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The effect of hydrogen charging on the mechanical behaviour of 5083 wrought aluminum alloy
Corrosion Science 49 (2007) 4443–4451 www.elsevier.com/locate/corsci
The effect of hydrogen charging on the mechanical behaviour of 5083 wrought aluminum alloy C.N. Panagopoulos *, E.P. Georgiou Laboratory of Physical Metallurgy, National Technical University of Athens, Zografos, 15780 Athens, Greece Received 28 September 2006; accepted 23 March 2007 Available online 6 August 2007
Abstract The effect of hydrogen cathodic charging on the mechanical behaviour of 5083 wrought aluminum alloy has been studied. Hardening of the surface layers of the examined alloy, due to hydrogen absorption, was observed. The tensile tests revealed that the ductility of 5083 wrought aluminum alloy decreased with increasing hydrogen charging time, for a constant value of charging current density, and with increasing charging current density, for a constant value of charging time. However, the ultimate tensile strength of the examined alloy was slightly affected by the hydrogen charging procedure. The cathodically charged 5083 wrought aluminum alloy exhibited brittle transgranular fracture at the surface layers and ductile intergranular fracture at the deeper layers of the alloy. 2007 Elsevier Ltd. All rights reserved. Keywords: A. Aluminum; B. Cathodic hydrogen charging; B. Tensile tests; B. Microhardness
1. Introduction Hydrogen may be introduced into the metallic materials during casting, acid pickling, welding and corrosion process. Hydrogen is mainly present in metallic materials in the form of small interstitial atoms. These atoms can move rapidly either by diffusion or by transportation through mobile line defects [1].
*
Corresponding author. Tel.: +32 10 7722171; fax: +32 10 7722119. E-mail address: [email protected] (C.N. Panagopoulos).
0010-938X/$ - see front matter 2007 Elsevier Ltd. All rights reserved. doi:10.1016/j.corsci.2007.03.047
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C.N. Panagopoulos, E.P. Georgiou / Corrosion Science 49 (2007) 4443–4451
Aluminum–magnesium alloys have been widely used in the aerospace, automotive, shipbuilding and construction industries by virtue of their light weight, fabricability, physical properties, corrosion resistance and low cost. However, many aluminum alloys have shown significant susceptibility to hydrogen embrittlement [2–5]. Some of the proposed theories concerning hydrogen incorporation in metals are presented below. The first theory suggested that molecular hydrogen dissolved in the surface layers of the metallic matrix through microcracks and voids, under the influence of high pressures. The increased hydrogen concentration gradient caused increasing stresses in the matrix, which in many cases exceed the fracture toughness of the material, causing cracks in the shape of blisters [6]. However, this model cannot explain hydrogen embrittlement from lowpressure gas. Troiano [7] suggested that hydrogen diffuses mainly through easy paths, such as grain boundaries, in the surface layers of the metallic materials. There the hydrogen in interstitial site reduces the atomic cohesive strength, creating localized microdefects. The increase in the concentration and the size of microdefects leads to the formation of microcracks. Other proposals suggest that hydrogen atoms diffuse in the lattice of the metallic material and interact with the existing dislocations. The interaction between the absorbed hydrogen and the dislocations enhances the dislocation mobility, and thus increase the ductility of the metallic material [8]. The last theory is known as hydrogen enhanced localized plasticity (HELP) mechanism and was suggested by H.K. Birnbaum and P. Sofronis. However, other investigators found that the concentrated hydrogen in the surface layers of the metallic materials leads to the formation of a severely hardened region [9]. The increase of the hardness of the surface layers is attributed to strain-hardening effects caused by the absorption of the hydrogen atoms, which reduce the dislocation mobility [9]. In any case, there is no unequivocal mechanism for hydrogen embrittlement. Although, 5083 wrought aluminum alloy has been used in various technological applications, however no study concerning the hydrogen absorption of this alloy has been published. Therefore, in this research paper, a thorough investigation of the effect of hydrogen charging on the mechanical behaviour of 5083 wrought aluminum alloy has been conducted. 2. Experimental procedure The material used in this study was a commercially supplied 5083 wrought aluminum alloy sheet. Its wt% chemical composition was the following: 95.2% Al, 3.5% Mg, 0.5% Mn, 0.3% Si, 0.26% Cr, 0.24% Fe. From this sheet, a number of microhardness (4 cm · 2 cm · 0.3 cm) and standard tensile specimens were produced. Prior to the cathodic hydrogen tests, the 5083 aluminum alloy specimens were polished by SiC papers with increasing finishes and 3 lm diamond paste. In order to ensure that all the specimens had similar surface topography, roughness experiments were conducted with the aid of a Marr Perthen Profilometer. The resulting roughness of the 5083 aluminum alloy specimens was approximately 0.05 lm. A stress relief procedure was also performed, by annealing the produced 5083 wrought aluminum specimens at 350 C for 2 h and slowly cooling at room temperature, in an automatic furnace with Ar+ atmosphere. Hydrogen was introduced in the aluminum specimens by electrolytic cathodic charging. This procedure was performed in a solution consisting 75% CH3OH, 22.4% H2O and 2.6%
C.N. Panagopoulos, E.P. Georgiou / Corrosion Science 49 (2007) 4443–4451
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H2SO4, poisoned with As2O3 as hydrogen recombination inhibitor, at room temperature. The charging currents employed were in the range of 15–90 mA/cm2, while the charging time varied from 2 to 30 h, respectively, with the use of graphite anodes. Mechanical testing was carried out immediately after charging, in order to minimize the loss of hydrogen. All tension tests were performed at a strain rate of 3.3 · 104 s1, at room temperature. Microhardness testing was performed using a Shimadju Vickers indenter, imposing 0.15 N for 15 s. The given values of the experimental parameters presented are the mean value of five independent experiments. However, taking into account the experimental errors, these are accountable as error bars in the graphs. The metallographical study of the surface of both the cathodically hydrogen charged and the fractured specimens was performed with the aid of a Jeol 6100 scanning electron microscope, which was connected with an electron dispersive X-ray analyzer (EDAX). A Siemens D 5000 X-ray diffractometer, using a copper filter, was also used for the structural study of the cathodically charged aluminum specimens.
3. Results and discussion Structural characterization of the examined 5083 wrought aluminum alloy, by means of X-ray diffraction, revealed the presence of aluminum- magnesium solid solution and Mg2Si intermetallic, Fig. 1. However, after cathodic hydrogen charging, aluminum hydride (AlH3) was also detected for intense charging conditions, Fig. 1. The presence of aluminum hydride in the surface layers of hydrogen charged aluminum alloy was a result of increased hydrogen concentration, which excided the solubility limit of hydrogen in the aluminum alloy. Microhardness tests on the surface layers of the cross section of charged specimens revealed increased surface hardening, due to hydrogen charging. Fig. 2 present clear 800 Al (200)
600
Intensity (cps)
(b) Hydrogen Charged
Al (220)
400
Al (311)
Al (111)
AlH3
Al (222)
200
Mg Si 2
(a) As Received
Si
0 10
20
30
40
50
60
70
80
90
2θ
Fig. 1. (a) XRD spectrogram of as received 5083 wrought aluminum alloy and (b) XRD spectrogram of hydrogen charged aluminum specimen, under 30 mA/cm2 current density and for 30 h.
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Fig. 2. Comparative graph of the effect of charging time and charging current density conditions on the surface hardness of 5083 wrought aluminum alloy.
evidence that the microhardness of the surface layers of wrought aluminum alloy 5083 increased with increasing hydrogen charging time, for a constant value of charging current density, and with increasing charging current density, for a constant value of charging time. The increase of the microhardness of the surface layers of the examined aluminum alloy, due to hydrogen charging, can be explained in terms of dislocation pinning mechanisms. It has been proved that solute hydrogen atoms often act as dislocation pinning sites, which increase the surface hardness of the hydrogen charged alloy [10]. Dislocations may be generated by hydrogen concentration gradient, but with increasing difficulty, due to increased resistance to new dislocations production by the existing dislocations. Thus, higher charging current density and higher charging time lead to increased hydrogen fugacity in the surface layers of the alloy and consequently to increased surface hardness, due to decreased dislocation mobility. However, after a certain point (charging current– charging time), the microhardness of the surface layers tended to reach a saturation level. This phenomenon is probably attributed to the fact that the aluminum matrix became saturated and no more hydrogen could dissolve in the aluminum matrix. A typical graph, connecting the specimens microhardness with the hydrogen depth in the surface layers of the alloy, is presented in Fig. 3. From the last figure, it can be said that the hardness in the surface layers of 5083 wrought aluminum appeared increased, in comparison to the hardness of the bulk material, due to the fact that the absorbed hydrogen atoms hinder the dislocation movement, contributing to the work hardening of the alloy. Taking into consideration that the distribution profile of the measured microhardness is similar to the distribution profile of hydrogen, the diffusion profile can be seen as an approximate result of Pick diffusion mechanism [11] of hydrogen in the cathodically charged 5083 wrought aluminum alloy. Therefore, the Diffusion coefficient can be calculated, by the following mathematical equation: 2
D ¼ x2 =ð4tÞ or D ¼ ð0:025 cmÞ =4 ð21; 600Þ ¼ 7 109 cm2 s1 or 7 1013 m2 s1
C.N. Panagopoulos, E.P. Georgiou / Corrosion Science 49 (2007) 4443–4451
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Fig. 3. Hardness profile of a charged aluminum specimen at 30 mA/cm2 for 6 h as a function diffusion depth.
where D is the diffusion coefficient (m2/s1), x is the diffusion length (m) of hydrogen in the surface layers of the examined aluminum specimen, obtained from Fig. 3, and t is the charging time (s). It should be noted that the diffusion coefficient of 5083 wrought aluminum alloy was found similar to other aluminum alloys [12]. Typical engineering stress–strain curve of an uncharged and hydrogen charged (I = 30 mA/cm2, t = 6 h) 5083 wrought aluminum alloy, are presented in Fig. 4. From Fig. 4, it could be said that hydrogen cathodic charging decreases slightly the ultimate tensile strength (UTS) and significantly the ductility of the examined alloy. The observed
Fig. 4. Typical stress–strain curves for as received and charged at 30 mA/cm2 for 15 h aluminum specimen.
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Fig. 5. SEM micrograph of hydrogen-induced blistering of a charged specimen at 30 mA/cm2 for 6 h.
slight decrease of the UTS of the hydrogen charged 5083 wrought aluminum alloy could be mainly attributed to the presence of blisters and microcracks on the surface of the alloy [13], during the hydrogen cathodic charging, as it could be seen in Fig. 5. On the other hand, the observed decrease of the ductility of 5083 wrought aluminum alloy could be attributed to the fact that the absorbed hydrogen atoms act hinder the movement of dislocations and thus decrease the ductility of the alloy. In addition, the presence of blisters should also be taken in consideration. In the first series of tensile experiments, the effect of hydrogen charging time on the UTS and ductility of 5083 wrought aluminum alloy, for a constant value of charging current density, was studied. Fig. 6 present the effect of hydrogen charging time on the UTS
Fig. 6. Graph presenting the effect of hydrogen charging time on the UTS and maximum tensile elongation of 5083 wrought aluminum alloy.
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Fig. 7. Graph presenting the effect of charging current density on the UTS and maximum tensile elongation of 5083 wrought aluminum alloy.
and ductility of the hydrogen charged alloy respectively. From the above graph, it could be said that the increase in hydrogen charging time brought about an independence of the UTS of aluminum alloy, while the ductility of the alloy was substantially decreased. The decrease of the ductility of hydrogen charged alloy could be explained in terms of the dislocation pinning mechanism and blister formation, described in previous paragraph. In the second series of tensile experiments, the effect of hydrogen charging current density on the UTS and ductility of 5083 wrought aluminum alloy, for a constant value of charging time, was studied. In Fig. 7, is seen that the increase of hydrogen charging current density decreased the ductility of the examined alloy. However, the UTS was found to be independent of the charging current density. The effect of the applied current density on the UTS and ductility of the hydrogen charged alloy was decreased in comparison to the effect of hydrogen charging time. The fracture mechanism of the as received and hydrogen charged 5083 wrought aluminum specimens, which were subjected to tensile testing tests, were studied with the aid of a scanning electron microscope. A typical morphology of the fractured surface of the as received and hydrogen charged aluminum specimens are shown in Fig. 8a and b. As it can be seen from the fractograph presented in Fig. 8a, the as received specimen exhibits ductile fracture, as it consists of dimples. On the other hand, the outer area of the charged specimen, Fig. 8b, exhibits brittle interganular fracture, while the inside area the fracture remains ductile. The thin brittle interganular fracture layer might be due to the pinning of dislocation by the concentrated hydrogen. However, as the hydrogen charging time increased, for the same value of hydrogen time current, the surface intergranular fracture layer appeared slightly increased, while the dimples in the inside area became smaller and shallower, Fig. 8c. This phenomenon must be attributed to the fact that the increasing charging time of 5083 wrought aluminum alloy resulted in higher diffusion of hydrogen and thus to a higher degree of surface brittleness of the charged specimen.
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Fig. 8. (a) SEM micrograph of the fractured surface of the as received aluminum specimen, (b) SEM micrograph of the fractured surface of a charged specimen at 30 mA/cm2 for 6 h and (c) SEM micrograph of the fractured surface of a charged specimen at 30 mA/cm2 for 30 h.
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4. Conclusions From the previous experimental investigation given, the following important deductions could be made: (1) As shown by X-ray diffraction results, AlH3 formation was detected on the surface layers for specific conditions. (2) Cathodic hydrogen charging resulted in an increase of the microhardness of the surface layers of wrought aluminum alloy 5083. (3) The tensile tests revealed that the ductility of 5083 wrought aluminum alloy decreased with increasing hydrogen charging time, for a constant value of charging current density. The ductility also decreased with increasing charging current density, for a constant value of charging time. (4) The ultimate tensile strength of the examined alloy was approximately independent of the used hydrogen charging conditions. (5) The cathodically charged 5083 wrought aluminum alloys exhibited brittle transgranular fracture at the surface layers, whereas ductile intergranular fracture was observed at the deeper layers of the same alloy. Acknowledgement The authors would like to thank the Greek Research Foundation (IKY) for financial support. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13]
P.J. Ferreira, I.M. Robertson, H.K. Birnbaum, Acta Mater. 47 (1999) 2991–2998. D. Nguyen, A.W. Thompson, I.M. Bernstein, Acta Metall. 35 (1987) 2417–2425. G. Itoh, M. Kanno, T. Hugiwara, T. Sakamoto, Acta Mater. 47 (1999) 3799–3809. W.K. Jang, S.S. Kim, K.S. Shin, Scripta Mater. 40 (1999) 503–508. C.N. Panagopoulos, P. Papapanayiotou, J. Mater. Sci. 30 (1995) 3449–3456. M. Bernstein, A.W. Thompson, American Society for Metals, Ohio, Metals Park, 1974. A.R. Troiano, Trans. Am. Soc. Met. 52 (1960) 54–80. H.K. Birnbaum, P. Sofronis, Mater. Sci. Eng. A 176 (1994) 191–202. J.W. Watson, Y.Z. Shen, M. Meshii, Metall. Trans. A 19 (1988) 2299–2304. P. Rozenak, B. Ladna, H.K. Birnbaum, J. Alloys Compd. 415 (2006) 134–142. C.N. Panagopoulos, A.S. El-Amoush, K.G. Georgarakis, J. Alloys Compd. 392 (2005) 159–164. T. Ishikawa, R.B. McLellan, Acta Metall. 34 (1986) 1091–1095. C.N. Panagopoulos, A.S. El-Amoush, P.E. Agathocleous, Corros. Sci. 40 (1998) 1837–1844.
The effect of hydrogen charging on the mechanical behaviour of 5083 wrought aluminum alloy C.N. Panagopoulos *, E.P. Georgiou Laboratory of Physical Metallurgy, National Technical University of Athens, Zografos, 15780 Athens, Greece Received 28 September 2006; accepted 23 March 2007 Available online 6 August 2007
Abstract The effect of hydrogen cathodic charging on the mechanical behaviour of 5083 wrought aluminum alloy has been studied. Hardening of the surface layers of the examined alloy, due to hydrogen absorption, was observed. The tensile tests revealed that the ductility of 5083 wrought aluminum alloy decreased with increasing hydrogen charging time, for a constant value of charging current density, and with increasing charging current density, for a constant value of charging time. However, the ultimate tensile strength of the examined alloy was slightly affected by the hydrogen charging procedure. The cathodically charged 5083 wrought aluminum alloy exhibited brittle transgranular fracture at the surface layers and ductile intergranular fracture at the deeper layers of the alloy. 2007 Elsevier Ltd. All rights reserved. Keywords: A. Aluminum; B. Cathodic hydrogen charging; B. Tensile tests; B. Microhardness
1. Introduction Hydrogen may be introduced into the metallic materials during casting, acid pickling, welding and corrosion process. Hydrogen is mainly present in metallic materials in the form of small interstitial atoms. These atoms can move rapidly either by diffusion or by transportation through mobile line defects [1].
*
Corresponding author. Tel.: +32 10 7722171; fax: +32 10 7722119. E-mail address: [email protected] (C.N. Panagopoulos).
0010-938X/$ - see front matter 2007 Elsevier Ltd. All rights reserved. doi:10.1016/j.corsci.2007.03.047
4444
C.N. Panagopoulos, E.P. Georgiou / Corrosion Science 49 (2007) 4443–4451
Aluminum–magnesium alloys have been widely used in the aerospace, automotive, shipbuilding and construction industries by virtue of their light weight, fabricability, physical properties, corrosion resistance and low cost. However, many aluminum alloys have shown significant susceptibility to hydrogen embrittlement [2–5]. Some of the proposed theories concerning hydrogen incorporation in metals are presented below. The first theory suggested that molecular hydrogen dissolved in the surface layers of the metallic matrix through microcracks and voids, under the influence of high pressures. The increased hydrogen concentration gradient caused increasing stresses in the matrix, which in many cases exceed the fracture toughness of the material, causing cracks in the shape of blisters [6]. However, this model cannot explain hydrogen embrittlement from lowpressure gas. Troiano [7] suggested that hydrogen diffuses mainly through easy paths, such as grain boundaries, in the surface layers of the metallic materials. There the hydrogen in interstitial site reduces the atomic cohesive strength, creating localized microdefects. The increase in the concentration and the size of microdefects leads to the formation of microcracks. Other proposals suggest that hydrogen atoms diffuse in the lattice of the metallic material and interact with the existing dislocations. The interaction between the absorbed hydrogen and the dislocations enhances the dislocation mobility, and thus increase the ductility of the metallic material [8]. The last theory is known as hydrogen enhanced localized plasticity (HELP) mechanism and was suggested by H.K. Birnbaum and P. Sofronis. However, other investigators found that the concentrated hydrogen in the surface layers of the metallic materials leads to the formation of a severely hardened region [9]. The increase of the hardness of the surface layers is attributed to strain-hardening effects caused by the absorption of the hydrogen atoms, which reduce the dislocation mobility [9]. In any case, there is no unequivocal mechanism for hydrogen embrittlement. Although, 5083 wrought aluminum alloy has been used in various technological applications, however no study concerning the hydrogen absorption of this alloy has been published. Therefore, in this research paper, a thorough investigation of the effect of hydrogen charging on the mechanical behaviour of 5083 wrought aluminum alloy has been conducted. 2. Experimental procedure The material used in this study was a commercially supplied 5083 wrought aluminum alloy sheet. Its wt% chemical composition was the following: 95.2% Al, 3.5% Mg, 0.5% Mn, 0.3% Si, 0.26% Cr, 0.24% Fe. From this sheet, a number of microhardness (4 cm · 2 cm · 0.3 cm) and standard tensile specimens were produced. Prior to the cathodic hydrogen tests, the 5083 aluminum alloy specimens were polished by SiC papers with increasing finishes and 3 lm diamond paste. In order to ensure that all the specimens had similar surface topography, roughness experiments were conducted with the aid of a Marr Perthen Profilometer. The resulting roughness of the 5083 aluminum alloy specimens was approximately 0.05 lm. A stress relief procedure was also performed, by annealing the produced 5083 wrought aluminum specimens at 350 C for 2 h and slowly cooling at room temperature, in an automatic furnace with Ar+ atmosphere. Hydrogen was introduced in the aluminum specimens by electrolytic cathodic charging. This procedure was performed in a solution consisting 75% CH3OH, 22.4% H2O and 2.6%
C.N. Panagopoulos, E.P. Georgiou / Corrosion Science 49 (2007) 4443–4451
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H2SO4, poisoned with As2O3 as hydrogen recombination inhibitor, at room temperature. The charging currents employed were in the range of 15–90 mA/cm2, while the charging time varied from 2 to 30 h, respectively, with the use of graphite anodes. Mechanical testing was carried out immediately after charging, in order to minimize the loss of hydrogen. All tension tests were performed at a strain rate of 3.3 · 104 s1, at room temperature. Microhardness testing was performed using a Shimadju Vickers indenter, imposing 0.15 N for 15 s. The given values of the experimental parameters presented are the mean value of five independent experiments. However, taking into account the experimental errors, these are accountable as error bars in the graphs. The metallographical study of the surface of both the cathodically hydrogen charged and the fractured specimens was performed with the aid of a Jeol 6100 scanning electron microscope, which was connected with an electron dispersive X-ray analyzer (EDAX). A Siemens D 5000 X-ray diffractometer, using a copper filter, was also used for the structural study of the cathodically charged aluminum specimens.
3. Results and discussion Structural characterization of the examined 5083 wrought aluminum alloy, by means of X-ray diffraction, revealed the presence of aluminum- magnesium solid solution and Mg2Si intermetallic, Fig. 1. However, after cathodic hydrogen charging, aluminum hydride (AlH3) was also detected for intense charging conditions, Fig. 1. The presence of aluminum hydride in the surface layers of hydrogen charged aluminum alloy was a result of increased hydrogen concentration, which excided the solubility limit of hydrogen in the aluminum alloy. Microhardness tests on the surface layers of the cross section of charged specimens revealed increased surface hardening, due to hydrogen charging. Fig. 2 present clear 800 Al (200)
600
Intensity (cps)
(b) Hydrogen Charged
Al (220)
400
Al (311)
Al (111)
AlH3
Al (222)
200
Mg Si 2
(a) As Received
Si
0 10
20
30
40
50
60
70
80
90
2θ
Fig. 1. (a) XRD spectrogram of as received 5083 wrought aluminum alloy and (b) XRD spectrogram of hydrogen charged aluminum specimen, under 30 mA/cm2 current density and for 30 h.
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Fig. 2. Comparative graph of the effect of charging time and charging current density conditions on the surface hardness of 5083 wrought aluminum alloy.
evidence that the microhardness of the surface layers of wrought aluminum alloy 5083 increased with increasing hydrogen charging time, for a constant value of charging current density, and with increasing charging current density, for a constant value of charging time. The increase of the microhardness of the surface layers of the examined aluminum alloy, due to hydrogen charging, can be explained in terms of dislocation pinning mechanisms. It has been proved that solute hydrogen atoms often act as dislocation pinning sites, which increase the surface hardness of the hydrogen charged alloy [10]. Dislocations may be generated by hydrogen concentration gradient, but with increasing difficulty, due to increased resistance to new dislocations production by the existing dislocations. Thus, higher charging current density and higher charging time lead to increased hydrogen fugacity in the surface layers of the alloy and consequently to increased surface hardness, due to decreased dislocation mobility. However, after a certain point (charging current– charging time), the microhardness of the surface layers tended to reach a saturation level. This phenomenon is probably attributed to the fact that the aluminum matrix became saturated and no more hydrogen could dissolve in the aluminum matrix. A typical graph, connecting the specimens microhardness with the hydrogen depth in the surface layers of the alloy, is presented in Fig. 3. From the last figure, it can be said that the hardness in the surface layers of 5083 wrought aluminum appeared increased, in comparison to the hardness of the bulk material, due to the fact that the absorbed hydrogen atoms hinder the dislocation movement, contributing to the work hardening of the alloy. Taking into consideration that the distribution profile of the measured microhardness is similar to the distribution profile of hydrogen, the diffusion profile can be seen as an approximate result of Pick diffusion mechanism [11] of hydrogen in the cathodically charged 5083 wrought aluminum alloy. Therefore, the Diffusion coefficient can be calculated, by the following mathematical equation: 2
D ¼ x2 =ð4tÞ or D ¼ ð0:025 cmÞ =4 ð21; 600Þ ¼ 7 109 cm2 s1 or 7 1013 m2 s1
C.N. Panagopoulos, E.P. Georgiou / Corrosion Science 49 (2007) 4443–4451
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Fig. 3. Hardness profile of a charged aluminum specimen at 30 mA/cm2 for 6 h as a function diffusion depth.
where D is the diffusion coefficient (m2/s1), x is the diffusion length (m) of hydrogen in the surface layers of the examined aluminum specimen, obtained from Fig. 3, and t is the charging time (s). It should be noted that the diffusion coefficient of 5083 wrought aluminum alloy was found similar to other aluminum alloys [12]. Typical engineering stress–strain curve of an uncharged and hydrogen charged (I = 30 mA/cm2, t = 6 h) 5083 wrought aluminum alloy, are presented in Fig. 4. From Fig. 4, it could be said that hydrogen cathodic charging decreases slightly the ultimate tensile strength (UTS) and significantly the ductility of the examined alloy. The observed
Fig. 4. Typical stress–strain curves for as received and charged at 30 mA/cm2 for 15 h aluminum specimen.
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Fig. 5. SEM micrograph of hydrogen-induced blistering of a charged specimen at 30 mA/cm2 for 6 h.
slight decrease of the UTS of the hydrogen charged 5083 wrought aluminum alloy could be mainly attributed to the presence of blisters and microcracks on the surface of the alloy [13], during the hydrogen cathodic charging, as it could be seen in Fig. 5. On the other hand, the observed decrease of the ductility of 5083 wrought aluminum alloy could be attributed to the fact that the absorbed hydrogen atoms act hinder the movement of dislocations and thus decrease the ductility of the alloy. In addition, the presence of blisters should also be taken in consideration. In the first series of tensile experiments, the effect of hydrogen charging time on the UTS and ductility of 5083 wrought aluminum alloy, for a constant value of charging current density, was studied. Fig. 6 present the effect of hydrogen charging time on the UTS
Fig. 6. Graph presenting the effect of hydrogen charging time on the UTS and maximum tensile elongation of 5083 wrought aluminum alloy.
C.N. Panagopoulos, E.P. Georgiou / Corrosion Science 49 (2007) 4443–4451
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Fig. 7. Graph presenting the effect of charging current density on the UTS and maximum tensile elongation of 5083 wrought aluminum alloy.
and ductility of the hydrogen charged alloy respectively. From the above graph, it could be said that the increase in hydrogen charging time brought about an independence of the UTS of aluminum alloy, while the ductility of the alloy was substantially decreased. The decrease of the ductility of hydrogen charged alloy could be explained in terms of the dislocation pinning mechanism and blister formation, described in previous paragraph. In the second series of tensile experiments, the effect of hydrogen charging current density on the UTS and ductility of 5083 wrought aluminum alloy, for a constant value of charging time, was studied. In Fig. 7, is seen that the increase of hydrogen charging current density decreased the ductility of the examined alloy. However, the UTS was found to be independent of the charging current density. The effect of the applied current density on the UTS and ductility of the hydrogen charged alloy was decreased in comparison to the effect of hydrogen charging time. The fracture mechanism of the as received and hydrogen charged 5083 wrought aluminum specimens, which were subjected to tensile testing tests, were studied with the aid of a scanning electron microscope. A typical morphology of the fractured surface of the as received and hydrogen charged aluminum specimens are shown in Fig. 8a and b. As it can be seen from the fractograph presented in Fig. 8a, the as received specimen exhibits ductile fracture, as it consists of dimples. On the other hand, the outer area of the charged specimen, Fig. 8b, exhibits brittle interganular fracture, while the inside area the fracture remains ductile. The thin brittle interganular fracture layer might be due to the pinning of dislocation by the concentrated hydrogen. However, as the hydrogen charging time increased, for the same value of hydrogen time current, the surface intergranular fracture layer appeared slightly increased, while the dimples in the inside area became smaller and shallower, Fig. 8c. This phenomenon must be attributed to the fact that the increasing charging time of 5083 wrought aluminum alloy resulted in higher diffusion of hydrogen and thus to a higher degree of surface brittleness of the charged specimen.
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Fig. 8. (a) SEM micrograph of the fractured surface of the as received aluminum specimen, (b) SEM micrograph of the fractured surface of a charged specimen at 30 mA/cm2 for 6 h and (c) SEM micrograph of the fractured surface of a charged specimen at 30 mA/cm2 for 30 h.
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4. Conclusions From the previous experimental investigation given, the following important deductions could be made: (1) As shown by X-ray diffraction results, AlH3 formation was detected on the surface layers for specific conditions. (2) Cathodic hydrogen charging resulted in an increase of the microhardness of the surface layers of wrought aluminum alloy 5083. (3) The tensile tests revealed that the ductility of 5083 wrought aluminum alloy decreased with increasing hydrogen charging time, for a constant value of charging current density. The ductility also decreased with increasing charging current density, for a constant value of charging time. (4) The ultimate tensile strength of the examined alloy was approximately independent of the used hydrogen charging conditions. (5) The cathodically charged 5083 wrought aluminum alloys exhibited brittle transgranular fracture at the surface layers, whereas ductile intergranular fracture was observed at the deeper layers of the same alloy. Acknowledgement The authors would like to thank the Greek Research Foundation (IKY) for financial support. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13]
P.J. Ferreira, I.M. Robertson, H.K. Birnbaum, Acta Mater. 47 (1999) 2991–2998. D. Nguyen, A.W. Thompson, I.M. Bernstein, Acta Metall. 35 (1987) 2417–2425. G. Itoh, M. Kanno, T. Hugiwara, T. Sakamoto, Acta Mater. 47 (1999) 3799–3809. W.K. Jang, S.S. Kim, K.S. Shin, Scripta Mater. 40 (1999) 503–508. C.N. Panagopoulos, P. Papapanayiotou, J. Mater. Sci. 30 (1995) 3449–3456. M. Bernstein, A.W. Thompson, American Society for Metals, Ohio, Metals Park, 1974. A.R. Troiano, Trans. Am. Soc. Met. 52 (1960) 54–80. H.K. Birnbaum, P. Sofronis, Mater. Sci. Eng. A 176 (1994) 191–202. J.W. Watson, Y.Z. Shen, M. Meshii, Metall. Trans. A 19 (1988) 2299–2304. P. Rozenak, B. Ladna, H.K. Birnbaum, J. Alloys Compd. 415 (2006) 134–142. C.N. Panagopoulos, A.S. El-Amoush, K.G. Georgarakis, J. Alloys Compd. 392 (2005) 159–164. T. Ishikawa, R.B. McLellan, Acta Metall. 34 (1986) 1091–1095. C.N. Panagopoulos, A.S. El-Amoush, P.E. Agathocleous, Corros. Sci. 40 (1998) 1837–1844.