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Effect of Al2O3 nano-particles on the microstructural stability of AZ31 Mg alloy after equal channel angular pressing
Materials Science and Engineering A 527 (2010) 2764–2771
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Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Effect of Al2 O3 nano-particles on the microstructural stability of AZ31 Mg alloy after equal channel angular pressing Y. Radi, R. Mahmudi ∗ School of Metallurgical and Materials Engineering, University of Tehran, Tehran, Iran
a r t i c l e
i n f o
Article history: Received 10 November 2009 Received in revised form 25 December 2009 Accepted 8 January 2010
Keywords: Magnesium alloy ECAP Thermal stability Nano-particle Grain growth
a b s t r a c t The microstructure and thermal stability of the equal channel angularly pressed (ECAPed) AZ31 magnesium alloy and its composites, reinforced with 0.5, 1, and 2 wt.% of Al2 O3 nano-particles, were investigated. The alumina nano-particles with an average diameter of 100 nm were added to AZ31 by a stir-casting method. After extruding, the materials were ECAPed at 500 K for 4 passes using route Bc, in which each sample was rotated 90◦ around its longitudinal axis between the passages. Textural studies revealed that nano-particles increased the intensity of the maximum orientations of the basal and prismatic planes in both extruded and ECAPed conditions, without changing their final positions. The microstructural stability of the ECAPed AZ31 and AZ31–1% Al2 O3 was examined by isochronal and isothermal annealing regimes in the temperature range of 473–740 K. The measured activation energies for isochronal grain growth showed three different values, depending on the temperature range investigated. In the low- and high-temperature ranges of 490–530 K and 700–740 K, the respective activation energies were 74.1 and 94.4 kJ/mol. These energies are respectively 0.8Qgb and 1.03Qgb , where Qgb is the activation for grain boundary diffusion. For the intermediate-temperature range of 530–700 K, however, the unexpectedly low activation energy of Q = 0.21Qgb was obtained. The same trend was observed for the particle-containing material with higher activation energies. To consider the pinning force of nanoparticles, the Burke’s model of grain growth during isothermal annealing was applied. Accordingly, two grain-growth regimes of the low-temperature region (<630 K) and the high-temperature region (>630 K) were identified. For the low-temperature region, the low activation energy of 32.1 kJ/mol may correspond to the energy for the reordering of grain boundaries in the fine-grained composite material. The activation energy of 110.9 kJ/mol found for temperatures higher than 630 K, lies between that for grain boundary diffusion and lattice diffusion of magnesium. The role of nano-particles in grain boundary pinning was also verified by the Hall–Petch plot. © 2010 Elsevier B.V. All rights reserved.
1. Introduction The application of magnesium alloys in automotive, aerospace, and electronic industries is growing due to their low density, good machineability and excellent castability [1]. Despite these advantages, however, Mg alloys have relatively low strength and poor formability, due to their hexagonal closed packed (hcp) structure with limited slip systems. Accordingly, many attempts have been made to address these drawbacks of wrought Mg alloys [1,2]. Microstructural refinement is an effective way for increasing strength and, in some cases, ductility of these alloys. This goal can be achieved by a variety of techniques including processing by equal channel angular pressing (ECAP) [3–5]. In this process, the cast or extruded material is repeatedly pressed in an L-shaped
∗ Corresponding author. Tel.: +98 21 6111 4137; fax: +98 21 8800 6076. E-mail address: [email protected] (R. Mahmudi). 0921-5093/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2010.01.029
die to undergo a very high shear strain deformation without any change in the cross-sectional area. In most investigations, because of the low ductility of Mg alloys, they are ECAPed at or above 473 K to avoid cracking. Processing at such temperatures can result in the absence of non-equilibrium grain boundaries with fine grain sizes of about 1 m [6], although the grain size of around 2.2 m with larger grains around 6 m after annealing has been reported in a superplastic ECAPed AZ31 [7]. The ECAPed microstructures usually consist of equiaxed grains with a high degree of homogeneity [3]. These structures are, however, susceptible to coarsening and instability at high temperatures, due to the high stored energy caused by severe plastic deformation [8,9]. Efforts have been made to develop new magnesium alloys having improved structural stability at high temperatures [10]. This has been achieved by alloying with different elements such as Zr, Ca, Sb, and rare earth (RE) elements. These elements can either react with the Mg matrix and other alloying elements of the base alloy to form thermally stable precipitates, or refine the microstructure by
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producing intermetallic particles which act as nucleation sites during solidification, both of which are beneficial to the improvement of mechanical properties [11–15]. The thermally stable intermetallics can decelerate grain boundary diffusion, resulting in enhanced stability of the microstructure in regions adjacent to grain boundaries. The improvement in thermal stability obtained by the alloying elements, however, might not be as high as those obtained by composite reinforcements. In fact, reinforcing magnesium alloys with the highly stable ceramic particles can further improve their properties [16–18]. The size of the reinforcing particles is very important, as the micron-sized particles usually cause low ductility, even though there is no reason for these particles to decrease tensile strength [19]. Dispersing thermally stable nano-sized oxide ceramic particulates in magnesium eliminates these problems and enhances mechanical properties [18,20]. The achieved thermal stability has been mainly attributed to the pinning effects of the nano-particles which reduce the mobility of the grain boundaries during hot deformation process and the subsequent annealing treatments [18,21,22]. However, a change in deformation mechanism by sub-micron particles [23] and activation of non-basal slip by nano-particles [24] have also been reported for the observed ductility improvements in magnesium alloys. The present work examines the simultaneous effects of alumina nano-particle additions and severe plastic deformation by ECAP on the restoration phenomena, microstructure and thermal stability of the AZ31 magnesium alloy. 2. Experimental procedure 2.1. Materials and processing The AZ31 magnesium alloy (Mg–3 wt.% Al–1 wt.% Zn–0.2 wt.% Mn) was chosen as the matrix and Al2 O3 nano-particles, with an average diameter of 100 nm, as reinforcements. The melt was prepared from high purity (99.8%) magnesium, aluminum, zinc, and manganese, melted in a graphite crucible placed in an electrical resistance furnace under the Foseco Magrex 36 covering flux held at 1023 K. Three different weight percents of nano-particles (0.5, 1 and 2%) were added to the molten base material using an Mg-8 wt.% Al2 O3 nano-composite master alloy developed at the University of British Columbia (UBC). Based on the previous experience [13–15], the melt was held at this temperature for 10 min and then stirred mechanically for 2 min using a stainless steel rod to provide a homogeneous composition. It was then poured into a steel mold preheated to 473 K, using a tilt-casting system to minimize casting defects and turbulence of the melt. The cast billets of 44 mm diameter were annealed at 673 K for 1 h and extruded to 11 mm × 11 mm bars at 638 K at a speed of 2.0–2.5 mm/s. The ECAP billets having dimensions of 10 mm × 10 mm × 80 mm were machined from the extruded bars. ECAP was performed at 503 K through a solid die having an angle between the two channels of ˚ = 90◦ and an angle of = 20◦ at the outer arc of curvature at the point of intersection of the two channels, as shown in Fig. 1. Repetitive pressing of the same sample was carried out with route Bc, in which each sample was rotated 90◦ around its longitudinal axis between the passages. This configuration leads to an imposed strain of about 100% on each passage through the die. Samples were sprayed with MoS2 lubricant and pressed at a speed of 1 mm/s for 4 passes. 2.2. Textural studies Texture samples were finely polished and immersed several times in a diluted solution of HCl to eliminate mechanical effect of the preparation procedure. Texture measurements were performed
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Fig. 1. Schematic diagram of the ECAP system used in this research.
using the Schulz reflection method in a Philips X’Pert diffractometer furnished with a close Eulerian cradle. The intensity distributions of the (0 0 0 2) and (1 0 1¯ 0) pole figures were measured from the planes parallel to the pressing direction for both extruded and ECAPed monolithic and composite materials. The measurement was performed using Co k␣ radiation at 50 kV with the sample tilt angle ranging from 0◦ to 90◦ . 2.3. Microstructural studies Both of the extruded and ECAPed bars were sliced into 5mm-thick discs to study the microstructure of the cross-sections perpendicular to the pressing direction. Optical and scanning electron microscopy (SEM) were used to examine the grain size and reinforcement distribution. All of the samples were polished by 0.05-m Al2 O3 and then etched with a solution of 10 ml acetic acid, 4.2 g picric acid, 20 ml H2 O and 50 ml ethanol for 3 s at room temperature. The microstructures of the extruded and ECAPed samples were examined using a Leitz optical microscope and a CamScanMV2300 scanning electron microscope (SEM), respectively. The Clemex vision professional image analysis program was used to measure the grain size and distribution according to the ASTM E112 standard. To have more representative grain size values, at least 400 grains were considered for each condition. Different samples were employed and different locations on each sample were pictured and the number of grains on each micrograph was taken as the weighting factor in the calculation of the final grain size. 2.4. Thermal stability tests The thermal stability of the monolithic AZ31 and AZ31–1% Al2 O3 composite samples was investigated by the isochronal and isothermal annealing of the ECAPed materials and measuring the grain size evolution. Isochronal annealing was performed for 30 min at thirteen equally spaced temperatures in the range 473–740 K, while isothermal annealing was conducted for 300–30,000 s at five equally spaced temperatures in the same temperature range. All
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annealing tests were conducted in an air-circulated furnace with an accuracy of ±2 K. 2.5. Microhardness The ECAPed AZ31 and AZ31–1% Al2 O3 samples were annealed at 13 equally spaced temperatures in the range 473–740 K for 30 min to produce a range of grain sizes. Vickers microhardness (Hv ) was measured on the plane perpendicular to the pressing direction in a Wolpert microhardness tester with a load of 200 g for 15 s. Each reading was an average of at least five separate measurements taken at random places on the surface of the specimens. The acquired data were used in the construction of the Hall–Petch diagrams to elucidate the role of nano-particles in the strengthening of the composite material. 3. Results and discussion 3.1. Microstructure and texture evolution Optical images of the extruded bars in the cross-sections perpendicular to pressing direction are shown in Fig. 2. The initial grain size of the monolithic AZ31 reduced from 19.9 to 13.7 m, as the nano-particle content increased to 2 wt.%. It can be depicted that all extruded materials show normal grain size distributions, with standard deviations in the range 0.24–0.41 m. Due to the relatively high extrusion temperature of 638 K (T ∼ 0.7Tm ), dynamic recrystallization (DRX) could occur, resulting in the observed grain refinement. DRX is based on the grain boundary migration [25], and thus particles having pinning drag could postpone this phenomenon so that finer grains could be obtained in the presence of second phase particles. SEM images of the ECAPed materials, also shown in Fig. 2, indicate that significant refinement has occurred after severe plastic deformation. The degree of the achieved refinement has increased with the Al2 O3 content and the normal grain size distribution has become more skewed after the addition of 2% nano-particles. The standard deviation of the measured grain sizes for the ECAPed conditions was in the range 0.17–0.28 m. It can be seen that some fine-grained colonies are surrounded by coarser grains in some regions of the microstructure. One possibility is that the nucleation of DRX was halted due to the pinning effect of particles or, the nucleation took place but the grain growth following nucleation did not developed to produce equiaxed larger grains. Therefore, increasing the alumina content to 2% in the ECAPed AZ31 could postpone the restoration phenomena and raise the DRX temperature. An increase in the area of fine-grained regions in the AZ31–2% Al2 O3 composite could support the idea of postponed DRX during ECAP. In contrast to AZ31–2% Al2 O3 , for the other two composites a rather more normal distribution of equiaxed and uni-sized grains could be observed. The (0 0 0 2) and (1 0 1¯ 0) pole figures of the AZ31 and AZ31–2% Al2 O3 samples in the extruded and ECAPed conditions are shown in Fig. 3. In the as-extruded materials (Fig. 3a and b), it is evident that {0 0 0 2} basal planes are mostly oriented parallel to the extrusion direction. After ECAP, however, both basal and prismatic planes are inclined about 45◦ to the extrusion axis, as shown in Fig. 3c and d. The comparison of the {0 0 0 2} pole figures of AZ31–2% Al2 O3 with that of AZ31 shows that the particles have not affected the final positions of the texture. However, the intensity of the maximum orientations in both of the extruded and ECAPed materials increased by about 30% after Al2 O3 addition. The same is true for the prismatic planes, for which the intensity increases by about 44 and 60% for the extruded and ECAPed conditions, respectively. It is further demonstrated that ECAP weakens the basal poles and strengthens the prismatic poles. The observed variations in the
textures of the materials could be indicative of various deformation mechanisms operating in different conditions of the studied materials. 3.2. Annealing and thermal stability It was shown that the grain size distribution of the AZ31–2% Al2 O3 composite was rather unusual and the texture evolution for this composite exhibited drastic changes. Static annealing of this material was carried out at 573 K for 30 min. The annealed microstructure, shown in Fig. 4a, consists of a bimodal grain structure with enlarged grains in both fine and coarse regions. Although the grain size distribution is near-normal within each region (Fig. 4b), the overall grain distribution is clearly bimodal with two distinct grain size distributions, as depicted in Fig. 4c. This behavior is in agreement with those previously reported for nano- and ultrafine-grained magnesium alloys [14,26,27]. Due to the non-uniform grain structure of AZ31–2% Al2 O3 , the measurement of the grain size is rather complicated and weighted average grain sizes should be used, as proposed by Rajasekhara et al. [28]. This is in contrast to the case of monolithic AZ31, in which the grain size measurement of the uniform equiaxed microstructure is not affected by such a complication. Therefore, AZ31–1% Al2 O3 having a normal grain size distribution was selected for the thermal stability studies. Isochronal annealing was conducted for 1800 s on the AZ31 and AZ31–1% Al2 O3 samples, which are hereafter referred to as “base” and “composite”, respectively. Thirteen various test temperatures were selected in the range 473–740 K to produce a wide range of grain sizes. The variation of grain size with annealing temperature, shown in Fig. 5a, indicates that although grain growth occurs in both materials, the composite material has a lower grain size than the base AZ31 at all temperatures. Normal grain growth in singlephase alloys under isochronal heat treatment conditions is well described by the following equation [29]:
Q
Dn − D0n = Ct exp −
RT
(1)
where D and D0 are the current and initial values of the average grain size, respectively, Q is the activation energy of grain growth, R is the gas constant, T is the temperature, t is the annealing time, n is the grain growth exponent, and C is a kinetic constant. When a parabolic relationship for grain growth is assumed, the n-value equals 2 and the slope of the plot ln(D2 − D02 ) against 1/T would yield the grain growth activation energy for the samples annealed at different temperatures for a fixed annealing time of 1800 s. This was done for the measured grain sizes, assuming the respective initial grain sizes (D0 ) are 6.4 and 5.3 m for the monolithic and composite materials, as shown in Fig. 5b. It is evident that there are three distinct temperature ranges with different activation energies. The activation energies of 74.1, 19.2, and 94.4 kJ/mol are found for the base material in the low-, intermediate-, and high-temperature ranges, respectively. It is worth comparing these activation energies with those for the lattice self-diffusion (Qlsd = 135 kJ/mol) [30] and grain boundary diffusion in pure Mg (Qgb = 92 kJ/mol) [31]. In the low-temperature range of 490–530 K, the activation energy of 74.1 kJ/mole is less than Qgb . In the high-temperature range of 700–740 K, however, the activation energy of 94.4 kJ/mol is less than Qlsd but very close to Qgb . In the intermediate-temperature range of 530–700 K, the activation energy of 19.2 kJ/mol is significantly lower than both Qlsd and Qgb , corresponding to about 0.21Qgb . These comparisons can be used in the identification of grain growth mechanisms in different temperature intervals. The abnormally low activation energy for grain growth at intermediate annealing temperatures has also been reported for an ECAPed Al–Mg alloy by Wang et al. [32]. Their measured
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Fig. 2. Microstructure and grain size distribution of extruded and ECAPed AZ31 and composites containing 0.5, 1.0 and 2.0 wt.% of nano-particles. Images for the extruded materials are taken by optical microscopy and those for the ECAPed materials are by SEM.
Fig. 3. (0 0 0 2) and (1 0 1¯ 0) pole figures of (a) extruded AZ31, (b) extruded AZ31–2% Al2 O3 , (c) ECAPed AZ31, and (d) ECAPed AZ31–2% Al2 O3 .
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Fig. 5. (a) Variation of grain size with annealing temperature for the base and AZ31–1% Al2 O3 composite samples isochronally annealed for 1800 s, and (b) plot of ln(D2 − D02 ) against 1/T to determine activation energies according to Eq. (1).
Fig. 4. (a) Bimodal grain structure of AZ31 containing 2 wt.% Al2 O3 nano-particles subjected to 4 passes of ECAP and subsequently annealed at 573 K for 3600 s, (b) grain size distribution for fine- and coarse-grained regions, and (c) overall grain size distribution.
activation energy of 0.36Qgb was attributed to the unrecrystallized microstructure with non-equilibrium grain boundaries containing a large number of extrinsic dislocations, which may show higher atomic mobility compared to the equilibrium grain boundaries. The results of structural stability of an ECAPed AZ31 obtained by Kim and Kim [8] are indicative of a Q = 0.27Qgb in the intermediatetemperature range of 523–673 K. The possibility of grain growth under the unrecrystallized condition was, however, rejected due to the fact that the non-equilibrium boundaries were not present in the as-ECAPed (recrystallized) microstructure. Using the concept of diffusivity, they argued that progressive decrease in dislocation density by enhanced recovery with increasing temperature may be the cause of the observed anomaly in the activation energy. Our results on the activation energy of grain growth are in agreement with those obtained by Kim and Kim [8]. The slight observed difference could be attributed to the higher ECAP temperature in the present work, which affects both grain size and distribution and thus the grain growth process. The same three temperature zones exist for the activation energy of composite material, the data of which are superimposed on those of the AZ31 shown in Fig. 5b. It is clear that the activation energies of the composite material are higher than those of the base alloy, being 117.9, 22.7, and 88.6 kJ/mol for the high-, intermediate, and low-temperature ranges, respectively. These higher activation energies for grain growth in this condition could be attributed to the presence of alumina nano-particles. In the low- and intermediatetemperature ranges, recovery is the dominant thermally activated
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[33] demonstrated that dispersed particles can pin the grain boundaries, in which case the traditional kinetic theory of grain growth expressed by Eq. (1) could not predict the growth behavior. According to Burke’s model, it is assumed that drag force is independent of grain size and hence grain growth rate is not controlled by the instantaneous grain size, D. Instead, the decreasing difference between ultimate limiting grain size (Dm ) and the changing value of the instantaneous grain size (D) can control the growth rate, and hence Burke’s model may be expressed as: D0 − D + ln Dm
D − D m 0 Dm − D
=
k0 t 2 Dm
Q
exp −
RT
(2)
where D0 is the initial grain size, D is the instantaneous grain size, Dm is the limiting ultimate grain size for the particular annealing temperature, and k0 is a constant. By differentiating Eq. (2), the basic growth rate equation is as: dD =k dt
Fig. 6. (a) Variation of grain size with annealing time for the AZ31–1% Al2 O3 composite samples isothermally annealed in the range 520–740 K, and (b) plot of dD/dt against 1/D to determine the k parameters in Eq. (3).
phenomenon. In this process, dislocation movement is the fundamental assumption [25] and thus an increase in the activation energy for grain growth of the composite material can be attributed to fact that dislocation movement has been decelerated by the nano-particles. At higher temperatures, pinning of grain boundaries is the possible mechanism for the increase of grain growth activation energy. It seems that nano-particles have been quite successful in this process, as the activation energy of 94.4 kJ/mol for the base material has increased to 117.9 kJ/mol for the composite material. The obtained activation energies for grain growth are not valid for long time annealing, because in the beginning stages dislocations should be annihilated and some amounts of energy is spent for this purpose. Therefore, the exact activation energy for grain growth should be obtained through isothermal annealing at long times. This would highlight the effects of nano-particles on the grain growth behavior of the composite material. It is expected that this type of annealing would cause no significant change on the grain growth results of the AZ31 alloy because the traditional theory can satisfactorily predict the grain growth behavior of monotonic alloys. Accordingly, isothermal annealing was only performed on the composite samples. Because of the long annealing times involved, the interactions between dislocations and particles are ignored and only the pinning effects of nano-particles on the grain boundary movement are considered. The variation of grain size with annealing time at different constant temperatures, shown in Fig. 6a, indicates that the traditional parabolic trend is not observed for the composite material. Burke
1
D
−
1 Dm
(3)
where k = k0 exp(−Q/RT). By plotting dD/dt against 1/D for each temperature, linear relationships with correlation factors greater than 0.98 is obtained, as shown in Fig. 6b. The slopes of these lines show an increase with increasing annealing temperature. Since k has an Arrhenius type relationship with temperature, plotting ln(k) versus 1/T would yield the activation energy of grain growth considering drag force of particles, as exhibited in Fig. 7. The slopes of the plotted lines show two different activation energies at the high- and low-temperature regimes, respectively. For the low-temperature region, the low activation energy of 32.1 kJ/mol may correspond to the energy for the reordering of grain boundaries in the finegrained material [34]. The activation energy of 110.9 kJ/mol found for temperatures higher than 630 K, however, lies between that for grain boundary diffusion and lattice diffusion of magnesium. These results indicate that the observed difference between the activation energies for grain growth in the low- and high-temperature ranges is higher than that obtained from the traditional grain growth theory. This can be attributed to the fact that dislocation annihilation is assumed to affect the activation energy in the traditional grain growth theory, while in the Burke’s model it is considered to be insignificant due to the long annealing times. More importantly, the pinning effect of particles included in the Burke’s analysis has affected the observed behavior. It is worth noting that the higher grain growth activation energy in the high-temperature regime is
Fig. 7. Plot of ln(k) against 1/T for the AZ31–1% Al2 O3 composite samples to calculate activation energies for the low- and high-temperature regimes according to Burke’s model.
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movement. This resistance can be resulted from a single or a combination of various strengthening mechanisms such as precipitation strengthening, strain hardening, and solid solution strengthening. The obtained H0 value of 38.8 for the composite material is higher than 36.7 obtained for the base alloy. The enhanced friction strength is attributed to the presence of alumina nano-particles that hinder dislocation movement during plastic deformation caused by the penetration of the indenter. The slope of the H–P lines, K, is the locking parameter which measures the relative hardening contribution of the grain boundaries. The higher K-value of 60.9 for the composite material, compared to 51.7 obtained for the base material, is indicative of the positive role of the nano-particles in strengthening the grain boundaries by a pinning mechanism. 4. Conclusions The microstructural stability of the equal channel angularly pressed (ECAPed) AZ31 magnesium alloy and its composites, reinforced with 0.5, 1, and 2 wt.% of Al2 O3 nano-particles, were investigated and the following conclusions were made:
Fig. 8. (a) Microhardness of the specimens isochronally annealed for 1800 s at different temperatures, and (b) Hall–Petch relationship for AZ31 and AZ31–1% Al2 O3 .
consistent with the results on ultrafine-grained aluminum alloy, reported by Roy et al. [34]. 3.3. Hall–Petch relationship The Hall–Petch (H–P) relationship, which relates the hardness and strength to the grain size, has frequently been used to investigate the contribution of grain boundary strengthening as well as other strengthening mechanisms to the overall strength and hardness [28,35]. The general form of the H–P relationship for hardness is given by: Hv = H0 + KD−1/2
(4)
where Hv is the measured microhardness, H0 is the lattice friction hardness, D is the average grain size, and K is a constant. In the present study, the H–P approach would help clarifying the role of nano-particles in the achieved strengthening. Accordingly, the microhardness of the ECAPed monolithic and composite materials was measured and plotted against annealing temperature in Fig. 8a. As can be seen, at all annealing temperatures the microhardness of the composite is higher than the base alloy. In Fig. 8b, the classic H–P relationship is plotted for AZ31 and AZ31–1% Al2 O3 . It is obvious that both the friction hardness (H0 ) and the locking parameter (K) are higher for the particle-containing material. The friction stress or hardness, determined from the intercept of the H–P line, represents the overall resistance of the crystal lattice to dislocation
1. The introduction of 2 wt.% Al2 O3 nano-particles into the base AZ31 alloy reduced the grain size from 19.9 to 13.7 m in the extruded state, and from 6.4 to 4.3 m in the ECAPed condition. Microstructural analysis showed that the combination of ECAP and nano-particle addition results in a more homogeneous grain size distribution. These particles increased the intensity of the maximum orientations of the basal and prismatic planes in both extruded and ECAPed conditions. 2. The isochronal grain growth showed three distinct regions, depending on the temperature range investigated. In the lowand high-temperature ranges of 490–530 K and 700–740 K, the respective activation energies were 74.1 and 94.4 kJ/mol. These energies are respectively 0.8Qgb and 1.03Qgb , where Qgb is the activation for grain boundary diffusion. For the intermediatetemperature range of 530–700 K, however, the unexpectedly low activation energy of Q = 0.21Qgb was obtained. The higher activation energy for the composite material in all three regions reflects the role of nano-particles in pinning grain boundaries as compared to the monolithic material. 3. According to Burke’s model for isothermal annealing, two graingrowth regions with significantly different activation energies of 32.1 and 110.9 kJ/mol were identified for the low- and hightemperature regimes, respectively. The relatively low activation energy of 32.1 kJ/mol for grain growth in the low-temperature regime can be attributed to the reordering and readjustment of grain boundaries. 4. Hall–Petch relationship constructed for the monolithic and composite materials showed higher values of H0 and K for the composite material. This implies that nano-particles have effectively contributed to grain boundary pinning and also to interacting with dislocations to increase the friction stress. Acknowledgement The authors wish to gratefully acknowledge Professor Tom Troczynski at the University of British Columbia (UBC) who provided the Mg–Al2 O3 nano-composite master alloys used to fabricate the materials used in this study [36]. References [1] B.L. Mordike, T. Ebert, Mater. Sci. Eng. A 302 (2001) 37–45. [2] I.J. Polmear, Light Alloys; Aluminum, Magnesium, Titanium, 2nd ed., Edward Arnold Co., London, 1989, pp. 169–189. [3] K. Xia, J.T. Wang, X. Wu, Mater. Sci. Eng. A 410–411 (2005) 324–327. [4] L. Jin, D. Jin, D. Moa, Mater. Sci. Eng. A 423 (2006) 247–252.
Y. Radi, R. Mahmudi / Materials Science and Engineering A 527 (2010) 2764–2771 [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22]
B.Q. Han, T.G. Langdon, Mater. Sci. Eng. A 410–411 (2005) 435–438. R.B. Figueiredo, T.G. Langdon, Mater. Sci. Eng. A 503 (2009) 141–144. R.B. Figueiredo, T.G. Langdon, J. Mater. Sci. 43 (2008) 7366–7371. H.K. Kim, W.J. Kim, Mater. Sci. Eng. A 385 (2004) 300–308. H.K. Kim, J. Mater. Sci. Lett. 39 (2004) 7101–7103. A.A. Luo, Int. Mater. Rev. 49 (2004) 13–30. Y. Miyahara, K. Matsubara, Z. Horita, T.G. Langdon, Metall. Mater. Trans. 36A (2005) 1705–1711. X. Gao, S.M. Zhu, B.C. Muddle, J.F. Nie, Scripta Mater. 53 (2005) 1321–1326. F. Kabirian, R. Mahmudi, Adv. Eng. Mater. 11 (2009) 189–193. M. Masoudpanah, R. Mahmudi, Mater. Sci. Eng. A 526 (2009) 22–30. F. Kabirian, R. Mahmudi, Metall. Mater. Trans. 40A (2009) 2190–2201. X.L. Zhong, W.L.E. Wong, M. Gupta, Acta Mater. 55 (2007) 6338–6344. S.F. Hassan, M. Gupta, Metall. Mater. Trans. 36A (2005) 2253–2258. M. Habibnejad-Korayem, R. Mahmudi, W.J. Poole, Mater. Sci. Eng. A 519 (2009) 198–203. S. Seshan, M. Jayamathy, S.V. Kailas, T.S. Srivatsan, Mater. Sci. Eng. A 363 (2003) 345–351. H. Ferkel, B.L. Mordike, Mater. Sci. Eng. A 298 (2001) 193–199. H. Cao, J.Y. Min, S.D. Wu, A.P. Xian, J.K. Shang, Mater. Sci. Eng. A 431 (2006) 86–91. M. Ferry, N.E. Hamilton, F.J. Humphreys, Acta Mater. 53 (2005) 1097–1109.
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[23] N. Stanford, M. Barnett, Scripta Mater. 58 (2008) 179–182. [24] C.S. Goh, J. Wei, L.C. Lee, M. Gupta, Acta Mater. 55 (2007) 5115–5121. [25] F.J. Humphreys, M. Hatherly, Recrystallization and Related Annealing Phenomena, Elsevier, Oxford, United Kingdom, 1996. [26] Q. Yang, A.K. Ghosh, Acta Mater. 54 (2006) 5147–5158. [27] Y.M. Wang, E. Ma, Acta Mater. 52 (2004) 1699–1709. [28] S. Rajasekhara, P.J. Ferreira, L.P. Karjalainen, A. Kyrolainen, Metall. Mater. Trans. 38A (2007) 1202–1210. [29] P.A. Beck, J.C. Kremer, L.J. Demer, M.L. Holzworth, Trans. TMS-AIME 175 (1948) 372–400. [30] P.G. Shewman, Trans. AIME 206 (1956) 918–922. [31] H.J. Frost, M.E. Ashby, Deformation Mechanism Maps, Pergamon Press, Oxford, 1982. [32] J. Wang, Y. Iwahashi, Z. Horita, M. Furukawa, M. Nemoto, R.Z. Valiev, F.G. Langdon, Acta Mater. 44 (1996) 2973–2982. [33] J.E. Burke, Trans. TMS-AIME 180 (1949) 73–91. [34] I. Roy, M. Chauhan, E.J. Lavernia, F.A. Mohamed, Metall. Mater. Trans. 37A (2006) 721–730. [35] R. Mahmudi, Scripta Metall. Mater. 32 (1995) 781–786. [36] S. Vasisht, A. Tsvetkov, T. Troczynski, AUTO21 Conference 2006, Barrie, Ontario, 2006.
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Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Effect of Al2 O3 nano-particles on the microstructural stability of AZ31 Mg alloy after equal channel angular pressing Y. Radi, R. Mahmudi ∗ School of Metallurgical and Materials Engineering, University of Tehran, Tehran, Iran
a r t i c l e
i n f o
Article history: Received 10 November 2009 Received in revised form 25 December 2009 Accepted 8 January 2010
Keywords: Magnesium alloy ECAP Thermal stability Nano-particle Grain growth
a b s t r a c t The microstructure and thermal stability of the equal channel angularly pressed (ECAPed) AZ31 magnesium alloy and its composites, reinforced with 0.5, 1, and 2 wt.% of Al2 O3 nano-particles, were investigated. The alumina nano-particles with an average diameter of 100 nm were added to AZ31 by a stir-casting method. After extruding, the materials were ECAPed at 500 K for 4 passes using route Bc, in which each sample was rotated 90◦ around its longitudinal axis between the passages. Textural studies revealed that nano-particles increased the intensity of the maximum orientations of the basal and prismatic planes in both extruded and ECAPed conditions, without changing their final positions. The microstructural stability of the ECAPed AZ31 and AZ31–1% Al2 O3 was examined by isochronal and isothermal annealing regimes in the temperature range of 473–740 K. The measured activation energies for isochronal grain growth showed three different values, depending on the temperature range investigated. In the low- and high-temperature ranges of 490–530 K and 700–740 K, the respective activation energies were 74.1 and 94.4 kJ/mol. These energies are respectively 0.8Qgb and 1.03Qgb , where Qgb is the activation for grain boundary diffusion. For the intermediate-temperature range of 530–700 K, however, the unexpectedly low activation energy of Q = 0.21Qgb was obtained. The same trend was observed for the particle-containing material with higher activation energies. To consider the pinning force of nanoparticles, the Burke’s model of grain growth during isothermal annealing was applied. Accordingly, two grain-growth regimes of the low-temperature region (<630 K) and the high-temperature region (>630 K) were identified. For the low-temperature region, the low activation energy of 32.1 kJ/mol may correspond to the energy for the reordering of grain boundaries in the fine-grained composite material. The activation energy of 110.9 kJ/mol found for temperatures higher than 630 K, lies between that for grain boundary diffusion and lattice diffusion of magnesium. The role of nano-particles in grain boundary pinning was also verified by the Hall–Petch plot. © 2010 Elsevier B.V. All rights reserved.
1. Introduction The application of magnesium alloys in automotive, aerospace, and electronic industries is growing due to their low density, good machineability and excellent castability [1]. Despite these advantages, however, Mg alloys have relatively low strength and poor formability, due to their hexagonal closed packed (hcp) structure with limited slip systems. Accordingly, many attempts have been made to address these drawbacks of wrought Mg alloys [1,2]. Microstructural refinement is an effective way for increasing strength and, in some cases, ductility of these alloys. This goal can be achieved by a variety of techniques including processing by equal channel angular pressing (ECAP) [3–5]. In this process, the cast or extruded material is repeatedly pressed in an L-shaped
∗ Corresponding author. Tel.: +98 21 6111 4137; fax: +98 21 8800 6076. E-mail address: [email protected] (R. Mahmudi). 0921-5093/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2010.01.029
die to undergo a very high shear strain deformation without any change in the cross-sectional area. In most investigations, because of the low ductility of Mg alloys, they are ECAPed at or above 473 K to avoid cracking. Processing at such temperatures can result in the absence of non-equilibrium grain boundaries with fine grain sizes of about 1 m [6], although the grain size of around 2.2 m with larger grains around 6 m after annealing has been reported in a superplastic ECAPed AZ31 [7]. The ECAPed microstructures usually consist of equiaxed grains with a high degree of homogeneity [3]. These structures are, however, susceptible to coarsening and instability at high temperatures, due to the high stored energy caused by severe plastic deformation [8,9]. Efforts have been made to develop new magnesium alloys having improved structural stability at high temperatures [10]. This has been achieved by alloying with different elements such as Zr, Ca, Sb, and rare earth (RE) elements. These elements can either react with the Mg matrix and other alloying elements of the base alloy to form thermally stable precipitates, or refine the microstructure by
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producing intermetallic particles which act as nucleation sites during solidification, both of which are beneficial to the improvement of mechanical properties [11–15]. The thermally stable intermetallics can decelerate grain boundary diffusion, resulting in enhanced stability of the microstructure in regions adjacent to grain boundaries. The improvement in thermal stability obtained by the alloying elements, however, might not be as high as those obtained by composite reinforcements. In fact, reinforcing magnesium alloys with the highly stable ceramic particles can further improve their properties [16–18]. The size of the reinforcing particles is very important, as the micron-sized particles usually cause low ductility, even though there is no reason for these particles to decrease tensile strength [19]. Dispersing thermally stable nano-sized oxide ceramic particulates in magnesium eliminates these problems and enhances mechanical properties [18,20]. The achieved thermal stability has been mainly attributed to the pinning effects of the nano-particles which reduce the mobility of the grain boundaries during hot deformation process and the subsequent annealing treatments [18,21,22]. However, a change in deformation mechanism by sub-micron particles [23] and activation of non-basal slip by nano-particles [24] have also been reported for the observed ductility improvements in magnesium alloys. The present work examines the simultaneous effects of alumina nano-particle additions and severe plastic deformation by ECAP on the restoration phenomena, microstructure and thermal stability of the AZ31 magnesium alloy. 2. Experimental procedure 2.1. Materials and processing The AZ31 magnesium alloy (Mg–3 wt.% Al–1 wt.% Zn–0.2 wt.% Mn) was chosen as the matrix and Al2 O3 nano-particles, with an average diameter of 100 nm, as reinforcements. The melt was prepared from high purity (99.8%) magnesium, aluminum, zinc, and manganese, melted in a graphite crucible placed in an electrical resistance furnace under the Foseco Magrex 36 covering flux held at 1023 K. Three different weight percents of nano-particles (0.5, 1 and 2%) were added to the molten base material using an Mg-8 wt.% Al2 O3 nano-composite master alloy developed at the University of British Columbia (UBC). Based on the previous experience [13–15], the melt was held at this temperature for 10 min and then stirred mechanically for 2 min using a stainless steel rod to provide a homogeneous composition. It was then poured into a steel mold preheated to 473 K, using a tilt-casting system to minimize casting defects and turbulence of the melt. The cast billets of 44 mm diameter were annealed at 673 K for 1 h and extruded to 11 mm × 11 mm bars at 638 K at a speed of 2.0–2.5 mm/s. The ECAP billets having dimensions of 10 mm × 10 mm × 80 mm were machined from the extruded bars. ECAP was performed at 503 K through a solid die having an angle between the two channels of ˚ = 90◦ and an angle of = 20◦ at the outer arc of curvature at the point of intersection of the two channels, as shown in Fig. 1. Repetitive pressing of the same sample was carried out with route Bc, in which each sample was rotated 90◦ around its longitudinal axis between the passages. This configuration leads to an imposed strain of about 100% on each passage through the die. Samples were sprayed with MoS2 lubricant and pressed at a speed of 1 mm/s for 4 passes. 2.2. Textural studies Texture samples were finely polished and immersed several times in a diluted solution of HCl to eliminate mechanical effect of the preparation procedure. Texture measurements were performed
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Fig. 1. Schematic diagram of the ECAP system used in this research.
using the Schulz reflection method in a Philips X’Pert diffractometer furnished with a close Eulerian cradle. The intensity distributions of the (0 0 0 2) and (1 0 1¯ 0) pole figures were measured from the planes parallel to the pressing direction for both extruded and ECAPed monolithic and composite materials. The measurement was performed using Co k␣ radiation at 50 kV with the sample tilt angle ranging from 0◦ to 90◦ . 2.3. Microstructural studies Both of the extruded and ECAPed bars were sliced into 5mm-thick discs to study the microstructure of the cross-sections perpendicular to the pressing direction. Optical and scanning electron microscopy (SEM) were used to examine the grain size and reinforcement distribution. All of the samples were polished by 0.05-m Al2 O3 and then etched with a solution of 10 ml acetic acid, 4.2 g picric acid, 20 ml H2 O and 50 ml ethanol for 3 s at room temperature. The microstructures of the extruded and ECAPed samples were examined using a Leitz optical microscope and a CamScanMV2300 scanning electron microscope (SEM), respectively. The Clemex vision professional image analysis program was used to measure the grain size and distribution according to the ASTM E112 standard. To have more representative grain size values, at least 400 grains were considered for each condition. Different samples were employed and different locations on each sample were pictured and the number of grains on each micrograph was taken as the weighting factor in the calculation of the final grain size. 2.4. Thermal stability tests The thermal stability of the monolithic AZ31 and AZ31–1% Al2 O3 composite samples was investigated by the isochronal and isothermal annealing of the ECAPed materials and measuring the grain size evolution. Isochronal annealing was performed for 30 min at thirteen equally spaced temperatures in the range 473–740 K, while isothermal annealing was conducted for 300–30,000 s at five equally spaced temperatures in the same temperature range. All
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annealing tests were conducted in an air-circulated furnace with an accuracy of ±2 K. 2.5. Microhardness The ECAPed AZ31 and AZ31–1% Al2 O3 samples were annealed at 13 equally spaced temperatures in the range 473–740 K for 30 min to produce a range of grain sizes. Vickers microhardness (Hv ) was measured on the plane perpendicular to the pressing direction in a Wolpert microhardness tester with a load of 200 g for 15 s. Each reading was an average of at least five separate measurements taken at random places on the surface of the specimens. The acquired data were used in the construction of the Hall–Petch diagrams to elucidate the role of nano-particles in the strengthening of the composite material. 3. Results and discussion 3.1. Microstructure and texture evolution Optical images of the extruded bars in the cross-sections perpendicular to pressing direction are shown in Fig. 2. The initial grain size of the monolithic AZ31 reduced from 19.9 to 13.7 m, as the nano-particle content increased to 2 wt.%. It can be depicted that all extruded materials show normal grain size distributions, with standard deviations in the range 0.24–0.41 m. Due to the relatively high extrusion temperature of 638 K (T ∼ 0.7Tm ), dynamic recrystallization (DRX) could occur, resulting in the observed grain refinement. DRX is based on the grain boundary migration [25], and thus particles having pinning drag could postpone this phenomenon so that finer grains could be obtained in the presence of second phase particles. SEM images of the ECAPed materials, also shown in Fig. 2, indicate that significant refinement has occurred after severe plastic deformation. The degree of the achieved refinement has increased with the Al2 O3 content and the normal grain size distribution has become more skewed after the addition of 2% nano-particles. The standard deviation of the measured grain sizes for the ECAPed conditions was in the range 0.17–0.28 m. It can be seen that some fine-grained colonies are surrounded by coarser grains in some regions of the microstructure. One possibility is that the nucleation of DRX was halted due to the pinning effect of particles or, the nucleation took place but the grain growth following nucleation did not developed to produce equiaxed larger grains. Therefore, increasing the alumina content to 2% in the ECAPed AZ31 could postpone the restoration phenomena and raise the DRX temperature. An increase in the area of fine-grained regions in the AZ31–2% Al2 O3 composite could support the idea of postponed DRX during ECAP. In contrast to AZ31–2% Al2 O3 , for the other two composites a rather more normal distribution of equiaxed and uni-sized grains could be observed. The (0 0 0 2) and (1 0 1¯ 0) pole figures of the AZ31 and AZ31–2% Al2 O3 samples in the extruded and ECAPed conditions are shown in Fig. 3. In the as-extruded materials (Fig. 3a and b), it is evident that {0 0 0 2} basal planes are mostly oriented parallel to the extrusion direction. After ECAP, however, both basal and prismatic planes are inclined about 45◦ to the extrusion axis, as shown in Fig. 3c and d. The comparison of the {0 0 0 2} pole figures of AZ31–2% Al2 O3 with that of AZ31 shows that the particles have not affected the final positions of the texture. However, the intensity of the maximum orientations in both of the extruded and ECAPed materials increased by about 30% after Al2 O3 addition. The same is true for the prismatic planes, for which the intensity increases by about 44 and 60% for the extruded and ECAPed conditions, respectively. It is further demonstrated that ECAP weakens the basal poles and strengthens the prismatic poles. The observed variations in the
textures of the materials could be indicative of various deformation mechanisms operating in different conditions of the studied materials. 3.2. Annealing and thermal stability It was shown that the grain size distribution of the AZ31–2% Al2 O3 composite was rather unusual and the texture evolution for this composite exhibited drastic changes. Static annealing of this material was carried out at 573 K for 30 min. The annealed microstructure, shown in Fig. 4a, consists of a bimodal grain structure with enlarged grains in both fine and coarse regions. Although the grain size distribution is near-normal within each region (Fig. 4b), the overall grain distribution is clearly bimodal with two distinct grain size distributions, as depicted in Fig. 4c. This behavior is in agreement with those previously reported for nano- and ultrafine-grained magnesium alloys [14,26,27]. Due to the non-uniform grain structure of AZ31–2% Al2 O3 , the measurement of the grain size is rather complicated and weighted average grain sizes should be used, as proposed by Rajasekhara et al. [28]. This is in contrast to the case of monolithic AZ31, in which the grain size measurement of the uniform equiaxed microstructure is not affected by such a complication. Therefore, AZ31–1% Al2 O3 having a normal grain size distribution was selected for the thermal stability studies. Isochronal annealing was conducted for 1800 s on the AZ31 and AZ31–1% Al2 O3 samples, which are hereafter referred to as “base” and “composite”, respectively. Thirteen various test temperatures were selected in the range 473–740 K to produce a wide range of grain sizes. The variation of grain size with annealing temperature, shown in Fig. 5a, indicates that although grain growth occurs in both materials, the composite material has a lower grain size than the base AZ31 at all temperatures. Normal grain growth in singlephase alloys under isochronal heat treatment conditions is well described by the following equation [29]:
Q
Dn − D0n = Ct exp −
RT
(1)
where D and D0 are the current and initial values of the average grain size, respectively, Q is the activation energy of grain growth, R is the gas constant, T is the temperature, t is the annealing time, n is the grain growth exponent, and C is a kinetic constant. When a parabolic relationship for grain growth is assumed, the n-value equals 2 and the slope of the plot ln(D2 − D02 ) against 1/T would yield the grain growth activation energy for the samples annealed at different temperatures for a fixed annealing time of 1800 s. This was done for the measured grain sizes, assuming the respective initial grain sizes (D0 ) are 6.4 and 5.3 m for the monolithic and composite materials, as shown in Fig. 5b. It is evident that there are three distinct temperature ranges with different activation energies. The activation energies of 74.1, 19.2, and 94.4 kJ/mol are found for the base material in the low-, intermediate-, and high-temperature ranges, respectively. It is worth comparing these activation energies with those for the lattice self-diffusion (Qlsd = 135 kJ/mol) [30] and grain boundary diffusion in pure Mg (Qgb = 92 kJ/mol) [31]. In the low-temperature range of 490–530 K, the activation energy of 74.1 kJ/mole is less than Qgb . In the high-temperature range of 700–740 K, however, the activation energy of 94.4 kJ/mol is less than Qlsd but very close to Qgb . In the intermediate-temperature range of 530–700 K, the activation energy of 19.2 kJ/mol is significantly lower than both Qlsd and Qgb , corresponding to about 0.21Qgb . These comparisons can be used in the identification of grain growth mechanisms in different temperature intervals. The abnormally low activation energy for grain growth at intermediate annealing temperatures has also been reported for an ECAPed Al–Mg alloy by Wang et al. [32]. Their measured
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Fig. 2. Microstructure and grain size distribution of extruded and ECAPed AZ31 and composites containing 0.5, 1.0 and 2.0 wt.% of nano-particles. Images for the extruded materials are taken by optical microscopy and those for the ECAPed materials are by SEM.
Fig. 3. (0 0 0 2) and (1 0 1¯ 0) pole figures of (a) extruded AZ31, (b) extruded AZ31–2% Al2 O3 , (c) ECAPed AZ31, and (d) ECAPed AZ31–2% Al2 O3 .
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Fig. 5. (a) Variation of grain size with annealing temperature for the base and AZ31–1% Al2 O3 composite samples isochronally annealed for 1800 s, and (b) plot of ln(D2 − D02 ) against 1/T to determine activation energies according to Eq. (1).
Fig. 4. (a) Bimodal grain structure of AZ31 containing 2 wt.% Al2 O3 nano-particles subjected to 4 passes of ECAP and subsequently annealed at 573 K for 3600 s, (b) grain size distribution for fine- and coarse-grained regions, and (c) overall grain size distribution.
activation energy of 0.36Qgb was attributed to the unrecrystallized microstructure with non-equilibrium grain boundaries containing a large number of extrinsic dislocations, which may show higher atomic mobility compared to the equilibrium grain boundaries. The results of structural stability of an ECAPed AZ31 obtained by Kim and Kim [8] are indicative of a Q = 0.27Qgb in the intermediatetemperature range of 523–673 K. The possibility of grain growth under the unrecrystallized condition was, however, rejected due to the fact that the non-equilibrium boundaries were not present in the as-ECAPed (recrystallized) microstructure. Using the concept of diffusivity, they argued that progressive decrease in dislocation density by enhanced recovery with increasing temperature may be the cause of the observed anomaly in the activation energy. Our results on the activation energy of grain growth are in agreement with those obtained by Kim and Kim [8]. The slight observed difference could be attributed to the higher ECAP temperature in the present work, which affects both grain size and distribution and thus the grain growth process. The same three temperature zones exist for the activation energy of composite material, the data of which are superimposed on those of the AZ31 shown in Fig. 5b. It is clear that the activation energies of the composite material are higher than those of the base alloy, being 117.9, 22.7, and 88.6 kJ/mol for the high-, intermediate, and low-temperature ranges, respectively. These higher activation energies for grain growth in this condition could be attributed to the presence of alumina nano-particles. In the low- and intermediatetemperature ranges, recovery is the dominant thermally activated
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[33] demonstrated that dispersed particles can pin the grain boundaries, in which case the traditional kinetic theory of grain growth expressed by Eq. (1) could not predict the growth behavior. According to Burke’s model, it is assumed that drag force is independent of grain size and hence grain growth rate is not controlled by the instantaneous grain size, D. Instead, the decreasing difference between ultimate limiting grain size (Dm ) and the changing value of the instantaneous grain size (D) can control the growth rate, and hence Burke’s model may be expressed as: D0 − D + ln Dm
D − D m 0 Dm − D
=
k0 t 2 Dm
Q
exp −
RT
(2)
where D0 is the initial grain size, D is the instantaneous grain size, Dm is the limiting ultimate grain size for the particular annealing temperature, and k0 is a constant. By differentiating Eq. (2), the basic growth rate equation is as: dD =k dt
Fig. 6. (a) Variation of grain size with annealing time for the AZ31–1% Al2 O3 composite samples isothermally annealed in the range 520–740 K, and (b) plot of dD/dt against 1/D to determine the k parameters in Eq. (3).
phenomenon. In this process, dislocation movement is the fundamental assumption [25] and thus an increase in the activation energy for grain growth of the composite material can be attributed to fact that dislocation movement has been decelerated by the nano-particles. At higher temperatures, pinning of grain boundaries is the possible mechanism for the increase of grain growth activation energy. It seems that nano-particles have been quite successful in this process, as the activation energy of 94.4 kJ/mol for the base material has increased to 117.9 kJ/mol for the composite material. The obtained activation energies for grain growth are not valid for long time annealing, because in the beginning stages dislocations should be annihilated and some amounts of energy is spent for this purpose. Therefore, the exact activation energy for grain growth should be obtained through isothermal annealing at long times. This would highlight the effects of nano-particles on the grain growth behavior of the composite material. It is expected that this type of annealing would cause no significant change on the grain growth results of the AZ31 alloy because the traditional theory can satisfactorily predict the grain growth behavior of monotonic alloys. Accordingly, isothermal annealing was only performed on the composite samples. Because of the long annealing times involved, the interactions between dislocations and particles are ignored and only the pinning effects of nano-particles on the grain boundary movement are considered. The variation of grain size with annealing time at different constant temperatures, shown in Fig. 6a, indicates that the traditional parabolic trend is not observed for the composite material. Burke
1
D
−
1 Dm
(3)
where k = k0 exp(−Q/RT). By plotting dD/dt against 1/D for each temperature, linear relationships with correlation factors greater than 0.98 is obtained, as shown in Fig. 6b. The slopes of these lines show an increase with increasing annealing temperature. Since k has an Arrhenius type relationship with temperature, plotting ln(k) versus 1/T would yield the activation energy of grain growth considering drag force of particles, as exhibited in Fig. 7. The slopes of the plotted lines show two different activation energies at the high- and low-temperature regimes, respectively. For the low-temperature region, the low activation energy of 32.1 kJ/mol may correspond to the energy for the reordering of grain boundaries in the finegrained material [34]. The activation energy of 110.9 kJ/mol found for temperatures higher than 630 K, however, lies between that for grain boundary diffusion and lattice diffusion of magnesium. These results indicate that the observed difference between the activation energies for grain growth in the low- and high-temperature ranges is higher than that obtained from the traditional grain growth theory. This can be attributed to the fact that dislocation annihilation is assumed to affect the activation energy in the traditional grain growth theory, while in the Burke’s model it is considered to be insignificant due to the long annealing times. More importantly, the pinning effect of particles included in the Burke’s analysis has affected the observed behavior. It is worth noting that the higher grain growth activation energy in the high-temperature regime is
Fig. 7. Plot of ln(k) against 1/T for the AZ31–1% Al2 O3 composite samples to calculate activation energies for the low- and high-temperature regimes according to Burke’s model.
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movement. This resistance can be resulted from a single or a combination of various strengthening mechanisms such as precipitation strengthening, strain hardening, and solid solution strengthening. The obtained H0 value of 38.8 for the composite material is higher than 36.7 obtained for the base alloy. The enhanced friction strength is attributed to the presence of alumina nano-particles that hinder dislocation movement during plastic deformation caused by the penetration of the indenter. The slope of the H–P lines, K, is the locking parameter which measures the relative hardening contribution of the grain boundaries. The higher K-value of 60.9 for the composite material, compared to 51.7 obtained for the base material, is indicative of the positive role of the nano-particles in strengthening the grain boundaries by a pinning mechanism. 4. Conclusions The microstructural stability of the equal channel angularly pressed (ECAPed) AZ31 magnesium alloy and its composites, reinforced with 0.5, 1, and 2 wt.% of Al2 O3 nano-particles, were investigated and the following conclusions were made:
Fig. 8. (a) Microhardness of the specimens isochronally annealed for 1800 s at different temperatures, and (b) Hall–Petch relationship for AZ31 and AZ31–1% Al2 O3 .
consistent with the results on ultrafine-grained aluminum alloy, reported by Roy et al. [34]. 3.3. Hall–Petch relationship The Hall–Petch (H–P) relationship, which relates the hardness and strength to the grain size, has frequently been used to investigate the contribution of grain boundary strengthening as well as other strengthening mechanisms to the overall strength and hardness [28,35]. The general form of the H–P relationship for hardness is given by: Hv = H0 + KD−1/2
(4)
where Hv is the measured microhardness, H0 is the lattice friction hardness, D is the average grain size, and K is a constant. In the present study, the H–P approach would help clarifying the role of nano-particles in the achieved strengthening. Accordingly, the microhardness of the ECAPed monolithic and composite materials was measured and plotted against annealing temperature in Fig. 8a. As can be seen, at all annealing temperatures the microhardness of the composite is higher than the base alloy. In Fig. 8b, the classic H–P relationship is plotted for AZ31 and AZ31–1% Al2 O3 . It is obvious that both the friction hardness (H0 ) and the locking parameter (K) are higher for the particle-containing material. The friction stress or hardness, determined from the intercept of the H–P line, represents the overall resistance of the crystal lattice to dislocation
1. The introduction of 2 wt.% Al2 O3 nano-particles into the base AZ31 alloy reduced the grain size from 19.9 to 13.7 m in the extruded state, and from 6.4 to 4.3 m in the ECAPed condition. Microstructural analysis showed that the combination of ECAP and nano-particle addition results in a more homogeneous grain size distribution. These particles increased the intensity of the maximum orientations of the basal and prismatic planes in both extruded and ECAPed conditions. 2. The isochronal grain growth showed three distinct regions, depending on the temperature range investigated. In the lowand high-temperature ranges of 490–530 K and 700–740 K, the respective activation energies were 74.1 and 94.4 kJ/mol. These energies are respectively 0.8Qgb and 1.03Qgb , where Qgb is the activation for grain boundary diffusion. For the intermediatetemperature range of 530–700 K, however, the unexpectedly low activation energy of Q = 0.21Qgb was obtained. The higher activation energy for the composite material in all three regions reflects the role of nano-particles in pinning grain boundaries as compared to the monolithic material. 3. According to Burke’s model for isothermal annealing, two graingrowth regions with significantly different activation energies of 32.1 and 110.9 kJ/mol were identified for the low- and hightemperature regimes, respectively. The relatively low activation energy of 32.1 kJ/mol for grain growth in the low-temperature regime can be attributed to the reordering and readjustment of grain boundaries. 4. Hall–Petch relationship constructed for the monolithic and composite materials showed higher values of H0 and K for the composite material. This implies that nano-particles have effectively contributed to grain boundary pinning and also to interacting with dislocations to increase the friction stress. Acknowledgement The authors wish to gratefully acknowledge Professor Tom Troczynski at the University of British Columbia (UBC) who provided the Mg–Al2 O3 nano-composite master alloys used to fabricate the materials used in this study [36]. References [1] B.L. Mordike, T. Ebert, Mater. Sci. Eng. A 302 (2001) 37–45. [2] I.J. Polmear, Light Alloys; Aluminum, Magnesium, Titanium, 2nd ed., Edward Arnold Co., London, 1989, pp. 169–189. [3] K. Xia, J.T. Wang, X. Wu, Mater. Sci. Eng. A 410–411 (2005) 324–327. [4] L. Jin, D. Jin, D. Moa, Mater. Sci. Eng. A 423 (2006) 247–252.
Y. Radi, R. Mahmudi / Materials Science and Engineering A 527 (2010) 2764–2771 [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22]
B.Q. Han, T.G. Langdon, Mater. Sci. Eng. A 410–411 (2005) 435–438. R.B. Figueiredo, T.G. Langdon, Mater. Sci. Eng. A 503 (2009) 141–144. R.B. Figueiredo, T.G. Langdon, J. Mater. Sci. 43 (2008) 7366–7371. H.K. Kim, W.J. Kim, Mater. Sci. Eng. A 385 (2004) 300–308. H.K. Kim, J. Mater. Sci. Lett. 39 (2004) 7101–7103. A.A. Luo, Int. Mater. Rev. 49 (2004) 13–30. Y. Miyahara, K. Matsubara, Z. Horita, T.G. Langdon, Metall. Mater. Trans. 36A (2005) 1705–1711. X. Gao, S.M. Zhu, B.C. Muddle, J.F. Nie, Scripta Mater. 53 (2005) 1321–1326. F. Kabirian, R. Mahmudi, Adv. Eng. Mater. 11 (2009) 189–193. M. Masoudpanah, R. Mahmudi, Mater. Sci. Eng. A 526 (2009) 22–30. F. Kabirian, R. Mahmudi, Metall. Mater. Trans. 40A (2009) 2190–2201. X.L. Zhong, W.L.E. Wong, M. Gupta, Acta Mater. 55 (2007) 6338–6344. S.F. Hassan, M. Gupta, Metall. Mater. Trans. 36A (2005) 2253–2258. M. Habibnejad-Korayem, R. Mahmudi, W.J. Poole, Mater. Sci. Eng. A 519 (2009) 198–203. S. Seshan, M. Jayamathy, S.V. Kailas, T.S. Srivatsan, Mater. Sci. Eng. A 363 (2003) 345–351. H. Ferkel, B.L. Mordike, Mater. Sci. Eng. A 298 (2001) 193–199. H. Cao, J.Y. Min, S.D. Wu, A.P. Xian, J.K. Shang, Mater. Sci. Eng. A 431 (2006) 86–91. M. Ferry, N.E. Hamilton, F.J. Humphreys, Acta Mater. 53 (2005) 1097–1109.
2771
[23] N. Stanford, M. Barnett, Scripta Mater. 58 (2008) 179–182. [24] C.S. Goh, J. Wei, L.C. Lee, M. Gupta, Acta Mater. 55 (2007) 5115–5121. [25] F.J. Humphreys, M. Hatherly, Recrystallization and Related Annealing Phenomena, Elsevier, Oxford, United Kingdom, 1996. [26] Q. Yang, A.K. Ghosh, Acta Mater. 54 (2006) 5147–5158. [27] Y.M. Wang, E. Ma, Acta Mater. 52 (2004) 1699–1709. [28] S. Rajasekhara, P.J. Ferreira, L.P. Karjalainen, A. Kyrolainen, Metall. Mater. Trans. 38A (2007) 1202–1210. [29] P.A. Beck, J.C. Kremer, L.J. Demer, M.L. Holzworth, Trans. TMS-AIME 175 (1948) 372–400. [30] P.G. Shewman, Trans. AIME 206 (1956) 918–922. [31] H.J. Frost, M.E. Ashby, Deformation Mechanism Maps, Pergamon Press, Oxford, 1982. [32] J. Wang, Y. Iwahashi, Z. Horita, M. Furukawa, M. Nemoto, R.Z. Valiev, F.G. Langdon, Acta Mater. 44 (1996) 2973–2982. [33] J.E. Burke, Trans. TMS-AIME 180 (1949) 73–91. [34] I. Roy, M. Chauhan, E.J. Lavernia, F.A. Mohamed, Metall. Mater. Trans. 37A (2006) 721–730. [35] R. Mahmudi, Scripta Metall. Mater. 32 (1995) 781–786. [36] S. Vasisht, A. Tsvetkov, T. Troczynski, AUTO21 Conference 2006, Barrie, Ontario, 2006.