ1 grain. Grain 1 and 3 are also found to be the Γ1 phase. As indicate in the figure, a grain (grain 4) about 100–200 nm in size lies between grain 2 and grain 5 of the outer layer. Fig. 2c shows a SADP acquired from position A at the interface of grains 4 and 5. Diffraction
patterns of a δ grain in the ½1120 δ zone and a Γ2 grain in the [011]Γ2 zone are identified. Dark field images (not shown) show that grain 4 is a Γ2 grain and grain 5 is a δ one. A well-defined OR is determined between them as ½1120 δ//[011]Γ2, (0001)δ//ð111Þ Γ2 and ð1100Þ δ//ð211Þ Γ2. In addition, a large portion of the interface is parallel to the (0001)δ// ð111Þ Γ2 planes with two short segments which are inclined approximately 70 degrees from (0001)δ as revealed in Fig. 2a. These segments are therefore the ð111Þ Γ2 plane. Fig. 2d shows the high-resolution image of position A, which reveals that the interface is a (0001)δ// ð111Þ Γ2 interface. Structural ledges of less than 10 nm high are observed
26
K.-K. Wang et al. / Materials Characterization 107 (2015) 23–28
at the interfaces with a ð111Þ Γ2 side surface, indicating that the {111}Γ2 planes possess a low interfacial energy. Although grain 1 and 3 are also Γ1 grains, no OR was found with the δ grain. The SADP in Fig. 2e taken from the Γ1/Γ2 interface (position B in Fig. 2a) shows that these two phases exhibit a cube-on-cube OR, that is, (100)Γ2//(100)Γ1 and [001]Γ2//[001]Γ1. The interface has two segments parallel to the ð111Þ Γ1/Γ2 and ð111Þ Γ1/Γ2 planes, respectively. The high-resolution image of position B is shown in Fig. 2f. Ledges of one lattice-plane height are also observed. No misfit dislocations are observed at both the δ/Γ2 and the Γ1/Γ2 interfaces. In sample B of poor uniformity the same ORs are observed too. Fig. 3a is a TEM BFI showing that the Γ1, Γ2 and δ phases are formed irregularly after galvannealing. The position of the original Fe–Zn interface can be identified unambiguously by the presence of a thin residual oxide layer. The coating is not uniform and the thickness varies from 4 to 8 μm, which is also evident in Fig. 1d. The Γ2 grain is approximately 3 μm in size with a δ grain lies next to it. Two smaller Γ1 grains are also present at the corners of the Γ2/δ interface. No OR was found between the δ and Γ2 grains here. However, the SADP in Fig. 3b taken
Fig. 3. Cross-sectional TEM (a) BFI and (b) SADP taken from the circle in (a) (zone axis = [011]Γ2 and [011]Γ1) for sample B.
from the circle in Fig. 3a shows the same cube-on-cube OR between Γ1 and Γ2 as that shown in Fig. 2d. Again, the interface is basically zigzagged ð111Þ Γ1/Γ2 and ð111Þ Γ1/Γ2 planes. Fig. 4a shows another BFI taken from an area at and beneath the original Fe–Zn interface (dashed line). Γ1 grains are formed extensively below the interface along the α-Fe grain boundaries to a depth of about 4 μm. This area is similar to the Γ1/Γ2 area in Fig. 1d. Two α-Fe grains in Fig.4a are thus surrounded by the Γ1 grains. The SADP of Fig. 4b taken from position A in Fig. 4a shows both the α-Fe and Γ1 patterns are in the [111] zone. The SADP in Fig. 4c, taken from position B after some tilting, shows that the α-Fe and Γ1 patterns are both in the [001] zone. Accordingly, both α-Fe/Γ1 grains exhibit the same cube-on-cube OR. The interfaces, however, are wavy and have no preferred crystallographic planes. The ORs of δ/Γ2 and Γ2/Γ1 observed in this study and the corresponding lattice mismatches are summarized in Table 2. The lattice mismatches are calculated by the coincidence site lattice model, which results in a good indication of the interfacial strain [20]. Crystals having a cube-on-cube ORs usually possess interfaces of the close-packed {111} planes which are believed to exhibit the lowest interfacial energy among the low index planes. This has been observed in Fig. 2f of the Γ1/Γ2 interface. The {111}/(0001) interface of the close-packed planes were also frequently observed for FCC/HCP crystals [20,21]. The high coincidence of the interfacial atoms across the interface can effectively reduce the interfacial energy. It is therefore evident that the extremely low lattice mismatches of 0.3–0.7% between the Γ1/Γ2 and δ/Γ2 interfaces allow excellent lattice coincidence at the interfaces. In Fig. 2d and f, no misfit dislocations are observed at the δ/Γ2 and Γ2/Γ1 interfaces, further confirming the excellent lattice coherence. Because of the OR and the very low interfacial energy between the Γ1/Γ2 interface, it is argued that the same OR and interface is also applicable to the δ/Γ1 interface which is, however, different from the (0001)δ//ð123Þ Γ1 one reported [12]. Considering the heterogeneous nucleation at the interface of two existing phases, the activation energy can be greatly reduced by the low interfacial energies of the nucleus. As a consequence, the energy barrier for the nucleation of Γ2 at the coherent δ/Γ1 interface should be very low. It thus explains why the Γ2 grain was observed to form between the δ and Γ1 (grain 2) in Fig. 2a. In the figure, the δ grain has no orientation relationship with grain 1 and 3, and the nucleation of the Γ2 grain at the interfaces is impeded due to the lack of OR and therefore the larger activation energy of nucleation. The cube-on-cube OR between cubic/cubic or cubic/tetragonal crystals having a lattice mismatch of 15% or less is predicted theoretically [22–24] and has indeed been frequently observed [25–27]. This is also confirmed by this research on the Γ2/Γ1 and α-Fe/Γ1 ORs. No OR, nevertheless, between Γ1 and Γ2 was reported previously, even though it appears obviously from Table 2 that the two phases have a low lattice misfit. The low lattice mismatch and excellent lattice coherence also allows the Γ1 phase to nucleate epitaxially at the α grain boundaries. This is indeed the case for sample B. Moreover, the formation of the Γ1 phase during the galvannealing of sample A might follow the same mechanism and exhibited a cube-on-cube OR with one of the α grains. Other ORs between the α-Fe and Γ1 phases, such as (110)α-Fe//(420)Γ1 and [001]α-Fe//½123 Γ1 in electrodeposited samples or ð112Þ α-Fe//ð111Þ Γ1 and [330]α-Fe 7° from [110]Γ1 in hot-dipped samples were reported [12,13]. However, these ORs are lack of consistence and reproducibility. Due to the strong reaction between Zn and Fe, the IMCs can grow rapidly and change the microstructure rapidly. A Γ1 grain and a nextneighbored α-Fe grain may be brought together after the original αFe grain was consumed by the Γ1 grain. In addition, no OR should be observed between the δ and Γ1 phases if the latter nucleated at the α grain boundaries instead of at the δ/α−Fe interface. The rare observation of the OR between the δ and Γ1 phases can be attributed to this reason. Although the α-Fe/Γ1 OR was observed, no well-defined interfaces are present between them (Fig. 4). This may be due to the much larger
K.-K. Wang et al. / Materials Characterization 107 (2015) 23–28
27
Table 2 ORs of observed in this study and the corresponding lattice mismatches. Phases δ/Γ2 Γ2/Γ1 α-Fe/Γ1
OR
Lattice mismatch (%)
½1120δ ==½011Γ2
0.7 8.1
ð0001Þδ ==ð111ÞΓ2 (100)Γ2//(100)Γ1 [001]Γ2//[001]Γ1 (100)α ‐ Fe//(100Γ1) [001]α ‐ Fe//[001]Γ1
0.3 4.9
lattice misfit of 4.9%. However, the lattice mismatch of 4.9% between αFe and Γ1 phases is still low for remaining interfacial coherency by forming misfit dislocations. Another possibility is that the growth of the Γ1 phase may be associated with the redistribution of the solute atoms, such as Si, which have a low solubility in Γ1 [28,29]. The accumulation of the solute atoms at the interface therefore changes the interfacial energy. ORs between various Fe–Zn IMCs were reported to be present in some TEM studies of the galvannealed coating [7,12,13] but not in others [11,14–17] of the galvannealed coating. The major reason can be attributed to the inhomogeneous local Fe–Zn reactions. Due to the presence of the inhibition layer and the residual oxides at the interface, the rapid interdiffusion of Fe and Zn occurs solely through such short cuts as grain and phase boundaries, which results in outburst at the positions and removes the overlaid layer(s). The outburst and the fast refill of liquid Zn break the presumed sequential growth of the intermetallic layers [30,31]. The TEM observation shown in Fig. 3a clearly demonstrated that the IMCs do not stack sequentially on substrate as Fe dissolves rapidly into Zn through the holes in the residual oxide film. Moreover, the inward movement of the interface during galvannealing consumed the original α-Fe grains at the surface which have a small size of 1–2 μm, destruct the ORs. In the present study of sample B, the Zn which flowed in through the holes in the oxide film further diffused relatively deep into α-Fe along the grain boundaries at the high GA temperature (833 K) and formed the Γ1 phase shown in Fig. 4a. The surface α-Fe grains were retained possibly because of the high Mn and Si contents inside the grains. The solubility of Si in the Γ1 phase was estimated to be as low as 0.25 at.% (or 0.11 wt.%) at 753 K [32]. The growth of Γ1 is dragged by the redistribution of Si at the interface. The OR between Γ1 and α-Fe was therefore reserved. The fast formation of the Γ1 phase at the α-Fe grain boundaries was also reported by Adachi and Kamei [11]. 4. Conclusions The microstructures of the GA coatings of two Mn–Si TRIP steels were studied by TEM. Crystallographic ORs between the α-Fe and Γ1 phases, the Γ1 and Γ2 phases, as well as the Γ2 and δ phases were characterized. Both the α-Fe/Γ1 phases and the Γ1/Γ2 phases exhibit a cube-oncube orientation relationship. The orientation relationship between Γ2 and δ phases is identified as ½1120 δ//[011]Γ2 and (0002)δ//ð111Þ Γ2. The interfaces of Γ2/δ and Γ1/Γ2 are basically the close-packed (111)Γ1, (111)Γ2 and (0001)δ planes. The transformation of Γ2 phase is facilitated by the reduced activation energy of heterogeneous nucleation by the coherent interface with both δ and Γ1 phases. Acknowledgments The NSYSU team acknowledges the Ministry of Science and Technology, R. O. C. and China Steel Corporation (CSC) for supporting the work Fig. 4. Cross sectional TEM (a) BFI and SADPs taken from (b) position A (zone axis = [111]Γ1 and [111]α-Fe) and (c) position B (zone axis = [001]Γ1 and [001]α-Fe) in (a) for sample B.
28
K.-K. Wang et al. / Materials Characterization 107 (2015) 23–28
under grant numbers MOST 103-2622-E-006-037 and 02T1D-RE030, respectively. The authors gratefully thank Dr. Luke Hsiung of the Lawrence Livermore National Laboratory for discussions as well as Mr. J.-S. Lin and Mr. J.-M. Huang of the CSC for helping with the rolling/ galvanizing processes.
References [1] E. Marquard, ASM Handbook, Vol. 3, Alloy Phase Diagrams, 10th ed. American Society for Metals, Materials Park, OH, 1992. [2] H.E. Swanson, E. Tatge, Standard X-Ray diffraction powder patterns, Natl. Bur. Stand. (U.S.) Circ. 359 (1953) 1. [3] P. Villars, L.D. Calvert, Pearson's Handbook of Crystallographic Data for Intermetallic Phases, Vols. 1–3, 2nd ed. American Society for Metals, Materials Park, OH, 1991. [4] C.H.E. Belin, R.C.H. Belin, Synthesis and crystal structure determinations in the Γ and δ phase domains of the iron–zinc system: electronic and bonding analysis of Fe13Zn39 and FeZn10, a subtle deviation from the Hume–Rothery standard, J. Solid State Chem. 151 (2000) 85–95. [5] C.E. Jordan, A.R. Marder, Fe–Zn phase formation in interstitial-free steels hot-dip galvanized at 450 °C: part I 0.00 wt% Al-Zn baths, J. Mater. Sci. 32 (1997) 5593–5602. [6] J.-I. Inagaka, M. Saturai, T. Watanabe, Alloying reactions in hot dip galvanizing and galvannealing processes, ISIJ Int. 35 (1995) 1388–1393. [7] H. Ohtsubo, T. Yagi, K. Nakai, Y. Ohmori, Crystallography of intermetallic interface layers in hot-dip galvanizing steel, ISIJ Int. 36 (1996) 317–323. [8] C.E. Jordan, A.R. Marder, Fe–Zn phase formation in interstitial-free steels hot-dip galvanized at 450 °C: part II 0.20 wt% Al–Zn baths, J. Mater. Sci. 32 (1997) 5603–5610. [9] M. Saito, Y. Uchida, T. Kittaka, Y. Hirose, Y. Hisamatsu, Formation behavior of alloy layer in initial stages of galvanizing, Tetsu to Hagane 77 (1991) 947–954. [10] H. Yamaguchi, Y. Hisamatsu, Reaction mechanism of the sheet galvanizing, Trans. ISIJ 19 (1979) 649–658. [11] Y. Adachi, K. Kamei, TEM study of diffusional alfa-Fe/Zn interface, Acta Metall. Mater. 43 (1995) 3189–3197. [12] Y. Lin, W.A. Chiou, M. Meshii, Atomic structures of interfaces in annealed electrogalvanized steel sheets, Proc. of 2th Intern. Conf. Zinc and Zinc Alloy Coated Steel Sheets, Verlag Stahleisen, Düsseldorf 1992, pp. 429–435. [13] C.S. Lin, M. Meshii, The effect of steel chemistry on the formation of Fe–Zn intermetallic compounds of galvanneal-coated steel sheets, Metall. Mater. Trans. B 25B (1994) 721–730. [14] F.H. Kao, W.C. Li, C.Y. Chen, C.Y. Huang, J.R. Yang, S.H. Wang, Cross-sectional observation of the intermetallic phase in a galvannealed steel, Mater. Sci. Eng. A A499 (2009) 45–48.
[15] J. Yu, J. Liu, W. Zhou, J. Zhang, J. Wu, Cross-sectional TEM observation of iron–zinc intermetallic Γ and Γ1 phases in commercial galvannealed IF steel sheets, Mater. Des. 28 (2007) 249–253. [16] T. Kato, M.H. Hong, K. Nunome, K. Sasaki, K. Kuroda, H. Saka, Cross-sectional TEM observation of multilayer structure of a galvannealed steel, Thin Solid Films 319 (1998) 132–139. [17] L. Cho, M.S. Kim, Y.H. Kim, B.C. De Cooman, Influence of gas atmosphere dew point on the galvannealing of CMnSi TRIP steel, Metall. Mater. Trans. A 44A (2013) 5081–5095. [18] K.K. Wang, C.W. Hsu, L.W. Chang, D. Gan, Key to successful galvannealing of CMnSi TRIP steels, Proc. of 9th Intern. Conf. Zinc and Zinc Alloy Coated Steel Sheets, The Chinese Society for Metals, Beijing 2013, pp. 68–72. [19] K.K. Wang, C.W. Hsu, L. Chang, D. Gan, K.C. Yang, Role of Al in Zn bath on the formation of the inhibition layer during hot-dip galvanizing for a 1.2Si–1.5Mn transformation-induced plasticity steel, Appl. Surf. Sci. 285P (2013) 458–468. [20] M.X. Zhang, P.M. Kalley, Crystallographic features of phase transformations in solids, Prog. Mater. Sci. 54 (2009) 1101–1170. [21] W.Z. Zhang, On the crystallography of precipitation, Prog. Mater. Sci. 50 (2005) 181–292. [22] A. Seki, K. Kamei, Epitaxial growth of Zn and Zn–Ni electrodeposits on steel sheets, ISIJ Int. 32 (1992) 1306–1310. [23] A. Van de Walle, Q. Hong, L. Miljacic, C. Balaji Gopal, S. Demers, G. Pomrehn, A. Kowalski, P. Tiwary, Ab initio calculation of anisotropic interfacial excess free energies, Phys. Rev. B89 (2014) 184101. [24] Z. Mao, C. Booth-Morrison, E. Plotnikov, D. Seidman, Effects of temperature and ferromagnetism on the γ-Ni/γ′-Ni3Al interfacial free energy from first principles calculations, J. Mater. Sci. 47 (2012) 7653–7659. [25] Y. Gao, K.L. Merkle, High-resolution electron microscopy of metal/metal and metal/ metal-oxide interfaces in the Ag/Ni and Au/Ni systems, J. Mater. Res. 5 (1990) 1995–2003. [26] P. Lu, F. Cosandey, Dislocation structures at Cu–MgO and Pd–MgO interfaces, Acta Metall. Mater. 40 (1992) S259–S266. [27] B.J. Kooi, H.B. Groen, J.T.M. De Hosson, Atomic structure of interfaces between Mn3O4 precipitates and Ag studied with HRTEM, Acta Mater. 45 (1997) 3587–3607. [28] M.S. Kozdras, P. Niessen, Silicon-induced destabilization of galvanized coatings in the sandelin peak region, Metallography 22 (1989) 253–267. [29] C. Sha, S. Liu, Y. Du, H. Xu, L. Zhang, Y. Liu, Experimental investigation and thermodynamic reassessment of the Fe–Si–Zn system, Calphad 34 (2010) 405–414. [30] M. Guttmann, Diffusive phase transformations in hot dip galvanizing, Mater. Sci. Forum 155–156 (1994) 527–548. [31] C.E. Jordan, A.R. Marder, The effect of iron oxide as an inhibition layer on iron–zinc reaction during hot-dip galvanizing, Metall. Mater. Trans. 29B (1998) 479–483. [32] J. Wang, X. Su, F. Yin, Z. Li, M. Zhao, The 480 °C and 405 °C isothermal sections of the phase diagram of Fe–Zn–Si ternary system, J. Alloys Compd. 399 (2005) 214–218.