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Effects of temper condition and corrosion on the fatigue performanceof a laser-welded Al–Cu–Mg–Ag (2139) alloy

A.T. Kermanidis a,*, A.D. Zervaki a, G.N. Haidemenopoulos a, Sp.G. Pantelakis b

a Department of Mechanical Engineering, University of Thessaly, GR-38334 Volos, Greeceb Laboratory of Technology and Strength of Materials (LTSM), Department of Mechanical Engineering and Aeronautics, University of Patras, 26500 Rion – Patras, Greece

a r t i c l e i n f o

Article history:Received 26 April 2009Accepted 13 July 2009Available online 16 July 2009

Keywords:C. Heat treatmentsD. WeldingE. Fatigue

a b s t r a c t

The effects of temper condition and corrosion on the fatigue behavior of a laser beam welded Al–Cu–Mg–Ag alloy (2139) have been investigated. Natural aging (T3 temper) and artificial aging (T8 temper) havebeen applied prior to welding. Corrosion testing has been performed by exposing the welded specimensto a salt spray medium for 720 h. Aging influences the corrosion behavior of laser welds. In the T3 temper,corrosion attack is in the form of pitting in the weld area, while in the T8 temper corrosion is in the formof pitting and intergranular corrosion in the base metal. In the latter case corrosion is attributed to thepresence of grain boundary precipitates. Corrosion degrades the fatigue behavior of 2139 welds. Thedegradation is equal for both the T3 and T8 tempers and for the corrosion exposure selected in this studycorresponds to a 52% reduction in fatigue limit. In both cases fatigue crack initiation is associated withcorrosion pits, which act as stress raisers. In the T3 temper, the fatigue crack initiation site is at the weldmetal/heat affected zone interface, while for the T8 temper the initiation site is at the base metal. Fatiguecrack initiation in uncorroded 2139 welds occurs at the weld toe at the root side, the weld reinforcementplaying a principal role as stress concentration site. The fatigue crack propagates through the partiallymelted zone and the weld metal in all cases. The findings in this paper present useful information forthe selection of appropriate heat treatment conditions, to facilitate control of the corrosion behavior inaluminium welds, which is of great significance for their fatigue performance.

� 2009 Elsevier Ltd. All rights reserved.

1. Introduction

Integral welded aluminium structures are a promising solutionfor aircraft airframes by being weight and cost efficient comparedto conventional riveted parts and simultaneously providing goodmechanical performance. Among the welding techniques, laserbeam welding (LBW) has been found to offer a concentrated,high-energy density heat source, increased welding speed, narrowheat affected zones as well as low residual stresses and distortions[1]. Critical for the application of laser welded structures in air-frames is their fatigue performance [2], which can be influencedby several factors, associated with the welding procedure. Thesteep thermal gradients imposed to the material during the laserwelding process lead to the generation of modified microstructuresdue to metallurgical changes in the vicinity of the weld [3]. Addi-tionally, the inhomogeneous plastic strains produced are the causeof residual stress fields generated in the near weld area [4]. Themicrostructure gradient as well as the residual stress gradient in

the heat affected zone can influence the fatigue performance ofthe weld significantly. Prediction of the microstructural changesdue to the welding process is a difficult task far from beingresolved. To fulfil this need several models have been proposed,which provide valuable input towards the prediction of the micro-structural evolution in welds [5–8]. In addition, the correlation ofthe modified microstructure to the weld’s fatigue characteristicsis also complex and far from being resolved. Thus, research activi-ties which focus on the fatigue behavior of aluminium welds aremostly experimental approaches that should be rather consideredas case studies referring to a specific material and specific weldingconditions.

Corrosion damage can be an additional significant cause of theweld’s fatigue property degradation when considering the longterm operation of an aircraft. Experimental investigations on theinteraction of corrosion and fatigue usually refer to fatigue andfatigue crack growth experiments performed in a certain corrosiveenvironment and not to tests performed on pre-corroded materi-als. Nevertheless, the behavior of pre-corroded material under fati-gue loads represents a different but more relevant situation, whenthe issue of structural integrity assessment in older airplanes isconcerned. Development of new corrosion maintenance policiesin aircraft includes methodologies for assessing the impact of

0261-3069/$ - see front matter � 2009 Elsevier Ltd. All rights reserved.doi:10.1016/j.matdes.2009.07.020

* Corresponding author. Tel.: +30 2421074014; fax: +30 2421074012.E-mail addresses: [email protected] (A.T. Kermanidis), [email protected] (A.D.

Zervaki), [email protected] (G.N. Haidemenopoulos), [email protected](Sp.G. Pantelakis).

Materials and Design 31 (2010) 42–49

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Materials and Design

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existing corrosion on structural integrity [9]. The damaging effectof prior corrosion on fatigue performance of aluminium alloyshas been studied by several researchers [10–14]. However, exper-imental information concerning fatigue behavior of corroded alu-minium welds is still lacking. The influence of corrosion on theweld’s fatigue performance can be associated mainly with: (i)introduction of corrosion notches (e.g. corrosion pits) which causestress concentrations and can act as fatigue crack initiation sitesand (ii) increased corrosion susceptibility of the modified micro-structure of the weld heat affected zone. Susceptibility of aircraftaluminium welds to environmental attack has been recognizedpresently for friction stir welds; in [15] corrosion tests on weldedaluminium alloy 7108 revealed that the weld area is susceptibleto localized corrosion, while evidence of pitting corrosion andintergranular attack on 2024 aluminium alloy welds has been re-ported in [16,17]. In addition, a few investigations have focusedon the SCC behavior of welded aluminium alloys [18,19]. Neverthe-less, presently there is a significant lack of information on the cor-rosion behavior of aluminium laser beam welds and its impact onfatigue performance, especially for the recently developed Al–Cu–Mg–Ag alloys.

The new Al–Cu–Mg–Ag alloys can offer improved mechanicalperformance (increased strength-to-weight ratios and fatigueresistance) relative to other alloys in the 2000 series [20]. It hasbeen demonstrated that controlled additions of Ag and Mg in Al–Cu alloys change the precipitation sequence. Instead of the h00

and h0 which form on the {1 0 0} planes in Al–Cu alloys a new pre-cipitate (the X-phase) is stimulated and forms as very thin hexag-onal shaped platelets on the {1 1 1} planes of the matrix [21,22].

As stated above, there is currently limited published informa-tion on the fatigue performance of weldments of the Al–Cu–Mg–Ag alloys. The aim of the present paper is to evaluate the effect ofthe aging treatment (temper) as well as the effect of corrosion onthe fatigue behavior of the Al–Cu–Mg–Ag laser beam welded al-loys. Two temper conditions have been selected, the T3 conditioncorresponding to solution heat treatment at the range of 480–520 �C for 10 min, cold work at a strain range of 2–4% followedby natural aging, with the T8 condition corresponding to artificialaging at the range of 150–180 �C for 16 h. In the T8 condition theage hardening has been enhanced by a certain amount of coldwork following quenching from the solution temperature. Saltspray has been selected as the corrosion medium for this study.The influence of corrosion exposure on the fatigue resistance ofthe welded samples has been evaluated in terms of fatigue S–Ncurves of uncorroded and pre-corroded welded coupons. Fracto-graphic analysis was conducted on the fractured samples to iden-tify fatigue crack initiation sites as well as crack path through theweld area.

2. Experimental

The Al–Cu–Mg–Ag butt-welded sheets were received fromGKSS research center (Germany) in the as welded condition withdimensions of 270 � 315 mm. Prior to welding the sheets weretempered to T3 and T8 condition. The thickness of the sheetswas 3.2 mm. The chemical composition of the base metal used inthis investigation corresponds to the 2139 alloy (Table 1). Fillerwire 4047 aluminium was used for the laser welding. A diode-pumped 3.3 kW Nd:YAG laser DY033 from Rofin-Sinar was em-ployed for the welding experiments. Laser power was adjusted at3000 W, and the focal point position was on the top of the sheet.Spot size in focus was 0.4 mm. The welding speed was 1.8 m/min. Filler wire 4047 of 1.2 mm diameter was used at a feed rateof 2.5 m/min. Surface protection was achieved by supplying 25 l/min He, while Ar was used for root protection at 15 l/min.

Machining of the fatigue specimens was performed according tothe ASTM E 466 specification (Fig. 1). All specimens were cut in thelongitudinal (L) orientation relative to the rolling direction. Fatiguetests were performed with a stress ratio R = 0.1 and a frequency of30 Hz on corroded and uncorroded welds for comparison.

Prior to fatigue testing the specimens were exposed to corrosiveenvironment in a salt spray chamber according to the ASTM B117standard. A continuous corrosion exposure in the salt fog for 720 hwas selected in order to cause sufficient corrosion damage for thefatigue performance investigation. The sodium chloride concentra-tion used for the accelerated corrosion exposure was 5% mass.Throughout the test it was ensured that inside the chamber thetemperature range was maintained at 35 ± 1 �C, while the pHwas maintained between 6.5–7.2. After exposure, the specimenswere cleaned according to ASTM G1 specification.

Standard metallographic techniques (including grinding andpolishing) were applied for the characterization of microstructures.Etching was performed by using Keller’s reagent. Fracture surfacesof fatigue specimens were examined in an SEM equipped with EDXdetector. Microhardness measurements in all weld sections wereperformed using a Vickers indenter and a load of 0.2 kg.

3. Results and discussion

3.1. Microstructural characterization of welds

A characteristic macrograph of the 2139 weld transverse sec-tion in T3 temper, revealing the different microstructures of basemetal (BM), heat affected zone (HAZ), partially melted zone(PMZ) and weld metal (WM) is shown in Fig. 2a–e. The PMZ is lo-cated adjacent to the fusion boundary, with a microstructureshown in Fig. 2d. It is characterized by a dark-etching eutecticphase and a light-etching a (Al-rich) band along the grain bound-aries (Fig. 3), as well as dark-etching eutectic constituents sur-rounded by the light-etching a phase in the grain interior. Theformation of this zone is expected during the welding of Al–Cualloys, which possess a wide solidification range.

A microhardness profile across the weld is shown in Fig. 4. Thehardness drop in the HAZ is associated with coarsening of the

Fig. 1. Fatigue specimen according to ASTM E 466. Dimensions in mm.

Table 1Chemical composition of 2139 base material (wt.%).

Si Fe Cu Mn Ag Mg Zn Ti Zr

0.04 0.06 4.79 0.3 0.34 0.45 <0.01 0.05 0.01

A.T. Kermanidis et al. / Materials and Design 31 (2010) 42–49 43


strengthening precipitates. Close to the PMZ as well as within thePMZ the hardness drop is associated with dissolution of thestrengthening phase and the formation of eutectic constituents.

The microstructure of 2139 weld transverse section in T8 tem-per can be best described with the aid of the respective microhard-ness profile across the weld bead is depicted in Fig. 5. There is agradual drop in the microhardness approaching from the basemetal to the PMZ and the weld metal, which can be attributed tothe dissolution and/or coarsening of the strengthening precipitates

[23–25]. The microhardness of the base metal is 142HV0.2 (pointA), drops to 123HV0.2 in point B, to 116 HV0.2 in the vicinity of pointC, while at the boundary with the PMZ has a value of 110HV0.2

(point D). The corresponding microstructures are shown inFig. 6a–d. The microstructure of the base metal (Fig. 6a, area A in

Fig. 2. (a) Varying 2139 material microstructure in the near weld area: microstructural gradient revealing, (b) base metal (BM), (c) heat affected zone (HAZ), (d) partiallymelted zone (PMZ) and (e) weld metal.

Fig. 3. Light-etching a (Al-rich) band along the grain boundaries in partially meltedzone. Fig. 4. Microhardness profile across the weld bead of 2139 T3 alloy.

44 A.T. Kermanidis et al. / Materials and Design 31 (2010) 42–49


Fig. 5) is characterized by the presence of grain boundary precipi-tates. In the outer region of the HAZ (area B in Fig. 5) there is sub-stantial coarsening of the grain boundary precipitates (Fig. 6b). Inthe inner region of the HAZ (area C in Fig. 5) this phase dissolvespartially (Fig. 6c) while at the HAZ/PMZ boundary (area D inFig. 5) the grain boundary precipitate dissolves completely(Fig. 6d). EDX analysis indicated that the particles appearing atthe grain boundaries were Cu-rich (Fig. 6e) without the presenceof Mg or Ag. Ideally the presence of Ag, promotes the intragranularprecipitation of the X-phase and decreases the amount of grainboundary precipitation. Nevertheless grain boundary precipitationin Al–Cu–Mg–Ag alloys tempered in the T8 condition has beenreported by other investigators as well [26]. The main differencebetween the T3 and T8 tempers is the presence of the grain bound-ary precipitates in the T8 temper.

3.2. Corrosion characterization

For the T3 temper, metallographic examination of the corrodedsamples showed high susceptibility of the laser beam welds to

Fig. 5. Microhardness profile across the weld bead for the 2139 T8 alloy.

Fig. 6. (a) Microstucture of the 2139 T8 base metal, 142 YV0.2 point A in Fig. 5, (b) HAZ microstructure 5 mm from the weld center line, 123 YV0.2 point B in Fig. 5, (c) HAZmicrostructure 3 mm from the weld center line, 116 YV0.2 point C in Fig. 5, (d) HAZ microstructure adjacent to PMZ, 110 YV0.2 point D in Fig. 5 and (e) EDX spectrumcorresponds to the grain boundaries particles shown in Fig. 6a–d.



corrosion. Localized corrosion attack was present in the PMZ in theform of extensive pitting followed by intergranular corrosion(Fig. 7a and b). Pit growth in the weld area prevailed in the PMZ,which due to the dissimilar microstructure, forms a localized gal-vanic couple leading to an anodic dissolution at the interfaceweld/HAZ (Fig. 7a). The average and maximum pit depths mea-sured in the weld area were 458 and 680 lm, respectively. Farfrom the weld the extent of corrosion damage was not significant.Pit depth measurements in the base metal area showed an averageand maximum pit depth of 64 and 100 lm, respectively.

For the T8 temper, the corrosion behavior of 2139 welds was re-versed in comparison to the T3 temper (Fig. 8a). The weld area wasfree of corrosion damage (Fig. 8b), while the base metal was signif-icantly damaged by corrosion. Corrosion damage evolved from pit-ting to intergranular corrosion as shown in Fig. 8c. The corrosion inthe base metal is attributed to the presence of the grain boundaryprecipitates, which increase the susceptibility of this alloy to inter-granular corrosion. The formation of these precipitates is associ-ated with Cu-depletion in the matrix adjacent to the grainboundaries [27]. It has been demonstrated that Cu-depleted zonesat the grain boundaries become anodic relative to the Cu-rich

zones in the grain interior, thus promoting intergranular corrosion[28]. Recently this Cu-depletion framework has been employed toexplain intergranular corrosion in Al–Cu–Mg–Ag alloys [27]. Theabsence of corrosion attack in the weld area for the T8 conditionis attributed to the dissolution of the grain boundary precipitates,as evident from the metallographic analysis presented above.

3.3. Fatigue behavior

The results of the fatigue experiments in alloy 2139 welds arepresented in Table 2 for the T3 and T8 tempers both for the cor-roded and uncorroded specimens. Cycles to failure are given as afunction of the maximum stress. The failure location is reportedas well. It should be noted that the machining of the fatigue spec-imens causes a relief of the weld residual stresses [29]. Therefore,the effect of these stresses on fatigue behavior is not taken intoaccount.

The experimental data were fitted to a three-parameter Weibuldistribution in order to obtain S–N curves. The results are shown inFig. 9. Corrosion, as expected, causes a significant degradation offatigue behavior. Calculated fixed fatigue limits at 107 cycles for

Fig. 7. Localized corrosion attack in 2139 T3 welds showing: (a) pitting corrosion at the PMZ and (b) intergranular corrosion.

Fig. 8. (a) Transverse section of corroded 2139 T8 weld showing weld area free from corrosion (b), while base metal suffers from intergranular corrosion (c).



uncorroded and corroded welds indicate a significant reduction of52% for the fatigue limit of the corroded material (55 MPa) relativeto the uncorroded material (115 MPa).

Comparing the fatigue curves of the uncorroded welds, no sig-nificant difference between the T3 and T8 tempers is observedespecially for stresses below 150 MPa. The same holds when wecompare the curves of the corroded welds. Only at stresses above100 MPa the welds of the T8 temper exhibit a slightly longer fati-gue life, a fact that will be explained below.

Fractographic analysis combined with metallography onsections of the weld was performed in order to identify the exactfatigue crack initiation sites and link with the corrosion sites andweld microstructure. Regarding the uncorroded specimens of bothT3 and T8 tempers, fractographic observations revealed that sec-ondary fatigue cracks initiate at the weld toe but only one maincrack causes fatigue fracture (Fig. 10a). These cracks initiate inthe PMZ and they propagate in two phases. During the first phasethey propagate through the PMZ in a transgranular mode(Fig. 10b). The fatigue crack propagation in the PMZ is also evidentfrom metallographic sections transverse to the fracture surface inthe middle of the crack front, shown in Fig. 10c. These figures showa strip of PMZ structure left in the fractured specimen, indicatingthat indeed the crack initiation was located in the PMZ of the weld.In the second phase of propagation the cracks enter the weld metalwhere they propagate until the critical size for the final fast frac-ture. The plane of fatigue crack propagation in the PMZ and theweld metal is perpendicular to the stress axis (mode I).

Inspection of the fracture surfaces revealed the typical appear-ance of a fatigue fracture area followed by a final fast fracture area,as shown in Fig. 11a. In Fig. 11b the fracture surface in the final fastfracture region is presented. This fracture surface is characteristicof a dendritic microstructure and confirms that the crack has prop-agated during the second phase through the weld metal. The weldreinforcement seemed to play a principal role in fatigue crack ini-tiation, since it imposes a considerable stress concentration. Thereinforcement height was measured to be around 0.4 mm, whilethe reinforcement angle was about 160�. These values indicate thatthe role of the reinforcement as a stress raiser in fatigue is signifi-cant [29]. It is reported that removing the reinforcement could re-sult in fatigue life extension, provided no other weld defects comeup to the weld surface [29,30].

Regarding the corroded specimens, the T3 and T8 welds show adifferent behavior as to the fatigue crack initiation site. For the T3temper, fatigue crack initiation is located in the weld area (as in theuncorroded welds), while for the T8 temper, fatigue crack initiationis located in the base metal. As fatigue is greatly influenced by cor-rosion, the fatigue crack initiation site coincides with the corrosionsites as described in the previous section. More specifically, fatiguecrack initiation in the welds of T3 temper was influenced by the

Table 2Fatigue test results on laser welded 2139 T3 and T8 specimens.

Material Maximum stress,rmax (MPa)

Cycles tofailure, Nf

Failure Salt sprayexposure (h)

2139 T8 150 770,819 Weld 0150 10,20,940 Weld125 30,52,866 Weld115 100,23,908 Run out

2139 T8 150 115,779 Base material 720150 84,619 Base material125 178,307 Base material125 200,000 Base material100 343,394 Base material100 360,226 Base material75 14,51,252 Base material50 100,60,028 Run out

2139 T3 175 240,310 Weld 0175 314,963 Weld150 12,88,705 Weld150 10,25,614 Weld125 102,67,443 Run out125 81,11,058 Weld115 101,63,699 Run out115 101,16,915 Run out

2139 T3 150 78,446 Weld 720125 169,660 Weld100 344,410 Weld75 100,02,000 Weld

Fig. 9. Fatigue results of uncorroded and pre-corroded laser welded 2139specimens.

Fig. 10. (a) Secondary fatigue crack initiation in PMZ and (b) strip of PMZ structure left in fractured fatigue specimen indicating propagation through the PMZ.



presence of corrosion pits at the PMZ (Fig. 12a). The PMZ is anodicrelative to the base metal and the galvanic cell created promotesthe evolution of localized corrosion damage which facilitates crackinitiation. The base metal is essentially not affected by corrosion. InFig. 12b an SEM image of the fatigue crack initiation site is shown.The site is related to corrosion damage consisting of multiple cor-rosion pits which have coalesced to form a larger defect. The stressconcentration induced by the corrosion pits in the weld dominatesthe fatigue life of the corroded sample. The large pits, which act asstress raisers reduce significantly the crack initiation stage andlead to degradation of the fatigue performance of the corrodedwelds. Due to the fact that at low fatigue stresses fatigue life ishighly dependent on the crack initiation stage, the fatigue behav-ior, which is reflected in the fatigue S–N data of Fig. 9, is justified.It shows that when the maximum stress is decreasing the failurelife difference between the corroded and uncorroded specimens in-creases up to the appearance of fatigue limit.

For the corroded welds of the T8 temper fracture by fatigue al-ways appears in the base metal. It initiates from corrosion pits atthe base metal area (Fig. 13), and propagates perpendicular tothe applied fatigue load, leading to fatigue life reduction comparedto the uncorroded welds. The reduction of total fatigue life isattributed to the reduction of fatigue crack initiation stage pro-moted by stress concentration due to existing pits at the base me-tal. These results justify the obtained fatigue behavior in the S–Ncurves shown in Fig. 9. The formation of grain boundary precipi-tates during the T8 temper moved the corrosion site to the baseplate, since these precipitates, which cause intergranular corrosion,dissolve in the HAZ, while they do not form in the weld metal. Thedifference in location of fatigue crack initiation in the T8 temper

might explain the slightly better fatigue behavior over the T3 tem-per at high stresses (above 100 MPa) since the base metal pos-sesses a higher strength relative to the weld area as it has beendemonstrated from the microhardness profiles of the welds.

The performed study presents evidence that temper conditionprior to welding in Al–Cu alloy influences the corrosion suscepti-bility of the weld. Additionally, corrosion susceptibility has beenfound to be determinant for the welds fatigue performance. Thefindings obtained indicate that selection of suitable temper condi-tion for aluminium alloys in aircraft applications may be utilizedfor control of the corrosion behavior of the welded component.Specifically, it was shown that by using the 2139 alloy in T8 temperas welded condition, corrosion damage can be isolated in the base

Fig. 12. (a) Fracture path for corroded 2139 T3 material and (b) fatigue crack initiation site. High magnification of region shown by arrow in Fig. 12a.

Fig. 13. Fatigue crack initiation from corrosion pits in 2139 T8 base material.Arrows show evidence of intergranular fracture.

Fig. 11. (a) Fracture surface of 2139 T3 uncorroded fatigue specimen and (b) fracture surface in the final fast fracture region.



material, whereas in the T3 as welded condition corrosion attack isconfined in the weld PMZ. The information obtained is importantas far as corrosion protection in the design of new aluminium alloywelds is concerned; i.e. by using suitable combination of materialtemper and corrosion surface coatings, corrosion damage effectsin Al–Cu laser welded components can be minimized.

4. Conclusions

The results presented above led to the following conclusionsregarding the effect of temper condition and corrosion on the fati-gue of Al–Cu–Mg–Ag alloy 2139 welds.

The temper condition of 2139 prior to laser welding influencesthe corrosion behavior of 2139 laser welds. In the T3 temper, cor-rosion attack is in the form of pitting in the weld area, while in theT8 temper corrosion is in the form of pitting and intergranular cor-rosion in the base metal. In the T8 condition corrosion is attributedto the presence of grain boundary precipitates in the base metal.

In both tempers corrosion damage has a deteriorating effect onfatigue behavior of 2139 welds. Stress concentration induced bycorrosion pits accelerates fatigue crack initiation and leads toreduction of fatigue life compared to uncorroded welds. Theobtained decrease in the fatigue limit is in the order of 52% withrespect to the uncorroded material in both T3 and T8 tempers. Inthe T3 temper, the fatigue crack initiation site is at the weld me-tal/HAZ interface, while for the T8 temper the initiation site is atthe base metal. Fatigue crack initiation in uncorroded 2139 weldsoccurs at the weld toe at the root side, the weld reinforcementplaying a principal role as stress concentration site. The fatiguecrack propagates through the PMZ and the weld metal in all cases.

Acknowledgments

This work has been supported by EU Wel-Air program underContract AST3-CT-2003-502832. The help of Mrs. Polina Taigani-dou with metallography and Mrs. Yiota Haidemenopoulos withSEM fractography is gratefully acknowledged.

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Author's personal copy - University of Thessaly · Author's personal copy existing corrosion on structural integrity[9]. The damaging effect of prior corrosion on fatigue performance