Effect of silicon on the microstructure and growth kinetics of intermetallic phases formed during hot-dip aluminizing of ferritic steel

Accepted Manuscript Effect of silicon on the microstructure and growth kinetics of intermetallic phases formed during hot-dip aluminizing of ferritic steel

B. Lemmens, H. Springer, I. De Graeve, J. De Strycker, D. Raabe, K. Verbeken PII: DOI: Reference:

S0257-8972(17)30288-8 doi: 10.1016/j.surfcoat.2017.03.040 SCT 22209

To appear in:

Surface & Coatings Technology

Received date: Revised date: Accepted date:

1 December 2016 22 February 2017 18 March 2017

Please cite this article as: B. Lemmens, H. Springer, I. De Graeve, J. De Strycker, D. Raabe, K. Verbeken , Effect of silicon on the microstructure and growth kinetics of intermetallic phases formed during hot-dip aluminizing of ferritic steel. The address for the corresponding author was captured as affiliation for all authors. Please check if appropriate. Sct(2017), doi: 10.1016/j.surfcoat.2017.03.040

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ACCEPTED MANUSCRIPT Effect of silicon on the microstructure and growth kinetics of intermetallic phases formed during hot-dip aluminizing of ferritic steel

Department of Materials Science and Engineering, Ghent University (UGent), Tech Lane Ghent

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B. Lemmens1,2 , H. Springer3, I. De Graeve1,2, J. De Strycker4, D. Raabe3, K. Verbeken1,*

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Science Park – Campus A, Technologiepark 903, B-9052 Zwijnaarde, Belgium Research group of Electrochemical and Surface Engineering (SURF), Vrije Universiteit Brussel

ArcelorMittal Global R&D Gent, J.F.Kennedylaan 3, B-9060 Zelzate, Belgium

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Max-Planck-Institut für Eisenforschung GmbH, 40237 Düsseldorf, Germany

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(VUB), Pleinlaan 2, B-1050 Brussels, Belgium

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1 Introduction

Dissimilar joining of aluminium (Al) and iron (Fe) based materials is of significant importance

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for engineering design. For structural applications the lower mass density of Al alloys is brought

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together with the higher strength and stiffness of Fe based steels, for example to reduce the weight of transportation systems in order to improve their performance and meet ever more stringent environmental legislation constraints [1]. Another important application is combining the superior corrosion resistance of Al – especially at elevated temperatures [2] – to the high strength of the steel substrates. This is widely realized by coating techniques, of which the most versatile, economic and commonly applied technique is dipping the steel substrate in a molten Al bath, referred to as hot-dip aluminizing (HDA) [3]. The main challenge in HDA, and dissimilar Al-steel joining in general, is the formation of inherently brittle intermetallic phases (IMPs) at the 1

ACCEPTED MANUSCRIPT interface between Al alloys and steel, as they are the main cause of cracking and spalling of the coating during forming operations [4]. In the case of pure Al coatings, the intermetallic seam consists of two different IMPs, namely the θ phase (Fe4Al13) adjacent to Al and the η phase (Fe2Al5) in between θ phase and steel. The η phase has been found to be the rate controlling component, rendering it of critical importance for the joint performance [5-9].

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Heumann and Dittrich [10] were among the first to report on the parabolic growth kinetics

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of the intermetallic layer, assuming it to solely consist of η phase. Denner and Jones [11] also indicated a linear dependence for the √𝑡 kinetics, but noted deviations from this relationship at

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longer reaction times. Eggeler et al. [12] detected the presence of a thin θ phase layer in between the Al and η layer. Fe saturation of the Al bath were shown to shift the interface between Al and

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intermetallic layer more outward compared to the pure Al, and the activation energies ΔH of for

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the growth of the reaction layer were about 155 kJ mol–1 for pure Al and 104 kJ mol–1 for Feenriched Al melts, respectively [7]. When maximum and not mean thicknesses were investigated,

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the values of the activation energies were about 20 kJ mol–1 lower.

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The main strategy to reduce the thickness of the IMP layers is the addition of silicon (Si) to the Al bath as it – additional to lowering the melting temperature and viscosity of the bath –

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suppresses the growth of the η phase [4,13,14]. Numerous investigations have been conducted on

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the underlying microstructural phenomena, but the mechanisms behind this growth retardation are not yet fully understood [13,15-21]. Nicholls [16] proposed that Si atoms might interact with the open crystallographic arrangement of the η phase, and consequently decrease the growth kinetics. Alternative hypotheses suggest a reduced activity of Al [15,21], the formation of a ternary phase acting as a diffusion barrier [20], an influence on the activity coefficient of the Si, or a reduction of the activation energy for the η phase formation [18,19]. Recently the presence of Si-enrichments at grain and phase boundaries of the η and θ phase could be detected by atom 2

ACCEPTED MANUSCRIPT probe tomography investigations, these enrichments might also affect the interdiffusion processes and thus growth kinetics of the IMP seam [22]. Apart from the reduction in η phase thickness, Si additions also cause a change in morphology of the interface from finger-like to a more planar structure, and the formation of ternary phases [18,23]. A reasonable understanding of the different types of Al-Fe-Si ternary phases is reflected by literature, but the phases have quite

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large and closely positioned homogeneity ranges. Despite some variation in nomenclature they

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are generally quoted as τ1 to τ10 [18,23]. Additionally the influence of the substrate must not be underestimated, the addition of Si to the steel [24] or whether it is austenitic can have an effect as

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well [25, 26].

It thus becomes clear that the Si-induced growth suppression of IMPs in dissimilar joining

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of Al and steel is a complex scenario of influences and consequences. Comparing and evaluating

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literature data in this field, however, is difficult since significantly different dipping temperatures and times, bath and substrate compositions, as well as procedures for calculating growth kinetics

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have been used in the different investigations. Consequently, a comprehensive and systematic

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study of the influence of varying Si concentrations and interdiffusion conditions on the build-up and growth kinetics of the IMP seam is of high interest. Therefore, this work aims to contribute to

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the understanding on the role of Si in the growth kinetics as well as the microscopic

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characteristics of the different intermetallic phases formed. After hot-dipping in Al baths with systematically varied Si contents, temperatures and dipping time, the thickness and morphology are analysed to elucidate the microstructural parameters which control the reaction kinetics of the intermetallic phase formation.

2 Materials and methods

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ACCEPTED MANUSCRIPT Hot-dipping experiments were performed with a Rhesca hot-dip simulator under an atmosphere of N2-5%H2 to prevent oxidation of the steel. DC06 low carbon European grade steel sheets (Fe0.01C-0.015Si-0.1Mn-0.01Cr-0.018Ni-0.05Al-0.01Cu-0.04Ti) of 150 by 100 mm and a thickness of around 1 mm were dipped in aluminium baths containing a Si contents of 0.3, 1, 3, 5 and 10 wt. %, respectively. A content of up to 0.3 wt. % Si in such a bath is currently considered in

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industrial applications as "pure" aluminium, as such this will be termed here likewise. After pre-

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heating the steel substrates at 800 °C for 60 s, several dipping times of 3, 60 and 300 s were applied. The dipping temperatures were varied between 670 and 725 °C, depending on the Si

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concentration the liquidus temperature decreases, as such lower temperatures could be applied. After withdrawal from the Al(-Si) baths the now coated sheets were cooled down to room

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temperature by a flux of N2 gas resulting in a cooling rate of approximately of 10 ° s–1.

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Cross sections of the aluminised steel samples were cut by spark erosion, ground and polished with standard metallographic techniques to a 1 µm finish. Scanning electron microscopy

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(SEM) was performed with a FEG SEM Quanta 450 equipped with a field emission gun and a

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TSL electron backscatter diffraction (EBSD) system. BSE characterization was performed with Carl Zeiss SEM Merlin. After the EBSD measurements a clean-up of the raw data was

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performed: clean-up with neighbour CI correlation with minimum CI of 0.1 and a neighbour

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orientation clean-up (level 3, tolerance 5 and minimum CI 0.1). The growth kinetics were investigated by thickness measurements on SEM-pictures with the ImageJ® software.

3 Results 3.1 Microstructure Representative EBSD results are shown in Fig. 1 as phase maps of the technically pure Al, Al-3, Al-5 and Al-10 wt.% Si experiments dipped for 60 s at temperatures of about 685 °C. Blue represents Al, yellow the θ phase, green the η phase and red Fe α (i.e. the steel substrate). The 4

ACCEPTED MANUSCRIPT ternary phase τ5 which is formed at higher Si contents is shown in light blue. Image quality data are superimposed on the phase maps to indicate grain boundaries. The η phase is thicker than the θ phase and contains the typical finger-like structures for low Si contents, while with increasing Si content the corresponding layer becomes thinner and the interface between the η phase and the steel converts to a more planar morphology. In the phase maps, the presence of some θ phase

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inclusions in the Al bath stemming from Fe-saturation of the Al bath, can be observed. From Al5wt.% Si on to higher Si-contents the ternary phase τ5 can be observed, placed between the Al

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and θ phase. This phase seems to occasionally grow in an irregular manner, as observed when

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comparing the left side to the right side for the Al-5wt.% Si case. For the 10 wt.% Si samples, the θ phase has nearly completely disappeared, and a clear τ5 layer is present, with an occasional

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irregularly grown τ5 grain (Fig. 2).

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In Fig. 3, a BSE image of pure HDA steel dipped for 3 s is depicted where on top the aluminium layer is visible in black, while the serrated θ layer with a thickness of around 2 µm is

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present below. Underneath the θ layer, the thicker η layer is observed. When zooming in on the θ

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layer (Fig. 3a) micro twins are observed as indicated by white arrows. Within the η phase at high magnification (Fig. 3b), a nano-scaled and regular wavy contrast is observed which might be

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linked to chemical variations and/or crystallographic ordering phenomena such as a

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superstructure. The latter might be corroborated by the presence of satellite diffraction spots observed in TEM diffractions of the η phase in previous works [18]. While twinning is a known feature in the structure of the θ phase [27-29], the wavy structure appears to have not been reported and systematically investigated to the author’s knowledge, and is the subject of future investigations. Fig. 4 shows corresponding results for experiments with Al-3 wt.% Si. Additional to the phases present in the pure Al experiments, the ternary τ1 phases (white) can be observed in the η phase. In Fig. 4a, the micro twinning in the θ phase can be observed more clearly while the wavy 5

ACCEPTED MANUSCRIPT features present in the η phase are more difficult to recognize (see Fig. 4b). When 10 wt. % Si was added to the Al-bath (Fig. 5.), the τ5-phase nearly completely replaces the θ phase, of which only small islands remain with micro twins (Fig. 5a). In the η phase the wavy features remain present (Fig. 5b).

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3.2 Growth kinetics

The effect of Si on the thickness of the intermetallic seam is summarized in Fig. 6. It becomes

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clear that the thickness of the η layer is decreasing while the θ (and τ5 for higher Si contents)

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remains rather stable. There is a linear relationship between the layer thickness and square root of time kinetics according to 𝑑 = 𝑘√𝑡 from which the k-values can be determined. These k-values 𝛥𝐻

), T the absolute temperature (in K) and ΔH the apparent activation energy for layer growth (in J

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are of the Arrhenius-type 𝑘~ 𝑒𝑥𝑝 (− 𝑅𝑇 ) with R being the universal gas constant (in J K–1 mol–

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mol-1), from which the ΔH values can be obtained. The plot to calculate the k-values for the η layer is shown in Fig. 7(a) for the different Al-Si alloys. The error bars indicate the standard

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deviation, which is rather large for the case of Al-0 wt.% Si, but reasonable in light of the widely

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known finger-like growth morphology characteristic of this phase [7,10,18]. The thickness of the η phase clearly decreases with increasing Si content, as does the standard deviation due to the

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more planar interface. In Fig. 7(b), the plot to calculate k-values is plotted for the θ phase. When the thickness values for experiments with Al-10wt.%Si are plotted, the τ5-phase was combined with that of the θ phase, as these two components cannot reliably be distinguished in SEM pictures. The ragged θ/Al interface and irregular τ5 growth diminish the accuracy for growth calculations of the θ phase further. The thickness of the θ and τ5 layer was not influenced by the addition of Si. For this reason, it was chosen to conduct the investigations on the growth kinetics exclusively on the η layer. In Fig. 7(c) the thickness of the η layer is plotted as a function of the 6

ACCEPTED MANUSCRIPT dipping temperature for 60s dipping time. Interestingly, the influence of the dipping temperature on the thickness of the η layer becomes smaller with increasing Si-content. For Al-10wt.%Si, the thickness of the η layer even seems to be decreasing with increasing bath temperature. The Arrhenius plot of the rate constants for the η phase is presented in Fig. 7(d). The apparent activation energies deduced from the respective slopes are summarized in Table I, where a

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decrease could be observed with rising Si content, and a negative value for the Al-10wt.%Si case.

4 Discussion

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For pure Al we observe an activation energy for η layer growth of around 224 kJ mol-1 which is slightly higher but still in good agreement with previously reported values of 190 kJ mol-1 [18]

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and 207 kJ mol-1 [19]. From Table I it can be observed that with increasing Si content the

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activation energy decreases. This dependence was also reported previously in literature, but only calculated from literature data [18] or for baths up to 3 wt.% Si [19]. A value for the Al-10wt.%Si

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case was not reported before and was found to be negative. Negative values of activation energies

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are possible when the reaction is for example comprised of several reactions (which in total conform the Arrhenius behaviour), where the enthalpic requirements may be predicting a faster

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growth, but the entropic considerations may also play an important role [30].

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Since the apparent activation energy for the Al-10wt.% Si case is negative, it causes the rate of the reaction to decrease with increasing temperature. This was also observed in Fig. 7(c) for the Al-10wt.% Si case which could explain the peculiar case of interdiffusion below the liquidus temperature of the Al-Si alloys with its opposite effect on the intermetallic seam, where the presence of Si was increasing the IMP layer thickness instead of reducing it as previously observed by Springer et al. [18]. This casts doubt onto the atomistic explanation of Nicholls [16] who attributed the decrease in growth of the η to a filling of the structural vacancies in the η phase by Si atoms. These structural vacancies in the η phase were assumed to accelerate the 7

ACCEPTED MANUSCRIPT growth along the c-axis. Additionally, the Si content within the η phase would be thought to increase with reducing growth rate, however, the EDX measurements reveal no rise in Si content in the η phase and remain at a level of around 2 at.%. Another hypothesis to explain the growth reduction in the η phase is a reduction of the activity coefficient of Al, but this was only valid in a limited composition range [13] and

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therefore this is not further considered here. Other possible explanations stated in literature which

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could influence the growth of the η layer are, Si enrichments at the grain and phase boundaries [22] or the formation of a ternary phase [20] which could act as a diffusion barrier. The presence

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of ternary phases could influence the growth of the η phase, either by its nucleation or by acting as a diffusion barrier. Here, the ternary phase τ1 was found at an addition of 3 wt.% Si while τ5

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was observed in the EBSD maps from 5 wt.% Si on, although a thin nanoscale layer could have

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already been present before which was not observed due to resolution constraints. The previously reported presence of Si enrichments in the grain and phase boundaries of the θ and η phase could

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be seen as nano-scaled ternary phases and act as a diffusion barrier reducing the growth rate. The

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τ1 phase was also found present in the η phase when experiments below the liquidus temperature of Al were performed [18]. As such the τ1 phase nucleation might not have an influence on the

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growth acceleration or reduction. On the other hand, in these lower temperature regime

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experiments, no τ5 phase was observed but a τ6 phase layer. As such the formation and the subsequently diffusion enhancer or barrier of the different ternary phases could have an influence, but due to their abundant presence the exact ground is difficult to pinpoint. As the apparent activation energy is negative, it can be assumed that the growth of the η phase is not build up of one reaction but several reactions. From this work however it cannot be seen what these reactions would be due to the complexity of the overall structure. Despite the fact that the growth kinetics describe clearly the retardation, the exact mechanism remains unresolved. The complex microstructural build-up of the reaction zone is 8

ACCEPTED MANUSCRIPT difficult to assess due to its hard and brittle components with similarity in chemical composition. As such more research is required to fully unravel the mechanism and the influence of Si which is behind the growth retardation of the η phase. Low temperature investigations and nano-scale analysis on the presence of the τ5 phase in lower Si cases as well as on its irregular growth should

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be conducted to further clarify this mechanism.

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5. Conclusion

The influence of Si addition to the Al bath on the intermetallic phase formation during HDA was

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investigated for temperatures between 670 and 725 °C and various dipping times. From the obtained results following conclusions can be drawn:

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1) New insights into the microstructure and growth kinetics of the intermetallic phases formed by

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hot-dip aluminizing could be obtained with EBSD and SEM investigations. 2) For lower Si contents, the IMP seam consisted of the major η phase and a thin θ phase. The η

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phase showed a nano-scaled regular structure of wave-like contrast while in the θ phase some

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microtwinning was observed. With increasing Si concentration τ1 and τ5 ternary phases appeared. τ1 was present in the η phase as inclusions while the τ5 layer was observed between the Al and θ

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layer from Al-5wt.% Si on.

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3) Apart from the additional phases, Si also influences the growth kinetics of the η phase, which can be described by parabolic laws. At temperatures around 685 °C, the addition of Si to aluminium bath results in a reduction of η phase thickness. Arrhenius plots could give insights on the activation energy for the growth of the η phase. The investigation revealed a reduction in apparent activation energy with increasing Si content, starting from around 224 kJ mol-1 for technically pure Al to around 140 kJ mol-1 for Al-1 and 3 wt.% Si and even -72 kJ mol-1 for Al10wt.% Si.

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ACCEPTED MANUSCRIPT 4) For the Al-10wt.% Si case the thickness of the η layer seems to decrease with increasing dipping temperature, this could explain the peculiar case of interdiffusion below the liquidus temperature of the Al-Si alloys with its opposite effect on the intermetallic seam, where the presence of Si was increasing the IMP layer thickness instead of reducing it.

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Acknowledgements

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Katja Angenendt and Monika Nellessen are gratefully acknowledged for their support in the BSE investigations. The authours also want to thank Linsey Lapeire and Aurélie Laureys for their

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support in the EBSD investigations.

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ACCEPTED MANUSCRIPT References [1] T.A. Barnes, I.R. Pashby, Joining techniques for aluminium spaceframes used in automobiles Part I - solid and liquid phase welding, J. Mater. Proces. Technol. 99 (2000) 62-71. [2] C.-J. Wang, S.-M. Chen, The high-temperature oxidation behavior of hot-dipping Al–Si coating on low carbon steel, Surf. Coat. Technol. 200 (2006) 6601–6605.

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ACCEPTED MANUSCRIPT [13] M. V. Akdeniz, A.O. Mekhrabov, T. Yilmaz, THE ROLE OF Si ADDITION ON THE INTERFACIAL INTERACTION IN Fe-AL DIFFUSION LAYER, Scr. Mater. 31 (1994) 12, 1723-1728. [14] G. Eggeler, W. Auer, H. Kaesche, On the influence of silicon on the growth of the alloy layer during hot dip aluminizing, J. Mater. Sci. 21(1986) 3348-3350. [15] S. Shankar, D. Apelian, Die soldering: mechanism of the interface reaction between molten

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ACCEPTED MANUSCRIPT coating on mild steel, Intermetal. 19 (2011) 1455-1460. [24] I. Infante Danzo, K. Verbeken, Y. Houbaert, Microstructure of hot dip coated Fe-Si steels, Thin Sol. Films 520 (2011) 1638-1644. [25] E. Frutos, J.L. González-Carrasco, C. Capdevila, J.A. Jiménez, Y. Houbaert, Development of hard intermetallic coatings on austenitic stainless steel by hot dipping in an Al-Si alloy, Surf. Coat. Technol. 203 (2009) 2916-2920.

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ACCEPTED MANUSCRIPT Tables Table I: Apparent activation energies for the formation of the η layer derived from the Arrhenius plots for the different Al-Si baths Activaton energy (kJ mol-1 )

Al-0wt.%Si

224

Al-1wt.%Si

142

Al-3wt.%Si

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Al-10wt.%Si

-72

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Si content (wt.%)

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ACCEPTED MANUSCRIPT List of figure captions Fig.1: EBSD phase maps of HDA dipped in pure Al (left), Al-3wt.%Si (middle), Al-5 wt.% Si (top right) and Al-10 wt.% Si (bottom right) for 60s around 685°C. The image quality of the different measurements was overlayed. Fig. 2: Irregular growth of the τ5 – phase in the Al-layer when the steel was dipped for 320s in an Al-10wt.%Si bath at 725°C.

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Fig.3: Overview of the intermetallic phases observed with BSE images in HDA steel dipped in pure aluminium. In (a) the arrows indicate the twins while in (b) the wavy structure in the η phase is shown.

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Fig.4: Overview of the intermetallic phases observed with BSE images in HDA steel dipped in Al-3wt.% Si. In (a) the arrows show the twins present in the θ phase and in (b) they indicate the wavy structure present in the η phase.

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Fig. 5: Overview of the intermetallic phases observed with BSE in HDA steel coated with Al-10 wt.% Si. In (a) the θ phase islands between the τ5 and η phase can be observed. (b) shows the presence of the wavy structure in the η phase. Fig. 6: Influence of Si on the thickness of the intermetallic layers dipped for 60s around 685°C

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Fig. 7: (a) Thickness of the η layer as a function of the square root of time for the different Al-Si alloys (b) Thickness of the θ (combined τ5 for higher Si contents) layer as a function of the square root of time (c) Evolution of the thickness of the η phase with dipping temperature for the different Al-Si alloys (d) Arrhenius plot of rate constants for the η phase

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ACCEPTED MANUSCRIPT Highlights Microstructure of hot-dip aluminized steel with rising Si content was investigated

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Si addition caused a reduction in η phase thickness and formation of ternary phases

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In the θ and η phase were respectively micro twins and a wavy structure observed

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Growth kinetics were investigated by dipping at various temperatures and times

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Activation energy for the η phase decreased with rising Si content in the Al bath

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