Influence of the alumina thickness at the interfaces on the fracture mechanisms of aluminium multilayer composites

Materials Science and Engineering A 496 (2008) 133–142

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Influence of the alumina thickness at the interfaces on the fracture mechanisms of aluminium multilayer composites a,∗ ´ ˜ a C.M. Cepeda-Jimenez , M. Pozuelo b , J.M. Garc´ıa-Infanta a , O.A. Ruano a , F. Carreno a b

Department of Physical Metallurgy, CENIM, CSIC, Av. Gregorio del Amo 8, 28040 Madrid, Spain Department of Materials Science and Engineering, 6531-G Boelter Hall, University of California, Los Angeles, CA 90095-1595, USA

a r t i c l e

i n f o

Article history: Received 20 February 2008 Received in revised form 30 April 2008 Accepted 5 May 2008 Keywords: Multilayer aluminium laminates Fracture mechanisms Delamination Interfacial strength Hot roll-bonding Alumina thickness

a b s t r a c t Three aluminium multilayer composites based in fifteen alternate layers of Al 6082 and pure Al 1050 have been produced by hot roll-bonding. Alumina layers of three different thicknesses were grown at the interfaces by an anodizing process. The influence of alumina thickness on the toughness of the multilayer laminates has been studied. These laminate composites have been tested at room temperature under Charpy impact test, three point bending test at low strain rate and shear tests on the interfaces. All three laminates exhibited higher impact toughness than the as-received Al 6082 alloy being highest for the one with the thicker alumina layer. The fracture mechanism of the laminate materials was shown to depend on the alumina thickness and on the imposed strain rate during the mechanical tests. At low strain rate, intrinsic toughening mechanisms were dominant, while at high strain rate extrinsic mechanisms like delamination were also activated increasing the impact toughness. © 2008 Elsevier B.V. All rights reserved.

1. Introduction Laminate metal composites (LMCs) consist of alternating metal or reinforced metal layers that are bonded with “sharp” interfaces. LMCs can dramatically improve many properties including fracture toughness, fatigue behaviour, impact behaviour, wear, corrosion, and damping capacity or they provide enhanced formability or ductility [1–5]. In many respects, impact toughness is the most interesting property to be enhanced because internal interfaces between layers in multilayer materials limit crack propagation through several mechanisms [6–11]. From the economical point of view, it is also worth to combine expensive materials with other low cost materials. Nowadays, the application of high strength age hardenable aluminium alloys is growing throughout the aerospace industry. Aluminium alloys are selected for their optimal combination of physical and mechanical properties [12]. One advantage is that light-metal components lead to an overall reduced weight and, thus, to reduced energy consumption. Another advantage, which may be just as important from an environmental point of view, is the fact that aluminium components may be recycled with relatively low energy demands [13].

∗ Corresponding author. Tel.: +34 91 5538900; fax: +34 91 5347425. ´ E-mail address: [email protected] (C.M. Cepeda-Jimenez). 0921-5093/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2008.05.015

Al–Mg–Si alloys have been studied extensively because of their technological importance and their exceptional increase in strength obtained by precipitation hardening [14–17]. The alloy selected in the present study, with the designation 6082, is optimized to be the strongest alloy in the family. The ease of shaping, low density and good weldability along with the low cost are factors that make this alloy commercially very attractive. Therefore, it can be of interest to obtain multilayer laminate materials combining Al 6082 alloy with pure aluminium by hot rolling, to improve toughness [6,7]. Since pure aluminium is exceptionally ductile, enhanced impact toughness can be achieved by an intrinsic toughening mechanism, or by an extrinsic toughness mechanism like interface delamination [18]. Furthermore, since pure aluminium is relatively inexpensive the multilayer material obtained can be economically very attractive. Joining of aluminium has some limitations. Most of these limitations arise from the formation of the protective oxide film (Al2 O3 ), which covers the aluminium-based materials [19]. This oxide film is spontaneously formed on aluminium when it is exposed in air. By an anodizing process, which is a classical approach to increase the tribological properties of aluminium alloys, alumina coatings are claimed to grow on the pure aluminium surface. This process allows studying the influence of the alumina thickness on the mechanical properties of the processed multilayer material. The alumina brittleness and the cracks induced during processing can affect the mechanical strength of the interfaces favouring delamination as a principal mechanism of crack arresting

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Table 1 Chemical composition of as-received aluminium alloys (weight percent) Alloy

Si

Fe

Cu

Mn

Mg

Cr

Zn

Ti

6082 “E” 1050 “H”

0.75 0.12

0.34 0.23

0.053 <0.005

0.48 0.036

0.86 <0.010

0.019 <0.010

0.032 <0.010

0.075 ± 0.005 0.017 ± 0.002

in multilayer materials, and thus may improve the extrinsic toughening. In this respect, the purpose of this work is to fabricate pure Al/Al–Mg–Si multilayer composites with different alumina-coating thickness at the interfaces and to study its influence on the impact toughness and the fracture mechanisms. 2. Experimental procedure 2.1. Materials and processing The aluminium alloys used in the present study were rolled Al–Mg–Si 6082-T6 alloy (termed “E”) and commercial pure Al 1050H24 (termed “H”) sheets of 2 mm in thickness. The initial width and length of the sheets was 60 and 100 mm, respectively. The composition in weight percentage of the alloys is included in Table 1 and some mechanical properties are summarized in Table 2. The as-received aluminium alloys present an extremely fine oxide film of approximately 2.5 nm when the product is recently manufactured [20]. This spontaneous aluminium oxide layer will be named “native alumina” in this study. Additionally, some Al 1050 sheets were anodized in 15 wt% sulphuric acid under an applied voltage of 20 V for two different times (30 s and 5 min) to generate two extra oxide film thicknesses. Prior to anodizing, the samples were cleaned with acetone. To ensure a constant and homogeneous temperature throughout the solution, forced convection is provided by electrolyte stirring. These oxide layers were formed at low electrolyte temperature (∼5 ◦ C) favouring rapid growth and reduced dissolution of the oxide, a process known as hard anodizing. The thickness of the alumina layers obtained by the anodizing process was determined by scanning electron microscopy analysis of the oxide cross-sections. Average and standard deviations of about 40 measurements were calculated. The alumina-coating thickness on anodized Al 1050 for 30 s and 5 min was about 0.30 ± 0.10 and 1.84 ± 0.40 ␮m respectively. After anodizing, three multilayer composites were considered as a function of the alumina layer thickness on the Al 1050 alloy. Each composite was a stack of 15 layers constituted by eight Al 6082 alloy layers and seven Al 1050 layers with different alumina-coating thickness. They were 30 mm thick and referenced in this work as AEH15 (with native alumina coating), AEH15-30s (anodized Al 1050 for 30 s) and AEH15-5min (anodized Al 1050 for 5 min). The stacked aluminium materials were welded by Tungsten Inert Gas (TIG) at their edges to avoid oxygen penetration and delamination during processing, and then hot-rolled at 465 ◦ C in

Table 2 Mechanical properties of as-received aluminium alloys Alloy

UTSa (MPa)

YSa (MPa)

HV

Elongationa (%)

6082-T6 “E” 1050-H24 “H”

315 105

255 75

108 44

10 10

UTS: ultimate tensile strength; Y: yield stress; HV: Vickers Hardness; T6: solution treating followed by quenching and finally age hardening; H24: work hardening followed by partial annealing (240 ◦ C). a Data provided by the alloy maker from tensile tests.

several passes without lubrication. The diameter of the rolls was 131 mm and the rolling speed was 346 mm/s. All rolling passes were parallel to the rolling direction of the as-received sheets. Without changing the rolling direction, the rolling process was repeated up to five cycles of four passes each, of about 4–8% reduction per pass with the sample being reheated at 465 ◦ C between series. The total reduction in thickness was 2.8:1, corresponding to an equivalent strain of ε = 1.2 (von Mises criterion). The resulting hot-rolled samples were in the form of a plate, of thickness about 11 mm, length about 340 mm and width about 60 mm. The average thickness of the aluminium layers in the AEH15 laminates was ∼725 ␮m. After hot rolling, and due to the high temperatures employed during the processing, it was necessary to carry out a heat treatment to improve the mechanical properties of the hardenable aluminium alloy included in the laminated material. The heat treatment that has been deemed optimal for the 6082 alloy is the T6 temper. This heat treatment involves solution treating the alloy at 530 ◦ C for 30 min, followed by rapid quenching in water and finally age hardening at 185 ◦ C for 6 h, which was applied to all rolled samples. 2.2. Microstructural determination Microstructure at the bond interfaces in the longitudinal– transversal (LT) orientation was observed by scanning electron microscopy (SEM) using a JEOL JSM 6500F instrument with field emission gun. The chemical compositions of aluminium alloys and laminate interfaces were examined by an electron probe microanalyzer (Oxford Inca) operating at 15 kV. Metallographic preparation involved methods of standard surface preparation. The specimens were electropolished in a 30% nitric acid solution in methanol at −25 ◦ C and 15 V. 2.3. Mechanical tests 2.3.1. Microhardness test Microhardness measurements were carried out around the laminate interfaces with a Vickers indenter with a load of 100 g during 15 s. Vickers microhardness values vs. distance to the interface were represented in order to observe the hardness gradient across the interface. The indentation sizes ranged between 12 ␮m for the Al 6082 alloy and 22 ␮m for the pure Al 1050. The distance to the interface was measured from the indentation centre using image analysis software. 2.3.2. Charpy test Two mm V-notched Charpy type testing specimens were mechanized with 10 mm × 10 mm × 55 mm dimensions from as-received monolithic Al 6082 (E) and Al 1050 (H), and AEH15 laminate materials containing different alumina-coating thicknesses. The samples were tested in the crack-arrester orientation. The notch was machined to end at an individual layer of the test sample such that the crack front advances through each layer interface sequentially during the test. Charpy impact testing was conducted on a pendulum impact tester using a maximum capacity of 294 J. The velocity of the striker in the impact instant was of 5.4 m/s, and the strain rate approximately 1.5 × 102 s−1 . Three samples of each material were tested.

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2.3.3. Three point bending test Bending tests have been selected because they deliver much more useful information about the mechanical properties of the composite laminates than the Charpy impact tests. The bending test allows determination of yield and maximum stresses, a graphical visualization of layer fracture and the different fracture mechanisms as testing proceeds, and finally, comparison of toughness values by comparison of the areas inside the stress–strain curves. For comparison with the Charpy impact test, three point bending tests were performed using two mm V-notched Charpy type specimens (10 mm × 10 mm × 55 mm) in the crack-arrester orientation, the loading span being 40 mm like in the Charpy test. This configuration is the most interesting, both technically and scientifically, to study the different fracture mechanisms operating during the bending test. The stress, , and the strain, ε, were converted from the recorded raw data according to the following relations [21]: =

3pl , 2ae2

ε=

6ed , l2

(1)

where a is the initial width of the sample, e is the initial thickness of the sample minus the V-notch depth, l is the span length between the supports, p is the force applied on the sample and d is the midspan displacement of the sample. Then, the nominal stress vs. nominal strain was represented in order to characterize the mechanical response to layer fracture and crack propagation across the laminate. Bending tests were performed using a Servosis universal test machine with load cell and time recorded by the data acquisition program and under displacement control at a rate of 0.04 mm/s, which corresponds to a strain rate about 1.2 × 10−3 s−1 , being five orders of magnitude lower than the strain rate in the Charpy test. At least two samples for each laminate were used to collect data. Fracture surface of selected specimens was examined by both macroscopic analysis and scanning electron microscopy to evaluate deformation micromechanisms and any interlayer debonding. 2.3.4. Shear test The bonding of aluminium surfaces is the critical step in the current process. The interface mechanical properties were measured by shear tests in a Servosis universal test machine (cross-head speed = 0.005 mm/s) using specimens of approximate dimensions 10 mm × 10 mm × 3 mm. The sample was clamped between two metal supports and the interface to test is located just outside the border of the tool and parallel to the load direction. Then, a square punch at a given gap distance is used to apply the shear load until failure of the interface. The shear stress, , and the shear strain, , are given by the expressions [22]: =

p , ae

 = tan(˛) =

d , lgap

(2)

where a, e, p and d have been already defined, ˛ is the shear angle and lgap is the span length between the supports and the mobile punch, corresponding to 0.35 mm in this study. Once the shear tests were carried out, the fracture surfaces were analyzed using SEM to assess more precisely the type of failure of the bonded layers. 3. Results and discussion 3.1. Microstructure The SEM microstructures were observed in the LT orientation. The microstructure of the as-received Al 6082 rolled sheet in the LT orientation is presented in Fig. 1a. The as-received material shows large grains (20–30 ␮m) that are slightly elongated and

Fig. 1. SEM micrographs showing the microstructure of LT orientation in as-received (a) Al 6082-T6 (E) and (b) Al 1050-H24 (H).

flattened parallel to the rolling direction. The Al 6082 alloy contains dispersoids and inclusions. The dispersoids were principally of AlMnSi phase (0.1 ␮m) and inclusions were mainly a mixture of ␤-AlFeSi and ␣-AlFeMnSi intermetallic phases (1–10 ␮m) distributed at grain boundaries, in addition to coarse Mg2 Si [23]. It is likely that during homogenization of the alloy, at about 540 ◦ C, the transformation of the ␤-AlFeSi phase into a more spheroidal ␣Al(FeMn)Si phase may occur. Additionally, the very fine dispersed precipitates shown in Fig. 1a are particles of ␤-Mg2 Si phase. Historically, alloys of the 6xxx series are said to have excess of silicon relative to Mg2 Si, due to the presumption that the pre-␤ hardening phase also has this composition. In this concept, the remaining Si is proposed to be present in other phases, such as Al–Fe–Si and Al–Si–Fe–Mn particles. In recent years, however, it has become evident that the composition of the ␤ phase contains Mg and Si in a ratio close to 1 rather than 2 [24]. The ideal composition of the ␤ phase is Mg5 Si6 , a phase with just the right ratio of Mg/Si to explain the excess amount of Si that optimizes the strength in 6082. Additionally, Fig. 1b shows the microstructure of the 1050 aluminium alloy received in the H24 condition, which was work hardened followed by partial annealing at 240 ◦ C. The as-received material presents an equiaxed grain structure with an average grain size of 2–3 ␮m. Additionally, insoluble iron-rich intermetallic particles were observed to be randomly distributed in the wrought sheet and ranged in size between 0.5 and 5 ␮m.

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Fig. 2. SEM micrographs at high magnification showing interfaces in the aluminium multilayer composites: (a) interface 3 in AEH15; (b) interface 2 in AEH15-30s; (c) interface 3 in AEH15-5min.

fracturing [25]. Therefore, the oxide layer from the surface is broken up into particles, or platelets aligned parallel to the rolling direction. Due to the cracks that open up in the alumina coating, the aluminium flows through the fractured alumina regions. The interface, therefore, is a combination of oxide fragments and bonded areas of “extruded” aluminium. Consequently, the cracking of the alumina coatings allows metal–metal contact and roll-bonding to take place. The spacing between the alumina fragments was somewhat variable and dependent on the initial alumina-coating thickness (Fig. 2). It has been observed that the spacing between fractured alumina platelets increase with the oxide film thickness. Also, it can be observed that thinner alumina coatings of AEH15 and AEH15-30s laminates (Fig. 2a and b) are broken into smaller fragments during the processing than for the thicker coatings (Fig. 2c). Mean length of alumina fragments measured by SEM was ∼2.8 and ∼17 ␮m for the AEH15-30s and AEH15-5min laminate respectively, which corresponds to an aspect ratio of 9.2 and 9.5 respectively. Le et al. [26] studied oxide films on Al 1200 and Al 1050 rolled surface for alumina thicknesses below 1.5 ␮m (including nano-scale films), and developed a model for cracking of the oxide layer during rolling. They predicted that the oxide segment length after rolling is directly proportional to the film thickness, with an aspect ratio of oxide segment length to film thickness approximately equal to 10, being small the effect of the rolling reduction on crack spacing. This is in good agreement with the aspect ratio of alumina fragments presented in this study. Furthermore, if it is assumed that the oxide is undeformable during the processing and the plate volume is conserved during deformation, then the new interfacial area due to the extruded aluminium between the cracked alumina must be equal for the three laminates, and only a function of the reduction in plate thickness. Additionally, the distribution of this new interfacial area would vary as a function of the alumina-coating thickness. Therefore, the final interfaces are made up of oxide fragments and newly generated fresh areas of extruded aluminium, which surface area fraction is the same for the three laminates. However, the spacing between the alumina platelets increases with the initial oxide film thickness and its distribution will play an important role in the mechanical properties of the interfaces. Finally, the microstructure of the Al 6082 (E) alloy in the three laminates is refined after rolling. It consists of subgrains finer than 2–3 ␮m contained in the elongated initial grains (Fig. 2). Fine precipitates, apparently Mg2 Si, can be observed homogeneously distributed both in grain boundaries and inside the grains. On the contrary, the microstructure of the Al 1050 (H) alloy (for instance, bottom layer in Fig. 2a) is coarser than that in the as-received state (Fig. 1b), due to the high temperatures reached during processing and the absence of precipitates inhibiting recrystallization. Fig. 2 shows also the presence of a grain-size gradient in the 1050 alloy, being this finer close to the interfaces. These grain-size gradients are attributed to diffusion of elements due to the high temperature and strain localization close to the interfaces during processing. 3.2. Mechanical tests

Fig. 2 shows the SEM micrographs of different interfaces in the AEH15 laminates with different alumina-coating thicknesses. The micrographs suggest a good bond since debonding or voids were not observed. However, further assessment of the integrity requires quantitative mechanical testing. White and bright particles, identified as Al2 O3 by microanalysis, are observed at the interfaces. The dark shadow present in the interface of the AEH15-5min laminate (Fig. 2c) is an image artefact due to repulsion of electrons. During rolling, the metal plastically deforms and extends, but the aluminium oxide is brittle and can respond to stress only by

3.2.1. Microhardness test According to microhardness measurements a hardness gradient has been observed across the laminates interfaces (Fig. 3). This gradient is mainly due to the diffusion of alloying elements (i.e. Si and Mg) as a result of high temperature and pressure during processing. The element diffusion distance is about 50 ␮m into the 6082 aluminium alloy (at the left side of the interface), and about 30 ␮m into the pure aluminium layer (at the right side of the interface). This gradient of microhardness values is similar in the three lam-

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Fig. 3. Vickers microhardness (100 g; 15 s) of the aluminium multilayer composites as a function of the distance to the interface.

inates considered and all studied interfaces. The higher extension of the hardness gradient into the Al 6082 alloy could be due to the fact that the pure aluminium is more easily extruded through the microcracks towards the Al 6082 than in the opposite direction. Additionally, the average hardness value for the Al 6082 alloy far from the interface was 120 HV, which is 11% higher than the hardness value for the as-received Al 6082 alloy. This increase in mechanical strength is attributed to grain refinement according to the Hall–Petch relationship. The presence of alloying elements in the Al 6082, that are able to precipitate, hinder grain growth during processing. On the contrary, the average hardness value for the Al 1050 alloy was 37 HV, being lower than that for the as-received Al 1050 alloy, which is attributed to an increase in its grain size due to the high processing temperatures. Finally, the slight difference in microhardness values between the laminates far from the interfaces, especially for the Al 6082, can be due to small differences in temperature during the processing or the thermal treatment of the composites. 3.2.2. Charpy test The results of Charpy impact test at room temperature in the crack-arrester orientation are reported in Table 3. The Charpy Vnotched (CVN) energy average value for the monolithic Al 6082 (E) was 360 kJ/m2 , while for the monolithic Al 1050 (H) was 333 kJ/m2 . The three laminates provide improved impact behaviour compared to monolithic materials. The improvement was found to increase significantly with the alumina layer thickness in the laminates. The impact value of the AEH15 and AEH15-30s laminates is 24 and 43% respectively higher than that for the as-received Al 6082 alloy. Finally, the Charpy energy value of the AEH15-5min laminate is 110% higher than that for the as-received Al 6082 alloy. Macrographs of Charpy-tested samples (Fig. 4) for the three laminates show different fracture behaviour. The macrographs illustrate the high ductility of the three laminate composites under Charpy impact loads as evidenced by extensive necking in the layers. The Table 3 Charpy V-notched (CVN) (kJ/m2 ) energy of as-received and laminate materials Material

CVN energy (kJ/m2 )

6082-T6 “E” 1050-H24 “H” AEH15 AEH15-30s AEH15-5min

360 333 447 513 757

Fig. 4. Macrographs of Charpy-tested fractured samples of the aluminium multilayer composites: (a) AEH15; (b) AEH15-30s; (c) AEH15-5min.

AEH15 laminate (containing native alumina layer) does not present debonding, despite extensive plastic strain of the aluminium layers, and the crack propagates without interruption sequentially in adjoining layers. Therefore, the increase in impact toughness for the AEH15 laminate can be due to the intrinsic toughening contribution from both the Al 6082 with higher hardness than in the as-received state, and the pure aluminium layers which were submitted to high temperatures during processing, resulting in recovery and elimination of internal stresses. On the contrary, AEH15-30s laminate shows several short delaminations between different layers, indi-

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cating that crack interruption and subsequent renucleation in the following layer can be easily produced. Finally, the macrograph corresponding to the Charpy-tested AEH15-5min laminate (Fig. 4c) shows several long delaminations with extensive plastic deformation of the aluminium layers, which enlarge the required energy to induce crack renucleation in the following layer. It is observed that the longer the extent of the debonding at the interface of the layers, the larger is the impact toughness. It can be concluded that the impact toughness increases with the alumina-coating thickness because the debonding of large alumina platelets produces high energy release, redirecting the crack perpendicularly towards the neighbour alumina particles and forcing their decohesion along the interface. A chain decohesion of alumina particles finally produces delamination. 3.2.3. Three point bending test Fig. 5 shows stress–strain curves obtained from three point bending tests for the as-received aluminium alloys and roll-bonded laminate materials in the crack-arrester orientation. This test provides additional information about crack propagation across the laminate and both intrinsic and extrinsic toughening mechanisms. The Al 6082 alloy presents a bending stress of 595 MPa and moderate ductility. In contrast, the pure aluminium presents low strength (237 MPa) but excellent plasticity. On the other hand, a maximum bending stress of 360 MPa was obtained for the AEH15 laminate. This value is slightly lower than that calculated by the rule of mixtures (372 MPa) due to the recovered microstructure and the increase in the grain size for the pure aluminium layers. Although the curve shape is very similar for the three laminates considered in this research, a small scatter in maximum strength is observed which is attributed to the notch position in the laminate, and also to the slight difference in temperature during the processing. It is interesting to note that the curves corresponding to the three laminate materials present higher ductility than the as-received Al 6082 alloy. The ductility of laminate composites cannot be predicted by the rule of mixtures. This is because the ductility of laminates is dependent on many variables such as the susceptibility of the less ductile layer to cracking, the contribution of interface cracking, the ease of delamination, the influence of layer thickness and the number of layers. Additionally, the bending curves for the three laminates show a similar pattern, with several and progressive load drops corresponding to cracking of the different layers. The pronounced stress

Fig. 6. Macrographs of bending tested fractured samples of the aluminium multilayer composites: (a) AEH15; (b) AEH15-30s; (c) AEH15-5min.

Fig. 5. Three point bending test of as-received aluminium alloys and of the aluminium multilayer composites with different alumina thickness in the interface.

drops of these laminates correspond to propagation of cracks in the Al 6082 layers until the crack is gradually arrested in the ductile layer by an intrinsic toughening mechanism. This causes the observed decrease in the slope of the curve after the drop. In general, and for the three laminates, the crack extends progressively across the laminate composite similarly; therefore no clear influence of the alumina-coating thickness on the bending mechanical behaviour has been found. Moreover, the area inside the –ε curve is associated with the material toughness. The area inside the bending curves of the AEH15, AEH15-30s and AEH15-5min laminates is 82, 40 and 108% higher, respectively, than that for the as-received Al

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6082 alloy. Therefore, these results differ from the trend obtained by Charpy impact test, where the impact toughness increases as a function of the alumina thickness, being 24, 41 and 110% higher, respectively, than that for the as-received Al 6082 alloy. Hence, other fracture mechanisms could be operating depending on the alumina thickness. Fig. 6 shows macrographs corresponding to fracture surfaces in bending of the AEH15 laminates with different alumina-coating thickness at the interfaces. Again, the laminate containing the native alumina layer (AEH15) does not present debonding, and extensive plastic deformation of the pure aluminium layer can be observed (Fig. 6a). The laminates containing an anodizing alumina layer (AEH15-30s in Fig. 6b and AEH15-5min in Fig. 6c) show much shorter delaminations than the laminates tested at high strain rate (Charpy tests). Therefore, the type of failure obtained with the laminate materials tested in bending at low strain rate is different from that obtained at high strain rate (105 times higher), being influenced by the operating deformation and fracture mechanisms. An analysis of fracture at higher magnification is given in Fig. 7, which shows fracture surfaces of the AEH15-30s (Fig. 7a) and AEH15-5min (Fig. 7b) laminates tested in bending in the arrester orientation. The AEH15-30s laminate shows evidence of interface debonding together with extensive plastic tearing of the pure aluminium layer. At this stage, deformation bands are apparent in the pure aluminium. On the contrary, the micrograph for the AEH15-5min laminate (Fig. 7b) shows equidistant joint points alternated with debonding patches due to alumina fragments decohesion. The pure aluminium layers in

AEH15-5min sample also present extensive plastic deformation according to its inherent intrinsic toughness that decreases the crack opening stress. Therefore, under low strain rate, the AEH155min laminate presents interfaces more difficult to debond than the AEH15-30s material, showing multiple bonded areas after testing. Fig. 8 shows SEM micrographs at low magnification of AEH1530s (Fig. 8a) and AEH15-5min (Fig. 8b) samples subjected to interrupted bending tests at 12% nominal strain. It can be observed that for the same deformation, debonding in the next interface to the notch position in the AEH15-30s laminate has been produced (crack bridging mechanism) while the interfaces in the AEH155min are undamaged. Therefore, it is confirmed that in spite of a thicker alumina coating, the AEH15-5min present stronger interfaces under low strain rate tests. It can be concluded that the mechanical behaviour of the interfaces depends on the strain rate during testing. The results have shown, for the AEH15-5min laminate, delamination (extrinsic mechanism) at high strain rate (Charpy test) and lack of delamination (intrinsic mechanism) at low strain rate (bending test). Additionally, the mechanical behaviour also depends on the spacing between the alumina platelets which in turn depends on the alumina thickness at a given rolling strain. It is our contention that the change from an intrinsic to an extrinsic mechanism depends on the presence of particles that generate microcracks able to coalesce. The coalescence of these microcracks depends on the distance between particles and on the applied strain rate.

Fig. 7. SEM micrographs at low magnification showing bending tested fractured surfaces: (a) AEH15-30s; (b) AEH15-5min. The white arrows indicate the applied load direction during the bending test, while the grey arrows point out the fragmented alumina layer.

Fig. 8. SEM micrographs at low magnification of bending tested samples subjected to an interrupted bending test at 12% nominal strain: (a) AEH15-30s; (b) AEH155min. The white arrows indicate the applied load direction during the bending test.

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Fig. 9. Shear test of as-received aluminium alloys and on different interfaces of the aluminium multilayer composites with different alumina thickness.

At high strain rates, the laminate deforms as a result of the Charpy pendulum impact. The local deformation leads to nucleation and growth of cracks. These cracks can propagate straight across the laminate, as in the AEH15 laminate, or ultimately may result in crack deflection at the interfaces, as observed experimentally for AEH15-5min, and in less extent for AEH15-30s. Brittle particles, such as alumina fragments present in the laminates

materials, supply the initiation sites for microcracks nucleation. Debonding of the big alumina platelets for the AEH15-5min laminate produces a high released energy, forcing the debonding of neighbour alumina platelets and resulting in delamination along the interface. In contrast, debonding of the nanometric alumina particles in the AEH15 laminate, alternated with many bonded areas, does not affect crack propagation across the laminate. Therefore, at high strain rate the failure mechanism on the interfaces comprises the nucleation and coalescence of debonded alumina particles ahead of the crack front, which depending on their size and distribution may result in delamination, improving the material toughness by an extrinsic fracture mechanism. At low strain rates the failure mechanism may be different depending on the alumina thickness of the laminates. Although extrinsic mechanisms may still operate, intrinsic fracture mechanisms are favoured at these low strain rates due to the increase in plasticity of the constituent materials. The behaviour of the AEH15 with nanometric size alumina particles is similar at different strain rates, but an increase in plasticity is found at low strain rates. Furthermore, the AEH15-5min laminate behaves similar to the laminate with native alumina (AEH15). This behaviour, however, is different to that observed at high strain rate where delamination occurred. This is because a large area between the alumina fragments exists allowing blunting of the microcracks, avoiding their coalescence and therefore delamination. In this way, the crack propagates in the stress direction across the laminate being arrested only by intrinsic fracture mechanisms. In contrast to the other two laminates, the AEH15-30s laminate showed microdelaminations

Fig. 10. SEM micrographs showing fractured surfaces from shear tests of (a) Al 6082 and (b) Al 1050 in AEH15-30s laminate; and (c) Al 6082 and (d) Al 1050 in AEH15-5min laminate.

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characteristic of an extrinsic fracture mechanism. In this case, particle decohesion along the interface could be produced because the short spacing between the alumina particles was smaller than the zone needed for crack blunting, inducing microdelaminations (Figs. 7a and 8a). Therefore, it is our contention that delamination occurs only for a narrow range of interparticular distances. It is worth noting that the compensated effect of shorter microdelaminations at low strain rate (lower extrinsic toughening than at high strain rates), and the higher ductility of the constituent materials (higher intrinsic toughening than at high strain rates), results in similar toughness increase in the AEH15-30s laminate at both strain rate conditions. This different mechanical behaviour for the laminate materials depending on the strain rate can be advantageous both for impact loads (high strain rate) and for fatigue applications (low strain rate). 3.2.4. Shear tests Finally, shear tests have been performed to characterize the mechanical properties of the interfaces (Fig. 9). In the description that follows, the interfaces are assigned numbers to indicate their location in the laminate (for example, i3 means the third interface from surface). The maximum shear stress of the Al 6082 (E) alloy is 154 MPa showing moderate elongation to failure. In contrast, the maximum shear stress of pure aluminium is only 62 MPa, but it presents much higher ductility. The interfaces of the AEH15 laminate containing native alumina layer are more ductile than the monolithic pure aluminium alloy, due to the different heating steps during processing. In this laminate, failure occurred outside the bond region, in the Al 1050 (H), indicating that the bond strength exceeds the fracture strength of the weaker component, an indication of excellent bonding. In contrast, the interfaces corresponding to the laminates with anodizing alumina layers show less ductility, although they have similar maximum stress than those of the AEH15 laminate and the as-received Al 1050 alloy. These interfaces show initially a ductile failure with cohesive failure of the weaker component (Al 1050 (H)), and finally catastrophic failure when the ductility of the interface is reached, being more premature for AEH15-5min. Thus, the interfaces in AEH15-5min under shearing stress are less ductile than those in AEH15-30s. It can be concluded that the interface toughness decreases when the alumina-coating thickness is increased. This is attributed to the higher volume fraction of brittle component in the tested bond, which is more susceptible to cracking and debonding during the shear test. Fig. 10 shows the SEM micrographs of shear fractured surfaces of AEH15-30s (Fig. 10a and b) and AEH15-5min (Fig. 10c and d) laminates. These fractured surfaces correspond to the zones where the maximum deformation of the interface was reached and thus interfacial failure occurred. In both laminates, the surface topography of the failed surface reveals perfectly the joint points between layers due to the extruded aluminium across the fractured alumina fragments. The SEM micrograph corresponding to the Al 6082 alloy in the AEH15-30s laminate (Fig. 10a) shows a homogeneous structure with multiple voids which indicates the alumina fragment position in the interface. The bright border of these voids corresponds to the extruded aluminium between the alumina fragments (bonded areas), which has been fractured during the shear test. Furthermore, the SEM micrograph of the Al 1050 alloy (Fig. 10b) shows a very ductile fractured surface indicating high bonding degree, according to the multiple joint points. On the contrary, the fractured surfaces corresponding to AEH15-5min laminate (Fig. 10c and d) present a more brittle behaviour. The Al 6082 surface, Fig. 10c, shows clearly large voids left by the alumina platelets and also bonded areas obtained during processing. The large bonded areas reflect the longer spaces between the cracked alumina for this thick alumina-coating laminate. Additionally, the

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alumina platelets can be observed on the Al 1050 fractured surface (Fig. 10d). It is worth noting that although the bonded area is the same for the three laminates considered, its distribution across the interface as a function of the initial alumina thickness determines the interface toughness. In summary, the alumina-coating thickness in the interface of the aluminium laminates affects the interface toughness and its fracture mechanisms. This influence was clearly demonstrated for the thickest coating. At low strain rate intrinsic fracture mechanisms are prevailing, while at high strain rate extrinsic mechanisms like delamination are also activated increasing the impact toughness. 4. Conclusions Three multilayer aluminium composites based in alternate Al 6082 and Al 1050 sheets have been processed by hot roll-bonding (ε = 1.2 comprising a total thickness reduction of 2.8:1). Three different alumina thicknesses in the interfaces have been considered and their influence on the fracture mechanisms of the multilayer laminates has been studied. The alumina coatings were grown by hard anodizing in sulphuric acid for two different times (30 s and 5 min). The major conclusions of the study are: 1. During rolling, aluminium oxide layers (native or anodized films) are fractured into fragments whose aspect ratio is independent of the initial alumina thickness. The aluminium matrix extrudes through the spaces opened up in the interface which are a function of the rolling reduction. Therefore, the bonded areas were similar for the three laminates although with different distribution along the interface. 2. All three laminates were found to exhibit higher impact toughness (Charpy tests) than the as-received Al 6082 alloy. In particular, the laminate containing the thickest alumina coating in the interface (AEH15-5min) exhibited toughness twice higher than that for the as-received Al 6082 alloy and delamination on the interfaces. In contrast, the laminate materials have shown different mechanical behaviour under slow strain rate loading (three point bending tests), and no clear influence of the alumina thickness has been found. Therefore, two fracture mechanisms depending on the strain rate conditions have been proposed. 3. The failure mechanism at high strain rate (Charpy test) comprises the nucleation and coalescence of cracks originated at debonded alumina fragments favouring extrinsic fracture mechanisms, such as delamination and crack renucleation. At low strain rate, intrinsic fracture mechanisms are favoured due to the higher ductility of the constituent materials. Delamination is avoided, if the distance between the alumina particles is larger than the crack blunting influence zone. Delamination occurs only for a narrow range of interparticle distances. Acknowledgements Financial support from CICYT (Project MAT2003-01172) is grate´ fully acknowledged. C.M. Cepeda-Jimenez and J.M. Garc´ıa-Infanta thank the Spanish Ministry of Education and Science for a Juan de la Cierva contract and a FPI fellowship respectively. We also thank ´ for the welding work, F.F. Gonzalez-Rodr´ ´ L. del Real-Alarcon ıguez for assistance during hot rolling and J. Chao-Hermida for assistance with the Charpy impact test. References [1] J. Wadsworth, D.R. Lesuer, Mater. Charact. 45 (2000) 289–313. [2] M.R. Abdullah, W.J. Cantwell, Compos. Sci. Technol. 66 (2006) 1682–1693.

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