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Effect of heat treatment on the intermetallic layer of cold sprayed aluminum coatings on magnesium alloy
Surface & Coatings Technology 205 (2011) 4665–4671
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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
Effect of heat treatment on the intermetallic layer of cold sprayed aluminum coatings on magnesium alloy Hengyong Bu a,⁎, Mohammed Yandouzi b, Chen Lu a, Bertrand Jodoin b a b
National Engineering Research Center of Light Alloy Net Forming, School of Materials Science and Engineering, Shanghai JiaoTong University, Shanghai 200240, PR China Mechanical Engineering Department, University of Ottawa, Ottawa, K1N 6N5 ON, Canada
a r t i c l e
i n f o
Article history: Received 23 January 2011 Accepted in revised form 3 April 2011 Available online 9 April 2011 Keywords: Aluminum coatings Magnesium alloys Heat treatment Cold spray Intermetallic phases
a b s t r a c t Dense and thick pure aluminum coatings were deposited on AZ91D-T4 magnesium substrates using the cold spray process. Heat treatments of the as-sprayed samples were carried out at 400 °C using different holding times. The feedstock powder, substrate and coating microstructures were examined using optical microscopy (OM), scanning electron microscopy (SEM), energy dispersive spectroscopy (EDS) as well as Vickers microhardness analysis. The results demonstrate that aluminum coatings having dense and uniform microstructure can be deposited successfully using a relatively large feedstock powder. It has been identified that the intermetallics Al3Mg2 (γ phase) and Mg17Al12 (β phase) were formed at the coating/substrate interface during heat treatment. The growth rate of these intermetallics follows the parabolic law and the γ phase has a higher growth rate than the β phase. The thickness of the Mg17Al12 and Al3Mg2 intermetallic layers has reached 83 μm and 149 μm, respectively. This result is almost 45% higher than what has been reported in the literature so far. This is attributed to the fact that T4 instead of as cast Mg alloy was used as substrate. In the T4 state, the Al concentration in the Mg matrix is higher, and thus intermetallic growth is faster as less enrichment is required to reach the critical level for intermetallic formation in the substrate. The AZ91D-T4 magnesium substrate contains single α phase with fine clusters/GP-zones which is considered beneficial for the intermetallic formation as well as the intimate contact between the coating/substrate interface and the deformed particles within the coating. © 2011 Elsevier B.V. All rights reserved.
1. Introduction Magnesium (Mg) alloys are attractive materials for a wide number of applications owing to their suitable mechanical properties, low density, as well as their vast reserve [1]. In addition, their high specific strength, good damping capacity and recyclability make them suitable solutions for the automotive, aerospace and electronic industries [1–3]. However, Mg alloys are susceptible to react with various metals or ions (such as chloride ions). Consequently, they have poor chemical and galvanic corrosion performance [1,4,5]. Furthermore, inevitable impact or scratch of Mg components will lead to significant mass loss due to its low hardness [6]. These poor properties increase the maintenance costs and limit the life cycle of Mg alloy components. One way to prevent these defects is to manufacture a coating which has better performance and isolate the Mg alloy matrix from the environment, without negative effects to the Mg-base alloy [7,8]. Several coating techniques such as anodizing vapor deposition, electroplating and chromate coating, to name a few,
⁎ Corresponding author. Tel.: +86 21 54742618; fax: +86 21 34202794. E-mail address: [email protected] (H. Bu). 0257-8972/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2011.04.018
have been used for magnesium protection. Nevertheless, they exhibit some limitations, such as low deposition efficiency, high power requirement or being environment unfriendly [8]. Cold spray is a coating technology which uses high pressure gas to accelerate solid powder particles beyond a critical velocity [9–13]. Upon impact with the substrate, the particles undergo severe plastic deformation and form a coating. Compared to other coating manufacturing methods such as thermal spray or electroplating, cold spray has demonstrated many advantages with the most important being that the coatings retain the original microstructure of the feedstock powders [10,11,14]. Using the cold spray process, previous studies have reported the possibility of protecting Mg-alloy using aluminum (Al) and Al-based alloy coatings [2,6,7,15–17]. It was found that Al coatings on ZE41A-T5 substrates have similar performances than bulk Al and can improve the corrosion resistance compared to the uncoated substrate [2,6]. In order to attain higher bonding strength and promote the overall coating performance, cold spray coatings have been heat treated. This has resulted in element diffusion to form intermetallic compounds within Al/Ni [18], Al/Ti [10,19] and Al/Mg [15,20] couples. Annealed cold spray Al coatings on magnesium substrates have revealed the formation of the intermetallic β and γ layers near the substrate/ coating interface. It has been reported that the later phases have similar corrosion resistance than Al alloys [15,20].
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The advantage of generating intermetallic compounds using cold spray combined with post heat treatment can be summarized as follows: First, it reduces shrinkage during subsequent annealing if the coatings have a high density, which is different from conventional sintering that exhibits large porosities [21]. Second, the formed intermetallic compounds are continuous and uniform onto the original substrate and exhibit higher hardness and better corrosion resistance [15]. Different intermetallic phases with different thicknesses have been reported for Mg/Al interdiffusion process during heat treatment. Shigematsu et al. [4] reported the formation of the Mg17Al12 phase with approximately 750 μm layer thickness after 1 h at 450 °C using powder metallurgy technology. Using pure Al powders with a mean diameter of 15 μm as feedstock in the cold spray process, Spencer et al. [15,20] have identified the existence of two intermetallic phases (Mg17Al12 and Al3Mg2) at the Mg/Al interface with a total thickness around 200 μm after 24 h at 400 °C. Liu et al. [22] have identified the presence of three intermetallic phases (MgAl, Mg3Al2 and Al3Mg2) during the vacuum diffusion bonding of the joint Mg and Al bulk materials annealing at 480 °C for 1 h. Consequently, it can be concluded that the type and thickness of intermetallic phases depend on the heat treatment conditions. Large feedstock powders have attracted much attention from the cold spray community as they have lower critical velocity and decreased the risk of explosion compared to fine powders [14,23–25]. As such, there is a considerable interest to determine and evaluate the intermetallic compounds fabricated by cold spray using relative large powders and subsequent heat treatment. The aim of this work is to use the cold spray process to spray Al onto Mg alloy using large Al feedstock powders. AZ91D-T4 magnesium substrate was selected since it contains a uniform solid solution α-phase with clusters/GP-zones and the aluminum concentration in the magnesium matrix is higher. Therefore intermetallic growth is faster because less enrichment is required to reach the critical level for intermetallic formation in the substrate. The effects of different heat treatments on the type of intermetallic phases and their thicknesses are evaluated and discussed.
2. Experimental procedures The coatings were produced at the University of Ottawa Cold Spray Laboratory. Details of the spray system can be found elsewhere [26,27]. Helium was used as both propellant gas and powder carrier gas, although nitrogen can be used if operating at higher pressure. The spraying parameters used in the current work are listed in Table 1. The feedstock powders were accelerated using a De Laval nozzle which has an expansion ratio of 10. Pure commercially available aluminum powder (Al-101, Centerline (Windsor) Ltd, Canada) was used as feedstock. The powder size distribution was measured by a laser particle size analyzer (Coulter LS, Fullerton, CA). Commercially available AZ91D-T4 magnesium substrates produced by High Pressure Die Casting (HPDC) were used after being submitted to a solution heat treatment at 400 °C for 12 h and followed by quenching in water. Before spraying the Al powders, the natural aging duration to T4 state substrates was no longer than 48 h and substrates with dimensions of 170 × 20 × 6 mm were grit blasted using (20 mesh – 24 grit) silica and cleaned using acetone. The substrates were then assembled on a twoaxis computer controlled moving system.
Table 1 Parameters for cold spray process used for spraying pure aluminum (Al-101) powders. Gas parameters Pressure (MPa)
Temperature (°C)
0.98
300
Spray distance (mm)
Number of pass
10
1
The coated samples were sectioned and prepared for microscopy analysis, following standard metallographic methods. The etching solutions for magnesium substrate and Al coating were 4% nitric acid + 96% ethanol and Kroll's reagent, respectively. The feedstock powder and coating microstructures were characterized using Optical Microscopy (OM) and Scanning Electron Microscopy (Zeiss EVO-MA10, LaB6 Analytical SEM) equipped with an Energy Disperse Spectroscopy (INCA X-act, Oxford, UK). Phase identification was investigated by X-ray diffraction (XRD) using a Philips X-Pert model 1830 X-ray diffractometer equipped with a graphite monochromator using Cu kα (λ = 0.15406 nm) radiation. Detailed scans were performed over a 20–100° 2θ range, 0.02° step width and 2 s per step acquisition time. Images of the cross sections were used to evaluate the coating porosity level using the commercially available Clemex imaging software. The powder and the coating microhardness were evaluated using a load of 25 g (HV0.025). All measurements used a dwell time of 15 s, using Duramin-2 Vickers hardness tester (Struers Inc., Cleveland, OH, USA). The reported hardness values are averages of at least 6 random measurements for each sample. The heat treatments of the as-sprayed coatings were carried out in a vacuum furnace under pressure of 5.0 × 10−4 bars. The samples were heated at constant heat treatment temperature (400 °C) for different period of time up to 20 h with a temperature fluctuation below 1 °C. 3. Results and discussion 3.1. Characteristics of pure Al powder The morphology of the Al-101 feedstock powder is shown in Fig. 1a. SEM observations of the free standing particles reveal the irregular morphology of the particles varying from spherical to elongated structures with different sizes. Fig. 1b presents an OM image of the polished cross section powder. The image shows a dendritic microstructure within the Al particle, caused by the rapid solidification rate experienced by the particle during the powder atomization process. The particle size distribution diagram obtained using laser size analyzer is shown in Fig. 1c. It indicates that the powder particles present a broad size distribution with a diameter ranging from 6 to 174 μm, and the particles larger than 50 μm represent approximately 40% of the particles. 3.2. Properties of the as sprayed coatings and substrate 3.2.1. Substrate The microstructure of the high pressure die casting AZ91D magnesium alloy is shown in Fig. 2a. It contains the primary α phase (Mg) matrix as well as the discontinuous intermetallic β phase (Mg17Al12) [28]. The later precipitates are located mostly at the α phase dendrite grain boundaries. After T4 heat treatment, the precipitates (β phase) were dissolved into the matrix and the microstructure of the Mg alloy displays a single Al supersaturated α phase as seen in Fig. 2b. Note that due to small size of clusters/GP zones that were formed inside the Al supersaturated grains during the natural aging, they could not be seen by using OM or SEM imaging technique. The XRD results of the as cast and T4 magnesium substrates are shown in Fig. 3, which also demonstrate that after solution treatment, the β phase peaks have disappeared. 3.2.2. Coating characteristics An SEM image under secondary electron mode of the as-sprayed coating cross section is shown in Fig. 4. The coating has a thickness of approximately 300 μm and limited porosity (less than 1%). The overall coating is dense with the exception of the top coating surface where limited porosity was observed. This is attributed to the fact that the large particles do not undergo enough plastic deformation upon impact and thus leave some porosity at the particle–particle
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Fig. 2. OM image showing (a) the microstructure of high pressure die casting AZ91D magnesium solid solution alloy α phase matrix and the β phase precipitates. (b) AZ91D magnesium alloy after T4 heat treatment (400 °C × 12 h and quenched in water).
Fig. 1. Al-101 feedstock powder: (a) SEM image shows the morphology of the free standing particles. (b) OM image of resin embedded powder cross section after polishing and weak etching revealing the dendrite microstructure of the Al particle. (c) Size distribution of the Al powder measured by laser size analyzer.
boundaries. In addition, the absence of significant peening effect that sometimes contributes to extra deformation of the particles already deposited leads to the increased porosity at the top of the coating. This phenomenon has already been described in other works [29]. The interface between the Al coating and the Mg substrate has almost no porosity, which means the particles near the interface had enough velocity to remove the surface oxide to expose fresh material, resulting in the intimate contact at the interface. This is considered beneficial for coating bonding strength and for enhanced diffusion of the elements during the subsequent heat treatment. The feedstock powder used in this study presents a large size distribution, ranging from less than 10 μm to over 170 μm. It is known that large powder particle size distributions have a detrimental effect on the coating quality because the large particles are harder to accelerate than the smaller ones while the latter may be affected by the bow shock wave present in front of the substrate [30]. Consequently, it is difficult for all of the particles to reach a similar high velocity, and thus it can be expected that the large feedstock powder particle size distribution used in this work will typically result in a large velocity distribution. As such, the particles which have a velocity lower than the critical velocity will rebound from the substrate surface or adhere to
the substrate with a low bonding strength and minimum deformation level, increasing the possibility of forming porosities. The morphology of the Al particles in the coating after etching is shown in Fig. 5. It can be seen that there were several large particles in the coating which were labeled by black arrow if their size was apparently larger than 50 μm, and the total area of the larger particles found in the coating was estimated to be 42%. Compared to the ratio of the particles larger than 50 μm in the feedstock powder, it is clear that the deposition efficiency is similar for all particle sizes in the current work. Since the produced coatings exhibit low porosity level, it is concluded that the disadvantages of large powder size distribution were alleviated by the choice of optimized spraying parameters.
Fig. 3. XRD results of the as cast and T4 treated AZ91D magnesium substrates.
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Fig. 4. SEM images of the cross section of Al coatings onto AZ91D magnesium.
Compared to the microhardness of Al feedstock powder of 39.6 ± 2.4 HV0.025, the hardness of the as sprayed coating was increased to 60± 3.4 HV0.025. It is concluded that the strain hardening effect due to the intense plastic deformation of the particles upon impact with the substrate played an important role on increasing the aluminum hardness as reported in other studies [23]. 3.3. Heat treatments of the coatings 3.3.1. Intermetallic phase Al–Mg binary phase diagram is shown in Fig. 6 [31]. It can be seen that all the Al–Mg intermetallic compound have a melting temperature lower than 450 °C and the Mg17Al12 phase has a wider composition range than Al3Mg2 phase at elevated temperature. The cold spray coating microstructure after the 4 h heat treatment at 400 °C is shown in Fig. 7a. Four distinctive zones can be observed. The first (the top layer) is the un-reacted as sprayed aluminum coating. The second and the third zones represent the intermetallic layers that formed at the coating/substrate interface. The fourth is the AZ91D magnesium substrate shown at the bottom. According to the Al–Mg binary phase diagram (Fig. 6), it is expected that the intermetallic compound near the substrate is Mg rich phase -Mg17Al12 (β phase) while the intermetallic layer present near the un-reacted Al coating will be the Al rich phase- Al3Mg2 (γ phase). Quantitative phase's composition was investigated using the EDS technique. It is accepted that in Al–Mg alloy systems, the composition ranges between 52–60 at.% Mg consist of a single β phase, while γ and R phases will be formed in the area range between 38.5–40.3 and 41.7–42.3 at.% Mg respectively. Note that the later, R phase can only exist between 320–370 °C, as it decomposes to β and γ phases [31].
Fig. 5. OM image of the etched as-sprayed coating showing the morphology of the Al particles and the arrows indicated the particles apparently larger than 50 μm.
Fig. 6. Al–Mg binary phase diagrams.
Three EDS analysis points were performed on the four different layers. According to the composition results, the intermetallic compounds near the magnesium substrate and the un-reacted aluminum coating were β phase (51.97–55.41 at.% Mg) and γ phase (39.23–39.27 at.% Mg), respectively. Moreover, EDS elemental line scan from the magnesium substrate side to the un-reacted aluminum coating side was conducted and the obtained quantitative composition is shown in Fig. 7b. It is observed that the β phase has a larger composition range compared to the γ phase and that the composition changes abruptly
Fig. 7. (a)SEM image of cold sprayed pure Al onto AZ91D- T4 substrate after heat treatment under vacuum at 400 °C- 4 h.revealing different intermetallic phases and (b) Quantitative EDS line scan as indicated by arrow in (a).
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at the interface of the two layers. It is worth noting that some porosity within the top layer of the γ phase was observed (arrows in Fig. 7a). Compared to Spencer [15] who also carried out heat treatment on the Al/Mg system, the volume of porosities on the coating cross section in the current study is smaller and only a few pores are observed. 3.3.2. Intermetallic thickness The intermetallic thickness variation with heat treatment time is shown in Fig. 8. It can be observed that the γ phase and β phase grow steadily for heat treatment at 400 °C under 16 h. For the 20 h heat treatment, the thickness of the γ phase layer increases at a much slower rate while the β phase does not show noticeable growth. These results are similar to previous studies [20], where a post annealing on Al/Mg at 413 °C was conducted, and found that the two intermetallic layers did not grow noticeably when the heat treatment time exceeded 24 h. Comparable conclusions can also be drawn from the linear regression growth curve shown in Fig. 9, where it can be observed that the slope of the γ phase is larger than β phase, and the former is almost 2.5 times higher than the latter. The curves reveal a linear trend that goes through the origin point, which demonstrates that the diffusion process obeys a parabolic rate law for heat treatment time no more than 20 h. Similar results have also been reported by Y. Funamizu and K. Watanabe [32] who suggested that the growth rate of the γ phase and β phase represent a linear relation with the square root of diffusion time when the heat treatment was performed at 425 °C and the slope of γ phase is almost 2.2 times of that for β phase. It implies that the intermetallics grow faster at the initial stage and then the overall concentration gradient decreased swiftly associated with the intermetallic thickness. It can be found in Fig. 8 that the migrating distance of Al atoms into Mg substrate is a little more than 60 μm, compared to less than 60 μm for the Mg atoms diffusing in the reverse direction when the heat treatment time is 4 h. This suggests that the aluminum diffuses more rapidly than magnesium in the system. Moreover, it is noted that when increasing the heat treatment time from 4 h to 20 h, the position of the β/γ interface is almost stable, which implies that these two intermetallics grow individually other than transform from one to the other. The intermetallic growth rate is mainly due to the atom diffusion rate in the two intermetallics. Thus, one of the reasons for higher growth rate of γ phase during the heat treatment is probably due to the higher diffusion rate of aluminum atoms in the two intermetallic phases [20,32]. Note that when two or more intermediate phases are formed in a diffusion zone in a binary system such Aluminum–Magnesium, many factors control the occurrence and growth of the various phases. The growth of a phase, in multiphase
Fig. 8. Thickness measurement of the intermetallic layers vs holding time of the heat treatment at 400 °C under vacuum. The horizontal zero axis corresponds to the original Mg/Al interface before the heat treatment.
Fig. 9. Growth curve of the intermetallic β- and γ-phases at 400 °C vs. square root diffusion time (h1/2).
diffusional growth is dependent on its own interdiffusion coefficient as well as those of its neighboring phases. Funamizu et al. [32] have attributed the difference in the thickness growth of the β and γ phase during heat treatment to the difference between the activation energy of layer growth and that of the interdiffusion of the two phases. The later was attributed to the difference in the composition range and temperature dependence of the two intermetallic phases as shown in the equilibrium phase diagram of the Al–Mg system (Fig. 6). High interdiffusion coefficient and low temperature dependence of the γ-phase composition contribute significantly in the growth of this intermetallic as compared to β-phase. More details about the different activation energies of the different phases within the Al–Mg system can be found elsewhere [32]. According to Fick's Second Law, the atom diffusion coefficient is mainly due to the concentration gradient and the temperature, and the diffusion distance is proportional to the square root of the heat treatment time. The thickness of the total intermetallics thicknessdiffusion distance-in this study obeys a parabolic rate law [32] (Fig. 9). However, Spencer et al. [15] suggested that the parabolic law is not applicable in the Al/Mg couple diffusion process because of the Kirkendall effect, which will lead to the condensation of vacancies near the edge of γ phase and then decrease the diffusion rate. It can be seen in Fig. 7a that there are few small pores near the edge of γ phase, as opposed to a large number of holes observed in Ref. [15] at the same position after diffusion, despite similar porosity level content in the as-sprayed coatings. The reason for the smaller porosities at the γ phase/un-reacted Al coating may be due to the large Al particles which were used as feedstock powder resulting in a coating with reduced number of Al/Al splats that would enhance the Mg diffusion in to the γ phase. As the thickness of the intermetallics obeys the parabolic law, increasing the heat treatment time or temperature to obtain thicker intermetallics is not a solution. First, the higher temperature will cause the melting of the intermetallic phases which were formed by thermal diffusion [4]; moreover, it will accelerate the grain growth of magnesium substrate which is detrimental to the mechanical performance based on the Hall–Petch relationship [33]. It has been recommended that the Al/Mg couple diffusion-induced phase should be performed between 400 °C to 436 °C [15], taking into account that the eutectic β phase will start local dissolution at 426 °C [5]. As such, the diffusion temperature was limited in a narrow extent and in the current study a temperature of 400 °C was selected, because it is the same temperature for AZ91D magnesium alloy solution treating. Consequently, during the diffusion heat treatment, the β phase will not precipitate and the substrate will remain a single α phase.
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When the as-sprayed coating samples are heat treated, the atoms can get enough energy and diffuse resulting in the formation of Al–Mg binary alloy band [20] between the Al coating and the Mg substrate. Generally, this binary band should contain αAl, αMg, γ phase and β phase with a concentration gradient according to the diffusion theory [22], but in fact it was found that it mainly consists of two different intermetallic compounds. Similar to other work [20], the layers of Mg matrix with saturated Al atoms and Al matrix with saturated Mg atoms were not detected in the OM and SEM images due to the limitation of the technique used in this work. Because Al/Mg diffusion is a reaction diffusion system, Al atoms and Mg atoms can react and form only one intermetallic compound when the atomic ratio arrive at a critical value at some particular point, according to the second law of thermodynamics. That is why the composition between the intermetallics and substrate/β phase and unreacted Al/γ phase will be changed abruptly rather than gradually. The typical optical microstructure of the as sprayed coating after heat treatment of 20 h is shown in Fig. 10. With a heat treatment time of 20 h at 400 °C, the β phase was 83 μm thick while the γ phase was 149 μm thick for a total thickness of 232 μm. Compared to other studies [20] which attained a total thickness of a little more than 160 μm below 413 °C for 24 h, the total intermetallic thickness is increased by 45%. The main reason for the thicker intermetallic layers found in the current work is believed to be that T4 instead of as-cast magnesium bulk [15] was used as substrates. To confirm this assumption, a comparison of the intermetallic thickness on T4 and as-cast AZ91D after 400 °C heat treatment for 4 h is shown in Fig. 11. The β phase and γ phase thicknesses were respectively 36.7 μm and 93.3 μm for T4 substrate and 30.8 μm and 83.2 μm for as-cast substrate. Generally, the aluminum content in the AZ91 magnesium bulk is near 9 wt.%. T4 substrates have a single solid solution α-phase with very fine clusters/ GP zones and Al atoms concentration is higher and more homogeneously distributed within Mg matrix. Hence intermetallic growth is faster as less enrichment is required to reach the critical level for intermetallic formation in the substrate. However, as-cast AZ91 substrates contain β phase precipitates and α phase Mg matrix. The β phase precipitates have aluminum content of 44.3%, leading to a magnesium matrix containing less than 9% aluminum because of beta precipitation. Thus, the as-cast state has more Al tied up in the beta phase and consequently less available in the Mg matrix to form intermetallics. Moreover, it is expected that the dense coating and the
Fig. 10. OM image of the cold sprayed Al coating onto AZ91D-T4 substrate after heat treatment at 400 °C for 20 h.
Fig. 11. Comparison of the intermetallic thicknesses formed on the coating which sprayed onto as cast and T4 substrates heat treated at 400 °C during 4 h.
intimate contact of the Al/Mg interface and the Al splats, which were dramatically influenced by the feedstock powders and the spraying conditions, were believed to have a positive effect on Al diffuse process. 3.3.3. Hardness The purpose of cold spray coatings and the intermetallics formed by post heat treatment is providing a protection of magnesium substrates from the environment. Any notable degradation of the substrate due to the heat treatment should be avoided. As such, microhardness testing was used to evaluate the mechanical properties of the Mg substrates. The microhardness of the Mg substrate and of the two intermetallic layers as a function of heat treatment time is shown in Fig. 12. The hardness value of as cast and T4 Mg substrates were 80.7 ± 1.8 HV0.025 and 73.7 ± 4.0 HV0.025, respectively. This hardness decrease was mainly due to the dissolution of the hard β phase, which acts as a reinforcement phase in the AZ magnesium alloy series, in the Mg matrix. After the diffusion heat treatment, the β phase was precipitated from the Mg matrix which saturated with Al atoms during the cooling process, and the hardness of Mg substrate increased to around 91 HV0.025. Despite the diffusion treatment being carried out for different holding time, the Mg substrate hardness was almost the same, because the substrate maintains a single α phase when the furnace temperature is kept at 400 °C and the β phase could not precipitate in this process. All the samples in this study were
Fig. 12. Microhardness results of the Mg substrates and the intermetallic layers. measured after 400 °C heat treatment for different period of time.
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cooled from 400 °C to room temperature according to the same cooling condition. Therefore, it is expected that the Mg substrates should have the same β phase content which is considered as a reinforcement phase in AZ series Mg alloys and could only precipitate in the cooling process. Moreover, the hardness of the γ phase was slightly larger than the β phase, 279.6 ± 13.7 HV0.025 and 260.5 ± 10.7 HV0.025 respectively. Similar values have been reported by Spencer et al. [15]. Note that the measured hardness is 2 times higher than the as cast AZ91D magnesium substrate. It is widely accepted that the higher hardness of the coating is considered helpful to attain a higher wear resistance [34]. 4. Conclusion The cold spray parameters and nozzle used in this study are appropriate for spraying large aluminum powders with wide size distribution and a dense coating is attained which has a porosity content less than 1%. The growth rate of the Al–Mg intermetallic phases during annealing at 400 °C obeys a parabolic law and the Alenrich phase (γ phase) grows much faster- around 2.5times- than the Mg-enrich phase (β phase). Increased diffusion time will lead to thicker intermetallics between un-reacted Al coating and Mg substrate. The total thickness of the intermetallics after 20 h diffusion is more than 230 μm, which is 45% thicker than thicknesses that have been reported in the literature. Considering the relative lower temperature and holding time used in this study, it is believed that the coating can provide better protection and less detrimental effect for the Mg substrate. Compared to the AZ91D magnesium alloy bulk, no matter as cast or T4 status, the intermetallic compounds which formed after cold spray and post heat treatment have a much higher hardness, increasing the possibility of promoting the wear resistance of magnesium alloy surface. As a result, the pure Al coatings deposited on Mg substrates using cold spray and post heat treatment can be treated as an alternative method for magnesium protection. Acknowledgments The authors would like to thank the Meridian Technologies (Ontario, Canada) for providing the AZ91D Magnesium materials and Shanghai JiaoTong University for the financial support given to this work.
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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
Effect of heat treatment on the intermetallic layer of cold sprayed aluminum coatings on magnesium alloy Hengyong Bu a,⁎, Mohammed Yandouzi b, Chen Lu a, Bertrand Jodoin b a b
National Engineering Research Center of Light Alloy Net Forming, School of Materials Science and Engineering, Shanghai JiaoTong University, Shanghai 200240, PR China Mechanical Engineering Department, University of Ottawa, Ottawa, K1N 6N5 ON, Canada
a r t i c l e
i n f o
Article history: Received 23 January 2011 Accepted in revised form 3 April 2011 Available online 9 April 2011 Keywords: Aluminum coatings Magnesium alloys Heat treatment Cold spray Intermetallic phases
a b s t r a c t Dense and thick pure aluminum coatings were deposited on AZ91D-T4 magnesium substrates using the cold spray process. Heat treatments of the as-sprayed samples were carried out at 400 °C using different holding times. The feedstock powder, substrate and coating microstructures were examined using optical microscopy (OM), scanning electron microscopy (SEM), energy dispersive spectroscopy (EDS) as well as Vickers microhardness analysis. The results demonstrate that aluminum coatings having dense and uniform microstructure can be deposited successfully using a relatively large feedstock powder. It has been identified that the intermetallics Al3Mg2 (γ phase) and Mg17Al12 (β phase) were formed at the coating/substrate interface during heat treatment. The growth rate of these intermetallics follows the parabolic law and the γ phase has a higher growth rate than the β phase. The thickness of the Mg17Al12 and Al3Mg2 intermetallic layers has reached 83 μm and 149 μm, respectively. This result is almost 45% higher than what has been reported in the literature so far. This is attributed to the fact that T4 instead of as cast Mg alloy was used as substrate. In the T4 state, the Al concentration in the Mg matrix is higher, and thus intermetallic growth is faster as less enrichment is required to reach the critical level for intermetallic formation in the substrate. The AZ91D-T4 magnesium substrate contains single α phase with fine clusters/GP-zones which is considered beneficial for the intermetallic formation as well as the intimate contact between the coating/substrate interface and the deformed particles within the coating. © 2011 Elsevier B.V. All rights reserved.
1. Introduction Magnesium (Mg) alloys are attractive materials for a wide number of applications owing to their suitable mechanical properties, low density, as well as their vast reserve [1]. In addition, their high specific strength, good damping capacity and recyclability make them suitable solutions for the automotive, aerospace and electronic industries [1–3]. However, Mg alloys are susceptible to react with various metals or ions (such as chloride ions). Consequently, they have poor chemical and galvanic corrosion performance [1,4,5]. Furthermore, inevitable impact or scratch of Mg components will lead to significant mass loss due to its low hardness [6]. These poor properties increase the maintenance costs and limit the life cycle of Mg alloy components. One way to prevent these defects is to manufacture a coating which has better performance and isolate the Mg alloy matrix from the environment, without negative effects to the Mg-base alloy [7,8]. Several coating techniques such as anodizing vapor deposition, electroplating and chromate coating, to name a few,
⁎ Corresponding author. Tel.: +86 21 54742618; fax: +86 21 34202794. E-mail address: [email protected] (H. Bu). 0257-8972/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2011.04.018
have been used for magnesium protection. Nevertheless, they exhibit some limitations, such as low deposition efficiency, high power requirement or being environment unfriendly [8]. Cold spray is a coating technology which uses high pressure gas to accelerate solid powder particles beyond a critical velocity [9–13]. Upon impact with the substrate, the particles undergo severe plastic deformation and form a coating. Compared to other coating manufacturing methods such as thermal spray or electroplating, cold spray has demonstrated many advantages with the most important being that the coatings retain the original microstructure of the feedstock powders [10,11,14]. Using the cold spray process, previous studies have reported the possibility of protecting Mg-alloy using aluminum (Al) and Al-based alloy coatings [2,6,7,15–17]. It was found that Al coatings on ZE41A-T5 substrates have similar performances than bulk Al and can improve the corrosion resistance compared to the uncoated substrate [2,6]. In order to attain higher bonding strength and promote the overall coating performance, cold spray coatings have been heat treated. This has resulted in element diffusion to form intermetallic compounds within Al/Ni [18], Al/Ti [10,19] and Al/Mg [15,20] couples. Annealed cold spray Al coatings on magnesium substrates have revealed the formation of the intermetallic β and γ layers near the substrate/ coating interface. It has been reported that the later phases have similar corrosion resistance than Al alloys [15,20].
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The advantage of generating intermetallic compounds using cold spray combined with post heat treatment can be summarized as follows: First, it reduces shrinkage during subsequent annealing if the coatings have a high density, which is different from conventional sintering that exhibits large porosities [21]. Second, the formed intermetallic compounds are continuous and uniform onto the original substrate and exhibit higher hardness and better corrosion resistance [15]. Different intermetallic phases with different thicknesses have been reported for Mg/Al interdiffusion process during heat treatment. Shigematsu et al. [4] reported the formation of the Mg17Al12 phase with approximately 750 μm layer thickness after 1 h at 450 °C using powder metallurgy technology. Using pure Al powders with a mean diameter of 15 μm as feedstock in the cold spray process, Spencer et al. [15,20] have identified the existence of two intermetallic phases (Mg17Al12 and Al3Mg2) at the Mg/Al interface with a total thickness around 200 μm after 24 h at 400 °C. Liu et al. [22] have identified the presence of three intermetallic phases (MgAl, Mg3Al2 and Al3Mg2) during the vacuum diffusion bonding of the joint Mg and Al bulk materials annealing at 480 °C for 1 h. Consequently, it can be concluded that the type and thickness of intermetallic phases depend on the heat treatment conditions. Large feedstock powders have attracted much attention from the cold spray community as they have lower critical velocity and decreased the risk of explosion compared to fine powders [14,23–25]. As such, there is a considerable interest to determine and evaluate the intermetallic compounds fabricated by cold spray using relative large powders and subsequent heat treatment. The aim of this work is to use the cold spray process to spray Al onto Mg alloy using large Al feedstock powders. AZ91D-T4 magnesium substrate was selected since it contains a uniform solid solution α-phase with clusters/GP-zones and the aluminum concentration in the magnesium matrix is higher. Therefore intermetallic growth is faster because less enrichment is required to reach the critical level for intermetallic formation in the substrate. The effects of different heat treatments on the type of intermetallic phases and their thicknesses are evaluated and discussed.
2. Experimental procedures The coatings were produced at the University of Ottawa Cold Spray Laboratory. Details of the spray system can be found elsewhere [26,27]. Helium was used as both propellant gas and powder carrier gas, although nitrogen can be used if operating at higher pressure. The spraying parameters used in the current work are listed in Table 1. The feedstock powders were accelerated using a De Laval nozzle which has an expansion ratio of 10. Pure commercially available aluminum powder (Al-101, Centerline (Windsor) Ltd, Canada) was used as feedstock. The powder size distribution was measured by a laser particle size analyzer (Coulter LS, Fullerton, CA). Commercially available AZ91D-T4 magnesium substrates produced by High Pressure Die Casting (HPDC) were used after being submitted to a solution heat treatment at 400 °C for 12 h and followed by quenching in water. Before spraying the Al powders, the natural aging duration to T4 state substrates was no longer than 48 h and substrates with dimensions of 170 × 20 × 6 mm were grit blasted using (20 mesh – 24 grit) silica and cleaned using acetone. The substrates were then assembled on a twoaxis computer controlled moving system.
Table 1 Parameters for cold spray process used for spraying pure aluminum (Al-101) powders. Gas parameters Pressure (MPa)
Temperature (°C)
0.98
300
Spray distance (mm)
Number of pass
10
1
The coated samples were sectioned and prepared for microscopy analysis, following standard metallographic methods. The etching solutions for magnesium substrate and Al coating were 4% nitric acid + 96% ethanol and Kroll's reagent, respectively. The feedstock powder and coating microstructures were characterized using Optical Microscopy (OM) and Scanning Electron Microscopy (Zeiss EVO-MA10, LaB6 Analytical SEM) equipped with an Energy Disperse Spectroscopy (INCA X-act, Oxford, UK). Phase identification was investigated by X-ray diffraction (XRD) using a Philips X-Pert model 1830 X-ray diffractometer equipped with a graphite monochromator using Cu kα (λ = 0.15406 nm) radiation. Detailed scans were performed over a 20–100° 2θ range, 0.02° step width and 2 s per step acquisition time. Images of the cross sections were used to evaluate the coating porosity level using the commercially available Clemex imaging software. The powder and the coating microhardness were evaluated using a load of 25 g (HV0.025). All measurements used a dwell time of 15 s, using Duramin-2 Vickers hardness tester (Struers Inc., Cleveland, OH, USA). The reported hardness values are averages of at least 6 random measurements for each sample. The heat treatments of the as-sprayed coatings were carried out in a vacuum furnace under pressure of 5.0 × 10−4 bars. The samples were heated at constant heat treatment temperature (400 °C) for different period of time up to 20 h with a temperature fluctuation below 1 °C. 3. Results and discussion 3.1. Characteristics of pure Al powder The morphology of the Al-101 feedstock powder is shown in Fig. 1a. SEM observations of the free standing particles reveal the irregular morphology of the particles varying from spherical to elongated structures with different sizes. Fig. 1b presents an OM image of the polished cross section powder. The image shows a dendritic microstructure within the Al particle, caused by the rapid solidification rate experienced by the particle during the powder atomization process. The particle size distribution diagram obtained using laser size analyzer is shown in Fig. 1c. It indicates that the powder particles present a broad size distribution with a diameter ranging from 6 to 174 μm, and the particles larger than 50 μm represent approximately 40% of the particles. 3.2. Properties of the as sprayed coatings and substrate 3.2.1. Substrate The microstructure of the high pressure die casting AZ91D magnesium alloy is shown in Fig. 2a. It contains the primary α phase (Mg) matrix as well as the discontinuous intermetallic β phase (Mg17Al12) [28]. The later precipitates are located mostly at the α phase dendrite grain boundaries. After T4 heat treatment, the precipitates (β phase) were dissolved into the matrix and the microstructure of the Mg alloy displays a single Al supersaturated α phase as seen in Fig. 2b. Note that due to small size of clusters/GP zones that were formed inside the Al supersaturated grains during the natural aging, they could not be seen by using OM or SEM imaging technique. The XRD results of the as cast and T4 magnesium substrates are shown in Fig. 3, which also demonstrate that after solution treatment, the β phase peaks have disappeared. 3.2.2. Coating characteristics An SEM image under secondary electron mode of the as-sprayed coating cross section is shown in Fig. 4. The coating has a thickness of approximately 300 μm and limited porosity (less than 1%). The overall coating is dense with the exception of the top coating surface where limited porosity was observed. This is attributed to the fact that the large particles do not undergo enough plastic deformation upon impact and thus leave some porosity at the particle–particle
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Fig. 2. OM image showing (a) the microstructure of high pressure die casting AZ91D magnesium solid solution alloy α phase matrix and the β phase precipitates. (b) AZ91D magnesium alloy after T4 heat treatment (400 °C × 12 h and quenched in water).
Fig. 1. Al-101 feedstock powder: (a) SEM image shows the morphology of the free standing particles. (b) OM image of resin embedded powder cross section after polishing and weak etching revealing the dendrite microstructure of the Al particle. (c) Size distribution of the Al powder measured by laser size analyzer.
boundaries. In addition, the absence of significant peening effect that sometimes contributes to extra deformation of the particles already deposited leads to the increased porosity at the top of the coating. This phenomenon has already been described in other works [29]. The interface between the Al coating and the Mg substrate has almost no porosity, which means the particles near the interface had enough velocity to remove the surface oxide to expose fresh material, resulting in the intimate contact at the interface. This is considered beneficial for coating bonding strength and for enhanced diffusion of the elements during the subsequent heat treatment. The feedstock powder used in this study presents a large size distribution, ranging from less than 10 μm to over 170 μm. It is known that large powder particle size distributions have a detrimental effect on the coating quality because the large particles are harder to accelerate than the smaller ones while the latter may be affected by the bow shock wave present in front of the substrate [30]. Consequently, it is difficult for all of the particles to reach a similar high velocity, and thus it can be expected that the large feedstock powder particle size distribution used in this work will typically result in a large velocity distribution. As such, the particles which have a velocity lower than the critical velocity will rebound from the substrate surface or adhere to
the substrate with a low bonding strength and minimum deformation level, increasing the possibility of forming porosities. The morphology of the Al particles in the coating after etching is shown in Fig. 5. It can be seen that there were several large particles in the coating which were labeled by black arrow if their size was apparently larger than 50 μm, and the total area of the larger particles found in the coating was estimated to be 42%. Compared to the ratio of the particles larger than 50 μm in the feedstock powder, it is clear that the deposition efficiency is similar for all particle sizes in the current work. Since the produced coatings exhibit low porosity level, it is concluded that the disadvantages of large powder size distribution were alleviated by the choice of optimized spraying parameters.
Fig. 3. XRD results of the as cast and T4 treated AZ91D magnesium substrates.
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Fig. 4. SEM images of the cross section of Al coatings onto AZ91D magnesium.
Compared to the microhardness of Al feedstock powder of 39.6 ± 2.4 HV0.025, the hardness of the as sprayed coating was increased to 60± 3.4 HV0.025. It is concluded that the strain hardening effect due to the intense plastic deformation of the particles upon impact with the substrate played an important role on increasing the aluminum hardness as reported in other studies [23]. 3.3. Heat treatments of the coatings 3.3.1. Intermetallic phase Al–Mg binary phase diagram is shown in Fig. 6 [31]. It can be seen that all the Al–Mg intermetallic compound have a melting temperature lower than 450 °C and the Mg17Al12 phase has a wider composition range than Al3Mg2 phase at elevated temperature. The cold spray coating microstructure after the 4 h heat treatment at 400 °C is shown in Fig. 7a. Four distinctive zones can be observed. The first (the top layer) is the un-reacted as sprayed aluminum coating. The second and the third zones represent the intermetallic layers that formed at the coating/substrate interface. The fourth is the AZ91D magnesium substrate shown at the bottom. According to the Al–Mg binary phase diagram (Fig. 6), it is expected that the intermetallic compound near the substrate is Mg rich phase -Mg17Al12 (β phase) while the intermetallic layer present near the un-reacted Al coating will be the Al rich phase- Al3Mg2 (γ phase). Quantitative phase's composition was investigated using the EDS technique. It is accepted that in Al–Mg alloy systems, the composition ranges between 52–60 at.% Mg consist of a single β phase, while γ and R phases will be formed in the area range between 38.5–40.3 and 41.7–42.3 at.% Mg respectively. Note that the later, R phase can only exist between 320–370 °C, as it decomposes to β and γ phases [31].
Fig. 5. OM image of the etched as-sprayed coating showing the morphology of the Al particles and the arrows indicated the particles apparently larger than 50 μm.
Fig. 6. Al–Mg binary phase diagrams.
Three EDS analysis points were performed on the four different layers. According to the composition results, the intermetallic compounds near the magnesium substrate and the un-reacted aluminum coating were β phase (51.97–55.41 at.% Mg) and γ phase (39.23–39.27 at.% Mg), respectively. Moreover, EDS elemental line scan from the magnesium substrate side to the un-reacted aluminum coating side was conducted and the obtained quantitative composition is shown in Fig. 7b. It is observed that the β phase has a larger composition range compared to the γ phase and that the composition changes abruptly
Fig. 7. (a)SEM image of cold sprayed pure Al onto AZ91D- T4 substrate after heat treatment under vacuum at 400 °C- 4 h.revealing different intermetallic phases and (b) Quantitative EDS line scan as indicated by arrow in (a).
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at the interface of the two layers. It is worth noting that some porosity within the top layer of the γ phase was observed (arrows in Fig. 7a). Compared to Spencer [15] who also carried out heat treatment on the Al/Mg system, the volume of porosities on the coating cross section in the current study is smaller and only a few pores are observed. 3.3.2. Intermetallic thickness The intermetallic thickness variation with heat treatment time is shown in Fig. 8. It can be observed that the γ phase and β phase grow steadily for heat treatment at 400 °C under 16 h. For the 20 h heat treatment, the thickness of the γ phase layer increases at a much slower rate while the β phase does not show noticeable growth. These results are similar to previous studies [20], where a post annealing on Al/Mg at 413 °C was conducted, and found that the two intermetallic layers did not grow noticeably when the heat treatment time exceeded 24 h. Comparable conclusions can also be drawn from the linear regression growth curve shown in Fig. 9, where it can be observed that the slope of the γ phase is larger than β phase, and the former is almost 2.5 times higher than the latter. The curves reveal a linear trend that goes through the origin point, which demonstrates that the diffusion process obeys a parabolic rate law for heat treatment time no more than 20 h. Similar results have also been reported by Y. Funamizu and K. Watanabe [32] who suggested that the growth rate of the γ phase and β phase represent a linear relation with the square root of diffusion time when the heat treatment was performed at 425 °C and the slope of γ phase is almost 2.2 times of that for β phase. It implies that the intermetallics grow faster at the initial stage and then the overall concentration gradient decreased swiftly associated with the intermetallic thickness. It can be found in Fig. 8 that the migrating distance of Al atoms into Mg substrate is a little more than 60 μm, compared to less than 60 μm for the Mg atoms diffusing in the reverse direction when the heat treatment time is 4 h. This suggests that the aluminum diffuses more rapidly than magnesium in the system. Moreover, it is noted that when increasing the heat treatment time from 4 h to 20 h, the position of the β/γ interface is almost stable, which implies that these two intermetallics grow individually other than transform from one to the other. The intermetallic growth rate is mainly due to the atom diffusion rate in the two intermetallics. Thus, one of the reasons for higher growth rate of γ phase during the heat treatment is probably due to the higher diffusion rate of aluminum atoms in the two intermetallic phases [20,32]. Note that when two or more intermediate phases are formed in a diffusion zone in a binary system such Aluminum–Magnesium, many factors control the occurrence and growth of the various phases. The growth of a phase, in multiphase
Fig. 8. Thickness measurement of the intermetallic layers vs holding time of the heat treatment at 400 °C under vacuum. The horizontal zero axis corresponds to the original Mg/Al interface before the heat treatment.
Fig. 9. Growth curve of the intermetallic β- and γ-phases at 400 °C vs. square root diffusion time (h1/2).
diffusional growth is dependent on its own interdiffusion coefficient as well as those of its neighboring phases. Funamizu et al. [32] have attributed the difference in the thickness growth of the β and γ phase during heat treatment to the difference between the activation energy of layer growth and that of the interdiffusion of the two phases. The later was attributed to the difference in the composition range and temperature dependence of the two intermetallic phases as shown in the equilibrium phase diagram of the Al–Mg system (Fig. 6). High interdiffusion coefficient and low temperature dependence of the γ-phase composition contribute significantly in the growth of this intermetallic as compared to β-phase. More details about the different activation energies of the different phases within the Al–Mg system can be found elsewhere [32]. According to Fick's Second Law, the atom diffusion coefficient is mainly due to the concentration gradient and the temperature, and the diffusion distance is proportional to the square root of the heat treatment time. The thickness of the total intermetallics thicknessdiffusion distance-in this study obeys a parabolic rate law [32] (Fig. 9). However, Spencer et al. [15] suggested that the parabolic law is not applicable in the Al/Mg couple diffusion process because of the Kirkendall effect, which will lead to the condensation of vacancies near the edge of γ phase and then decrease the diffusion rate. It can be seen in Fig. 7a that there are few small pores near the edge of γ phase, as opposed to a large number of holes observed in Ref. [15] at the same position after diffusion, despite similar porosity level content in the as-sprayed coatings. The reason for the smaller porosities at the γ phase/un-reacted Al coating may be due to the large Al particles which were used as feedstock powder resulting in a coating with reduced number of Al/Al splats that would enhance the Mg diffusion in to the γ phase. As the thickness of the intermetallics obeys the parabolic law, increasing the heat treatment time or temperature to obtain thicker intermetallics is not a solution. First, the higher temperature will cause the melting of the intermetallic phases which were formed by thermal diffusion [4]; moreover, it will accelerate the grain growth of magnesium substrate which is detrimental to the mechanical performance based on the Hall–Petch relationship [33]. It has been recommended that the Al/Mg couple diffusion-induced phase should be performed between 400 °C to 436 °C [15], taking into account that the eutectic β phase will start local dissolution at 426 °C [5]. As such, the diffusion temperature was limited in a narrow extent and in the current study a temperature of 400 °C was selected, because it is the same temperature for AZ91D magnesium alloy solution treating. Consequently, during the diffusion heat treatment, the β phase will not precipitate and the substrate will remain a single α phase.
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When the as-sprayed coating samples are heat treated, the atoms can get enough energy and diffuse resulting in the formation of Al–Mg binary alloy band [20] between the Al coating and the Mg substrate. Generally, this binary band should contain αAl, αMg, γ phase and β phase with a concentration gradient according to the diffusion theory [22], but in fact it was found that it mainly consists of two different intermetallic compounds. Similar to other work [20], the layers of Mg matrix with saturated Al atoms and Al matrix with saturated Mg atoms were not detected in the OM and SEM images due to the limitation of the technique used in this work. Because Al/Mg diffusion is a reaction diffusion system, Al atoms and Mg atoms can react and form only one intermetallic compound when the atomic ratio arrive at a critical value at some particular point, according to the second law of thermodynamics. That is why the composition between the intermetallics and substrate/β phase and unreacted Al/γ phase will be changed abruptly rather than gradually. The typical optical microstructure of the as sprayed coating after heat treatment of 20 h is shown in Fig. 10. With a heat treatment time of 20 h at 400 °C, the β phase was 83 μm thick while the γ phase was 149 μm thick for a total thickness of 232 μm. Compared to other studies [20] which attained a total thickness of a little more than 160 μm below 413 °C for 24 h, the total intermetallic thickness is increased by 45%. The main reason for the thicker intermetallic layers found in the current work is believed to be that T4 instead of as-cast magnesium bulk [15] was used as substrates. To confirm this assumption, a comparison of the intermetallic thickness on T4 and as-cast AZ91D after 400 °C heat treatment for 4 h is shown in Fig. 11. The β phase and γ phase thicknesses were respectively 36.7 μm and 93.3 μm for T4 substrate and 30.8 μm and 83.2 μm for as-cast substrate. Generally, the aluminum content in the AZ91 magnesium bulk is near 9 wt.%. T4 substrates have a single solid solution α-phase with very fine clusters/ GP zones and Al atoms concentration is higher and more homogeneously distributed within Mg matrix. Hence intermetallic growth is faster as less enrichment is required to reach the critical level for intermetallic formation in the substrate. However, as-cast AZ91 substrates contain β phase precipitates and α phase Mg matrix. The β phase precipitates have aluminum content of 44.3%, leading to a magnesium matrix containing less than 9% aluminum because of beta precipitation. Thus, the as-cast state has more Al tied up in the beta phase and consequently less available in the Mg matrix to form intermetallics. Moreover, it is expected that the dense coating and the
Fig. 10. OM image of the cold sprayed Al coating onto AZ91D-T4 substrate after heat treatment at 400 °C for 20 h.
Fig. 11. Comparison of the intermetallic thicknesses formed on the coating which sprayed onto as cast and T4 substrates heat treated at 400 °C during 4 h.
intimate contact of the Al/Mg interface and the Al splats, which were dramatically influenced by the feedstock powders and the spraying conditions, were believed to have a positive effect on Al diffuse process. 3.3.3. Hardness The purpose of cold spray coatings and the intermetallics formed by post heat treatment is providing a protection of magnesium substrates from the environment. Any notable degradation of the substrate due to the heat treatment should be avoided. As such, microhardness testing was used to evaluate the mechanical properties of the Mg substrates. The microhardness of the Mg substrate and of the two intermetallic layers as a function of heat treatment time is shown in Fig. 12. The hardness value of as cast and T4 Mg substrates were 80.7 ± 1.8 HV0.025 and 73.7 ± 4.0 HV0.025, respectively. This hardness decrease was mainly due to the dissolution of the hard β phase, which acts as a reinforcement phase in the AZ magnesium alloy series, in the Mg matrix. After the diffusion heat treatment, the β phase was precipitated from the Mg matrix which saturated with Al atoms during the cooling process, and the hardness of Mg substrate increased to around 91 HV0.025. Despite the diffusion treatment being carried out for different holding time, the Mg substrate hardness was almost the same, because the substrate maintains a single α phase when the furnace temperature is kept at 400 °C and the β phase could not precipitate in this process. All the samples in this study were
Fig. 12. Microhardness results of the Mg substrates and the intermetallic layers. measured after 400 °C heat treatment for different period of time.
H. Bu et al. / Surface & Coatings Technology 205 (2011) 4665–4671
cooled from 400 °C to room temperature according to the same cooling condition. Therefore, it is expected that the Mg substrates should have the same β phase content which is considered as a reinforcement phase in AZ series Mg alloys and could only precipitate in the cooling process. Moreover, the hardness of the γ phase was slightly larger than the β phase, 279.6 ± 13.7 HV0.025 and 260.5 ± 10.7 HV0.025 respectively. Similar values have been reported by Spencer et al. [15]. Note that the measured hardness is 2 times higher than the as cast AZ91D magnesium substrate. It is widely accepted that the higher hardness of the coating is considered helpful to attain a higher wear resistance [34]. 4. Conclusion The cold spray parameters and nozzle used in this study are appropriate for spraying large aluminum powders with wide size distribution and a dense coating is attained which has a porosity content less than 1%. The growth rate of the Al–Mg intermetallic phases during annealing at 400 °C obeys a parabolic law and the Alenrich phase (γ phase) grows much faster- around 2.5times- than the Mg-enrich phase (β phase). Increased diffusion time will lead to thicker intermetallics between un-reacted Al coating and Mg substrate. The total thickness of the intermetallics after 20 h diffusion is more than 230 μm, which is 45% thicker than thicknesses that have been reported in the literature. Considering the relative lower temperature and holding time used in this study, it is believed that the coating can provide better protection and less detrimental effect for the Mg substrate. Compared to the AZ91D magnesium alloy bulk, no matter as cast or T4 status, the intermetallic compounds which formed after cold spray and post heat treatment have a much higher hardness, increasing the possibility of promoting the wear resistance of magnesium alloy surface. As a result, the pure Al coatings deposited on Mg substrates using cold spray and post heat treatment can be treated as an alternative method for magnesium protection. Acknowledgments The authors would like to thank the Meridian Technologies (Ontario, Canada) for providing the AZ91D Magnesium materials and Shanghai JiaoTong University for the financial support given to this work.
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