corrosion-induced hydrogen embrittlement in aluminum · pdf file1. introduction the structural...

Corrosion Science 48 (2006) 1209–1224

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Corrosion-induced hydrogen embrittlementin aluminum alloy 2024

H. Kamoutsi a, G.N. Haidemenopoulos a,*,V. Bontozoglou a, S. Pantelakis b

a Department of Mechanical and Industrial Engineering, University of Thessaly, 38334 Volos, Greeceb Department of Mechanical and Aeronautical Engineering, University of Patras, 26110 Rio, Greece

Received 18 August 2004; accepted 13 May 2005Available online 19 August 2005

Abstract

The present paper focuses on the observed corrosion-induced embrittlement of alloy 2024and tries to answer the key question on whether the observed embrittlement is attributed tohydrogen uptake and trapping in the material. Hydrogen is produced during the corrosionprocess and is being trapped in distinct energy states, which correspond to different micro-structural sites. The formation of a hydrogen-affected zone beneath the corrosion layer is sup-ported by fractographic analysis. Removal of the corrosion layer leads to complete restorationof yield strength but only partial restoration of ductility. Additional heat treatment to releasethe trapped hydrogen leads only to complete restoration of ductility.� 2005 Elsevier Ltd. All rights reserved.

Keywords: A. Aluminium; A. 2024-T351; B. AFM; B. SEM; C. Exfoliation corrosion; C. Hydrogenembrittlement; C. Hydrogen trapping; C. Fractography

0010-938X/$ - see front matter � 2005 Elsevier Ltd. All rights reserved.doi:10.1016/j.corsci.2005.05.015

* Corresponding author. Tel.: +30 24210 74061/74062; fax: +30 24210-74061.E-mail address: [email protected] (G.N. Haidemenopoulos).

mailto:[email protected]

1210 H. Kamoutsi et al. / Corrosion Science 48 (2006) 1209–1224

1. Introduction

The structural integrity of aging aircraft structures can be affected by corrosion.As the time of an aircraft in service increases, there is a growing probability that cor-rosion will interact with other forms of damage, such as single fatigue cracks or mul-tiple-site damage. The aging aircraft may have accumulated corrosion damage overthe service life and its residual strength depends on possible degradation stemmingfrom corrosion-induced embrittling mechanisms. One characteristic example wherefailure was attributed to multi-site damage (MSD) has been the Aloha Airlines acci-dent in 1988. Damage was attributed to growth and linkage of multiple fatiguecracks, emanating from rivet holes [1]. Recent investigations on fire fighting planesof the Hellenic Aerospace Industry has also shown considerable corrosion damagearound rivet holes [2].

There are two key questions regarding this issue: (1) Is there a corrosion-induceddegradation of ductility, which in turn degrades damage tolerance and the residualstrength of aerostructures? and (2) What is the underlying corrosion-induced embrit-tling mechanism? The answer to the first question has been given by a series ofexperiments [3–5], involving mechanical testing of pre-corroded alloy 2024. It wasshown that (i) degradation of ductility and of fatigue life increases with corrosionexposure time and (ii) removal of the corrosion layer restores strength but notductility. These results were attributed to the operation of a bulk corrosion-inducedembrittlement mechanism, and it was suggested that hydrogen might be a possibleunderlying cause.

Other researchers have also considered hydrogen as an embrittlement mecha-nism in Al-alloys. Studies by Scamans et al. [6] of Al embrittlement in humidair, argued for a major role of hydrogen. In particular, the intergranular crack pathand the reversibility of the phenomenon (recovery of ductility after degassing) sup-ported a hydrogen rather than an anodic dissolution mechanism. Also, Scamansand Tuck [7] measured hydrogen permeability and stress corrosion resistance ofthe Al–Mg–Zn alloy, as functions of quench rate and aging treatment, and foundsimilar trends. Speidel [8] reviewed results up to 1992, mainly for Al–Mg–Zn alloys.More recently, Young and Scully [9] considered the kinetics of crack growth ofaluminum alloy 7050 in a humid air environment and confirmed that hydrogenembrittlement was the controlling mechanism. Jones [10] summarized evidencefor hydrogen uptake and its contribution to crack growth for the low-strength alloy5083.

Hydrogen is produced by surface corrosion reactions and part of it is absorbed inatomic form into the material [11]. In particular, the production of atomic hydrogenby a single-electron transfer process, according to the reaction

H2O + e�!OH�+ H ð1Þ

makes water an aggressive environment for aluminium alloys [9]. Absorbed hydro-gen diffuses towards the interior of the material and may be retained at various pref-erential locations. More specifically, it has been shown [12,13] that lattice defects(vacancies, dislocations, grain boundaries) and precipitates provide a variety of trap-

H. Kamoutsi et al. / Corrosion Science 48 (2006) 1209–1224 1211

ping sites. Hydrogen traps are mechanistically classified as reversible and irreversible[14], depending on the steepness of the energy barrier needed to be overcome byhydrogen to escape from them.

Thermal desorption has been successfully used to study hydrogen diffusion andtrapping in pure aluminium [15,16], Al–Cu, Al–Mg2Si [13] and Al–Li alloys [17].It has also been combined with accelerated corrosion tests in order to characterizecorrosion and hydrogen absorption in alloy 2024 [18,19]. In the last two works,the existence of multiple trapping states was verified and the quantity and evolutionpattern of hydrogen was discussed. The goal of the present study is to link hydrogenuptake and trapping to material embrittlement.

2. Experimental procedures

The material used for the present study was alloy 2024-T351, supplied in thick-nesses 1.6–3.0 mm. The chemical composition of the alloy (wt.%) is: Al–4.35Cu–1.5Mg–0.64Mn–0.5Si–0.5Fe. Exfoliation corrosion testing (EXCO) was performedaccording to ASTM specification G34-90 [20]. It included exposure at 25 ± 0.5 �C,in a solution containing 234 g NaCl, 50 g KNO3 and 6.3 ml concentrated HNO3

(70 wt.%) diluted to 1 L of distilled water. Exposure times in the EXCO solution ran-ged from 15 min to 96 h. Specimen cleaning after removal from the corrosive solu-tion involved soaking in concentrated HNO3 for 5 min, rinsing in water, then inacetone and thoroughly drying in a purge of warm air. The cleaning process aimedat the complete removal of corrosion products and its duration did not exceed10 min.

The early stages of corrosion were studied by AFM, while the later stages bySEM and metallographic sectioning. An in-house thermal desorption/gas chroma-tography system was employed in order to determine hydrogen being trapped inthe alloy during corrosion. The corroded specimens were introduced in a thermo-stated furnace and held under inert nitrogen purge. A constant heating rate of5 �C/min was applied, and the amount of hydrogen evolved was measured as afunction of specimen temperature. The following measures were taken to guaranteethat detected hydrogen originates from outgassing from the alloy microstructure: Inorder to exclude possible effect of containment, we conducted blind experimentswith all the containment in place but no specimens, and detected no measurableamount of hydrogen. In order to exclude possible effect of corrosion products, weused the above described meticulous cleaning procedure. Optical microscopy beforeand after cleaning confirmed removal of all visible corrosion products. Further,hydrogen measurements of (uncorroded or slightly corroded) specimens beforeand after cleaning produced identical results. This is taken to confirm lack of inter-ference of both the cleaning method and the amount of corrosion products withhydrogen measurement. Finally, to exclude possible effect of carrier gas, we haveused chromatographic purity N2, and have regularly confirmed the absence of O2

from the containment atmosphere. Calibration of the measurements was performed


by the use of standard mixtures (1000 ppm H2 in N2) certified by specialty gasproviders.

Hydrogen profiling was also performed, by taking hydrogen measurements frommultiple corroded specimens that were subjected to progressively deeper material re-moval. The removal was performed by grinding on SiC abrasive wheels of 1000 grid.Five identical specimens from the 2.4 mm thick plate were exposed to the EXCO cor-rosion test for 24 h. The first specimen was measured for hydrogen immediately aftercleaning, while the others were ground on all six sides by 50 lm, 100 lm, 200 lm and350 lm respectively.

Microhardness testing was performed in corroded and reference specimens, whichwere heated to various temperatures simulating the thermal desorption cycle ofhydrogen measurements. These data were set in perspective with the characteristicsof hydrogen trapping states. Microhardness measurements in the corroded speci-mens were taken after mechanical removal of the corroded layer. In order to mini-mize possible hardening effect of the removal process, a slow speed, low-pressurepolishing technique was employed. Tensile testing of corroded specimens was per-formed according to ASTM E8m-94a [21] specification. For the tests a 200 KNZwick universal testing machine and a servohydraulic MTS 250 KN machine wereused. The strain rate was 10�4 s�1. The tensile tests were conducted in the Universityof Partas and have been previously reported [3,22]. The results were adopted here.Fractography of tensile specimens was performed using scanning electron micro-scopy in order to identify the mode of fracture and provide evidence for the relationof hydrogen uptake to the deterioration of ductility.

3. Results and discussion

3.1. Microstructural description of corrosion

Corrosion in this alloy starts in the form of pitting. Pits begin to appear as early asafter 15 min of exposure and are mainly located at intersections of cracks in the pro-tective surface oxide layer (Fig. 1a). With exposure time pits become deeper and startto be connected by a network of intergranular corrosion paths (Fig. 1b). This processof pit-to-pit interactions leads to pit clustering and coalescence (Fig. 1c). From thatpoint on, corrosion does not penetrate much deeper but instead spreads beneath thesurface and causes exfoliation of surface layers (Fig. 1d).

The above type of damage evolution—and more specifically the intergranular net-work following pit growth—is responsible for the transport of corrosive solutiondeep into the material, where it reacts producing hydrogen. Thus, hydrogen is gen-erated at the front of the corrosion layer (for example at the bottom of pits) andspreads to the adjacent unaffected material establishing a hydrogen diffusion zonebelow the corrosion zone. The depth of attack, i.e. the deepest location of the corro-sion front as a function of time, is determined by metallographic sections. The valueafter 24 h of exposure to EXCO is 350 lm, as shown in Fig. 1e.


3.2. Hydrogen evolution with heating

Results of thermal desorption of hydrogen from corroded specimens are plottedin Fig. 2a. The pattern of evolution is a strong function of the specimens� tempera-ture, indicating the existence of multiple trapping sites. Different curves correspondto specimens with varying exposure time to the exfoliation solution. It is evident thatthe quantity of hydrogen increases drastically with exposure. Fig. 2b is a magnifica-tion of the y-axis and shows that a similar dependence on exposure time also holdsfor the low-temperature part of the spectrum.

Four major peaks are identified, in agreement with the experiments of Charitidouet al. [19]. The onset of the peaks labelled as T2, T3 and T4 occurs respectively at200, 410 and 495 (±10) �C. The existence of critical temperatures, below which no

Fig. 1. (a) Alloy 2024-T3 (LT plane), 15 min in the corrosive solution EXCO. 3D depiction of corrosion,initiation of pitting (arrow A) (50 · 50 lm scan). (b) Alloy 2024-T3 (LT plane), for 4 h in the corrosivesolution exco. Initiation of intergranular corrosion (arrow A). (c) SEM topographies of Alloy 2024 (LTplane) exhibiting pit-to-pit interactions and pit clustering after exposure in EXCO for (a) 8 and (b) 12 h.(d) Alloy 2024, 24 h EXCO exposure, showing exfoliation (LT direction). (e) Determination of depth ofattack (LT direction).

Fig. 1 (continued )


hydrogen evolution takes place, classifies these traps as irreversible [14]. Trappingstate T1 is found to release hydrogen continuously at temperatures close to ambientand is considered a reversible trap.

The total quantity of hydrogen liberated from each trap is estimated by integrat-ing the area under the respective peak. Results of total hydrogen quantity as a func-tion of exposure time in the corrosive solution are shown in Fig. 3. The threeirreversible traps saturate, indicating depletion of available trapping sites, in contrastto the reversible trap, T1. Trapping state T4 saturates first and also has the highestdesorption temperature. These observations support the conclusion that T4 has thehighest binding energy and thus is energetically favored.

The tentative microstructural origin of the above traps has been previously dis-cussed [19]. Trapping state T1 has been associated with interstitial lattice sites, inagreement with detailed measurements of Young and Scully [16]. Trap T2 has beenassociated with the semi-coherent interfaces of the strengthening phases or the inco-herent interfaces of dispersoids and the matrix lattice. Trap T3 has been attributed tothe formation of Mg hydride, as proposed by Scamans and Tuck [7]. Finally,

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a

b

Fig. 2. (a) Desorption of H2 in specimens of aluminium alloy 2024-T3 for continuous heating up to600 �C. Thickness of material 1.8 mm. (b) Magnification of (a) in the region up to 1.0 lgH2/min.


trapping state T4 starts to release hydrogen at a temperature coinciding with thedissolution of the strengthening Al2CuMg (S) phase. The ability of all the abovemicrostructural elements to serve as hydrogen traps has been demonstrated by Sai-toh et al. [13] using tritium autoradiography.

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0 20 40 60 80 100 1200 20 40 60 80 100 120

Fig. 3. Amount of hydrogen desorbed from the four trapping states (T1–T4) as a function of corrosionexposure time.


A closer look at Fig. 2a indicates the possible existence of additional trapping sites,manifested by two shoulders between traps T2 and T3. The amount of hydrogen asso-ciated with these states is evidently too small compared with the amount related to themain traps T3–T4, and the former are masked by the latter. The origin of these minorstates may be related to dislocations and vacancies respectively of the matrix, as hasbeen demonstrated by thermal desorption measurements in high purity aluminum [16].

Another key question addressed by the present experiments is how deep into thematerial does hydrogen penetrate. Sequential sectioning of specimens exposed toEXCO for 24 h provides the hydrogen profile. The quantity of hydrogen remainingin the specimens as a function of the depth of material removed is depicted in Fig. 4.The amount of hydrogen in the uncorroded material is also shown, and indicatesthat the hydrogen affected zone extends beyond 350 lm, which is the corrosion depthof attack shown in Fig. 1e. Based on the above data, an estimate of the depth ofhydrogen penetration is 400–500 lm.

There are some fundamental difficulties in the application of diffusion analysis tointerpret/predict the above results. First, hydrogen is generated at the front of the

-50 0 50 100 150 200 250 300 350 400 4500

20

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µgH

2

Depth (µm)

uncorroded

Fig. 4. Total amount of H2 (lg) for different removal depths. 24 h exposure to EXCO solution.


corrosion zone and part of it diffuses towards the unaffected material, establishing ahydrogen-affected zone. However, as the experiment progresses, the boundary be-tween the two zones is continuously displaced towards the interior: hydrogen-affected material becomes corroded (possibly releasing the trapped hydrogen), whilenew material is affected. Second, because of the highly non-uniform shape of the cor-rosion zone (see Fig. 1e), uncorroded but hydrogen-affected material is removed, to-gether with corrosion products, during sectioning. Thus, the amounts of hydrogenshown in Fig. 4 correspond mostly to the deeper regions of the affected zone. Third,there is a four-order-of-magnitude scatter (10�11–10�15 m2/s) in reported values ofthe diffusivity of hydrogen in aluminum [23,24].

Keeping the above complications in mind, we may use the experimentally deter-mined penetration depth of 400–500 lm to estimate by naive transient diffusion ana-lysis an effective diffusivity of

Deff ¼d2

16t� 10�13 m2=s ð2Þ

where d = hydrogen penetration depth and t = diffusion time (� exposure time24 h).

This value should be interpreted as a trap-affected diffusivity, because the speci-mens in the present work are initially free of hydrogen and become charged duringthe corrosion process. Recent experiments by Young and Scully [16] have helped ex-plain the large scatter of H2 diffusivity values, and in particular have shown that theactivation energy of dislocation traps in pure polycrystalline aluminium leads to aroom-temperature diffusivity of order of 10�13. As discussed above, the main trap-ping states of the presently considered aluminum alloy 2024 act similarly to disloca-tions, in that they produce elastic stress fields at the interfaces of the strengtheningphases and dispersoids with the matrix lattice. In view of the above, the value of


diffusivity estimated by the experimentally determined hydrogen penetration depthappears reasonable.

3.3. Microhardness measurements

The hydrogen spectra presented in the previous paragraphs provide some cluesabout trapping states, based on the temperatures where hydrogen evolution takesplace. However, during the thermal cycle of hydrogen desorption the material under-goes microstructural changes that affect its hardness. A comparison between the hard-ness of corroded and uncorroded material that experience the same heating history isshown in Fig. 5, and differences are set in perspective with the hydrogen spectra.

As expected, the hardness of the uncorroded material decreases with temperatureup to 400 �C due to overaging (a slight increase at 250 �C is attributed to a secondaryageing peak), while it rises again above 400 �C due to dissolution and reprecipitation.The corroded material exhibits a similar behavior in general. However there are twodistinct differences in the hardness values. The corroded material appears softer thanthe uncorroded up to 300 �C and harder above 350 �C. The first temperature regionincorporates states T1 and T2. Hydrogen in the T1 state is arguably trapped in inter-stitial sites while T2 hydrogen is trapped at particle/matrix interfaces. Recent re-search indicated that hydrogen might aid dislocation motion by altering thebinding energy between dislocations and solute atoms [25,26]. On the other hand,hydrogen at matrix/particle interfaces might alter the interaction energy between dis-locations and particles for coherency hardening in a way similar to that proposed forlattice dislocation. The result is that hydrogen in T1 and T2 states probably aidsdislocation motion in overcoming either the lattice resistance or semicoherent parti-cle resistance and the corroded material appears softer.

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Fig. 5. Microhardness profile versus temperature and hydrogen spectrum for alloy 2024.


The second temperature region (above 300 �C) incorporates states T3 and T4. It ispresently argued that T3 hydrogen is trapped as Mg hydride, while T4 hydrogen istrapped in the main hardening phase (h or S). As is apparent from Fig. 5, hydrogenaffects the dissolution and reprecipitation process or even the overaging (coarsening)process in the alloy and the corroded alloy appears harder. However more work isneeded to clarify this behavior.

3.4. Mechanical testing and fractography

In order to establish the effect of corrosion and hydrogen trapping on strengthand ductility, tensile testing of corroded samples [22] and corresponding fracto-graphic analysis were undertaken. All corroded specimens were exposed for 24 hand the following conditions were considered: Corroded specimens (C), corrodedspecimens from which the corrosion layer was subsequently removed (CR), corrodedspecimens heat treated at 495 �C (CH), and corroded specimens with both corrosionremoval and heat treatment (CRH). The reference material consisted of uncorrodedspecimens subjected to the respective material removal and heat treatment (condi-tions U, UR, UH and URH respectively).

Elongation to fracture and energy density are the measurements used to assesstensile ductility of the specimen. The energy density, U, is defined as:

U ¼Z ef

0

r � de ð3Þ

and represents the energy per unit volume absorbed for the deformation and fractureof material. Energy density is presently evaluated as the area under the engineeringstress–strain curve as an approximation, since no significant necking is observed.

Tensile test results are given in Table 1. The tensile results for each test series rep-resent the average value of four tensile tests. The deviations typically did not exceed1% for ultimate tensile strength (UTS), 2.5% for yield strength, 6% for elongationand 5% for energy density. The data are presented as percentages of the respectivereference values in Fig. 6. Corroded material (condition C) exhibits a yield strength92% and an elongation 39% that of the uncorroded material. Heating of the speci-men at 495 �C (condition CH) does not affect yield strength, however it restores

Table 1Results of tensile tests

Condition Yield strength (MPa) UTS (MPa) Elongation (%) Energy density (MJ/m3)

U 376 489 15 72UH 327 496 17 81UR 358 472 16 74URH 296 447 14 59C 349 425 6 25CH 303 433 10 40CR 363 468 11 49CRH 287 438 14 57

Ref C CH CR CRH0

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80

100%

Valu

e of

pro

pert

y

Condition

Yield strength UTS Elongation Energy Density

Fig. 6. The results of Table 1 presented as percentages of the respective reference values.

Fig. 7. Fracture surfaces of condition C (corroded 24 h) specimen. Individual photos correspond todifferent modes of fracture from the specimen surface (on start of corrosion) through the centre of thespecimen.



ductility to 56% of the uncorroded heat treated (condition UH). Removal of the cor-rosion layer (condition CR) restores the yield strength to 100%, while ductility isonly partially restored to 66% that of the reference value. Removal of the corrosionlayer followed by heating at 495 �C fully restores both the yield strength and elonga-tion to the reference value (condition CRH).

The tensile test results are interpreted as follows: Corrosion affects a surface layerof the material by introducing defects such as pitting, exfoliation and intergranular

Fig. 8. Fracture surfaces of condition CR (corroded 24 h + removal of 350 lm) specimen. Magnificationof the embrittled region and, transition to ductile fracture.


cracks. These defects degrade the strength of the material and lead to a deteriorationof yield strength, due to a decrease of the load carrying capacity of the corrodedlayer as well as due to stress concentration caused by corrosion notches. Removalof material up to the maximum depth of attack removes all the above defects andthus restores strength. Ductility is also degraded by corrosion, but is only partiallyrestored by removal of the corroded layer. This constitutes evidence of a bulk mech-anism of embrittlement. Subsequent heating to a temperature where all hydrogentraps are activated and release hydrogen fully restores ductility, implicating hydro-gen embrittlement as the responsible mechanism.

Additional evidence in favor of a hydrogen mechanism comes from the results offractographic analysis of tension specimens (condition C) presented in Fig. 7. Themost interesting finding is that a quasicleavage (embrittled) zone always appears be-tween intergranular corrosion and the ductile uncorroded matrix. This observation isconsistent with the existence of a hydrogen affected zone, which degrades the ductil-ity of the material even when the corrosion layer has been removed.

The proposed mechanism of hydrogen induced embrittlement of alloy 2024 issketched in Fig. 8. Corrosion creates damage, which penetrates in a non-uniformway from the surface to the interior of the material, resulting (for 24 h exposure inEXCO) in a depth of attack of 350 lm. Below the corrosion layer, a hydrogenzone is established due to the diffusion of hydrogen released by the corrosion reac-tions (Fig. 8a). Removal of a uniform layer equal to the depth of attack, removesall the corrosion damage and all hydrogen trapped in the 350 lm surface layer inthe alloy. The result is an elongation increase to 66% of the reference value. How-ever, part of the hydrogen zone still remains (Fig. 8b). Heating of the alloy to495 �C liberates all hydrogen (Fig. 8c) and restores ductility back to the originalvalue.

4. Conclusions

The experiments performed in this work lead to the following conclusions regard-ing corrosion-induced hydrogen embrittlement in aircraft Al-alloys:

1. Corrosion damage starts with pitting and proceeds to pit-to-pit interactions, inter-granular attack and exfoliation.

2. Hydrogen is produced during the corrosion process and is being trapped in dis-tinct states in the interior of the material.

3. The temperature of hydrogen evolution and the variation of hardness with heat-ing provide indirect evidence of the microstructural nature of hydrogen traps.

4. The hydrogen front advances with the corrosion front and establishes a hydrogen-affected zone beneath the corrosion layer.

5. Removal of the corrosion layer leads to complete restoration of yield strength butonly partial restoration of ductility. Removal of the corrosion layer and heating toactivate all hydrogen traps leads not only to complete restoration of strength butalso to complete restoration of ductility.


6. Detailed fractographic analysis shows the existence of a quasicleavage transitionzone between the intergranular corrosion zone and the ductile corrosion-un-affected material. This quasicleavage zone has been embrittled by hydrogen diffu-sion and trapping.

These results constitute evidence of hydrogen embrittlement in aluminum alloy 2024.

Acknowledgment

Part of this work has been financially supported by the Greek Secretariat ofResearch and Technology (GSRT) and by AirBus.

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