- Email: [email protected]
Nanomechanics of Mg–Al intermetallic compounds
Surface & Coatings Technology 204 (2010) 2118–2122
Contents lists available at ScienceDirect
Surface & Coatings Technology j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s u r f c o a t
Nanomechanics of Mg–Al intermetallic compounds M.-X. Zhang a,b,⁎, H. Huang a, K. Spencer b, Y.-N. Shi c a b c
School of Mechanical and Mining Engineering, The University of Queensland, Brisbane Qld 4072, Australia ARC Centre of Excellence for Design in Light Metals, Australia Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China
a r t i c l e
i n f o
Article history: Received 6 July 2009 Accepted in revised form 22 November 2009 Available online 29 November 2009 Keywords: Magnesium alloy Intermetallic compound Nanoindentation Kinetic metallization Cold spray
a b s t r a c t Cold spraying of pure Al powder on a pure Mg substrate together with subsequent post-spray annealing treatment produced Mg17Al12 (β-phase) and Mg2Al3 (γ-phase) intermetallic layers on the surface of the substrate. These layers showed significantly better nanomechanical properties, including the reduced elastic modulus and nanohardness, which were determined using nanoindentation, than commercial purity Mg and AZ91 alloys. Combined with their improved corrosion resistance, it is believed that both the γ-phase and the β-phase layers can provide effective protection of Mg alloys from wear and corrosion. The effect of postspray annealing process on the formation of thick, uniform and dense intermetallic layers on pure Mg substrate was also investigated. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Low wear resistance and high corrosive activity of Mg alloys significantly limit their structural applications, even though the low densities of these alloys has attracted increasing interests in automotive and aeronautical industries. A number of previous studies [1–6] have showed that surface alloying treatment with Al, or so called aluminising, offers an effective approach in improving both wear and corrosion resistance due to the increase in volume fraction of the Mg17Al12 (β-phase) in the surface area. As one of the two equilibrium Mg–Al intermetallic compounds, the β-phase not only has higher hardness, which may lead to an increase in the wear resistance, but can also act as an anodic barrier and thus, inhibit the overall corrosion of the alloy [7,8]. Hence, the layer with higher volume fraction of β-phase will have better surface durability and the best result can be achieved with a continuous intermetallic coating. However, obtaining continuous intermetallic layers on an Mg substrate is difficult using techniques like packed powder diffusion. Recently Spencer and Zhang [9] reported a new technique to generate such coatings on Mg–Al–Zn alloys through post-spray annealing of Al cold sprayed coatings. The intermetallic coatings exhibits improved hardness and corrosion resistance compared to the substrate [9]. However, the Zn content in the substrate may affect the diffusion process of Al in Mg, and then affect the formation of intermetallic coatings. This was not discussed and the coatings were not fully characterised in the previous study [9]. In the present work, in order
⁎ Corresponding author. Tel./fax: +61 7 3346 8709; +61 7 33467015. E-mail address: [email protected] (M.-X. Zhang). 0257-8972/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2009.11.031
to eliminate the effect of Zn, a pure Mg substrate has been used so that only binary Mg–Al intermetallic compound coatings can be produced through reaction diffusion with Al. The effect of post-spray annealing on the formation of the intermetallic layers has also been investigated. Furthermore, although similar intermetallic compounds have been reported in an Mg–Al diffusion couple [10], their mechanical and chemical behaviour was not evaluated. In addition, nanoindentation test was used to characterise the nanomechanical properties, including the reduced modulus and nanohardness, of the β-phase and γ-phase (Mg2Al3) formed on a pure Mg substrate. 2. Experimental Commercial purity Mg was used as the substrate material. Pure Al powder with an average particle size of 15 µm was deposited on the 20 × 20 × 8 mm substrates using an Inovati Kinetic Metallization (KM) system at 125 °C with helium as carrying gas at a pressure of 620 kPa in order to increase the impact velocity. KM is similar in principle to conventional cold spray with the distinction that it uses a convergent nozzle to accelerate the gas to sonic velocity (supersonic velocity is used in conventional cold spray) [11]. The mass flow rate of the powder mixtures injected into the gas stream was 15 g min− 1. The nozzle standoff distance from the substrate was 12 mm and the traverse speed was 50 mm min− 1. The post-spray annealing treatments were carried out within a temperature range from 360 °C to 430 °C for 24 h to investigate the effect of annealing temperatures. In order to understand the influence of annealing time on the formation of intermetallic layers, the post-spray treatment was also carried out at 413 °C for various times. This selection of 413 °C is because this temperature is the typical solution treatment temperature for AZ91
M.-X. Zhang et al. / Surface & Coatings Technology 204 (2010) 2118–2122
alloy and maximum Al concentration can be obtained in the solid solution without melting due to the formation of eutectic structure. After annealing, the portion of the Al coating that remained unreacted was readily removed, which was followed by further polishing of the intermetallic layers with the γ-phase layer on the top and the β-phase layer adjacent to the substrate. To expose the β-phase layer, the γ-phase layer has to be removed as well. Nanoindentation tests were performed on the polished β and γ layer using a Hysitron Triboindenter®. A threesided Berkovich indenter with tip radius of 50 nm was used. The indentation load was fixed at 1000 µN and the loading/unloading rate remained constant at 100 μN per second. There was a stabilized time of 5 s between the loading and unloading stages. An atomic force microscope (AFM) was used to examine the surface topographies prior to and after each indentation. The reduced elastic modulus Er and hardness nHV were calculated from the load–displacement curves [12]. Er was determined upon the slope (S) of the unloading curve using the following equation, which is in fact a measure of the elastic property of the surface [13]. Er =
pffiffiffi πS pffiffiffi 2 A
where A is the projected area after the indentation. For comparison, the nanomechanical properties of the pure Mg substrate and the widely used AZ91D alloy after solution treatment at 413 °C for 6 h
2119
followed by ageing at 168 °C for 6 h were also measured. The thickness of the intermetallic layers was measured from the cross sections of samples using optical microscopy. In order to identify the phases formed at the coating/substrate interface, X-ray diffraction was conducted in a Bruker D8 Advance diffractometer with a Cu-Kα target. The operation voltage was 40 V and the scanning speed was 0.5°/min with step size of 0.05°. 3. Results and discussion Micrographs of typical intermetallic layers before and after removing the unreacted portion of the Al KM coating are shown in Fig. 1(a) and (b). According to the Mg–Al binary phase diagram the top layer is the γ-phase (Mg2Al3) and the second layer is the β-phase (Mg17Al12). There is a thin layer of supersaturated solid solution Al in Mg below the β layer, but it is not visible in the optical micrographs. Figs. 2(a) and (b) show the X-ray diffraction spectra taken from γ-phase layer and β-phase layer, respectively. Indexing the spectra confirmed the phases. To keep the figure clear, it only shows indices of the major peaks in the figures. In Fig. 2 (a) a few β-phase peaks can also be observed due to the penetration of the X-ray through the thin areas of the γ layer. Similarly, Mg peaks also appear in Fig. 2(b) due to the same reason. The variations of thickness of the intermetallic layers with the post-spray annealing time (from 5 to 50 h) are shown in Fig. 3(a), which indicates the growth rate of the layers. All the three curves
Fig. 1. Typical optical micrographs showing both the γ-phase layer and the β-phase layer before removing the un-reacted Al deposition (a) and after removing the Al deposition (b).
2120
M.-X. Zhang et al. / Surface & Coatings Technology 204 (2010) 2118–2122
Fig. 2. X-ray diffraction spectra taken from the γ-phase layer (a) and β-phase layer (b).
Fig. 3. Variations of the thickness of the Mg–Al intermetallic layers with the post-spray annealing temperature (a) and time (b). Annealing at different temperature was performed for 24 h and the annealing for various times was carried out at 413 °C.
M.-X. Zhang et al. / Surface & Coatings Technology 204 (2010) 2118–2122
2121
Fig. 4. (a) microhardness profile verse the distance from the edge of the γ-layer. The intermetallic layers were produced through post-spray annealing at 413 °C for 24 h. (b) the load–displacement curves obtained from the nanoindentation tests on γ-phase, β-phase, pure Mg and AZ91D alloy as aged at 168 °C for 6 h.
corresponding to the two individual layers and the total layer consist of two sections with 24 h as the separation point. Annealing treatments longer than 24 h do not lead to significant thickening of the intermetallic layers. Although the thickness of the β-phase layer slightly rises, the γ-phase layer is thinned to some extent. This leads to the total thickness of the two layers remaining constant. For annealing times less than 24 h, the layers grow at a nearly constant rate, even though the γ-phase layer grows much faster than the β-phase layer. The higher growth rate of γ-phase is due to the more rapid diffusion of Al atoms than Mg atoms in Al. During annealing, Al atoms diffuse into the Mg substrate from the KM coating, and the Mg atoms diffuse to the coating from the substrate. This leads to the formation of an Al–Mg binary alloy band between the Al coating and the Mg substrate [14]. This binary band contains two discrete regions that correspond to the compositions of β-phase and γ-phase in the phase diagram. As the compositions of both γ and β phases are within a range according to the Mg–Al phase diagram, there is a slight gradient in concentration within both γ and β phase bands. Funamizu and Watanabe's work [14] has concluded that Al atoms diffuse more rapidly than Mg atoms. As one of the results of the faster diffusion of Al, within the binary band, the intermetallic with the higher fraction of Al (γ phase) will grow faster than the intermetallic rich in Mg (β phase). However, after 24 h the growth rate is reduced to nearly zero. This is attributed to the formation of the intermetallic layers that separate the Al coating and the Mg substrate. Further thickening of the layers is arrested due to the formation of porosity, limiting the supply of Al. The variations of the thickness of the intermetallic layers with annealing treatment temperature (from 360 °C up to 430 °C) are shown in Fig. 3(b). Although the γ-phase layer is always thicker than that of the β-phase layer, the total and individual thickness of the two layers increase with the rise of the annealing temperatures from 360 °C to 430 °C for 24 h annealing. This annealing time leads to the thickest intermetallic layers as shown in Fig. 3(a). Since the Mg–Al eutectic temperature is 437 °C, to prevent the substrate and the KM coating from melting, the annealing temperature must be kept below this temperature. The result shown in Fig. 3(b) can also be understood in terms of the inter-diffusion process. As Al atoms diffuse much faster than Mg [14], the γ-phase layer is always thicker than that of β-phase layer. But, higher temperatures also result in faster diffusion of both Al and Mg, and therefore the thickness of the individual γ and β layers and the total layer increase with temperatures. Another predicted result of the faster diffusion of Al is that the pores caused by the Kirkendall effect always form at Al side in between the intermetallic layers and the Al coating, as confirmed by Fig. 1(a). In addition, the results in Fig. 3 indicate that the optimised annealing process to obtain the thickest intermetallic layers is a heat
treatment from 400–413 °C for 24 h. This temperature range is suitable for solution treatment of most AZ series Mg alloys. Thus, the two processes can be combined to make application of this treatment more practical. Fig. 4(a) is the microhardness profile of the intermetallic layers vs distance from the top edge of the γ-phase layer as shown in Fig. 1(b). In terms of the distance, the microhardness variation is divided into three segments corresponding to γ-phase, β-phase and pure Mg substrate. The hardness of the γ-phase is slightly higher than that of the β-phase and the overall hardness of the intermetallic layers is approximately 9 times higher than that of the pure Mg substrate. This indicates the possibility of higher sliding wear resistance of the intermetallic layers. This mechanical behaviour can also be well evaluated by nanoindentation due to the difficulty of extracting tensile samples from the intermetallic compounds. Fig. 4(b) shows typical nanoindentation load–displacement curves for the γ-phase layer, the β-phase layer, the pure Mg and the AZ91D alloy as aged at 168 °C after solution treatment at 413 °C. The reduced elastic modulus (Er) and the nano-hardness (nHV) were determined from the load–displacement curves and the results are listed in Table 1. For comparison, the mechanical properties of the pure Mg were measured by nanoindentation as well. The average reduced modulus of the pure Mg is 47.7 GPa. Given that the elastic modulus (Ei) and the Poisson's ratio (νi) of the diamond indenter are 1141 GPa and 0.07 [13], respectively, and the Poisson's ratio of the pure Mg (ν) is 0.29, the elastic modulus of the pure Mg calculated using Olive– Pharr formula [12] is 45.6 GPa, which agrees very well with the nominal value of 45 GPa [15]. This indicates that the reduced modulus determined by the nanoindentation gave a good measure of the elastic property of the material. We have also computed the elastic moduli of the two intermetallic compounds and the AZ91D alloy using the nominal value of the Poisson's ratio of 0.29. As shown in Table 1, the two intermetallic compounds have very close elastic moduli, which are significantly higher than those of the pure Mg and the AZ91D alloy. Higher elastic modulus implies stronger bonding in the intermetallic compounds. As the difference in electronegativity (EN) between Mg and Al is 0.3 (ENMg = 1.31 and ENAl = 1.61) [16], there is no ionic bonding in these compounds. Thus, the bond of the γ-phase Table 1 Nanomechanical properties of the γ-phase, the β-phase, pure Mg and AZ91D alloy as aged. Data was determined using nanoindentation. Nanomechanics
γ-phase
β-phase
Pure Mg
AZ91D
Er (GPa) E (GPa) nHV (GPa)
66.8 ± 3.5 65.0 ± 3.6 4.40 ± 0.3
68.6 ± 5.0 66.8 ± 5.2 4.35 ± 0.3
47.7 ± 3.0 45.6 ± 3.0 0.90 ± 0.05
51.6 ± 3.5 49.5 ± 3.5 1.24 ± 0.1
2122
M.-X. Zhang et al. / Surface & Coatings Technology 204 (2010) 2118–2122
Fig. 5. AFM images of the nanoindents. (a) γ-phase; (b) β-phase; (c) pure Mg.
and the β-phase should be a mixture of metallic and covalent bonds, which is responsible for the higher modulus of the intermetallic compounds. The curves in Fig. 4(b) also show significant smaller residual depths of penetration of the intermetallic compounds compared to the pure Mg and the AZ91D alloy, indicating their much higher hardness. This is attributed to the portion of covalent bond in the intermetallic compounds compared to the metals that are full metallic bonding. Figs. 5(a), (b) and (c) show typical AFM images of the nanoindents on the γ layer, the β layer and the pure Mg, respectively. The images in Figs. 5(a) and (b) show no cracks surrounding the indents on both γ and β phases though they are expected to be brittle. There are no apparent pile-ups or sink-ins observed around the indent edges of the two intermetallic compounds and the pure Mg, suggesting that plastic flows under the indentation loading were similar for the three types of materials. Recently, Spencer and Zhang [9] investigated the corrosion resistance of the Mg–Al intermetallic compound coatings on Mg–Al– Zn based magnesium alloys compared with other commercially available alloys, such as AZ91E at solution treatment and ageing condition. It was shown that only very minor discolouration was observed on the γ-phase layer while extensive pitting occurred on the AZ91E alloy after 48 h immersion testing in 5 wt.%NaCl solution at room temperature. The electrochemical behaviour of the intermetallic compounds was evaluated using anodic polarisation. Both the γphase and the β-phase have similar electrochemical characteristics to those of pure Al and Al–Si alloys, showing stable passivation over a wide potential range, and low corrosion current density. Neither pure Mg nor AZ91E passivates under similar conditions, and the two alloys have much higher corrosion current density and a less noble open circuit potential than the intermetallic compounds. This implies the intermetallic layers have much higher corrosion resistance than that of the substrate alloys. 4. Summary In summary, the present results show that both the γ-phase and the β-phase have very similar nanomechanical properties, including
reduced elastic modulus and hardness, determined using nanoindentation. The much greater reduced elastic modulus and the 9 times higher microhardness of the intermetallic compounds than the pure Mg substrate lead to a possibility of improvement in wear resistance. Together with the previous observation that both the γ and the β layers significantly improve corrosion resistance, it is considered that producing such intermetallic layers on Mg alloys can provide effective protection of the Mg substrate from both wear and corrosion.
Acknowledgement The authors Zhang and Spencer would like to thank Australian Research Council (ARC) Centre of Excellence for Design in Light Metals for funding support. HH would also like to acknowledge the financial support from ARC under discovery program.
References [1] I. Shigematsu, M. Nakamura, N. Saitou, K. Shimojima, J. Mater. Sci. Let. 19 (2000) 437. [2] H. Huo, F. Wang, Y. Li, J. Xu, Corr. Sci. Prot. Technol. 13 (2001) 484. [3] M.-X. Zhang, P.M. Kelly, J. Mater. Res. 17 (2002) 2477. [4] M.-X. Zhang, Y.-N. Shi, H.Q. Sun, P.M. Kelly, J. Nanosci. Nanotechnol. 8 (2008) 2724. [5] H.Q. Sun, Y.N. Shi, M.-X. Zhang, K. Lu, Surf. Coat. Technol. 202 (2008) 3947. [6] L. Zhu, G. Song, Surf. Coat. Technol. 200 (2006) 2834. [7] N. Pebere, C. Riera, F. Dabosi, Electrochim. Acta 35 (1990) 555. [8] G.L. Song, A. Atrens, Adv. Eng. Mater. 1 (1999) 11. [9] K. Spencer, M.-X. Zhang, Scripta Mater. 61 (2009) 44. [10] I.N. Sneddon, Int. J. Eng. Sci. 3 (1965) 47. [11] J. Karthikegam, adv, Mater. Processes 164 (2006) 12. [12] W.C. Oliver, G.M. Pharr, J. Mater. Res. 19 (2004) 3. [13] H. Huang, K. Winchester, Y. Liu, X.Z. Hu, C.A. Musca, J.M. Dell, L. Faraone, J. Micromech. Microeng. 15 (2005) 608. [14] Y. Funamizu, K. Watanabe, Trans. JIM 13 (1972) 278. [15] M.W. Toaz, E.J. Ripling, Trans. AIME 206 (1956) 936. [16] J.P. Schaffer, A. Saxena, S.D. Antolovich, T.H. Sanders Jr., S.B. Warner, The Science and Design of Engineering Materials, The McGraw-Hill Companies, Inc., Taipei, 1999, p. 770.
Contents lists available at ScienceDirect
Surface & Coatings Technology j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s u r f c o a t
Nanomechanics of Mg–Al intermetallic compounds M.-X. Zhang a,b,⁎, H. Huang a, K. Spencer b, Y.-N. Shi c a b c
School of Mechanical and Mining Engineering, The University of Queensland, Brisbane Qld 4072, Australia ARC Centre of Excellence for Design in Light Metals, Australia Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China
a r t i c l e
i n f o
Article history: Received 6 July 2009 Accepted in revised form 22 November 2009 Available online 29 November 2009 Keywords: Magnesium alloy Intermetallic compound Nanoindentation Kinetic metallization Cold spray
a b s t r a c t Cold spraying of pure Al powder on a pure Mg substrate together with subsequent post-spray annealing treatment produced Mg17Al12 (β-phase) and Mg2Al3 (γ-phase) intermetallic layers on the surface of the substrate. These layers showed significantly better nanomechanical properties, including the reduced elastic modulus and nanohardness, which were determined using nanoindentation, than commercial purity Mg and AZ91 alloys. Combined with their improved corrosion resistance, it is believed that both the γ-phase and the β-phase layers can provide effective protection of Mg alloys from wear and corrosion. The effect of postspray annealing process on the formation of thick, uniform and dense intermetallic layers on pure Mg substrate was also investigated. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Low wear resistance and high corrosive activity of Mg alloys significantly limit their structural applications, even though the low densities of these alloys has attracted increasing interests in automotive and aeronautical industries. A number of previous studies [1–6] have showed that surface alloying treatment with Al, or so called aluminising, offers an effective approach in improving both wear and corrosion resistance due to the increase in volume fraction of the Mg17Al12 (β-phase) in the surface area. As one of the two equilibrium Mg–Al intermetallic compounds, the β-phase not only has higher hardness, which may lead to an increase in the wear resistance, but can also act as an anodic barrier and thus, inhibit the overall corrosion of the alloy [7,8]. Hence, the layer with higher volume fraction of β-phase will have better surface durability and the best result can be achieved with a continuous intermetallic coating. However, obtaining continuous intermetallic layers on an Mg substrate is difficult using techniques like packed powder diffusion. Recently Spencer and Zhang [9] reported a new technique to generate such coatings on Mg–Al–Zn alloys through post-spray annealing of Al cold sprayed coatings. The intermetallic coatings exhibits improved hardness and corrosion resistance compared to the substrate [9]. However, the Zn content in the substrate may affect the diffusion process of Al in Mg, and then affect the formation of intermetallic coatings. This was not discussed and the coatings were not fully characterised in the previous study [9]. In the present work, in order
⁎ Corresponding author. Tel./fax: +61 7 3346 8709; +61 7 33467015. E-mail address: [email protected] (M.-X. Zhang). 0257-8972/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2009.11.031
to eliminate the effect of Zn, a pure Mg substrate has been used so that only binary Mg–Al intermetallic compound coatings can be produced through reaction diffusion with Al. The effect of post-spray annealing on the formation of the intermetallic layers has also been investigated. Furthermore, although similar intermetallic compounds have been reported in an Mg–Al diffusion couple [10], their mechanical and chemical behaviour was not evaluated. In addition, nanoindentation test was used to characterise the nanomechanical properties, including the reduced modulus and nanohardness, of the β-phase and γ-phase (Mg2Al3) formed on a pure Mg substrate. 2. Experimental Commercial purity Mg was used as the substrate material. Pure Al powder with an average particle size of 15 µm was deposited on the 20 × 20 × 8 mm substrates using an Inovati Kinetic Metallization (KM) system at 125 °C with helium as carrying gas at a pressure of 620 kPa in order to increase the impact velocity. KM is similar in principle to conventional cold spray with the distinction that it uses a convergent nozzle to accelerate the gas to sonic velocity (supersonic velocity is used in conventional cold spray) [11]. The mass flow rate of the powder mixtures injected into the gas stream was 15 g min− 1. The nozzle standoff distance from the substrate was 12 mm and the traverse speed was 50 mm min− 1. The post-spray annealing treatments were carried out within a temperature range from 360 °C to 430 °C for 24 h to investigate the effect of annealing temperatures. In order to understand the influence of annealing time on the formation of intermetallic layers, the post-spray treatment was also carried out at 413 °C for various times. This selection of 413 °C is because this temperature is the typical solution treatment temperature for AZ91
M.-X. Zhang et al. / Surface & Coatings Technology 204 (2010) 2118–2122
alloy and maximum Al concentration can be obtained in the solid solution without melting due to the formation of eutectic structure. After annealing, the portion of the Al coating that remained unreacted was readily removed, which was followed by further polishing of the intermetallic layers with the γ-phase layer on the top and the β-phase layer adjacent to the substrate. To expose the β-phase layer, the γ-phase layer has to be removed as well. Nanoindentation tests were performed on the polished β and γ layer using a Hysitron Triboindenter®. A threesided Berkovich indenter with tip radius of 50 nm was used. The indentation load was fixed at 1000 µN and the loading/unloading rate remained constant at 100 μN per second. There was a stabilized time of 5 s between the loading and unloading stages. An atomic force microscope (AFM) was used to examine the surface topographies prior to and after each indentation. The reduced elastic modulus Er and hardness nHV were calculated from the load–displacement curves [12]. Er was determined upon the slope (S) of the unloading curve using the following equation, which is in fact a measure of the elastic property of the surface [13]. Er =
pffiffiffi πS pffiffiffi 2 A
where A is the projected area after the indentation. For comparison, the nanomechanical properties of the pure Mg substrate and the widely used AZ91D alloy after solution treatment at 413 °C for 6 h
2119
followed by ageing at 168 °C for 6 h were also measured. The thickness of the intermetallic layers was measured from the cross sections of samples using optical microscopy. In order to identify the phases formed at the coating/substrate interface, X-ray diffraction was conducted in a Bruker D8 Advance diffractometer with a Cu-Kα target. The operation voltage was 40 V and the scanning speed was 0.5°/min with step size of 0.05°. 3. Results and discussion Micrographs of typical intermetallic layers before and after removing the unreacted portion of the Al KM coating are shown in Fig. 1(a) and (b). According to the Mg–Al binary phase diagram the top layer is the γ-phase (Mg2Al3) and the second layer is the β-phase (Mg17Al12). There is a thin layer of supersaturated solid solution Al in Mg below the β layer, but it is not visible in the optical micrographs. Figs. 2(a) and (b) show the X-ray diffraction spectra taken from γ-phase layer and β-phase layer, respectively. Indexing the spectra confirmed the phases. To keep the figure clear, it only shows indices of the major peaks in the figures. In Fig. 2 (a) a few β-phase peaks can also be observed due to the penetration of the X-ray through the thin areas of the γ layer. Similarly, Mg peaks also appear in Fig. 2(b) due to the same reason. The variations of thickness of the intermetallic layers with the post-spray annealing time (from 5 to 50 h) are shown in Fig. 3(a), which indicates the growth rate of the layers. All the three curves
Fig. 1. Typical optical micrographs showing both the γ-phase layer and the β-phase layer before removing the un-reacted Al deposition (a) and after removing the Al deposition (b).
2120
M.-X. Zhang et al. / Surface & Coatings Technology 204 (2010) 2118–2122
Fig. 2. X-ray diffraction spectra taken from the γ-phase layer (a) and β-phase layer (b).
Fig. 3. Variations of the thickness of the Mg–Al intermetallic layers with the post-spray annealing temperature (a) and time (b). Annealing at different temperature was performed for 24 h and the annealing for various times was carried out at 413 °C.
M.-X. Zhang et al. / Surface & Coatings Technology 204 (2010) 2118–2122
2121
Fig. 4. (a) microhardness profile verse the distance from the edge of the γ-layer. The intermetallic layers were produced through post-spray annealing at 413 °C for 24 h. (b) the load–displacement curves obtained from the nanoindentation tests on γ-phase, β-phase, pure Mg and AZ91D alloy as aged at 168 °C for 6 h.
corresponding to the two individual layers and the total layer consist of two sections with 24 h as the separation point. Annealing treatments longer than 24 h do not lead to significant thickening of the intermetallic layers. Although the thickness of the β-phase layer slightly rises, the γ-phase layer is thinned to some extent. This leads to the total thickness of the two layers remaining constant. For annealing times less than 24 h, the layers grow at a nearly constant rate, even though the γ-phase layer grows much faster than the β-phase layer. The higher growth rate of γ-phase is due to the more rapid diffusion of Al atoms than Mg atoms in Al. During annealing, Al atoms diffuse into the Mg substrate from the KM coating, and the Mg atoms diffuse to the coating from the substrate. This leads to the formation of an Al–Mg binary alloy band between the Al coating and the Mg substrate [14]. This binary band contains two discrete regions that correspond to the compositions of β-phase and γ-phase in the phase diagram. As the compositions of both γ and β phases are within a range according to the Mg–Al phase diagram, there is a slight gradient in concentration within both γ and β phase bands. Funamizu and Watanabe's work [14] has concluded that Al atoms diffuse more rapidly than Mg atoms. As one of the results of the faster diffusion of Al, within the binary band, the intermetallic with the higher fraction of Al (γ phase) will grow faster than the intermetallic rich in Mg (β phase). However, after 24 h the growth rate is reduced to nearly zero. This is attributed to the formation of the intermetallic layers that separate the Al coating and the Mg substrate. Further thickening of the layers is arrested due to the formation of porosity, limiting the supply of Al. The variations of the thickness of the intermetallic layers with annealing treatment temperature (from 360 °C up to 430 °C) are shown in Fig. 3(b). Although the γ-phase layer is always thicker than that of the β-phase layer, the total and individual thickness of the two layers increase with the rise of the annealing temperatures from 360 °C to 430 °C for 24 h annealing. This annealing time leads to the thickest intermetallic layers as shown in Fig. 3(a). Since the Mg–Al eutectic temperature is 437 °C, to prevent the substrate and the KM coating from melting, the annealing temperature must be kept below this temperature. The result shown in Fig. 3(b) can also be understood in terms of the inter-diffusion process. As Al atoms diffuse much faster than Mg [14], the γ-phase layer is always thicker than that of β-phase layer. But, higher temperatures also result in faster diffusion of both Al and Mg, and therefore the thickness of the individual γ and β layers and the total layer increase with temperatures. Another predicted result of the faster diffusion of Al is that the pores caused by the Kirkendall effect always form at Al side in between the intermetallic layers and the Al coating, as confirmed by Fig. 1(a). In addition, the results in Fig. 3 indicate that the optimised annealing process to obtain the thickest intermetallic layers is a heat
treatment from 400–413 °C for 24 h. This temperature range is suitable for solution treatment of most AZ series Mg alloys. Thus, the two processes can be combined to make application of this treatment more practical. Fig. 4(a) is the microhardness profile of the intermetallic layers vs distance from the top edge of the γ-phase layer as shown in Fig. 1(b). In terms of the distance, the microhardness variation is divided into three segments corresponding to γ-phase, β-phase and pure Mg substrate. The hardness of the γ-phase is slightly higher than that of the β-phase and the overall hardness of the intermetallic layers is approximately 9 times higher than that of the pure Mg substrate. This indicates the possibility of higher sliding wear resistance of the intermetallic layers. This mechanical behaviour can also be well evaluated by nanoindentation due to the difficulty of extracting tensile samples from the intermetallic compounds. Fig. 4(b) shows typical nanoindentation load–displacement curves for the γ-phase layer, the β-phase layer, the pure Mg and the AZ91D alloy as aged at 168 °C after solution treatment at 413 °C. The reduced elastic modulus (Er) and the nano-hardness (nHV) were determined from the load–displacement curves and the results are listed in Table 1. For comparison, the mechanical properties of the pure Mg were measured by nanoindentation as well. The average reduced modulus of the pure Mg is 47.7 GPa. Given that the elastic modulus (Ei) and the Poisson's ratio (νi) of the diamond indenter are 1141 GPa and 0.07 [13], respectively, and the Poisson's ratio of the pure Mg (ν) is 0.29, the elastic modulus of the pure Mg calculated using Olive– Pharr formula [12] is 45.6 GPa, which agrees very well with the nominal value of 45 GPa [15]. This indicates that the reduced modulus determined by the nanoindentation gave a good measure of the elastic property of the material. We have also computed the elastic moduli of the two intermetallic compounds and the AZ91D alloy using the nominal value of the Poisson's ratio of 0.29. As shown in Table 1, the two intermetallic compounds have very close elastic moduli, which are significantly higher than those of the pure Mg and the AZ91D alloy. Higher elastic modulus implies stronger bonding in the intermetallic compounds. As the difference in electronegativity (EN) between Mg and Al is 0.3 (ENMg = 1.31 and ENAl = 1.61) [16], there is no ionic bonding in these compounds. Thus, the bond of the γ-phase Table 1 Nanomechanical properties of the γ-phase, the β-phase, pure Mg and AZ91D alloy as aged. Data was determined using nanoindentation. Nanomechanics
γ-phase
β-phase
Pure Mg
AZ91D
Er (GPa) E (GPa) nHV (GPa)
66.8 ± 3.5 65.0 ± 3.6 4.40 ± 0.3
68.6 ± 5.0 66.8 ± 5.2 4.35 ± 0.3
47.7 ± 3.0 45.6 ± 3.0 0.90 ± 0.05
51.6 ± 3.5 49.5 ± 3.5 1.24 ± 0.1
2122
M.-X. Zhang et al. / Surface & Coatings Technology 204 (2010) 2118–2122
Fig. 5. AFM images of the nanoindents. (a) γ-phase; (b) β-phase; (c) pure Mg.
and the β-phase should be a mixture of metallic and covalent bonds, which is responsible for the higher modulus of the intermetallic compounds. The curves in Fig. 4(b) also show significant smaller residual depths of penetration of the intermetallic compounds compared to the pure Mg and the AZ91D alloy, indicating their much higher hardness. This is attributed to the portion of covalent bond in the intermetallic compounds compared to the metals that are full metallic bonding. Figs. 5(a), (b) and (c) show typical AFM images of the nanoindents on the γ layer, the β layer and the pure Mg, respectively. The images in Figs. 5(a) and (b) show no cracks surrounding the indents on both γ and β phases though they are expected to be brittle. There are no apparent pile-ups or sink-ins observed around the indent edges of the two intermetallic compounds and the pure Mg, suggesting that plastic flows under the indentation loading were similar for the three types of materials. Recently, Spencer and Zhang [9] investigated the corrosion resistance of the Mg–Al intermetallic compound coatings on Mg–Al– Zn based magnesium alloys compared with other commercially available alloys, such as AZ91E at solution treatment and ageing condition. It was shown that only very minor discolouration was observed on the γ-phase layer while extensive pitting occurred on the AZ91E alloy after 48 h immersion testing in 5 wt.%NaCl solution at room temperature. The electrochemical behaviour of the intermetallic compounds was evaluated using anodic polarisation. Both the γphase and the β-phase have similar electrochemical characteristics to those of pure Al and Al–Si alloys, showing stable passivation over a wide potential range, and low corrosion current density. Neither pure Mg nor AZ91E passivates under similar conditions, and the two alloys have much higher corrosion current density and a less noble open circuit potential than the intermetallic compounds. This implies the intermetallic layers have much higher corrosion resistance than that of the substrate alloys. 4. Summary In summary, the present results show that both the γ-phase and the β-phase have very similar nanomechanical properties, including
reduced elastic modulus and hardness, determined using nanoindentation. The much greater reduced elastic modulus and the 9 times higher microhardness of the intermetallic compounds than the pure Mg substrate lead to a possibility of improvement in wear resistance. Together with the previous observation that both the γ and the β layers significantly improve corrosion resistance, it is considered that producing such intermetallic layers on Mg alloys can provide effective protection of the Mg substrate from both wear and corrosion.
Acknowledgement The authors Zhang and Spencer would like to thank Australian Research Council (ARC) Centre of Excellence for Design in Light Metals for funding support. HH would also like to acknowledge the financial support from ARC under discovery program.
References [1] I. Shigematsu, M. Nakamura, N. Saitou, K. Shimojima, J. Mater. Sci. Let. 19 (2000) 437. [2] H. Huo, F. Wang, Y. Li, J. Xu, Corr. Sci. Prot. Technol. 13 (2001) 484. [3] M.-X. Zhang, P.M. Kelly, J. Mater. Res. 17 (2002) 2477. [4] M.-X. Zhang, Y.-N. Shi, H.Q. Sun, P.M. Kelly, J. Nanosci. Nanotechnol. 8 (2008) 2724. [5] H.Q. Sun, Y.N. Shi, M.-X. Zhang, K. Lu, Surf. Coat. Technol. 202 (2008) 3947. [6] L. Zhu, G. Song, Surf. Coat. Technol. 200 (2006) 2834. [7] N. Pebere, C. Riera, F. Dabosi, Electrochim. Acta 35 (1990) 555. [8] G.L. Song, A. Atrens, Adv. Eng. Mater. 1 (1999) 11. [9] K. Spencer, M.-X. Zhang, Scripta Mater. 61 (2009) 44. [10] I.N. Sneddon, Int. J. Eng. Sci. 3 (1965) 47. [11] J. Karthikegam, adv, Mater. Processes 164 (2006) 12. [12] W.C. Oliver, G.M. Pharr, J. Mater. Res. 19 (2004) 3. [13] H. Huang, K. Winchester, Y. Liu, X.Z. Hu, C.A. Musca, J.M. Dell, L. Faraone, J. Micromech. Microeng. 15 (2005) 608. [14] Y. Funamizu, K. Watanabe, Trans. JIM 13 (1972) 278. [15] M.W. Toaz, E.J. Ripling, Trans. AIME 206 (1956) 936. [16] J.P. Schaffer, A. Saxena, S.D. Antolovich, T.H. Sanders Jr., S.B. Warner, The Science and Design of Engineering Materials, The McGraw-Hill Companies, Inc., Taipei, 1999, p. 770.