Effects of Carbon Contents in Steels on Alloy Layer Growth during Hot-dip Aluminum Coating

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1 , pp Effects of Carbon Contents in Steels on Alloy Layer Growth during Hot-dip Aluminum Coating T. SASAKI, T. YAKOU, K. MOCHIDUKI and K. ICHINOSE 1) Department of Engineering, Yokohama National University, 79-5 Tokowadai, Hodogaya-ku, Yokohama-shi, Kanagawa Japan. 1) Department of Engineering, Tokyo Denki University, 2-2 Nishiki-cho, Kanda, Chiyoda-ku Tokyo Japan. (Received on May 30, 2005; accepted on August 25, 2005) Hot-dip aluminum coating of hypo-eutectoid steels containing mass% carbon were performed, and the alloy layers formed in the coating were investigated. In the hot-dip aluminum coating at immersion temperatures ranging from 700 to 850 C, the alloy layers on the steels consisted of a single phase of the intermetallic compound Fe 2 Al 5. The thickness of the alloy layer increased in proportion to the increasing square root of the immersion time (t 1/2 ) for immersion temperatures lower than 800 C for the whole base steel. On the other hand, for immersion temperatures higher than 800 C, the thickness of the alloy layer on the 0.05 mass% C steel and 0.45 mass% C steels exhibited a negative deviation from the linear relationship. The growth rate constant k decreased as the carbon concentration of the base steel increased up to 0.8 %, above which k had a constant value. The reaction activation energies for the base steel in this study were approximately kj/mol. The alloy layer/base steel interfaces were serrated, and the serration width decreased with increasing carbon concentration of the base steel. In addition, the serration width had a larger value in the immersion temperature range wherein the pro-eutectoid ferrite content in the base steel was larger. KEY WORDS: hot-dip aluminum coating; carbon steel; diffusion; aluminizing; intermetallic compound; alloy layer; interface; heat treatment. 1. Introduction Aluminizing of steel has received considerable attention because of the superior corrosion resistance and reflectivity of aluminized steel, as well as the lower toxicity by inhaling fumes, when compared to galvanizing. In particular, hot-dip aluminizing is widely used in industry as an economical means of coating large steel products. By this treatment, aluminum-rich intermetallic compounds such as Fe 2 Al 5 or FeAl 3 are formed at the steel/aluminum interface by reaction. These intermetallic compounds have high hardness values of (750HV 1000HV 1 11) )and good wear resistance, however they cause cracking and peeling because of their brittle nature. Therefore, control of the growth of these intermetallic compounds is important during the hot-dipping process. The intermetallic compound layers formed by aluminizing have been termed the alloy layer. The alloy layer thickness depends on various factors such as immersion time, temperature, composition of the molten aluminum bath, and the alloying elements present in the base steel. It is well known that the addition of Si, Cu and Be in the aluminum bath decrease the alloy layer thickness ) Many theoretical explanations have been proposed for the kinetics of the alloy layer formation ) However, most previous studies have been confined to pure iron or low carbon steels. Ueda et al. have reported that the presence of C, Si, Mn, Cr, Cr or Ni in the base steel decreases the alloy layer thickness ) The effects of carbon concentration in the base steel are not fully understood because the microstructures of the steels vary, and change with carbon concentration or temperature in the Fe C system. In this study, hot-dip aluminizing of steels with various carbon concentrations was performed in order to investigate the effects of carbon concentration on the formation of the alloy layers during reaction diffusion of aluminum on steels. 2. Experimental Procedures Carbon steels containing 0.05 mass%, 0.45 mass%, and 0.88 mass% carbon were used in this study. Their chemical compositions are given in Table 1. These steels are called C05-steel, C45-steel, and C88-steel, respectively in the following discussion. The steels were annealed at C for 1 h, followed by air cooling to room temperature. Specimens measuring 25 mm 50 mm 2 mm were cut, followed by polishing of the 25 mm 50 mm faces using #320 abrasive paper (mean grain size: 48 mm). Hot-dip aluminum coating was performed by pure aluminum bath, which does not contain a flux or additional elements, in order to investigate the effect of carbon in the base steel on diffusion of aluminum. Industrial pure aluminum with a purity of 99.7 % was melted in a graphite ISIJ

2 Table 1. Chemical composition of the steels in the experiment (mass%). Fig. 1. Cross-sectional micrographs of C05-steel specimens immersed in molten aluminum at temperatures ranging from 700 up to 850 C. crucible (capacity: m 3 ). The degreased specimen was immersed in molten aluminum, followed by lifting at constant speed of approximately 20 mm/s and cooling in air. The immersion times t were 10 s, 300 s, 600 s, s and s, and the temperatures of the aluminum bath T D (immersion temperature) were 700, 750, 800 and 850 C, respectively. In order to examine the alloy layers using optical microscopy formed by the aluminizing, the specimens were cut and the cross-sections were polished, followed by etching using a solution of 5 ml HNO 3 95 ml C 2 H 5 OH. 3. Results and Discussion 3.1. Alloy Layers Formed by Hot-dip Aluminizing Figure 1 shows cross-sectional micrographs of the C05- steel specimens immersed in molten aluminum at different temperatures and with varying immersion times. For all the other cases, an alloy layer was formed between the surface aluminum and the base steel, and the alloy layer/steel interface was serrated. Figure 2 shows cross-sectional micrographs of the C45- steel samples subjected to the aluminizing treatment. The interfaces in the C45-steel specimens are more linear than for the C05-steel specimens. Figure 3 shows similar crosssectional micrographs of C88-steel specimens, which have high carbon concentrations, subjected to the aluminizing treatment. For these C88-steel specimens, the interfaces are more linear as well. In order to identify the intermetallic compounds formed in the alloy layers, the specimens were subjected to X-ray diffraction measurements after polishing down to half the thickness of the alloy layer. Figures 4(a) and 4(b) show the X-ray diffraction data for C05-steel specimens for T D 700 C, t s and T D 850 C, t s, respectively. The results clearly indicate the formation of the intermetallic compound Fe 2 Al 5 over the whole immersion temperature range. Furthermore, it was confirmed that the alloy layers were single phase Fe 2 Al 5 from X-ray diffraction results for the C45-steel and C88-steel specimens as shown in Figs. 4(c), 4(d) and 4(e), 4(f) respectively. These results unambiguously indicate that the coatings formed on the surfaces of the aluminized specimens subjected to the hot-dip aluminizing treatment at immersion temperatures ranging from 700 to 850 C consisted of aluminum and single phase Fe 2 Al 5 alloy layers ISIJ 1888

3 ISIJ International, Vol. 45 (2005), No. 12 Fig. 2. Cross-sectional micrographs of C45-steels immersed in molten aluminum in the temperature range of 700 to 850 C. Fig. 3. Cross-sectional micrographs of C88-steels immersed in molten aluminum in the temperature range of 700 to 850 C ISIJ

4 Fig. 4. X-ray diffraction data for the alloy layers of the hot-dip aluminum-coated steel specimens. The open circle symbol indicate diffraction peaks for Fe 2 Al 5. Fig. 5. Cross-sectional micrograph in wide range of C05-steel immersed at 700 C for 3600 s Effects of Carbon Concentration on Alloy Layer Thickness As shown in Figs.1 3, the thickness of alloy layer in each specimen is greatly different. Therefore, the influence of the coating conditions and C contents on the thickness of alloy layer was investigated. Figure 4 shows a cross-sectional micrograph of the specimen in wide range. The thickness of aluminum layer varies by the approximately twice in the specimen. This is thought to be due to the condition of molten aluminum bath in lifting the specimen and/or the effect of lifting speed. However, the thickness of alloy layer was approximately constant. Seemingly, in the lower immersion temperature, the thickness of aluminum layer tends to increase with increasing the immersion time. Moreover, this tendency is more pronounced in the base steel having higher carbon contents. The kinetics of the alloy layer formation during the aluminizing treatment have been previously explained based on the reaction diffusion at the base steel/alloy layer interface and the molten aluminum/alloy layer interface ) If one assumes that the alloy layer grows by diffusion control of Al atoms and Fe atoms, the thickness of the alloy layer x (in mm) after immersion time t (s) can be expressed by the following equation. 15,17) x kt 1/2...(1) where k is the rate constant (mms 1/2 ). The rate constant k is given by Eq. (2) calculated from the Al concentration at the aluminum/alloy layer interface, C Al1, the Al concentration in the alloy layer at the alloy layer/base steel interface, C Al2, and the Fe concentrations at the alloy layer/base steel interface and the aluminum/alloy layer interface, C Fe1 and C Fe2, respectively. k D C Al1 2 C Al CAl1 Al2 D C Fe1 C Fe C 12 /...(2) Here, D Al and D Fe are the diffusion coefficients for Al and Fe respectively. Thus, the Arrhenius relationship for the rate constant k similar to D is: EA k k 0 exp...(3) RT Here, k 0, E A, R and T are the frequency factor, the reaction activation energy (kj mol 1 ), the gas constant, and the room temperature (K), respectively. Fe1 Fe ISIJ 1890

5 Fig. 7. Arrhenius plot of the rate constant k for the growth of an alloy layer in molten aluminum. The numbers in parentheses are the reaction activation energy, E A values. Fig. 6. Thickness of the alloy layer formed by the hot-dip aluminizing treatment plotted as a function of the square root of the immersion time (t 1/2 ) for the various steel specimens: (a) C05-steel, (b) C45-steel, and (c) C88- steel. Fig. 8. Relationship between the rate constant k and the carbon concentration C for the different steel samples from this study compared with data from literature. Figures 6(a) 6(c) show the relationship between the mean thickness of alloy layer x formed in the coatings of the specimens and the square root of immersion time t 1/2. In the case of C05-steel, as shown in Fig. 6(a), the alloy layer thickness x at immersion temperatures lower than 800 C increases linearly with increasing t 1/2. However, the value of x at immersion temperatures higher than 850 C shows a maximum and then decreases. This tendency, wherein x decreases at longer immersion times, is more remarkable at higher immersion temperatures T D. Yeremenko et al. theoretically explained this negative deviation from the linear relationship as due to the dissolution of the alloy layer into the molten aluminum at higher immersion temperatures. 15) It is conjectured that in this study the dissolution of the specimen into the aluminum bath increases as the immersion temperature rises. In the case of the C45-steel specimens, as shown in Fig. 6(b), the dissolution of the alloy layer decreased when compared to the C05-steel. In contrast, in the case of the C88- steel specimens with the highest carbon concentration, x follows a linear relationship over the whole measured temperature range, as shown in Fig. 6(c). Figure 7 shows the Arrhenius plot of the rate constant k obtained from the slope of the linear relationships in Fig. 6. For the conditions under which the alloy layer dissolved, the rate constant k can be obtained by a collinear approximation of the immersion time t lower than which x has a maximum value. The rate constant k calculated from the results reported by Ueda et al. 21) wherein commercial pure iron, Fe 2mass%C (Fe 2C) and Fe 3mass%C (Fe 3C) specimens were immersed in molten aluminum (purity: 99.8 %), are shown by dashed lines in Fig. 7. As shown in Fig. 7, (ln k) is proportional to 1/T, for all specimens and it decreases with increasing carbon concentration of the base steel. The reaction activation energy, E A, values are approximately kj. The microstructures of the base steels used in this study transform from a ferrite cementite structure (a q) or a ferrite austenite structure (a g), to a single phase of austenite (g) in the immersion temperature range from 700 to 850 C. It is estimated that these transformations with a corresponding change in crystal structure influence the growth of the alloy layers. The linear relationship between (ln k) and 1/T is probably due to the smaller effect of the microstructure of the base steel comparatively on the mean thickness x of alloy layer. The relationship between the rate constant k and the carbon concentration C for the specimens from this study is shown in Fig. 8. The rate constant k decreases rapidly as C increases, up to 0.8 mass%, above which it is almost constant. Ueda et al. suggested that the growth of the alloy layer was prevented by the pro-eutectoid cementite. Based ISIJ

6 caused by the difference in the crystal structure between the ferrite and austenite phases, and the undulation increased with increasing pro-eutectoid ferrite in the base steel. Fig. 9. Relationship between the serration width of the steel/ alloy layer interface and the square root of the immersion time (t 1/2 ). 4. Conclusions Alloy layers formed on the surfaces of steels, which had varying carbon concentrations, by hot-dip aluminizing were investigated with respect to the growth kinetics. The following conclusions were derived from the study. (1) The alloy layers formed by the hot-dip aluminum coating over the whole measured temperature range consisted of a single phase of Fe 2 Al 5. (2) At immersion temperatures lower than 800 C, the thickness of the alloy layer decreased as the immersion temperature increased or the carbon concentration of the base steel decreased. (3) The growth rate of the alloy layers (rate constant k) obtained from the results in the range of shorter immersion times follows a linear relationship, and the reaction activation energy was approximately kj/mol. (4) The dissolution of the specimen material decreased with increasing carbon concentration in the base steel. (5) The rate constant k decreased with increasing carbon concentration of the base steel up to 0.8 %, beyond which it had a constant value. (6) The alloy layer/base steel interface was serrated, and serration width was smaller in the base steel with a higher carbon concentration. In addition, the serration width had a larger value in the immersion temperature range wherein the pro-eutectoid ferrite content in the base steel was larger. on this, it is considered that the alloy layer grew preferentially in the pro-eutectoid ferrite, and consequently, k for the specimens with carbon concentrations lower than 0.8 mass% C increased The Interface of the Alloy Layer and Base Steel In the cross-sectional micrographs shown in Fig. 1, the alloy layer/steel interface was observed to be serrated, whereas the interface was smooth in the specimens with higher carbon concentrations. The serration width W is defined as the difference between the maximum thickness xmax and the minimum thickness x min. Figure 9 shows the serration width plotted as a function of the square root of immersion time t 1/2. For the C05-steel, shown in Fig. 8(a), W increases as t increases in the temperature range of 700 to 800 C, above which it shows a maximum at approximately t 1/2 40. In addition, W at immersion temperatures of 750 and 800 C tends to be large. A similar tendency is observed in the C45-steel specimens. The microstructures of the C05-steel and C45-steel specimens have an (a g) structure in the temperature range of 750 to 800 C. In the temperature of 700 C and 850 C, they have (a q) and g structures. In the C88-steel specimens, for the microstructure which shows a g structure at temperatures above 750 C (shown in Fig. 8(c)), W is approximately 50 mm over the whole temperature range, and it decreases as the immersion temperature rises. Therefore, it is conjectured that the serration width was REFERENCES 1) D. O. Gittings, D. H. Rowland and J. O. Mack: Trans., Am. Soc. Met., 43 (1951), ) G. M. Bedford and J. B. Boustead: Met. Technol., 1 (1974), ) A. Bahadur and O. N. Mohanty: Mater. Trans., JIM, 32 (1991), ) R. W. Richards, R. D. Jones, P. D. Clements and H. Clarke: Int. Mater. Rev., 39 (1994), ) A. Bahadur and O. N. Mohanty: Mater. Trans., JIM, 36 (1995), ) K. Bouche, F. Barbier and A. Coulet: Mater. Sci. Eng. A, 249A (1998), ) H. R. Shahverdi, M. R. Ghomashchi, S. Shabestari and J. Hejazi: 124 (2002), ) Y. -J. Li, J. Wang and X. Holly: Mater. Sci. Technol., 19 (2003), ) L. Yajiang, Z. Yonglan and L. Yuxian: J. Mater. Sci., 30 (1995), ) K. Nishida and T. Narita: J. Jpn. Inst. Met., 35 (1971), ) S. Kobayashi and T. Yakou: Mater. Sci. Eng. A, A338 (2002), ) G. Eggler, W. Auer and H. Kaesche: J. Mater. Sci., 21 (1986), ) N. A. El-Mahallawy, M. A. Taha, M. A. Shady, A. R. El-Sissi, A. N. Attia and W. Reif: Mater. Sci. Technol., 13 (1997), ) S. H. Hwang, J. H. Song and Y.-S. Kim: Mater. Sci. A, A390 (2005), ) V. N. Yeremenko, Y. V. Natanzon and V. I. Dybkov: J. Mater. Sci., 16 (1981), ) G. Eggler, H. Vogel, J. Friedrich and H. Kaesche: Pract. Met., 22 (1985), ) V. I. Dybkov: J. Mater. Sci., 21 (1986), ) V. I. Dybkov: J. Mater. Sci., 25 (1990), ) A. Bouayad, Ch. Gerometta, A. Belkebir and A. Ambari: Mater. Sci. Eng. A, A363 (2003), ) Y. Ueda and M. Ninomi: J. Jpn. Inst. Met., 42 (1978), ) M. Niinomi, Y. Ueda and M. Sano: Mater. Trans., JIM, 23 (1982), ISIJ 1892