- Email: [email protected]
Residual stress and damage development in the aluminium alloy EN AW-6061 particle reinforced with Al2O3 under thermal fatigue loading
Materials Science and Engineering A 501 (2009) 6–15
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Residual stress and damage development in the aluminium alloy EN AW-6061 particle reinforced with Al2 O3 under thermal fatigue loading Arne M. Klaska a , Tilmann Beck b,∗ , Alexander Wanner a , Detlef Löhe a a b
Universität Karlsruhe (TH), Institut für Werkstoffkunde I, Kaiserstraße 12, 76131 Karlsruhe, Germany Forschungszentrum Jülich, IEF-2, 52425 Jülich, Germany
a r t i c l e
i n f o
Article history: Received 7 January 2008 Received in revised form 25 July 2008 Accepted 6 October 2008 Keywords: Thermal fatigue Metal matrix composite Damage evolution Residual stresses Aluminium alloy EN AW-6061
a b s t r a c t Thermal fatigue (TF) tests were performed on the aluminium alloy EN AW-6061-T6, non-reinforced and reinforced with 15 and 22 vol.% Al2 O3 particles, respectively. The MMCs were produced by stir casting followed by hot extrusion. Thermal fatigue tests were performed using a 6 kW diode laser with a beam area of about 8 mm × 8 mm focused on the centre of one flat side of disc shaped specimens. The reverse side was either attached to a water-cooled aluminium plate or directly cooled by water. The maximum temperature Tmax of the irradiated side was varied between 573 and 773 K. The heating rate was 50 K/s. Residual stresses in the matrix alloy were measured by X-ray diffraction using the sin2 -method after T6 heat treatment and after defined temperature cycles. Initial residual compressive stresses between −20 and −65 MPa result from the machining processes before T6 heat treatment. During the first temperature cycle the residual stresses in all materials change to tension at almost all Tmax . The peak value of the residual stresses reaches 50 up to 65 MPa and is nearly independent from Tmax . Damage evolution was observed by light optical microscope and SEM after the same cycles as the residual stress measurements. Grain boundary reliefs arise and increase in all materials with increasing number of TF cycles and intergranular damage of EN AW-6061 and the matrix alloy of the MMCs is observed. Close to the particles, damage is more pronounced due to thermal and mechanical mismatch of the phases. Four mechanisms causing damage and residual stress development could be identified: thermally induced global deformation due to inhomogeneous distribution of temperature, thermally induced local deformation due to coefficient of thermal expansion (CTE) mismatch (different ˛th of both phases), mechanically induced local deformation due to different deformation behaviour of both phases and overaging. In the non-reinforced alloy global deformation is the dominant mechanism while in the MMCs also local mechanisms are significant. © 2008 Elsevier B.V. All rights reserved.
1. Introduction Today, Al-based MMCs are widely used in applications like brake discs and drums, pistons, structural components, etc. because of their high specific strength, high specific stiffness and low coefficient of thermal expansion (CTE) compared with the nonreinforced basic alloy [1]. Their mechanical properties under quasi-static as well as isothermal cyclic loading have been investigated extensively [2–4] and also creep behaviour of MMCs is well understood. Under isothermal cyclic loading particle reinforced MMCs may show strongly different lifetimes in stress and strain controlled tests. In stress controlled mode the fatigue strength of MMCs
∗ Corresponding author. Tel.: +49 2461 61 4425; fax: +49 2461 61 6464. E-mail addresses: [email protected] (A.M. Klaska), [email protected] (T. Beck), [email protected] (A. Wanner), [email protected] (D. Löhe). 0921-5093/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2008.10.009
(e.g. EN AW-6061-T6 reinforced with 20 vol.% Al2 O3 particles (6061/Al2 O3 /20p ) [4]) is usually higher as compared to the nonreinforced matrix alloy. However, the lifetime not only depends on particle volume fraction but also on chemical composition and size of reinforcements, strength of interfaces, and on the heat treatment state, grain size and dislocation density of the matrix alloy [3,5,6]. In strain controlled mode, the higher Young’s modulus of the MMCs causes larger stress amplitudes which result in a reduced LCF and HCF lifetime of particle reinforced MMCs [4,7,8]. Furthermore, the strength of reinforcements, interfacial precipitations, grain sizes and strength of interfaces influence the fatigue strength and lifetimes. However, the impact of interfacial strength on the fatigue strength is not completely understood up to now, because high as well as low interfacial strength may correlate with high fatigue strength [4,9,10]. Different papers of Chang et al., Mishra et al., Li and Langdon, and Kim et al. give an overview of creep mechanisms in MMCs concerning parameters like particle distribution, matrix structure, creep exponent and stress [11–16]. Usually, at higher loadings a
A.M. Klaska et al. / Materials Science and Engineering A 501 (2009) 6–15
threshold theory is applied regarding dislocation climbing as well as the development of Orowan rings between particles and the influence of relaxation of stresses at the interfaces [15,17,18]. In 6061/Al2 O3 /22p dislocation creep is the dominating creep mechanism in case of tensile creep at 573 K and 15–50 MPa. Dislocation climbing occurs at higher loadings (>50 MPa), grain boundary sliding is observed at higher temperatures and sufficiently small stresses. In the last case, the load transfer from the matrix to the particles is interfered resulting in higher creep rates of the MMC compared to the non-reinforced basic alloy. This is attributed to thermally activated development and annihilation of dislocations, accelerated precipitation kinetics of the matrix material, diffusion at interfaces and is further promoted by small matrix grains. In general, failure of 6061/Al2 O3 /22p under creep loading starts with delamination of interfaces. At creep under pressure, dislocation climbing is the dominant mechanism instead of dislocation slipping [19–21]. In 6061/Al2 O3 /10-22p , thermo-cyclic creep is observed at constant loadings and alternating temperature. The creep deformation decreases with increasing particle volume fraction and increasing matrix strength. This is caused by complex dislocation movement around particles resulting in asymmetric strain fields during heating and cooling. Cyclic creep is accelerated compared to isothermal creep due to cyclic initiation of primary creep. It is assumed, that elastic–plastic deformation at minimum temperature and interfacial diffusion at maximum temperature are further dominant mechanisms [22,23]. MMCs are widely used for different applications, e.g. brake discs and drums for trains (German ICE 2 and 3 [1]) or cars. Aluminiumbased metal matrix composites are also applied for pistons and many other engine parts. Mostly, SiC, Al2 O3 − and TiB2 ceramics are used for these reinforcements [1–29]. All of them undergo thermal fatigue (TF) loading during service. Nevertheless, the deformation, lifetime and damage behaviour during thermal fatigue loading including the complex interactions between reinforcement and matrix [9,30] still have to be determined in detail. These interactions result in additional mechanical loading due to the thermal mismatch between matrix and reinforcements (CTE mismatch) and therefore, cause enormous plastic deformations of the matrix close to the interfaces. Relaxation of the induced stress fields at high temperature due to creep processes is possible and further thermally induced loadings arise during cooling. This thermal fatigue loading frequently causes microcracks at brittle phases at the interfaces (e.g. spinel in Mg-rich matrix alloys) and afterwards, delamination of interfaces and decreasing mechanical strength [22,31,32]. Accordingly, the present work identifies and characterises the mechanisms of residual stress development and the damage evolution during TF loading in EN AW-6061-T6, non-reinforced (6061) and reinforced with 15 and 22 vol.% Al2 O3 particles (6061/Al2 O3 /15p , 6061/Al2 O3 /22p ).
7
Table 1 Parameters of T6 heat treatment. Material
Solution annealing
6061 6061/Al2 O3 /15p 6061/Al2 O3 /22p
Quenching
803 K/30 min 833 K/30 min
Artificial ageing 433 K/24 h
293 K/water
433 K/8 h
magnesium in the matrix material of the MMC due to MgA12 O4 spinel growth at the particle/matrix interfaces (Fig. 1) during T6 heat treatment. This is verified by SEM investigations [34]. Table 2 shows the chemical composition of 6061 and the MMCs matrix alloys. Extensive metallographic analysis of phases and microstructure (Figs. 2 and 3) revealed inhomogeneous grain size distributions (Figs. 4 and 5) and the formation of sub-grains in all materials (e.g. Fig. 2) and, furthermore, the tendency of reinforcements to agglomerate [34]. Disc shaped samples with a thickness of 10 and a diameter of 25 mm were taken from the extruded rods for TF testing. TF tests with maximum temperatures of 573, 673 and 773 K were performed by using a 6 kW diode laser. The beam was focused on the surface of the specimen with a spot size of about 8 mm × 8 mm (Fig. 6). The temperature at the specimen’s surface was measured by a pyrometer in one colour mode. The heating and cooling rates were set to 50 K/s by the means of an adaptive controller. In the first test series, a water-cooled aluminium plate was attached to the reverse side of the specimens (Fig. 7). In the second series, cooling was done by a water bath covering half of the specimens’ height (Fig. 8). After performing certain numbers of temperature cycles (0, 1, 10, 100, 1000 and in two cases up to 10,000) surface residual stress measurements have been done in the non-reinforced alloy and the matrices of the MMCs by means of the sin2 -method using 1 mm aperture and 15 distance angles (−70◦ ≤ ≥ 70◦ ) in an X-ray diffractometer class Karlsruhe. Table 3 shows the X-ray elastic constants for Cr K␣ radiation and the used accelerating voltage and current. The measurements were performed at the (3 1 1) matrix diffraction peak. Furthermore, the penetration depth of the X-ray has to be considered to ensure that the purely biaxial stresses existing at the surface of the MMCs are measured [35]. Table 4 gives the resulting penetration depths in aluminium and ␣-alumina. As the maximum depth is not higher than 10 m, which is significantly smaller than the particle spacing (e.g. Fig. 1), the assumption of a plane stress state is valid. The standard deviations are shown by
2. Materials and experimental procedure 6061/Al2 O3 /15p and 6061/Al2 O3 /22p were manufactured by Alcan Inc. (Canada) via the Duralcan stir casting process [33]. These MMCs as well as 6061 were extruded at the ARC Leichtmetallkompetenzzentrum Ranshofen GmbH (Austria) with a deformation ratio of ϕ = 4 and subsequently subjected to a T6 (Table 1) heat treatment. Chemical analysis of the matrix material by means of ICP1 gas analysis for the reinforced and of spark emission spectroscopy for the non-reinforced alloy showed a smaller fraction of
1
Inductive coupled plasma.
Fig. 1. Reinforcements afflicted with MgAl2 O4 spinel.
8
A.M. Klaska et al. / Materials Science and Engineering A 501 (2009) 6–15
Table 2 Chemical composition of EN AW-6061 after ICP and spark emission spectroscopy, respectively. Material
Chemical composition of the matrix alloys (wt.%)
6061 6061/Al2 O3 /15p 6061/Al2 O3 /22p
Al
Mg
Si
Cu
Cr
Fe
Ti
Mn
Ni
Bal. Bal. Bal.
0.83 0.43 0.23
0.641 0.62 0.66
0.288 0.42 0.21
0.088 0.212 0.188
0.047 0.14 0.16
0.022 0.013 0.004
0.002 0.15 0.04
<0.001 0.02 <0.001
Fig. 2. Overview of the microstructure of 6061 (cross section).
Fig. 3. Overview of the microstructure of 6061/Al2 O3 /15p (cross section).
error bars in Figs. 9–14 and are calculated by linear regressions from the 2–sin2 relation which is exemplarily shown in Fig. 15. The data points show a small scatter around the linear regression. This is attributed to limited grain statistics, whereas texture effects could not be detected. In order to consider possible surface textures, which might influence the residual stress measurements, texture analyses have been performed at all materials giving evidence that due to the low deformation ratio during extrusion surface textures are negligible in all test materials. Exclusively, residual stress measurements are performed in the non-reinforced alloy and the matrices of the MMCs. Due to a high intensity of the scattered radiation of the matrices, local residual stress measurement inside the reinforcements could not be obtained.
The damage evolution was investigated in 6061 (Tmax = 573 K) and the MMCs (Tmax = 573 and 773 K) under TF loading at the initially polished specimens’ surfaces by light optical and scanning electron microscope. Similar to the residual stress measurements the tests were stopped after certain numbers of cycles for metallographic investigation. 3. Results 3.1. Residual stresses Figs. 9–14 show the development of surface residual stresses measured in the matrix of the MMCs and the non-reinforced alloy during thermal fatigue tests at different maximum temperatures.
Fig. 4. Grain size distribution in cross section of 6061/Al2 O3 /15p .
A.M. Klaska et al. / Materials Science and Engineering A 501 (2009) 6–15
9
Fig. 5. Grain size distribution in 6061/Al2 O3 /22p .
Fig. 6. Dimension of heat spot and site on specimen.
For each maximum temperature the results are presented individually for the different modes of cooling to clarify the influence of the particle volume fraction on the development of residual stresses. After machining followed by T6 heat treatment residual compressive stresses are measured in all materials. Obviously, thermal cycling generally results in residual tensile stresses independent from Tmax . Furthermore, the development of residual stress levels is qualitatively similar in all materials, but is strongly influenced by Tmax .
Fig. 8. Test setup with water bath.
3.1.1. Tmax = 573 K As shown by the roughly similar development of residual stresses versus cycles in Figs. 9 and 10 the general deformation behaviour of both MMCs in TF loading is nearly independent from the mode of cooling. Between N = 10 and 100, the residual stresses in 6061/Al2 O3 /15p are nearly 0. Considerable residual tensile stresses occur at N = 1000. During further TF cycling the residual tensile stresses slightly decrease. In 6061/Al2 O3 /22p , residual tensile stresses already appear after N = 1 with a slight difference between Table 3 X-ray elastic constants of (3 1 1) lattice planes in aluminium and X-ray generator parameters. Young’s modulus (GPa) Accelerating voltage (kV) Current (mA)
69.13 0.353 40 50
Table 4 Penetration depth of Cr K␣ radiation in aluminium and ␣-alumina.
Fig. 7. Test setup with water-cooled aluminium plate.
Material
Penetration depths (2R = 112◦ → ˝ = 56◦ and 0◦ < < 70◦ )
Aluminium ␣-Alumina
3.5–10.3 m 3.7–10.8 m
10
A.M. Klaska et al. / Materials Science and Engineering A 501 (2009) 6–15
Fig. 9. Residual stresses measured after TF at Tmax = 573 K—plate cooling.
Fig. 11. Residual stresses measured after TF at Tmax = 673 K—plate cooling.
the two modes of cooling. Afterwards, the residual stress level keeps nearly constant (40–50 MPa) for the cooling plate (Fig. 9) to N = 1000 and slightly increases until N = 10 and then keeps that level up to N = 1000 for the water bath. However, the residual stresses in the non-reinforced alloy increase faster when the specimen is directly cooled with water (Fig. 10), where the highest level of residual tensile stresses is reached after the first cycle and then keeps-considering the scatter of the residual stress measurementnearly constant (25–40 MPa) during further TF loading. With the means of the cooling plate (Fig. 9), the residual stresses grow more slowly and reach the tensile regime between 10 and 100 TF cycles. Nevertheless, all materials are following the same trend by switching from residual compressive to tensile stresses during continued TF loading. In general, the highest amount of residual tensile stresses is obtained for 6061/Al2 O3 /22p , while the lowest residual stresses occur in 6061/Al2 O3 /15p . Added trend lines in Fig. 9 clarify increasing residual stresses in the non-reinforced alloy and 6061/Al2 O3 /15p whereas slightly decreasing residual stresses occur in 6061/Al2 O3 /22p between N = 1 and 1000. For the water bath (Fig. 10) increasing residual stresses in this cycle interval are measured for the MMCs and a nearly constant residual stress level is apparent in the case of 6061.
contrast to the results at Tmax = 573 K, a distinct difference between the modes of cooling appears for the MMCs. For both MMCs the residual tensile stresses are higher (up to 60 MPa) after cooling in the water bath (40–60 MPa) in Fig. 12, whereas the cooling plate (Fig. 11) results in residual stresses between 5 and 30 MPa. Contemporaneously, the residual tensile stresses slightly decrease after the first cycle in case of the water bath as well, whereas a slight increase is observed in case of the cooling plate. For the non-reinforced alloy cooled by the water bath, residual stresses increase to 65 MPa after 100 cycles, whereas cooling on the plate results in residual stresses of up to 45 MPa. Trend lines show slightly decreasing residual stresses for the non-reinforced material cooled by the plate (Fig. 11) and the MMCs cooled by the water bath (Fig. 12) within the measured range of TF cycles. However, increasing residual stresses are measured in the plate-cooled MMCs (Fig. 11) and in the non-reinforced alloy cooled by the water bath (Fig. 12).
3.1.2. Tmax = 673 K Residual stresses in all materials switch from compression (−20 to −50 MPa) to tension within the first TF cycle (Figs. 11 and 12). In
3.1.3. Tmax = 773 K In this case (Figs. 13 and 14), residual tensile stresses (10–40 MPa) appear after the first cycle for all materials. But in contrast to Tmax = 573 and 673 K, in the non-reinforced alloy no differences between the different modes of cooling are observed within the precision of measurement. Especially, after the first TF cycle the residual stresses are independent from particle volume fraction and mode of cooling. For the MMCs, higher residual tensile
Fig. 10. Residual stresses measured after TF at Tmax = 573 K—water bath cooling.
Fig. 12. Residual stresses measured after TF at Tmax = 673 K—water bath cooling.
A.M. Klaska et al. / Materials Science and Engineering A 501 (2009) 6–15
11
Fig. 13. Residual stresses measured after TF at Tmax = 773 K—plate cooling.
stress levels occur when the water bath is used for cooling. A systematic correlation between the particle volume fraction and the mode of cooling is not detectable. However, slightly higher residual stresses occur by using the water bath for cooling in all materials. Also, the trend lines in Figs. 13 and 14 clarify constant or slightly decreasing residual stresses for cooling with the aluminium plate (Fig. 13) and slightly increasing residual stresses for cooling with the water bath (Fig. 14). In addition, Fig. 14 indicates increasing residual stresses with increasing particle volume fraction.
Fig. 16. Polished surface of 6061/Al2 O3 /15p before TF cycling.
3.2. Damage evolution Starting from a flat, polished surface without any reliefs, neither from polishing nor from grain boundaries (Fig. 16, MMC and Fig. 17, non-reinforced), during the first TF cycle grain boundary reliefs become apparent in all investigated materials independent from Tmax (Fig. 18, 6061/Al2 O3 /15p ). The intensity of these grain boundary displacements grows with increasing number of TF cycles and with increasing particle volume fraction Vp . Fig. 19 exemplarily shows a light optical microscopy image of the surface of 6061/Al2 O3 /15p after 100 TF cycles at Tmax = 573 K. In addition, slip bands were observed after the first TF cycle especially in the nonreinforced alloy (Fig. 20) but also in matrix regions close to particles of the MMC-samples. In 6061 large-scale deformation of matrix regions near grain boundaries can also be observed, which are shown in Fig. 21 after 100 TF cycles. Up to this number of cycles, grain boundary slipping and matrix deformation increase faster in the MMCs than in the non-reinforced alloy (Figs. 19 and 21). After 1000 TF cycles, in the non-reinforced alloy some transgranular cracks as well as intergranular cracks inside an area of concentrated damage can be observed (Fig. 22). However, the overall level of slip bands and grain boundary reliefs keeps nearly constant between
Fig. 14. Residual stresses measured after TF at Tmax = 773 K—water bath cooling.
Fig. 15. 2 plotted vs. sin2
for 6061/Al2 O3 /22p , 573 K, N = 10, cooled in water.
Fig. 17. Polished surface of 6061 before TF cycling.
12
A.M. Klaska et al. / Materials Science and Engineering A 501 (2009) 6–15
Fig. 21. Surface of 6061 upon 100 TF cycles at Tmax = 573 K. Fig. 18. Grain boundary relief at surface of 6061/Al2 O3 /15p upon 1 TF cycle at Tmax = 573 K.
100 and 1000 TF cycles. In 6061/Al2 O3 /15p a few intergranular cracks occur after 1000 TF cycles (Fig. 23), which can clearly be identified after 10,000 cycles (Fig. 24) as well as delamination at the interfaces. No intergranular cracks are observed in 6061/Al2 O3 /22p , but a large number of particle fractures are detectable (Fig. 25). By increasing Tmax for the MMCs to 773 K the damage of the MMCs as well as overaging of the matrix alloy is accelerated. Massive damage already occurs after 100 TF cycles as shown by coarse precipitations inside the matrix grains as well as damage by pronounced grain boundary reliefs (Fig. 26), which are close to intergranular cracks. Particle failure (Fig. 26) also occurs in 6061/Al2 O3 /15p . Further SEM investigations showed massive interfacial displacement and delamination (Fig. 27). Further cycling results in interfacial and intergranular cracks. 3.3. Discussion TF loading results in four mechanisms influencing the development of residual stresses and damage: Fig. 19. Surface of 6061/Al2 O3 /15p upon 100 TF cycles at Tmax = 573 K.
Fig. 20. Surface of 6061 upon 1 TF cycle at Tmax = 573 K.
Fig. 22. Intergranular crack in 6061 upon 1000 TF cycles at Tmax = 573 K.
A.M. Klaska et al. / Materials Science and Engineering A 501 (2009) 6–15
Fig. 23. First intergranular cracks in 6061/Al2 O3 /15p upon 1000 TF cycles at Tmax = 573 K.
Fig. 24. Intergranular cracks and capacious damage in 6061/Al2 O3 /15p upon 10,000 TF cycles at Tmax = 573 K.
Fig. 25. Particle failure in 6061/Al2 O3 /22p upon 1000 TF cycles at Tmax = 573 K.
13
Fig. 26. Grain boundary relief, precipitations and particle fracture in 6061/Al2 O3 /15p upon 100 TF cycles at Tmax = 773 K.
(1) Thermally induced, global deformation as a result of inhomogeneous temperature distribution over large length scales within the specimen due to localized transient heating inside the beam focus of the laser and cooling of the specimen from the opposite side [22,32]. This results in macroscopic elastic–plastic strains in the whole specimen and residual tensile stresses after cooling inside the laser irradiated area due to plastic compression (in case of the MMCs only inside the matrices) during heating [36,37,9]. These deformations develop both in the MMC and in the non-reinforced alloy. This effect is enforced by higher temperature gradients inside the specimens during heating and cooling due to higher differences of local thermal strain and by higher Tmax due to increased plastic compression (of the matrices for the MMCs) and relaxation processes inside the aluminium alloy during heating. The effect of global thermally induced deformation is diminished by increasing Vp due to decreasing coefficient of thermal expansion of the composites which results in lower strain differences between areas of high and low temperature. Contemporaneously, increasing Vp yield steeper temperature gradients inside the specimens resulting in higher macroscopic residual stresses in the matrices. However, the last effect is at least partially compensated by simultaneously increased high temperature strength.
Fig. 27. Interfacial displacement in 6061/Al2 O3 /15p upon 100 TF cycles at Tmax = 773 K.
14
A.M. Klaska et al. / Materials Science and Engineering A 501 (2009) 6–15
(2) Overaging of the aluminium alloy, which decreases the amount of residual stresses by reducing the yield stress at room temperature and higher temperatures. This results in higher local plastic deformation at all temperatures and stress relaxation during heating. Therefore, a decrease in elastic strains and stresses occurs. Although, the measured residual stresses are well below the matrix’ yield stress, overaging results in a reduction of the residual stresses due to the integral measurement method which averages local stresses below and above the yield stress. The progress of overaging accelerates with increasing Vp and Tmax and this results in a faster decline of the amount of residual stresses [36–45]. (3) Thermally induced, local deformation inside the matrices of the MMCs near the interfaces due to CTE mismatch, which results in bulk-averaged residual tensile stresses in the matrix after cooling [36,37]. Going more into detail, high local residual tensile stresses near the interfaces are facing slight residual compressive stresses far from the interfaces [30]. The effect of thermally induced local deformation on the averaged residual stresses increases with increasing Vp due to an increasing amount of deformations adaptation between particles and matrix. Increasing Tmax shows a similar effect in consequence of generally higher thermally induced plastic deformations [9,10,25,35,46,47]. (4) Mechanically induced, local deformation occurring inside the matrices of the MMCs near the interfaces due to different stiffness of matrix and particles and different deformation behaviour [1,49,50]. These deformations directly depend on (1). Even in case of large plastic deformation of the composite, only the matrix of the MMCs is plastically deformed because the ceramic reinforcements only undergo elastic strains. Tensile deformation, which develops in the laser irradiated zone during cooling (1) results in local residual compressive stresses inside the matrix near the interfaces of the composite. Increasing Tmax escalates the plastic compression of the matrices during heating and yields increasing plastic tensile strains in the matrices during cooling. Also an increasing Vp rises the influence of this mechanism since less matrix volume is available in case of similar plastic deformation. Accordingly, both increasing Tmax and increasing Vp cause higher local residual compressive stresses in the matrices. The residual stress development in the matrix by (3) and (4) is strongly influenced by geometrical necessary dislocations [48]. For the non-reinforced alloy global deformation (1) and overaging (2) are the only relevant mechanisms because only macroscopic stresses and strains are obtained during TF cycling. Accordingly, increasing Tmax and direct water cooling result in a faster development of higher residual tensile stresses due to global deformation. On the other hand, overaging reduces the residual tensile stress level. Therefore, the difference between the modes of cooling, which is observable at Tmax = 573 and 673 K (Figs. 9–12), vanishes at Tmax = 773 K (Figs. 13 and 14). In case of the MMCs global deformation (1) and overaging (2) are superimposed by the local deformation mechanisms (3) and (4). Upon TF at Tmax = 573 K the larger high temperature strength of 6061/Al2 O3 /15p results in less pronounced global deformation and therefore in a slower increase of residual stresses than in the nonreinforced alloy. The slight decrease of residual stresses between 1000 and 10,000 cycles is due to overaging (2) of the matrix alloy. However, for 6061/Al2 O3 /22p also showing greater high temperature strength than 6061 the local deformation (3) obviously is very pronounced and residual tensile stresses are measured after the first TF cycle. The general development of the residual stresses
is nearly independent from the mode of cooling which strongly influences the global deformation (Figs. 9 and 10). Nevertheless, in the case of 6061/Al2 O3 /22p there is a remarkable difference in the residual stress level after the first TF cycle between plate and water bath cooling which cannot be thoroughly explained with the available experimental data. At 673 K a strong influence of the modes of cooling on the residual tensile stresses of the MMCs is observable (Figs. 11 and 12). Obviously, both thermally induced, global (1) and local deformations (3) determine the residual stress level. Again, for Tmax = 773 K the effect of these mechanisms cannot be separated (Figs. 13 and 14), because on the one hand, global deformation (1) increases due to higher temperature differences resulting in higher residual tensile stresses. The importance of this mechanism becomes apparent by the significant difference in the residual stress level for different modes of cooling. On the other hand, thermally induced, local deformation (3) is also increasing with Tmax . Starting from an almost stress-free state at Tmax , the surrounding matrix shrinks onto the reinforcements during cooling. This also results in locally high residual tensile stresses. At the same time, due to the high Tmax , overaging of the matrix alloy (2) reduces the residual stresses. This mechanism also seems to be responsible for the equalisation of the residual stress level in the MMCs and the non-reinforced alloy. The influence of mechanically induced, local deformation (4) on the development of residual stresses is very low, because no residual compressive stresses are observed during TF loading except at 573 K in 6061/Al2 O3 /15p during the first 100 cycles. Nevertheless this mechanism exists [48–50]. The damage development in 6061 is governed by global deformation (1). Because of large grain sizes, a huge fraction of this deformation occurs by dislocation movement which becomes obvious by slip lines (Fig. 21). However, grain boundary sliding also plays a major role (Fig. 22). In addition, diffusion processes result in overaging and clustering of Mg2 Si precipitations at grain boundaries. Finally, the notch effect of these precipitations contributes to intergranular cracking (Fig. 22). Also, the observed grain boundary relief in the MMCs (Fig. 23) is caused by (1). The stronger development of the grain boundary relief compared to 6061 is due to smaller grains which give raise to grain boundary slipping. In general, damage close to agglomerations of particles (Figs. 24 and 25) is more pronounced than far from agglomerations due to the notch effect of the particles. However, no damage is observed at the interfaces for Tmax = 573 K. In this case, the effect of the CTE mismatch obviously is not sufficient for inducing damage at the interfaces. Nevertheless, at higher maximum temperature (773 K) thermally (3) and mechanically induced local deformation result in sufficient stresses at the interfaces causing delamination of interfaces and particle fractures. In addition, slipping of grain boundaries close to agglomerations of particles also is responsible for more pronounced reliefs at Tmax = 773 K (Figs. 24 and 27) [20,21,51,52]. Particle fracture in 6061/Al2 O3 /22p occurs because of increasing influence of local deformations due to CTE mismatch (3) compared to the lower particle volume fraction. In addition to the increase in volume fraction, the higher fraction of large particles supports this mechanism. 4. Summary and conclusion Thermal fatigue tests were carried out by means of a 6 kW diode laser on the aluminium alloy 6061 and on the MMCs 6061/Al2 O3 /15p and 6061/Al2 O3 /22p . After defined temperature cycles, residual stresses were measured by X-ray diffraction and damage evolution was investigated by light optical microscope. The following results were obtained:
A.M. Klaska et al. / Materials Science and Engineering A 501 (2009) 6–15
1. Using a water bath for cooling during TF cycling of 6061 results in higher residual tensile stresses due to the larger temperature gradient than occurring when using the cooling plate. 2. Thermally induced global deformation (1) and overaging (2) are the dominant mechanisms in 6061 for residual stress development. 3. The most important mechanisms for residual stress development in the MMCs are thermally induced global (1) and local (3) deformation, which both result in residual tensile stresses. Overaging (2) only reduces the residual stress level, while the effect of mechanically induced local deformation (4) is negligible. 4. Concerning the residual stress development, 6061/Al2 O3 /15p shows the largest high temperature strength by slowest growth of residual stresses. 5. Damage starts with a grain boundary relief independent from material and maximum temperature. The relief is more pronounced close to particles in the MMCs due to notch effects and complex interactions between matrix and particles. 6. In 6061 the damage process ends with initiation of intergranular cracks usually starting at precipitations, while in the MMCs the damage development depends on particle volume fraction and maximum temperature. However, intergranular cracks are observed independent of maximum temperature, while delamination of interfaces occurs at higher Tmax . Acknowledgments The presented results were achieved within the scope of the project “Thermal and thermal mechanical fatigue loading of particle reinforced aluminium alloys”, which was funded by the German Research Foundation (DFG) under grant-#Lo 347/13-2. The authors thanks go to the DFG for financial support and the research partners at the Institute of Materials at the Ruhr-Universitaet Bochum Prof. M. Pohl and S. Heimann for their cooperation. References [1] K.U. Kainer, in: K.U. Kainer (Ed.), Metal Matrix Composites, Wiley-VCH, Weinheim, 2006, p. 54. [2] A. Miserez, A. Mortensen, Acta Mater. 52 (2004) 5331–5345. [3] H. Biermann, O. Hartmann, in: K.U. Kainer (Ed.), Metal Matrix Composites, Wiley-VCH, Weinheim, 2006, pp. 173–196. [4] C. Bosi, G.L. Garagnani, R. Tovo, M. Vedani, Int. J. Mater. Prod. Technol. 17 (3/4) (2002) 228–242. [5] Z.Z. Chen, K. Tokaji, Mater. Lett. 58 (2004) 2314–2321. [6] S.C. Tjong, G.S. Wang, Mater. Sci. Eng. A 386 (2004) 48–53. [7] O. Hartmann, K. Herrmann, H. Biermann, Adv. Eng. Mater. 6 (7) (2004) 477–485. [8] M. Papakyriacou, H.R. Mayer, S.E. Stanzl-Tschegg, M. Gröschl, Int. J. Fatigue 18 (7) (1996) 475–481. ˜ [9] E. Carreno-Morelli, Metal. Mater. Trans. A 35 (2004) 25–35. [10] H.M.A. Winand, A.F. Whitehouse, P.J. Withers, Mater. Sci. Eng. A 284 (2000) 103–113. [11] N. Chawla, K.K. Chawla, Metal Matrix Composites, Springer Science + Business Media, New York, 2006. [12] R.S. Mishra, JOM 51 (11) (1999) 65–68.
15
[13] R.S. Mishra, Scripta Metal. Mater. 2 (1992) 309–313. [14] R.S. Mishra, A.B. Pandey, A.K. Mukherjee, Mater. Sci. Eng. A 201 (1995) 205– 210. [15] Y. Li, T.G. Langdon, Acta Mater. 47 (12) (1999) 3395–3403. [16] W.J. Kim, Y.S. Lee, S.J. Moon, S.H. Hong, Mater. Sci. Technol. 16 (2000) 675–680. [17] S. Spigarelli, M. Cabibbo, E. Evangelista, T.G. Langdon, Mater. Sci. Eng. A 328 (2002) 39–47. [18] H. Dieringa, Y. Huang, P. Maier, N. Hort, K.U. Kainer, Mater. Sci. Eng. A 410–411 (2005) 85–88. [19] N. Matsuda, H. Tazawa, S. Ishikawa, M. Saitoh, Mater. Sci. Eng. A 319–321 (2001) 726–729. [20] G.C. Requena, D. Telfser, C. Hörist, H.P. Degischer, Mater. Sci. Technol. 8 (2002) 515–521. [21] C. Broeckmann, Kriechen partikelverstärkter metallischer Werkstoffe, Habilitationsschrift, Ruhr-Universität Bochum, 2001. [22] R.J. S Arsenault, Proceedings of the 26th International Symposium on Automotive Technology Automation (ISATA), Aachen, Germany, 1993, pp. 517–524. [23] A. Flaig, Thermal cycling creep of a fibre reinforced aluminum alloy, Dissertation, Universitaet Stuttgart, 2000. [24] E. Köhler, J. Niehues, B. Sommer, Giesserei-Praxis 5 (2005) 167–176. [25] Q.B. Ouyang, W.L. Wang, D. Zhang, G.D. Zhang, Key Eng. Mater. 313 (2006) 31–36. [26] O. Beffort, Feinstbearbeitung technischer Oberflächen-6, Internationales IWFKolloquium, Egerkingen, Schweiz, 2002. [27] K. Weinert, M. Buschka, M. Liedschulte, D. Biermann, U. Huber, J. Niehues, in: K.U. Kainer (Ed.), Metallische Verbundwerkstoffe, Wiley-VCH, Weinheim, 2003, pp. 207–212. [28] W. Henning, E. Köhler, Maschinenmarkt 101 (51) (1995) 44–47. [29] A.I. Nussbaum, Light Met. Age 2 (1997) 54–58. [30] C.G. Kang, J.H. Lee, S.W. Youn, J.K. Oh, J. Mater. Process. Technol. 166 (2005) 173–182. [31] I. Dutta, Acta Mater. 48 (2000) 1055–1074. [32] H. Sehitoglu, ASM Handbook 19, Fatigue and Fracture, vol. 4, ASM International, Ohio, 1996, pp. 527–556. [33] M. Kok, J. Mater. Process. Technol. 161 (2005) 381–387. [34] A.M. Klaska, A. Ohl, E. Kerscher, D. Löhe, Prakt. Metallogr. (2007) 33–51. [35] A.M. Klaska, T. Beck, A. Wanner, D. Löhe, Materialwiss. Werkstofftech. 37 (8) (2006) 637–649. [36] P. Fernández, R. Fernández, G. Gonzáles-Doncel, G. Bruno, Scripta Mater. 52 (2005) 793–797. [37] P. Fernández, G. Bruno, G. Gonzáles-Doncel, Compos. Sci. Technol. 66 (2006) 1738–1748. [38] F. Delmas, M.J. Casanove, P. Lours, A. Couret, A. Coujou, Mater. Sci. Eng. A 373 (2004) 80–89. [39] C. Cayron, P.A. Buffat, Acta Mater. 48 (2000) 2639–2653. [40] C.Y. Barlow, N. Hansen, Acta Metall. Mater. 39 (8) (1991) 1971–1979. [41] A. Kandil, Metall 58 (11) (2004) 721–726. [42] L. Wang, N. Shi, R.J. Arsenault, in: P.K. Liaw, J.R. Weertman, H.L. Marcus, J.S. Sauter (Eds.), Fine Symposium, 1991, pp. 129–149. [43] Y. Le, H. Wang, Y. Zhang, Z. Wang, J. Adv. Mater. 37 (3) (2005) 66–72. [44] W.M. Zhong, E. Goiffon, G. L’Espérance, M. Suéry, J.J. Blandin, Mater. Sci. Eng. A 214 (1996) 84–92. [45] S. Das, S. Das, K. Das, J. Mater. Sci. 41 (2006) 5402–5406. [46] H. Kehler, R. Bode, H. Haddenhorst, Metall 49 (3) (1995) 191–195. [47] G. Bruno, R. Fernández, G. González-Doncel, Mater. Sci. Forum 490–491 (2005) 539–544. [48] M.F. Ashby, in: A. Kelly, R.B. Nicholson (Eds.), Strengthening Methods in Crystals, Elsevier, Amsterdam, 1971, pp. 137–192. [49] T.W. Clyne, P.J. Withers, An Introduction to Metal Matrix Composites, University Press, Cambridge, 1993. [50] N. Hansen, C.Y. Barlow, in: S. Suresh, A. Mortensen, A. Needleman (Eds.), Fundamentals of Metal Matrix Composites, Butterworth-Heinemann, Stoneham, 1993, pp. 109–118. [51] Z.Y. Ma, S.C. Tjong, Compos. Sci. Technol. 59 (1999) 737–747. [52] G. Cseh, J. Bär, H.-J. Gudladt, J. Lendvai, A. Juhász, Mater. Sci. Eng. A 272 (1999) 145–151.
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Residual stress and damage development in the aluminium alloy EN AW-6061 particle reinforced with Al2 O3 under thermal fatigue loading Arne M. Klaska a , Tilmann Beck b,∗ , Alexander Wanner a , Detlef Löhe a a b
Universität Karlsruhe (TH), Institut für Werkstoffkunde I, Kaiserstraße 12, 76131 Karlsruhe, Germany Forschungszentrum Jülich, IEF-2, 52425 Jülich, Germany
a r t i c l e
i n f o
Article history: Received 7 January 2008 Received in revised form 25 July 2008 Accepted 6 October 2008 Keywords: Thermal fatigue Metal matrix composite Damage evolution Residual stresses Aluminium alloy EN AW-6061
a b s t r a c t Thermal fatigue (TF) tests were performed on the aluminium alloy EN AW-6061-T6, non-reinforced and reinforced with 15 and 22 vol.% Al2 O3 particles, respectively. The MMCs were produced by stir casting followed by hot extrusion. Thermal fatigue tests were performed using a 6 kW diode laser with a beam area of about 8 mm × 8 mm focused on the centre of one flat side of disc shaped specimens. The reverse side was either attached to a water-cooled aluminium plate or directly cooled by water. The maximum temperature Tmax of the irradiated side was varied between 573 and 773 K. The heating rate was 50 K/s. Residual stresses in the matrix alloy were measured by X-ray diffraction using the sin2 -method after T6 heat treatment and after defined temperature cycles. Initial residual compressive stresses between −20 and −65 MPa result from the machining processes before T6 heat treatment. During the first temperature cycle the residual stresses in all materials change to tension at almost all Tmax . The peak value of the residual stresses reaches 50 up to 65 MPa and is nearly independent from Tmax . Damage evolution was observed by light optical microscope and SEM after the same cycles as the residual stress measurements. Grain boundary reliefs arise and increase in all materials with increasing number of TF cycles and intergranular damage of EN AW-6061 and the matrix alloy of the MMCs is observed. Close to the particles, damage is more pronounced due to thermal and mechanical mismatch of the phases. Four mechanisms causing damage and residual stress development could be identified: thermally induced global deformation due to inhomogeneous distribution of temperature, thermally induced local deformation due to coefficient of thermal expansion (CTE) mismatch (different ˛th of both phases), mechanically induced local deformation due to different deformation behaviour of both phases and overaging. In the non-reinforced alloy global deformation is the dominant mechanism while in the MMCs also local mechanisms are significant. © 2008 Elsevier B.V. All rights reserved.
1. Introduction Today, Al-based MMCs are widely used in applications like brake discs and drums, pistons, structural components, etc. because of their high specific strength, high specific stiffness and low coefficient of thermal expansion (CTE) compared with the nonreinforced basic alloy [1]. Their mechanical properties under quasi-static as well as isothermal cyclic loading have been investigated extensively [2–4] and also creep behaviour of MMCs is well understood. Under isothermal cyclic loading particle reinforced MMCs may show strongly different lifetimes in stress and strain controlled tests. In stress controlled mode the fatigue strength of MMCs
∗ Corresponding author. Tel.: +49 2461 61 4425; fax: +49 2461 61 6464. E-mail addresses: [email protected] (A.M. Klaska), [email protected] (T. Beck), [email protected] (A. Wanner), [email protected] (D. Löhe). 0921-5093/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2008.10.009
(e.g. EN AW-6061-T6 reinforced with 20 vol.% Al2 O3 particles (6061/Al2 O3 /20p ) [4]) is usually higher as compared to the nonreinforced matrix alloy. However, the lifetime not only depends on particle volume fraction but also on chemical composition and size of reinforcements, strength of interfaces, and on the heat treatment state, grain size and dislocation density of the matrix alloy [3,5,6]. In strain controlled mode, the higher Young’s modulus of the MMCs causes larger stress amplitudes which result in a reduced LCF and HCF lifetime of particle reinforced MMCs [4,7,8]. Furthermore, the strength of reinforcements, interfacial precipitations, grain sizes and strength of interfaces influence the fatigue strength and lifetimes. However, the impact of interfacial strength on the fatigue strength is not completely understood up to now, because high as well as low interfacial strength may correlate with high fatigue strength [4,9,10]. Different papers of Chang et al., Mishra et al., Li and Langdon, and Kim et al. give an overview of creep mechanisms in MMCs concerning parameters like particle distribution, matrix structure, creep exponent and stress [11–16]. Usually, at higher loadings a
A.M. Klaska et al. / Materials Science and Engineering A 501 (2009) 6–15
threshold theory is applied regarding dislocation climbing as well as the development of Orowan rings between particles and the influence of relaxation of stresses at the interfaces [15,17,18]. In 6061/Al2 O3 /22p dislocation creep is the dominating creep mechanism in case of tensile creep at 573 K and 15–50 MPa. Dislocation climbing occurs at higher loadings (>50 MPa), grain boundary sliding is observed at higher temperatures and sufficiently small stresses. In the last case, the load transfer from the matrix to the particles is interfered resulting in higher creep rates of the MMC compared to the non-reinforced basic alloy. This is attributed to thermally activated development and annihilation of dislocations, accelerated precipitation kinetics of the matrix material, diffusion at interfaces and is further promoted by small matrix grains. In general, failure of 6061/Al2 O3 /22p under creep loading starts with delamination of interfaces. At creep under pressure, dislocation climbing is the dominant mechanism instead of dislocation slipping [19–21]. In 6061/Al2 O3 /10-22p , thermo-cyclic creep is observed at constant loadings and alternating temperature. The creep deformation decreases with increasing particle volume fraction and increasing matrix strength. This is caused by complex dislocation movement around particles resulting in asymmetric strain fields during heating and cooling. Cyclic creep is accelerated compared to isothermal creep due to cyclic initiation of primary creep. It is assumed, that elastic–plastic deformation at minimum temperature and interfacial diffusion at maximum temperature are further dominant mechanisms [22,23]. MMCs are widely used for different applications, e.g. brake discs and drums for trains (German ICE 2 and 3 [1]) or cars. Aluminiumbased metal matrix composites are also applied for pistons and many other engine parts. Mostly, SiC, Al2 O3 − and TiB2 ceramics are used for these reinforcements [1–29]. All of them undergo thermal fatigue (TF) loading during service. Nevertheless, the deformation, lifetime and damage behaviour during thermal fatigue loading including the complex interactions between reinforcement and matrix [9,30] still have to be determined in detail. These interactions result in additional mechanical loading due to the thermal mismatch between matrix and reinforcements (CTE mismatch) and therefore, cause enormous plastic deformations of the matrix close to the interfaces. Relaxation of the induced stress fields at high temperature due to creep processes is possible and further thermally induced loadings arise during cooling. This thermal fatigue loading frequently causes microcracks at brittle phases at the interfaces (e.g. spinel in Mg-rich matrix alloys) and afterwards, delamination of interfaces and decreasing mechanical strength [22,31,32]. Accordingly, the present work identifies and characterises the mechanisms of residual stress development and the damage evolution during TF loading in EN AW-6061-T6, non-reinforced (6061) and reinforced with 15 and 22 vol.% Al2 O3 particles (6061/Al2 O3 /15p , 6061/Al2 O3 /22p ).
7
Table 1 Parameters of T6 heat treatment. Material
Solution annealing
6061 6061/Al2 O3 /15p 6061/Al2 O3 /22p
Quenching
803 K/30 min 833 K/30 min
Artificial ageing 433 K/24 h
293 K/water
433 K/8 h
magnesium in the matrix material of the MMC due to MgA12 O4 spinel growth at the particle/matrix interfaces (Fig. 1) during T6 heat treatment. This is verified by SEM investigations [34]. Table 2 shows the chemical composition of 6061 and the MMCs matrix alloys. Extensive metallographic analysis of phases and microstructure (Figs. 2 and 3) revealed inhomogeneous grain size distributions (Figs. 4 and 5) and the formation of sub-grains in all materials (e.g. Fig. 2) and, furthermore, the tendency of reinforcements to agglomerate [34]. Disc shaped samples with a thickness of 10 and a diameter of 25 mm were taken from the extruded rods for TF testing. TF tests with maximum temperatures of 573, 673 and 773 K were performed by using a 6 kW diode laser. The beam was focused on the surface of the specimen with a spot size of about 8 mm × 8 mm (Fig. 6). The temperature at the specimen’s surface was measured by a pyrometer in one colour mode. The heating and cooling rates were set to 50 K/s by the means of an adaptive controller. In the first test series, a water-cooled aluminium plate was attached to the reverse side of the specimens (Fig. 7). In the second series, cooling was done by a water bath covering half of the specimens’ height (Fig. 8). After performing certain numbers of temperature cycles (0, 1, 10, 100, 1000 and in two cases up to 10,000) surface residual stress measurements have been done in the non-reinforced alloy and the matrices of the MMCs by means of the sin2 -method using 1 mm aperture and 15 distance angles (−70◦ ≤ ≥ 70◦ ) in an X-ray diffractometer class Karlsruhe. Table 3 shows the X-ray elastic constants for Cr K␣ radiation and the used accelerating voltage and current. The measurements were performed at the (3 1 1) matrix diffraction peak. Furthermore, the penetration depth of the X-ray has to be considered to ensure that the purely biaxial stresses existing at the surface of the MMCs are measured [35]. Table 4 gives the resulting penetration depths in aluminium and ␣-alumina. As the maximum depth is not higher than 10 m, which is significantly smaller than the particle spacing (e.g. Fig. 1), the assumption of a plane stress state is valid. The standard deviations are shown by
2. Materials and experimental procedure 6061/Al2 O3 /15p and 6061/Al2 O3 /22p were manufactured by Alcan Inc. (Canada) via the Duralcan stir casting process [33]. These MMCs as well as 6061 were extruded at the ARC Leichtmetallkompetenzzentrum Ranshofen GmbH (Austria) with a deformation ratio of ϕ = 4 and subsequently subjected to a T6 (Table 1) heat treatment. Chemical analysis of the matrix material by means of ICP1 gas analysis for the reinforced and of spark emission spectroscopy for the non-reinforced alloy showed a smaller fraction of
1
Inductive coupled plasma.
Fig. 1. Reinforcements afflicted with MgAl2 O4 spinel.
8
A.M. Klaska et al. / Materials Science and Engineering A 501 (2009) 6–15
Table 2 Chemical composition of EN AW-6061 after ICP and spark emission spectroscopy, respectively. Material
Chemical composition of the matrix alloys (wt.%)
6061 6061/Al2 O3 /15p 6061/Al2 O3 /22p
Al
Mg
Si
Cu
Cr
Fe
Ti
Mn
Ni
Bal. Bal. Bal.
0.83 0.43 0.23
0.641 0.62 0.66
0.288 0.42 0.21
0.088 0.212 0.188
0.047 0.14 0.16
0.022 0.013 0.004
0.002 0.15 0.04
<0.001 0.02 <0.001
Fig. 2. Overview of the microstructure of 6061 (cross section).
Fig. 3. Overview of the microstructure of 6061/Al2 O3 /15p (cross section).
error bars in Figs. 9–14 and are calculated by linear regressions from the 2–sin2 relation which is exemplarily shown in Fig. 15. The data points show a small scatter around the linear regression. This is attributed to limited grain statistics, whereas texture effects could not be detected. In order to consider possible surface textures, which might influence the residual stress measurements, texture analyses have been performed at all materials giving evidence that due to the low deformation ratio during extrusion surface textures are negligible in all test materials. Exclusively, residual stress measurements are performed in the non-reinforced alloy and the matrices of the MMCs. Due to a high intensity of the scattered radiation of the matrices, local residual stress measurement inside the reinforcements could not be obtained.
The damage evolution was investigated in 6061 (Tmax = 573 K) and the MMCs (Tmax = 573 and 773 K) under TF loading at the initially polished specimens’ surfaces by light optical and scanning electron microscope. Similar to the residual stress measurements the tests were stopped after certain numbers of cycles for metallographic investigation. 3. Results 3.1. Residual stresses Figs. 9–14 show the development of surface residual stresses measured in the matrix of the MMCs and the non-reinforced alloy during thermal fatigue tests at different maximum temperatures.
Fig. 4. Grain size distribution in cross section of 6061/Al2 O3 /15p .
A.M. Klaska et al. / Materials Science and Engineering A 501 (2009) 6–15
9
Fig. 5. Grain size distribution in 6061/Al2 O3 /22p .
Fig. 6. Dimension of heat spot and site on specimen.
For each maximum temperature the results are presented individually for the different modes of cooling to clarify the influence of the particle volume fraction on the development of residual stresses. After machining followed by T6 heat treatment residual compressive stresses are measured in all materials. Obviously, thermal cycling generally results in residual tensile stresses independent from Tmax . Furthermore, the development of residual stress levels is qualitatively similar in all materials, but is strongly influenced by Tmax .
Fig. 8. Test setup with water bath.
3.1.1. Tmax = 573 K As shown by the roughly similar development of residual stresses versus cycles in Figs. 9 and 10 the general deformation behaviour of both MMCs in TF loading is nearly independent from the mode of cooling. Between N = 10 and 100, the residual stresses in 6061/Al2 O3 /15p are nearly 0. Considerable residual tensile stresses occur at N = 1000. During further TF cycling the residual tensile stresses slightly decrease. In 6061/Al2 O3 /22p , residual tensile stresses already appear after N = 1 with a slight difference between Table 3 X-ray elastic constants of (3 1 1) lattice planes in aluminium and X-ray generator parameters. Young’s modulus (GPa) Accelerating voltage (kV) Current (mA)
69.13 0.353 40 50
Table 4 Penetration depth of Cr K␣ radiation in aluminium and ␣-alumina.
Fig. 7. Test setup with water-cooled aluminium plate.
Material
Penetration depths (2R = 112◦ → ˝ = 56◦ and 0◦ < < 70◦ )
Aluminium ␣-Alumina
3.5–10.3 m 3.7–10.8 m
10
A.M. Klaska et al. / Materials Science and Engineering A 501 (2009) 6–15
Fig. 9. Residual stresses measured after TF at Tmax = 573 K—plate cooling.
Fig. 11. Residual stresses measured after TF at Tmax = 673 K—plate cooling.
the two modes of cooling. Afterwards, the residual stress level keeps nearly constant (40–50 MPa) for the cooling plate (Fig. 9) to N = 1000 and slightly increases until N = 10 and then keeps that level up to N = 1000 for the water bath. However, the residual stresses in the non-reinforced alloy increase faster when the specimen is directly cooled with water (Fig. 10), where the highest level of residual tensile stresses is reached after the first cycle and then keeps-considering the scatter of the residual stress measurementnearly constant (25–40 MPa) during further TF loading. With the means of the cooling plate (Fig. 9), the residual stresses grow more slowly and reach the tensile regime between 10 and 100 TF cycles. Nevertheless, all materials are following the same trend by switching from residual compressive to tensile stresses during continued TF loading. In general, the highest amount of residual tensile stresses is obtained for 6061/Al2 O3 /22p , while the lowest residual stresses occur in 6061/Al2 O3 /15p . Added trend lines in Fig. 9 clarify increasing residual stresses in the non-reinforced alloy and 6061/Al2 O3 /15p whereas slightly decreasing residual stresses occur in 6061/Al2 O3 /22p between N = 1 and 1000. For the water bath (Fig. 10) increasing residual stresses in this cycle interval are measured for the MMCs and a nearly constant residual stress level is apparent in the case of 6061.
contrast to the results at Tmax = 573 K, a distinct difference between the modes of cooling appears for the MMCs. For both MMCs the residual tensile stresses are higher (up to 60 MPa) after cooling in the water bath (40–60 MPa) in Fig. 12, whereas the cooling plate (Fig. 11) results in residual stresses between 5 and 30 MPa. Contemporaneously, the residual tensile stresses slightly decrease after the first cycle in case of the water bath as well, whereas a slight increase is observed in case of the cooling plate. For the non-reinforced alloy cooled by the water bath, residual stresses increase to 65 MPa after 100 cycles, whereas cooling on the plate results in residual stresses of up to 45 MPa. Trend lines show slightly decreasing residual stresses for the non-reinforced material cooled by the plate (Fig. 11) and the MMCs cooled by the water bath (Fig. 12) within the measured range of TF cycles. However, increasing residual stresses are measured in the plate-cooled MMCs (Fig. 11) and in the non-reinforced alloy cooled by the water bath (Fig. 12).
3.1.2. Tmax = 673 K Residual stresses in all materials switch from compression (−20 to −50 MPa) to tension within the first TF cycle (Figs. 11 and 12). In
3.1.3. Tmax = 773 K In this case (Figs. 13 and 14), residual tensile stresses (10–40 MPa) appear after the first cycle for all materials. But in contrast to Tmax = 573 and 673 K, in the non-reinforced alloy no differences between the different modes of cooling are observed within the precision of measurement. Especially, after the first TF cycle the residual stresses are independent from particle volume fraction and mode of cooling. For the MMCs, higher residual tensile
Fig. 10. Residual stresses measured after TF at Tmax = 573 K—water bath cooling.
Fig. 12. Residual stresses measured after TF at Tmax = 673 K—water bath cooling.
A.M. Klaska et al. / Materials Science and Engineering A 501 (2009) 6–15
11
Fig. 13. Residual stresses measured after TF at Tmax = 773 K—plate cooling.
stress levels occur when the water bath is used for cooling. A systematic correlation between the particle volume fraction and the mode of cooling is not detectable. However, slightly higher residual stresses occur by using the water bath for cooling in all materials. Also, the trend lines in Figs. 13 and 14 clarify constant or slightly decreasing residual stresses for cooling with the aluminium plate (Fig. 13) and slightly increasing residual stresses for cooling with the water bath (Fig. 14). In addition, Fig. 14 indicates increasing residual stresses with increasing particle volume fraction.
Fig. 16. Polished surface of 6061/Al2 O3 /15p before TF cycling.
3.2. Damage evolution Starting from a flat, polished surface without any reliefs, neither from polishing nor from grain boundaries (Fig. 16, MMC and Fig. 17, non-reinforced), during the first TF cycle grain boundary reliefs become apparent in all investigated materials independent from Tmax (Fig. 18, 6061/Al2 O3 /15p ). The intensity of these grain boundary displacements grows with increasing number of TF cycles and with increasing particle volume fraction Vp . Fig. 19 exemplarily shows a light optical microscopy image of the surface of 6061/Al2 O3 /15p after 100 TF cycles at Tmax = 573 K. In addition, slip bands were observed after the first TF cycle especially in the nonreinforced alloy (Fig. 20) but also in matrix regions close to particles of the MMC-samples. In 6061 large-scale deformation of matrix regions near grain boundaries can also be observed, which are shown in Fig. 21 after 100 TF cycles. Up to this number of cycles, grain boundary slipping and matrix deformation increase faster in the MMCs than in the non-reinforced alloy (Figs. 19 and 21). After 1000 TF cycles, in the non-reinforced alloy some transgranular cracks as well as intergranular cracks inside an area of concentrated damage can be observed (Fig. 22). However, the overall level of slip bands and grain boundary reliefs keeps nearly constant between
Fig. 14. Residual stresses measured after TF at Tmax = 773 K—water bath cooling.
Fig. 15. 2 plotted vs. sin2
for 6061/Al2 O3 /22p , 573 K, N = 10, cooled in water.
Fig. 17. Polished surface of 6061 before TF cycling.
12
A.M. Klaska et al. / Materials Science and Engineering A 501 (2009) 6–15
Fig. 21. Surface of 6061 upon 100 TF cycles at Tmax = 573 K. Fig. 18. Grain boundary relief at surface of 6061/Al2 O3 /15p upon 1 TF cycle at Tmax = 573 K.
100 and 1000 TF cycles. In 6061/Al2 O3 /15p a few intergranular cracks occur after 1000 TF cycles (Fig. 23), which can clearly be identified after 10,000 cycles (Fig. 24) as well as delamination at the interfaces. No intergranular cracks are observed in 6061/Al2 O3 /22p , but a large number of particle fractures are detectable (Fig. 25). By increasing Tmax for the MMCs to 773 K the damage of the MMCs as well as overaging of the matrix alloy is accelerated. Massive damage already occurs after 100 TF cycles as shown by coarse precipitations inside the matrix grains as well as damage by pronounced grain boundary reliefs (Fig. 26), which are close to intergranular cracks. Particle failure (Fig. 26) also occurs in 6061/Al2 O3 /15p . Further SEM investigations showed massive interfacial displacement and delamination (Fig. 27). Further cycling results in interfacial and intergranular cracks. 3.3. Discussion TF loading results in four mechanisms influencing the development of residual stresses and damage: Fig. 19. Surface of 6061/Al2 O3 /15p upon 100 TF cycles at Tmax = 573 K.
Fig. 20. Surface of 6061 upon 1 TF cycle at Tmax = 573 K.
Fig. 22. Intergranular crack in 6061 upon 1000 TF cycles at Tmax = 573 K.
A.M. Klaska et al. / Materials Science and Engineering A 501 (2009) 6–15
Fig. 23. First intergranular cracks in 6061/Al2 O3 /15p upon 1000 TF cycles at Tmax = 573 K.
Fig. 24. Intergranular cracks and capacious damage in 6061/Al2 O3 /15p upon 10,000 TF cycles at Tmax = 573 K.
Fig. 25. Particle failure in 6061/Al2 O3 /22p upon 1000 TF cycles at Tmax = 573 K.
13
Fig. 26. Grain boundary relief, precipitations and particle fracture in 6061/Al2 O3 /15p upon 100 TF cycles at Tmax = 773 K.
(1) Thermally induced, global deformation as a result of inhomogeneous temperature distribution over large length scales within the specimen due to localized transient heating inside the beam focus of the laser and cooling of the specimen from the opposite side [22,32]. This results in macroscopic elastic–plastic strains in the whole specimen and residual tensile stresses after cooling inside the laser irradiated area due to plastic compression (in case of the MMCs only inside the matrices) during heating [36,37,9]. These deformations develop both in the MMC and in the non-reinforced alloy. This effect is enforced by higher temperature gradients inside the specimens during heating and cooling due to higher differences of local thermal strain and by higher Tmax due to increased plastic compression (of the matrices for the MMCs) and relaxation processes inside the aluminium alloy during heating. The effect of global thermally induced deformation is diminished by increasing Vp due to decreasing coefficient of thermal expansion of the composites which results in lower strain differences between areas of high and low temperature. Contemporaneously, increasing Vp yield steeper temperature gradients inside the specimens resulting in higher macroscopic residual stresses in the matrices. However, the last effect is at least partially compensated by simultaneously increased high temperature strength.
Fig. 27. Interfacial displacement in 6061/Al2 O3 /15p upon 100 TF cycles at Tmax = 773 K.
14
A.M. Klaska et al. / Materials Science and Engineering A 501 (2009) 6–15
(2) Overaging of the aluminium alloy, which decreases the amount of residual stresses by reducing the yield stress at room temperature and higher temperatures. This results in higher local plastic deformation at all temperatures and stress relaxation during heating. Therefore, a decrease in elastic strains and stresses occurs. Although, the measured residual stresses are well below the matrix’ yield stress, overaging results in a reduction of the residual stresses due to the integral measurement method which averages local stresses below and above the yield stress. The progress of overaging accelerates with increasing Vp and Tmax and this results in a faster decline of the amount of residual stresses [36–45]. (3) Thermally induced, local deformation inside the matrices of the MMCs near the interfaces due to CTE mismatch, which results in bulk-averaged residual tensile stresses in the matrix after cooling [36,37]. Going more into detail, high local residual tensile stresses near the interfaces are facing slight residual compressive stresses far from the interfaces [30]. The effect of thermally induced local deformation on the averaged residual stresses increases with increasing Vp due to an increasing amount of deformations adaptation between particles and matrix. Increasing Tmax shows a similar effect in consequence of generally higher thermally induced plastic deformations [9,10,25,35,46,47]. (4) Mechanically induced, local deformation occurring inside the matrices of the MMCs near the interfaces due to different stiffness of matrix and particles and different deformation behaviour [1,49,50]. These deformations directly depend on (1). Even in case of large plastic deformation of the composite, only the matrix of the MMCs is plastically deformed because the ceramic reinforcements only undergo elastic strains. Tensile deformation, which develops in the laser irradiated zone during cooling (1) results in local residual compressive stresses inside the matrix near the interfaces of the composite. Increasing Tmax escalates the plastic compression of the matrices during heating and yields increasing plastic tensile strains in the matrices during cooling. Also an increasing Vp rises the influence of this mechanism since less matrix volume is available in case of similar plastic deformation. Accordingly, both increasing Tmax and increasing Vp cause higher local residual compressive stresses in the matrices. The residual stress development in the matrix by (3) and (4) is strongly influenced by geometrical necessary dislocations [48]. For the non-reinforced alloy global deformation (1) and overaging (2) are the only relevant mechanisms because only macroscopic stresses and strains are obtained during TF cycling. Accordingly, increasing Tmax and direct water cooling result in a faster development of higher residual tensile stresses due to global deformation. On the other hand, overaging reduces the residual tensile stress level. Therefore, the difference between the modes of cooling, which is observable at Tmax = 573 and 673 K (Figs. 9–12), vanishes at Tmax = 773 K (Figs. 13 and 14). In case of the MMCs global deformation (1) and overaging (2) are superimposed by the local deformation mechanisms (3) and (4). Upon TF at Tmax = 573 K the larger high temperature strength of 6061/Al2 O3 /15p results in less pronounced global deformation and therefore in a slower increase of residual stresses than in the nonreinforced alloy. The slight decrease of residual stresses between 1000 and 10,000 cycles is due to overaging (2) of the matrix alloy. However, for 6061/Al2 O3 /22p also showing greater high temperature strength than 6061 the local deformation (3) obviously is very pronounced and residual tensile stresses are measured after the first TF cycle. The general development of the residual stresses
is nearly independent from the mode of cooling which strongly influences the global deformation (Figs. 9 and 10). Nevertheless, in the case of 6061/Al2 O3 /22p there is a remarkable difference in the residual stress level after the first TF cycle between plate and water bath cooling which cannot be thoroughly explained with the available experimental data. At 673 K a strong influence of the modes of cooling on the residual tensile stresses of the MMCs is observable (Figs. 11 and 12). Obviously, both thermally induced, global (1) and local deformations (3) determine the residual stress level. Again, for Tmax = 773 K the effect of these mechanisms cannot be separated (Figs. 13 and 14), because on the one hand, global deformation (1) increases due to higher temperature differences resulting in higher residual tensile stresses. The importance of this mechanism becomes apparent by the significant difference in the residual stress level for different modes of cooling. On the other hand, thermally induced, local deformation (3) is also increasing with Tmax . Starting from an almost stress-free state at Tmax , the surrounding matrix shrinks onto the reinforcements during cooling. This also results in locally high residual tensile stresses. At the same time, due to the high Tmax , overaging of the matrix alloy (2) reduces the residual stresses. This mechanism also seems to be responsible for the equalisation of the residual stress level in the MMCs and the non-reinforced alloy. The influence of mechanically induced, local deformation (4) on the development of residual stresses is very low, because no residual compressive stresses are observed during TF loading except at 573 K in 6061/Al2 O3 /15p during the first 100 cycles. Nevertheless this mechanism exists [48–50]. The damage development in 6061 is governed by global deformation (1). Because of large grain sizes, a huge fraction of this deformation occurs by dislocation movement which becomes obvious by slip lines (Fig. 21). However, grain boundary sliding also plays a major role (Fig. 22). In addition, diffusion processes result in overaging and clustering of Mg2 Si precipitations at grain boundaries. Finally, the notch effect of these precipitations contributes to intergranular cracking (Fig. 22). Also, the observed grain boundary relief in the MMCs (Fig. 23) is caused by (1). The stronger development of the grain boundary relief compared to 6061 is due to smaller grains which give raise to grain boundary slipping. In general, damage close to agglomerations of particles (Figs. 24 and 25) is more pronounced than far from agglomerations due to the notch effect of the particles. However, no damage is observed at the interfaces for Tmax = 573 K. In this case, the effect of the CTE mismatch obviously is not sufficient for inducing damage at the interfaces. Nevertheless, at higher maximum temperature (773 K) thermally (3) and mechanically induced local deformation result in sufficient stresses at the interfaces causing delamination of interfaces and particle fractures. In addition, slipping of grain boundaries close to agglomerations of particles also is responsible for more pronounced reliefs at Tmax = 773 K (Figs. 24 and 27) [20,21,51,52]. Particle fracture in 6061/Al2 O3 /22p occurs because of increasing influence of local deformations due to CTE mismatch (3) compared to the lower particle volume fraction. In addition to the increase in volume fraction, the higher fraction of large particles supports this mechanism. 4. Summary and conclusion Thermal fatigue tests were carried out by means of a 6 kW diode laser on the aluminium alloy 6061 and on the MMCs 6061/Al2 O3 /15p and 6061/Al2 O3 /22p . After defined temperature cycles, residual stresses were measured by X-ray diffraction and damage evolution was investigated by light optical microscope. The following results were obtained:
A.M. Klaska et al. / Materials Science and Engineering A 501 (2009) 6–15
1. Using a water bath for cooling during TF cycling of 6061 results in higher residual tensile stresses due to the larger temperature gradient than occurring when using the cooling plate. 2. Thermally induced global deformation (1) and overaging (2) are the dominant mechanisms in 6061 for residual stress development. 3. The most important mechanisms for residual stress development in the MMCs are thermally induced global (1) and local (3) deformation, which both result in residual tensile stresses. Overaging (2) only reduces the residual stress level, while the effect of mechanically induced local deformation (4) is negligible. 4. Concerning the residual stress development, 6061/Al2 O3 /15p shows the largest high temperature strength by slowest growth of residual stresses. 5. Damage starts with a grain boundary relief independent from material and maximum temperature. The relief is more pronounced close to particles in the MMCs due to notch effects and complex interactions between matrix and particles. 6. In 6061 the damage process ends with initiation of intergranular cracks usually starting at precipitations, while in the MMCs the damage development depends on particle volume fraction and maximum temperature. However, intergranular cracks are observed independent of maximum temperature, while delamination of interfaces occurs at higher Tmax . Acknowledgments The presented results were achieved within the scope of the project “Thermal and thermal mechanical fatigue loading of particle reinforced aluminium alloys”, which was funded by the German Research Foundation (DFG) under grant-#Lo 347/13-2. The authors thanks go to the DFG for financial support and the research partners at the Institute of Materials at the Ruhr-Universitaet Bochum Prof. M. Pohl and S. Heimann for their cooperation. References [1] K.U. Kainer, in: K.U. Kainer (Ed.), Metal Matrix Composites, Wiley-VCH, Weinheim, 2006, p. 54. [2] A. Miserez, A. Mortensen, Acta Mater. 52 (2004) 5331–5345. [3] H. Biermann, O. Hartmann, in: K.U. Kainer (Ed.), Metal Matrix Composites, Wiley-VCH, Weinheim, 2006, pp. 173–196. [4] C. Bosi, G.L. Garagnani, R. Tovo, M. Vedani, Int. J. Mater. Prod. Technol. 17 (3/4) (2002) 228–242. [5] Z.Z. Chen, K. Tokaji, Mater. Lett. 58 (2004) 2314–2321. [6] S.C. Tjong, G.S. Wang, Mater. Sci. Eng. A 386 (2004) 48–53. [7] O. Hartmann, K. Herrmann, H. Biermann, Adv. Eng. Mater. 6 (7) (2004) 477–485. [8] M. Papakyriacou, H.R. Mayer, S.E. Stanzl-Tschegg, M. Gröschl, Int. J. Fatigue 18 (7) (1996) 475–481. ˜ [9] E. Carreno-Morelli, Metal. Mater. Trans. A 35 (2004) 25–35. [10] H.M.A. Winand, A.F. Whitehouse, P.J. Withers, Mater. Sci. Eng. A 284 (2000) 103–113. [11] N. Chawla, K.K. Chawla, Metal Matrix Composites, Springer Science + Business Media, New York, 2006. [12] R.S. Mishra, JOM 51 (11) (1999) 65–68.
15
[13] R.S. Mishra, Scripta Metal. Mater. 2 (1992) 309–313. [14] R.S. Mishra, A.B. Pandey, A.K. Mukherjee, Mater. Sci. Eng. A 201 (1995) 205– 210. [15] Y. Li, T.G. Langdon, Acta Mater. 47 (12) (1999) 3395–3403. [16] W.J. Kim, Y.S. Lee, S.J. Moon, S.H. Hong, Mater. Sci. Technol. 16 (2000) 675–680. [17] S. Spigarelli, M. Cabibbo, E. Evangelista, T.G. Langdon, Mater. Sci. Eng. A 328 (2002) 39–47. [18] H. Dieringa, Y. Huang, P. Maier, N. Hort, K.U. Kainer, Mater. Sci. Eng. A 410–411 (2005) 85–88. [19] N. Matsuda, H. Tazawa, S. Ishikawa, M. Saitoh, Mater. Sci. Eng. A 319–321 (2001) 726–729. [20] G.C. Requena, D. Telfser, C. Hörist, H.P. Degischer, Mater. Sci. Technol. 8 (2002) 515–521. [21] C. Broeckmann, Kriechen partikelverstärkter metallischer Werkstoffe, Habilitationsschrift, Ruhr-Universität Bochum, 2001. [22] R.J. S Arsenault, Proceedings of the 26th International Symposium on Automotive Technology Automation (ISATA), Aachen, Germany, 1993, pp. 517–524. [23] A. Flaig, Thermal cycling creep of a fibre reinforced aluminum alloy, Dissertation, Universitaet Stuttgart, 2000. [24] E. Köhler, J. Niehues, B. Sommer, Giesserei-Praxis 5 (2005) 167–176. [25] Q.B. Ouyang, W.L. Wang, D. Zhang, G.D. Zhang, Key Eng. Mater. 313 (2006) 31–36. [26] O. Beffort, Feinstbearbeitung technischer Oberflächen-6, Internationales IWFKolloquium, Egerkingen, Schweiz, 2002. [27] K. Weinert, M. Buschka, M. Liedschulte, D. Biermann, U. Huber, J. Niehues, in: K.U. Kainer (Ed.), Metallische Verbundwerkstoffe, Wiley-VCH, Weinheim, 2003, pp. 207–212. [28] W. Henning, E. Köhler, Maschinenmarkt 101 (51) (1995) 44–47. [29] A.I. Nussbaum, Light Met. Age 2 (1997) 54–58. [30] C.G. Kang, J.H. Lee, S.W. Youn, J.K. Oh, J. Mater. Process. Technol. 166 (2005) 173–182. [31] I. Dutta, Acta Mater. 48 (2000) 1055–1074. [32] H. Sehitoglu, ASM Handbook 19, Fatigue and Fracture, vol. 4, ASM International, Ohio, 1996, pp. 527–556. [33] M. Kok, J. Mater. Process. Technol. 161 (2005) 381–387. [34] A.M. Klaska, A. Ohl, E. Kerscher, D. Löhe, Prakt. Metallogr. (2007) 33–51. [35] A.M. Klaska, T. Beck, A. Wanner, D. Löhe, Materialwiss. Werkstofftech. 37 (8) (2006) 637–649. [36] P. Fernández, R. Fernández, G. Gonzáles-Doncel, G. Bruno, Scripta Mater. 52 (2005) 793–797. [37] P. Fernández, G. Bruno, G. Gonzáles-Doncel, Compos. Sci. Technol. 66 (2006) 1738–1748. [38] F. Delmas, M.J. Casanove, P. Lours, A. Couret, A. Coujou, Mater. Sci. Eng. A 373 (2004) 80–89. [39] C. Cayron, P.A. Buffat, Acta Mater. 48 (2000) 2639–2653. [40] C.Y. Barlow, N. Hansen, Acta Metall. Mater. 39 (8) (1991) 1971–1979. [41] A. Kandil, Metall 58 (11) (2004) 721–726. [42] L. Wang, N. Shi, R.J. Arsenault, in: P.K. Liaw, J.R. Weertman, H.L. Marcus, J.S. Sauter (Eds.), Fine Symposium, 1991, pp. 129–149. [43] Y. Le, H. Wang, Y. Zhang, Z. Wang, J. Adv. Mater. 37 (3) (2005) 66–72. [44] W.M. Zhong, E. Goiffon, G. L’Espérance, M. Suéry, J.J. Blandin, Mater. Sci. Eng. A 214 (1996) 84–92. [45] S. Das, S. Das, K. Das, J. Mater. Sci. 41 (2006) 5402–5406. [46] H. Kehler, R. Bode, H. Haddenhorst, Metall 49 (3) (1995) 191–195. [47] G. Bruno, R. Fernández, G. González-Doncel, Mater. Sci. Forum 490–491 (2005) 539–544. [48] M.F. Ashby, in: A. Kelly, R.B. Nicholson (Eds.), Strengthening Methods in Crystals, Elsevier, Amsterdam, 1971, pp. 137–192. [49] T.W. Clyne, P.J. Withers, An Introduction to Metal Matrix Composites, University Press, Cambridge, 1993. [50] N. Hansen, C.Y. Barlow, in: S. Suresh, A. Mortensen, A. Needleman (Eds.), Fundamentals of Metal Matrix Composites, Butterworth-Heinemann, Stoneham, 1993, pp. 109–118. [51] Z.Y. Ma, S.C. Tjong, Compos. Sci. Technol. 59 (1999) 737–747. [52] G. Cseh, J. Bär, H.-J. Gudladt, J. Lendvai, A. Juhász, Mater. Sci. Eng. A 272 (1999) 145–151.