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Coupled effects of substrate microstructure and sulphuric acid anodizing on fatigue life of a 2017A aluminum alloy
Materials and Design 86 (2015) 723–734
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Coupled effects of substrate microstructure and sulphuric acid anodizing on fatigue life of a 2017A aluminum alloy C. Fares a,b,⁎, L. Hemmouche a, M.A. Belouchrani a, A. Amrouche c, D. Chicot d, E.S. Puchi-Cabrera e,1 a
Laboratoire Génie des Matériaux, Ecole Militaire Polytechnique, BP17C Bordj El Bahri, Alger, Algeria Laboratoire Physique Théorique et Physique des Matériaux LPTPM, Université Hassiba Ben Bouali, BP151 Hai Essalam, 02000 Chlef, Algeria Laboratoire de Génie Civil et géo Environnement LGCgE, EA 4515, Faculté des Sciences Appliquées FSA Béthune, Université d'Artois, France d Université Lille Nord de France, USTL, LML, CNRS, UMR 8107, F-59650 Villeneuve d'Ascq, France e School of Metallurgical Engineering and Materials Science, Faculty of Engineering, Universidad Central de Venezuela, 47885, Los Chaguaramos, Caracas 1040, Venezuela b c
a r t i c l e
i n f o
Article history: Received 13 April 2015 Received in revised form 18 July 2015 Accepted 21 July 2015 Available online 30 July 2015 Keywords: Sulphuric acid anodizing Recovery annealing Age hardening Fatigue Microstructure 2017A aluminum alloys
a b s t r a c t The present study has been conducted in order to investigate the coupled effects of substrate microstructure and the sulphuric acid anodizing process, on the fatigue life of a 2017A aluminum alloy, by means of plane bending fatigue tests (R = −1). The effect of two different tempers, namely artificially aging (T6) and recovery annealing, prior to anodizing, was studied. The microstructural analysis of the substrate and anodic films was carried out by means of different techniques. It has been determined that the anodic film gives rise to a decrease in the fatigue properties of the heat treated alloy at all stress levels. However, at low stresses, the reduction in fatigue life is more pronounced for the annealed material. The effect of substrate microstructure on the morphology of the anodic layer and on the reduction of the fatigue life of the alloy is also discussed. © 2015 Elsevier Ltd. All rights reserved.
1. Introduction Fatigue behavior is an important selection criterion for the use of aluminum alloys in aerospace applications. Very often, for the aluminum alloys of the 2000 series, the main requirements for the components include a high fatigue strength, which is achieved by a proper material selection and heat treatment, in order to develop fine strengthening precipitates [1]. However, these precipitates usually cause localized corrosion such as pitting corrosion and intergranular corrosion, due to the potential difference between the precipitates and the matrix [2–7]. Al–Cu alloys are generally anodized by electrolysis to increase their corrosion and wear resistance and provide better adhesion for paint primers [8–10]. However, the improvement in the fatigue performance of these alloys could be severely impaired after the anodizing process. This discrepancy classically appears to be associated to the brittle and porous nature of the oxide layer [11–13], as well as the tensile residual stresses induced during the anodizing process [13–15]. Much of the
⁎ Corresponding author at: Laboratoire Génie des Matériaux, Ecole Militaire Polytechnique, BP17C Bordj El Bahri, Alger, Algeria. E-mail addresses: [email protected] (C. Fares), [email protected] (E.S. Puchi-Cabrera). 1 Currently at: Université Lille Nord de France, UVHC, LAMIH, UMR 8201, F-59313 Valenciennes, France.
http://dx.doi.org/10.1016/j.matdes.2015.07.120 0264-1275/© 2015 Elsevier Ltd. All rights reserved.
available literature provides evidence that the reduction in fatigue life is not only due to the well-known brittle properties of the aluminum oxide, but also to the film thickness [15–18], type of anodizing process [14,15,18,19], surface pretreatment prior anodizing [20–22] and sealing treatment after anodizing [23,24]. However, the reported data available in the literature regarding the combined effect of substrate microstructure and anodizing on the reduction in fatigue life is quite limited. In previous works conducted on the 2017A aluminum alloy [25–27], the morphological changes that occur in the anodic film as a consequence of the changes in substrate microstructure after heat treatments of artificial aging and annealing, have been investigated. One of the results of this study is the major role of intermetallic compounds (size and dispersion) on the oxide morphology and composition. As a result of the substrate microstructural changes, the alloying elements are incorporated into the oxide layer either as intermetallic compounds or oxides, as well as the Cu enrichment of the film–substrate interface after the aging treatment. Thus, the present work has been carried out in order to study the coupled effects of substrate microstructure and Sulphuric Acid Anodizing (SAA) on the fatigue behavior of 2017A aluminum alloy. The motivation for this work arises from the need to understand the different factors, which govern the film fracture mechanism. Fatigue cracks can be initiated from a range of possible sites, depending on the morphology and composition of the oxide layer, which in turn depends significantly on the substrate microstructure. Therefore, the combined effect of alloy
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microstructure and anodizing on fatigue performance has been investigated by SAA the substrate previously subjected to two different heat treatments, including a recovery-annealing and artificial age-hardening. Thus, the fatigue performance, of the heat treated and anodized specimens, has been investigated by means of plane bending fatigue tests conducted at 25 Hz. Also, a fractographic analysis of some selected samples after fatigue testing has been carried out, in order to determine the role of the anodic layer on the nucleation of fatigue cracks. 2. Experimental procedure 2.1. Materials preparation Fig. 1. Fatigue specimens' geometry and orientation regarding the initial rolled plate.
The material investigated in the present work is a 2017A-T4 aluminum alloy, whose chemical composition is given in Table 1. The material was provided as a rolled plate of approximately 4 mm in thickness. All the specimens were machined with their axis parallel to the longitudinal (L) direction of the plate. Fig. 1 illustrates the sketch of the fatigue specimens according to the DIN 50142 standard. All the dimensions shown are given in mm. The specimens were divided in two different groups. The samples corresponding to the first group were annealed at 300 °C for 30 min and cooled to room temperature. The specimens corresponding to the second group were solution-heat treated in a furnace with Ar atmosphere for 30 min at 500 ± 5 °C, which was followed by water quenching (b 40 °C). Subsequently, the samples were artificially aged for 10 h at 160 °C. The average grain size, determined by means of the mean linear intercept method, were of approximately 45 and 30 μm, for annealed and aged states, respectively. 2.2. Anodizing and sealing In order to investigate the effect of anodizing on the fatigue behavior of the alloy, a number of specimens were first surface treated by means of SAA and then sealed in boiling water. Prior to the anodizing process, the samples were subjected to different etching treatments in order to produce a chemically clean surface. The specimens were degreased and pickled in an alkaline solution (15 g/l of NaOH) at 65 °C for 3 min. Subsequently, these were rinsed thoroughly in deionized water and immersed in a sulphuric acid/chromic acid mixture (H2SO4: 180 ml/l, CrO3: 65 g/l) for 8 min at 60 °C. Finally, the samples were rinsed and dried in a warm air stream. SAA was accomplished in a thermostatically controlled electrochemical cell (±2 °C). All the specimens were individually anodized at 20 °C in a sulphuric acid solution (H2SO4: 200 g/l) at a constant voltage of 12 V for 30 min. Finally, after SAA, the specimens were sealed in boiling water at 97 °C for 30 min. The thickness of the anodic film was measured by means of both optical and SEM techniques. The average thicknesses of the anodic films were found to be approximately 14 and 18 μm, for the aging and annealing heat treatments, respectively. The mechanical properties of the alloy prior and after SAA, as determined by tensile tests, are given in Table 2. 2.3. Fatigue tests In order to determine the S–N curves, plane bending fatigue tests (R = −1) were conducted at a frequency of 25 Hz and at room temperature. The fatigue tests were carried out employing a 100 kN servo-
Table 1 Chemical composition of the 2017A aluminum alloy. Element
Cu
Fe
Si
Mn
Mg
Zn
Cr
Ti
Al
Composition (wt.%)
3.890 0.478 0.840 0.783 0.806 0.061 0.04 0.077 –
hydraulic SHENK-PWS testing machine under load control. The tests were set to run up to either the complete specimen fracture or to a maximum of 107 loading cycles, in the cases that fracture did not occur.
2.4. Microstructural and EPMA analysis After the SAA treatment, the morphology of both the substrate and the anodic film were investigated using an Electron Probe MicroAnalysis (EPMA) CAMECA SX100. All the samples were cross-sectioned, embedded in an epoxy resin, polished and carbon coated with a Bal-Tec SCD005 sputter coater. Both secondary and back-scattered electron images were obtained at 20 kV and 10 nA. Intensity profiles were obtained at 15 kV and 15 nA, whereas the WDS-mappings were carried out with an acceleration voltage of 15 kV, a probe current of 40 nA, a step size of 1 μm and a dwell time of 20 msec per measuring point. For both profiles and mappings, Wavelength Dispersive X-ray Spectrometry was employed. A thallium acid phthalate crystal (TAP) was employed to measure the Al and Mg Kα X-rays, a lithium fluoride crystal (LiF) for Cu and Mn Kα X-rays and a Ni–C multilayer crystal (referred to as PC2) to detect the O Kα X-rays. Scanning electron microscopy (SEM) was used to analyze the film morphology and fracture surface of the fatigue specimens. The analysis was conducted on a Hitachi (S-3000N) microscope. X-ray diffraction patterns were recorded on a Bruker-AXS type D8 Advance spectrometer with CoKα radiation of 0.178897 nm. The diffractometer scans were performed in a θ–2θ geometry, with a range of 2θ of 20°–55°, a step size of 0.02° and a counting time of 52 s/step, which corresponds to penetration depths of about 8 μm in the anodic layer. Crystalline phases were identified by comparison with the JCPDS database included in the software. The Raman spectra were recorded using a Horiba Jobin-Yvon LabRam Aramis coupled with an Olympus microscope with a 100× objective. The 632.8 nm red line from a Helium–Neon (HeNe) laser with a power of 50 mW was used as the excitation source. All Raman spectra were registered at room temperature between 100 and 1400 cm−1 in five (05) stages with two (02) accumulations each 20 s.
Table 2 Mechanical properties of the 2017A prior and after anodizing. Before anodizing
After anodizing
Heat treatment
Aging
Annealing
Aging
Annealing
Yield strength (Re, MPa) Ultimate tensile strength (Rm, MPa) Elongation (%) Elastic modulus (E, GPa)
400 ± 8.1 466 ± 9
338 ± 7 414 ± 8.2
378 ± 7.8 444 ± 8.6
316 ± 6.9 392 ± 7.9
6 74
9 74
4 73.6
5 73.6
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3. Results and discussion 3.1. Substrate and anodic film microstructures Figs. 2 and 3 show the distribution of intermetallic particles in the 2017A aluminum alloy both after annealing and age hardening. For both heat treatments, the main constituent particles (Fig. 2a and b) are Al–Cu and Al–Cu–Mg. The coarse particles Al–Cu–Mn–Fe– Si(Mg) appear bright in the backscatter image, whereas the surrounding aluminum matrix appears in gray (Fig. 3). The objective of the annealing treatment, which was the elimination of the internal stresses while keeping the initial microstructure, was not reached. It can be observed that, for this state, both Al–Cu and Al–Cu– Mg particles are widely spaced, coarsened and dispersed (Fig. 2a). On the other hand, the effect of age hardening (involving solution treatment at 500 °C, quenching and artificial aging for 10 h at 160 °C) on the microstructure of the 2017A aluminum alloy is shown in Fig. 2b. Here, it is clearly seen that the alloying elements trapped in solution precipitate to form a uniform distribution of very fine particles. The second phase θ-Al2Cu particles, though small, are brought closer. Spherical particles, identified as S-Al2CuMg, have a larger size than the Al2Cu particles and are relatively scarce in the alloy (Fig. 2b). The various analytical techniques employed in this study, allowed the characterization of the oxide layer obtained by SAA. The incorporation of
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the alloying elements into the oxide layer, such as Cu and Mg were established by EPMA [25,26] in agreement with the findings reported earlier concerning a series of metastable sputtering-deposited Al–Cu films [9,10]. Copper has a strong influence on the corrosion of aluminum alloys, where it is present in at matrix regions or as a constituent of secondary phases. When aluminum–copper alloys are immersed in a corrosive environment, a complex electrochemical behavior can be generated due to differences in reactivities of the various phases, leading to galvanic coupling between them [29–31]. Al2Cu particles reveal a cathodic behavior [30], so they are oxidized preferentially when the alloy is polarized above 5–6 V (SCE). Conversely, the S-phase (Al–Cu–Mg precipitate) [29–31], act as anode. They are selectively dissolved by relatively small anodic polarization in sulphuric and tartaric acid electrolytes [28]. Fig. 4 shows the qualitative profiles, along the line (1–2) (see Fig. 3) connecting the inside of the sample (1) to the edge (2) (oxide layers for both heat treatments). For the annealed state, Cu and Mg were incorporated into the oxide layer with the absence of a pre-existing enrichment of the alloying elements at the interface. In the case of the aging heat treatment, the incorporation of Cu particles into the anodic film occurred with the enrichment of the alloy at the film–substrate interface. In contrast, there was not any requirement of prior enrichment of Mg at the interface. Cu enrichment at the interface, which is characteristic of selective dissolution of intermetallic compounds, coincides with the decrease of the oxygen curve (Fig. 4b). This result, confirmed by Raman
Fig. 2. X-ray maps by EPMA of Al, Cu, Mn and Mg in the substrate after: a) annealing treatment, b) age hardening treatment. For both heat treatments, the main constituent particles are observed to be Al–Cu and Al–Cu–Mg. After annealing (panel a) the particles are widely spaced, coarsened and dispersed. On the contrary, after the age hardening treatment (panel b) the alloying elements trapped in solution precipitate to form a uniform distribution of very fine particles.
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Fig. 2 (continued).
spectroscopy (Fig. 5), is explained by the copper oxide formation in detriment of the alumina formation. Fig. 5 shows the obtained Raman spectra in the analyzed zone (Fig. 5a) after several trials. Two spectra are detected: alumina and cuprite. Fig. 5b shows a strong analogy with that of cuprite (Cu2O). The
most characteristic vibrations of cuprite (222, 299 and between 500 and 600 cm−1) are also observable. The common largest and no intense bands located at the region between 200 and 800 cm−1 reveal the presence of Al2O3 (Fig. 5c). The band at 980 cm−1 is one of the intense vibrations to be distinguished in Fig. 5c. It is due to the Al_O stretching
Fig. 3. SEM micrograph showing AlCuMnFeSi(Mg) particles in the substrate and cavities after Al2Cu particles dissolution in the film of the anodized 2017A aluminum alloy after: a) annealing treatment, b) age hardening treatment. The coarse particles Al–Cu–Mn–Fe–Si(Mg) appear bright in the backscatter image, whereas the surrounding aluminum matrix appears in gray.
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Fig. 4. X-rays profiles of O, Al and Cu for the anodized samples obtained by EPMA after different heat treatments. a) After annealing, Cu and Mg have been incorporated into the oxide layer with the absence of a pre-existing enrichment of the alloying elements at the interface. Cu enrichment at the interface occurs concurrently with the decrease of the oxygen curve. b) After age hardening, the incorporation of Cu particles into the anodic film occurred with the enrichment of the alloy at the film–substrate interface. However, no prior enrichment of Mg at the interface has occurred.
vibration, which is generally shifted from 1067 cm−1 to 790 cm−1 for the hydrated alumina. Similar results have been reported by Thomas et al. [31]. The Cartography of the obtained Raman spectra (Fig. 6a and b) showed the copper distribution and the presence of alumina in the analyzed zone (Fig. 5a). For both heat treatments, the copper oxide accumulates at the film–substrate interface. This accumulation of copper at the interface and its oxidation is in agreement with the literature [9,10,28]. It has been found that the oxidation of copper, which is dissolved by galvanic coupling between the matrix and second phases, will occur when a critical concentration is reached, in the enriched
copper layer. The copper enrichment occurs when the Gibbs free energy per equivalent for the formation of the copper oxide is higher than that for formation of alumina [9,10]. As a consequence of the formation of the copper oxide at the interface, the film thickness (after 30 min of SAA) in the aging state (14 μm) is slightly smaller than that corresponding to the annealing state (18 μm). The copper oxide accumulated at the film/matrix interface, during SAA, is present as Tenorite (CuO) in the DRX spectra (Fig. 7) and as cuprite (Cu2O) in the Raman spectra. This difference may be due to the CuO oxide reduction under the laser radiation applied in Raman spectroscopy [32].
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Fig. 5. Raman spectra of the anodized aged sample (a) Zone of analysis. (b) Copper oxide spectrum obtained at the interface. The most characteristic vibrations of cuprite (222, 299 and between 500 and 600 cm−1) can be observed. (c) Alumina spectrum obtained at the oxide layer, indicating the presence of the typical bands located at the region between 200 and 800 cm−1. The band at 980 cm−1, due to the Al_O stretching vibration, can also be clearly distinguished.
In Fig. 7, we can also observe strong peaks corresponding to the intermetallic compounds of Al2Cu, identified from the JCPDS #2-1309 data. The presence of Al2Cu in the oxide layer can be explained by its cathodic behavior during SAA. The peak intensity of these particles, corresponding to annealing state, is less than that corresponding to the aging state (Fig. 7). The oxide boundary of alumina moves faster towards the particles, as current distribution changes in favor of the less resistive particles [33]. The low resistance of the coarse Al2Cu particles favors
their oxidation and dissolution during SAA, hence, increasing the porosity of the oxide (Fig. 3a). 3.2. S–N curves For both metallurgical states, the static mechanical properties, such as elastic modulus, yield stress and fracture strength were not found to vary as a consequence of the SAA treatment, in agreement with the
Fig. 6. Raman spectra for anodized samples showing accumulation of copper oxide at the interface after: (a) age hardening and (b) annealing. Copper oxide is shown in blue whereas alumina is shown in red. For both heat treatments, the copper oxide accumulates at the film–substrate interface.
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Fig. 7. XRD patterns of the anodic films grown after the annealing and age hardening heat treatments showing the presence of Al2Cu particles and CuO in the oxide layer. The copper oxide accumulated at the film/matrix interface is present as Tenorite (CuO). Strong peaks corresponding to the intermetallic compounds of Al2Cu can be observed. The peak intensity of these particles, corresponding to annealing state, is less than that corresponding to the aging state.
Fig. 8. S–N curves for the heat treated (age hardening and annealing) 2017A aluminum alloy samples prior and after anodizing. The fatigue life of the age hardened alloy is observed to be greater than that of the annealed material. SAA leads to a significant decrease in the fatigue life of the alloy, in comparison with the heat treated samples prior to SAA, which is less pronounced as the maximum alternating stress increases.
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Fig. 9. SEM image showing the fracture surface of an age hardened specimen. Individual micrographs corresponding to different fracture modes observed from the specimen surface to the center: (a) specimen fracture surface, (b) distinct fracture modes observed, (c) ductile fracture zone, and (d) brittle fracture zone with cracks nucleation sites. At the specimen surface the fracture is brittle. Below the surface there is a zone of quasi-cleavage fracture. Further below, the fracture appears to be entirely ductile, as characterized by the presence of dimples.
findings earlier reported by Gröber et al. [34]. However, as can be seen from Table 2, ductility is significantly affected by SAA. The elongation exhibits a reduction for the anodized samples, similar to that reported by Shih et al. [24], indicating that anodizing impairs the capability of the material to deform plastically. Since the flow stress is determined by the interaction of dislocations with the coherent precipitates, the grain size is not expected to have a significant effect on the yield strength. Also, the elastic modulus is observed to be approximately the same for both alloys. The results of the plane bending fatigue tests of the heat treated samples prior and after to SAA are shown in Fig. 8. The curves corresponding to the aged hardened (with fine particles) and annealed (with coarse particles) specimens of the 2017A aluminum alloy show that the effect of microstructure at both high and low maximum alternating stresses is the same. The fatigue life of the age hardened alloy is consistently greater than that of the annealed material. As indicated above, in the aged alloy the yield strength is controlled by the interaction of dislocations with precipitates and the mean precipitate spacing is much smaller than the grain size [1]. Therefore, it is not expected that grain size plays a significant role in determining the yield stress of the material. For both heat treatments (Fig. 8), SAA leads to an important decrease in the fatigue life of the alloy, as compared to the heat treated
samples prior to SAA. However, this decrease is less pronounced as the maximum alternating stress increases. One possible explanation for this phenomenon could be the difference in the relative importance of crack nucleation versus crack propagation in the high and low maximum alternating stress regimes. Indeed, fatigue life is essentially divided into two regions: crack nucleation (dominant at low stresses) and crack propagation (dominant at high stresses). The anodic film would be expected to influence more significantly the fatigue crack nucleation stage and hence fatigue lives, at low stresses. The obtained results are in agreement with those of Nie et al. [19]. For a fatigue life of 106 cycles, there is a decrease of approximately 49% in fatigue strength for the age hardening condition and 64% for annealing condition. As discussed above, in general the decrease in fatigue life is related to the film morphology, which in turn is influenced by substrate microstructure. The observed reduction results mainly from the intermetallic compounds, which are formed within the film and to the formation of the copper oxide at the interface, in the case of the age hardening condition. During SAA the Al-rich matrix dissolves preferentially around the Al2Cu particles, which gives rise to the detachment of these particles and therefore, to their incorporation to the film and their dissolution during anodizing producing voids. At the film–substrate interface, the accumulation of the
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Fig. 10. SEM image showing fracture surface of an annealed specimen. Individual micrographs corresponding to different fracture modes observed from the specimen surface to the center: (a) Specimen fracture surface. (b) Brittle fracture zone illustrating a fractured constituent particle from which several fatigue cracks started. (c) Ductile fracture zone. (d) Fatigue crack growth. In general, fatigue cracks were observed to nucleate at constituent particles of Al–Cu–Mn–Fe–Si(Mg).
brittle copper oxide favors the delamination. For the annealed condition, the fatigue life reduction is more pronounced and it has been associated to the presence of cavities, which resulted as a consequence of the incorporation of the coarse Al 2 Cu particles into the film and their dissolution during anodizing.
3.3. Fatigue fractographic analysis Several specimens tested at a stress of 100 MPa were examined by SEM after failure, in order to understand the fatigue fracture mechanisms in the heat treated samples, prior and after anodizing. For both heat treatments, the fracture surface appearance varies from the surface to the center of the sample. At the specimen surface the fracture is brittle, as shown in Figs. 9 and 10. Immediately below, there is a zone of quasi-cleavage fracture [35] (Fig. 9) and further below the fracture appears to be entirely ductile (characterized by the presence of dimples), as shown in Figs. 9 and 10. Fig. 10 shows a fractured constituent particle from which several fatigue cracks started (Fig. 10b). The fatigue cracks nucleated at constituent particles of Al–Cu–Mn–Fe–Si(Mg), as determined by the EDS techniques.
The specimens subjected to anodic oxidation exhibit different fracture features. Figs. 11 and 12 illustrate typical fracture surfaces corresponding to anodized samples, both after the age hardening and annealing treatments. As can be observed in Fig. 11, the fracture surfaces of the specimens after age hardening and anodizing are similar to those of the heat treated samples prior to SAA. In this case, the quasi-cleavage fracture zone (Fig. 11c) is greater than those observed in the samples prior to SAA (Fig. 9b). The samples tested at 100 MPa exhibited fracture characteristics similar to those of the ductile materials. For the anodized specimens, multiple crack nucleation sites were found on the fractured surface, as can be seen in Figs. 11b and 12a. Shiozawa et al. [36] have also reported that the number of fatigue crack initiation sites increased for the anodized specimens as compared to untreated ones. The fractographic analysis conducted in the present work clearly illustrates the occurrence of these multi-site fatigue crack initiation phenomenon from the film and the subsequent propagation of such cracks towards the substrate (Fig. 11b). For the annealed and anodized samples tested at lower stresses, a smaller dimple fracture zone and a larger brittle fracture area was noticed (Fig. 12a). The SEM examination showed that in some regions, partial delamination of the anodic film occurred along
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Fig. 11. SEM images showing the fracture surface of an age hardened and anodized specimen. The individual micrographs correspond to different fracture modes observed from the specimen surface through the center. (a) Specimen fracture zone. (b) Cracks nucleation sites and fatigue crack growth. Multiple crack nucleation sites are present on the fractured surface. (c) Distinct fractures modes observed. The quasi-cleavage fracture zone is greater than that present in the samples prior to SAA (Fig. 9b). (d) Ductile fracture zone.
many areas, leading to the subsequent fracture of the material at the interface, as can be seen in Figs. 12a, b and c, in agreement with the findings earlier reported by Nie et al. [19]. The delamination of the film observed in Fig. 12c is related to the presence of intermetallic compounds and to the accumulation of the copper oxide particles at the film–substrate interface, as shown in Fig. 6. Fig. 12a and c also show the presence of many cavities with a large surface area just beneath the anodized film. These cavities appear to have coalesced, giving rise to the formation of a single large avoid. In general, the specimen surface showed an intergranular fracture mode. As a consequence of the propagation of multiple cracks, the ductile fracture zone (DFZ) was created (Fig. 12a). The DFZ is mostly located at the center of the fracture surface. 4. Conclusions The present study focused on the coupled effect of substrate microstructure and anodizing on the mechanical properties of a 2017A-T4 aluminum alloy especially the fatigue performance. SAA does not have
any significant effect on the static mechanical properties of the alloy, with the exception of ductility, which is significantly impaired. At high alternating stresses, the presence of the anodic film for both microstructural conditions does not seem to have a pronounced effect on fatigue crack growth. Therefore, under these alternating conditions, anodizing gives rise to a slight decrease in fatigue life. However, the results obtained from fatigue testing at low alternating stresses showed that anodizing reduces significantly the fatigue life of the alloy for both microstructural conditions. The decrease in fatigue life has been attributed to the preferential dissolution of the matrix around the cathodic particles, such as Al2Cu which gave rise to the formation of cavities. Whatever the reason of cavities initiation, these acted as stress concentrators and promoted the nucleation of many small fatigue cracks. The fatigue life reduction can also be associated to the brittle nature of the copper oxide particles accumulated at the interface. The decrease in fatigue life observed in the specimens subjected to annealing is more pronounced than that determined for the samples subjected to age hardening. For the latter condition, the intermetallic particles are coarser and therefore, they leave larger cavities after their dissolution during
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Fig. 12. SEM images showing the fracture surface of an annealed and anodized specimen. Individual micrographs corresponding to different fracture modes observed from the specimen surface through the center: (a) Specimen surface fracture. A smaller dimple fracture zone and a larger brittle fracture area can be noticed. (b) Micrograph showing film detachment. (c) Micrograph showing large cavities under the detached film. The delamination of the film is related to the presence of intermetallic compounds and to the accumulation of the copper oxide particles at the film–substrate interface. (d) Propagation of multiple cracks in the large cavity. Many cavities with a large surface area were observed to be located just beneath the anodized film.
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[8] M. Saenz de Miera, M. Curioni, P. Skeldon, G.E. Thompson, Modelling the anodizing behaviour of aluminium alloys in sulphuric acid through alloy analogues, Corros. Sci. 50 (2008) 3410–3415. [9] T. Hashimoto, X. Zhou, P. Skeldon, G.E. Thompson, Structure of the copper-enriched layer introduced by anodic oxidation of copper-containing aluminium alloy, Electrochimica Acta (2015) (in press). [10] S.J. Garcia-Vergara, K. El Khazmi, P. Skeldon, G.E. Thompson, Influence of copper on the morphology of porous anodic alumina, Corros. Sci. 48 (2006) 2937–2946. [11] A.M. Cree, G.W. Weidmann, Effect of anodised coatings on fatigue crack growth rates in aluminum alloy, Surf. Eng. 13 (1997) 51–55. [12] A.M. Cree, G.W. Weidmann, The fracture and fatigue properties of anodised aluminium alloy, Trans. Inst. Met. Finish. 75 (1997) 199–202. [13] A.M. Cree, S.V. Hainsworth, G.W. Weidmann, The fracture and fatigue properties of anodised aluminium alloy, Trans. Inst. Met. Finish. 84 (2006) 246–251. [14] J.A.M. DeCamargo, H.J. Cornelis, V.M.O.H. Cioffi, M.Y.P. Costa, Coating residual stress effects on fatigue performance of 7050-T7451 aluminum alloy, Surf. Coat. Technol. 201 (2007) 9448–9455. [15] B. Lonyuk, I. Apachitei, J. Duszczyk, The effect of oxide coatings on fatigue properties of 7475-T6 aluminium alloy, Surf. Coat. Technol. 201 (2007) 8688–8694. [16] E. Cirik, K. Genel, Effect of anodic oxidation on fatigue performance of 7075-T6 alloy, Surf. Coat. Technol. 202 (2008) 5190–5201. [17] R. Sadeler, Effect of a commercial hard anodizing on the fatigue property of a 2014T6 aluminium alloy, J. Mater. Sci. 41 (2006) 5803–5809. [18] W. Bensalah, K. Elleuch, M. Feki, M. De Petris-Wery, H.F. Ayedi, Mechanical failure of anodized film of aluminium in bending, Mater. Des. 30 (2009) 3141–3149.
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[19] B. Nie, Z. Zhang, Z. Zhao, Q. Zhong, Effect of anodizing treatment on the very high cycle fatigue behavior of 2A12-T4 aluminum alloy, Mater. Des. 50 (2013) 1005–1010. [20] M. Chaussumier, C. Mabru, M. Shahzad, R. Chieragatti, F. Rezai-Aria, A predictive fatigue life model for anodized 7050 aluminium alloy, Inter. J. of Fatigue 48 (2013) 205–213. [21] M. Chaussumier, C. Mabru, R. Chieragatti, M. Shahzad, Fatigue life model for 7050 chromic anodized aluminium alloy, Procedia Engineering 66 (2013) 300–312. [22] M. Shahzad, M. Chaussumier, R. Chieragatti, C. Mabru, F. Rezai-Aria, Surface characterization and influence of anodizing process on fatigue life of Al 7050 alloy, Mater. Des. 32 (2011) 3328–3335. [23] X. Zhao, G. Wei, Y. Yu, Y. Guo, A. Zhang, An analysis of mechanical properties of anodized aluminum film at high stress, Surf. Rev. Letters 22 (1550002) (2015) 1–7. [24] T.-S. Shih, T.-H. Lee, Y.-J. Jhou, The effects of anodization treatment on the microstructure and fatigue behavior of 7075-T73 aluminum, Alloy Mater Trans. 55 (2014) 1280–1285. [25] C. Fares, M.A. Belouchrani, S. Bellayer, T. Boukharouba, A. Britah, Influence of intermetallic compounds and metallurgical state of the 2017A aluminium alloy on the morphology of alumina films developed by anodic oxidation, J. Tribology Surf. Eng. 2 (2011) 239–251. [26] C. Fares, A. Boudiaf, M.A. Belouchrani, T. Boukharouba, Microstructural characterization of oxide layer developed by sulphuric anodisation on 2017A alloys, Int. J. Mater. Res. 104 (2013) 1–6. [27] L. Hemmouche, C. Fares, M.A. Belouchrani, Influence of heat treatments and anodization on fatigue life of 2017A alloy, Eng. Fail. Anal. 35 (2013) 554–561.
[28] M. Curioni, Saenz De Miera M, Skeldon P, Thompson GE, Ferguson J. Macroscopic and local filming behavior of AA2024T3 aluminum alloy during anodizing in sulfuric acid electrolyte, J. Electrochem. Soc 155 (2008) C387–C395. [29] M. Bethencourt, F.J. Botana, M.J. Cano, M. Marcos, J.M. Sánchez-Amaya, L. GonzálezRovira, Behaviour of the alloy AA2017 in aqueous solutions of NaCl. Part I: Corrosion mechanisms, Corr. Sci. 51 (2009) 518–524. [30] C. Blanc, A. Freulon, M.C. Lafont, Y. Kihn, G. Mankowski, Modelling the corrosion behaviour of Al2CuMg coarse particles in copper-rich aluminium alloys, Corr. Sci. 48 (2006) 3838–3851. [31] P.V. Thomas, V. Ramakrishnan, V.K. Vaidyan, Oxidation studies of aluminium thin films by Raman spectroscopy, Thin Solid Films. 170 (1989) 35–40. [32] X. Zhou, G.E. Thomson, P. Skeldon, K. Shimizu, G.C. Wood, The valence state of copper in anodic films formed on Al-1 at.% Cu alloy, Corr. Sci. 47 (2005) 1299–1306. [33] L.E. Fratila-Apachitei, H. Terryn, P. Skeldon, G.E. Thompson, J. Duszczyk, L. Katgerman, Influence of substrate microstructure on the growth of anodic oxide layers, Electrochim. Acta 49 (2004) 1127–1140. [34] D. Gröber, W. Georgi, M. Sieber, I. Scharf, R.J. Hellmig, E. Leidich, T. Lampke, P. Mayr, The effect of anodising on the fatigue performance of self-tapping aluminium screws, Inter. J. Fatigue 75 (2015) 108–114. [35] AlTh Kermanidis, P.V. Petroyiannis, SpG Pantelakis, Fatigue and damage tolerance behaviour of corroded 2024 T351 aircraft aluminum alloy, Theoretical and Applied Fracture Mechanics 43 (2005) 121–132. [36] K. Shiozawa, H. Kbayashi, M. Terada, A. Matsui, Effect of anodized coatings on fatigue strength in aluminium alloy, 5th International Conference of Computational and Experimental Methods, 6 2001, pp. 397–406.
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Coupled effects of substrate microstructure and sulphuric acid anodizing on fatigue life of a 2017A aluminum alloy C. Fares a,b,⁎, L. Hemmouche a, M.A. Belouchrani a, A. Amrouche c, D. Chicot d, E.S. Puchi-Cabrera e,1 a
Laboratoire Génie des Matériaux, Ecole Militaire Polytechnique, BP17C Bordj El Bahri, Alger, Algeria Laboratoire Physique Théorique et Physique des Matériaux LPTPM, Université Hassiba Ben Bouali, BP151 Hai Essalam, 02000 Chlef, Algeria Laboratoire de Génie Civil et géo Environnement LGCgE, EA 4515, Faculté des Sciences Appliquées FSA Béthune, Université d'Artois, France d Université Lille Nord de France, USTL, LML, CNRS, UMR 8107, F-59650 Villeneuve d'Ascq, France e School of Metallurgical Engineering and Materials Science, Faculty of Engineering, Universidad Central de Venezuela, 47885, Los Chaguaramos, Caracas 1040, Venezuela b c
a r t i c l e
i n f o
Article history: Received 13 April 2015 Received in revised form 18 July 2015 Accepted 21 July 2015 Available online 30 July 2015 Keywords: Sulphuric acid anodizing Recovery annealing Age hardening Fatigue Microstructure 2017A aluminum alloys
a b s t r a c t The present study has been conducted in order to investigate the coupled effects of substrate microstructure and the sulphuric acid anodizing process, on the fatigue life of a 2017A aluminum alloy, by means of plane bending fatigue tests (R = −1). The effect of two different tempers, namely artificially aging (T6) and recovery annealing, prior to anodizing, was studied. The microstructural analysis of the substrate and anodic films was carried out by means of different techniques. It has been determined that the anodic film gives rise to a decrease in the fatigue properties of the heat treated alloy at all stress levels. However, at low stresses, the reduction in fatigue life is more pronounced for the annealed material. The effect of substrate microstructure on the morphology of the anodic layer and on the reduction of the fatigue life of the alloy is also discussed. © 2015 Elsevier Ltd. All rights reserved.
1. Introduction Fatigue behavior is an important selection criterion for the use of aluminum alloys in aerospace applications. Very often, for the aluminum alloys of the 2000 series, the main requirements for the components include a high fatigue strength, which is achieved by a proper material selection and heat treatment, in order to develop fine strengthening precipitates [1]. However, these precipitates usually cause localized corrosion such as pitting corrosion and intergranular corrosion, due to the potential difference between the precipitates and the matrix [2–7]. Al–Cu alloys are generally anodized by electrolysis to increase their corrosion and wear resistance and provide better adhesion for paint primers [8–10]. However, the improvement in the fatigue performance of these alloys could be severely impaired after the anodizing process. This discrepancy classically appears to be associated to the brittle and porous nature of the oxide layer [11–13], as well as the tensile residual stresses induced during the anodizing process [13–15]. Much of the
⁎ Corresponding author at: Laboratoire Génie des Matériaux, Ecole Militaire Polytechnique, BP17C Bordj El Bahri, Alger, Algeria. E-mail addresses: [email protected] (C. Fares), [email protected] (E.S. Puchi-Cabrera). 1 Currently at: Université Lille Nord de France, UVHC, LAMIH, UMR 8201, F-59313 Valenciennes, France.
http://dx.doi.org/10.1016/j.matdes.2015.07.120 0264-1275/© 2015 Elsevier Ltd. All rights reserved.
available literature provides evidence that the reduction in fatigue life is not only due to the well-known brittle properties of the aluminum oxide, but also to the film thickness [15–18], type of anodizing process [14,15,18,19], surface pretreatment prior anodizing [20–22] and sealing treatment after anodizing [23,24]. However, the reported data available in the literature regarding the combined effect of substrate microstructure and anodizing on the reduction in fatigue life is quite limited. In previous works conducted on the 2017A aluminum alloy [25–27], the morphological changes that occur in the anodic film as a consequence of the changes in substrate microstructure after heat treatments of artificial aging and annealing, have been investigated. One of the results of this study is the major role of intermetallic compounds (size and dispersion) on the oxide morphology and composition. As a result of the substrate microstructural changes, the alloying elements are incorporated into the oxide layer either as intermetallic compounds or oxides, as well as the Cu enrichment of the film–substrate interface after the aging treatment. Thus, the present work has been carried out in order to study the coupled effects of substrate microstructure and Sulphuric Acid Anodizing (SAA) on the fatigue behavior of 2017A aluminum alloy. The motivation for this work arises from the need to understand the different factors, which govern the film fracture mechanism. Fatigue cracks can be initiated from a range of possible sites, depending on the morphology and composition of the oxide layer, which in turn depends significantly on the substrate microstructure. Therefore, the combined effect of alloy
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microstructure and anodizing on fatigue performance has been investigated by SAA the substrate previously subjected to two different heat treatments, including a recovery-annealing and artificial age-hardening. Thus, the fatigue performance, of the heat treated and anodized specimens, has been investigated by means of plane bending fatigue tests conducted at 25 Hz. Also, a fractographic analysis of some selected samples after fatigue testing has been carried out, in order to determine the role of the anodic layer on the nucleation of fatigue cracks. 2. Experimental procedure 2.1. Materials preparation Fig. 1. Fatigue specimens' geometry and orientation regarding the initial rolled plate.
The material investigated in the present work is a 2017A-T4 aluminum alloy, whose chemical composition is given in Table 1. The material was provided as a rolled plate of approximately 4 mm in thickness. All the specimens were machined with their axis parallel to the longitudinal (L) direction of the plate. Fig. 1 illustrates the sketch of the fatigue specimens according to the DIN 50142 standard. All the dimensions shown are given in mm. The specimens were divided in two different groups. The samples corresponding to the first group were annealed at 300 °C for 30 min and cooled to room temperature. The specimens corresponding to the second group were solution-heat treated in a furnace with Ar atmosphere for 30 min at 500 ± 5 °C, which was followed by water quenching (b 40 °C). Subsequently, the samples were artificially aged for 10 h at 160 °C. The average grain size, determined by means of the mean linear intercept method, were of approximately 45 and 30 μm, for annealed and aged states, respectively. 2.2. Anodizing and sealing In order to investigate the effect of anodizing on the fatigue behavior of the alloy, a number of specimens were first surface treated by means of SAA and then sealed in boiling water. Prior to the anodizing process, the samples were subjected to different etching treatments in order to produce a chemically clean surface. The specimens were degreased and pickled in an alkaline solution (15 g/l of NaOH) at 65 °C for 3 min. Subsequently, these were rinsed thoroughly in deionized water and immersed in a sulphuric acid/chromic acid mixture (H2SO4: 180 ml/l, CrO3: 65 g/l) for 8 min at 60 °C. Finally, the samples were rinsed and dried in a warm air stream. SAA was accomplished in a thermostatically controlled electrochemical cell (±2 °C). All the specimens were individually anodized at 20 °C in a sulphuric acid solution (H2SO4: 200 g/l) at a constant voltage of 12 V for 30 min. Finally, after SAA, the specimens were sealed in boiling water at 97 °C for 30 min. The thickness of the anodic film was measured by means of both optical and SEM techniques. The average thicknesses of the anodic films were found to be approximately 14 and 18 μm, for the aging and annealing heat treatments, respectively. The mechanical properties of the alloy prior and after SAA, as determined by tensile tests, are given in Table 2. 2.3. Fatigue tests In order to determine the S–N curves, plane bending fatigue tests (R = −1) were conducted at a frequency of 25 Hz and at room temperature. The fatigue tests were carried out employing a 100 kN servo-
Table 1 Chemical composition of the 2017A aluminum alloy. Element
Cu
Fe
Si
Mn
Mg
Zn
Cr
Ti
Al
Composition (wt.%)
3.890 0.478 0.840 0.783 0.806 0.061 0.04 0.077 –
hydraulic SHENK-PWS testing machine under load control. The tests were set to run up to either the complete specimen fracture or to a maximum of 107 loading cycles, in the cases that fracture did not occur.
2.4. Microstructural and EPMA analysis After the SAA treatment, the morphology of both the substrate and the anodic film were investigated using an Electron Probe MicroAnalysis (EPMA) CAMECA SX100. All the samples were cross-sectioned, embedded in an epoxy resin, polished and carbon coated with a Bal-Tec SCD005 sputter coater. Both secondary and back-scattered electron images were obtained at 20 kV and 10 nA. Intensity profiles were obtained at 15 kV and 15 nA, whereas the WDS-mappings were carried out with an acceleration voltage of 15 kV, a probe current of 40 nA, a step size of 1 μm and a dwell time of 20 msec per measuring point. For both profiles and mappings, Wavelength Dispersive X-ray Spectrometry was employed. A thallium acid phthalate crystal (TAP) was employed to measure the Al and Mg Kα X-rays, a lithium fluoride crystal (LiF) for Cu and Mn Kα X-rays and a Ni–C multilayer crystal (referred to as PC2) to detect the O Kα X-rays. Scanning electron microscopy (SEM) was used to analyze the film morphology and fracture surface of the fatigue specimens. The analysis was conducted on a Hitachi (S-3000N) microscope. X-ray diffraction patterns were recorded on a Bruker-AXS type D8 Advance spectrometer with CoKα radiation of 0.178897 nm. The diffractometer scans were performed in a θ–2θ geometry, with a range of 2θ of 20°–55°, a step size of 0.02° and a counting time of 52 s/step, which corresponds to penetration depths of about 8 μm in the anodic layer. Crystalline phases were identified by comparison with the JCPDS database included in the software. The Raman spectra were recorded using a Horiba Jobin-Yvon LabRam Aramis coupled with an Olympus microscope with a 100× objective. The 632.8 nm red line from a Helium–Neon (HeNe) laser with a power of 50 mW was used as the excitation source. All Raman spectra were registered at room temperature between 100 and 1400 cm−1 in five (05) stages with two (02) accumulations each 20 s.
Table 2 Mechanical properties of the 2017A prior and after anodizing. Before anodizing
After anodizing
Heat treatment
Aging
Annealing
Aging
Annealing
Yield strength (Re, MPa) Ultimate tensile strength (Rm, MPa) Elongation (%) Elastic modulus (E, GPa)
400 ± 8.1 466 ± 9
338 ± 7 414 ± 8.2
378 ± 7.8 444 ± 8.6
316 ± 6.9 392 ± 7.9
6 74
9 74
4 73.6
5 73.6
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3. Results and discussion 3.1. Substrate and anodic film microstructures Figs. 2 and 3 show the distribution of intermetallic particles in the 2017A aluminum alloy both after annealing and age hardening. For both heat treatments, the main constituent particles (Fig. 2a and b) are Al–Cu and Al–Cu–Mg. The coarse particles Al–Cu–Mn–Fe– Si(Mg) appear bright in the backscatter image, whereas the surrounding aluminum matrix appears in gray (Fig. 3). The objective of the annealing treatment, which was the elimination of the internal stresses while keeping the initial microstructure, was not reached. It can be observed that, for this state, both Al–Cu and Al–Cu– Mg particles are widely spaced, coarsened and dispersed (Fig. 2a). On the other hand, the effect of age hardening (involving solution treatment at 500 °C, quenching and artificial aging for 10 h at 160 °C) on the microstructure of the 2017A aluminum alloy is shown in Fig. 2b. Here, it is clearly seen that the alloying elements trapped in solution precipitate to form a uniform distribution of very fine particles. The second phase θ-Al2Cu particles, though small, are brought closer. Spherical particles, identified as S-Al2CuMg, have a larger size than the Al2Cu particles and are relatively scarce in the alloy (Fig. 2b). The various analytical techniques employed in this study, allowed the characterization of the oxide layer obtained by SAA. The incorporation of
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the alloying elements into the oxide layer, such as Cu and Mg were established by EPMA [25,26] in agreement with the findings reported earlier concerning a series of metastable sputtering-deposited Al–Cu films [9,10]. Copper has a strong influence on the corrosion of aluminum alloys, where it is present in at matrix regions or as a constituent of secondary phases. When aluminum–copper alloys are immersed in a corrosive environment, a complex electrochemical behavior can be generated due to differences in reactivities of the various phases, leading to galvanic coupling between them [29–31]. Al2Cu particles reveal a cathodic behavior [30], so they are oxidized preferentially when the alloy is polarized above 5–6 V (SCE). Conversely, the S-phase (Al–Cu–Mg precipitate) [29–31], act as anode. They are selectively dissolved by relatively small anodic polarization in sulphuric and tartaric acid electrolytes [28]. Fig. 4 shows the qualitative profiles, along the line (1–2) (see Fig. 3) connecting the inside of the sample (1) to the edge (2) (oxide layers for both heat treatments). For the annealed state, Cu and Mg were incorporated into the oxide layer with the absence of a pre-existing enrichment of the alloying elements at the interface. In the case of the aging heat treatment, the incorporation of Cu particles into the anodic film occurred with the enrichment of the alloy at the film–substrate interface. In contrast, there was not any requirement of prior enrichment of Mg at the interface. Cu enrichment at the interface, which is characteristic of selective dissolution of intermetallic compounds, coincides with the decrease of the oxygen curve (Fig. 4b). This result, confirmed by Raman
Fig. 2. X-ray maps by EPMA of Al, Cu, Mn and Mg in the substrate after: a) annealing treatment, b) age hardening treatment. For both heat treatments, the main constituent particles are observed to be Al–Cu and Al–Cu–Mg. After annealing (panel a) the particles are widely spaced, coarsened and dispersed. On the contrary, after the age hardening treatment (panel b) the alloying elements trapped in solution precipitate to form a uniform distribution of very fine particles.
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Fig. 2 (continued).
spectroscopy (Fig. 5), is explained by the copper oxide formation in detriment of the alumina formation. Fig. 5 shows the obtained Raman spectra in the analyzed zone (Fig. 5a) after several trials. Two spectra are detected: alumina and cuprite. Fig. 5b shows a strong analogy with that of cuprite (Cu2O). The
most characteristic vibrations of cuprite (222, 299 and between 500 and 600 cm−1) are also observable. The common largest and no intense bands located at the region between 200 and 800 cm−1 reveal the presence of Al2O3 (Fig. 5c). The band at 980 cm−1 is one of the intense vibrations to be distinguished in Fig. 5c. It is due to the Al_O stretching
Fig. 3. SEM micrograph showing AlCuMnFeSi(Mg) particles in the substrate and cavities after Al2Cu particles dissolution in the film of the anodized 2017A aluminum alloy after: a) annealing treatment, b) age hardening treatment. The coarse particles Al–Cu–Mn–Fe–Si(Mg) appear bright in the backscatter image, whereas the surrounding aluminum matrix appears in gray.
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Fig. 4. X-rays profiles of O, Al and Cu for the anodized samples obtained by EPMA after different heat treatments. a) After annealing, Cu and Mg have been incorporated into the oxide layer with the absence of a pre-existing enrichment of the alloying elements at the interface. Cu enrichment at the interface occurs concurrently with the decrease of the oxygen curve. b) After age hardening, the incorporation of Cu particles into the anodic film occurred with the enrichment of the alloy at the film–substrate interface. However, no prior enrichment of Mg at the interface has occurred.
vibration, which is generally shifted from 1067 cm−1 to 790 cm−1 for the hydrated alumina. Similar results have been reported by Thomas et al. [31]. The Cartography of the obtained Raman spectra (Fig. 6a and b) showed the copper distribution and the presence of alumina in the analyzed zone (Fig. 5a). For both heat treatments, the copper oxide accumulates at the film–substrate interface. This accumulation of copper at the interface and its oxidation is in agreement with the literature [9,10,28]. It has been found that the oxidation of copper, which is dissolved by galvanic coupling between the matrix and second phases, will occur when a critical concentration is reached, in the enriched
copper layer. The copper enrichment occurs when the Gibbs free energy per equivalent for the formation of the copper oxide is higher than that for formation of alumina [9,10]. As a consequence of the formation of the copper oxide at the interface, the film thickness (after 30 min of SAA) in the aging state (14 μm) is slightly smaller than that corresponding to the annealing state (18 μm). The copper oxide accumulated at the film/matrix interface, during SAA, is present as Tenorite (CuO) in the DRX spectra (Fig. 7) and as cuprite (Cu2O) in the Raman spectra. This difference may be due to the CuO oxide reduction under the laser radiation applied in Raman spectroscopy [32].
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Fig. 5. Raman spectra of the anodized aged sample (a) Zone of analysis. (b) Copper oxide spectrum obtained at the interface. The most characteristic vibrations of cuprite (222, 299 and between 500 and 600 cm−1) can be observed. (c) Alumina spectrum obtained at the oxide layer, indicating the presence of the typical bands located at the region between 200 and 800 cm−1. The band at 980 cm−1, due to the Al_O stretching vibration, can also be clearly distinguished.
In Fig. 7, we can also observe strong peaks corresponding to the intermetallic compounds of Al2Cu, identified from the JCPDS #2-1309 data. The presence of Al2Cu in the oxide layer can be explained by its cathodic behavior during SAA. The peak intensity of these particles, corresponding to annealing state, is less than that corresponding to the aging state (Fig. 7). The oxide boundary of alumina moves faster towards the particles, as current distribution changes in favor of the less resistive particles [33]. The low resistance of the coarse Al2Cu particles favors
their oxidation and dissolution during SAA, hence, increasing the porosity of the oxide (Fig. 3a). 3.2. S–N curves For both metallurgical states, the static mechanical properties, such as elastic modulus, yield stress and fracture strength were not found to vary as a consequence of the SAA treatment, in agreement with the
Fig. 6. Raman spectra for anodized samples showing accumulation of copper oxide at the interface after: (a) age hardening and (b) annealing. Copper oxide is shown in blue whereas alumina is shown in red. For both heat treatments, the copper oxide accumulates at the film–substrate interface.
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Fig. 7. XRD patterns of the anodic films grown after the annealing and age hardening heat treatments showing the presence of Al2Cu particles and CuO in the oxide layer. The copper oxide accumulated at the film/matrix interface is present as Tenorite (CuO). Strong peaks corresponding to the intermetallic compounds of Al2Cu can be observed. The peak intensity of these particles, corresponding to annealing state, is less than that corresponding to the aging state.
Fig. 8. S–N curves for the heat treated (age hardening and annealing) 2017A aluminum alloy samples prior and after anodizing. The fatigue life of the age hardened alloy is observed to be greater than that of the annealed material. SAA leads to a significant decrease in the fatigue life of the alloy, in comparison with the heat treated samples prior to SAA, which is less pronounced as the maximum alternating stress increases.
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Fig. 9. SEM image showing the fracture surface of an age hardened specimen. Individual micrographs corresponding to different fracture modes observed from the specimen surface to the center: (a) specimen fracture surface, (b) distinct fracture modes observed, (c) ductile fracture zone, and (d) brittle fracture zone with cracks nucleation sites. At the specimen surface the fracture is brittle. Below the surface there is a zone of quasi-cleavage fracture. Further below, the fracture appears to be entirely ductile, as characterized by the presence of dimples.
findings earlier reported by Gröber et al. [34]. However, as can be seen from Table 2, ductility is significantly affected by SAA. The elongation exhibits a reduction for the anodized samples, similar to that reported by Shih et al. [24], indicating that anodizing impairs the capability of the material to deform plastically. Since the flow stress is determined by the interaction of dislocations with the coherent precipitates, the grain size is not expected to have a significant effect on the yield strength. Also, the elastic modulus is observed to be approximately the same for both alloys. The results of the plane bending fatigue tests of the heat treated samples prior and after to SAA are shown in Fig. 8. The curves corresponding to the aged hardened (with fine particles) and annealed (with coarse particles) specimens of the 2017A aluminum alloy show that the effect of microstructure at both high and low maximum alternating stresses is the same. The fatigue life of the age hardened alloy is consistently greater than that of the annealed material. As indicated above, in the aged alloy the yield strength is controlled by the interaction of dislocations with precipitates and the mean precipitate spacing is much smaller than the grain size [1]. Therefore, it is not expected that grain size plays a significant role in determining the yield stress of the material. For both heat treatments (Fig. 8), SAA leads to an important decrease in the fatigue life of the alloy, as compared to the heat treated
samples prior to SAA. However, this decrease is less pronounced as the maximum alternating stress increases. One possible explanation for this phenomenon could be the difference in the relative importance of crack nucleation versus crack propagation in the high and low maximum alternating stress regimes. Indeed, fatigue life is essentially divided into two regions: crack nucleation (dominant at low stresses) and crack propagation (dominant at high stresses). The anodic film would be expected to influence more significantly the fatigue crack nucleation stage and hence fatigue lives, at low stresses. The obtained results are in agreement with those of Nie et al. [19]. For a fatigue life of 106 cycles, there is a decrease of approximately 49% in fatigue strength for the age hardening condition and 64% for annealing condition. As discussed above, in general the decrease in fatigue life is related to the film morphology, which in turn is influenced by substrate microstructure. The observed reduction results mainly from the intermetallic compounds, which are formed within the film and to the formation of the copper oxide at the interface, in the case of the age hardening condition. During SAA the Al-rich matrix dissolves preferentially around the Al2Cu particles, which gives rise to the detachment of these particles and therefore, to their incorporation to the film and their dissolution during anodizing producing voids. At the film–substrate interface, the accumulation of the
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Fig. 10. SEM image showing fracture surface of an annealed specimen. Individual micrographs corresponding to different fracture modes observed from the specimen surface to the center: (a) Specimen fracture surface. (b) Brittle fracture zone illustrating a fractured constituent particle from which several fatigue cracks started. (c) Ductile fracture zone. (d) Fatigue crack growth. In general, fatigue cracks were observed to nucleate at constituent particles of Al–Cu–Mn–Fe–Si(Mg).
brittle copper oxide favors the delamination. For the annealed condition, the fatigue life reduction is more pronounced and it has been associated to the presence of cavities, which resulted as a consequence of the incorporation of the coarse Al 2 Cu particles into the film and their dissolution during anodizing.
3.3. Fatigue fractographic analysis Several specimens tested at a stress of 100 MPa were examined by SEM after failure, in order to understand the fatigue fracture mechanisms in the heat treated samples, prior and after anodizing. For both heat treatments, the fracture surface appearance varies from the surface to the center of the sample. At the specimen surface the fracture is brittle, as shown in Figs. 9 and 10. Immediately below, there is a zone of quasi-cleavage fracture [35] (Fig. 9) and further below the fracture appears to be entirely ductile (characterized by the presence of dimples), as shown in Figs. 9 and 10. Fig. 10 shows a fractured constituent particle from which several fatigue cracks started (Fig. 10b). The fatigue cracks nucleated at constituent particles of Al–Cu–Mn–Fe–Si(Mg), as determined by the EDS techniques.
The specimens subjected to anodic oxidation exhibit different fracture features. Figs. 11 and 12 illustrate typical fracture surfaces corresponding to anodized samples, both after the age hardening and annealing treatments. As can be observed in Fig. 11, the fracture surfaces of the specimens after age hardening and anodizing are similar to those of the heat treated samples prior to SAA. In this case, the quasi-cleavage fracture zone (Fig. 11c) is greater than those observed in the samples prior to SAA (Fig. 9b). The samples tested at 100 MPa exhibited fracture characteristics similar to those of the ductile materials. For the anodized specimens, multiple crack nucleation sites were found on the fractured surface, as can be seen in Figs. 11b and 12a. Shiozawa et al. [36] have also reported that the number of fatigue crack initiation sites increased for the anodized specimens as compared to untreated ones. The fractographic analysis conducted in the present work clearly illustrates the occurrence of these multi-site fatigue crack initiation phenomenon from the film and the subsequent propagation of such cracks towards the substrate (Fig. 11b). For the annealed and anodized samples tested at lower stresses, a smaller dimple fracture zone and a larger brittle fracture area was noticed (Fig. 12a). The SEM examination showed that in some regions, partial delamination of the anodic film occurred along
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Fig. 11. SEM images showing the fracture surface of an age hardened and anodized specimen. The individual micrographs correspond to different fracture modes observed from the specimen surface through the center. (a) Specimen fracture zone. (b) Cracks nucleation sites and fatigue crack growth. Multiple crack nucleation sites are present on the fractured surface. (c) Distinct fractures modes observed. The quasi-cleavage fracture zone is greater than that present in the samples prior to SAA (Fig. 9b). (d) Ductile fracture zone.
many areas, leading to the subsequent fracture of the material at the interface, as can be seen in Figs. 12a, b and c, in agreement with the findings earlier reported by Nie et al. [19]. The delamination of the film observed in Fig. 12c is related to the presence of intermetallic compounds and to the accumulation of the copper oxide particles at the film–substrate interface, as shown in Fig. 6. Fig. 12a and c also show the presence of many cavities with a large surface area just beneath the anodized film. These cavities appear to have coalesced, giving rise to the formation of a single large avoid. In general, the specimen surface showed an intergranular fracture mode. As a consequence of the propagation of multiple cracks, the ductile fracture zone (DFZ) was created (Fig. 12a). The DFZ is mostly located at the center of the fracture surface. 4. Conclusions The present study focused on the coupled effect of substrate microstructure and anodizing on the mechanical properties of a 2017A-T4 aluminum alloy especially the fatigue performance. SAA does not have
any significant effect on the static mechanical properties of the alloy, with the exception of ductility, which is significantly impaired. At high alternating stresses, the presence of the anodic film for both microstructural conditions does not seem to have a pronounced effect on fatigue crack growth. Therefore, under these alternating conditions, anodizing gives rise to a slight decrease in fatigue life. However, the results obtained from fatigue testing at low alternating stresses showed that anodizing reduces significantly the fatigue life of the alloy for both microstructural conditions. The decrease in fatigue life has been attributed to the preferential dissolution of the matrix around the cathodic particles, such as Al2Cu which gave rise to the formation of cavities. Whatever the reason of cavities initiation, these acted as stress concentrators and promoted the nucleation of many small fatigue cracks. The fatigue life reduction can also be associated to the brittle nature of the copper oxide particles accumulated at the interface. The decrease in fatigue life observed in the specimens subjected to annealing is more pronounced than that determined for the samples subjected to age hardening. For the latter condition, the intermetallic particles are coarser and therefore, they leave larger cavities after their dissolution during
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Fig. 12. SEM images showing the fracture surface of an annealed and anodized specimen. Individual micrographs corresponding to different fracture modes observed from the specimen surface through the center: (a) Specimen surface fracture. A smaller dimple fracture zone and a larger brittle fracture area can be noticed. (b) Micrograph showing film detachment. (c) Micrograph showing large cavities under the detached film. The delamination of the film is related to the presence of intermetallic compounds and to the accumulation of the copper oxide particles at the film–substrate interface. (d) Propagation of multiple cracks in the large cavity. Many cavities with a large surface area were observed to be located just beneath the anodized film.
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