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Effect of long term annealing on the microstructure and tensile properties of hpdc AZ91 Mg alloy
Journal of Alloys and Compounds 467 (2009) 271–277
Effect of long term annealing on the microstructure and tensile properties of hpdc AZ91 Mg alloy D.G. Leo Prakash a,∗ , Doris Regener b a
b
Department of Engineering Science, University of Oxford, Parks Road, OX1 3PJ, Oxford, UK Institut f¨ur Werkstoff- und F¨ugetechnik, Otto-von-Guericke-Universit¨at Magdeburg, PF 4120, 39016, Magdeburg, Germany Received 16 November 2007; accepted 13 December 2007 Available online 23 December 2007
Abstract This work is aimed to obtain the effect of long term annealing (LTA) on the tensile properties and on the microstructure of a high pressure die cast (hpdc) AZ91 Mg alloy. Tensile tests were performed on as-cast and LTA for 1000 h at 150 and 200 ◦ C hpdc castings and the effect of LTA on the tensile properties such as yield strength (YS), ultimate tensile strength (UTS), ductility and fracture strain (FS) are confirmed. Additionally, explanations on micro failure mode of the material are presented to explain the influence of different micro features on failure of the material. The average size, area fraction and clustering tendency of m (massive) and c+d (continuous + discontinuous) Mg17 Al12 particles are quantified and their effects on LTA are obtained. The results confirm that the UTS, YS, ductility and FS are mainly influenced by the area fraction, size distribution and spatial arrangement of  particles. © 2007 Elsevier B.V. All rights reserved. Keywords: AZ91 alloy; High pressure die casting; Long term annealing; Microstructure; Micro–macro interactions
1. Introduction The past decade has witnessed significant growth in the consumption of magnesium and its alloys, mainly due to the automobile manufacturer’s requirement to reduce the weight of the vehicle in order to reduce the fuel utilization [1]. In addition to this, it is more useful to fulfil the aeronautical and electrical industry needs. The most common die cast magnesium alloy is AZ91. AZ91 accounts for more than 50% of all high pressure die castings [2] and this is due to its good strength, ductility and castability [3]. The typical microstructure of hpdc AZ91 alloy reveals a cored dendrite microstructure with primary ␣Mg solid solution, surrounded by a divorced eutectic region. The intermetallic -phase (Mg17 Al12 ) particles are embedded in this eutectic and grain boundary region. Additionally, the material contains processing defects like inclusions, pore bands, gas and shrinkage pores. These microstructural features have a complex geometry. Their locations and arrangements are often
∗
Corresponding author. Tel.: +44 1865 273108. E-mail address: [email protected] (D.G.L. Prakash).
0925-8388/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2007.12.046
non-uniform and usually strong spatial correlations exist. This multi length scale micro features cause multiple fracture micromechanisms, which, affect the fracture path and mechanical properties of this material. This also depends on the spatial arrangement, size distribution and shape of different microstructural feature. Automotive components of Mg alloys may be exposed to moderate temperatures in the range of 60–200 ◦ C. The thermal resistance of hpdc magnesium alloys is not investigated significantly. Therefore, it is important to know the impact of low temperature annealing on the mechanical properties of hpdc castings. Suman [4] performed LTA of AZ91 and AM60 castings at 120 ◦ C for 7, 30, 60 and 90 days and found that the YS increased (round tensile bar with 6.3 mm reduced section) by up to 40 MPa. He also stated that the YS decrease when the LTA time is more than 60 days. Bowles et al. [5] performed a similar study of the same materials for up to 208 days (120 ◦ C) and obtained an increase in YS of 3 MPa for 5 mm thick castings and 6 MPa for 2 mm thick castings. Same authors [5] also stated the decrease in ductility for increasing LTA time. Basner et al. [6] and Westengen et al. [7] studied LTA of AM seri-
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ous alloys and found increase in yield strength and decrease in ductility. Various authors reported the presence of discontinuous (d ) and continuous (c ) [4–6,8] precipitates in directly aged AZ91 casting. However, very limited information is documented on the quantitative characterization of  phase. NMR spectroscopy was used recently for the bulk quantification of  phase [5,9]. Aung and Zhou [10] quantitatively measured the volume fraction of bulk  phase in aged (for 8, 16 and 26 h in T4 and T6 conditions) AZ91D alloy at 200 ◦ C and obtained the effect of aging on corrosion behaviour. Present authors (Prakash et al. [11]) quantitatively characterized the different m and c+d phases from the microstructure using an image processing technique and discussed the microstructural developments during annealing. However, LTA condition-microstructure-mechanical property correlation is not sufficiently investigated. Therefore, this highlights the need for a systematic microstructural quantitative analysis to obtain the variation of micro quantities with respect to the LTA conditions of the castings and their relation to the tensile properties. Additionally, understanding the microscopic failure modes of the material is also equally important to predict the influence of the micro features on fracture behaviour, which is strongly related to the mechanical properties. Therefore, the aim of this work was to perform the same by quantitative characterization of microstructure and correlating the micro quantities with the tensile properties for different LTA conditions. 2. Experimental procedure The investigated AZ91 Mg alloy was produced by cold chamber hpdc machine GDK 200 in the form of plates with the dimension of 200 mm × 53 mm × 10 mm. The chemical composition of the produced alloy is presented in Table 1. All the castings were produced with the same conditions and more attention was not paid on the processing conditions as the scope of the present study was micro–macro correlations. These hpdc plates were further long term annealed (a condition designated as T5) in air at temperatures of 150 and 200 ◦ C for 1000 h. Tensile specimens of a cross-section of 10 mm × 10 mm with 50 mm gauge length were machined from the plates for the tensile analysis. Uniaxial tensile test was performed on these specimens at a constant strain rate of 10−4 s−1 in a computer controlled servohydraulic test machine at room temperature. The specimens from each tensile test were taken and polished by standard methods and further introduced to optical and scanning electron microscopy for microstructural characterization. A microstructural area of 25 mm2 (a quarter of the cross-section) was grabbed at 100× as continuous microstructural frames with an optical microscopy from the unetched cross-section. This microstructural frame was used to create the microstructural montage and it was further introduced to image processing to quantify the gas and shrinkage microporosity. The unetched crosssection was further etched (with a combination of picric acid (6 g), water (10 ml), acetic acid (5 ml) and ethanol (100 ml)) to reveal the  phase. The microstructural area of 1.86 mm2 was grabbed from the etched surface at 1000× for the quantification of m phase by image processing. The same location was viewed in SEM at 4000× and the microstructures were grabbed for the quantitative
characterization of c+d particle. The location of the m phase was computed as Cartesian coordinates by using the microstructures obtained for m quantification. These (X,Y) coordinates of m phase were used to identify the same phases in the SEM microstructures grabbed at 4000× to separate the m phases from the c+d phases. The separated m and c+d phases underwent microstructural quantification studies. A montage of 0.74 × 10−2 mm2 from the SEM (4000×) is used to quantify the c+d phase. The clustering tendency of these features was explained by comparing their spatial arrangement (nearest neighbour distance) from the present montage with the expected random arrangement. The values of clustering tendency below 1 indicate clustering, equal to 1; random, equal to 2.15; uniform distributions. The details of montage creation and the quantification procedure of microporosity and  phase are presented elsewhere [11,12,13]. In addition, in situ tensile analysis coupled with SEM was performed to understand the microscopic failure mode of the material.
3. Results and discussion This section explains the variations in tensile properties and microstructural quantities with respect to the LTA, failure modes of the material and the micro–macro correlations of the hpdc AZ91 magnesium alloy. 3.1. Macro properties The stress versus strain curves obtained from the tensile tests of the as-cast and LTA castings are shown in Fig. 1. Around 20 specimens were investigated in each case and few examples were presented in Fig. 1. No necking was observed in the deformed and failed specimens and the different cases showed a similar flow of deformation except the variations in hardening slope. A steep increase in hardening is observed in the LTA-150 ◦ C material compared to other cases and LTA-200 ◦ C admits for the lower hardening slope. The as-cast and LTA alloy mainly differ in the magnitudes of YS, UTS and FS. However, the difference in the FS is more predominant, particularly this could be of the variations in the arrangement of microporosity and the brittle  particles. These alloys exhibit no yield point and the 0.2% proof strength was taken as an indication of the yield point. Fig. 2a and b gives the variations in average YS (Fig. 2a) and UTS (Fig. 2b) of as-cast and LTA castings. These results clearly show the effect of LTA on the hpdc AZ91 alloy and there is a significant difference in the quantities with respect to the LTA temperature. The average YS of LTA-150 ◦ C (137.5 N/mm2 ) is high compared to the 200 ◦ C (111 N/mm2 ) case and the same of as-cast material (129.6 N/mm2 ) lie in-between these two cases.
Table 1 Chemical composition of hpdc AZ91 Mg alloy Alloy
% Al
% Mn
% Zn
% Si
% Cu
% Ni
% Fe
AZ91
9.3
0.12
0.79
0.02
0.0007
0.0006
0.0046
Fig. 1. Stress vs. strain results of the tensile experiments.
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Fig. 2. Average YS (a), UTS (b) and FS (c) of as-cast and LTA castings.
In the case of average UTS, as-cast material is having high value (179.3 N/mm2 ) compared to the LTA castings. As similar to YS, LTA-200 ◦ C (134.5 N/mm2 ) declare low UTS values compared to LTA-150 ◦ C (155 N/mm2 ) case. The increasing strain hardening behaviour provides better UTS for the as-cast material compared to LTA castings. The average FS values of as-cast and LTA castings are shown in Fig. 2c. It is high in the as-cast (1.39) material compared to annealed castings and LTA-150 ◦ C (0.34) hold the low value, and LTA-200 ◦ C is in-between (0.96) these two cases. It shows the strong effect of LTA on the FS compared to YS and UTS. LTA casting fails early compared to as-cast material indicating that the elongation of the LTA castings is low compared to the as-cast material. The average elongation of the as-cast, LTA-150 ◦ C and LTA-200 ◦ C castings is 3.17%, 0.53% and 1.41%, respectively. The present study concludes that the mechanical properties (strength and ductility) vary significantly with LTA conditions. YS of LTA-150 ◦ C have an increase of 5.8% compared to ascast and 19% compared to LTA-200 ◦ C. In case of UTS, as-cast material hold the maximum and it is an increase in 13.5% of LTA150 ◦ C and 25% of LTA-200 ◦ C. A significant drop in ductility is observed in LTA castings and the drop is 83% in LTA-150 ◦ C and 55% in LTA-200 ◦ C castings compared to as-cast materials. These results provide an expectation microstructural variation in different LTA castings and such changes would be responsible for the variations in mechanical properties with respect to the annealing conditions. 3.2. Microstructure and failure modes Metallographic examinations showed that the HPDC castings show a distinct fine grained surface layer and a coarse grained
Fig. 3. Typical as-cast microstructure of HPDC AZ91 magnesium alloy which shows the primary (␣), eutectic (␣) and brittle intermetallic  phases of the material.
interior. Fig. 3 is a typical microstructure of hpdc AZ91 alloy which shows ␣-primary, ␣-eutectic and brittle intermetallic  phases of the material. c and d precipitates were expected to precipitate in the LTA specimens. No microstructural changes except the size variation of m were detected in optical microscopy investigations using magnifications up to 1000×. However, these microstructural developments could be analysed by SEM microscopy at higher magnifications (above 2000×). Fig. 4 shows the SEM images of the LTA specimens treated at different temperatures. Fig. 4a and b shows the microstructural developments in LTA150 ◦ C and LTA-200 ◦ C castings, and Fig. 4c and d shows the clear appearance of continuous and discontinuous precipitation.
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Fig. 4. SEM images shows the microstructural changes due to long term annealing (a) 150 ◦ C (b) 200 ◦ C and appearance of (c) continuous and (d) discontinuous precipitation.
It shows that the d precipitates appears as irregular  lamellae/globular and the c precipitates are short aligned rod shaped precipitates. No symmetry was observed in the arrangement of d and c phases with respect to the arrangement of m . c and d precipitates are arranged in the grain boundary and in the supersaturated eutectic region. The microstructure of the specimen annealed at 150 ◦ C (Fig. 4a), both c and d precipitations are present and particularly the number density of c precipitations are very high and closely spaced; and around 80% of the matrix is invaded by both of these phases. Microstructure of the LTA-200 ◦ C (Fig. 4b) shows high number density of d precipitation and low number density of c precipitations arranged in the eutectic region. Starink [14] investigated the ageing of Mg–Al alloys and reported the c precipitation occurring throughout entire grains. However, the present examination shows that the c precipitation only occurs in localised regions and particularly the grain centres are relatively free of precipitates. The size of m phase is increasing with increasing annealing temperature. Additionally, the shape of m changes from irregular to regular for increasing annealing temperature. Fig. 5 illustrates an example of the size distribution of m (Fig. 5a) and c+d (Fig. 5b) particles in LTA materials. The result
shows that the number density of c+d particles in the LTA150 ◦ C casting is significantly higher than LTA-200 ◦ C case. This explains the high number of c+d (particularly c ) phase development in the material annealed at 150 ◦ C compared to LTA-200 ◦ C case. The low number density of bigger c+d particles (see Fig. 5b) and high number density of bigger m particles (see Fig. 5a) in the LTA-200 ◦ C casting compared to LTA-150 ◦ C casting confirm the agglomeration of c and d phases with neighbouring m particles. An example of area fraction of m and c+d phases in LTA castings are showed in Fig. 5c. The results show that the area fraction of m phase is very high in the annealed specimens compared to as-cast specimens and it increases as the annealing temperature increases, particularly it is very high in the case of 200 ◦ C due to high diffusion of Al atoms towards the m phase and agglomeration of neighbouring particles. These results explain that the nucleation rate of c and/or d phase increases as annealing temperature increases until a critical annealing temperature, which is around 150 ◦ C and there is a domination of agglomeration of neighbouring c+d particles beyond this critical annealing temperature. A similar trend is stated by Celotto [15] where the number of c precipitates per unit volume in aged AZ91 castings decreases from 1.0 × 1012 /mm3 at 70 ◦ C to 1.5 × 109 /mm3 at 300 ◦ C. Celotto
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275
Fig. 5. Size distribution (a and b) and area fraction (c) of m and c+d particles.
[15] also stated that there is a dramatic increase in the number density of precipitates at temperatures lower than 150 ◦ C. The results also show that the discontinuous and continuous precipitation reactions occur competitively over a range of annealing temperatures. The lack of slip systems of hpdc magnesium alloys causes less deformation due to the reduction of dislocation motion between the neighbouring grains, which leads to a grain boundary cleavage. It is concluded that AZ91 follows intergranular brittle failure and, cleavage and quasi-cleavage are the most common fracture modes [16]. The crack initiation and growth from shrinkage pores, damage of brittle  particles and grain boundary cleavage are found to be the primary failure modes of the present material from the in situ tensile analysis [16]. Examples of these micro mechanisms are shown in Fig. 6. The predominant intergranular failure of hpdc AZ91 alloy is also due to the arrangement of shrinkage pores and  particles in the intergranular region [16]. The effect of inclusion on fracture was found to be negligible due to its scarce presence in the investigated material. Regener et al. [17] also observed the  particle cracking in Mg alloys during the in situ tensile analysis. In addition, Yoo et al. [18] documented that damages in massive  phase in the grain boundaries supports the grain boundary fracture. The different failure modes of hpdc AZ91 alloy under tensile loading and the extent effect of shrinkage pores and  particles on failure are explained in the previous contribution of authors [16]. The microstructural damages or cracks have strong effect on the macro behaviour of the material [19–21]. The above explanations notice the greater influence of pores (1–300 m) and  particles (1–4 m) on fracture behaviour
of hpdc AZ91 alloy. Particularly, the size, shape and arrangement of these inhomogenities are the important parameters that influence the mechanical properties by altering the fracture behaviour. Complete attention was only paid to  particle as it is the only microstructural variable in the investigated material. Detailed quantification of these inhomogenities was performed to understand the interaction between the microstructure and the mechanical properties, which is presented in the next subdivision. 3.3. Micro quantities and their effects The microstructural quantification results of as-cast and LTA castings and their corresponding tensile properties are shown in Table 2. It confirms the significant variation in the area fraction, average size and clustering tendency of m and c+d particles with respect to the LTA conditions of the casting. The average size, variation in area fraction and clustering tendency of microporosity, and average grain size of the analysed material is 31.2 m2 , 1–1.8%, 0.59–0.7 and 7.72 m, respectively. The quantification results confirm that the area fraction of  phase in as-cast material is low compared to LTA castings and it increases with increasing annealing temperature mainly due to the nucleation, diffusion and agglomeration reactions as explained. The area fraction of c+d particles in LTA-150 ◦ C castings is around two times of LTA-200 ◦ C castings. The high rate of fine and closely spaced c particle nucleation in LTA150 ◦ C provides low average size, high clustering tendency and high area fraction for LTA-150 ◦ C castings. Compared to LTA-150 ◦ C case, LTA-200 ◦ C afford high average size, low
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Fig. 6. Microscopic failure modes (a) crack initiation and growth from shrinkage pore, (b) damage of  particles and (c) grain boundary fracture. Table 2 Comparison of micro quantities and macro properties Casting type
AC LTA 150 ◦ C 200 ◦ C
Area fraction (%)
Average size (m2 )
Clustering tendency
m
c+d
m
c+d
m
c+d
–
1.98
–
0.63–0.70
5.7–7.4 2.1–4.4
3.35 5.73
0.045 0.09
0.99–1.07 1–1.09
2.9–5 6.3–8.1 12.5–14.9
clustering tendency and low area fraction of c+d particles due to agglomeration process. Besides, high average size, low clustering tendency (random arrangement) and high area fraction of m particles in LTA castings compared to as-cast material are observed due to the diffusion and agglomeration reactions. The increase in m compared to as-cast material is around two times in LTA-150 ◦ C and around three times in LTA-200 ◦ C. The similar trend is obtained in the average size of m particles. With respect to the changes of micro quantities a significant difference in UTS, YS and FS between as-cast and LTA castings is observed. Better YS of LTA-150 ◦ C castings compared to other
UTS (N/mm2 )
YS (N/mm2 )
–
165–190
122–135
0.38–0.43 0.84–0.9
143–166 128–147
131–144 106–121
FS (%)
1–1.9 0.28–0.69 0.71–1.4
cases is observed due to the increased nucleation of c particles, which act as a strengthener of this material. This is due to the fine c precipitates in the eutectic and grain boundary region resulting in the arrest of dislocation motion. The high number density of fine and low number density of bigger m particles (see Fig. 5a) in as-cast material compared to LTA-200 ◦ C would be the reason for better YS in as-cast material. This confirms that the strengthening effect of  particles decreases when its size increases. However, different trend is noticed in the case of UTS where as-cast material have the better property due to the high strain
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hardening rate compared to LTA castings. This is because of easy dislocation motion in the as-cast material due to the low number density and area fraction of  particles. Free from c+d particles is an additional supportive factor in this case. The high number density and area fraction of  particles (c+d + m ) in LTA castings compared to as-cast material cause a decrease in UTS due to the drop in strain hardening rate. However, the high probability of arresting the dislocation movement in LTA-150 ◦ C compared to LTA-200 ◦ C castings provides better UTS for LTA150 ◦ C by a steep increase in hardening slope. Besides, high probability of damage in bigger brittle m particles of LTA200 ◦ C castings softens the material, which decreases the UTS. The FS values shows that the LTA-150 ◦ C case fails early compared to other cases and the same is better in as-cast material. This depends on the fracture behaviour of the material and the important attribute is the grain boundary region as the material fails by intergranular brittle failure due to the lack of dislocation motion from a grain to other. It is already confirmed that the arrangement of  particles in the grain boundary region also speed up the failure process [16]. The arrangement of brittle m and c+d particles in the grain boundary and in the eutectic region weakens the grain boundary region and has significant effect on this as it increases the area fraction of grain boundary region, which inhibits transmission of slip across the boundaries. These particles support the crack initiation and propagation in the grain boundary region due to the development of high stress concentration in the particles during deformation and this is severe in the case of LTA-150 ◦ C as the grain boundary region is almost occupied by the m and c+d particles. As-cast material holds better FS values compared to LTA-200 ◦ C case due to the low size, area fraction and number density of the  particles. In the case of LTA-200 ◦ C, the probability of damage is high as it is having high number density of bigger m particles, which also speeds up the failure process. These explanations also confirm the results of ductility of different castings. 4. Summery and conclusions The present study concludes that the mechanical properties (strength and ductility) vary significantly with LTA conditions. It also confirms the significant variation in the area fraction, average size and clustering tendency of m and c+d particles with respect to the LTA conditions of the casting and their effect on LTA conditions. The area fraction of bulk  phase in as-cast material is low compared to LTA castings and it increases with increasing annealing temperature due to the nucleation of c and d particles and diffusion of Al atoms towards  particles during LTA. As-cast material holds better UTS and ductility compared to the LTA castings. LTA-150 ◦ C offers the vigorous
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nucleation of fine closely spaced c+d particles, which provides better YS compared to as-cast and LTA-200 ◦ C castings and c particles act as a strengthener of annealed hpdc AZ91 alloy. The agglomeration of different  particles takes place in the LTA200 ◦ C castings and provides poor YS and UTS compared to LTA-150 ◦ C castings. The high strain hardening rate provides better ductility for LTA-200 ◦ C compared to the LTA-150 ◦ C castings. The increase in size and area fraction of  particles in the grain boundary region supports the failure process and leads to early failure in LTA castings. References [1] D. Magers, J. Willekens, in: B.L. Mordike, K.U. Kainer (Eds.), Magnesium Alloys and Their Applications, Wolfsburg, Germany, 1998, pp. 105– 112. [2] S. Schumann, F. Friedrich, in: B.L. Mordike, K.U. Kainer (Eds.), Magnesium Alloys and Their Applications, Wolfsburg, Germany, 1998, pp. 3–13. [3] D. Brungs, in: B.L. Mordike, K.U. Kainer (Eds.), Magnesium Alloys and Their Applications, Wolfsburg, Germany, 1998, pp. 113–117. [4] C. Suman, SAE International Congress and Exhibition, SAE Technical Paper Series, Paper No. 890207 (1989). [5] A.L. Bowles, T.J. Bastow, C.J. Davidson, J.R. Griffiths, P.D.D. Rodrigo, Magnesium Technology, The Minerals Metals and Materials Society, Nashville, USA, 2000, pp. 295–300. [6] T.G. Basner, M. Evans, D.J. Sakkinen, Magnesium Properties and Applications for Automobiles, Society of Automotive Engineers, Detroit, USA, 1993, pp. 59–64. [7] H. Westengen, L.-Y. Wei, T. Aune, D. Albright, in: B.L. Mordike, K.U. Kainer (Eds.), Magnesium Alloys and Their Applications, Wolfsburg, Germany, 1998, pp. 209–214. [8] D.J. Sakkinen, M. Evans, Proceedings of the 17th International Die Casting Congress and Exposition, Cleveland USA, 1993, pp. 305–313. [9] G. Song, A.L. Bowles, D.H. StJohn, Mater. Sci. Eng. A 366 (2004) 74– 86. [10] N.N. Aung, W. Zhou, J. Appl. Electrochem. 32 (2002) 1397–1401. [11] D.G.L. Prakash, D. Regener, W.J.J. Vorster, Effect of long term annealing on the microstructure of hpdc AZ91 Mg alloy: a quantitative analysis by image processing, Comput. Mater. Sci., submitted for publication. [12] D.G.L. Prakash, D. Regener, Prakt. Metallogr. 42 (2005) 555–575. [13] D.G.L. Prakash, D. Regener, J. Alloys Compd 461 (2008) 139–146. [14] M.J. Starink, J. Mater. Sci. 32 (1997) 4061–4070. [15] S. Celotto, Acta Mater. 48 (2000) 1775–1787. [16] .G.L. Prakash, D. Regener, W.J.J. Vorster, Sci. Eng. A 488 (2008) 303–310. [17] D. Regener, E. Schick, I. Wagner, H. Heyse, Materialwiss. Werkst. 30 (1999) 525–532. [18] M.S. Yoo, Y.C. Kim, S. Ahn, N.J. Kim, Mater. Sci. Forum 419–422 (2003) 419–424. [19] A.K. Vasudevan, O. Richmond, F. Zok, J.D. Embury, Mater. Sci. Eng. A. 107 (1989) 63–69. [20] Y. Brechet, J. Newell, S. Tao, J.D. Embury, Scrip. Metall. et Mater. 28 (1993) 47–51. [21] J.D. S Embury, Procedings of Twelfth Risø International Symposium on Materials Science, Risø National Laboratory, Roskilde, Denmark, 1991, pp. 317–322.
Effect of long term annealing on the microstructure and tensile properties of hpdc AZ91 Mg alloy D.G. Leo Prakash a,∗ , Doris Regener b a
b
Department of Engineering Science, University of Oxford, Parks Road, OX1 3PJ, Oxford, UK Institut f¨ur Werkstoff- und F¨ugetechnik, Otto-von-Guericke-Universit¨at Magdeburg, PF 4120, 39016, Magdeburg, Germany Received 16 November 2007; accepted 13 December 2007 Available online 23 December 2007
Abstract This work is aimed to obtain the effect of long term annealing (LTA) on the tensile properties and on the microstructure of a high pressure die cast (hpdc) AZ91 Mg alloy. Tensile tests were performed on as-cast and LTA for 1000 h at 150 and 200 ◦ C hpdc castings and the effect of LTA on the tensile properties such as yield strength (YS), ultimate tensile strength (UTS), ductility and fracture strain (FS) are confirmed. Additionally, explanations on micro failure mode of the material are presented to explain the influence of different micro features on failure of the material. The average size, area fraction and clustering tendency of m (massive) and c+d (continuous + discontinuous) Mg17 Al12 particles are quantified and their effects on LTA are obtained. The results confirm that the UTS, YS, ductility and FS are mainly influenced by the area fraction, size distribution and spatial arrangement of  particles. © 2007 Elsevier B.V. All rights reserved. Keywords: AZ91 alloy; High pressure die casting; Long term annealing; Microstructure; Micro–macro interactions
1. Introduction The past decade has witnessed significant growth in the consumption of magnesium and its alloys, mainly due to the automobile manufacturer’s requirement to reduce the weight of the vehicle in order to reduce the fuel utilization [1]. In addition to this, it is more useful to fulfil the aeronautical and electrical industry needs. The most common die cast magnesium alloy is AZ91. AZ91 accounts for more than 50% of all high pressure die castings [2] and this is due to its good strength, ductility and castability [3]. The typical microstructure of hpdc AZ91 alloy reveals a cored dendrite microstructure with primary ␣Mg solid solution, surrounded by a divorced eutectic region. The intermetallic -phase (Mg17 Al12 ) particles are embedded in this eutectic and grain boundary region. Additionally, the material contains processing defects like inclusions, pore bands, gas and shrinkage pores. These microstructural features have a complex geometry. Their locations and arrangements are often
∗
Corresponding author. Tel.: +44 1865 273108. E-mail address: [email protected] (D.G.L. Prakash).
0925-8388/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2007.12.046
non-uniform and usually strong spatial correlations exist. This multi length scale micro features cause multiple fracture micromechanisms, which, affect the fracture path and mechanical properties of this material. This also depends on the spatial arrangement, size distribution and shape of different microstructural feature. Automotive components of Mg alloys may be exposed to moderate temperatures in the range of 60–200 ◦ C. The thermal resistance of hpdc magnesium alloys is not investigated significantly. Therefore, it is important to know the impact of low temperature annealing on the mechanical properties of hpdc castings. Suman [4] performed LTA of AZ91 and AM60 castings at 120 ◦ C for 7, 30, 60 and 90 days and found that the YS increased (round tensile bar with 6.3 mm reduced section) by up to 40 MPa. He also stated that the YS decrease when the LTA time is more than 60 days. Bowles et al. [5] performed a similar study of the same materials for up to 208 days (120 ◦ C) and obtained an increase in YS of 3 MPa for 5 mm thick castings and 6 MPa for 2 mm thick castings. Same authors [5] also stated the decrease in ductility for increasing LTA time. Basner et al. [6] and Westengen et al. [7] studied LTA of AM seri-
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ous alloys and found increase in yield strength and decrease in ductility. Various authors reported the presence of discontinuous (d ) and continuous (c ) [4–6,8] precipitates in directly aged AZ91 casting. However, very limited information is documented on the quantitative characterization of  phase. NMR spectroscopy was used recently for the bulk quantification of  phase [5,9]. Aung and Zhou [10] quantitatively measured the volume fraction of bulk  phase in aged (for 8, 16 and 26 h in T4 and T6 conditions) AZ91D alloy at 200 ◦ C and obtained the effect of aging on corrosion behaviour. Present authors (Prakash et al. [11]) quantitatively characterized the different m and c+d phases from the microstructure using an image processing technique and discussed the microstructural developments during annealing. However, LTA condition-microstructure-mechanical property correlation is not sufficiently investigated. Therefore, this highlights the need for a systematic microstructural quantitative analysis to obtain the variation of micro quantities with respect to the LTA conditions of the castings and their relation to the tensile properties. Additionally, understanding the microscopic failure modes of the material is also equally important to predict the influence of the micro features on fracture behaviour, which is strongly related to the mechanical properties. Therefore, the aim of this work was to perform the same by quantitative characterization of microstructure and correlating the micro quantities with the tensile properties for different LTA conditions. 2. Experimental procedure The investigated AZ91 Mg alloy was produced by cold chamber hpdc machine GDK 200 in the form of plates with the dimension of 200 mm × 53 mm × 10 mm. The chemical composition of the produced alloy is presented in Table 1. All the castings were produced with the same conditions and more attention was not paid on the processing conditions as the scope of the present study was micro–macro correlations. These hpdc plates were further long term annealed (a condition designated as T5) in air at temperatures of 150 and 200 ◦ C for 1000 h. Tensile specimens of a cross-section of 10 mm × 10 mm with 50 mm gauge length were machined from the plates for the tensile analysis. Uniaxial tensile test was performed on these specimens at a constant strain rate of 10−4 s−1 in a computer controlled servohydraulic test machine at room temperature. The specimens from each tensile test were taken and polished by standard methods and further introduced to optical and scanning electron microscopy for microstructural characterization. A microstructural area of 25 mm2 (a quarter of the cross-section) was grabbed at 100× as continuous microstructural frames with an optical microscopy from the unetched cross-section. This microstructural frame was used to create the microstructural montage and it was further introduced to image processing to quantify the gas and shrinkage microporosity. The unetched crosssection was further etched (with a combination of picric acid (6 g), water (10 ml), acetic acid (5 ml) and ethanol (100 ml)) to reveal the  phase. The microstructural area of 1.86 mm2 was grabbed from the etched surface at 1000× for the quantification of m phase by image processing. The same location was viewed in SEM at 4000× and the microstructures were grabbed for the quantitative
characterization of c+d particle. The location of the m phase was computed as Cartesian coordinates by using the microstructures obtained for m quantification. These (X,Y) coordinates of m phase were used to identify the same phases in the SEM microstructures grabbed at 4000× to separate the m phases from the c+d phases. The separated m and c+d phases underwent microstructural quantification studies. A montage of 0.74 × 10−2 mm2 from the SEM (4000×) is used to quantify the c+d phase. The clustering tendency of these features was explained by comparing their spatial arrangement (nearest neighbour distance) from the present montage with the expected random arrangement. The values of clustering tendency below 1 indicate clustering, equal to 1; random, equal to 2.15; uniform distributions. The details of montage creation and the quantification procedure of microporosity and  phase are presented elsewhere [11,12,13]. In addition, in situ tensile analysis coupled with SEM was performed to understand the microscopic failure mode of the material.
3. Results and discussion This section explains the variations in tensile properties and microstructural quantities with respect to the LTA, failure modes of the material and the micro–macro correlations of the hpdc AZ91 magnesium alloy. 3.1. Macro properties The stress versus strain curves obtained from the tensile tests of the as-cast and LTA castings are shown in Fig. 1. Around 20 specimens were investigated in each case and few examples were presented in Fig. 1. No necking was observed in the deformed and failed specimens and the different cases showed a similar flow of deformation except the variations in hardening slope. A steep increase in hardening is observed in the LTA-150 ◦ C material compared to other cases and LTA-200 ◦ C admits for the lower hardening slope. The as-cast and LTA alloy mainly differ in the magnitudes of YS, UTS and FS. However, the difference in the FS is more predominant, particularly this could be of the variations in the arrangement of microporosity and the brittle  particles. These alloys exhibit no yield point and the 0.2% proof strength was taken as an indication of the yield point. Fig. 2a and b gives the variations in average YS (Fig. 2a) and UTS (Fig. 2b) of as-cast and LTA castings. These results clearly show the effect of LTA on the hpdc AZ91 alloy and there is a significant difference in the quantities with respect to the LTA temperature. The average YS of LTA-150 ◦ C (137.5 N/mm2 ) is high compared to the 200 ◦ C (111 N/mm2 ) case and the same of as-cast material (129.6 N/mm2 ) lie in-between these two cases.
Table 1 Chemical composition of hpdc AZ91 Mg alloy Alloy
% Al
% Mn
% Zn
% Si
% Cu
% Ni
% Fe
AZ91
9.3
0.12
0.79
0.02
0.0007
0.0006
0.0046
Fig. 1. Stress vs. strain results of the tensile experiments.
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Fig. 2. Average YS (a), UTS (b) and FS (c) of as-cast and LTA castings.
In the case of average UTS, as-cast material is having high value (179.3 N/mm2 ) compared to the LTA castings. As similar to YS, LTA-200 ◦ C (134.5 N/mm2 ) declare low UTS values compared to LTA-150 ◦ C (155 N/mm2 ) case. The increasing strain hardening behaviour provides better UTS for the as-cast material compared to LTA castings. The average FS values of as-cast and LTA castings are shown in Fig. 2c. It is high in the as-cast (1.39) material compared to annealed castings and LTA-150 ◦ C (0.34) hold the low value, and LTA-200 ◦ C is in-between (0.96) these two cases. It shows the strong effect of LTA on the FS compared to YS and UTS. LTA casting fails early compared to as-cast material indicating that the elongation of the LTA castings is low compared to the as-cast material. The average elongation of the as-cast, LTA-150 ◦ C and LTA-200 ◦ C castings is 3.17%, 0.53% and 1.41%, respectively. The present study concludes that the mechanical properties (strength and ductility) vary significantly with LTA conditions. YS of LTA-150 ◦ C have an increase of 5.8% compared to ascast and 19% compared to LTA-200 ◦ C. In case of UTS, as-cast material hold the maximum and it is an increase in 13.5% of LTA150 ◦ C and 25% of LTA-200 ◦ C. A significant drop in ductility is observed in LTA castings and the drop is 83% in LTA-150 ◦ C and 55% in LTA-200 ◦ C castings compared to as-cast materials. These results provide an expectation microstructural variation in different LTA castings and such changes would be responsible for the variations in mechanical properties with respect to the annealing conditions. 3.2. Microstructure and failure modes Metallographic examinations showed that the HPDC castings show a distinct fine grained surface layer and a coarse grained
Fig. 3. Typical as-cast microstructure of HPDC AZ91 magnesium alloy which shows the primary (␣), eutectic (␣) and brittle intermetallic  phases of the material.
interior. Fig. 3 is a typical microstructure of hpdc AZ91 alloy which shows ␣-primary, ␣-eutectic and brittle intermetallic  phases of the material. c and d precipitates were expected to precipitate in the LTA specimens. No microstructural changes except the size variation of m were detected in optical microscopy investigations using magnifications up to 1000×. However, these microstructural developments could be analysed by SEM microscopy at higher magnifications (above 2000×). Fig. 4 shows the SEM images of the LTA specimens treated at different temperatures. Fig. 4a and b shows the microstructural developments in LTA150 ◦ C and LTA-200 ◦ C castings, and Fig. 4c and d shows the clear appearance of continuous and discontinuous precipitation.
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Fig. 4. SEM images shows the microstructural changes due to long term annealing (a) 150 ◦ C (b) 200 ◦ C and appearance of (c) continuous and (d) discontinuous precipitation.
It shows that the d precipitates appears as irregular  lamellae/globular and the c precipitates are short aligned rod shaped precipitates. No symmetry was observed in the arrangement of d and c phases with respect to the arrangement of m . c and d precipitates are arranged in the grain boundary and in the supersaturated eutectic region. The microstructure of the specimen annealed at 150 ◦ C (Fig. 4a), both c and d precipitations are present and particularly the number density of c precipitations are very high and closely spaced; and around 80% of the matrix is invaded by both of these phases. Microstructure of the LTA-200 ◦ C (Fig. 4b) shows high number density of d precipitation and low number density of c precipitations arranged in the eutectic region. Starink [14] investigated the ageing of Mg–Al alloys and reported the c precipitation occurring throughout entire grains. However, the present examination shows that the c precipitation only occurs in localised regions and particularly the grain centres are relatively free of precipitates. The size of m phase is increasing with increasing annealing temperature. Additionally, the shape of m changes from irregular to regular for increasing annealing temperature. Fig. 5 illustrates an example of the size distribution of m (Fig. 5a) and c+d (Fig. 5b) particles in LTA materials. The result
shows that the number density of c+d particles in the LTA150 ◦ C casting is significantly higher than LTA-200 ◦ C case. This explains the high number of c+d (particularly c ) phase development in the material annealed at 150 ◦ C compared to LTA-200 ◦ C case. The low number density of bigger c+d particles (see Fig. 5b) and high number density of bigger m particles (see Fig. 5a) in the LTA-200 ◦ C casting compared to LTA-150 ◦ C casting confirm the agglomeration of c and d phases with neighbouring m particles. An example of area fraction of m and c+d phases in LTA castings are showed in Fig. 5c. The results show that the area fraction of m phase is very high in the annealed specimens compared to as-cast specimens and it increases as the annealing temperature increases, particularly it is very high in the case of 200 ◦ C due to high diffusion of Al atoms towards the m phase and agglomeration of neighbouring particles. These results explain that the nucleation rate of c and/or d phase increases as annealing temperature increases until a critical annealing temperature, which is around 150 ◦ C and there is a domination of agglomeration of neighbouring c+d particles beyond this critical annealing temperature. A similar trend is stated by Celotto [15] where the number of c precipitates per unit volume in aged AZ91 castings decreases from 1.0 × 1012 /mm3 at 70 ◦ C to 1.5 × 109 /mm3 at 300 ◦ C. Celotto
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Fig. 5. Size distribution (a and b) and area fraction (c) of m and c+d particles.
[15] also stated that there is a dramatic increase in the number density of precipitates at temperatures lower than 150 ◦ C. The results also show that the discontinuous and continuous precipitation reactions occur competitively over a range of annealing temperatures. The lack of slip systems of hpdc magnesium alloys causes less deformation due to the reduction of dislocation motion between the neighbouring grains, which leads to a grain boundary cleavage. It is concluded that AZ91 follows intergranular brittle failure and, cleavage and quasi-cleavage are the most common fracture modes [16]. The crack initiation and growth from shrinkage pores, damage of brittle  particles and grain boundary cleavage are found to be the primary failure modes of the present material from the in situ tensile analysis [16]. Examples of these micro mechanisms are shown in Fig. 6. The predominant intergranular failure of hpdc AZ91 alloy is also due to the arrangement of shrinkage pores and  particles in the intergranular region [16]. The effect of inclusion on fracture was found to be negligible due to its scarce presence in the investigated material. Regener et al. [17] also observed the  particle cracking in Mg alloys during the in situ tensile analysis. In addition, Yoo et al. [18] documented that damages in massive  phase in the grain boundaries supports the grain boundary fracture. The different failure modes of hpdc AZ91 alloy under tensile loading and the extent effect of shrinkage pores and  particles on failure are explained in the previous contribution of authors [16]. The microstructural damages or cracks have strong effect on the macro behaviour of the material [19–21]. The above explanations notice the greater influence of pores (1–300 m) and  particles (1–4 m) on fracture behaviour
of hpdc AZ91 alloy. Particularly, the size, shape and arrangement of these inhomogenities are the important parameters that influence the mechanical properties by altering the fracture behaviour. Complete attention was only paid to  particle as it is the only microstructural variable in the investigated material. Detailed quantification of these inhomogenities was performed to understand the interaction between the microstructure and the mechanical properties, which is presented in the next subdivision. 3.3. Micro quantities and their effects The microstructural quantification results of as-cast and LTA castings and their corresponding tensile properties are shown in Table 2. It confirms the significant variation in the area fraction, average size and clustering tendency of m and c+d particles with respect to the LTA conditions of the casting. The average size, variation in area fraction and clustering tendency of microporosity, and average grain size of the analysed material is 31.2 m2 , 1–1.8%, 0.59–0.7 and 7.72 m, respectively. The quantification results confirm that the area fraction of  phase in as-cast material is low compared to LTA castings and it increases with increasing annealing temperature mainly due to the nucleation, diffusion and agglomeration reactions as explained. The area fraction of c+d particles in LTA-150 ◦ C castings is around two times of LTA-200 ◦ C castings. The high rate of fine and closely spaced c particle nucleation in LTA150 ◦ C provides low average size, high clustering tendency and high area fraction for LTA-150 ◦ C castings. Compared to LTA-150 ◦ C case, LTA-200 ◦ C afford high average size, low
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Fig. 6. Microscopic failure modes (a) crack initiation and growth from shrinkage pore, (b) damage of  particles and (c) grain boundary fracture. Table 2 Comparison of micro quantities and macro properties Casting type
AC LTA 150 ◦ C 200 ◦ C
Area fraction (%)
Average size (m2 )
Clustering tendency
m
c+d
m
c+d
m
c+d
–
1.98
–
0.63–0.70
5.7–7.4 2.1–4.4
3.35 5.73
0.045 0.09
0.99–1.07 1–1.09
2.9–5 6.3–8.1 12.5–14.9
clustering tendency and low area fraction of c+d particles due to agglomeration process. Besides, high average size, low clustering tendency (random arrangement) and high area fraction of m particles in LTA castings compared to as-cast material are observed due to the diffusion and agglomeration reactions. The increase in m compared to as-cast material is around two times in LTA-150 ◦ C and around three times in LTA-200 ◦ C. The similar trend is obtained in the average size of m particles. With respect to the changes of micro quantities a significant difference in UTS, YS and FS between as-cast and LTA castings is observed. Better YS of LTA-150 ◦ C castings compared to other
UTS (N/mm2 )
YS (N/mm2 )
–
165–190
122–135
0.38–0.43 0.84–0.9
143–166 128–147
131–144 106–121
FS (%)
1–1.9 0.28–0.69 0.71–1.4
cases is observed due to the increased nucleation of c particles, which act as a strengthener of this material. This is due to the fine c precipitates in the eutectic and grain boundary region resulting in the arrest of dislocation motion. The high number density of fine and low number density of bigger m particles (see Fig. 5a) in as-cast material compared to LTA-200 ◦ C would be the reason for better YS in as-cast material. This confirms that the strengthening effect of  particles decreases when its size increases. However, different trend is noticed in the case of UTS where as-cast material have the better property due to the high strain
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hardening rate compared to LTA castings. This is because of easy dislocation motion in the as-cast material due to the low number density and area fraction of  particles. Free from c+d particles is an additional supportive factor in this case. The high number density and area fraction of  particles (c+d + m ) in LTA castings compared to as-cast material cause a decrease in UTS due to the drop in strain hardening rate. However, the high probability of arresting the dislocation movement in LTA-150 ◦ C compared to LTA-200 ◦ C castings provides better UTS for LTA150 ◦ C by a steep increase in hardening slope. Besides, high probability of damage in bigger brittle m particles of LTA200 ◦ C castings softens the material, which decreases the UTS. The FS values shows that the LTA-150 ◦ C case fails early compared to other cases and the same is better in as-cast material. This depends on the fracture behaviour of the material and the important attribute is the grain boundary region as the material fails by intergranular brittle failure due to the lack of dislocation motion from a grain to other. It is already confirmed that the arrangement of  particles in the grain boundary region also speed up the failure process [16]. The arrangement of brittle m and c+d particles in the grain boundary and in the eutectic region weakens the grain boundary region and has significant effect on this as it increases the area fraction of grain boundary region, which inhibits transmission of slip across the boundaries. These particles support the crack initiation and propagation in the grain boundary region due to the development of high stress concentration in the particles during deformation and this is severe in the case of LTA-150 ◦ C as the grain boundary region is almost occupied by the m and c+d particles. As-cast material holds better FS values compared to LTA-200 ◦ C case due to the low size, area fraction and number density of the  particles. In the case of LTA-200 ◦ C, the probability of damage is high as it is having high number density of bigger m particles, which also speeds up the failure process. These explanations also confirm the results of ductility of different castings. 4. Summery and conclusions The present study concludes that the mechanical properties (strength and ductility) vary significantly with LTA conditions. It also confirms the significant variation in the area fraction, average size and clustering tendency of m and c+d particles with respect to the LTA conditions of the casting and their effect on LTA conditions. The area fraction of bulk  phase in as-cast material is low compared to LTA castings and it increases with increasing annealing temperature due to the nucleation of c and d particles and diffusion of Al atoms towards  particles during LTA. As-cast material holds better UTS and ductility compared to the LTA castings. LTA-150 ◦ C offers the vigorous
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nucleation of fine closely spaced c+d particles, which provides better YS compared to as-cast and LTA-200 ◦ C castings and c particles act as a strengthener of annealed hpdc AZ91 alloy. The agglomeration of different  particles takes place in the LTA200 ◦ C castings and provides poor YS and UTS compared to LTA-150 ◦ C castings. The high strain hardening rate provides better ductility for LTA-200 ◦ C compared to the LTA-150 ◦ C castings. The increase in size and area fraction of  particles in the grain boundary region supports the failure process and leads to early failure in LTA castings. References [1] D. Magers, J. Willekens, in: B.L. Mordike, K.U. Kainer (Eds.), Magnesium Alloys and Their Applications, Wolfsburg, Germany, 1998, pp. 105– 112. [2] S. Schumann, F. Friedrich, in: B.L. Mordike, K.U. Kainer (Eds.), Magnesium Alloys and Their Applications, Wolfsburg, Germany, 1998, pp. 3–13. [3] D. Brungs, in: B.L. Mordike, K.U. Kainer (Eds.), Magnesium Alloys and Their Applications, Wolfsburg, Germany, 1998, pp. 113–117. [4] C. Suman, SAE International Congress and Exhibition, SAE Technical Paper Series, Paper No. 890207 (1989). [5] A.L. Bowles, T.J. Bastow, C.J. Davidson, J.R. Griffiths, P.D.D. Rodrigo, Magnesium Technology, The Minerals Metals and Materials Society, Nashville, USA, 2000, pp. 295–300. [6] T.G. Basner, M. Evans, D.J. Sakkinen, Magnesium Properties and Applications for Automobiles, Society of Automotive Engineers, Detroit, USA, 1993, pp. 59–64. [7] H. Westengen, L.-Y. Wei, T. Aune, D. Albright, in: B.L. Mordike, K.U. Kainer (Eds.), Magnesium Alloys and Their Applications, Wolfsburg, Germany, 1998, pp. 209–214. [8] D.J. Sakkinen, M. Evans, Proceedings of the 17th International Die Casting Congress and Exposition, Cleveland USA, 1993, pp. 305–313. [9] G. Song, A.L. Bowles, D.H. StJohn, Mater. Sci. Eng. A 366 (2004) 74– 86. [10] N.N. Aung, W. Zhou, J. Appl. Electrochem. 32 (2002) 1397–1401. [11] D.G.L. Prakash, D. Regener, W.J.J. Vorster, Effect of long term annealing on the microstructure of hpdc AZ91 Mg alloy: a quantitative analysis by image processing, Comput. Mater. Sci., submitted for publication. [12] D.G.L. Prakash, D. Regener, Prakt. Metallogr. 42 (2005) 555–575. [13] D.G.L. Prakash, D. Regener, J. Alloys Compd 461 (2008) 139–146. [14] M.J. Starink, J. Mater. Sci. 32 (1997) 4061–4070. [15] S. Celotto, Acta Mater. 48 (2000) 1775–1787. [16] .G.L. Prakash, D. Regener, W.J.J. Vorster, Sci. Eng. A 488 (2008) 303–310. [17] D. Regener, E. Schick, I. Wagner, H. Heyse, Materialwiss. Werkst. 30 (1999) 525–532. [18] M.S. Yoo, Y.C. Kim, S. Ahn, N.J. Kim, Mater. Sci. Forum 419–422 (2003) 419–424. [19] A.K. Vasudevan, O. Richmond, F. Zok, J.D. Embury, Mater. Sci. Eng. A. 107 (1989) 63–69. [20] Y. Brechet, J. Newell, S. Tao, J.D. Embury, Scrip. Metall. et Mater. 28 (1993) 47–51. [21] J.D. S Embury, Procedings of Twelfth Risø International Symposium on Materials Science, Risø National Laboratory, Roskilde, Denmark, 1991, pp. 317–322.