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Growth kinetics of Al–Fe intermetallic compounds during annealing treatment of friction stir lap welds
M A TE RI A L S C HA RACT ER I ZA TI O N 90 ( 20 1 4 ) 1 2 1–1 2 6
Available online at www.sciencedirect.com
ScienceDirect www.elsevier.com/locate/matchar
Growth kinetics of Al–Fe intermetallic compounds during annealing treatment of friction stir lap welds M. Movahedia,⁎, A.H. Kokabia , S.M. Seyed Reihania , H. Najafib , S.A. Farzadfarc , W.J. Chengd , C.J. Wangd a
Department of Materials Science and Engineering, Sharif University of Technology, P.O. Box 11365-9466, Azadi Ave., Tehran, Iran Institute of Condensed Matter Physics (ICMP), EPFL, CH-1015 Lausanne, Switzerland c McGill University, Department of Materials Engineering, Montreal, QC H3A 2B2, Canada d Department of Mechanical Engineering, National Taiwan University of Science and Technology, Taipei 10607, Taiwan, ROC b
AR TIC LE D ATA
ABSTR ACT
Article history:
In this study, we explored the growth kinetics of the Al–Fe intermetallic (IM) layer at the
Received 12 December 2013
joint interface of the St-12/Al-5083 friction stir lap welds during post-weld annealing
Received in revised form
treatment at 350, 400 and 450 °C for 30 to 180 min. Optical microscope (OM), field emission
10 January 2014
gun scanning electron microscope (FEG-SEM) and transmission electron microscope (TEM)
Accepted 24 January 2014
were employed to investigate the structure of the weld zone. The thickness and composition of the IM layers were evaluated using image analysis system and electron back-scatter diffraction (EBSD), respectively. Moreover, kernel average misorientation
Keywords:
(KAM) analysis was performed to evaluate the level of stored energy in the as-welded
Friction stir welding
state. The results showed that the growth kinetics of the IM layer was not governed by a
Annealing
parabolic diffusion law. Presence of the IM compounds as well as high stored energy near
Intermetallics
the joint interface of the as-welded sample was recognized to be the origin of the observed
Diffusion
deviation from the parabolic diffusion law.
Transmission electron microscopy
© 2014 Elsevier Inc. All rights reserved.
Kernel average misorientation
1. Introduction Joining steel to aluminum alloys can be employed in producing steel–aluminum dissimilar parts for various industrial fields such as transportation, shipbuilding and aerospace. However, producing strong joints between steel and aluminum alloys by conventional fusion welding processes is challenging due to obstacles resulting from significant difference in the melting points of steel and aluminum alloys as well as the formation of “thick” and brittle Al–Fe intermetallic (IM) compounds at the joint interface. Low heat-input in solid state welding processes such as Friction Stir (FS) welding makes them a more applicable
candidate for aluminum to steel joining [1–3]. Although the thickness of the Al–Fe IM layer (DIM) plays a key role in obtaining high quality joints, available data on the relation between DIM and the joint strength are limited and contradictory. In this regard, some studies have shown that an IM layer with a thickness below a critical value (2 μm [4], 8 μm [5] and 10 μm [6,7]) has no detrimental effect on the joint strength, and even may improve the joint quality. On the contrary, some other works indicate that the formation of IM phases, regardless of their thicknesses, reduces the joint strength [2,8]. An annealing treatment is usually performed after solid state welding of dissimilar metals to promote metallurgical
⁎ Corresponding author. Tel.: + 98 21 66165224; fax: +98 21 66005717. E-mail addresses: [email protected] (M. Movahedi), [email protected] (A.H. Kokabi), [email protected] (S.M. Seyed Reihani), [email protected] (H. Najafi), [email protected] (S.A. Farzadfar), [email protected] (W.J. Cheng), [email protected] (C.J. Wang). 1044-5803/$ – see front matter © 2014 Elsevier Inc. All rights reserved. http://dx.doi.org/10.1016/j.matchar.2014.01.023
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bond at the joint interface, reduce the residual stresses and modify the microstructure of the weld zone. Since the composition and thickness of the IM layer are affected by annealing conditions, it is crucial to control the annealing temperature and time. Consequently, it is the point of interest to investigate the Al–Fe IM growth kinetics during annealing treatment. This can provide some specific knowledge for anticipating DIM as a result of various heat treatment conditions. There are several published studies on the phase recognition and growth kinetics of the Al–Fe IM in the solid iron/liquid aluminum couples [9–17]. Yousaf et al. [14] studied the effects of the compositions of aluminizing bath, aluminizing temperatures (675–950 °C) and dipping times (1–10 min) on the morphology and growth of the IM layer in the carbon steel samples hot-dip aluminized with pure Al and Al–11 wt.% Cu alloy. Their results showed that during the aluminizing process, Al5Fe2 IM layer was developed at the interface of the aluminum and the steel substrate. Additionally, in both melts, Al5Fe2 thickness increased with enhancement of the aluminizing temperature at the range of 675–775 °C, while above this range, it was found constant. Likewise, Al5Fe2 thickness increased rapidly with the increase in the dipping time up to 4 min and beyond 4 min, it increased very slowly. Yousaf et al. [14] did not present any quantitative model for the growth kinetics of the IM layer with the dipping time. The majority of the researches on the solid iron/liquid aluminum couples [9–13,15–17] indicate that a “parabolic” law exists between DIM and the solid/liquid reaction time (t). In fact, the mean thickness of the Al–Fe IM layer has a linear relationship with the square root of time: DIM ¼ k t0:5
ð1Þ
where, k is the parabolic growth rate depending on the interdiffusion coefficient of Al and Fe atoms in the IM phase. A good agreement between the experimental results and the Eq. (1) indicates that the growth rate of the IM layer is governed by the “volume diffusion” of Al and Fe atoms, i.e. volume diffusion is the dominant mechanism contributing to the growth of the IM layer. Furthermore, various Al–Fe IM compounds have been reported to be formed at the interface. The growth kinetics of the Al–Fe IM compounds formed by aluminizing of carbon steel was investigated by Kobayashi and Yakou [15]. They reported that the coating layer consisted of a single phase of Al5Fe2 at the temperatures lower than 1273 K, while the IM compounds of AlFe and AlFe3 were formed at the temperatures higher than 1273 K. They also suggested that the AlFe and AlFe3 IM layers had a diffusion-controlled growth because the thickness of the layers demonstrated a linear relationship with the square root of diffusion time. Springer et al. [16] studied the formation of the Al–Fe IM layer between a low carbon steel and two aluminum alloys (commercially pure aluminum and Al–5 wt.% Si) at the temperature of 675 °C. Their results showed that the total width of the IM layers was governed mainly by the parabolic diffusion controlled growth. Moreover, they stated that adding Si to Al melts resulted in a reduced growth rate of the IM layer as compared to pure Al melts. This effect could be attributed to the incorporation of Si into the IM layer and then decrease in the diffusion rate of aluminum atoms in Al5Fe2, by occupying the vacancies in the c-axis of the crystal structure of Al5Fe2. Cheng
and Wang [17] examined the formation of the IM layer in the aluminide/nickel duplex coating on mild steel by hot-dipping nickel pre-plated mild steel in a molten pure aluminum bath at 670 °C with different immersion times. They showed that as the dipping time increased to 120 s, the nickel layer disappeared and the phase constitution of the IM layer changed, as well. The main part of the IM layer became Al5Fe2 with a small amount of Al3Fe. They also reported that the growth of the Al5Fe2 followed the parabolic law. There are few works on the growth kinetics of the Al–Fe IM at the interface of the solid iron/solid aluminum couples [18–21]. Tricario et al. [20] studied the effects of heat treatment at 100 to 500 °C for 5 to 25 min on the explosion-welded joints. Their results indicated the increase of the Al − Fe IM (identified as Al3Fe and Al5Fe2 on the aluminum side and steel side, respectively) layer thickness due to the increase in both temperature and time. They also concluded that below 500 °C, the temperature was the main factor influencing the growth of the IM phases at the aluminum–steel interface while the effect of time was negligible. However, time showed a strong impact on the IM growth in the samples annealed at 500 °C. The effects of the temperature (220–345 °C) and duration time (15–60 min) during post-weld heat treatment on the lap joints of 5A02 aluminum alloy to 304 stainless steel sheets were investigated by Dong et al. [21]. They employed gas tungsten arc welding with Zn–15%Al flux-cored wire to join the sheets. They reported that the diffusion of Fe and Al into the weld/steel interface controls the thickness of the IM layer and thus, the thickness of the interfacial layer significantly increased when the heat treatment temperature was 345 °C or when the duration time was 60 min. Tricario et al. [20] and Dong et al. [21] did not explore the agreement between the Al–Fe IM growth kinetics with the parabolic diffusion law. Moreover, Naoi and Kajihara [18] investigated the interdiffusion in the Al–Fe couples at 550 to 640 °C for times up to 120 h. They showed that the only phase formed at the interface was Al5Fe2 which exhibited a parabolic growth behavior. Only limited experimental data discussing the effects of the annealing treatment on the growth kinetics of the IM layers at the interface of the aluminum/steel FS welded joints are available in the literature. Springer et al. [19] investigated the effects of the annealing treatment at the range of 200–600 °C for 9–64 min on the formation and growth kinetics of the IM layers in FS ‘butt’ welded samples composed of a low carbon steel and two aluminum alloys (pure Al and Al–5 wt.% Si). Their results indicated that the total thickness of the Al–Fe IM layer was governed by the parabolic growth of the dominant phase (i.e. Al5Fe2). They also reported the formation of an extremely narrow layer of Al3Fe at the interface. Moreover, they noted that variations in the tool rotation and travel speeds, which influence plastic deformation and consequently the amount of the stored energy, could affect atomic diffusion and IM phase nucleation at the interface during annealing. Therefore, the main purpose of the present study is to provide a new insight into the growth kinetics of the Al–Fe IM thickness during the post-weld annealing of the Al-5083/ St-12 FS ‘lap’ welds.
M A TE RI A L S C HA RACT ER I ZA TI O N 90 ( 20 1 4 ) 1 2 1–1 2 6
2. Material and Methods Al-5083 (Al–4.45 wt.% Mg, annealed with an average grain size of ~24 μm) and St-12 (Fe–< 0.1 wt.% C, as received with a rolled structure) sheets with the thicknesses of 3 mm and 1 mm, respectively, were used as the base materials. The joint design is illustrated in Fig. 1-a. The welding tool made of H-13 tool steel had a shoulder with diameter of 16 mm and a conical pin with height of 3 mm, tip diameter of 3 mm and root diameter of 4.5 mm. The tilt angle of the tool axis was fixed at 3°. At first, the rotating pin was gently pushed to the Al-5083 sheet until the pin tip entered 0.2–0.3 mm into the St-12 sheet. Then, the tool started to move along the joint. The travel and rotation speeds of the welding tool were chosen to be 23 cm·min−1 and 750 rpm, respectively, to lower the welding heat input and avoid the formation of any IM layer at the joint interface. Thus, it is reasonable to assume that the growth of the Al–Fe IM layers occurred merely during post-weld annealing. To prepare the samples, the welds were cut perpendicular to the welding direction and annealed at 350, 400 and 450 °C for 30 to 180 min. All investigations on the IM layers and microstructural features were performed inside the region shown in Fig. 1-b. To observe and characterize the Al–
Fig. 1 – a) Joint design (dimensions in millimeter) and b) transverse section of the welded joint. Investigations were made on the regions inside the rectangle, and c) FEG-SEM image and EBSD point analysis results from the IM layers.
123
Fe IM phases, a field emission gun scanning electron microscope (FEG-SEM) and an electron back-scatter diffractometer (EBSD) were employed. The electron back-scatter patterns (EBSPs) were indexed by INCA Crystal software via comparison to the simulated EBSPs. The simulated patterns were created according to the crystallographic structures obtained from the Inorganic Crystal Structure Database (ICSD). Six Kikuchi reference lines and 1.5° angular tolerance were employed to index the EBSPs. An image analysis system was used to measure the total thickness of the IM layers. Forty measurements were taken for each sample and the average value was reported. The microstructures of the regions adjacent to the joint interface were observed by optical and transmission electron microscopy (TEM), using a Philips CM-20 equipment (200 kV) and also a FEI Tecnai Osiris Scanning TEM (300 kV). To prepare TEM foils, the specimens were first cut from the cross sections of the weld interface. The slices were subsequently thinned by mechanical polishing down to ~40 μm, followed by ion milling (3–4 keVAr+) for final thinning to electron transparency.
3. Results and Discussion Fig. 1-c shows the results of EBSD point analysis on the joint interface of the sample annealed at 400 °C for 180 min. Two intermetallics, Al5Fe2 and Al3Fe, were detected. Al5Fe2 was the main part of the IM layer and was formed next to the St-12 base metal. Due to the small thickness of the Al5Fe2 and Al3Fe layers (especially that of Al3Fe), it was not possible to accurately measure the thickness of each individual layer. The total thickness of both Al5Fe2 and Al3Fe layers was therefore reported as DIM. DIM at the joint interface of the samples annealed at various temperatures and times is presented in Fig. 2. At the annealing condition of 350 °C/30 min, a discontinuous IM layer was formed at the interface. However, with increasing annealing time up to 90 and 180 min, a continuous and very narrow (<1 μm) IM layer was formed. Annealing treatment at 400 and 450 °C led to an increase in DIM, as the thickness reached up to 7.8 and 10.2 μm for the annealing conditions of 400 °C/180 min and 450 °C/180 min, respectively. As previously mentioned, if the growth of the IM layer with annealing time (at a given temperature) is governed by volume diffusion, a parabolic relationship between DIM and annealing time (t) will exist. In other words, there must be a linear relationship between DIM and t0.5. Fig. 2-a exhibits the variation of DIM with t0.5 as well as the “linear” trend lines with the highest R-squared (closest R2 to 1) which pass through both the data points and the origin (0,0). It should be emphasized that the IM thickness at the interface of the as-welded sample was zero and the fitting curve should therefore pass through the origin. Referring to Fig. 2-a, the R2 values of the trend lines are −0.67, 0.767 and 0.813 for annealing temperatures of 350, 400 and 450 °C, respectively. The low values of R2 indicate the misfit between the linear trend lines and the experimental data. It could therefore be deduced that the growth of the IM layers at the joint interface is not governed by volume diffusion. Yamada et al. [22,23] and Tanaka et al. [24] showed that if grain growth occurs during annealing, the decrease in the grain boundary diffusion due to the reduction in boundary
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annealing. A time exponent equal or close to 0.5 is then expected. In order to evaluate the degree of the grain boundary diffusion in the investigated samples, the microstructures of the St-12 and Al-5083 stirred zones near the joint interface of the as-welded samples were explored, and the results are presented in Fig. 3-a and b. Both Al and steel have extremely small grain sizes (4 μm and < 100 nm, respectively), and therefore large boundary areas. A significant atomic diffusion across grain boundaries and a time exponent close to 0.5 are thus expected, especially at lower temperatures. However, a comparison between the time exponents at 450 and 350 °C reveals that such a trend is not present in the results obtained in this work, i.e. the time exponent decreases from 0.352 at 450 °C to 0.195 at 350 °C. Moreover, the time exponent at 400 °C, i.e. 1.334, is significantly larger than 0.5. It could suggest that the growth of the IM layer is not governed by volume and/or grain boundary diffusion at this temperature and some other mechanisms may also be operative. Further microstructural analysis performed on the FS welded zone revealed that the following phenomena could contribute, in part, to the observed deviation from the parabolic diffusion law:
Fig. 2 – Intermetallic layer thickness vs. a) square root of annealing time and b) annealing time. The linear and power trend lines were plotted in (a) and (b), respectively. area will lead to a continuous alteration in the diffusion status and deviation from the parabolic law. In that case, a time exponent below 0.5 is obtained. A power-law curve was therefore fitted to the data acquired in this work to evaluate the mentioned mechanism (Fig. 2-b). The exponents of 0.195, 1.334, and 0.352 were obtained for the annealing temperatures of 350, 400 and 450 °C, respectively. As can be seen, the exponents at 350 and 450 °C are below 0.5 and the exponent at 400 °C is much higher than 0.5. It is also shown [22–24] that at lower annealing temperatures, the contribution of grain boundary diffusion is more prominent than that of volume diffusion. As well, the rate of grain growth could diminish, and boundary area would be approximately constant during
(i) Formation of the intermetallic compounds near the joint interface of the as-welded samples (and inside the St-12 sheet): Fig. 4-a shows the IM compounds forming a stringer-like pattern. During welding, the rotating pin stirs and mixes the layers of aluminum and steel at the joint interface. Due to the relatively high temperature at the weld zone and also short diffusion distance between the layers, the Al–Fe IM compounds are formed as a result of the interdiffusion of Al and Fe atoms. During annealing, once the growing layers of the Al–Fe IM at the joint interface reach to these stringer-like intermetallics, the growth kinetics can abruptly change. (ii) Presence of high stored energy in the joint region of the as-welded samples: If there is a high stored energy in the joint region of the as-welded samples, this stored energy may be gradually released during post-weld annealing. The conditions of
Fig. 3 – The grains structures of a) Al-5083 (optical microscope) and b) St-12 (TEM) stirred zones exactly near the joint interface of the as-welded samples.
M A TE RI A L S C HA RACT ER I ZA TI O N 90 ( 20 1 4 ) 1 2 1–1 2 6
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growth of Al3Fe layer does not obey the parabolic law. Hence, the growth of the total thickness, containing both Al5Fe2 and Al3Fe compounds, is expected to deviate from a parabolic law. It should be emphasized that the mentioned phenomenon is not specific to FS welded samples, and it can be observed in other Al–Fe couples [9,10].
4. Conclusion In this research, the growth kinetics of the Al–Fe intermetallic layers at the joint interface of the St-12/Al-5083 friction stir lap welds during post-weld annealing at the temperature range of 350 to 450 °C for 30 to 180 min was investigated. The following conclusions can be drawn from the present work:
Fig. 4 – a) TEM elemental map of layered structure near the joint interface and inside St-12 stirred zone. The atomic percentages were determined using energy dispersive spectroscopy (EDS). KAM maps of b) Al-5083 and c) St-12 regions inside the stirred zones of the as-welded samples.
nucleation and growth of the IM compounds, e.g. in the case of Al3Fe [19], can therefore be altered during annealing, resulting in a deviation from the parabolic diffusion law. However, the question is whether a friction-stirred region in which dynamic recrystallization occurs [25], could contain a high amount of stored energy? In order to evaluate the level of stored energy in the as-welded state, kernel average misorientation (KAM) analysis was performed, and the corresponding maps for Al-5083 and St-12 stirred zones are presented in Fig. 4-b and c, respectively. The KAM value corresponds to the average misorientation angle between the crystallographic orientations of a given point and that of its neighbor points. In other words, the KAM value shows the local strain value, i.e. high KAM values (larger than 3° [26]) indicate a high level of stored energy [27]. As can be seen in Fig. 4-b, a high level of stored energy is present in Al-5083 (orange and red zones), which could affect the nucleation and growth rate of the Al3Fe compound during annealing. (iii) Non-parabolic growth of Al3Fe based on literature data: Bouche et al. [10] and Bouayad et al. [9] showed that Al5Fe2 and Al3Fe compounds were formed during aluminizing of carbon steels. They also found a linear relationship between the thickness of Al5Fe2 and t0.5 as well as between the thickness of Al3Fe and t, where t was the aluminizing time. This indicates that the
1 The intermetallic layer at the joint interface was composed of two intermetallic compounds, Al5Fe2 and Al3Fe. Al5Fe2 was the main fraction of the intermetallic layer and was formed next to the St-12 base metal. However, Al3Fe was formed adjacent to the Al-5083 base metal with fewer thickness compared to the Al5Fe2. 2 Given the graphs between DIM and square root of annealing time (t0.5) for each temperature and the misfit between the linear trend lines and the experimental data, it was concluded that the growth of the IM layers at the joint interface was not governed by volume diffusion. 3 The growth of the IM layer was not also governed by grain boundary diffusion. However, both Al and steel stirred zones near the joint interface of the as-welded samples had extremely small grain sizes (4 μm and <100 nm, respectively), and therefore large boundary areas. 4 Detailed microstructural analysis showed that the following phenomena could result in the observed deviation from the parabolic diffusion law: (i) Formation of the intermetallic compounds near the joint interface of the as-welded samples, (ii) presence of high stored energy in the joint region of the as-welded samples, and (iii) non-parabolic growth of Al3Fe.
Acknowledgment The authors would like to thank Dr. Duncan T.L. Alexander for TEM analysis of the samples at Interdisciplinary Centre for Electron Microscopy (CIME), EPFL, Lausanne, Switzerland.
REFERENCES [1] Lee WB, Schmuecker M, Mercardo UA, Biallas G, Jung SB. Interfacial reaction in steel–aluminum joints made by friction stir welding. Scr Mater 2006;55:355–8. [2] Tanaka T, Morishige T, Hirata T. Comprehensive analysis of joint strength for dissimilar friction stir welds of mild steel to aluminum alloys. Scr Mater 2009;61:756–9. [3] Qiu R, Shi H, Zhang K, Tu Y, Iwamoto C, Satonaka S. Interfacial characterization of joint between mild steel and aluminum alloy welded by resistance spot welding. Mater Charact 2010;61:684–8.
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[4] Movahedi M, Kokabi AH, Seyed Reihani SM, Najafi H. Effect of tool travel and rotation speeds on weld zone defects and joint strength of aluminum–steel lap joints made by friction stir welding. Sci Technol Weld Joining 2012;17:162–7. [5] Bozzi S, Helbert-Etter AL, Baudin T, Criqui B, Kerbiguet JG. Intermetallic compounds in Al6016/IF–steel friction stir spot welds. Mater Sci Eng A 2010;527:4505–9. [6] Padmanabhan R, Oliveira MC, Menezes LF. Deep drawing of aluminium–steel tailor-welded blanks. Mater Des 2008;29:154–60. [7] Lee CY, Choi DH, Yeon YM, Jung SB. Dissimilar friction stir spot welding of low carbon steel and Al–Mg alloy by formation of IMCs. Sci Technol Weld Joining 2009;14:216–20. [8] Movahedi M, Kokabi AH, Seyed Reihani SM. Investigation on the bond strength of Al-1100/St-12 roll bonded sheets, optimization and characterization. Mater Des 2011;32:3143–9. [9] Bouayad A, Gerometta C, Belkebir A, Ambari A. Kinetic interactions between solid iron and molten aluminium. Mater Sci Eng A 2003;363:53–61. [10] Bouché K, Barbier F, Coulet A. Intermetallic compound layer growth between solid iron and molten aluminium. Mater Sci Eng A 1998;249:167–75. [11] Shahverdi HR, Ghomashchi MR, Shabestari S, Hejazi J. Microstructural analysis of interfacial reaction between molten aluminium and solid iron. J Mater Proc Technol 2002;124:345–52. [12] Agudo L, Eyidi D, Schmaranzer C, Arenholz E, Jank N, Bruckner J, et al. Intermetallic Fe × Aly-phases in a steel/Al-alloy fusion weld. J Mater Sci 2007;42:4205–14. [13] Cheng WJ, Wang CJ. Effect of silicon on the formation of intermetallic phases in aluminide coating on mild steel. Intermetallics 2011;19:1455–60. [14] Yousaf M, Iqbal J, Ajmal M. Variables affecting growth and morphology of the intermetallic layer (Fe2Al5). Mater Charact 2011;62:517–25. [15] Kobayashi S, Yakou T. Control of intermetallic compound layers at interface between steel and aluminum by diffusion-treatment. Mater Sci Eng A 2002;338:44–53. [16] Springer H, Kostka A, Payton EJ, Raabe D, Kaysser-Pyzalla A, Eggeler G. On the formation and growth of intermetallic
[17]
[18]
[19]
[20]
[21]
[22]
[23]
[24]
[25] [26]
[27]
phases during interdiffusion between low-carbon steel and aluminum alloys. Acta Mater 2011;59:1586–600. Cheng WJ, Wang CJ. Characterization of intermetallic layer formation in aluminide/nickel duplex coating on mild steel. Mater Charact 2012;69:63–70. Naoi D, Kajihara M. Growth behavior of Fe2Al5 during reactive diffusion between Fe and Al at solid-state temperatures. Mater Sci Eng A 2007;459:375–82. Springer H, Kostka A, dos Santos JF, Raabe D. Influence of intermetallic phases and Kirkendall-porosity on the mechanical properties of joints between steel and aluminium alloys. Mater Sci Eng A 2011;528:4630–42. Tricarico L, Spina R, Sorgente D, Brandizzi M. Effects of heat treatments on mechanical properties of Fe/Al explosion-welded structural transition joints. Mater Des 2009;30:2693–700. Dong H, Liao C, Yang L, Dong C. Effects of post-weld heat treatment on dissimilar metal joint between aluminum alloy and stainless steel. Mater Sci Eng A 2012;550:423–8. Yamada T, Miura K, Kajihara M, Kurokawa N, Sakamoto K. Formation of intermetallic compound layers in Sn/Au/Sn diffusion couple during annealing at 433 K. J Mater Sci 2004;39:2327–34. Yamada T, Miura K, Kajihara M, Kurokawa N, Sakamoto K. Kinetics of reactive diffusion between Au and Sn during annealing at solid-state temperatures. Mater Sci Eng A 2005;390:118–26. Tanaka Y, Kajihara M, Watanabe Y. Growth behavior of compound layers during reactive diffusion between solid Cu and liquid Al. Mater Sci Eng A 2007;445–446:355–63. Mishra RS, Ma ZY. Friction stir welding and processing. Mater Sci Eng R 2005;50:1–78. Farzadfar SA, Martin É, Sanjari M, Essadiqi E, Wells MA, Yue S. On the deformation, recrystallization and texture of hot-rolled Mg–2.9Y and Mg–2.9Zn solid solution alloys—a comparative study. Mater Sci Eng A 2012;534:209–19. Britton TB, Birosca S, Preuss M, Wilkinson AJ. Electron backscatter diffraction study of dislocation content of a macrozone in hot-rolled Ti–6Al–4V alloy. Scr Mater 2010;62:639–42.
Available online at www.sciencedirect.com
ScienceDirect www.elsevier.com/locate/matchar
Growth kinetics of Al–Fe intermetallic compounds during annealing treatment of friction stir lap welds M. Movahedia,⁎, A.H. Kokabia , S.M. Seyed Reihania , H. Najafib , S.A. Farzadfarc , W.J. Chengd , C.J. Wangd a
Department of Materials Science and Engineering, Sharif University of Technology, P.O. Box 11365-9466, Azadi Ave., Tehran, Iran Institute of Condensed Matter Physics (ICMP), EPFL, CH-1015 Lausanne, Switzerland c McGill University, Department of Materials Engineering, Montreal, QC H3A 2B2, Canada d Department of Mechanical Engineering, National Taiwan University of Science and Technology, Taipei 10607, Taiwan, ROC b
AR TIC LE D ATA
ABSTR ACT
Article history:
In this study, we explored the growth kinetics of the Al–Fe intermetallic (IM) layer at the
Received 12 December 2013
joint interface of the St-12/Al-5083 friction stir lap welds during post-weld annealing
Received in revised form
treatment at 350, 400 and 450 °C for 30 to 180 min. Optical microscope (OM), field emission
10 January 2014
gun scanning electron microscope (FEG-SEM) and transmission electron microscope (TEM)
Accepted 24 January 2014
were employed to investigate the structure of the weld zone. The thickness and composition of the IM layers were evaluated using image analysis system and electron back-scatter diffraction (EBSD), respectively. Moreover, kernel average misorientation
Keywords:
(KAM) analysis was performed to evaluate the level of stored energy in the as-welded
Friction stir welding
state. The results showed that the growth kinetics of the IM layer was not governed by a
Annealing
parabolic diffusion law. Presence of the IM compounds as well as high stored energy near
Intermetallics
the joint interface of the as-welded sample was recognized to be the origin of the observed
Diffusion
deviation from the parabolic diffusion law.
Transmission electron microscopy
© 2014 Elsevier Inc. All rights reserved.
Kernel average misorientation
1. Introduction Joining steel to aluminum alloys can be employed in producing steel–aluminum dissimilar parts for various industrial fields such as transportation, shipbuilding and aerospace. However, producing strong joints between steel and aluminum alloys by conventional fusion welding processes is challenging due to obstacles resulting from significant difference in the melting points of steel and aluminum alloys as well as the formation of “thick” and brittle Al–Fe intermetallic (IM) compounds at the joint interface. Low heat-input in solid state welding processes such as Friction Stir (FS) welding makes them a more applicable
candidate for aluminum to steel joining [1–3]. Although the thickness of the Al–Fe IM layer (DIM) plays a key role in obtaining high quality joints, available data on the relation between DIM and the joint strength are limited and contradictory. In this regard, some studies have shown that an IM layer with a thickness below a critical value (2 μm [4], 8 μm [5] and 10 μm [6,7]) has no detrimental effect on the joint strength, and even may improve the joint quality. On the contrary, some other works indicate that the formation of IM phases, regardless of their thicknesses, reduces the joint strength [2,8]. An annealing treatment is usually performed after solid state welding of dissimilar metals to promote metallurgical
⁎ Corresponding author. Tel.: + 98 21 66165224; fax: +98 21 66005717. E-mail addresses: [email protected] (M. Movahedi), [email protected] (A.H. Kokabi), [email protected] (S.M. Seyed Reihani), [email protected] (H. Najafi), [email protected] (S.A. Farzadfar), [email protected] (W.J. Cheng), [email protected] (C.J. Wang). 1044-5803/$ – see front matter © 2014 Elsevier Inc. All rights reserved. http://dx.doi.org/10.1016/j.matchar.2014.01.023
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bond at the joint interface, reduce the residual stresses and modify the microstructure of the weld zone. Since the composition and thickness of the IM layer are affected by annealing conditions, it is crucial to control the annealing temperature and time. Consequently, it is the point of interest to investigate the Al–Fe IM growth kinetics during annealing treatment. This can provide some specific knowledge for anticipating DIM as a result of various heat treatment conditions. There are several published studies on the phase recognition and growth kinetics of the Al–Fe IM in the solid iron/liquid aluminum couples [9–17]. Yousaf et al. [14] studied the effects of the compositions of aluminizing bath, aluminizing temperatures (675–950 °C) and dipping times (1–10 min) on the morphology and growth of the IM layer in the carbon steel samples hot-dip aluminized with pure Al and Al–11 wt.% Cu alloy. Their results showed that during the aluminizing process, Al5Fe2 IM layer was developed at the interface of the aluminum and the steel substrate. Additionally, in both melts, Al5Fe2 thickness increased with enhancement of the aluminizing temperature at the range of 675–775 °C, while above this range, it was found constant. Likewise, Al5Fe2 thickness increased rapidly with the increase in the dipping time up to 4 min and beyond 4 min, it increased very slowly. Yousaf et al. [14] did not present any quantitative model for the growth kinetics of the IM layer with the dipping time. The majority of the researches on the solid iron/liquid aluminum couples [9–13,15–17] indicate that a “parabolic” law exists between DIM and the solid/liquid reaction time (t). In fact, the mean thickness of the Al–Fe IM layer has a linear relationship with the square root of time: DIM ¼ k t0:5
ð1Þ
where, k is the parabolic growth rate depending on the interdiffusion coefficient of Al and Fe atoms in the IM phase. A good agreement between the experimental results and the Eq. (1) indicates that the growth rate of the IM layer is governed by the “volume diffusion” of Al and Fe atoms, i.e. volume diffusion is the dominant mechanism contributing to the growth of the IM layer. Furthermore, various Al–Fe IM compounds have been reported to be formed at the interface. The growth kinetics of the Al–Fe IM compounds formed by aluminizing of carbon steel was investigated by Kobayashi and Yakou [15]. They reported that the coating layer consisted of a single phase of Al5Fe2 at the temperatures lower than 1273 K, while the IM compounds of AlFe and AlFe3 were formed at the temperatures higher than 1273 K. They also suggested that the AlFe and AlFe3 IM layers had a diffusion-controlled growth because the thickness of the layers demonstrated a linear relationship with the square root of diffusion time. Springer et al. [16] studied the formation of the Al–Fe IM layer between a low carbon steel and two aluminum alloys (commercially pure aluminum and Al–5 wt.% Si) at the temperature of 675 °C. Their results showed that the total width of the IM layers was governed mainly by the parabolic diffusion controlled growth. Moreover, they stated that adding Si to Al melts resulted in a reduced growth rate of the IM layer as compared to pure Al melts. This effect could be attributed to the incorporation of Si into the IM layer and then decrease in the diffusion rate of aluminum atoms in Al5Fe2, by occupying the vacancies in the c-axis of the crystal structure of Al5Fe2. Cheng
and Wang [17] examined the formation of the IM layer in the aluminide/nickel duplex coating on mild steel by hot-dipping nickel pre-plated mild steel in a molten pure aluminum bath at 670 °C with different immersion times. They showed that as the dipping time increased to 120 s, the nickel layer disappeared and the phase constitution of the IM layer changed, as well. The main part of the IM layer became Al5Fe2 with a small amount of Al3Fe. They also reported that the growth of the Al5Fe2 followed the parabolic law. There are few works on the growth kinetics of the Al–Fe IM at the interface of the solid iron/solid aluminum couples [18–21]. Tricario et al. [20] studied the effects of heat treatment at 100 to 500 °C for 5 to 25 min on the explosion-welded joints. Their results indicated the increase of the Al − Fe IM (identified as Al3Fe and Al5Fe2 on the aluminum side and steel side, respectively) layer thickness due to the increase in both temperature and time. They also concluded that below 500 °C, the temperature was the main factor influencing the growth of the IM phases at the aluminum–steel interface while the effect of time was negligible. However, time showed a strong impact on the IM growth in the samples annealed at 500 °C. The effects of the temperature (220–345 °C) and duration time (15–60 min) during post-weld heat treatment on the lap joints of 5A02 aluminum alloy to 304 stainless steel sheets were investigated by Dong et al. [21]. They employed gas tungsten arc welding with Zn–15%Al flux-cored wire to join the sheets. They reported that the diffusion of Fe and Al into the weld/steel interface controls the thickness of the IM layer and thus, the thickness of the interfacial layer significantly increased when the heat treatment temperature was 345 °C or when the duration time was 60 min. Tricario et al. [20] and Dong et al. [21] did not explore the agreement between the Al–Fe IM growth kinetics with the parabolic diffusion law. Moreover, Naoi and Kajihara [18] investigated the interdiffusion in the Al–Fe couples at 550 to 640 °C for times up to 120 h. They showed that the only phase formed at the interface was Al5Fe2 which exhibited a parabolic growth behavior. Only limited experimental data discussing the effects of the annealing treatment on the growth kinetics of the IM layers at the interface of the aluminum/steel FS welded joints are available in the literature. Springer et al. [19] investigated the effects of the annealing treatment at the range of 200–600 °C for 9–64 min on the formation and growth kinetics of the IM layers in FS ‘butt’ welded samples composed of a low carbon steel and two aluminum alloys (pure Al and Al–5 wt.% Si). Their results indicated that the total thickness of the Al–Fe IM layer was governed by the parabolic growth of the dominant phase (i.e. Al5Fe2). They also reported the formation of an extremely narrow layer of Al3Fe at the interface. Moreover, they noted that variations in the tool rotation and travel speeds, which influence plastic deformation and consequently the amount of the stored energy, could affect atomic diffusion and IM phase nucleation at the interface during annealing. Therefore, the main purpose of the present study is to provide a new insight into the growth kinetics of the Al–Fe IM thickness during the post-weld annealing of the Al-5083/ St-12 FS ‘lap’ welds.
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2. Material and Methods Al-5083 (Al–4.45 wt.% Mg, annealed with an average grain size of ~24 μm) and St-12 (Fe–< 0.1 wt.% C, as received with a rolled structure) sheets with the thicknesses of 3 mm and 1 mm, respectively, were used as the base materials. The joint design is illustrated in Fig. 1-a. The welding tool made of H-13 tool steel had a shoulder with diameter of 16 mm and a conical pin with height of 3 mm, tip diameter of 3 mm and root diameter of 4.5 mm. The tilt angle of the tool axis was fixed at 3°. At first, the rotating pin was gently pushed to the Al-5083 sheet until the pin tip entered 0.2–0.3 mm into the St-12 sheet. Then, the tool started to move along the joint. The travel and rotation speeds of the welding tool were chosen to be 23 cm·min−1 and 750 rpm, respectively, to lower the welding heat input and avoid the formation of any IM layer at the joint interface. Thus, it is reasonable to assume that the growth of the Al–Fe IM layers occurred merely during post-weld annealing. To prepare the samples, the welds were cut perpendicular to the welding direction and annealed at 350, 400 and 450 °C for 30 to 180 min. All investigations on the IM layers and microstructural features were performed inside the region shown in Fig. 1-b. To observe and characterize the Al–
Fig. 1 – a) Joint design (dimensions in millimeter) and b) transverse section of the welded joint. Investigations were made on the regions inside the rectangle, and c) FEG-SEM image and EBSD point analysis results from the IM layers.
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Fe IM phases, a field emission gun scanning electron microscope (FEG-SEM) and an electron back-scatter diffractometer (EBSD) were employed. The electron back-scatter patterns (EBSPs) were indexed by INCA Crystal software via comparison to the simulated EBSPs. The simulated patterns were created according to the crystallographic structures obtained from the Inorganic Crystal Structure Database (ICSD). Six Kikuchi reference lines and 1.5° angular tolerance were employed to index the EBSPs. An image analysis system was used to measure the total thickness of the IM layers. Forty measurements were taken for each sample and the average value was reported. The microstructures of the regions adjacent to the joint interface were observed by optical and transmission electron microscopy (TEM), using a Philips CM-20 equipment (200 kV) and also a FEI Tecnai Osiris Scanning TEM (300 kV). To prepare TEM foils, the specimens were first cut from the cross sections of the weld interface. The slices were subsequently thinned by mechanical polishing down to ~40 μm, followed by ion milling (3–4 keVAr+) for final thinning to electron transparency.
3. Results and Discussion Fig. 1-c shows the results of EBSD point analysis on the joint interface of the sample annealed at 400 °C for 180 min. Two intermetallics, Al5Fe2 and Al3Fe, were detected. Al5Fe2 was the main part of the IM layer and was formed next to the St-12 base metal. Due to the small thickness of the Al5Fe2 and Al3Fe layers (especially that of Al3Fe), it was not possible to accurately measure the thickness of each individual layer. The total thickness of both Al5Fe2 and Al3Fe layers was therefore reported as DIM. DIM at the joint interface of the samples annealed at various temperatures and times is presented in Fig. 2. At the annealing condition of 350 °C/30 min, a discontinuous IM layer was formed at the interface. However, with increasing annealing time up to 90 and 180 min, a continuous and very narrow (<1 μm) IM layer was formed. Annealing treatment at 400 and 450 °C led to an increase in DIM, as the thickness reached up to 7.8 and 10.2 μm for the annealing conditions of 400 °C/180 min and 450 °C/180 min, respectively. As previously mentioned, if the growth of the IM layer with annealing time (at a given temperature) is governed by volume diffusion, a parabolic relationship between DIM and annealing time (t) will exist. In other words, there must be a linear relationship between DIM and t0.5. Fig. 2-a exhibits the variation of DIM with t0.5 as well as the “linear” trend lines with the highest R-squared (closest R2 to 1) which pass through both the data points and the origin (0,0). It should be emphasized that the IM thickness at the interface of the as-welded sample was zero and the fitting curve should therefore pass through the origin. Referring to Fig. 2-a, the R2 values of the trend lines are −0.67, 0.767 and 0.813 for annealing temperatures of 350, 400 and 450 °C, respectively. The low values of R2 indicate the misfit between the linear trend lines and the experimental data. It could therefore be deduced that the growth of the IM layers at the joint interface is not governed by volume diffusion. Yamada et al. [22,23] and Tanaka et al. [24] showed that if grain growth occurs during annealing, the decrease in the grain boundary diffusion due to the reduction in boundary
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annealing. A time exponent equal or close to 0.5 is then expected. In order to evaluate the degree of the grain boundary diffusion in the investigated samples, the microstructures of the St-12 and Al-5083 stirred zones near the joint interface of the as-welded samples were explored, and the results are presented in Fig. 3-a and b. Both Al and steel have extremely small grain sizes (4 μm and < 100 nm, respectively), and therefore large boundary areas. A significant atomic diffusion across grain boundaries and a time exponent close to 0.5 are thus expected, especially at lower temperatures. However, a comparison between the time exponents at 450 and 350 °C reveals that such a trend is not present in the results obtained in this work, i.e. the time exponent decreases from 0.352 at 450 °C to 0.195 at 350 °C. Moreover, the time exponent at 400 °C, i.e. 1.334, is significantly larger than 0.5. It could suggest that the growth of the IM layer is not governed by volume and/or grain boundary diffusion at this temperature and some other mechanisms may also be operative. Further microstructural analysis performed on the FS welded zone revealed that the following phenomena could contribute, in part, to the observed deviation from the parabolic diffusion law:
Fig. 2 – Intermetallic layer thickness vs. a) square root of annealing time and b) annealing time. The linear and power trend lines were plotted in (a) and (b), respectively. area will lead to a continuous alteration in the diffusion status and deviation from the parabolic law. In that case, a time exponent below 0.5 is obtained. A power-law curve was therefore fitted to the data acquired in this work to evaluate the mentioned mechanism (Fig. 2-b). The exponents of 0.195, 1.334, and 0.352 were obtained for the annealing temperatures of 350, 400 and 450 °C, respectively. As can be seen, the exponents at 350 and 450 °C are below 0.5 and the exponent at 400 °C is much higher than 0.5. It is also shown [22–24] that at lower annealing temperatures, the contribution of grain boundary diffusion is more prominent than that of volume diffusion. As well, the rate of grain growth could diminish, and boundary area would be approximately constant during
(i) Formation of the intermetallic compounds near the joint interface of the as-welded samples (and inside the St-12 sheet): Fig. 4-a shows the IM compounds forming a stringer-like pattern. During welding, the rotating pin stirs and mixes the layers of aluminum and steel at the joint interface. Due to the relatively high temperature at the weld zone and also short diffusion distance between the layers, the Al–Fe IM compounds are formed as a result of the interdiffusion of Al and Fe atoms. During annealing, once the growing layers of the Al–Fe IM at the joint interface reach to these stringer-like intermetallics, the growth kinetics can abruptly change. (ii) Presence of high stored energy in the joint region of the as-welded samples: If there is a high stored energy in the joint region of the as-welded samples, this stored energy may be gradually released during post-weld annealing. The conditions of
Fig. 3 – The grains structures of a) Al-5083 (optical microscope) and b) St-12 (TEM) stirred zones exactly near the joint interface of the as-welded samples.
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growth of Al3Fe layer does not obey the parabolic law. Hence, the growth of the total thickness, containing both Al5Fe2 and Al3Fe compounds, is expected to deviate from a parabolic law. It should be emphasized that the mentioned phenomenon is not specific to FS welded samples, and it can be observed in other Al–Fe couples [9,10].
4. Conclusion In this research, the growth kinetics of the Al–Fe intermetallic layers at the joint interface of the St-12/Al-5083 friction stir lap welds during post-weld annealing at the temperature range of 350 to 450 °C for 30 to 180 min was investigated. The following conclusions can be drawn from the present work:
Fig. 4 – a) TEM elemental map of layered structure near the joint interface and inside St-12 stirred zone. The atomic percentages were determined using energy dispersive spectroscopy (EDS). KAM maps of b) Al-5083 and c) St-12 regions inside the stirred zones of the as-welded samples.
nucleation and growth of the IM compounds, e.g. in the case of Al3Fe [19], can therefore be altered during annealing, resulting in a deviation from the parabolic diffusion law. However, the question is whether a friction-stirred region in which dynamic recrystallization occurs [25], could contain a high amount of stored energy? In order to evaluate the level of stored energy in the as-welded state, kernel average misorientation (KAM) analysis was performed, and the corresponding maps for Al-5083 and St-12 stirred zones are presented in Fig. 4-b and c, respectively. The KAM value corresponds to the average misorientation angle between the crystallographic orientations of a given point and that of its neighbor points. In other words, the KAM value shows the local strain value, i.e. high KAM values (larger than 3° [26]) indicate a high level of stored energy [27]. As can be seen in Fig. 4-b, a high level of stored energy is present in Al-5083 (orange and red zones), which could affect the nucleation and growth rate of the Al3Fe compound during annealing. (iii) Non-parabolic growth of Al3Fe based on literature data: Bouche et al. [10] and Bouayad et al. [9] showed that Al5Fe2 and Al3Fe compounds were formed during aluminizing of carbon steels. They also found a linear relationship between the thickness of Al5Fe2 and t0.5 as well as between the thickness of Al3Fe and t, where t was the aluminizing time. This indicates that the
1 The intermetallic layer at the joint interface was composed of two intermetallic compounds, Al5Fe2 and Al3Fe. Al5Fe2 was the main fraction of the intermetallic layer and was formed next to the St-12 base metal. However, Al3Fe was formed adjacent to the Al-5083 base metal with fewer thickness compared to the Al5Fe2. 2 Given the graphs between DIM and square root of annealing time (t0.5) for each temperature and the misfit between the linear trend lines and the experimental data, it was concluded that the growth of the IM layers at the joint interface was not governed by volume diffusion. 3 The growth of the IM layer was not also governed by grain boundary diffusion. However, both Al and steel stirred zones near the joint interface of the as-welded samples had extremely small grain sizes (4 μm and <100 nm, respectively), and therefore large boundary areas. 4 Detailed microstructural analysis showed that the following phenomena could result in the observed deviation from the parabolic diffusion law: (i) Formation of the intermetallic compounds near the joint interface of the as-welded samples, (ii) presence of high stored energy in the joint region of the as-welded samples, and (iii) non-parabolic growth of Al3Fe.
Acknowledgment The authors would like to thank Dr. Duncan T.L. Alexander for TEM analysis of the samples at Interdisciplinary Centre for Electron Microscopy (CIME), EPFL, Lausanne, Switzerland.
REFERENCES [1] Lee WB, Schmuecker M, Mercardo UA, Biallas G, Jung SB. Interfacial reaction in steel–aluminum joints made by friction stir welding. Scr Mater 2006;55:355–8. [2] Tanaka T, Morishige T, Hirata T. Comprehensive analysis of joint strength for dissimilar friction stir welds of mild steel to aluminum alloys. Scr Mater 2009;61:756–9. [3] Qiu R, Shi H, Zhang K, Tu Y, Iwamoto C, Satonaka S. Interfacial characterization of joint between mild steel and aluminum alloy welded by resistance spot welding. Mater Charact 2010;61:684–8.
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[4] Movahedi M, Kokabi AH, Seyed Reihani SM, Najafi H. Effect of tool travel and rotation speeds on weld zone defects and joint strength of aluminum–steel lap joints made by friction stir welding. Sci Technol Weld Joining 2012;17:162–7. [5] Bozzi S, Helbert-Etter AL, Baudin T, Criqui B, Kerbiguet JG. Intermetallic compounds in Al6016/IF–steel friction stir spot welds. Mater Sci Eng A 2010;527:4505–9. [6] Padmanabhan R, Oliveira MC, Menezes LF. Deep drawing of aluminium–steel tailor-welded blanks. Mater Des 2008;29:154–60. [7] Lee CY, Choi DH, Yeon YM, Jung SB. Dissimilar friction stir spot welding of low carbon steel and Al–Mg alloy by formation of IMCs. Sci Technol Weld Joining 2009;14:216–20. [8] Movahedi M, Kokabi AH, Seyed Reihani SM. Investigation on the bond strength of Al-1100/St-12 roll bonded sheets, optimization and characterization. Mater Des 2011;32:3143–9. [9] Bouayad A, Gerometta C, Belkebir A, Ambari A. Kinetic interactions between solid iron and molten aluminium. Mater Sci Eng A 2003;363:53–61. [10] Bouché K, Barbier F, Coulet A. Intermetallic compound layer growth between solid iron and molten aluminium. Mater Sci Eng A 1998;249:167–75. [11] Shahverdi HR, Ghomashchi MR, Shabestari S, Hejazi J. Microstructural analysis of interfacial reaction between molten aluminium and solid iron. J Mater Proc Technol 2002;124:345–52. [12] Agudo L, Eyidi D, Schmaranzer C, Arenholz E, Jank N, Bruckner J, et al. Intermetallic Fe × Aly-phases in a steel/Al-alloy fusion weld. J Mater Sci 2007;42:4205–14. [13] Cheng WJ, Wang CJ. Effect of silicon on the formation of intermetallic phases in aluminide coating on mild steel. Intermetallics 2011;19:1455–60. [14] Yousaf M, Iqbal J, Ajmal M. Variables affecting growth and morphology of the intermetallic layer (Fe2Al5). Mater Charact 2011;62:517–25. [15] Kobayashi S, Yakou T. Control of intermetallic compound layers at interface between steel and aluminum by diffusion-treatment. Mater Sci Eng A 2002;338:44–53. [16] Springer H, Kostka A, Payton EJ, Raabe D, Kaysser-Pyzalla A, Eggeler G. On the formation and growth of intermetallic
[17]
[18]
[19]
[20]
[21]
[22]
[23]
[24]
[25] [26]
[27]
phases during interdiffusion between low-carbon steel and aluminum alloys. Acta Mater 2011;59:1586–600. Cheng WJ, Wang CJ. Characterization of intermetallic layer formation in aluminide/nickel duplex coating on mild steel. Mater Charact 2012;69:63–70. Naoi D, Kajihara M. Growth behavior of Fe2Al5 during reactive diffusion between Fe and Al at solid-state temperatures. Mater Sci Eng A 2007;459:375–82. Springer H, Kostka A, dos Santos JF, Raabe D. Influence of intermetallic phases and Kirkendall-porosity on the mechanical properties of joints between steel and aluminium alloys. Mater Sci Eng A 2011;528:4630–42. Tricarico L, Spina R, Sorgente D, Brandizzi M. Effects of heat treatments on mechanical properties of Fe/Al explosion-welded structural transition joints. Mater Des 2009;30:2693–700. Dong H, Liao C, Yang L, Dong C. Effects of post-weld heat treatment on dissimilar metal joint between aluminum alloy and stainless steel. Mater Sci Eng A 2012;550:423–8. Yamada T, Miura K, Kajihara M, Kurokawa N, Sakamoto K. Formation of intermetallic compound layers in Sn/Au/Sn diffusion couple during annealing at 433 K. J Mater Sci 2004;39:2327–34. Yamada T, Miura K, Kajihara M, Kurokawa N, Sakamoto K. Kinetics of reactive diffusion between Au and Sn during annealing at solid-state temperatures. Mater Sci Eng A 2005;390:118–26. Tanaka Y, Kajihara M, Watanabe Y. Growth behavior of compound layers during reactive diffusion between solid Cu and liquid Al. Mater Sci Eng A 2007;445–446:355–63. Mishra RS, Ma ZY. Friction stir welding and processing. Mater Sci Eng R 2005;50:1–78. Farzadfar SA, Martin É, Sanjari M, Essadiqi E, Wells MA, Yue S. On the deformation, recrystallization and texture of hot-rolled Mg–2.9Y and Mg–2.9Zn solid solution alloys—a comparative study. Mater Sci Eng A 2012;534:209–19. Britton TB, Birosca S, Preuss M, Wilkinson AJ. Electron backscatter diffraction study of dislocation content of a macrozone in hot-rolled Ti–6Al–4V alloy. Scr Mater 2010;62:639–42.