aluminum 6082 t6

12 Metallurgical Science and Technology

INTRODUCTION

Aluminium alloys have been more and more extensively uti-lised in structural applications and transportation industry[1,2] for their light weight and attractive mechanical proper-ties achieved by thermal treatments. In the Al-Mg-Si alloys,that are widely used in welded structures, the strengtheningmechanism is associated to a well known sequence of pre-cipitation in four distinct stages, during aging after solutionheat treatment and quenching, as validated by numerous ex-perimental works [3,4,5]. This precipitation may significantlyincrease hardness and tensile strength of the Al-Mg-Si alloysdepending on precipitate structure, size and distribution [6].In the utilisation of these alloys, one difficulty to be over-come is the general reduction, as compared to the parentmaterial, of mechanical properties of welded joints. This isconsequent upon the weaker strength of the WM and deterio-ration in the HAZ due to the welding thermal cycles that mayprofoundly affect the initial precipitation structure.The precipitate sequence and morphology in these alloys are

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*S. Missori and **A. Sili*Dipartimento Ingegneria Meccanica - Università di Roma - Tor Vergata - Italy

**Dipartimento Chimica Industriale e Ingegneria dei Materiali - Università di Messina - Italy

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well known in literature [7]. In particular, a process modelbased on established principles of phase equilibria, precipi-tate coarsening and dislocation precipitate interactions, wasdeveloped by Shercliff and Ashby to describe the ageing ofthe age hardening Al alloys [8,9]; moreover the works ofMyhr and Grong [10,11] deal with the microstructural sta-bility of the 6082-T6 Al alloy at elevated temperatures andshow some applications to the prediction of strength losses inthe HAZ of fusion welds. But the studies concerning theeffects produced by the heat of welding, which are peculiarof both the process and alloy considered, are not exhaustive;in fact a variety of phenomena, such as overaging associatedto precipitation process [12] or liquation cracking [13], maybe induced and the resulting metallurgical and mechanicalproperties deterioration can be controlled by the alloy start-ing temper conditions or, when it is possible, by the postweldaging procedure [14].In this paper the microstructural evolution and mechanical

Vol. 18 (1) (2000) 13

characteristics of joints welded with GMAW procedure, madeof plates of 6082 aluminium alloy in the T6 conditions, areinvestigated. The alloy here considered, belonging to the class6xxx (Al-Mg-Si), is the European near equivalent of 6061(the American variant). It was reinforced by solution at 535°Cand water quenching treatment, followed by final aging at175°C for 8-10 hours (treatment reported as T6). In a previous paper [15], the microstructural evolution ofthis alloy, due to welding, was investigated by means of SEMobservations and microanalysis measurements and correlated

to the results of tensile and microhardness tests. In the presentpaper the results of additional Vickers microhardness testsand other mechanical tests, such as fatigue rotating bendingtests and Charpy V impact tests, are reported and analysedon the basis of the microstructural evolution. In particulardiagrams of fatigue stress - cycles for parent and weldedmaterial are compared and discussed. Values of fracturetoughness are estimated in terms of K

IC, by empirical rela-

tions, from impact test energy data. Moreover fractured sur-face observations through SEM fractography are showed.

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In the welding trials are utilised plate, 10 mm thick, made ofaluminium alloy type 6082-T6 (AlSiMg0.9Mn0.7), quenchedand thermally aged.As shown in fig. 1, this material is characterised by an alu-minium matrix with the presence of a rounded and elongatedshape constituent (1-10 µm length), identified by SEMmicroanalysis as (Fe,Mn)

3SiAl

12, and of extremely fine par-

ticles and micropores (~ 0.1-0.5 µm size), uniformly distrib-

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uted. The particles were assumed to be the phase Mg2Si, not

completed dissolved during the solution treatment, and themicropores were attributed to their removal during polishingand etching operations.GMAW process was performed in four passes on bevelledplates, with the joint geometry showed in fig. 2. A filler wirehaving a composition type 4043 (~ 5% Si), which is one ofthe materials commonly recommended for obtaining a crack-

TABLE 1 - Characteristics of materials (*)

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TABLE 2 - Parameters of GMAW procedure

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The experimental work includes:

a) Tensile tests on both welded and unwelded specimens, with20x10 mm2 rectangular section and 80 mm gage length.

b) Vickers microhardness tests (500g, 10 s) along traverseson the cross welded section, in order to distinguish theeffects of the various passes. Moreover Vickersmicrohardness measurements were performed along thespecimens previously fractured by fatigue rotating bend-ing test.

c) Fatigue rotating bending tests on unnotched specimens ofcylindrical shape, diameter 9 mm, smoothly finished. Anumber of 17 specimens taken from parent metal and 15from welded sample were submitted to a standard rotat-ing bending test (ratio of the minimum stress amplitude tothe maximum stress amplitude R = -1). The load was im-posed by four symmetrically located bearings, in order toobtain a pure bending along the central portion of the speci-men.

d) SEM observation and fractography of both unwelded andwelded fatigue specimens.

e) Charpy V impact tests on both unwelded and welded speci-mens, 7.5 mm thick, standard notched. Welded sampleswere divided into two groups: the first one with notchlocation in the WM and the second one with notch loca-tion in the HAZ, centered at about 6 mm from fusion line.

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Results of tensile tests on both unwelded and welded sam-ples, performed in a previous paper [15], are here mentioned.These tests showed a remarkable reduction of both tensilestrength (from ~270 MPa down to 160-170 MPa) and elon-gation (from ~ 9% down to ~ 6%) as compared with nominal

values of the parent metal. Five samples of the six fracturedin the HAZ, the sixth one fractured in the WM (table 3). Inthe HAZ, rupture is localised at a certain distance (about 6mm) from fusion line.

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The Vickers microhardness survey along four traverses(named 1...4 in fig. 3) on the cross weld section showed thefollowing results:• in the WM the average hardness was always about 60

HV;• in the HAZ, from the initial value of about 100 HV meas-

ured in the parent metal, different behaviours for the vari-ous traverses are observed, characterised by a general re-duction, with a minimum around at 60 HV recorded at adistance variable in the range 5 to 7 mm from the weldmetal boundary line, in the region where the rupture dur-ing tensile test was observed.

The Vickers microhardness survey along the welded speci-mens fractured by rotating bending test showed the follow-ing results (fig. 4):

TABLE 3 - Tensile tests results

N. Sample Tensile Location Elongationstrength of rupture (%)(MPa)

1 Base material 276 - 8.92 269 - 8.83 177 HAZ 7.04 178 HAZ 6.55 As welded 180 HAZ 5.56 154 HAZ 6.37 165 WM 5.68 158 HAZ 6.3

free melt zone [16], was utilised. The compositions of bothparent and filler materials are reported in table 1. The GMAWprocess parameters are given in table 2.

The storage time between welding operations and mechani-cal tests was of some weeks.

Vol. 18 (1) (2000) 15

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• in the WM higher values(about 80 HV) are found nearto the external surface (line5, 6, where higher values ofstress are applied) with anincrement of 20 HV in com-parison with values measuredon samples not submitted tofatigue tests. On the otherhand no increment of hard-ness is recorded in the near-to-axis zone (no stress ap-plied).

• in the HAZ and in the parentmetal, where the appliedstress are very lower than theyield strength, no remarkablechanges are observed, ascompared to the correspond-ing values in not-cycled sam-ples (line 3, 4 of fig. 3).


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The results of fatigue test are reported in the stress - cyclesdiagram of fig. 5 and show in comparison fatigue strength ofwelded and parent metal. The experienced values of fatiguerupture stress fall in the range 130 to 280 MPa for unweldedspecimens and in the range 70 to 100 MPa for welded speci-mens, when the number of cycles is ranging between 105 and2x106. All of welded specimens fractured in the weld metal.

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SEM metallographic examination showed in the WM solidi-fication cells of aluminium matrix, surrounded by Al-Si eu-tectic particles, and in the HAZ the same structures whichcharacterised the parent metal (rounded and elongated shapeconstituent, fine particles and many micropores).Fractured surface appearance under fatigue bending tests areshown in fig. 6 and 7, for both the cases of unwelded andwelded specimens. In welded specimens fracture occurred inall cases within the melted zone.

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Charpy V impact tests showed that minimum values ofabsorbed energy were obtained in the specimens havingnotch location in the WM (~ 6.9 J), thus exhibiting a remark-able reduction of toughness. In the unwelded specimensthe representative value is about 10.6 J, while in the weldedspecimens having notch location in the HAZ a higher aver-age value (~13.0 J) was recorded (table 4).

Vol. 18 (1) (2000) 17

TABLE 4 - Charpy impact test data, estimated yield strength and calculated fracture toughness values

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Metallographic observations and SEM microanalysis meas-urements in the parent metal and HAZ, performed in [15],have shown the presence of two different constituents andmany micropores in the aluminium matrix. Both these con-stituents are initially present in the parent metal and they cannot be associated to thermal alteration consequent upon weld-ing.The microstructural evolution of the Al-Si-Mg alloys duringwelding and, in particular, the reduction of tensile strengthand the location of the minimum hardness values in the HAZare to be related to the mechanisms of deterioration of con-stituents initially present in the alloy, due to over aging in theportion of the metal submitted to temperatures high enoughto produce the evolution of precipitate, according to a wellstudied sequential formation [12]; so the hardness distribu-tion in the HAZ of the Al-Mg-Si alloys depends on the inter-play between dissolution and reprecipitation, which are com-peting processes [10].Considering the distribution of maximum temperaturesreached during welding in the HAZ [17], as well known, themore distant the zone from fusion line, the lower is the peaklocal temperature experienced by the alloy. Thus, the fol-lowing heat affected sub-zones have been envisaged.• A resolubilization zone, in the portion of metal undergo-

ing a temperature greater of 500 °C, with complete dis-solution of hardening phases and possible partial fusionin the vicinity of melt zone. Here a partial recovery ofhardness is occurred, possibly due to some degree of natu-ral aging during the storage time (some weeks) betweenwelding operations and mechanical tests.

• An over-aging zone in the portion of the metal submittedto temperatures lower than required to resolubilize the pre-cipitates, but sufficient to produce the transformation β”(the strengthening metastable precipitate) → β’ (the in-termediate metastable phase) → β (the equilibrium phase),in the range 500-380°C, and β” → β’, in the range 380-240°C, where 380° and 240°C are the assumed superiorlimits of existence of phases β’ and β” respectively.

• A slightly affected zone, experiencing peak temperatureslower than 240°C, in which supposedly no transforma-

tions of phases have been occurred, except some possiblecoalescence of constituent β”.

The above said transformations can explain also the locationof the region exhibiting the minimum hardness and the loca-tion of fracture of the tensile samples, which is in most casesin the HAZ, at a distance from the fusion line variable in therange 5 to 7 mm, according to the location of the hardnesstraverse. Here the strengthening phase β” may be transformed,with consequent loss of hardness and strength. This transfor-mation occurs when material experiences a temperature be-tween 240-380°C, which is believed to be the interval of ex-istence of phase β’.The results of fatigue tests, which report compared fatiguestrength of welded and parent metal (fig. 5), show rather re-duced values of fatigue resistance for welded specimens. Itis interesting to notice that, in spite of the fact that tensiletests generally gave rise to a fracture in the HAZ, all thewelded fatigue specimens fractured in the melted zone, start-ing from outer surface.As far as we consider hardness values after a number of cy-cles, it appears that weld metal cyclically hardens (measuredincrement of about 20 HV) near the external surface of sam-ples, where the applied stress reaches the highest values, closeto the yield strength. Both parent metal and HAZ, submit-ted to stresses very lower than the yield strength, keep theirinitial hardness.SEM images of fatigue fracture surface allow to identify theinitiation, propagation and final rupture zones (see fig. 6 and7). The aspect of fracture surface in the unwelded samplesshows an initiation zone (lower part of the figure 6a) whichsuggests a multiple crack origin, near to the outer surface,that formed a single crack front progressing toward the cen-tre (fig. 6b). The propagation zone (middle part) exhibitscleavage steps (fig. 6c); striations are not distinctly visible.The final fracture zone (upper part) has a slant section, sug-gesting dimples (fig. 6d).The aspect of fracture surface in the welded samples showsan initiation zone (lower/left part of the figure 7a) without anevident crack origin. The propagation zone (middle part ofthe section) exhibits cleavage steps; some striations are alsovisible (fig. 7b). Quite numerous pores, coming from weld-ing process, are visible. The lower fatigue resistance of welded

(*) Estimated using the regression formula (2)(**) Calculated from empirical relation (1)


REFERENCES

1. Stol, I., Selecting manufacturing processes for automotive aluminumspace frames, Weld. J., 73 (2), (1994), pp. 57-65.

2. Irving, B., Building tomorrow’s Automobiles, Weld. J. 74 (8), (1995),pp. 29-34.

3. Lutts, A., Pre-precipitation in Al-Mg-Ge and Al-Mg-Si, ActaMetallurgica, 9 (6), (1961), pp. 577-586.

4. Panceri, C., and Federighi, T., A resistometric study of precipitation in analuminum 1.4% Mg

2Si alloy, J. of Institute of Metals, 94, (1961), pp. 94-107.

5. Ceresara, S., Di Russo, E., Fiorini, P., Giarda, A., Effect of Si excess onthe aging behaviour of Al-Mg

2Si 0.8% alloy, Mat. Sci. Eng., 5, (1970),

pp. 220-227.6. Hatch, J. E., (editor), Aluminium - Properties and physical metallurgy,

Am. Soc. Metals, Metals Park, Ohio (1984)7. Miyauchi, T., Fujikawa, S., Hirano, K., Precipitation process of Al-Mg-

Si alloys by aging, J. of Japan Institute of Light Metals, 21 (9), (1971), pp. 595.8. Shercliff, H. R., Ashby, M. F., A process model for age hardening of

aluminium alloys-I. The model, Acta metall. mater., 38 (1990), pp. 1789-18029. Shercliff, H. R., Ashby M. F., A process model for age hardening of

aluminium alloys-II. Applications of the model, Acta metall. mater., 38(1990), pp. 1803-1812

CONCLUSIONS

a) Tensile strength of welded joints of 6082-T6 Al alloy, un-der the experienced welding conditions, undergo a remark-able reduction of the initial value. The residual strengthof the welded joint is around 60% of the parent metal; thisis consistent with indications of design norms such as in[19], assessing a reduction of allowable design stress of57% for this class of welded joints.

b) In the HAZ both tensile strength and hardness reduce to aminimum at a distance from the weld fusion line of about6 mm, presumably due to over-aging consequent to the trans-formation of the strengthening metastable precipitate.

c) In most of the experienced cases, fracture in tensile tests islocated in the HAZ; on the contrary, fatigue test speci-

specimens is likely to be referred, in addition to the lowermechanical properties of melt zone, to the negative influenceof porosity (fig.7c), which is responsible of a local stressraise for notch effect. The final fracture zone has a slantsection, with cleavage step appearance (fig. 7d). Charpy V impact tests showed that minimum values ofabsorbed energy were obtained in the melt zone (~ 6.9 J). Inthe parent metal the representative value is about 10.6 J, whilein the HAZ a higher average value (13.0 J) was recorded.The behaviour of the weld metal can be related both to thereduced toughness of the Al-Si5 alloy and to the presence ofsome porosity in the solidification structure. Instead, HAZtoughness is higher than in parent metal, probably due to theabove described evolution of the β” precipitate, which givesrise to a structure with reduced strength, but more ductile.Moreover, the fracture toughness values of parent metal, WMand HAZ can be evaluated by using the following empiricalrelation [18]:

K CVNIC

YS YSσ σ

= −

2

0 64 0 01. .(1)

where KIC

(MPa m ) is the plain-strain fracture toughness,

CVN (J) is the Charpy V impact test value and σYS ( MPa)

is the yield strength.Values of yield strength of the base metal, WM and HAZwere roughly estimated according to Vickers hardness (HV)by the use of the following regression formula [11]:σ

YS = 3 HV – 48.1 (2)

Results of such evaluation are reported in table 4. It can benoticed that the minimum values of toughness and K

IC are

obtained in the WM, where toughness is negatively influ-enced by the presence of some porosity. Base metal and HAZhave the same K

IC values, because in the HAZ low σ

YS val-

ues are balanced by high CVN data.

mens fractured in the WM in all cases. Fatigue fracturein the welded specimens occurred earlier than in theunwelded specimens, due to the presence of some poros-ity and to reduced mechanical properties (tensile strengthand toughness) in the WM. Cyclic behaviour after theapplication of numerous cycles of loading, reveals thatweld metal cyclically hardens (measured increment ofabout 20 HV), while both parent metal and HAZ (sub-mitted to stresses lower than the yield strength) keep theirinitial hardness.

d) Fracture toughness KIC

, evaluated from Charpy V testimpact energy data through an empirical relation, exhib-its the minimum value in WM and the same values inHAZ and in parent metal.

10. Myhr, O. R., Grong, ∅., Process modelling applied to 6082-T6 alu-minium weldments-I. Reaction kinetics, Acta metall. mater., 39 (1991),pp. 2693-2702

11. Myhr, O. R., Grong, ∅., Process modelling applied to 6082-T6 alu-minium weldments-II. Applications of model, Acta metall. mater., 39(1991), pp. 2703-2708

12. Malin, V., Study of metallurgical phenomena in the HAZ of 6061-T6aluminum welded joints, Weld. J., 74(9), (1995), pp. 305s-318s.

13. Miyazaki, M., Nishio, K., Katoh, M., Mukae, S., and Kerr, W., Quanti-tative investigation of heat-affected zone cracking in aluminum alloyA6061, Weld. J., 69 (9), (1990), pp. 362s-371s.

14. Kluken, A. O., Bj∅rneklett, B., A study of mechanical properties foraluminium GMA weldments, Weld. J., 76(2), (1997), pp. 39-44

15. Missori, S., and Sili, A., Microstructural and mechanical properties of6082-T6 aluminium alloy welds, Proceedings of the 30th InternationalSymposium on Automotive Technology & Automation, Florence, 16th-19th June 1997, pp. 241-246

16. Irving, B., Welding the four most popular aluminum alloys, Weld. J., 73(2), (1994), pp.51-55

17. Myhr, O. R., Grong, ∅., Dimensionless maps for heat flow analyses infusion welding, Acta metall. mater., 38 (1990), pp. 449-460

18. Barsom, J. M., and Rolfe, S.T., ASTM STP 466 (1970), p. 281.19. American Society of Mechanical Engineers-New York-ASME Pressure

Vessel Code Sect. VIII.

aluminum 6082 t6

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