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SiC composites
Journal of Alloys and Compounds 479 (2009) 451–456
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Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jallcom
Effect of reinforcement geometry on precipitation kinetics of powder metallurgy AA2009/SiC composites ˜ P. Rodrigo ∗ , P. Poza, V. Utrilla, A. Urena Departamento de Ciencia e Ingeniería de Materiales, ESCET, Universidad Rey Juan Carlos, c/Tulipán s/n, 28933 Móstoles, Madrid, Spain
a r t i c l e
i n f o
Article history: Received 26 November 2008 Received in revised form 19 December 2008 Accepted 23 December 2008 Available online 4 January 2009 Keywords: Metal-matrix composites Precipitation Microstructure Calorimetry Transmission electron microscopy (TEM)
a b s t r a c t The effect of reinforcement morphology on metastable precipitation in AA2009 matrix composite reinforced with 15% volume of SiC has been studied. AA2009 is a special formulation Al–Cu–Mg alloy for discontinuously reinforced aluminium (DRA) whose ageing behaviour and precipitation kinetics are not well defined. The ageing kinetics for this material have been evaluated at two different ageing temperatures (170 and 190 ◦ C). Hardness peaks in the two different precipitation sequences existing in the matrix alloy have been identified. The response of composites to the ageing treatments was evaluated by means of Vickers hardness and micro-hardness measurements on polished specimens and was also studied by differential scanning calorimetry (DSC). © 2009 Elsevier B.V. All rights reserved.
1. Introduction Reinforcing aluminium matrices with short ceramic reinforcements (particles or whiskers) has been extensively used over the last 20 years to develop a promising category of aluminium matrix composites (AMCs). The ceramic reinforcements improve the stiffness, the wear and the creep resistance and, to a minor extent, the strength of the metallic matrices. However, massive commercialization of Al/SiCp composites has been hindered both by technological barriers and high processing cost [1]. Beside, most of the aluminium alloys used as matrices are able to be hardened by ageing; for this reason, the reinforcement influence on the precipitation behaviour, under different heat treatment conditions, is an important goal to optimise the composite performance. Age hardening in AMCs does not only depend on the matrix alloy characteristics but also on the shape, size and distribution of reinforcements into the matrix [2,3]. Although this topic has been extensively studied it is necessary to extend the actual knowledge to new alloys developed for AMCs. Precipitation kinetics in AMCs are usually faster than in the corresponding unreinforced alloys. In addition, the matrices of the AMC are harder than the unreinforced counterparts. For this reasons, ageing response of aluminium alloys and AMCs continues being reason for research: in composites with different reinforcements percentages [4], determination of the aging precipitation by transmission electron microscopy (TEM) and differential scan-
∗ Corresponding author. Tel.: +34 914887073; fax: +34 914888150. E-mail address: [email protected] (P. Rodrigo). 0925-8388/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2008.12.114
ning calorimetry (DSC) [5], effects of different alloy components in precipitation behaviour [6] numerical models to predict the microstructural evolution [7,8], influence of strain rates on precipitate thickening [9]. This different behaviour has been attributed [10] to the change in dislocation and vacancies densities in the matrix due to the reinforcement presence. The dislocation density increases in AMC, mainly close to the ceramic particles, due to thermal stresses that appear in the composites during cooling from the processing temperature. This is a consequence of the thermal expansion coefficient mismatch between aluminium alloy and ceramic reinforcement. Thermal stresses relief is reached by means of dislocations emission from the matrix/reinforcement interface towards the matrix. On the contrary, the number of vacancies diminishes in the matrix as the reinforcement content is increased, because grain boundaries, dislocations and, mainly, matrix/reinforcement interfaces act as vacancies sinks. The nucleation of incoherent precipitates is favoured by the presence of reinforcements due to the higher defect density. Al–Cu–Mg alloys show two different aging sequences, depending on the Cu/Mg ratio, and under overageing conditions, two equilibrium phases can be precipitated: (Al2 Cu) and S (Al2 CuMg). Both precipitation sequences could take place simultaneously or separately [11]:
SSS → GP zones → → → (Al2 Cu)
(1)
SSS → co-clusters/GPB zones → S → S /S(Al2 CuMg)
(2)
The objective of this investigation is to analyse the precipitation behaviour of the AA2009 aluminium alloy, developed exclusively
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Table 1 Nominal composition of the AA2009 matrix alloy (wt%). Alloy
Cu
Mg
Zn
Si
Fe
O
Other (each)
Other (total)
Al
AA2009
3.2–4.4
1.0–1.6
0.10
0.25
0.07
0.6
0.05
0.15
Balance
to be discontinuously reinforced with SiC for aerospace applications. Its main characteristic is a very low impurity (Fe and Si) content. AA2009/SiC is attractive for its high strength, stiffness and thermal stability, being an ideal substitute for titanium used at moderate temperatures. The mechanical behaviour of Al 2009 matrix composites has been studied with special interest on hot formability, deformation behaviour, tensile properties and superplasticity [12–16]. The effect of SiC particulate size on the hardening response of AA2009 was studied by Sannino and Rack [17]; this paper concluded that the stage of age hardening consists of heterogeneous nucleation of S /S. However, the influence of reinforcement shape has not been studied and this is the main aim of this paper. For it, the ageing behaviour and precipitation kinetic are analysed by Vickers hardness and micro-hardness measurements and DSC.
route described by Geiger and Walker [18]. Materials were supplied in the form of extruded plates in the F-temper, with 20 mm thickness for the material reinforced with whiskers and 13 mm for the particle-reinforced composite. 2.2. Ageing treatments Artificial ageing treatments (T6) were applied to the as-received composites. Firstly, specimens of both materials were solution annealed. The solution treatments
2. Materials and experimental technique 2.1. Materials Two materials based on an AA2009 alloy, with a chemical composition showed in Table 1, reinforced with 15 vol.% SiC were studied in this investigation. One of them was reinforced with SiC particles (AA2009/SiC/15p) while whiskers were used in the other one (AA2009/SiC/15w). Both composites were manufactured by Advanced Composites Materials Corporation (Greer, USA) with a powder metallurgy (PM)
Fig. 1. SEM images of parent composites showing banding and clustering: (a) AA2009/SiC/15w and (b) AA2009/SiC/15p.
Fig. 2. TEM images of (Al2 Cu) and S (Al2 CuMg) precipitates associated to SiC cluster ¯ and (b) zones: (a) AA2009/SiC/15w composite with ED pattern of Al2 CuMg [1 0 1] AA2009/SiC/15p composite, including ED pattern of Al2 Cu [1 1¯ 1].
P. Rodrigo et al. / Journal of Alloys and Compounds 479 (2009) 451–456
453
were carried out at 500 ± 5 ◦ C. AA2009/SiC/15w was treated for 2 h while the treatment was extended up to 4 h for AA2009/SiC/15p because the greater amount of precipitated phases detected for this composite in the as-received conditions. The quenching media was a mixture of water–ice (0 ◦ C). Subsequently, artificial ageing treatments were carried out into a silicone oil bath for times in the range of 15 min to 33 h, at two different ageing temperatures: 170 ± 1 and 190 ± 1 ◦ C. The ageing response of both composites was evaluated by means of Vickers hardness according to UNE-EN ISO 6507-1 [19] standard and micro-hardness measurements on polished specimens, considering the hardness for each condition as the average value of at least 10 measurements. Two different microhardness Vickers testers were used. For the whisker reinforced composite, micro-hardness measurements were carried out using an Akasi MVK-E3 equipment with a load of 25 g applied for 15 s. In the case of the particle-reinforced composite, a Buehler Micromet 2103 machine was used with a load of 50 g for 12 s. The hardness tests were performed with a Vickers hardness tester (Wolpert TESTOR 2100) applying 500 g for 12 s. The influence of the reinforcement morphology on the precipitation sequences was also studied by DSC. These tests were conducted on composite specimens quenched in water (20 ◦ C) and mixture water–ice (0 ◦ C) using a Mettler Toledo DSC 822. All DSC tests were carried out under an argon atmosphere with scanning rates of 10 and 30 ◦ C/min. The DSC scans started at 25 ◦ C and finished at 500 ◦ C. At least two tests were made in each condition and material. The microstructure of these materials was studied using environmental scanning electron microscopy (ESEM) and TEM, as well as energy dispersive X-ray microanalysis (EDX). Metallographic samples were ground on emery paper up to 1200 grade and polished with diamond paste of 3 m particle size and observed using a Phillips XL30 ESEM equipped with EDX. TEM samples were examined using a Philips Tecnai 20 transmission electron microscope employing selected area (SAD) pattern and nano-beam-diffraction (NBD) pattern as well as EDX analysis.
3. Results and discussion 3.1. Microstructure of as-received composites Both studied composites showed a heterogeneous distribution of SiC reinforcements inside the aluminium alloy matrix, being visible the characteristic preferential orientation produced by the extrusion processes applied to manufacture the composite plates from the PM pieces (Fig. 1). Inhomogeneities in the reinforcement distribution (banding and clustering) are distinguished. These defects in the ceramic phase distribution are more important in the composite reinforced with whiskers due to their higher aspect ratios which favour their alignment during plastic deformation. Whiskers lengths were in the range of 5 to 20 m with average diameters of 0.5 m (l/d = 10–40). The composites studied were received in the as-extruded condition, consequently intermetallic compounds of great sizes, preferentially nucleated on the reinforcement surfaces, were observed. TEM images and electron diffraction (ED) patterns of both parent materials showed the equilibrium phases: orthorhombic S phase, Al2 CuMg (Cmcm a = 0.400 nm, b = 0.923 nm, c = 0.714 nm) (Fig. 2a) and tetragonal phase, Al2 Cu (I4/mcm, structure with a = 0.6067 nm and c = 0.4877 nm) (Fig. 2b). is the majority phase, forming larger precipitates with sizes close to 1 m. In addition, some oxide inclusions of smaller sizes, mainly of aluminium and magnesium, were also observed; these were formed from the oxide layer which surrounds the aluminium powders used in the composite manufacture. The incorporation of ceramic reinforcements into an aluminium alloy always induces residual stresses from mismatching in the coefficients of thermal expansion (CTE␣–SiC = 3.8 × 10−6 K−1 and CTEAl = 23 × 10−6 K−1 ). These stresses are concentrated in the metallic matrix surrounding the discontinuous reinforcements (particles or whiskers) which could be plastically deformed. Plastic deformation is especially intense in the matrix close to the reinforcement corners. Other places of intense plastic deformation are the reinforcement clusters, where a greater density of dislocations generated in the matrix is usually observed. The main difference observed in the matrix microstructure of both composites is related with the presence of small Mg2 Si precipitates (hexagonal structure, P62c) (Fig. 3a) and large Ferich intermetallics (Al7 Cu2 Fe, tetragonal P4/mmc, a = 0.6336 nm,
Fig. 3. TEM images of (a) Mg2 Si formed on the SiCp surface, including its ED pattern [1 1 1] and (b) an intermetallic compound identified as Al7 Cu2 Fe, including its ED ¯ pattern [1 0 1].
c = 1.4870 nm) (Fig. 3b) identified by means of TEM-ED in the AA2009/SiC/15p composite. Although the proportion of these precipitates is very low, because the amount of Fe in the matrix is limited (0.073%), its presence could influence the aging response of the particle-reinforced composite, since these intermetallics are not dissolved during the solution heat treatment. This diminishes the amount of solubilised Cu in the matrix and, therefore, the percentage of this element available for precipitation hardening. The presence of Mg2 Si on the SiCp surface and Al7 Cu2 Fe in the matrix would indicate that the AA2009 powders used for the
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Fig. 4. Age hardening curves: (a) AA2009/SiC/15p composite aged at 170 ◦ C, (b) AA2009/SiC/15w aged at 170 ◦ C, (c) AA2009/SiC/15p aged at 190 ◦ C and (d) AA2009/SiC/15w aged at 190 ◦ C.
AA2009/SiC/15p composite manufacturing would contain higher proportion of Fe and Si. This last element is combined with Mg, which has a high tendency to migrate from the matrix to the particles surfaces, precipitating in form of Mg2 Si. 3.2. Ageing behaviour The hardness and micro-hardness curves obtained for both aged composites at 170 and 190 ◦ C are quite similar and two hardening peaks are distinguished. The first one, formed after several increases and decreases of hardness, appears after very short ageing times. The second one, which is better defined, is followed by a gradual decrease of the hardness, which can be understood as an overageing effect from this point. Experimental hardness (HV5) and microhardness (HV0.05 and HV0.0025) curves for both composites are plotted in Fig. 4. The main effect of the increase in aging temperature (from 170 to 190 ◦ C) is a reduction on times to reach the maximum hardnesses, which means the acceleration of the precipitation kinetics. It has also been observed that as the ageing temperature increased the peak hardness values diminished. Several explanations for the first stage of hardening in Al–Cu–Mg alloys are proposed by different authors. Formation of GPB zones in this aluminium alloy family has been studied by several authors since 1960 until now, this zones can coexist with S , S and S phases [20]. Ringer et al. [21] have found, applying atom probe film ion microscopy (APFIM), that GPB zones are associated to the formation of co-clusters Cu/Mg during the initial stages of the artificial ageing. Reich et al. [22] suggested that it is due to solute-dislocations interaction. Ratchev et al. [23] attributed this effect to heterogeneous formation of a phase S on dislocations helices. The reinforcement morphology has also influence on the hardening response of these composites. The acceleration of the precipitation kinetics, as the ageing temperature is increased,
depends on the reinforcement shape as can be observed for the two maxima of curves plotted in Fig. 4. The composite reinforced with SiC whiskers reaches both hardening peaks at 170 ◦ C after 1 and 10 h, whereas the particle-reinforced material needs 6 and 20–24 h to acquire similar hardness at the same temperature. However, when the ageing temperature is increased up to 190 ◦ C, the first hardness peak was detected after 30 min in both materials and the second maximum was reached after 15 h in the particle-reinforced composite, while the whiskers reinforced one only needed 6 h. These results show that AA2009/SiC/15p has greater sensitivity to the ageing temperature than AA2009/SiC/15w. This behaviour could be explained considering a higher dislocation density close to the whiskers than that observed close to the particles; as a consequence, precipitation kinetics of the hardening phases are accelerated. 3.3. Precipitation kinetics Results from DSC scans of both materials, subjected to different quenching treatments and scanning rates are showed in Figs. 5 and 6, respectively. DSC curves for both composites present four different thermal effects. Two exothermic effects, marked with (I) and (III), and other two endothermic ones, identified as (II) and (IV). An increase in heating rates (from 10 to 30 ◦ C/min) causes the displacement of the transformation peaks towards higher temperatures. When the heating rate is high (30 ◦ C/min), the influence of the quenching rate is not significant and all peaks appear to the same temperatures. Exothermic effect (I) corresponds to the formation of GP/GPB zones and the effect (II) to their solubilisation [19,24,25]. This effect (I) appears at lower temperature in the whisker reinforced material than in the particulate composite, when 10 ◦ C/min was applied as heating rate.
P. Rodrigo et al. / Journal of Alloys and Compounds 479 (2009) 451–456
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Fig. 5. DSC thermograms of as-quenched AA2009/SiC/15w using heating rates of 10 and 30 ◦ C/min: (a) quenched in water and (b) quenched in water–ice mixture.
Fig. 6. DSC thermograms of as-quenched AA2009/SiC/15p using heating rates of 10 and 30 ◦ C/min: (a) quenched in water and (b) quenched in water–ice mixture.
Effect (III) is constituted by two exothermic overlapped peaks; this can be clearly distinguished from the differential curve of heat flow because the appearance of a well-defined shoulder. The first peak of the doublet (IIIA) has been explained considering two possibilities: it could be attributed to the formation of phase [26,27] and to the formation of S phase [28–30]. The second peak of the doublet (IIIB) has also several explanations in Al–Cu–Mg alloys: the formation of + S phases, only phase [30,31] or two distinct S phase precipitates [32]. However, it can be accepted that the first peak (IIIA) corresponds to the formation of S phase and the second one (IIIB) to the phase formation attending to the transformation temperatures (Table 2). The formation of S and phases was also confirmed by TEM; Fig. 7 shows a TEM image along the 1 0 0 zone axis of the matrix of AA2009/SiCp composite treated at 170 ◦ C for 20 h and the SAED pattern where S , and phases were identified. Effect (IV)
would be associated to the dissolution of the phases formed in (III). The fact that the presence of SiC whiskers transform the doublet observed in the particulate composite into a single exothermic effect (III) could be explained considering that the incorporation of SiCw accelerates the kinetic of precipitation of phase, and both exothermic effects (IIIA and IIIB) overlap to the same temperature. The height of the exothermic peak increases, which indicates a greater amount of formed phases. If both phases precipitate approximately at the same temperature, the height of the exothermic effect will be greater than that corresponding to two separated peaks. Finally, the two different precipitation sequences (1 and 2) which can occur, separately or simultaneously, in the Al–Cu–Mg alloys depending on the content in Cu and on the Cu/Mg relation are considered. The proportions of and S phases in the composite matrix should be approximately 2.1% and 5.3%, respectively,
Table 2 Temperature (in ◦ C) of exothermic transformations detected by DSC in parent composites. Material
Quenching media Water
AA2009/SiC/15w Water–ice Water AA2009/SiC/15p Water–ice
Heating rate (◦ C/min)
GPZ/GPBZ formation
IIIA effect
IIIB effect
10 30 10 30
68.2 88.8 73.4 88.8
264.8 292.7 261.4 273.3
– – 273.0 295.8
10 30 10 30
80.0 91.6 73.4 90.7
263.4 287.9 263.9 288.1
284.4 306.3 274.1 307.5
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the cooling rate during quenching in the precipitation kinetics depends on the morphology of the reinforcement. (4) Two exothermic transformations were found in both studied composites by means DSC curves. First of them is associated to the simultaneous formation of GP and GPB zones. The second one is constituted by exothermic doublets formed for the overlapped precipitation of S and phases. (5) The exothermic effect formed by the overlapping of S and phases precipitation is very sensitive to the cooling media used for quenching the AA2009/SiC/15w composite. Both exothermic transformations are not completely overlapped in specimens quenched in water–ice. However, the exothermic transformation did not form a doublet in water quenched specimens. Acknowledgements Authors wish to thank Ministerio de Educación y Ciencia for economical support present research (project MAT2004-06018) and Comunidad de Madrid through the program ESTRUMAT-CM (reference MAT/77). References
Fig. 7. AA2009/SiCp composite aged at 170 ◦ C for 20 h. Detail of the hardening precipitates, S , and phases in [0 0 1]␣ orientation.
considering the Cu (3.5 wt%) and Mg (1.1 wt%) content in AA2009. Small changes in the Cu/Mg ratio have a great influence in the amount of formed equilibrium phases and a local increase in the Mg concentration produce a considerable reduction of the amount of Al2 Cu formed. For that reason, the formation of Al7 Cu2 Fe precipitates detected only in the composite reinforced with particles could affect to the proportion of and S phases formed in this composite, favouring the formation of S because of the removing of the dissolved Cu from the matrix. This is in agreement with the doublet observed in the particle-reinforced material where the first peak is assigned to the formation of S . 4. Conclusions (1) Reinforcement morphology modifies the precipitation behaviour of the AA2009 matrix composites. AA2009/SiC/15p composite exhibits Fe-rich intermetallic phases, Al7 Cu2 Fe, and Mg2 Si precipitated close to the SiC/Al interfaces. This has an influence on the ageing kinetics of this material as the Cu/Mg rate is modified with respect to AA2009/SiC/15w composite. (2) Ageing kinetics of AA2009/SiC/15p composite are highly influenced by the ageing temperature. Times necessary to reach ageing peaks in this composite are larger than those observed in the material reinforced with 15 vol.% of SiC whiskers at 170 ◦ C. An increase in the ageing temperature up to 190 ◦ C reduces the ageing times more effectively in the case of the particlereinforced composite than in the whiskers one. (3) As the heating rate increases, the exothermic transformation is displaced towards higher temperatures. The influence of
[1] F. Ortega-Celaya, M.I. Pech-Canul, J. López-Cuevas, J.C. Rendón-Ángeles, M.A. Pech-Canul, J. Mater. Process. Technol. 183 (2007) 368–373. [2] A. Borrego, G. González-Doncel, Mater. Sci. Eng. A 245 (1998) 10–18. [3] J.T. Wang, M. Furukawa, Z. Horita, M. Nemoto, Y. Ma, T.G. Langdon, Metall. Mater. Trans. A 26 (1995) 581–587. [4] Sharmilee Pal, R. Mitra, V.V. Bhanuprasad, Mater. Sci. Eng. A 480 (2008) 496–505. [5] D. Ovono Ovono, I. Guillot, D. Massinon, J. Alloy Compd. 432 (2007) 241–246. [6] D. Bakavos, P.B. Prangnell, B. Bes, F. Eberl, Mater. Sci. Eng. A 491 (2008) 214– 223. [7] I.N. Khan, M.J. Starink, J.L. Yan, Mater. Sci. Eng. A 472 (2008) 66–74. [8] W. Sha, Mater. Design 28 (2007) 1524–1530. [9] Z. Gao, X. Zhang, M. Chena, J. Alloys Compd. (2008), doi:10.1016/j.jallcom. 2008.08.055. [10] S. Suresh, K.K. Chawla, in: S. Suresh, A. Mortensen, A. Needleman (Eds.), Fundamentals of Metal Matrix Composites, Butterworth-Heinemann, 1993, pp. 119–136. [11] P.I. Gouma, D.J. Lloyd, M.J. Mills, Mater. Sci. Eng. A 319–321 (2001) 439–442. [12] K.C. Chan, G.Q. Tong, L. Gao, J. Mater. Process. Technol. 74 (1998) 142–148. [13] B.L. Xiao, J.Z. Fan, X.F. Tian, W.Y. Zhang, L.K. Shi, J. Mater. Sci. 40 (2005) 5757–5762. [14] T.S. Srivatsan, Meslet Al-Hajri, C. Smith, M. Petraroli, Mater. Sci. Eng. A 346 (2003) 91–100. [15] T.S. Srivatsan, V. Meslet Al-Hajri, K. Vasudevan, Int. J. Fatigue 27 (2005) 357–371. [16] B.Q. Han, K.C. Chan, J. Mater. Sci. Lett. 16 (1997) 827–829. [17] A.P. Sannino, H.J. Rack, J. Mater. Sci. 30 (1995) 4316–4322. [18] A.L. Geiger, J.A. Walker, JOM 43 (1991) 8–15. [19] UNE-EN ISO 6507-1. European standard: Metallic materials. Vickers hardness test. Part I: Test method (ISO 6507-1:2005). [20] A. Charai, T. Walther, C. Alfonso, A.-M. Zahra, C.Y. Zahra, Acta Mater. 48 (10) (2000) 2751–2764. [21] S.P. Ringer, K. Hono, T. Saksai, I.J. Polmear, Scripta Mater. 36 (5) (1997) 517– 521. [22] L. Reich, S.P. Ringer, K. Hono, Phil. Mag. Lett. 79 (9) (1999) 639–648. [23] P. Ratchev, B. Verlinden, P. De Smet, P. Van Houtte, Acta Mater. 46 (10) (1998) 3523–3533. [24] A.M. Zahra, C.Y. Zahra, C. Alfonso, A. Charaï, Scripta Mater. 39 (11) (1998) 1553–1558. [25] A. Hayoune, D. Hamana, J. Alloys Compd. (2008), doi:10.1016/j.jallcom. 2008.06.070. [26] K.K. Chawla, A.H. Esmaeili, A.K. Datye, A.K. Vasudevan, Scripta Metall. Mater. 25 (1991) 1315–1319. [27] T.S. Kim, T.H. Kim, K.H. Oh, H.I. Lee, J. Mater. Sci. 27 (10) (1992) 2599–2605. [28] S. Arakawa, T. Hatayama, K. Matsugi, O. Yanagisawa, Scripta Mater. 42 (2000) 755–760. [29] V.K. Varma, Y.R. Mahajan, V.V. Kutumbarao, Scripta Mater. 37 (1997) 485– 489. [30] Z. Chen, Z. Zheng, Scripta Mater. 50 (2004) 1067–1071. [31] Z.Y. Ma, J. Bi, Y.X. Lu, M. Luo, Y.X. Gao, Scripta Metall. Mater. 29 (1993) 225–229. [32] S.C. Wang, M.J. Starink, Acta Mater. 55 (2007) 933–941.
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jallcom
Effect of reinforcement geometry on precipitation kinetics of powder metallurgy AA2009/SiC composites ˜ P. Rodrigo ∗ , P. Poza, V. Utrilla, A. Urena Departamento de Ciencia e Ingeniería de Materiales, ESCET, Universidad Rey Juan Carlos, c/Tulipán s/n, 28933 Móstoles, Madrid, Spain
a r t i c l e
i n f o
Article history: Received 26 November 2008 Received in revised form 19 December 2008 Accepted 23 December 2008 Available online 4 January 2009 Keywords: Metal-matrix composites Precipitation Microstructure Calorimetry Transmission electron microscopy (TEM)
a b s t r a c t The effect of reinforcement morphology on metastable precipitation in AA2009 matrix composite reinforced with 15% volume of SiC has been studied. AA2009 is a special formulation Al–Cu–Mg alloy for discontinuously reinforced aluminium (DRA) whose ageing behaviour and precipitation kinetics are not well defined. The ageing kinetics for this material have been evaluated at two different ageing temperatures (170 and 190 ◦ C). Hardness peaks in the two different precipitation sequences existing in the matrix alloy have been identified. The response of composites to the ageing treatments was evaluated by means of Vickers hardness and micro-hardness measurements on polished specimens and was also studied by differential scanning calorimetry (DSC). © 2009 Elsevier B.V. All rights reserved.
1. Introduction Reinforcing aluminium matrices with short ceramic reinforcements (particles or whiskers) has been extensively used over the last 20 years to develop a promising category of aluminium matrix composites (AMCs). The ceramic reinforcements improve the stiffness, the wear and the creep resistance and, to a minor extent, the strength of the metallic matrices. However, massive commercialization of Al/SiCp composites has been hindered both by technological barriers and high processing cost [1]. Beside, most of the aluminium alloys used as matrices are able to be hardened by ageing; for this reason, the reinforcement influence on the precipitation behaviour, under different heat treatment conditions, is an important goal to optimise the composite performance. Age hardening in AMCs does not only depend on the matrix alloy characteristics but also on the shape, size and distribution of reinforcements into the matrix [2,3]. Although this topic has been extensively studied it is necessary to extend the actual knowledge to new alloys developed for AMCs. Precipitation kinetics in AMCs are usually faster than in the corresponding unreinforced alloys. In addition, the matrices of the AMC are harder than the unreinforced counterparts. For this reasons, ageing response of aluminium alloys and AMCs continues being reason for research: in composites with different reinforcements percentages [4], determination of the aging precipitation by transmission electron microscopy (TEM) and differential scan-
∗ Corresponding author. Tel.: +34 914887073; fax: +34 914888150. E-mail address: [email protected] (P. Rodrigo). 0925-8388/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2008.12.114
ning calorimetry (DSC) [5], effects of different alloy components in precipitation behaviour [6] numerical models to predict the microstructural evolution [7,8], influence of strain rates on precipitate thickening [9]. This different behaviour has been attributed [10] to the change in dislocation and vacancies densities in the matrix due to the reinforcement presence. The dislocation density increases in AMC, mainly close to the ceramic particles, due to thermal stresses that appear in the composites during cooling from the processing temperature. This is a consequence of the thermal expansion coefficient mismatch between aluminium alloy and ceramic reinforcement. Thermal stresses relief is reached by means of dislocations emission from the matrix/reinforcement interface towards the matrix. On the contrary, the number of vacancies diminishes in the matrix as the reinforcement content is increased, because grain boundaries, dislocations and, mainly, matrix/reinforcement interfaces act as vacancies sinks. The nucleation of incoherent precipitates is favoured by the presence of reinforcements due to the higher defect density. Al–Cu–Mg alloys show two different aging sequences, depending on the Cu/Mg ratio, and under overageing conditions, two equilibrium phases can be precipitated: (Al2 Cu) and S (Al2 CuMg). Both precipitation sequences could take place simultaneously or separately [11]:
SSS → GP zones → → → (Al2 Cu)
(1)
SSS → co-clusters/GPB zones → S → S /S(Al2 CuMg)
(2)
The objective of this investigation is to analyse the precipitation behaviour of the AA2009 aluminium alloy, developed exclusively
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P. Rodrigo et al. / Journal of Alloys and Compounds 479 (2009) 451–456
Table 1 Nominal composition of the AA2009 matrix alloy (wt%). Alloy
Cu
Mg
Zn
Si
Fe
O
Other (each)
Other (total)
Al
AA2009
3.2–4.4
1.0–1.6
0.10
0.25
0.07
0.6
0.05
0.15
Balance
to be discontinuously reinforced with SiC for aerospace applications. Its main characteristic is a very low impurity (Fe and Si) content. AA2009/SiC is attractive for its high strength, stiffness and thermal stability, being an ideal substitute for titanium used at moderate temperatures. The mechanical behaviour of Al 2009 matrix composites has been studied with special interest on hot formability, deformation behaviour, tensile properties and superplasticity [12–16]. The effect of SiC particulate size on the hardening response of AA2009 was studied by Sannino and Rack [17]; this paper concluded that the stage of age hardening consists of heterogeneous nucleation of S /S. However, the influence of reinforcement shape has not been studied and this is the main aim of this paper. For it, the ageing behaviour and precipitation kinetic are analysed by Vickers hardness and micro-hardness measurements and DSC.
route described by Geiger and Walker [18]. Materials were supplied in the form of extruded plates in the F-temper, with 20 mm thickness for the material reinforced with whiskers and 13 mm for the particle-reinforced composite. 2.2. Ageing treatments Artificial ageing treatments (T6) were applied to the as-received composites. Firstly, specimens of both materials were solution annealed. The solution treatments
2. Materials and experimental technique 2.1. Materials Two materials based on an AA2009 alloy, with a chemical composition showed in Table 1, reinforced with 15 vol.% SiC were studied in this investigation. One of them was reinforced with SiC particles (AA2009/SiC/15p) while whiskers were used in the other one (AA2009/SiC/15w). Both composites were manufactured by Advanced Composites Materials Corporation (Greer, USA) with a powder metallurgy (PM)
Fig. 1. SEM images of parent composites showing banding and clustering: (a) AA2009/SiC/15w and (b) AA2009/SiC/15p.
Fig. 2. TEM images of (Al2 Cu) and S (Al2 CuMg) precipitates associated to SiC cluster ¯ and (b) zones: (a) AA2009/SiC/15w composite with ED pattern of Al2 CuMg [1 0 1] AA2009/SiC/15p composite, including ED pattern of Al2 Cu [1 1¯ 1].
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were carried out at 500 ± 5 ◦ C. AA2009/SiC/15w was treated for 2 h while the treatment was extended up to 4 h for AA2009/SiC/15p because the greater amount of precipitated phases detected for this composite in the as-received conditions. The quenching media was a mixture of water–ice (0 ◦ C). Subsequently, artificial ageing treatments were carried out into a silicone oil bath for times in the range of 15 min to 33 h, at two different ageing temperatures: 170 ± 1 and 190 ± 1 ◦ C. The ageing response of both composites was evaluated by means of Vickers hardness according to UNE-EN ISO 6507-1 [19] standard and micro-hardness measurements on polished specimens, considering the hardness for each condition as the average value of at least 10 measurements. Two different microhardness Vickers testers were used. For the whisker reinforced composite, micro-hardness measurements were carried out using an Akasi MVK-E3 equipment with a load of 25 g applied for 15 s. In the case of the particle-reinforced composite, a Buehler Micromet 2103 machine was used with a load of 50 g for 12 s. The hardness tests were performed with a Vickers hardness tester (Wolpert TESTOR 2100) applying 500 g for 12 s. The influence of the reinforcement morphology on the precipitation sequences was also studied by DSC. These tests were conducted on composite specimens quenched in water (20 ◦ C) and mixture water–ice (0 ◦ C) using a Mettler Toledo DSC 822. All DSC tests were carried out under an argon atmosphere with scanning rates of 10 and 30 ◦ C/min. The DSC scans started at 25 ◦ C and finished at 500 ◦ C. At least two tests were made in each condition and material. The microstructure of these materials was studied using environmental scanning electron microscopy (ESEM) and TEM, as well as energy dispersive X-ray microanalysis (EDX). Metallographic samples were ground on emery paper up to 1200 grade and polished with diamond paste of 3 m particle size and observed using a Phillips XL30 ESEM equipped with EDX. TEM samples were examined using a Philips Tecnai 20 transmission electron microscope employing selected area (SAD) pattern and nano-beam-diffraction (NBD) pattern as well as EDX analysis.
3. Results and discussion 3.1. Microstructure of as-received composites Both studied composites showed a heterogeneous distribution of SiC reinforcements inside the aluminium alloy matrix, being visible the characteristic preferential orientation produced by the extrusion processes applied to manufacture the composite plates from the PM pieces (Fig. 1). Inhomogeneities in the reinforcement distribution (banding and clustering) are distinguished. These defects in the ceramic phase distribution are more important in the composite reinforced with whiskers due to their higher aspect ratios which favour their alignment during plastic deformation. Whiskers lengths were in the range of 5 to 20 m with average diameters of 0.5 m (l/d = 10–40). The composites studied were received in the as-extruded condition, consequently intermetallic compounds of great sizes, preferentially nucleated on the reinforcement surfaces, were observed. TEM images and electron diffraction (ED) patterns of both parent materials showed the equilibrium phases: orthorhombic S phase, Al2 CuMg (Cmcm a = 0.400 nm, b = 0.923 nm, c = 0.714 nm) (Fig. 2a) and tetragonal phase, Al2 Cu (I4/mcm, structure with a = 0.6067 nm and c = 0.4877 nm) (Fig. 2b). is the majority phase, forming larger precipitates with sizes close to 1 m. In addition, some oxide inclusions of smaller sizes, mainly of aluminium and magnesium, were also observed; these were formed from the oxide layer which surrounds the aluminium powders used in the composite manufacture. The incorporation of ceramic reinforcements into an aluminium alloy always induces residual stresses from mismatching in the coefficients of thermal expansion (CTE␣–SiC = 3.8 × 10−6 K−1 and CTEAl = 23 × 10−6 K−1 ). These stresses are concentrated in the metallic matrix surrounding the discontinuous reinforcements (particles or whiskers) which could be plastically deformed. Plastic deformation is especially intense in the matrix close to the reinforcement corners. Other places of intense plastic deformation are the reinforcement clusters, where a greater density of dislocations generated in the matrix is usually observed. The main difference observed in the matrix microstructure of both composites is related with the presence of small Mg2 Si precipitates (hexagonal structure, P62c) (Fig. 3a) and large Ferich intermetallics (Al7 Cu2 Fe, tetragonal P4/mmc, a = 0.6336 nm,
Fig. 3. TEM images of (a) Mg2 Si formed on the SiCp surface, including its ED pattern [1 1 1] and (b) an intermetallic compound identified as Al7 Cu2 Fe, including its ED ¯ pattern [1 0 1].
c = 1.4870 nm) (Fig. 3b) identified by means of TEM-ED in the AA2009/SiC/15p composite. Although the proportion of these precipitates is very low, because the amount of Fe in the matrix is limited (0.073%), its presence could influence the aging response of the particle-reinforced composite, since these intermetallics are not dissolved during the solution heat treatment. This diminishes the amount of solubilised Cu in the matrix and, therefore, the percentage of this element available for precipitation hardening. The presence of Mg2 Si on the SiCp surface and Al7 Cu2 Fe in the matrix would indicate that the AA2009 powders used for the
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Fig. 4. Age hardening curves: (a) AA2009/SiC/15p composite aged at 170 ◦ C, (b) AA2009/SiC/15w aged at 170 ◦ C, (c) AA2009/SiC/15p aged at 190 ◦ C and (d) AA2009/SiC/15w aged at 190 ◦ C.
AA2009/SiC/15p composite manufacturing would contain higher proportion of Fe and Si. This last element is combined with Mg, which has a high tendency to migrate from the matrix to the particles surfaces, precipitating in form of Mg2 Si. 3.2. Ageing behaviour The hardness and micro-hardness curves obtained for both aged composites at 170 and 190 ◦ C are quite similar and two hardening peaks are distinguished. The first one, formed after several increases and decreases of hardness, appears after very short ageing times. The second one, which is better defined, is followed by a gradual decrease of the hardness, which can be understood as an overageing effect from this point. Experimental hardness (HV5) and microhardness (HV0.05 and HV0.0025) curves for both composites are plotted in Fig. 4. The main effect of the increase in aging temperature (from 170 to 190 ◦ C) is a reduction on times to reach the maximum hardnesses, which means the acceleration of the precipitation kinetics. It has also been observed that as the ageing temperature increased the peak hardness values diminished. Several explanations for the first stage of hardening in Al–Cu–Mg alloys are proposed by different authors. Formation of GPB zones in this aluminium alloy family has been studied by several authors since 1960 until now, this zones can coexist with S , S and S phases [20]. Ringer et al. [21] have found, applying atom probe film ion microscopy (APFIM), that GPB zones are associated to the formation of co-clusters Cu/Mg during the initial stages of the artificial ageing. Reich et al. [22] suggested that it is due to solute-dislocations interaction. Ratchev et al. [23] attributed this effect to heterogeneous formation of a phase S on dislocations helices. The reinforcement morphology has also influence on the hardening response of these composites. The acceleration of the precipitation kinetics, as the ageing temperature is increased,
depends on the reinforcement shape as can be observed for the two maxima of curves plotted in Fig. 4. The composite reinforced with SiC whiskers reaches both hardening peaks at 170 ◦ C after 1 and 10 h, whereas the particle-reinforced material needs 6 and 20–24 h to acquire similar hardness at the same temperature. However, when the ageing temperature is increased up to 190 ◦ C, the first hardness peak was detected after 30 min in both materials and the second maximum was reached after 15 h in the particle-reinforced composite, while the whiskers reinforced one only needed 6 h. These results show that AA2009/SiC/15p has greater sensitivity to the ageing temperature than AA2009/SiC/15w. This behaviour could be explained considering a higher dislocation density close to the whiskers than that observed close to the particles; as a consequence, precipitation kinetics of the hardening phases are accelerated. 3.3. Precipitation kinetics Results from DSC scans of both materials, subjected to different quenching treatments and scanning rates are showed in Figs. 5 and 6, respectively. DSC curves for both composites present four different thermal effects. Two exothermic effects, marked with (I) and (III), and other two endothermic ones, identified as (II) and (IV). An increase in heating rates (from 10 to 30 ◦ C/min) causes the displacement of the transformation peaks towards higher temperatures. When the heating rate is high (30 ◦ C/min), the influence of the quenching rate is not significant and all peaks appear to the same temperatures. Exothermic effect (I) corresponds to the formation of GP/GPB zones and the effect (II) to their solubilisation [19,24,25]. This effect (I) appears at lower temperature in the whisker reinforced material than in the particulate composite, when 10 ◦ C/min was applied as heating rate.
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Fig. 5. DSC thermograms of as-quenched AA2009/SiC/15w using heating rates of 10 and 30 ◦ C/min: (a) quenched in water and (b) quenched in water–ice mixture.
Fig. 6. DSC thermograms of as-quenched AA2009/SiC/15p using heating rates of 10 and 30 ◦ C/min: (a) quenched in water and (b) quenched in water–ice mixture.
Effect (III) is constituted by two exothermic overlapped peaks; this can be clearly distinguished from the differential curve of heat flow because the appearance of a well-defined shoulder. The first peak of the doublet (IIIA) has been explained considering two possibilities: it could be attributed to the formation of phase [26,27] and to the formation of S phase [28–30]. The second peak of the doublet (IIIB) has also several explanations in Al–Cu–Mg alloys: the formation of + S phases, only phase [30,31] or two distinct S phase precipitates [32]. However, it can be accepted that the first peak (IIIA) corresponds to the formation of S phase and the second one (IIIB) to the phase formation attending to the transformation temperatures (Table 2). The formation of S and phases was also confirmed by TEM; Fig. 7 shows a TEM image along the 1 0 0 zone axis of the matrix of AA2009/SiCp composite treated at 170 ◦ C for 20 h and the SAED pattern where S , and phases were identified. Effect (IV)
would be associated to the dissolution of the phases formed in (III). The fact that the presence of SiC whiskers transform the doublet observed in the particulate composite into a single exothermic effect (III) could be explained considering that the incorporation of SiCw accelerates the kinetic of precipitation of phase, and both exothermic effects (IIIA and IIIB) overlap to the same temperature. The height of the exothermic peak increases, which indicates a greater amount of formed phases. If both phases precipitate approximately at the same temperature, the height of the exothermic effect will be greater than that corresponding to two separated peaks. Finally, the two different precipitation sequences (1 and 2) which can occur, separately or simultaneously, in the Al–Cu–Mg alloys depending on the content in Cu and on the Cu/Mg relation are considered. The proportions of and S phases in the composite matrix should be approximately 2.1% and 5.3%, respectively,
Table 2 Temperature (in ◦ C) of exothermic transformations detected by DSC in parent composites. Material
Quenching media Water
AA2009/SiC/15w Water–ice Water AA2009/SiC/15p Water–ice
Heating rate (◦ C/min)
GPZ/GPBZ formation
IIIA effect
IIIB effect
10 30 10 30
68.2 88.8 73.4 88.8
264.8 292.7 261.4 273.3
– – 273.0 295.8
10 30 10 30
80.0 91.6 73.4 90.7
263.4 287.9 263.9 288.1
284.4 306.3 274.1 307.5
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the cooling rate during quenching in the precipitation kinetics depends on the morphology of the reinforcement. (4) Two exothermic transformations were found in both studied composites by means DSC curves. First of them is associated to the simultaneous formation of GP and GPB zones. The second one is constituted by exothermic doublets formed for the overlapped precipitation of S and phases. (5) The exothermic effect formed by the overlapping of S and phases precipitation is very sensitive to the cooling media used for quenching the AA2009/SiC/15w composite. Both exothermic transformations are not completely overlapped in specimens quenched in water–ice. However, the exothermic transformation did not form a doublet in water quenched specimens. Acknowledgements Authors wish to thank Ministerio de Educación y Ciencia for economical support present research (project MAT2004-06018) and Comunidad de Madrid through the program ESTRUMAT-CM (reference MAT/77). References
Fig. 7. AA2009/SiCp composite aged at 170 ◦ C for 20 h. Detail of the hardening precipitates, S , and phases in [0 0 1]␣ orientation.
considering the Cu (3.5 wt%) and Mg (1.1 wt%) content in AA2009. Small changes in the Cu/Mg ratio have a great influence in the amount of formed equilibrium phases and a local increase in the Mg concentration produce a considerable reduction of the amount of Al2 Cu formed. For that reason, the formation of Al7 Cu2 Fe precipitates detected only in the composite reinforced with particles could affect to the proportion of and S phases formed in this composite, favouring the formation of S because of the removing of the dissolved Cu from the matrix. This is in agreement with the doublet observed in the particle-reinforced material where the first peak is assigned to the formation of S . 4. Conclusions (1) Reinforcement morphology modifies the precipitation behaviour of the AA2009 matrix composites. AA2009/SiC/15p composite exhibits Fe-rich intermetallic phases, Al7 Cu2 Fe, and Mg2 Si precipitated close to the SiC/Al interfaces. This has an influence on the ageing kinetics of this material as the Cu/Mg rate is modified with respect to AA2009/SiC/15w composite. (2) Ageing kinetics of AA2009/SiC/15p composite are highly influenced by the ageing temperature. Times necessary to reach ageing peaks in this composite are larger than those observed in the material reinforced with 15 vol.% of SiC whiskers at 170 ◦ C. An increase in the ageing temperature up to 190 ◦ C reduces the ageing times more effectively in the case of the particlereinforced composite than in the whiskers one. (3) As the heating rate increases, the exothermic transformation is displaced towards higher temperatures. The influence of
[1] F. Ortega-Celaya, M.I. Pech-Canul, J. López-Cuevas, J.C. Rendón-Ángeles, M.A. Pech-Canul, J. Mater. Process. Technol. 183 (2007) 368–373. [2] A. Borrego, G. González-Doncel, Mater. Sci. Eng. A 245 (1998) 10–18. [3] J.T. Wang, M. Furukawa, Z. Horita, M. Nemoto, Y. Ma, T.G. Langdon, Metall. Mater. Trans. A 26 (1995) 581–587. [4] Sharmilee Pal, R. Mitra, V.V. Bhanuprasad, Mater. Sci. Eng. A 480 (2008) 496–505. [5] D. Ovono Ovono, I. Guillot, D. Massinon, J. Alloy Compd. 432 (2007) 241–246. [6] D. Bakavos, P.B. Prangnell, B. Bes, F. Eberl, Mater. Sci. Eng. A 491 (2008) 214– 223. [7] I.N. Khan, M.J. Starink, J.L. Yan, Mater. Sci. Eng. A 472 (2008) 66–74. [8] W. Sha, Mater. Design 28 (2007) 1524–1530. [9] Z. Gao, X. Zhang, M. Chena, J. Alloys Compd. (2008), doi:10.1016/j.jallcom. 2008.08.055. [10] S. Suresh, K.K. Chawla, in: S. Suresh, A. Mortensen, A. Needleman (Eds.), Fundamentals of Metal Matrix Composites, Butterworth-Heinemann, 1993, pp. 119–136. [11] P.I. Gouma, D.J. Lloyd, M.J. Mills, Mater. Sci. Eng. A 319–321 (2001) 439–442. [12] K.C. Chan, G.Q. Tong, L. Gao, J. Mater. Process. Technol. 74 (1998) 142–148. [13] B.L. Xiao, J.Z. Fan, X.F. Tian, W.Y. Zhang, L.K. Shi, J. Mater. Sci. 40 (2005) 5757–5762. [14] T.S. Srivatsan, Meslet Al-Hajri, C. Smith, M. Petraroli, Mater. Sci. Eng. A 346 (2003) 91–100. [15] T.S. Srivatsan, V. Meslet Al-Hajri, K. Vasudevan, Int. J. Fatigue 27 (2005) 357–371. [16] B.Q. Han, K.C. Chan, J. Mater. Sci. Lett. 16 (1997) 827–829. [17] A.P. Sannino, H.J. Rack, J. Mater. Sci. 30 (1995) 4316–4322. [18] A.L. Geiger, J.A. Walker, JOM 43 (1991) 8–15. [19] UNE-EN ISO 6507-1. European standard: Metallic materials. Vickers hardness test. Part I: Test method (ISO 6507-1:2005). [20] A. Charai, T. Walther, C. Alfonso, A.-M. Zahra, C.Y. Zahra, Acta Mater. 48 (10) (2000) 2751–2764. [21] S.P. Ringer, K. Hono, T. Saksai, I.J. Polmear, Scripta Mater. 36 (5) (1997) 517– 521. [22] L. Reich, S.P. Ringer, K. Hono, Phil. Mag. Lett. 79 (9) (1999) 639–648. [23] P. Ratchev, B. Verlinden, P. De Smet, P. Van Houtte, Acta Mater. 46 (10) (1998) 3523–3533. [24] A.M. Zahra, C.Y. Zahra, C. Alfonso, A. Charaï, Scripta Mater. 39 (11) (1998) 1553–1558. [25] A. Hayoune, D. Hamana, J. Alloys Compd. (2008), doi:10.1016/j.jallcom. 2008.06.070. [26] K.K. Chawla, A.H. Esmaeili, A.K. Datye, A.K. Vasudevan, Scripta Metall. Mater. 25 (1991) 1315–1319. [27] T.S. Kim, T.H. Kim, K.H. Oh, H.I. Lee, J. Mater. Sci. 27 (10) (1992) 2599–2605. [28] S. Arakawa, T. Hatayama, K. Matsugi, O. Yanagisawa, Scripta Mater. 42 (2000) 755–760. [29] V.K. Varma, Y.R. Mahajan, V.V. Kutumbarao, Scripta Mater. 37 (1997) 485– 489. [30] Z. Chen, Z. Zheng, Scripta Mater. 50 (2004) 1067–1071. [31] Z.Y. Ma, J. Bi, Y.X. Lu, M. Luo, Y.X. Gao, Scripta Metall. Mater. 29 (1993) 225–229. [32] S.C. Wang, M.J. Starink, Acta Mater. 55 (2007) 933–941.