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Evaluating the mechanical properties of a thermomechanically processed unmodified A356 Al alloy employing shear punch testing method
Materials and Design 43 (2013) 419–425
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Evaluating the mechanical properties of a thermomechanically processed unmodified A356 Al alloy employing shear punch testing method N. Haghdadi, A. Zarei-Hanzaki ⇑, Ali A. Roostaei, A.R. Hemmati School of Metallurgy and Materials Engineering, College of Engineering, University of Tehran, Tehran, Iran
a r t i c l e
i n f o
Article history: Received 26 May 2012 Accepted 1 July 2012 Available online 7 July 2012 Keywords: Aluminum alloy Thermomechanical processing Shear punch Mechanical properties Quality index
a b s t r a c t The effects of thermomechanical processing (TMP) parameters on the microstructure evolution and final mechanical properties of an unmodified A356 Al alloy have been investigated. The evaluation of mechanical properties was carried out using shear punch testing (SPT) method. The applied TMP cycles encompassed a set of isothermal hot compression tests in the temperature range of 420–540 °C under various strain rates of 0.001, 0.01 and 0.1 s1. The results indicate that increasing both the TMP temperature and strain rate has enhanced the room temperature strength of the experimental alloy. Furthermore, while ductility follows a similar improving trend with increasing the TMP temperature, it is deteriorated by increasing the strain rate. The obtained results are reasoned debating the role of solute atoms and the changes in Si particles shape and size. The present work also made use of quality index (Q) to measure the mechanical performance of the TMPed unmodified A356 Al alloy. Ó 2012 Elsevier Ltd. All rights reserved.
1. Introduction In industry, aerospace and automotive in particular, there is an ongoing quest to reduce the weight and cost of finished parts. In this respect, of importance is exploiting the materials with unique features such as low density and recyclability [1]. This implies that the substitution of aluminum alloys for steels and magnesium alloys is a momentous aim considering the high weight of steel parts and the recycling difficulties of magnesium products [2]. The demand of the aforementioned industries to produce costefficient integral components of complex geometry has led to an increased utilization of Al–Si cast alloys, which possess excellent castability, weldability and pressure tightness, together with high wear and corrosion resistance [3]. Nevertheless, the Al–Si alloys in as-cast condition suffer from an inadequate level of mechanical performance. As is widely reported, the mechanical properties of the Al–Si cast alloys are greatly affected by the size, distribution and morphology of the eutectic Si particles [4]. In the as-cast condition, the microstructure of the alloy contains brittle acicular Si particles with sharp edges. The presence of such coarse plate-like particles would inevitably degrade the mechanical properties of the alloy, particularly the ductility [5]. This may limit the practical use of Al–Si cast alloys to non-critical applications where inadequate performance would not result in catastrophic failure [6,7].
⇑ Corresponding author. Tel.: +98 21 61114167; fax: +98 21 88006076. E-mail address: [email protected] (A. Zarei-Hanzaki). 0261-3069/$ - see front matter Ó 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.matdes.2012.07.001
Accordingly, there has been a continuous effort to enhance the mechanical properties of Al–Si cast alloys by controlling the Si particles size and morphology. The latter has been attained through different approaches such as chemical modification [8,9], quench modification [10] and long-time heat treatment [7,11]. The chemical modification was performed through the addition of an elemental modifier (e.g., Sr, Na and Sb) to the alloy melt, as a result of which the Si morphology in the microstructure changed from acicular plate-like to fibrous form [7–9]. As an example for the application of this approach, sodium fluoride salts associated with T4 treatment have been widely used in the foundry in order to modify the eutectic Si-rich particles [12,13]. However, the chemical modification was accompanied by some shortcomings such as increased porosity and environmental safety concerns [14]. In the quench modification technique, applying high cooling rates shifted Al–Si eutectic point to the lower temperatures, and thus led to the formation of finer Si particles [10]. The downside associated with this technique was the high thermal stresses induced during quenching [15]. Silicon spheroidizing treatment (SST) was another method used by a number of researchers to improve the mechanical properties of Al–Si cast alloys by holding them at high temperatures [7,11]. In this method, the driving force for the spheroidization of the Si particles was provided by a reduction in the surface energy [7]. It was, however, reported that the cast unmodified Al–Si alloys are resistant to the SST [16]. This was related to the lack of interfacial instabilities and shape perturbations in such alloys, which were essential for the initial fragmentation of the Si particles [7]. In addition, according to the literature [17,18], the rate of Si particle coarsening in the microstructure of
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unmodified alloys is about 2.5 times higher than that of the modified ones. This makes the precise control of heat treatment process too difficult to achieve the optimum morphology of the Si particles. Taking the disadvantages of the above-mentioned approaches into consideration, thermomechanical processing (TMP) has been put into practice by researchers as an alternative efficient method to control the morphology of the Si particles in the microstructure of the Al–Si alloys [19–23]. In this context, Zarei-Hanzaki et al. [20] studied the room temperature mechanical characteristics of a Srmodified thixocast A356 Al alloy after hot compression. They showed that the simultaneous presence of deformation and heat influences the morphology of the Si particles, thereby resulting in higher strength without sacrificing the ductility. In other investigations by Ding et al. [22] and Mallapur et al. [23], similar enhancements in the mechanical properties of an Al–12Si–0.2Mg alloy and an A356 Al alloy after hot extrusion and hot forging processes were reported, respectively. However, as regards the possibility of enhancing the room temperature mechanical properties of Al–Si alloys through TMP, there is a lack of systematic research conducted on the unmodified alloys. That being the case, the present study aims to investigate the effects of TMP parameters, i.e., strain rate and temperature, on the microstructure evolution and final mechanical properties of unmodified Al–7Si–0.4Mg alloy. 2. Experimental procedure The experimental material was an unmodified A356 cast Al alloy with the chemical composition given in Table 1. In order to apply the predetermined thermomechanical processing (TMP) cycles, hot compression tests were employed. To this end, cylindrical hot compression specimens, 7 mm in diameter and 11 mm in height, were machined from the cast ingot. The specimens were then subjected to compressive deformation at the temperatures of 420, 460, 500 and 540 °C under the strain rates of 0.001, 0.01 and 0.1 s1 (in accordance with ASTM E209 [24]). This was carried out using an Instron 4208 universal testing machine equipped with a programmable resistance furnace and a tool geometry that allows for the quenching of the specimens thereupon. Prior to the tests, the specimens were held at the deformation temperature for 5 min to ensure a homogeneous temperature distribution. A very thin mica plate was used to reduce the friction and prevent the adhesion of the specimens to the anvils. At the end of straining, the specimens were rapidly quenched in warm water (60 °C to minimize quenching stresses and specimen distortion [25]). Scanning electron microscopy (SEM, CamScan, MV2300, England) was used to examine the microstructures of as-cast and TMPed specimens. Preparatory to that, the specimens were mechanically polished and etched using Keller solution. Mechanical properties were evaluated using a shear punch testing (SPT) rig with a punch of 3 ± 0.005 mm diameter and receivinghole of 3.04 ± 0.005 mm diameter. The SPT specimens were wirecut into 600 lm-thick slices. With the purpose of removing oxides and surface flaws, the slices surfaces were mechanically polished to the final thickness of 300 ± 15 lm. On each slice, the SPT was repeated a minimum of three times. The tests were carried out at the room temperature with the constant crosshead speed of 0.2 mm/ min. The load–displacement data were converted to shear stress– displacement data using the following equation. The related data
acquisition and result analyses have been thoroughly explained elsewhere [20,26–29].
s¼
PF ¼ Cr 2prt
where P is the applied load, F is the friction load, t is the specimen thickness, r is the average of punch and receiving-hole radii and C is the correlation coefficient (C = 0.67 for Al–Mg–Si alloys [20]). Hardness measurements were also performed using a Brinell hardness tester under the load of 10 kg. The average of at least five hardness measurements for each specimen was reported as its hardness value. 3. Results and discussion 3.1. Microstructure evolution The SEM micrograph of A356 Al alloy in as-cast condition is depicted in Fig. 1. It is evident that the microstructure of the experimental alloy consists of coarse acicular silicon particles with a mean diameter of 10.3 lm embedded in a dendritic aluminum matrix. The mechanical properties of the cast A356 Al alloy are listed in Table 2. As is seen, the alloy possesses insufficient strength and very poor ductility in the as-cast condition. Practically, the coarse plate-like nature of the Si particles leads to premature crack initiation, thereby degrading the strength and ductility of the alloy [7]. Fig. 2 shows the microstructures of specimens after hot compression tests at the temperature of 540 °C under different strain rates. As is observed, the Si particles are refined and evenly redistributed in the aluminum matrix. The average diameters of the Si particles in the microstructures of the specimens TMPed under strain rates of 0.001, 0.01 and 0.1 s1 are measured to be 4.1, 3.6 and 3.4 lm, respectively, which shows a declining trend with increasing the strain rate. It is also to be noted that at lower strain rates, a higher degree of spheroidization has occurred. Furthermore, at the constant strain rate of 0.01 s1, the microstructure evolution with TMP temperature is depicted in Fig. 3. As is seen, the degree of Si spheroidization as well as the average diameter of the Si particles increases with increasing the TMP temperature. The average diameters of the Si particles in the microstructures of the specimens TMPed at 420 and 500 °C (under the strain rate of 0.01 s1) are measured to be 2.4 and 3.3 lm, respectively. The observed changes in the Si particles morphology due to the
Table 1 The chemical composition of the experimental A356 Al alloy. Element
Si
Mg
Fe
Cu
Mn
Al
wt.%
7.2
0.4
0.24
0.03
0.01
Bal.
ð1Þ
Fig. 1. The SEM micrograph of cast A356 Al alloy.
N. Haghdadi et al. / Materials and Design 43 (2013) 419–425 Table 2 The mechanical properties of Al–7Si–0.4Mg alloy in as-cast condition. Hardness (BHN)
UTS (MPa)
Elongation (%)
59
174
4
temperature and strain rate variations can be justified in terms of the mechanism by which the Si particles are spheroidized.
3.1.1. Si spheroidization As is well established, during holding the A356 Al alloy specimens at high temperatures, the Si particles undergo fragmentation, spheroidization and ultimately coarsening, as the holding time increases [7]. In the absence of any external loading, the thermal disintegration of Si plates, which acts as the governing factor in the break-up process, begins at local crystal defects (Rayleigh’s criteria) [7]. These defects are the common morphological faults such as terminations, kinks and striations as well as holes and fissures in the plates [7]. In addition, the strains induced by coefficient of thermal expansion (CTE) mismatch between the silicon and aluminum phases during heat-up can lead to the brittle fracture of the eutectic Si particles [30]. The effects of the imposed strain from external loading on the fragmentation step may be best explained by considering Al–Si alloys as in situ composites. Under the external loading, strain
421
incompatibility develops in the interface of the Si particles and aluminum matrix, because of their different elastic constants and deformation behavior [31]. This, in turn, would lead to the cracking of the Si particles even at low strains. The plastic strain also generates a high density of lattice defects, which can act as favorite sites for crack initiation and the subsequent fracture of the Si particles. The spheroidization and coarsening of the eutectic silicon particles are driven by a decrease in interfacial energy and a decrease in elastic strain energy [32]. The capillary forces due to the curvature difference between the edge and the flat interface trigger mass transport. As is well known, the spheroidization is a diffusional process controlled by the migration of Si atoms through aluminum matrix [7]. Similar to any other diffusional processes, the spheroidizing is a function of time and temperature. Therefore, on the one hand, this is the main reason for the observed increase in the spheroidization degree of the Si particles with decreasing the strain rate (increasing the process time) (see Fig. 2). On the other hand, the higher degrees of spheroidization at higher temperatures, which can also be deduced from microstructures in Fig. 3, seem rational. Nevertheless, applying higher temperatures and lower strain rates may lead to the coarsening of the Si particles [19]. The coarsening may be due to the supply of Si atoms diffusing out of the cores of dendrites and/or through Ostwald ripening mechanism in which larger particles grow at the expense of the smaller ones [7,33].
Fig. 2. The SEM microstructures of hot compressed specimens at 540 °C, under strain rates of (a) 0.001, (b) 0.01 and (c) 0.1 s1. The arrows show some of the Si particles present in the microstructures.
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Fig. 3. The SEM microstructures of hot compressed specimens under the strain rate of 0.01 s1, at the temperatures of (a) and (b) 420 °C, (c) and (d) 500 °C. The arrows show some of the Si particles distributed in the microstructure.
3.2. Mechanical properties As is well documented, the mechanical properties of A356 Al alloys are strongly dependent on the a (Al)-phase characteristics as well as the size, morphology and distribution of the Si particles [7]. The influences of the thermomechanical parameters (i.e., strain rate and temperature) on the mechanical properties of TMPed A356 Al alloy are elaborately accounted for in the following sections. It is to be mentioned that the UTS and the elongation values are calculated from their shear counterparts (ultimate shear strength and shear elongation, respectively) applying the correlation coefficient of 0.67, as was explained earlier in this paper. A typical shear punch test curve obtained for the specimen TMPed at 460 °C under the strain rate of 0.1 s1 is depicted in Fig. 4a. 3.2.1. Strength The variations of ultimate tensile strength (UTS) with TMP temperature under different strain rates are plotted in Fig. 4b. For comparison, the as-cast related value is also included in the chart. In general, regardless of the applied TMP conditions, the TMPed specimens possess significantly higher UTS values than that of the ascast one. A closer observation of Fig. 4b, however, reveals that the specimens processed at higher temperatures and strain rates exhibit the higher UTS values. During TMP at high temperatures, the cast material experiences homogenization. The microstructure homogeneity increases as the
TMP temperature rises. This, in turn, would end to eliminating the segregation of alloying elements in the microstructure [34]. The segregation of solute elements resulting from dendritic solidification is reported to adversely affect the mechanical properties of the experimental material [34]. Furthermore, at higher temperatures, higher amounts of Si and Mg enter into solid solution [34] and result in the increased natural age-hardenability of the alloy. As is well known, the solute atoms may also serve as barriers to the movement of dislocations and thereby rendering appreciable influence on the strength properties of the alloy. In addition, upon quenching of the specimens, which were TMPed at higher temperatures, more amounts of fine Si precipitates deposit at the Si–Al interface, and thus improve the Si–Al interfacial bonding strength [35]. Meanwhile, the effect of strain rate on the UTS may be rationalized considering the role of tangled dislocation structures in hindering the dislocation movement, which is more prominent under higher strain rates [36]. Apart from this, the lack of dislocation annihilation during deformation with high strain rate contributes to the formation of ultrafine substructures thereby increasing the alloy strength. It is also reported that the rate of damage accumulation in the Al–Si alloys is a strong function of Si particles shape and size [5]. Coarse acicular Si particles with high shape factor are considered as preferred sites for premature crack initiation during deformation. This is related to the high levels of stress concentration
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(a) 240
around the lengthy Si particles [37]. In contrast, the Si particles with spherical shape exhibit a higher resistance to debonding under the imposed straining [38]. As the Si spheroidization is more pronounced at higher TMP temperatures [39], the observed enhancement in the UTS values with increasing the TMP temperature seems reasonable. Moreover, concerning the effect of strain rate on the strength, the size of the Si particles plays a vital role. In this regard, Gurland and Plateau [40] proposed the following equation to estimate the required critical stress for developing a crack in a particle of size D:
Shear Stress (MPa)
200 160 120 80 40
"
0 0
0.05
0.1
0.15
0.2
0.25
0.3
Displacement (mm)
(b)
0.1(1/s)
0.01(1/s)
UTS (MPa)
400 300 200 100 As-cast
420
460
500
540
Temperature (oC) Fig. 4. (a) A typical shear strength–displacement curve obtained by shear punch testing of examined alloy after TMP at 460 °C under the strain rate of 0.1 s1, and (b) The variations of UTS with different TMP conditions along with the as-cast value.
140 0.1(1/s)
0.01(1/s)
0.001(1/s)
HB (N.mm-2)
125 110 95
ð2Þ
Dq2
1
r ¼ r0 þ kk2
ð3Þ
where r is the strength, k and r0 are constants and k is the mean particle interspacing. In addition to all the above-mentioned factors, the finer grain size of the microstructure due to the breaking of Al dendrites could result in enhanced mechanical properties through Hall–Petch strengthening mechanism [41,42]. It is also reported that dendrite fineness associated with a reduction in dendrite spacing could have a positive influence on the mechanical characteristics of cast structures [43–45]. The hardness variations with TMP parameters are represented in Fig. 5. It is obvious that the hardness values change in the same way as the UTS results (compare with Fig. 4). The same interpretations stated to explain the UTS variations remain valid for the hardness evolution, as well.
80 65 50
420
460
500
540
o
Temperature ( C) Fig. 5. Brinell hardness values for the experimental alloy after TMP under different conditions.
25 0.1(1/s)
0.01(1/s)
0.001(1/s)
20
Elongation (%)
#12
where q is the stress concentration factor, E is the weighted average of the elastic moduli of the particle and matrix and c is the interfacial energy of the crack. Taking the Eq. (2) into consideration and the fact that at a given temperature, the size of the Si particles becomes smaller with increasing the strain rate [39], the higher UTS value at higher strain rate would be rationalized. From another point of view, having smaller Si particles is associated with a decrease in particle interspacing, which, in turn, leads to mechanical strengthening through Orowan dislocation bowing mechanism:
0.001(1/s)
500
0
r¼
Ec
15 10 5 0
As-cast
420
460
500
540
3.2.2. Ductility Fig. 6 illustrates the effect of TMP parameters on the tensile elongation of the experimental alloy. For comparison, the elongation value of the cast specimen is also included in the chart. A dramatic increase in the elongation value of the cast specimen is observed upon TMP under any conditions. However, according to Fig. 6, the higher ductility values are obtained at higher temperatures but lower strain rates. The acceleration of homogenization and the dissolution/morphological changes of brittle intermetallic compounds with increasing the temperature are believed to be reasons for the observed behavior [46,47]. Moreover, as is well accepted, the A356 Al alloys with rounder Si particles exhibit higher ductility values due to the slow accumulation of damage during deformation [37]. In this regard, the microstructural studies in the present work showed a higher degree of spheroidization at higher temperatures and lower strain rates. Fig. 7, accordingly, depicts the variations of Si particle roundness under different TMP conditions where roundness is calculated using the following equation:
Particle roundness ¼
4 particle Area 100 Particle perimeter
ð4Þ
o
Temperature ( C) Fig. 6. Elongation to failure for the cast and TMPed A356 Al alloy under various conditions.
It is apparent from Figs. 6 and 7 that the elongation and roundness variations with TMP temperature and strain rate follow a similar trend. This gives support to the idea of attributing the higher
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100 0.01(1/s)
0.1(1/s)
0.001(1/s)
Quality Index (MPa)
Si Particle Roundness
0.1(1/s)
80 60 40 20 0
0.01(1/s)
0.001(1/s)
600 500 400 300 200
420
460
500
540
As-cast
420
Fig. 7. The variations of Si particle roundness with different TMP routes.
ductility values of the A356 Al alloy at higher temperatures and lower strain rates to the higher spheroidization degree of the Si particles. Furthermore, at higher temperatures and lower strain rates, local stresses around the eutectic Si particles subside due to the ease of the dislocation movement. Generally, due to plastic incompatibility and differences in the coefficient of thermal expansion between the Si particles and aluminum matrix, local stresses build up at the particle/matrix interface, which is dictated by an increase in the density of dislocations therein. Moreover, from the metallurgical point of view, it is well known that the Al-rich (a) phase and fibrous or needle-like Si particles have dissimilar growth behavior with Si particles growing from the liquid in a faceted manner (smooth growth interface), while the a-Al phase solidifies with surfaces that are rough. Due to this, some localized deformations are expected to occur in their boundary, which is not perfectly conformed [12,13,44,45]. This seems to be another reason for the local stress build-up at the interface. In view of their thermally activated nature, dislocation annihilation, climbing and cross slip are more encouraged at higher temperatures. This, in turn, gives rise to stress relaxation and the reduction of local stress incompatibilities at the matrix/particle interface, which ultimately render a positive influence on the ductility of the experimental alloy [48]. Besides contributing to the higher degree of dynamic recovery, the effect of lower strain rate at a given temperature may be realized considering the size of the Si particles. As was previously discussed, under lower strain rates, the lesser degree of particle fragmentation and the longer available time for the diffusion of Si atoms lead to the greater average size of the Si particles [39]. As a result, the interspacing between the Si particles increases. Hence, a larger volume of the aluminum matrix would be available between particles to blunt the interface-originated cracks, and thereby improving the elongation [49]. 3.2.3. Quality index Towards a comprehensive view of the mechanical performance of the Al–Si casting alloys, the empirical parameter of quality index, Q, has been employed by many researchers to date [50]. For instance, Ceschini et al. [51] took advantage of the Q-parameter in order to draw a correlation between the tensile properties and solidification microstructure of an A357 Al alloy. In addition, the quality index was used by Shabestari and Shahri [52] to investigate the optimum modification and solidification conditions of an A356 Al alloy, which gives the best tensile properties. In similar fashion, this study utilizes the Q-parameter to evaluate the relationship between the TMP parameters and the mechanical quality of the experimental A356 Al alloy. The quality index involves a linear combination of two tensile parameters, i.e., UTS and elongation, each of which becomes weighed by different coefficients of C and K, respectively:
460
500
540
o
Temperature ( C)
o
Temperature ( C)
Fig. 8. Quality index values for the cast and TMPed A356 Al alloy under various conditions.
Q ¼ C UTS þ K log ð%EÞ
ð5Þ
Drouzy et al. [53] described the quality index for Al–Si alloys by the following expression in which fixed numbers replace the weight coefficients:
Q ¼ UTS ðMPaÞ þ 150 log ð%EÞ
ð6Þ
Fig. 8 shows the calculated Q-values of the cast and TMPed specimens under various conditions. As is seen, the highest Q-values are obtained for the specimens processed at higher temperatures and strain rates. The observed trend is in line with the UTS variations in Fig. 4. It is also interesting to note from Fig. 8 that the highest Q-value (corresponding to the specimen processed at 540 °C under the strain rate of 0.1 s1) is calculated to be about 630 MPa. This value is much higher than the highest values reported in the literature for the chemically modified and heat-treated Al–Si alloys [51,52].
4. Conclusions In the present study, the room temperature mechanical properties and microstructure evolution of a thermomechanically processed unmodified A356 Al alloy have been investigated. Based on this research, the following outcomes are reported: (1) With the aid of TMP, the spheroidization of the Si particles was occurred at relatively low temperatures and durations. In consequence, the TMPed A356 Al alloy acquired a higher strength without sacrificing its ductility. (2) Higher UTS values were obtained upon TMP at higher temperatures and strain rates. The more homogeneous matrix microstructure containing rounder Si particles at high temperatures was considered as one good reason for the observed behavior. In addition, increasing the strain rate results in having smaller Si particles which, in turn, leads to mechanical strengthening through Orowan dislocation bowing mechanism. (3) Brinell hardness variations for the TMPed specimens followed the same trend as the UTS variations. (4) Increasing the TMP temperature and/or decreasing the strain rate led to higher ductility values, mainly due to the spheroidization of the Si particles together with the restoration of the deformed microstructure. (5) The overall mechanical performance of the TMPed unmodified A356 Al alloy was assessed using quality index. The highest Qvalues were obtained for the specimens processed at higher temperatures and strain rates.
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Evaluating the mechanical properties of a thermomechanically processed unmodified A356 Al alloy employing shear punch testing method N. Haghdadi, A. Zarei-Hanzaki ⇑, Ali A. Roostaei, A.R. Hemmati School of Metallurgy and Materials Engineering, College of Engineering, University of Tehran, Tehran, Iran
a r t i c l e
i n f o
Article history: Received 26 May 2012 Accepted 1 July 2012 Available online 7 July 2012 Keywords: Aluminum alloy Thermomechanical processing Shear punch Mechanical properties Quality index
a b s t r a c t The effects of thermomechanical processing (TMP) parameters on the microstructure evolution and final mechanical properties of an unmodified A356 Al alloy have been investigated. The evaluation of mechanical properties was carried out using shear punch testing (SPT) method. The applied TMP cycles encompassed a set of isothermal hot compression tests in the temperature range of 420–540 °C under various strain rates of 0.001, 0.01 and 0.1 s1. The results indicate that increasing both the TMP temperature and strain rate has enhanced the room temperature strength of the experimental alloy. Furthermore, while ductility follows a similar improving trend with increasing the TMP temperature, it is deteriorated by increasing the strain rate. The obtained results are reasoned debating the role of solute atoms and the changes in Si particles shape and size. The present work also made use of quality index (Q) to measure the mechanical performance of the TMPed unmodified A356 Al alloy. Ó 2012 Elsevier Ltd. All rights reserved.
1. Introduction In industry, aerospace and automotive in particular, there is an ongoing quest to reduce the weight and cost of finished parts. In this respect, of importance is exploiting the materials with unique features such as low density and recyclability [1]. This implies that the substitution of aluminum alloys for steels and magnesium alloys is a momentous aim considering the high weight of steel parts and the recycling difficulties of magnesium products [2]. The demand of the aforementioned industries to produce costefficient integral components of complex geometry has led to an increased utilization of Al–Si cast alloys, which possess excellent castability, weldability and pressure tightness, together with high wear and corrosion resistance [3]. Nevertheless, the Al–Si alloys in as-cast condition suffer from an inadequate level of mechanical performance. As is widely reported, the mechanical properties of the Al–Si cast alloys are greatly affected by the size, distribution and morphology of the eutectic Si particles [4]. In the as-cast condition, the microstructure of the alloy contains brittle acicular Si particles with sharp edges. The presence of such coarse plate-like particles would inevitably degrade the mechanical properties of the alloy, particularly the ductility [5]. This may limit the practical use of Al–Si cast alloys to non-critical applications where inadequate performance would not result in catastrophic failure [6,7].
⇑ Corresponding author. Tel.: +98 21 61114167; fax: +98 21 88006076. E-mail address: [email protected] (A. Zarei-Hanzaki). 0261-3069/$ - see front matter Ó 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.matdes.2012.07.001
Accordingly, there has been a continuous effort to enhance the mechanical properties of Al–Si cast alloys by controlling the Si particles size and morphology. The latter has been attained through different approaches such as chemical modification [8,9], quench modification [10] and long-time heat treatment [7,11]. The chemical modification was performed through the addition of an elemental modifier (e.g., Sr, Na and Sb) to the alloy melt, as a result of which the Si morphology in the microstructure changed from acicular plate-like to fibrous form [7–9]. As an example for the application of this approach, sodium fluoride salts associated with T4 treatment have been widely used in the foundry in order to modify the eutectic Si-rich particles [12,13]. However, the chemical modification was accompanied by some shortcomings such as increased porosity and environmental safety concerns [14]. In the quench modification technique, applying high cooling rates shifted Al–Si eutectic point to the lower temperatures, and thus led to the formation of finer Si particles [10]. The downside associated with this technique was the high thermal stresses induced during quenching [15]. Silicon spheroidizing treatment (SST) was another method used by a number of researchers to improve the mechanical properties of Al–Si cast alloys by holding them at high temperatures [7,11]. In this method, the driving force for the spheroidization of the Si particles was provided by a reduction in the surface energy [7]. It was, however, reported that the cast unmodified Al–Si alloys are resistant to the SST [16]. This was related to the lack of interfacial instabilities and shape perturbations in such alloys, which were essential for the initial fragmentation of the Si particles [7]. In addition, according to the literature [17,18], the rate of Si particle coarsening in the microstructure of
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unmodified alloys is about 2.5 times higher than that of the modified ones. This makes the precise control of heat treatment process too difficult to achieve the optimum morphology of the Si particles. Taking the disadvantages of the above-mentioned approaches into consideration, thermomechanical processing (TMP) has been put into practice by researchers as an alternative efficient method to control the morphology of the Si particles in the microstructure of the Al–Si alloys [19–23]. In this context, Zarei-Hanzaki et al. [20] studied the room temperature mechanical characteristics of a Srmodified thixocast A356 Al alloy after hot compression. They showed that the simultaneous presence of deformation and heat influences the morphology of the Si particles, thereby resulting in higher strength without sacrificing the ductility. In other investigations by Ding et al. [22] and Mallapur et al. [23], similar enhancements in the mechanical properties of an Al–12Si–0.2Mg alloy and an A356 Al alloy after hot extrusion and hot forging processes were reported, respectively. However, as regards the possibility of enhancing the room temperature mechanical properties of Al–Si alloys through TMP, there is a lack of systematic research conducted on the unmodified alloys. That being the case, the present study aims to investigate the effects of TMP parameters, i.e., strain rate and temperature, on the microstructure evolution and final mechanical properties of unmodified Al–7Si–0.4Mg alloy. 2. Experimental procedure The experimental material was an unmodified A356 cast Al alloy with the chemical composition given in Table 1. In order to apply the predetermined thermomechanical processing (TMP) cycles, hot compression tests were employed. To this end, cylindrical hot compression specimens, 7 mm in diameter and 11 mm in height, were machined from the cast ingot. The specimens were then subjected to compressive deformation at the temperatures of 420, 460, 500 and 540 °C under the strain rates of 0.001, 0.01 and 0.1 s1 (in accordance with ASTM E209 [24]). This was carried out using an Instron 4208 universal testing machine equipped with a programmable resistance furnace and a tool geometry that allows for the quenching of the specimens thereupon. Prior to the tests, the specimens were held at the deformation temperature for 5 min to ensure a homogeneous temperature distribution. A very thin mica plate was used to reduce the friction and prevent the adhesion of the specimens to the anvils. At the end of straining, the specimens were rapidly quenched in warm water (60 °C to minimize quenching stresses and specimen distortion [25]). Scanning electron microscopy (SEM, CamScan, MV2300, England) was used to examine the microstructures of as-cast and TMPed specimens. Preparatory to that, the specimens were mechanically polished and etched using Keller solution. Mechanical properties were evaluated using a shear punch testing (SPT) rig with a punch of 3 ± 0.005 mm diameter and receivinghole of 3.04 ± 0.005 mm diameter. The SPT specimens were wirecut into 600 lm-thick slices. With the purpose of removing oxides and surface flaws, the slices surfaces were mechanically polished to the final thickness of 300 ± 15 lm. On each slice, the SPT was repeated a minimum of three times. The tests were carried out at the room temperature with the constant crosshead speed of 0.2 mm/ min. The load–displacement data were converted to shear stress– displacement data using the following equation. The related data
acquisition and result analyses have been thoroughly explained elsewhere [20,26–29].
s¼
PF ¼ Cr 2prt
where P is the applied load, F is the friction load, t is the specimen thickness, r is the average of punch and receiving-hole radii and C is the correlation coefficient (C = 0.67 for Al–Mg–Si alloys [20]). Hardness measurements were also performed using a Brinell hardness tester under the load of 10 kg. The average of at least five hardness measurements for each specimen was reported as its hardness value. 3. Results and discussion 3.1. Microstructure evolution The SEM micrograph of A356 Al alloy in as-cast condition is depicted in Fig. 1. It is evident that the microstructure of the experimental alloy consists of coarse acicular silicon particles with a mean diameter of 10.3 lm embedded in a dendritic aluminum matrix. The mechanical properties of the cast A356 Al alloy are listed in Table 2. As is seen, the alloy possesses insufficient strength and very poor ductility in the as-cast condition. Practically, the coarse plate-like nature of the Si particles leads to premature crack initiation, thereby degrading the strength and ductility of the alloy [7]. Fig. 2 shows the microstructures of specimens after hot compression tests at the temperature of 540 °C under different strain rates. As is observed, the Si particles are refined and evenly redistributed in the aluminum matrix. The average diameters of the Si particles in the microstructures of the specimens TMPed under strain rates of 0.001, 0.01 and 0.1 s1 are measured to be 4.1, 3.6 and 3.4 lm, respectively, which shows a declining trend with increasing the strain rate. It is also to be noted that at lower strain rates, a higher degree of spheroidization has occurred. Furthermore, at the constant strain rate of 0.01 s1, the microstructure evolution with TMP temperature is depicted in Fig. 3. As is seen, the degree of Si spheroidization as well as the average diameter of the Si particles increases with increasing the TMP temperature. The average diameters of the Si particles in the microstructures of the specimens TMPed at 420 and 500 °C (under the strain rate of 0.01 s1) are measured to be 2.4 and 3.3 lm, respectively. The observed changes in the Si particles morphology due to the
Table 1 The chemical composition of the experimental A356 Al alloy. Element
Si
Mg
Fe
Cu
Mn
Al
wt.%
7.2
0.4
0.24
0.03
0.01
Bal.
ð1Þ
Fig. 1. The SEM micrograph of cast A356 Al alloy.
N. Haghdadi et al. / Materials and Design 43 (2013) 419–425 Table 2 The mechanical properties of Al–7Si–0.4Mg alloy in as-cast condition. Hardness (BHN)
UTS (MPa)
Elongation (%)
59
174
4
temperature and strain rate variations can be justified in terms of the mechanism by which the Si particles are spheroidized.
3.1.1. Si spheroidization As is well established, during holding the A356 Al alloy specimens at high temperatures, the Si particles undergo fragmentation, spheroidization and ultimately coarsening, as the holding time increases [7]. In the absence of any external loading, the thermal disintegration of Si plates, which acts as the governing factor in the break-up process, begins at local crystal defects (Rayleigh’s criteria) [7]. These defects are the common morphological faults such as terminations, kinks and striations as well as holes and fissures in the plates [7]. In addition, the strains induced by coefficient of thermal expansion (CTE) mismatch between the silicon and aluminum phases during heat-up can lead to the brittle fracture of the eutectic Si particles [30]. The effects of the imposed strain from external loading on the fragmentation step may be best explained by considering Al–Si alloys as in situ composites. Under the external loading, strain
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incompatibility develops in the interface of the Si particles and aluminum matrix, because of their different elastic constants and deformation behavior [31]. This, in turn, would lead to the cracking of the Si particles even at low strains. The plastic strain also generates a high density of lattice defects, which can act as favorite sites for crack initiation and the subsequent fracture of the Si particles. The spheroidization and coarsening of the eutectic silicon particles are driven by a decrease in interfacial energy and a decrease in elastic strain energy [32]. The capillary forces due to the curvature difference between the edge and the flat interface trigger mass transport. As is well known, the spheroidization is a diffusional process controlled by the migration of Si atoms through aluminum matrix [7]. Similar to any other diffusional processes, the spheroidizing is a function of time and temperature. Therefore, on the one hand, this is the main reason for the observed increase in the spheroidization degree of the Si particles with decreasing the strain rate (increasing the process time) (see Fig. 2). On the other hand, the higher degrees of spheroidization at higher temperatures, which can also be deduced from microstructures in Fig. 3, seem rational. Nevertheless, applying higher temperatures and lower strain rates may lead to the coarsening of the Si particles [19]. The coarsening may be due to the supply of Si atoms diffusing out of the cores of dendrites and/or through Ostwald ripening mechanism in which larger particles grow at the expense of the smaller ones [7,33].
Fig. 2. The SEM microstructures of hot compressed specimens at 540 °C, under strain rates of (a) 0.001, (b) 0.01 and (c) 0.1 s1. The arrows show some of the Si particles present in the microstructures.
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Fig. 3. The SEM microstructures of hot compressed specimens under the strain rate of 0.01 s1, at the temperatures of (a) and (b) 420 °C, (c) and (d) 500 °C. The arrows show some of the Si particles distributed in the microstructure.
3.2. Mechanical properties As is well documented, the mechanical properties of A356 Al alloys are strongly dependent on the a (Al)-phase characteristics as well as the size, morphology and distribution of the Si particles [7]. The influences of the thermomechanical parameters (i.e., strain rate and temperature) on the mechanical properties of TMPed A356 Al alloy are elaborately accounted for in the following sections. It is to be mentioned that the UTS and the elongation values are calculated from their shear counterparts (ultimate shear strength and shear elongation, respectively) applying the correlation coefficient of 0.67, as was explained earlier in this paper. A typical shear punch test curve obtained for the specimen TMPed at 460 °C under the strain rate of 0.1 s1 is depicted in Fig. 4a. 3.2.1. Strength The variations of ultimate tensile strength (UTS) with TMP temperature under different strain rates are plotted in Fig. 4b. For comparison, the as-cast related value is also included in the chart. In general, regardless of the applied TMP conditions, the TMPed specimens possess significantly higher UTS values than that of the ascast one. A closer observation of Fig. 4b, however, reveals that the specimens processed at higher temperatures and strain rates exhibit the higher UTS values. During TMP at high temperatures, the cast material experiences homogenization. The microstructure homogeneity increases as the
TMP temperature rises. This, in turn, would end to eliminating the segregation of alloying elements in the microstructure [34]. The segregation of solute elements resulting from dendritic solidification is reported to adversely affect the mechanical properties of the experimental material [34]. Furthermore, at higher temperatures, higher amounts of Si and Mg enter into solid solution [34] and result in the increased natural age-hardenability of the alloy. As is well known, the solute atoms may also serve as barriers to the movement of dislocations and thereby rendering appreciable influence on the strength properties of the alloy. In addition, upon quenching of the specimens, which were TMPed at higher temperatures, more amounts of fine Si precipitates deposit at the Si–Al interface, and thus improve the Si–Al interfacial bonding strength [35]. Meanwhile, the effect of strain rate on the UTS may be rationalized considering the role of tangled dislocation structures in hindering the dislocation movement, which is more prominent under higher strain rates [36]. Apart from this, the lack of dislocation annihilation during deformation with high strain rate contributes to the formation of ultrafine substructures thereby increasing the alloy strength. It is also reported that the rate of damage accumulation in the Al–Si alloys is a strong function of Si particles shape and size [5]. Coarse acicular Si particles with high shape factor are considered as preferred sites for premature crack initiation during deformation. This is related to the high levels of stress concentration
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(a) 240
around the lengthy Si particles [37]. In contrast, the Si particles with spherical shape exhibit a higher resistance to debonding under the imposed straining [38]. As the Si spheroidization is more pronounced at higher TMP temperatures [39], the observed enhancement in the UTS values with increasing the TMP temperature seems reasonable. Moreover, concerning the effect of strain rate on the strength, the size of the Si particles plays a vital role. In this regard, Gurland and Plateau [40] proposed the following equation to estimate the required critical stress for developing a crack in a particle of size D:
Shear Stress (MPa)
200 160 120 80 40
"
0 0
0.05
0.1
0.15
0.2
0.25
0.3
Displacement (mm)
(b)
0.1(1/s)
0.01(1/s)
UTS (MPa)
400 300 200 100 As-cast
420
460
500
540
Temperature (oC) Fig. 4. (a) A typical shear strength–displacement curve obtained by shear punch testing of examined alloy after TMP at 460 °C under the strain rate of 0.1 s1, and (b) The variations of UTS with different TMP conditions along with the as-cast value.
140 0.1(1/s)
0.01(1/s)
0.001(1/s)
HB (N.mm-2)
125 110 95
ð2Þ
Dq2
1
r ¼ r0 þ kk2
ð3Þ
where r is the strength, k and r0 are constants and k is the mean particle interspacing. In addition to all the above-mentioned factors, the finer grain size of the microstructure due to the breaking of Al dendrites could result in enhanced mechanical properties through Hall–Petch strengthening mechanism [41,42]. It is also reported that dendrite fineness associated with a reduction in dendrite spacing could have a positive influence on the mechanical characteristics of cast structures [43–45]. The hardness variations with TMP parameters are represented in Fig. 5. It is obvious that the hardness values change in the same way as the UTS results (compare with Fig. 4). The same interpretations stated to explain the UTS variations remain valid for the hardness evolution, as well.
80 65 50
420
460
500
540
o
Temperature ( C) Fig. 5. Brinell hardness values for the experimental alloy after TMP under different conditions.
25 0.1(1/s)
0.01(1/s)
0.001(1/s)
20
Elongation (%)
#12
where q is the stress concentration factor, E is the weighted average of the elastic moduli of the particle and matrix and c is the interfacial energy of the crack. Taking the Eq. (2) into consideration and the fact that at a given temperature, the size of the Si particles becomes smaller with increasing the strain rate [39], the higher UTS value at higher strain rate would be rationalized. From another point of view, having smaller Si particles is associated with a decrease in particle interspacing, which, in turn, leads to mechanical strengthening through Orowan dislocation bowing mechanism:
0.001(1/s)
500
0
r¼
Ec
15 10 5 0
As-cast
420
460
500
540
3.2.2. Ductility Fig. 6 illustrates the effect of TMP parameters on the tensile elongation of the experimental alloy. For comparison, the elongation value of the cast specimen is also included in the chart. A dramatic increase in the elongation value of the cast specimen is observed upon TMP under any conditions. However, according to Fig. 6, the higher ductility values are obtained at higher temperatures but lower strain rates. The acceleration of homogenization and the dissolution/morphological changes of brittle intermetallic compounds with increasing the temperature are believed to be reasons for the observed behavior [46,47]. Moreover, as is well accepted, the A356 Al alloys with rounder Si particles exhibit higher ductility values due to the slow accumulation of damage during deformation [37]. In this regard, the microstructural studies in the present work showed a higher degree of spheroidization at higher temperatures and lower strain rates. Fig. 7, accordingly, depicts the variations of Si particle roundness under different TMP conditions where roundness is calculated using the following equation:
Particle roundness ¼
4 particle Area 100 Particle perimeter
ð4Þ
o
Temperature ( C) Fig. 6. Elongation to failure for the cast and TMPed A356 Al alloy under various conditions.
It is apparent from Figs. 6 and 7 that the elongation and roundness variations with TMP temperature and strain rate follow a similar trend. This gives support to the idea of attributing the higher
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100 0.01(1/s)
0.1(1/s)
0.001(1/s)
Quality Index (MPa)
Si Particle Roundness
0.1(1/s)
80 60 40 20 0
0.01(1/s)
0.001(1/s)
600 500 400 300 200
420
460
500
540
As-cast
420
Fig. 7. The variations of Si particle roundness with different TMP routes.
ductility values of the A356 Al alloy at higher temperatures and lower strain rates to the higher spheroidization degree of the Si particles. Furthermore, at higher temperatures and lower strain rates, local stresses around the eutectic Si particles subside due to the ease of the dislocation movement. Generally, due to plastic incompatibility and differences in the coefficient of thermal expansion between the Si particles and aluminum matrix, local stresses build up at the particle/matrix interface, which is dictated by an increase in the density of dislocations therein. Moreover, from the metallurgical point of view, it is well known that the Al-rich (a) phase and fibrous or needle-like Si particles have dissimilar growth behavior with Si particles growing from the liquid in a faceted manner (smooth growth interface), while the a-Al phase solidifies with surfaces that are rough. Due to this, some localized deformations are expected to occur in their boundary, which is not perfectly conformed [12,13,44,45]. This seems to be another reason for the local stress build-up at the interface. In view of their thermally activated nature, dislocation annihilation, climbing and cross slip are more encouraged at higher temperatures. This, in turn, gives rise to stress relaxation and the reduction of local stress incompatibilities at the matrix/particle interface, which ultimately render a positive influence on the ductility of the experimental alloy [48]. Besides contributing to the higher degree of dynamic recovery, the effect of lower strain rate at a given temperature may be realized considering the size of the Si particles. As was previously discussed, under lower strain rates, the lesser degree of particle fragmentation and the longer available time for the diffusion of Si atoms lead to the greater average size of the Si particles [39]. As a result, the interspacing between the Si particles increases. Hence, a larger volume of the aluminum matrix would be available between particles to blunt the interface-originated cracks, and thereby improving the elongation [49]. 3.2.3. Quality index Towards a comprehensive view of the mechanical performance of the Al–Si casting alloys, the empirical parameter of quality index, Q, has been employed by many researchers to date [50]. For instance, Ceschini et al. [51] took advantage of the Q-parameter in order to draw a correlation between the tensile properties and solidification microstructure of an A357 Al alloy. In addition, the quality index was used by Shabestari and Shahri [52] to investigate the optimum modification and solidification conditions of an A356 Al alloy, which gives the best tensile properties. In similar fashion, this study utilizes the Q-parameter to evaluate the relationship between the TMP parameters and the mechanical quality of the experimental A356 Al alloy. The quality index involves a linear combination of two tensile parameters, i.e., UTS and elongation, each of which becomes weighed by different coefficients of C and K, respectively:
460
500
540
o
Temperature ( C)
o
Temperature ( C)
Fig. 8. Quality index values for the cast and TMPed A356 Al alloy under various conditions.
Q ¼ C UTS þ K log ð%EÞ
ð5Þ
Drouzy et al. [53] described the quality index for Al–Si alloys by the following expression in which fixed numbers replace the weight coefficients:
Q ¼ UTS ðMPaÞ þ 150 log ð%EÞ
ð6Þ
Fig. 8 shows the calculated Q-values of the cast and TMPed specimens under various conditions. As is seen, the highest Q-values are obtained for the specimens processed at higher temperatures and strain rates. The observed trend is in line with the UTS variations in Fig. 4. It is also interesting to note from Fig. 8 that the highest Q-value (corresponding to the specimen processed at 540 °C under the strain rate of 0.1 s1) is calculated to be about 630 MPa. This value is much higher than the highest values reported in the literature for the chemically modified and heat-treated Al–Si alloys [51,52].
4. Conclusions In the present study, the room temperature mechanical properties and microstructure evolution of a thermomechanically processed unmodified A356 Al alloy have been investigated. Based on this research, the following outcomes are reported: (1) With the aid of TMP, the spheroidization of the Si particles was occurred at relatively low temperatures and durations. In consequence, the TMPed A356 Al alloy acquired a higher strength without sacrificing its ductility. (2) Higher UTS values were obtained upon TMP at higher temperatures and strain rates. The more homogeneous matrix microstructure containing rounder Si particles at high temperatures was considered as one good reason for the observed behavior. In addition, increasing the strain rate results in having smaller Si particles which, in turn, leads to mechanical strengthening through Orowan dislocation bowing mechanism. (3) Brinell hardness variations for the TMPed specimens followed the same trend as the UTS variations. (4) Increasing the TMP temperature and/or decreasing the strain rate led to higher ductility values, mainly due to the spheroidization of the Si particles together with the restoration of the deformed microstructure. (5) The overall mechanical performance of the TMPed unmodified A356 Al alloy was assessed using quality index. The highest Qvalues were obtained for the specimens processed at higher temperatures and strain rates.
N. Haghdadi et al. / Materials and Design 43 (2013) 419–425
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