Hot Deformation of AA6082 Containing FineIntermetallic Particles

CECILIA POLETTI, TOMASZ WOJCIK, and CHRISTOF SOMMITSCH

Hot deformation of AA6082 aluminum alloy was studied by compression tests carried outbetween 573 K and 823 K (300 �C and 550 �C) under a wide range of strain rates. Light opticaland scanning electron microscopy were used to study the as-received microstructure, whichconsisted of elongated, partially recrystallized grains containing fine Mg2Si and AlFeMnSiparticles. The hot-deformed material showed the effects of dynamic recovery, i.e., small lowangle grain boundary formation and dislocation pinning by fine particles. The flow data wereused to calculate the constitutive equations, obtaining high values of n exponent. This behaviorwas attributed to the interaction of particles with dislocations during hot deformation.Threshold stresses were introduced to adjust the constitutive equation to a n exponent value of 5at high stresses and a value of 3 in the low stresses range, which was related to dislocations’climbing and sliding and thus to dynamic recovery. The threshold values were related to thedetachment stresses in close connection with the precipitation state which was a function of thedeformation temperature.

DOI: 10.1007/s11661-012-1487-8� The Minerals, Metals & Materials Society and ASM International 2012

I. INTRODUCTION

A. Hot Deformation in Aluminum Alloys

HOT deformation of Al-Mg-Si alloys is of industrialimportance due to their increasing use in structural lightparts as a consequence of their high specific bendingstiffness and strength. During its processing, large ingotsare hot rolled in a multistep mill after homogenizationand before aging. Homogenization of AA6082 is carriedout in order to spheroidize plate-like particles and toprecipitate dispersoids.[1] Afterward and during hotdeformation, dynamic recovery (DRV)[2] is the mainrestoration mechanism that takes place in aluminumalloys due to the high stacking fault energy[3] of thesematerials. If large strains are reached, dynamic recoverycan be followed by grain pinching off which producesnew grains by geometric dynamic recrystallization(gDRX).[4–6] All these dynamic processes are alsoinfluenced by the interaction between dispersed particlesand dislocations or boundaries. In this way, smallprecipitates can pin or retard the movement of disloca-tions, grain and subgrain boundaries by means of theZener-drag effect as shown in many works such as inReference 7. On the other hand, large particles can

promote dynamic recrystallization,[8] although this effectis not commonly observed in aluminum alloys.The DRV occurs by developing a substructure and

reaching a constant subgrain size (diameter) dss duringhot deformation when steady state flow is reached. Thesize of the subgrains can be correlated to the Zener–Hollomon (Z) parameter as used in Reference 9

ln dSS ¼ B� C lnZ ½1�

where B and C are empirical constants.

B. Constitutive Equations

The phenomenological constitutive equations areused to describe the interdependence of the stress rwith the strain rate _e and temperature T during hotdeformation. The Z parameter is defined as

Z ¼ _e exp Q=RTð Þ ½2�

where Q is the apparent energy of activation and R isthe universal gas constant. Furthermore, Z can berelated with the flow stress and temperature using theEq. [3] for low stresses or the universal Eq. [4] for awide range of stress values.[10]

APrnP¼_e exp QP=RTð Þ ¼ ZP ½3�

A sinh arð Þn¼ _e exp Q=RTð Þ ¼ Z ½4�

with A, b, a, and n as material constants and P as asubscript related to the power law.In this work, the Eqs. [3] and [4] were replaced by the

modified ones as used in Reference 11, in which thenormalized stress replaces the stress

CECILIA POLETTI, Associate Professor, is with the Institute forMaterials Science and Welding, Graz University of Technology,Kopernikusgasse 24/I, 8010 Graz, Austria. Contact e-mail: cecilia.po-letti@tugraz.at TOMASZ WOJCIK, Project Assistant, is with theInstitute of Materials Science and Technology, Vienna University ofTechnology, Favoritenstrasse 9-11, 1040 Vienna, Austria. CHRISTOFSOMMITSCH, Professor, is with the Institute for Materials Scienceand Welding, Graz University of Technology, and also with theChristian Doppler Laboratory for Materials Modeling and Simula-tion, Graz University of Technology.

Manuscript submitted March 15, 2012.Article published online October 23, 2012

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 44A, MARCH 2013—1577

Ap1rG

� �np1¼ _e exp Q=RTð Þ ¼ Zp ½5�

A1 sinh a1rG

� �� �n1¼ _e exp Q=RTð Þ ¼ Z ½6�

in which the subscript 1 appears to differentiate be-tween the equations where the reduced stress is usedand the others. The shear modulus G depends on thetemperature for aluminum alloys[12,13] as follows:[14]

G MPað Þ ¼ 30220� 16T ½7�

In the constitutive equations, the n exponents can becorrelated to the deformation mechanism[15] and thus tothe development of the microstructure. In this way, n =1 is related to diffusional flow, meaning vacancies’movement at grain boundaries (Coble creep) or in thebulk (Nabarro-Herring), n = 2 describes the grainboundary sliding, and n between 3 and 5 is related todislocation creep.[9] In the case of n = 8, the proposedbehavior of the microstructure is the invariant creepmechanism,[16] in which the second hard phases such asdispersoids or particles can pin the subgrain movement,and deformation takes place without any change in themicrostructure.[17]

The objective of this work is to study the hotdeformation mechanisms of a hot-rolled AA6082 alu-minum alloy containing large amount of intermetallicphases taking place during hot deformation in a widerange of temperatures and strain rates. The goal is tofind a correlation of the constitutive equations with thedeveloped microstructure during hot deformation withT and _e.

II. EXPERIMENTAL

A. Material

The alloy used in this study was AA6082 aluminumalloy with the following nominal chemical compositionin wt pct: Mg = 0.75 to 0.85, Si = 1.1 to 1.2, Mn =0.45 to 0.055, Fe maximum = 0.3, Cu maximum = 0.1,and balanced with aluminum. The plate was homoge-nized at 773 K (500 �C) with a standard procedure,rolled in 5 steps at a starting temperature of 823 K(550 �C), and cooled down in air.

B. Compression Tests

Cylindrical samples of 20-mm length and 15-mmdiameter were machined with a Rastagaev shape withthe main axis perpendicular to the rolling direction. Thecompression tests to determine the flow curves wereperformed using a Servotest machine in a range oftemperatures from 573 K to 823 K (300 �C to 550 �C)and between 0.001 and 500 s�1 of constant strain rate.

Some compression tests at specific deformationparameters were carried out in a Gleeble� 1500 servohydraulic machine on cylindrical specimens of 15-mmlength and 10-mm diameter to produce samples formetallographic studies. The cylinders were heated by the

ohmic resistance of the material to the electrical current.To avoid static restoration mechanisms after hot defor-mation which can lead to misinterpretations of thedeveloped microstructure,[18,19] water quenching wasapplied in situ immediately after hot compression.

C. Metallography

The samples were cut along the compression axisbefore and after deformation tests and then embeddedfor metallographic preparation. All the samples wereground and polished up to 0.04 lm in SiO2 colloidalsolution. The microstructure of the as-received condi-tion was observed by means of a light optical micro-scope (LOM) and polarized light after electro-etchingthe surface with a Barker’s solution for 3 minutes. TheEBSD measurements were carried out for the as-received condition as well for samples deformed withthe Gleeble� 1500 machine by means of a field emissiongun—environmental scanning electron microscope(FEG-ESEM) FEI device, at 20 kV and spot sizenumber 6. Micrographs were also taken using the backscattered electron mode (BSE) to obtain contrast fromdifferent phases as a function of the elements’ distribu-tion. Subgrain sizes were determined from the EBSDmeasurements as described elsewhere.[12]

The angular sensitivity of the EBSD measurements islimited up to 0.5 deg, as reported by F. J. Humphreys inReference 20, of the misorientation angles obtained byEBSD compared to other methods,[21] leading to anoverestimation of the high angle grain boundary (HAGB)content. For this reason, some subgrain boundary mis-orientations were measured by means of transmissionelectron microscopy (TEM) techniques. Angle subgrainboundary below the limit of detection by the EBSDmethod was determined using Kikuchi diffraction inTEM. For this purpose, one subgrain was tilted into [001]zone axis and a diffraction pattern of this subgrain wastaken using convergent electron beam. Kikuchi diffrac-tion patterns were obtained on all neighboring subgrains.Due to the fact that the position of the Kikuchi lines isvery sensitive to the orientation angle of the observedsubgrain, small deviations from the exact zone axis couldbe determined.[22] The angle was calculated using theradius of the ZOLZ (zero-order Laue zone) measured inthe diffraction pattern and the known radius of theEwald-sphere. This procedure was repeated on severalregions of intersecting subgrain boundaries. The analysisof Kikuchi patterns was carried out using ‘‘jems’’.[23]

Finally, EDX analysis were carried out in the STEMmode in order to determine the chemical composition ofsome dispersed phases.

D. Differential Scanning Calorimetry (DSC)

DSC tests were carried out with a TMA 2940 CEthermal analysis equipment from TA Instruments. Theas-received sample was heated at 10 K/minutes to 823 K(550 �C) and tested against a pure aluminum samplesuch as in Reference 24. Specific heat flow was plotted asa function of the temperature; endothermic peaks(down) represent the dissolution of phases while

1578—VOLUME 44A, MARCH 2013 METALLURGICAL AND MATERIALS TRANSACTIONS A

exhothermic ones (up) are related to precipitationphenomena.

III. RESULTS

A. Microstructure Before Deformation

Figure 1 shows the microstructure of the hot-rolledsample in the as-received condition. The grains present aplate-like shape and are oriented in the rolling direction(Figure 1(a)). Figure 1(b) depicts the IPF result fromthe EBSD measurement of the RD plane. The grainboundaries (represented by misorientations larger than15 deg) are plotted with thicker lines than the subgrainones (1 to 15 deg). The microstructure consists ofaround 60 vol. pct of recrystallized grains and 40 vol. pctof grains which present a substructure with an averagesubgrain size of 7.5 lm. The mean grain size is around320 lm and is obtained from LOM pictures under theASTM E112—10 norm. Figure 1(c) and d show FEG-SEM pictures in BSE mode in the as-received conditionand after exposure at 773 K (500 �C) during 30 minutes,respectively. Fe, Mn, and Si containing aluminideintermetallic phases (AlFeMnSi) are shown bright, while

Mg2Si fine needles (Figure 1(c)) and coarse precipitates(Figure 1(d)) are dark. All the precipitates are confirmedby EDX local measurements. Image analysis using‘‘ImageJ 1.45 s’’ software results in about 3 vol. pct ofAlFeMnSi with a mean radius of 80 nm, a median of70 nm, and 1 vol. pct of Mg2Si particles with a meanand median radius of 60 nm and a mean ferret minimumof 30 nm.

B. DSC Results

Figure 2 shows the thermograph of the as-receivedmaterial tested at a heating rate of 10 K/ minutes, andit was interpreted according to the literature forsimilar Al-Mg-Si phases.[25] The low temperature peakin the range of 423 K to 493 K (150 �C to 220 �C)represents the dissolution of GPZones formed duringnatural aging. At around 543 K (270 �C), the meta-stable phase b¢¢ precipitates and transforms into themore stable b¢ phase at 573 K (300 �C) by precipita-tion of the excess of Si. The phase b¢ transforms intothe most stable and large phase b at around 743 K(470 �C). The stable phases b and Si dissolve thenabove 773 K (500 �C).

Fig. 1—AA6082 rolled material in the as-received condition showing (a) disk shape grains (LOM after electro-etching with Baker’s), (b) partiallyrecrystallized grains (inverse polar figure, subgrain boundaries shown as thinner lines than grain boundaries) and (c, d) Mg2Si (black) andAlFeMnSi (white) in the as-received condition and after 30 min at 773 K (500 �C), respectively (FEG-SEM, BSE mode).

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C. Flow Curves

The flow curves at 573 K, 673 K, and 773 K (300 �C,400 �C, and 500 �C) for all the strain rates are shown inFigure 3. In general and after a small plastic strain isreached, the curves present steady-state flow for alltemperatures and strain rates, except for 773 K (500 �C)in which some softening is observed.

D. Microstructure After Hot Deformation

TEM pictures and EDX measurements show thepresence of AlFeMnSi phases in Figure 4. Figure 5shows the inverse pole figures of four samples after hotdeformation at different parameters. The subgrain sizeincreases with increasing temperature and decreasing thestrain rate. The micro- and substructures are heteroge-neous. The hot rolling microstructure remains basicallyunchanged: grains are compressed even more and

Fig. 2—DSC thermograph of the as-received material showing disso-lution and precipitation of phases (exo up).

Fig. 3—Flow curves at strain rates from 0.001 to 500 s�1 and deformation temperatures of (a, b) 573 K and 673 K (300 �C and 400 �C) and(c) 773 K (500 �C) showing steady state and softening flows, respectively.

Fig. 4—(a) Intermetallic phase of around 200-nm wide composed by Al, Fe, Mn, and Si, as shown in (b) EDS measurement. (STEM).

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heterogeneous bands are observed. In some grains, anew substructure is recognized in which the averagesubgrain size depends only on the temperature and thestrain rate. Furthermore, serrated grain boundaries areobserved at low temperatures and high strain rates.Samples deformed at and above 773 K (500 �C) addi-tionally show some grain boundary pinching off, withthe consequent formation of new grains by geometricdynamic recrystallization (gDRX), and were shown inReference 14.

The precipitate state and the subgrains of the samplesafter deformation at low and high temperatures areshown in Figure 6. The intermetallic aluminides of Feand Mn are observed bright, in a range size of some100 nm to around 1 lm. The Mg2Si particles observedin the as-received condition in a needle shape are

retained at 573 K (300 �C), while they are lesser (in vol.pct) and coarser after deformation at 773 K (500 �C).Small AlMnFeSi intermetallic particles appear decorat-ing the grain and subgrain boundaries for both temper-atures, while at 573 K (300 �C) of deformation, smallneedles of Mg2Si are arranged at the boundaries andwithin grains.Figure 7 shows the subgrain structure of samples

deformed at 773 K (500 �C) at two different strain ratesobtained by the TEM technique. The here shownmisorientations among subgrains are in the order of0.3 ± 0.2 deg and thus cannot be detected by EBSD. Itcan be observed that at lower strain rates, largermisorientations among boundaries are measured. Fur-thermore, small intermetallic particles are located atthese subgrain boundaries, which appear bent at these

Fig. 5—Inverse pole figures of EBSD measurements of samples deformed at (a) 1 s�1 and 823 K (550 �C) and (b) 0.001 s�1 and 573 K (300 �C).High angle grain boundaries are shown in black.

Fig. 6—FEG-SEM micrographs of samples after deformation showing Mg2Si phase (dark), iron, and AlFeMnSi (bright) at (a) 573 K (300 �C)and 0.1 s�1 and (b) 773 K (500 �C) and 0.001 s�1.

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 44A, MARCH 2013—1581

zones. A detail in Figure 8 shows dislocations attachedto the intermetallic phases after hot deformation at773 K (500 �C) and 0.001 s�1.

Misorientation distribution is plotted in Figure 9 as afunction of the temperature of deformation and strainrate. Grain misorientation values between 40 and 50 degare more frequent after deformation at higher temper-atures. This is related to the fact that at lowertemperatures, relatively smaller subgrains form than athigh temperatures, increasing the relative amount of lowangle grain boundaries.

E. Constitutive Equations

The flow peak stresses are used to build the consti-tutive equations. Using the creep equation, an exponentvalue np[1] of 11.75 is obtained, very close to the values

obtained in Reference 26 for the same alloy in theannealed condition and in Reference 27 for 6063. Due tothe large range of temperature and strain rates used forthe hot deformation tests, Eq. [6] is used to obtain abetter correlation of the deformation parameters withthe stress values. The general correlation found betweenstress r and Z parameter is

Z ¼ A1 sinh 500rG

� �h i7½8�

From the Arrhenius dependence of Z, the apparentactivation energy Q using Eq. [8] is found to be125 kJ mol�1 K�1, which is close to the self-diffusion

Fig. 7—TEM figures of two samples deformed at 773 K (500 �C) and at (a) 1 s�1 and (b) 0.001 s�1 showing subgrain boundaries with misorien-tations below 2 deg. The intermetallic particles (black) can be seen blocking dislocations and subgrain boundaries.

Fig. 8—TEM picture of the sample deformed at 773 K (500 �C) and0.001 s�1 showing detached dislocations at the intermetallic particles(arrows).

Fig. 9—Misorientation distribution as a function of the deformationparameters.

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of aluminum at the bulk, reported to be between 120and 144 kJ mol�1 K�1.[28]

IV. DISCUSSION

The large value of n1 calculated with the constitutiveequation in Eq. [6] (n1 = 7) should be related to thepinning of dislocation movement by the fine dispersoidspresent in the material. If a threshold stress is taken intoaccount, then Eq. [6] can be rewritten as

A2 sinh ar� rth

G

� �n� �¼ _e eðQ=RTÞ ¼ Z ½9�

as was done in the work of Spigarelli et al.[29] and inReference 30 for the same alloy after overaging.

The introduction of a threshold stress in the presentwork resulted in two equations:

a) At low logZ (up to 15), the equation can be writtenas

A3 sinh 500r� rth

G

� �� �3¼ _e eð125 kJmol�1K�1=RTÞ ½10�

b) In the high Z values range (between 15 and 32), theexponent parameter changes from n = 3 to n = 5 andthen

A4 sinh 500r� rth

G

� �� �5¼ _e eð125 kJmol�1K�1=RTÞ ½11�

The values of n = 3 and n = 5 are related to viscous-glide and climb controlled deformation, respectively,which are known to be rate-controlling mechanisms inthese alloys.[31–33] The microstructure revealed by theEBSD measurements shows mainly subgrain formation.No substructure development in some grains is found byEBSD, although subgrains with misorientations below 1deg are detected using TEM methods. The calculatedvalue of the activation energy for high temperaturedeformation is 125 kJ/mol, which is reasonably close tothe activation energy for self-diffusion in Al (143 kJ/mol) and to the activation energy for diffusion of Mgatoms in Al (135 kJ/mol), i.e., to the theoretic values forn = 5 and n = 3 regimes, respectively.

The rth values are summarized in Table I for all thetemperatures. An Arrhenius-type correlation as found inReference 34 could not be determined. The temperatureshows instead a good linear correlation with thethreshold stresses. The reduced stresses were defined asa function of the temperature in Reference 28, althoughin that case, the precipitation state did not change withthe temperature of deformation because the material

was overaged. The threshold values found in this workare much higher than that found in Reference 35 for thesame type of alloy in which AlMnFeSi particles are notpresent, and the drag effect is only imposed by theMg2Si particles. In the present work, a jump of thevalues between 673 K and 773 K (400 �C and 500 �C)shows the influence of the precipitation state; above773 K (500 �C), the stable b phase dissolves, as observedin Reference 28. The correlation for the low temperatureand the high temperature ranges was not subdividedbecause only 2 values for each range were determined.The mean subgrain size dss [lm] of recovered grains

measured by means of EBSD can be correlated well(R2 fitting value of 0.98) with the Zener–Hollomonparameter using Q = 125 kJ mol�1, as shown inFigures 10 and 11(a)), obtaining

ln dss ¼ 4:8� 0:122 lnZ ½12�

The mean subgrain size can be correlated (R2 = 0.98)to the reduced stress values and the following relation-ship can be written:

r=G ¼ 0:001þ 0:006=dss lm½ � ½13�

The subgrain size dependence in aluminum can becorrelated with the stress for many materials as[36]

dssb¼ k

rG

� ��r½14�

For the case of aluminum alloys, usually, k = 20 andr = 1,[34] and in the presence of particles, r is replacedby the effective stress re = r � rth.

[37] The measuredsubgrain size is good correlated with the effective stressaccording to Eqs. [14], [15], as shown in Figure 11(b).Additionally, the results obtained in this work agreewith those obtained after creep tests of the same type ofalloy in Reference 33.Finally, experimental evidence of very low misorien-

tation angles (in the range of 0.3 deg) and the dragging

Table I. Threshold Stress and Relaxation Factor j as a

Function of the Temperature

Temperature (�C) rth (MPa) j from Fig. 12

300 21.5 0.69400 20 0.69500 16 0.20550 15.5 0.18

Fig. 10—Zener–Hollomon parameter as a function of Eqs. [6], [8],and [9] showing a good correlation.

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 44A, MARCH 2013—1583

effect of the dislocation by particles demand a physicalinterpretation of the threshold stress.

A. Physical Interpretation of the Threshold Stress

The threshold stress introduced in Eqs. [10], [11], and[12] is interpreted in Reference 38 as the value belowwhich creep deformation is negligible mainly due to thepresence of fine particles which slow down the disloca-tions and boundaries’ movement. The adjustment of then1 exponent of the constitutive equation to physicalrelevant values by including a rth was presented in manyworks.[39,40] The threshold stress rth can be related to thea thermal detachment stress rd as proposed in Reference41 and interpreted as the stress necessary to detachdislocations from the particles.

rd ¼ r0

ffiffiffiffiffiffiffiffiffiffiffiffiffi1� j2p

½15�

j represents the relaxation factor and varies between 0for maximal attraction to 1 for no attachment, and r0

is the Orowan stress which can be calculated as

r0 ¼3Gb

2ðk� rÞ ½16�

with G as the shear modulus (30.22-16T GPa), b as theBurger’s vector (0.286 nm at room temperature andvarying with the temperature by thermal expansion asshown in Reference 42) and r as the radius of the par-ticle. The particle interspacing k can be calculatedusing Eq. [17]:

k ¼ 2r

ffiffiffiffiffiffip6fv

r½17�

where fv is the volume fraction of the particles.

In the studied case, the AlFeMnSi intermetallicphases represent around 3 vol. pct and their form andsizes are not modified in the studied temperature andstrain rate range. The stable Mg2Si phase observed inthe as-received material condition starts dissolving (seeendothermic peak in the DSC diagram) at around 773 K(500 �C). Due to high heating rates, dissolution of thisphase occurs during hot deformation. It means that atlower temperatures (in this case 573 K to 673 K (300 �Cto 400 �C)), the 1 vol. pct of Mg2Si particles canadditionally block the boundaries’ movement. Softeningobserved in the flow curves at high temperatures shouldbe related with the dissolution of this phase. It was alsoobserved that at high temperatures and low strain rate,pinching off of the original grain boundaries at largestrains combined with large subgrains produced somenew grains by geometric dynamic recrystallization. Thiseffect could be additionally related to the softening atthese deformation parameters.[13]

Additionally, and because we are dealing with twotypes of particles below 723 K (450 �C), the Eqs. [16],[17] were modified according to Reference 43. Theeffective mean radius rA replaces the radius r inEqs. [16], [17] and can be calculated as follows:

rA ¼ffiffiffi2

3

r Pi fv;ir

2iP

i fv;iri½18�

in which the number of particles in a unit of volume(nv,i) in the original equation was replaced by thefraction of particles (fv,i) in the present work.Thus, the dependence of the rd with the j value was

calculated for the cases of the low and high temperaturesrange, meaning with and without the influence of fineMg2Si particles, respectively, as shown in Figure 12. 65 nm(median radius) was used for 3 vol. pct of AlFeMnSi

Fig. 11—(a) Mean subgrain size as a function of the Zener–Hollomon parameter, (b) subgrain size: measured (open square), predicted byEqs. [14], [15] (circles), and measured[35] (open triangles) as a function of the reduced effective stress values.

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particles and 30 nm (ferret minimum) for the 1 vol. pct ofMg2Si needles.

If directly correlating the threshold factor for eachtemperature and precipitate state condition, the varia-tion of the relaxation factor j with the deformationparameters is obtained from Figure 12 and shown inTable I. From this, it is deduced that the interaction ofthe dislocations with particles is a function mainly of theprecipitation state, which governs the Orowan stresses.Although the detachment stresses decrease by increasingtemperature, a larger interaction between particles anddislocations occurs in the higher temperature rangeswith the remaining AlFeMnSi particles and due to thelower Orowan stresses.

The experimental evidence of boundaries and dislo-cations’ dragging in the TEM picture shown in Figure 8for 773 K (500 �C) is in agreement with the concepts ofdetachment stresses and relaxation j parameter.

V. CONCLUSIONS

The formability of a commercial age hardenablealuminum alloy with high content of fine dispersoidswas studied by means of hot compression tests, metal-lography after in situ water quenching and constitutiveequations.

High values of np1 exponent 11.75 for the creep lawand n1 = 7 for the universal one were related to astrengthening effect of the fine dispersoids. A thresholdstress as a function of the precipitation state wasintroduced in the constitutive equations, which resultsin a physically more correct n1 exponent value of 5 athigh stresses and of 3 in the lower stresses range. Thesevalues are related to dislocation climbing and slidingand thus to dynamic recovery.

Based on experimental evidence, the threshold stresswas explained with the concepts of the detachment

stress, related to the interaction of dislocations toparticles such as AlFeMnSi intermetallics in the wholerange of temperatures and additionally of Mg2Si pre-cipitates below 723 K (450 �C). Thus, although thethreshold stresses increased with decreasing tempera-ture, the relaxation factor showed maximal attractionbetween particles and dislocations in the 773 K to 823 K(500 �C to 550 �C) temperature range. This can beexplained by the lack of Mg2Si particles’ contribution tothe Orowan stresses in this temperature range. Thedetachment of dislocations was supported at 773 K(500 �C) by TEM investigations.

ACKNOWLEDGMENTS

The authors would like to thank AMAG for the provi-sion of the material and for carrying out the compressiontests at the RWTH in Aachen, and the Institute of Mate-rials Science and Technology and USTEM (both fromthe Vienna University of Technology) for the Gleebleand the FEG-SEM EBSD facilities, respectively. Thework was partially financed by the Austrian Agency ofResearch funds under the project number [P22238-N22]and by the Christian Doppler research society.

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Fig. 12—Detachment stress rd as a function of relaxation factor jfor the two cases: low temperature range with the presence of bothAlFeMnSi aluminides and fine Mg2Si precipitates (open symbols)and high temperature range with the presence only of aluminides(full symbols). Additionally, the threshold stresses, rs are plotted toobtain the corresponding j.

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