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The negative effect of solution treatment on the age hardening of A356 alloy
Materials Science & Engineering A 566 (2013) 112–118
Contents lists available at SciVerse ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
The negative effect of solution treatment on the age hardening of A356 alloy H.C. Long a, J.H. Chen a, C.H. Liu a,n, D.Z. Li b, Y.Y. Li b a b
College of Materials Science and Engineering, Hunan University, Changsha 410082, China Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 30 October 2012 Received in revised form 10 December 2012 Accepted 13 December 2012 Available online 12 January 2013
The effect of solution treatment on the age hardening of an A356 (Al–7.0Si–0.3Mg) alloy was investigated by hardness measurements and electron microscopy. This study presents experimental evidences that a prolonged solution treatment has a detrimental influence on age hardening due to the porosity in the cast A356 alloy. The enrichment of Mg on the internal surface of the porosity was observed. Nanoscale particles of magnesium oxides exist in the samples that have been solution treated with an extended duration. Our results indicate that Mg atoms would diffuse toward the porosity pre-existed in the A356 alloy during solution treatment. This phenomenon is consistent with the degraded aging kinetics and hardening potential of the A356 alloys solution treated for a long time, since the reduction of Mg solutes in the Al matrix shall lead to a decrease of the volume fraction of the Mg–Si-containing strengthening precipitates, which are mainly the monoclinic b00 phase. Our findings emphasize the significance of controlling porosities for the A356 alloys and may provide useful information for designing thermal treatment parameters to reach the optimum balance between strength and ductility. & 2013 Elsevier B.V. All rights reserved.
Keywords: A356 alloy Porosity Aging Solution treatment Electron microscopy
1. Introduction A356 alloys (Al–Si–Mg) are extensively used as structural materials in many industries, for instance, automotive and aerospace, for their excellent castability, high strength-to-density ratio, good fatigue and corrosion resistance. The cast components are generally heat treated to the T6 condition, including solution treatment, water quenching, natural and artificial aging. Extensive research efforts have focused on the influence of solution treatment on the microstructure and thus mechanical property of A356 alloy [1–6]. The main objective of solution treatment in the range of 540– 550 1C for A356 alloy is to thermally alter silicon particle characteristics and dissolve Mg2Si particles to achieve a homogeneous solid solution [1]. Dissolution of Mg2Si and homogenization of the matrix, which occur within 15 min at 540 1C in A356 alloys [2], can be quickly reached. However, it usually takes a relatively long time to change the morphology of silicon from a polyhedral to globular structure, which is highly responsible for improving ductility [4]. In the foundry industry, it was frequently stated that a component made of dendritic A356 alloy should be solution treated at 540 1C for no less than 6 h [5,6]. The nano-particles precipitated out from the supersaturated solid solution containing excess magnesium and silicon can render intense age hardening to the A356 alloy [7,8]. Therefore, the artificial aging are always applied after solution treatment followed by water quenching. In a word, the A356 alloy
castings are generally heat treated to obtain a desired combination of strength and ductility. During solidification of cast A356 alloys, porosity was inevitably introduced owing to the precipitation of hydrogen from liquid solution (gas porosity) or by shrinkage (shrinkage porosity), and more usually by a combination of these effects [9–11]. Even nowadays, it is still hard to avoid porosity in A356 alloys, especially when produced as large or complex cast products. And porosity maybe the most persistent and common complaint of casting users and contributes directly to customer concerns about reliability and quality [11,12]. Deleterious impacts that porosity brings are usually believed to be caused by the reduction in effective area by pore volume fraction and by stress concentrations at voids leading to premature failure [13–16]. Tremendous researches have been done concerning the effect of porosity with different type, morphology, and size on the mechanical properties of Al–Si–Mg alloys [13–17]. However, little attention has yet been paid to the evolution of porosity whose content may change during solution treatment at a relatively high temperature. Furthermore, it is also unclear what this change brings to the age hardening of A356 alloy. In the present work, it was found that the extending of solution treatment depressed the age hardening of A356 alloy. The origin of this effect was clarified through deliberate investigation by scanning electron microscopy (SEM) and transmission electron microscopy (TEM). 2. Experimental procedure
n
Corresponding author. Tel.: þ86 731 88664009; fax: þ86 731 88664010. E-mail address: [email protected] (C.H. Liu).
0921-5093/$ - see front matter & 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.12.093
The A356 alloy, with chemical composition Al–7.0Si–0.3Mg (wt%), used in this study was prepared by permanent mold cast
H.C. Long et al. / Materials Science & Engineering A 566 (2013) 112–118
process. All samples were made by cutting out pieces 10 10 1 mm3 from the as-cast specimen. The solution treatment was performed in an air-recirculating furnace at 550 1C for three different durations: 30 min, 2 h and 10 h. The samples were then transferred into the water at room temperature to apply quenching. In order to study the influence of solution treatment time on the age hardening response in A356 alloy, artificial aging was conducted subsequently in an oil bath at 180 1C from 10 min to 48 h. Vickers micro-hardness was determined using a load of 4.9 N and a dwell time of 10 s. Each datum was the average value of at least 10 indentations. The error bars, which were given as standard deviations, were used to show variability in the measures. For the purpose of exploring the composition change of the porosities during solution treatment, we had taken an ‘‘in-situ’’ observation procedure. To be specific, we adopted energy dispersive X-ray spectroscopy (EDS) line-scanning to observe the variation of element concentration at the same place around a typical porosity with solution treatment. Firstly, a specimen was prepared by grinding with SiC abrasive paper and mechanical polishing. After cleaning by ethanol and drying, the specimen was then investigated in the SEM. The element concentration profiles across a region near the porosities owning special features were recorded. Thirdly, the specimen was taken out from the SEM apparatus and put in the solution treatment furnace at 550 1C for 0.5 h. Subsequently, the specimen was rapidly cooled and taken back to the SEM again for element concentration detection at the same region near the same porosities with those in the previous observations. Similarly, the element concentration across a region near the porosities in the specimen solution treated for 2 h and 8 h were also recorded by the above procedure. The successive steps of ‘‘in-situ’’ observation are shown in Fig. 1. Through comparison of the acquired element concentrations with the original ones, we got the composition change of the porosity with the prolonging of solution treatment time. Following grinding and punching, thin-foil TEM specimens were prepared by electro-polishing in an electrolyte consisted of 1/3 HNO3 in methanol. The temperature was kept below 25 1C during polishing. In addition, the TEM specimens were also used for characterizing the morphology of Si particles by SEM. As these specimens were deeply etched, more details about the Si particles were obtained. The SEM micrographs were achieved through an FEI QUANTA 200 microscope in the secondary electron mode. This instrument was equipped with an EDAX energy dispersive X-ray spectrometer for composition analysis. All the TEM observations
Fig. 1. Successive steps of ‘‘in-situ’’ observation.
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mentioned above were performed using a JEOL JEM-3010 microscope, operated at 300 kV.
3. Results 3.1. The evolution of eutectic Si particles during solution treatment Typical microstructures of A356 cast alloys solution treated with various times are shown in Fig. 2. The as-cast alloy comprises mainly the primary a-Al and the connected eutectic silicon with a coral-like or acicular morphology (Fig. 2a). The solution treatment for 30 min at 550 1C has a slight effect on the morphology of the eutectic silicon particles (Fig. 2b). When the solution time increased to 2 h, the aspect ratios of the silicon particles just decreased a few (Fig. 2c). Though the eutectic silicon had spheroidized and coarsened to some extent, most of the eutectic silicon particles were still of the platelets or longish rod-like morphology. A remarkable change of silicon particles occurred when the solution treatment time extended to 10 h (Fig. 2d). The initial fiber-like Si particles underwent necking and were broken down into smaller fragments, and the aspect ratio of the silicon particles decreased notably. And the spacing between the particles increased evidently. It has been proved that a structure containing large and elongated particles but small dislocation slip distances (within the dendrite cells) exhibited a high particle cracking rate and thus a low ductility [18]. Therefore, especially for unmodified A356 alloy, structure refinement by means of solution treatment for enough time is very important in improving the ductility. 3.2. The influence of solution treatment on the age hardening response The hardness curves of the alloys solution treated at 550 1C (either for 30 min, 2 h, or 10 h) during artificial aging can be seen in Fig. 3. Each curve can be divided roughly into three stages. At the first stage, there is a rapid increase in hardness with aging time. Then the hardness goes steadily up until it reached the peak with an aging time of 8 h 10 h. At the final stage, the hardness of the alloy gradually decreases as a result of over-aging. The hardness evolution with aging time in the present work is consistent with the results many researchers had reported [19,20]. However, the age hardening potential dropped off significantly with the increasing of solution treatment time. As shown in Fig. 3, by increasing solution time from 30 min to 2 h, the peak-aging hardness decreased from 118 to 102 HV. With further increase in solution time up to10 h, the peak hardness decreased to 95 HV. The average hardness at a given aging time fell about 5 to 20 HV when the solution treatment time was increased from 30 min to 2 h. The trend was also found when the solution time was extended from 2 h to 10 h. The microstructure formed by aging for 8 h (corresponding to the peak-strength) for the specimen solution treated for 30 min is given in Fig. 4. In Fig. 4a, a bright-field TEM image, numerous fine and needle-like precipitates were observed to be aligned along the o1004 directions of the Al matrix. Fig. 4b shows a typical highresolution TEM image of the precipitates in Fig. 4a. After determination of the lattice parameters, this precipitate was identified to be monoclinic b00 (Mg5Si6) fully coherent with the matrix [21]. This kind of structure was also observed in the peak-aged specimens solution treated for 2 h and 10 h. Looking similar in TEM images, the microstructures of these peak-aged samples are not shown. Fig. 2 exhibits that the size of Si particles increased and their aspect ratio decreased with increasing solution treatment time. This would result in a slight decrease in hardness for the specimen solution treated for a long time [22]. As shown in Fig. 3, there
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Fig. 2. Microstructures of the unmodified A356 alloy obtained from (a) as-cast specimen, (b) specimen solution treated at 550 1C for 30 min, (c) specimen solution treated at 550 1C for 2 h, and (d) specimen solution treated at 550 1C for 10 h.
3.3. Enrichment of Mg near the porosity
Fig. 3. Age-hardening curves of A356 alloys solution treated with different times.
is only a small difference (3–6 HV) between the quenched hardness. Except for impacts from silicon particles, the hardening effect is determined by the amount of Mg–Si containing precipitates (b00 ) generated via employing an aging treatment to a-Al matrix of A356 alloys [23]. Therefore, the suppressed age hardening response due to a long solution time was primarily attributed to the decline of solutes used to form strengthening precipitates. Since only Mg/Si elements are included in the precipitates and there is excess Si in the A356 alloy to form saturated Si solution, the Mg loss in the a-Al matrix devotes to the decline of solutes.
Following the procedure described in Section 2, the composition change of the porosity with the prolonging of solution treatment time was determined. Distribution of porosity in the as-cast sample is shown in Fig. 5a. There were several shrinkage porosities in the as-cast sample which had an irregular shape, corresponding to the shape of the interdendritic region [9,10]. Fig. 5b is the partial enlarged drawing of a typical porosity in Fig. 5a and shows a region at the interface between the porosity and a-Al matrix. The spectra measured by EDS spot analysis at the point indicated by a cross-like symbol in Fig. 5b is shown in Fig. 5c. It can be seen that the primary elements on the surface of the porosity in the as-cast specimen included Al, Si, O, Mg and the percentage of Mg concentration was very small comparatively. The EDS element line profiles of Mg corresponding to the same place (red line in Fig. 5b) of the same specimen, but with different solution treatment time, are shown in Fig. 5d. For as-cast condition, it is clear that very low Mg concentration was detected both in the porosity and the a-Al matrix, albeit a small peak existed at the interface between them. Because of its high reactivity, magnesium oxidizes readily in liquid and solid states. Hence, the small peak signal might be yielded from an amorphous magnesium oxide which was most likely to be one kind of nonmetallic inclusions formed during foundry processes [24]. Strikingly, the Mg concentration in the porosity and surrounding area increased significantly when the as-cast specimen was solution treated for 30 min. Moreover, this phenomenon became more obvious with solution treatment time extending. Mg enrichment around porosities during solution treatment is evident by the present finding. The intensifying of Mg content near the porosity will result in a reduction of effective solutes used for
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Fig. 4. (a) TEM micrograph recorded in [001]Al zone axis of typical Mg–Si-containing strengthening precipitates (b00 precursors) obtained from the specimen solution treated for 30 min and aged for 8 h at 180 1C. (b) High-resolution TEM micrograph recorded in [001]Al zone axis of the cross-section of a needle-like precipitate in Fig. 4a, the unit-cell was overlaid in the image.
Fig. 5. (a) Distribution of porosity in the as-cast sample, (b) a partial enlarged drawing of a typical shrinkage porosity in (a), which showed the interface region between the porosity and a-Al matrix, (c) EDS spot analysis results detected from a point, indicated by red cross symbol in the porosity of (b), (d) EDS line profiles, along the designated red lines, of Mg concentration detected from a specimen solution treated with different time. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article).
forming Mg–Si-containing strengthening precipitates in the a-Al matrix during the following aging treatment. 3.4. Existence of MgO nanoscale particles in the solution treated A356 alloy In order to exclude the impact of surface oxidation on the conclusion, TEM and EDS were used for further clarifying the phenomenon mentioned above. Fig. 6 shows the TEM micrograph of a specimen solution treated at 550 1C for 10 h. It is surprising
that flakes in Fig. 6 were frequently observed on the edge of the penetrated hole in the TEM thin-foil. From the ring-like selected area diffraction (SAED) pattern (inset in Fig. 6), it is inferred that plenty of nanoscale particles existed in the specimen. Fig. 7 presents the magnified SAED pattern and the corresponding EDS spectra obtained from a certain area in Fig. 6. The several sharp concentric rings observed in this pattern suggest polycrystalline nature of this region. The d-spacing determined from the ring diameter exhibits in the lower right corner of Fig. 7. Analysis of the diffraction pattern and interplanar spacing reveal that the
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nanoscale particles correspond to a cubic system with interference indices (200) and (220) as the brightest rings. The calculation is in good agreement with the data of MgO PDF card. The estimation strongly recommends that these nanoscale particles are MgO crystals. This conclusion is also confirmed by the EDS result which shows Mg and O elements account for a quite large proportion of the content. So the flake in Fig. 6 may be oxide bifilm (MgO and Al2O3) formed around the porosity during solution treatment. This kind of flake containing MgO nanoscale particles was also abundant in the specimen solution treated at 550 1C for 2 h. A few MgO particles were also found in the specimen solution treated at 550 1C for 30 min, while none was observed in the as-cast specimen. On the basis of these facts, we can arrive at the conclusion that these observed MgO nanoparticles are highly correlated with migration of Mg toward the porosity during solution treatment. The formation of MgO nanoparticles will be discussed below.
Fig. 6. TEM micrograph of a region in the specimen solution treated at 550 1C for 10 h, the set is the corresponding SAED pattern.
4. Discussion In order to explain the phenomenon discovered in our study, a schematic diagram (Fig. 8) is proposed to illustrate the mechanism of Mg loss in A356 alloys during solution treatment. Once the as-cast specimen is solution treated at a relatively high temperature, magnesium and a certain amount of silicon would dissolve in the a-Al matrix to produce a homogeneous solid solution within a short time. It was shown that multiple precipitating Si nanotwins could form on the eutectic Si particles in Al–Si alloys upon annealing [25]. We surmise that these Si particles are surrounded by a-Al matrix with a loose interface, which can assist the rapid diffusion of dissolved atoms. As shown in the middle of Fig. 8, the gray shaded oval particles represent eutectic Si particles. We use dotted lines to represent the interfaces between eutectic Si particles and the a-Al matrix. According to Kerte´sz et al., a prolonged high-temperature heat treatment of Al– Mg–Si alloys produced Mg enrichment on the surface [26]. As Mg atoms are outdiffusing from the Al matrix during solution treatment, the amount of hardening precipitates formed during aging treatment will decrease and thus lead to a decline of the aged hardness. However, the part of Mg enriched on the surface is limited and cannot explain the huge difference between the age hardening curves listed in Fig. 3. It can also be seen from Fig. 5d that the Mg content variation at the matrix near the porosity was much smaller than that in porosity. Two isolated dark irregular polygons in the oval at the left of Fig. 8 represent shrinkage porosity (the irregular one formed corresponding to the shape of the interdendritic region) and gas porosity (the smoother one that is approximately rounded), respectively. During solidification of cast A356 alloys, the oxidation of molten aluminum occurred and Al2O3 film formed at the internal surface of the micro-void [16]. Owing to the fact that the formation enthalpy of MgO is far lower than Al2O3 [27], the chemical affinity of Mg with the oxygen in the oxide film at the internal surface of the micro-void may be the driving force for Mg diffusion to the micro-voids, especially when the gas content inside a micro-void in A356 alloy is restricted from outer atmosphere. The thermal analysis performed by differential scanning calorimetry proved that even the stable b (Mg2Si) phase in a-Al was not stable above 520 1C [7], so all the Mg–Si phases would
Fig. 7. Analysis of polycrystalline electron diffraction pattern and the EDS spectra achieved from a part in Fig. 6.
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Fig. 8. Schematic illustration of the Mg loss process during solution treatment.
diffuse into the aluminum matrix during solution treatment. But the MgO formed near the porosities should be very stable, as the melting temperature of MgO is about 2800 1C, far exceeding the solution temperature. For A356 alloy, the magnesium depletion during solution treatment can get worse as a result of the ubiquity of porosity in the casting alloy. The thermally activated Mg atoms will diffuse through the loose interfaces mentioned above toward the porosities. Though the mechanism for the formation of MgO nanocrystalline remains unclear and needs further research, it is yet a definite evidence of Mg enrichment inside the A356 alloys during solution treatment. Because of shortened diffusion distance, the Mg enrichment in porosity is easier than on the sample surface. Therefore, Mg enrichment in porosity during solution treatment is the main reason of magnesium depletion of a-Al matrix for A356 cast alloy with a lot of porosities. Mg concentration in porosity would significantly increase during solution treatment, as described in Sections 3.3 and 3.4. Mg enrichment in porosity will result in Mg loss in Al matrix and influence the precipitation kinetics in the following aging treatment. Consequently, the alloy may lose its expected hardness seriously if the alloy with a lot of porosities was solution treated for a long time. In view of the evil influence mentioned above, more attention should be paid to the control of the quantity and size of porosities to enhance the mechanical property of A356 alloy. Otherwise, solution treatment time should be restricted in some range to reduce the loss of Mg under the premise of guaranteeing the requirements of ductility. Based on the discussion above, it is reasonable for us to draw the conclusion that a specimen containing more porosity will obtain lower hardness for lower Mg concentration used for forming Mg–Si-containing strengthening precipitates.
5. Conclusions The present study has demonstrated the microstructure evolution and the age hardening characteristic of an A356 alloy in relation with solution treatment. From the obtained results, the following can be concluded. (1) Apart from the spheroidization of the Si particles during solution treatment, thermally activated Mg atoms can diffuse out from the Al-matrix and enrich near the porosity preexisted in the alloy. (2) A prolonged solution treatment has a negative influence on the age hardening due to the existence of porosity in the cast A356 alloy. The monoclinic b00 precipitates play a crucial role in the age hardening response of the alloy. (3) The diffusion of Mg atoms toward porosities during solution treatment may lead to a significant reduction of the Mg solutes available to form Mg–Si-containing strengthening
precipitates in the Al-matrix upon aging treatment. The Mg enrichment around porosity becomes more severe with the increase of solution treatment time. (4) Nanoscale particles of magnesium oxides were found in the specimen solution treated for an extended time, while none was observed in the as-cast specimen. The formation of MgO nanoparticles is correlated with the Mg enrichment in porosity occurred during solution treatment. (5) This study indicates that controlling the porosity formed in the casting process is crucial for enhancing the mechanical properties of the A356 alloy. For Al–Si–Mg alloys with abundant porosities, the solution treatment time should be limited to an optimum to reach a balance between strength and ductility of the alloy.
Acknowledgments This work is financially supported by the National Basic Research (973) Program of China (No. 2009CB623704); the National Natural Science Foundation of China (No. 51171063); Instrumental Innovation Foundation of Hunan Province (No. 2011TT1003); the Aid Program for Science and Technology Innovative Research Team in Higher Educational Institutions of Hunan Province, Hunan Provincial Innovation Foundation For Postgraduate and Gatan China Scholarship.
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Contents lists available at SciVerse ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
The negative effect of solution treatment on the age hardening of A356 alloy H.C. Long a, J.H. Chen a, C.H. Liu a,n, D.Z. Li b, Y.Y. Li b a b
College of Materials Science and Engineering, Hunan University, Changsha 410082, China Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang 110016, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 30 October 2012 Received in revised form 10 December 2012 Accepted 13 December 2012 Available online 12 January 2013
The effect of solution treatment on the age hardening of an A356 (Al–7.0Si–0.3Mg) alloy was investigated by hardness measurements and electron microscopy. This study presents experimental evidences that a prolonged solution treatment has a detrimental influence on age hardening due to the porosity in the cast A356 alloy. The enrichment of Mg on the internal surface of the porosity was observed. Nanoscale particles of magnesium oxides exist in the samples that have been solution treated with an extended duration. Our results indicate that Mg atoms would diffuse toward the porosity pre-existed in the A356 alloy during solution treatment. This phenomenon is consistent with the degraded aging kinetics and hardening potential of the A356 alloys solution treated for a long time, since the reduction of Mg solutes in the Al matrix shall lead to a decrease of the volume fraction of the Mg–Si-containing strengthening precipitates, which are mainly the monoclinic b00 phase. Our findings emphasize the significance of controlling porosities for the A356 alloys and may provide useful information for designing thermal treatment parameters to reach the optimum balance between strength and ductility. & 2013 Elsevier B.V. All rights reserved.
Keywords: A356 alloy Porosity Aging Solution treatment Electron microscopy
1. Introduction A356 alloys (Al–Si–Mg) are extensively used as structural materials in many industries, for instance, automotive and aerospace, for their excellent castability, high strength-to-density ratio, good fatigue and corrosion resistance. The cast components are generally heat treated to the T6 condition, including solution treatment, water quenching, natural and artificial aging. Extensive research efforts have focused on the influence of solution treatment on the microstructure and thus mechanical property of A356 alloy [1–6]. The main objective of solution treatment in the range of 540– 550 1C for A356 alloy is to thermally alter silicon particle characteristics and dissolve Mg2Si particles to achieve a homogeneous solid solution [1]. Dissolution of Mg2Si and homogenization of the matrix, which occur within 15 min at 540 1C in A356 alloys [2], can be quickly reached. However, it usually takes a relatively long time to change the morphology of silicon from a polyhedral to globular structure, which is highly responsible for improving ductility [4]. In the foundry industry, it was frequently stated that a component made of dendritic A356 alloy should be solution treated at 540 1C for no less than 6 h [5,6]. The nano-particles precipitated out from the supersaturated solid solution containing excess magnesium and silicon can render intense age hardening to the A356 alloy [7,8]. Therefore, the artificial aging are always applied after solution treatment followed by water quenching. In a word, the A356 alloy
castings are generally heat treated to obtain a desired combination of strength and ductility. During solidification of cast A356 alloys, porosity was inevitably introduced owing to the precipitation of hydrogen from liquid solution (gas porosity) or by shrinkage (shrinkage porosity), and more usually by a combination of these effects [9–11]. Even nowadays, it is still hard to avoid porosity in A356 alloys, especially when produced as large or complex cast products. And porosity maybe the most persistent and common complaint of casting users and contributes directly to customer concerns about reliability and quality [11,12]. Deleterious impacts that porosity brings are usually believed to be caused by the reduction in effective area by pore volume fraction and by stress concentrations at voids leading to premature failure [13–16]. Tremendous researches have been done concerning the effect of porosity with different type, morphology, and size on the mechanical properties of Al–Si–Mg alloys [13–17]. However, little attention has yet been paid to the evolution of porosity whose content may change during solution treatment at a relatively high temperature. Furthermore, it is also unclear what this change brings to the age hardening of A356 alloy. In the present work, it was found that the extending of solution treatment depressed the age hardening of A356 alloy. The origin of this effect was clarified through deliberate investigation by scanning electron microscopy (SEM) and transmission electron microscopy (TEM). 2. Experimental procedure
n
Corresponding author. Tel.: þ86 731 88664009; fax: þ86 731 88664010. E-mail address: [email protected] (C.H. Liu).
0921-5093/$ - see front matter & 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.12.093
The A356 alloy, with chemical composition Al–7.0Si–0.3Mg (wt%), used in this study was prepared by permanent mold cast
H.C. Long et al. / Materials Science & Engineering A 566 (2013) 112–118
process. All samples were made by cutting out pieces 10 10 1 mm3 from the as-cast specimen. The solution treatment was performed in an air-recirculating furnace at 550 1C for three different durations: 30 min, 2 h and 10 h. The samples were then transferred into the water at room temperature to apply quenching. In order to study the influence of solution treatment time on the age hardening response in A356 alloy, artificial aging was conducted subsequently in an oil bath at 180 1C from 10 min to 48 h. Vickers micro-hardness was determined using a load of 4.9 N and a dwell time of 10 s. Each datum was the average value of at least 10 indentations. The error bars, which were given as standard deviations, were used to show variability in the measures. For the purpose of exploring the composition change of the porosities during solution treatment, we had taken an ‘‘in-situ’’ observation procedure. To be specific, we adopted energy dispersive X-ray spectroscopy (EDS) line-scanning to observe the variation of element concentration at the same place around a typical porosity with solution treatment. Firstly, a specimen was prepared by grinding with SiC abrasive paper and mechanical polishing. After cleaning by ethanol and drying, the specimen was then investigated in the SEM. The element concentration profiles across a region near the porosities owning special features were recorded. Thirdly, the specimen was taken out from the SEM apparatus and put in the solution treatment furnace at 550 1C for 0.5 h. Subsequently, the specimen was rapidly cooled and taken back to the SEM again for element concentration detection at the same region near the same porosities with those in the previous observations. Similarly, the element concentration across a region near the porosities in the specimen solution treated for 2 h and 8 h were also recorded by the above procedure. The successive steps of ‘‘in-situ’’ observation are shown in Fig. 1. Through comparison of the acquired element concentrations with the original ones, we got the composition change of the porosity with the prolonging of solution treatment time. Following grinding and punching, thin-foil TEM specimens were prepared by electro-polishing in an electrolyte consisted of 1/3 HNO3 in methanol. The temperature was kept below 25 1C during polishing. In addition, the TEM specimens were also used for characterizing the morphology of Si particles by SEM. As these specimens were deeply etched, more details about the Si particles were obtained. The SEM micrographs were achieved through an FEI QUANTA 200 microscope in the secondary electron mode. This instrument was equipped with an EDAX energy dispersive X-ray spectrometer for composition analysis. All the TEM observations
Fig. 1. Successive steps of ‘‘in-situ’’ observation.
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mentioned above were performed using a JEOL JEM-3010 microscope, operated at 300 kV.
3. Results 3.1. The evolution of eutectic Si particles during solution treatment Typical microstructures of A356 cast alloys solution treated with various times are shown in Fig. 2. The as-cast alloy comprises mainly the primary a-Al and the connected eutectic silicon with a coral-like or acicular morphology (Fig. 2a). The solution treatment for 30 min at 550 1C has a slight effect on the morphology of the eutectic silicon particles (Fig. 2b). When the solution time increased to 2 h, the aspect ratios of the silicon particles just decreased a few (Fig. 2c). Though the eutectic silicon had spheroidized and coarsened to some extent, most of the eutectic silicon particles were still of the platelets or longish rod-like morphology. A remarkable change of silicon particles occurred when the solution treatment time extended to 10 h (Fig. 2d). The initial fiber-like Si particles underwent necking and were broken down into smaller fragments, and the aspect ratio of the silicon particles decreased notably. And the spacing between the particles increased evidently. It has been proved that a structure containing large and elongated particles but small dislocation slip distances (within the dendrite cells) exhibited a high particle cracking rate and thus a low ductility [18]. Therefore, especially for unmodified A356 alloy, structure refinement by means of solution treatment for enough time is very important in improving the ductility. 3.2. The influence of solution treatment on the age hardening response The hardness curves of the alloys solution treated at 550 1C (either for 30 min, 2 h, or 10 h) during artificial aging can be seen in Fig. 3. Each curve can be divided roughly into three stages. At the first stage, there is a rapid increase in hardness with aging time. Then the hardness goes steadily up until it reached the peak with an aging time of 8 h 10 h. At the final stage, the hardness of the alloy gradually decreases as a result of over-aging. The hardness evolution with aging time in the present work is consistent with the results many researchers had reported [19,20]. However, the age hardening potential dropped off significantly with the increasing of solution treatment time. As shown in Fig. 3, by increasing solution time from 30 min to 2 h, the peak-aging hardness decreased from 118 to 102 HV. With further increase in solution time up to10 h, the peak hardness decreased to 95 HV. The average hardness at a given aging time fell about 5 to 20 HV when the solution treatment time was increased from 30 min to 2 h. The trend was also found when the solution time was extended from 2 h to 10 h. The microstructure formed by aging for 8 h (corresponding to the peak-strength) for the specimen solution treated for 30 min is given in Fig. 4. In Fig. 4a, a bright-field TEM image, numerous fine and needle-like precipitates were observed to be aligned along the o1004 directions of the Al matrix. Fig. 4b shows a typical highresolution TEM image of the precipitates in Fig. 4a. After determination of the lattice parameters, this precipitate was identified to be monoclinic b00 (Mg5Si6) fully coherent with the matrix [21]. This kind of structure was also observed in the peak-aged specimens solution treated for 2 h and 10 h. Looking similar in TEM images, the microstructures of these peak-aged samples are not shown. Fig. 2 exhibits that the size of Si particles increased and their aspect ratio decreased with increasing solution treatment time. This would result in a slight decrease in hardness for the specimen solution treated for a long time [22]. As shown in Fig. 3, there
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Fig. 2. Microstructures of the unmodified A356 alloy obtained from (a) as-cast specimen, (b) specimen solution treated at 550 1C for 30 min, (c) specimen solution treated at 550 1C for 2 h, and (d) specimen solution treated at 550 1C for 10 h.
3.3. Enrichment of Mg near the porosity
Fig. 3. Age-hardening curves of A356 alloys solution treated with different times.
is only a small difference (3–6 HV) between the quenched hardness. Except for impacts from silicon particles, the hardening effect is determined by the amount of Mg–Si containing precipitates (b00 ) generated via employing an aging treatment to a-Al matrix of A356 alloys [23]. Therefore, the suppressed age hardening response due to a long solution time was primarily attributed to the decline of solutes used to form strengthening precipitates. Since only Mg/Si elements are included in the precipitates and there is excess Si in the A356 alloy to form saturated Si solution, the Mg loss in the a-Al matrix devotes to the decline of solutes.
Following the procedure described in Section 2, the composition change of the porosity with the prolonging of solution treatment time was determined. Distribution of porosity in the as-cast sample is shown in Fig. 5a. There were several shrinkage porosities in the as-cast sample which had an irregular shape, corresponding to the shape of the interdendritic region [9,10]. Fig. 5b is the partial enlarged drawing of a typical porosity in Fig. 5a and shows a region at the interface between the porosity and a-Al matrix. The spectra measured by EDS spot analysis at the point indicated by a cross-like symbol in Fig. 5b is shown in Fig. 5c. It can be seen that the primary elements on the surface of the porosity in the as-cast specimen included Al, Si, O, Mg and the percentage of Mg concentration was very small comparatively. The EDS element line profiles of Mg corresponding to the same place (red line in Fig. 5b) of the same specimen, but with different solution treatment time, are shown in Fig. 5d. For as-cast condition, it is clear that very low Mg concentration was detected both in the porosity and the a-Al matrix, albeit a small peak existed at the interface between them. Because of its high reactivity, magnesium oxidizes readily in liquid and solid states. Hence, the small peak signal might be yielded from an amorphous magnesium oxide which was most likely to be one kind of nonmetallic inclusions formed during foundry processes [24]. Strikingly, the Mg concentration in the porosity and surrounding area increased significantly when the as-cast specimen was solution treated for 30 min. Moreover, this phenomenon became more obvious with solution treatment time extending. Mg enrichment around porosities during solution treatment is evident by the present finding. The intensifying of Mg content near the porosity will result in a reduction of effective solutes used for
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Fig. 4. (a) TEM micrograph recorded in [001]Al zone axis of typical Mg–Si-containing strengthening precipitates (b00 precursors) obtained from the specimen solution treated for 30 min and aged for 8 h at 180 1C. (b) High-resolution TEM micrograph recorded in [001]Al zone axis of the cross-section of a needle-like precipitate in Fig. 4a, the unit-cell was overlaid in the image.
Fig. 5. (a) Distribution of porosity in the as-cast sample, (b) a partial enlarged drawing of a typical shrinkage porosity in (a), which showed the interface region between the porosity and a-Al matrix, (c) EDS spot analysis results detected from a point, indicated by red cross symbol in the porosity of (b), (d) EDS line profiles, along the designated red lines, of Mg concentration detected from a specimen solution treated with different time. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article).
forming Mg–Si-containing strengthening precipitates in the a-Al matrix during the following aging treatment. 3.4. Existence of MgO nanoscale particles in the solution treated A356 alloy In order to exclude the impact of surface oxidation on the conclusion, TEM and EDS were used for further clarifying the phenomenon mentioned above. Fig. 6 shows the TEM micrograph of a specimen solution treated at 550 1C for 10 h. It is surprising
that flakes in Fig. 6 were frequently observed on the edge of the penetrated hole in the TEM thin-foil. From the ring-like selected area diffraction (SAED) pattern (inset in Fig. 6), it is inferred that plenty of nanoscale particles existed in the specimen. Fig. 7 presents the magnified SAED pattern and the corresponding EDS spectra obtained from a certain area in Fig. 6. The several sharp concentric rings observed in this pattern suggest polycrystalline nature of this region. The d-spacing determined from the ring diameter exhibits in the lower right corner of Fig. 7. Analysis of the diffraction pattern and interplanar spacing reveal that the
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nanoscale particles correspond to a cubic system with interference indices (200) and (220) as the brightest rings. The calculation is in good agreement with the data of MgO PDF card. The estimation strongly recommends that these nanoscale particles are MgO crystals. This conclusion is also confirmed by the EDS result which shows Mg and O elements account for a quite large proportion of the content. So the flake in Fig. 6 may be oxide bifilm (MgO and Al2O3) formed around the porosity during solution treatment. This kind of flake containing MgO nanoscale particles was also abundant in the specimen solution treated at 550 1C for 2 h. A few MgO particles were also found in the specimen solution treated at 550 1C for 30 min, while none was observed in the as-cast specimen. On the basis of these facts, we can arrive at the conclusion that these observed MgO nanoparticles are highly correlated with migration of Mg toward the porosity during solution treatment. The formation of MgO nanoparticles will be discussed below.
Fig. 6. TEM micrograph of a region in the specimen solution treated at 550 1C for 10 h, the set is the corresponding SAED pattern.
4. Discussion In order to explain the phenomenon discovered in our study, a schematic diagram (Fig. 8) is proposed to illustrate the mechanism of Mg loss in A356 alloys during solution treatment. Once the as-cast specimen is solution treated at a relatively high temperature, magnesium and a certain amount of silicon would dissolve in the a-Al matrix to produce a homogeneous solid solution within a short time. It was shown that multiple precipitating Si nanotwins could form on the eutectic Si particles in Al–Si alloys upon annealing [25]. We surmise that these Si particles are surrounded by a-Al matrix with a loose interface, which can assist the rapid diffusion of dissolved atoms. As shown in the middle of Fig. 8, the gray shaded oval particles represent eutectic Si particles. We use dotted lines to represent the interfaces between eutectic Si particles and the a-Al matrix. According to Kerte´sz et al., a prolonged high-temperature heat treatment of Al– Mg–Si alloys produced Mg enrichment on the surface [26]. As Mg atoms are outdiffusing from the Al matrix during solution treatment, the amount of hardening precipitates formed during aging treatment will decrease and thus lead to a decline of the aged hardness. However, the part of Mg enriched on the surface is limited and cannot explain the huge difference between the age hardening curves listed in Fig. 3. It can also be seen from Fig. 5d that the Mg content variation at the matrix near the porosity was much smaller than that in porosity. Two isolated dark irregular polygons in the oval at the left of Fig. 8 represent shrinkage porosity (the irregular one formed corresponding to the shape of the interdendritic region) and gas porosity (the smoother one that is approximately rounded), respectively. During solidification of cast A356 alloys, the oxidation of molten aluminum occurred and Al2O3 film formed at the internal surface of the micro-void [16]. Owing to the fact that the formation enthalpy of MgO is far lower than Al2O3 [27], the chemical affinity of Mg with the oxygen in the oxide film at the internal surface of the micro-void may be the driving force for Mg diffusion to the micro-voids, especially when the gas content inside a micro-void in A356 alloy is restricted from outer atmosphere. The thermal analysis performed by differential scanning calorimetry proved that even the stable b (Mg2Si) phase in a-Al was not stable above 520 1C [7], so all the Mg–Si phases would
Fig. 7. Analysis of polycrystalline electron diffraction pattern and the EDS spectra achieved from a part in Fig. 6.
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Fig. 8. Schematic illustration of the Mg loss process during solution treatment.
diffuse into the aluminum matrix during solution treatment. But the MgO formed near the porosities should be very stable, as the melting temperature of MgO is about 2800 1C, far exceeding the solution temperature. For A356 alloy, the magnesium depletion during solution treatment can get worse as a result of the ubiquity of porosity in the casting alloy. The thermally activated Mg atoms will diffuse through the loose interfaces mentioned above toward the porosities. Though the mechanism for the formation of MgO nanocrystalline remains unclear and needs further research, it is yet a definite evidence of Mg enrichment inside the A356 alloys during solution treatment. Because of shortened diffusion distance, the Mg enrichment in porosity is easier than on the sample surface. Therefore, Mg enrichment in porosity during solution treatment is the main reason of magnesium depletion of a-Al matrix for A356 cast alloy with a lot of porosities. Mg concentration in porosity would significantly increase during solution treatment, as described in Sections 3.3 and 3.4. Mg enrichment in porosity will result in Mg loss in Al matrix and influence the precipitation kinetics in the following aging treatment. Consequently, the alloy may lose its expected hardness seriously if the alloy with a lot of porosities was solution treated for a long time. In view of the evil influence mentioned above, more attention should be paid to the control of the quantity and size of porosities to enhance the mechanical property of A356 alloy. Otherwise, solution treatment time should be restricted in some range to reduce the loss of Mg under the premise of guaranteeing the requirements of ductility. Based on the discussion above, it is reasonable for us to draw the conclusion that a specimen containing more porosity will obtain lower hardness for lower Mg concentration used for forming Mg–Si-containing strengthening precipitates.
5. Conclusions The present study has demonstrated the microstructure evolution and the age hardening characteristic of an A356 alloy in relation with solution treatment. From the obtained results, the following can be concluded. (1) Apart from the spheroidization of the Si particles during solution treatment, thermally activated Mg atoms can diffuse out from the Al-matrix and enrich near the porosity preexisted in the alloy. (2) A prolonged solution treatment has a negative influence on the age hardening due to the existence of porosity in the cast A356 alloy. The monoclinic b00 precipitates play a crucial role in the age hardening response of the alloy. (3) The diffusion of Mg atoms toward porosities during solution treatment may lead to a significant reduction of the Mg solutes available to form Mg–Si-containing strengthening
precipitates in the Al-matrix upon aging treatment. The Mg enrichment around porosity becomes more severe with the increase of solution treatment time. (4) Nanoscale particles of magnesium oxides were found in the specimen solution treated for an extended time, while none was observed in the as-cast specimen. The formation of MgO nanoparticles is correlated with the Mg enrichment in porosity occurred during solution treatment. (5) This study indicates that controlling the porosity formed in the casting process is crucial for enhancing the mechanical properties of the A356 alloy. For Al–Si–Mg alloys with abundant porosities, the solution treatment time should be limited to an optimum to reach a balance between strength and ductility of the alloy.
Acknowledgments This work is financially supported by the National Basic Research (973) Program of China (No. 2009CB623704); the National Natural Science Foundation of China (No. 51171063); Instrumental Innovation Foundation of Hunan Province (No. 2011TT1003); the Aid Program for Science and Technology Innovative Research Team in Higher Educational Institutions of Hunan Province, Hunan Provincial Innovation Foundation For Postgraduate and Gatan China Scholarship.
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