Contributions of phase and structural transformations in multicomponent Al-Mg alloys to the linear and nonlinear mechanisms of anelasticity

ISSN 0031�918X, The Physics of Metals and Metallography, 2014, Vol. 115, No. 2, pp. 192–201. © Pleiades Publishing, Ltd., 2014.Original Russian Text © I.S. Golovin, A.S. Bychkov, A.V. Mikhailovskaya, S.V. Dobatkin, 2014, published in Fizika Metallov i Metallovedenie, 2014, Vol. 115, No. 2, pp. 204–214.

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1. INTRODUCTION

It has been shown in [1] that the precipitates of theβ phase (Al3Mg2) in binary Al–Mg alloys substantiallysuppress the grain�boundary relaxation. The dissolu�tion of the β phase and, as a consequence, the enrich�ment of the solid solution by magnesium atoms andthe precipitation of the metastable β' phase upon thesubsequent heating decrease the dislocation mobilityin the solid solution. Changes in the content of theβ phase by changing the composition of the alloys andthe regimes of their heat treatment (e.g., the regime oftheir quenching) makes it possible to control the leveland the ranges of the action of these mechanisms ofanelasticity in the Al–Mg alloys. The most promisingway of improving the complex of the service and tech�nological characteristics of structural aluminum alloysis the formation of an ultrafine�grained structure insemi�finished products. A decrease in the grain sizeleads not only to an improvement in the room�tem�

perature mechanical properties, but also to theappearance of the effect of superplasticity in somealloys [2, 3].

An effective method of obtaining an ultrafine�grained (UFG) structure is an additional alloying ofthe binary alloys by transition metals and the use ofsevere plastic deformation (SPD) [4, 5]. The UFGstructure produced by SPD has high strength proper�ties and satisfactory plasticity. However, the UFGstructure after SPD is thermally unstable because of asignificant excessive energy stored upon plastic defor�mation, which is mainly concentrated in the nonequi�librium grain boundaries of large extension. Becauseof the nonequilibrium state of the boundaries with anenhanced fractality, the grain�boundary relaxation,which usually occurs via elastic deformation on flatregions of the grain boundaries, in the SPD structuresproves often to be suppressed significantly or even isabsent at all [6]. Upon the heating of aluminum alloys,

Contributions of Phase and Structural Transformations in Multicomponent Al–Mg Alloys to the Linear

and Nonlinear Mechanisms of AnelasticityI. S. Golovina, A. S. Bychkova, A. V. Mikhailovskayaa, and S. V. Dobatkina, b

aNational University of Science and Technology MISiS, Leninskii pr. 4, Moscow, 119049 RussiabBaikov Institute of Metallurgy and Materials Science, Russian Academy of Sciences,

Leninskii pr. 49, Moscow, 119991 Russiae�mail: [email protected]

Received April 23, 2013; in final form, June 11, 2013

Abstract—The effects of the processes of severe plastic deformation (SPD), recrystallization, and precipita�tion of the β phase in multicomponent alloys of the Al–5Mg–Mn–Cr and Al–(4–5%)Mg–Mn–Zn–Sc sys�tems on the mechanisms of grain�boundary relaxation and dislocation�induced microplasticity have been stud�ied in some detail. To stabilize the ultrafine�grained structure and prevent grain growth, dispersed Al–transi�tion�metal particles, such as Al3Zr, Al6Mn, Al7Cr, Al6(Mn,Cr), Al18Cr2Mg3 have been used. We have specialinterest in alloys with additions of scandium, which forms compounds of the Al3Sc type and favors the pre�cipitation of finer particles compared to the aluminides of other transition metals. After SPD, Al–(4–5%)Mg–Mn–Zr–Sc alloys exhibit an enhanced recrystallization temperature. The general features of the dislocationand grain�boundary anelasticity that have been established for the binary Al–Mg alloys are retained; i.e.,(1) the decrease in the dislocation density in the process of recrystallization of cold�worked alloys leads to theformation of a pseudo�peak in the curves of the temperature dependences of internal friction (TDIF) and toa decrease in the critical amplitude of deformation corresponding to the onset of dislocation motion in a stressfield; (2) the precipitation of the β phase suppresses the grain�boundary relaxation; (3) the dissolution of theβ phase, the passage of the magnesium atoms into the solid solution, and the precipitation of the β' phaseupon heating hinder the motion of dislocations; (4) the coarsening of the highly dispersed particles contain�ing Zr and Sc increases the dislocation mobility. The grain�boundary relaxation and dislocation–impurityinteraction and their temperature dependences, as well as processes of the additional alloying of the binaryalloys by Mn, Cr, Zr, and Sc, have been estimated quantitatively.

Keywords: Al–Mg�based multicomponent alloys, relaxation and hysteresis mechanisms of anelasticity

DOI: 10.1134/S0031918X14020082

STRENGTH AND PLASTICITY

THE PHYSICS OF METALS AND METALLOGRAPHY Vol. 115 No. 2 2014

CONTRIBUTIONS OF PHASE AND STRUCTURAL TRANSFORMATIONS 193

the grains grow rapidly while the strength propertiesdecrease. After annealing, which makes the grainboundaries more equilibrium, there again can developa process of the elastic relaxation of stresses, which isregistered as a thermally activated relaxation peak inthe TDIF curves [7].

To stabilize the UFG structure and prevent graingrowth, disperse aluminum–transition�metal parti�cles, such as Al3Zr, Al6Mn, Al7Cr, Al6(Mn,Cr) [8, 9]and Al18Cr2Mg3 [10, 11], are used, which are ther�mally stable and have a high density in the unit volumeof the aluminum matrix. From this point of view, theuse of scandium additions, which forms an Al3Sc com�pound with aluminum, is an effective method; theirparticles precipitate from the solid solution in a moredisperse form as compared to those of the aluminidesof other transition metals, such as Mn, Cr, Zr [12–14].However, the stability of Al3Sc is not as high as the alu�minides of zirconium, manganese, or chromium and,upon heating, they coarsen rapidly. To enhance thethermal stability, scandium is introduced into the alu�minum alloys together with zirconium, which is dis�solved in the Al3Sc phase. The formation of anultrafine�grained structure is also necessary to ensurethe effect of superplasticity, which is achieved uponthe fabrication of articles by the method of superplas�tic formation, which makes it possible to obtain hol�low articles with an intricate shape. To obtain super�plastic magnalium alloys, it is reasonable to use bothzirconium and scandium [15–17] and additives ofchromium and manganese simultaneously [18–21],which makes it possible to produce a stable UFGstructure necessary for subsequent superplasticitydeformation.

This work was aimed at studying the effects ofalloying the binary Al–(4–5%)Mg alloys by Mn, Cr,Zr, and Sc, as well as the effects of SPD on the dislo�cation and grain�boundary anelasticity in multicom�ponent alloys in wide ranges of temperatures, frequen�cies, and amplitudes of vibrations.

2. EXPERIMENTAL

Alloys of the Al–5Mg–Mn–Cr, Al–4Mg–Mn–Zr–Sc, and Al–5Mg–Mn–Zr–Sc systems (see table)have been analyzed. The methods of the investigationof the amplitude and temperature dependences ofinternal friction (ADIF and TDIF, respectively) usinga DMA Q800 (TA Instruments) dynamic analyzer, thetechnological features of the preparation of multicom�ponent alloys and samples from these alloys, as well asthe methods of analyzing their structure, have beendescribed in [7, 22, 23]. Alloys that contain zirconiumand scandium were subjected to SPD by the method ofequal�channel angular pressing (ECAP) at a tempera�ture of 300°С using samples with dimensions of 10 ×10 × 70 mm3 with the angle of the channel intersectionequal to 90° with different numbers of passages (N)through the ВС route (after six passages, the true defor�mation is ~6.8). The methods of preparation, as well asthe structure and properties of these alloys, have beendescribed in [24, 25].

The microstructure was studied using an Axiovert200M Mat optical microscope (OM). The polishedsections were prepared by mechanical lapping and pol�ishing; then, electrolytic polishing in a chloroalcoholelectrolyte at a voltage of 15–20 V for 7–8 s was used. Toreveal the grain structure, the samples were subjected toanodic oxidation in a 10% aqueous solution of thehydroborofluoric acid. The thin foils were prepared byelectropolishing disks in a chloroalcohol electrolyte usinga Struers Tenupol unit and were studied using a JEOL2000�CX transmission electron microscope (TEM).

3. RESULTS AND DISCUSSION

The effects of the β phase on the grain�boundaryand dislocation anelasticity that were revealed inbinary Al–Mg alloys (to 12% Mg) make it possible toanalyze (according to a similar scheme) the amplitudeand temperature dependences of internal friction inmore complex multicomponent alloys. The additionalalloying by transition metals was chosen to obtain the

Compositions of the alloys

Alloys Mg Mn Cr Zr Sc

Al–4Mg 4

Al–5Mg 5

Al–5Mg–Mn 5 0.6

Al–5Mg–Mn–Cr 5 0.6 0.25

Al–4Mg–Mn–Zr 4 1.5 0.4

Al–4Mg–Mn–Zr–Sc 4 1.5 0.4 0.4

Al–5Mg–Mn–Zr–Sc 4.6 0.6 0.1 0.2

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maximum amount of the second�phase particles forthe maximum stabilization of the UFG structure,including that obtained after ECAP.

3.1. Multicomponent Al–Mg–Mn–Cr Alloys

The additional alloying of the Al–Mg alloys bymanganese and chromium leads to the formation of afairly fine, stable grain structure after deformation andsubsequent recrystallization (Figs. 1a, 1b) due to theformation of insoluble disperse particles of theAl6(Mn,Cr) phase, which hampers grain growth up tosubsolidus temperatures employed for the superplasticdeformation of the alloy [21]. After 30�min annealingat 550°С, the average grain size in the alloy without

chromium was 13 ± 3 μm and, in the alloy with 0.25%Cr, it was 9 ± 1 μm. In the alloy with manganese butwithout chromium, the size of Al6Mn particles in thecold�worked state was about 100 nm; chromium leadsto a decrease in the size of the Al6(Mn,Cr) particles to≈50 nm (estimated without the allowance for the foilthickness, Figs. 1c, 1d). Alloying substantiallyimproves the superplasticity characteristics [19, 26].The TDIF curves of these alloys have been studied in[19, 22].

Figure 2 displays the ADIF curves of samplesobtained after quenching from 550°С via the proce�dure described in [1] and the results of their graphicalanalysis and the values of the critical parameters. Theinflection in the dependences of the slope of the ADIF

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Fig. 1. (a, b) Initial grain structure (optical microscopy (OM)) and (c, d) disperse particles of intermetallic compounds of tran�sition metals (TEM) in (a, c) Al–Mg–Mn and (b, d) Al–Mg–Mn–Cr alloys after annealing at 550°С for 30 min; and in(e) Al–Mg–Mn–Zr–Sc and (f) Al–Mg–Mn–Zr alloys after annealing at 350°С for 2 h and 500°С for 6 h, respectively (OM).



curves (tanα), which characterizes the transition fromthe vibrations of dislocations between the pinningpoints to their motion over the crystal, and in thedependences of the critical stress (εcr) required for theonset of dislocation motion through a system of pointobstacles on the reciprocal temperature (ln(tanα) andln(εcr) vs. 1/T, respectively) in these alloys is observed,just as in the binary Al–5%Mg alloy, at a temperatureclose to the temperature of the dissolution of theβ phase. The values of the critical amplitude in multi�component alloys are higher (the corresponding val�ues of εcr for the binary alloy Al–5Mg are drawn inFig. 2 by a dashed line for comparison). Judging fromthe experimental results, the temperature dependenceof the parameters of the ADIF and the effective forceof friction of dislocations upon their motion in thesubstitutional solid solution of the Al–5Mg–Mn–Cralloys did not change qualitatively compared to thebinary alloy with a similar content of manganese. Thissuggests that disperse particles, such as Al6(Mn,Cr),can efficiently affect the dislocation mobility in theprocess of plastic deformation [27], whereas in theelastic region of loading the mobility of dislocations is

more efficiently affected both by the alloying of thesolid solution and by the precipitates of the metastableβ' phase.

3.2.Multicomponent Al–Mg–Mn–Zr–Sc Alloys

These alloys were investigated in two states: (1)annealed after casting and (2) after ECAP. Theannealing of the alloys was performed via two regimes:at 500°С for 6 h for the Al–4Mg–Mn–Zr and at350°С for 2 h for the Al–4Mg–Mn–Zr–Sc alloy. Thetemperature and time parameters of the annealingswere chosen so as to obtain products of the decompo�sition of the solid solution of the maximum density,i.e., with the minimum interparticle spacings [28].After annealing, the alloys had a nondendritic struc�ture with grain size of ~20 μm in the Al–4Mg–Mn–Zr–Sc alloy and ~30 μm in the Al–4Mg–Mn–Zralloy (Figs. 1e, 1f).

In the TDIF curves for these alloys in the annealedstate, just as in binary alloys with 4% Mg, a peak of IFis observed that is caused by grain�boundary relaxation(Figs. 3a, 3b). Just as in binary alloys [1], the presence

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Fig. 2. (a, b) ADIF curves for annealed at 550°C and quenched alloys (a) Al–5Mg–0.6Mn and (b) Al–5Mg–0.6Mn–0.25Cr.(c, d) Dependences of ln(tanα) and ln(εcr2) and of the equilibrium content of the β phase on the reciprocal temperature of mea�surements of the ADIF (1/T (K–1)) for the alloys (c) Al–5Mg–0.6Mn and (d) Al–5Mg–0.6Mn–0.25Cr at various temperaturesin the range of 30– 550°C and frequency f = 3 Hz. Dashed line indicates analogous dependences for the binary alloy Al–5Mgaccording to the data of [1] for comparison. Vertical dashed line shows solvus temperature.

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of the β phase substantially suppresses the grain�boundary anelasticity. The additional alloying ele�ments in the Al–4Mg alloy decrease the degree ofgrain�boundary relaxation as compared to the binaryalloy.

The ADIF curves of the Al–4Mg–Mn–Zr andAl⎯4Mg–Mn–Zr–Sc alloys measured at differenttemperatures are given in Figs. 4a and 4b. An analysisof the temperature dependence of the main parame�ters of the ADIF (tanα and εcr) is presented in Fig. 4c.The performed analysis reveals a substantial differencein the mobilities of dislocations in these alloys com�pared to their prototype—the binary alloy with4% Mg—at a temperature above the solvus tempera�ture of the β phase (the data for the binary alloy areshown in Fig. 4c by dashed lines). The temperaturedependence of the values of tanα, which reflects themagnitude of the energy dissipated upon the vibrationsof the dislocation segments in the field of applied peri�odic stresses, behaves qualitatively similar to that for

the binary Al–4Mg alloy: upon the transition from thetwo�phase (α + β) into the single�phase (α) region,the values of tanα begin increasing rapidly withincreasing temperature. However, the rate of thisgrowth of tanα in multicomponent alloys is less thanthat in the binary alloy, since the highly disperse parti�cles of the Al3(Sc,Zr) and Al6Mn phases additionallyhamper the elastic bowing of dislocations in the fieldof applied periodic vibrations. The values of the criti�cal amplitude εcr at which the dislocation breaks awayfor the pinning points and begins moving over the crys�tal also begin decreasing at the temperatures above thetemperature of dissolution of the β phase. The rate ofdecrease in the critical amplitude εcr with increasingtemperature is substantially lower than that of thebinary alloy.

The electron�microscopic analysis of the annealedAl–Mg–Mn–Zr–Sc after ECAP revealed the forma�tion of a mainly submicrocrystalline structure(Fig. 5a). The average size of grains was 860 ± 60 nm.

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The presence of high�angle grain boundaries followsfrom the character of the electron�diffraction patternsand also from the presence of a fringe contrast at grainboundaries (Fig. 5a). A subgrain structure with low�angle misorientations of the elements was alsoobserved, the dimensions of which correspond to thesizes of grains with high�angle boundaries. Particles ofAl3(Sc,Zr) and Al6Mn with an average size of 5–25 and10–60 nm, respectively, were observed in the struc�ture. Coarse particles of intermetallic compounds ofrounded shape up to 150 nm in size and particles ofplatelet shape (50–150) × (150–350) nm in size werealso present, which were formed upon annealing.

In the Al–Mg–Mn–Zr alloy after ECAP of anannealed sample there is formed a structure similar tothat of the Al–Mg–Mn–Zr–Sc alloy, but with agreater size of grains, which reaches 1240 ± 70 nm(Fig. 5b). Al3Zr and Al6Mn particles are also present.The Al3Zr particles are distributed nonuniformly overthe volume of the sample and are encountered in theform of agglomerates. The size of the Al3Zr particles isfrom 15 to 65 nm; that of the Al6Mn particles, from 10to 60 nm. Just as in the case of the Al–Mg–Mn–Zr–Sc alloy after annealing and ECAP, in the Al–Mg–Mn–Zr alloy there were observed coarser particles of inter�metallic compounds of rounded shape (70–180 nm insize), of square shape (up to 250 nm in size), and ofplatelet shape ((70–214) × (210–890) nm in size).

After ECAP of the Al–4Mg–Mn–Zr andAl⎯4Mg–Mn–Zr–Sc alloys, upon continuous heat�ing, the recrystallization temperature Тr proves to besignificantly higher than the temperatures at which thegrain�boundary relaxation in the annealed state cantake place (Figs. 3c, 3d). After ECAP with N = 6, therecrystallization pseudo�peak of IF in the Al–Mg–Mn–Zr alloy lies at a temperature of ~470°С (seeFig. 3c) and, in the Al–Mg–Mn–Zr–Sc alloy, it liesat a temperature of about 570°С (Fig. 3d) upon con�tinuous heating at a rate of 1 K/min. The mechanism ofthe formation of this IF pseudopeak has been describedin [1, see Fig. 5]. Near 240°С in the TDIF curves ofSPD samples, there is an inflection (Figs. 3c, 3d, upperinset). The position of this inflection is independent ofthe vibration frequency and, in fact, coincides with thetemperature of the dissolution of the β phase; conse�quently, this effect is due to the phase transformationrather than to the thermally activated relaxation pro�cess. After ECAP, no grain�boundary IF peak isobserved in the samples. This is because of the non�equilibrium structure of grain boundaries after SPD,which suppresses the possibility of grain�boundarysliding [7, 29]. The grain�boundary peak is onlyformed in the TDIF curve in the course of repeatedmeasurements of the ECAP sample heated during thefirst measurement to a temperature of 570°С, i.e.,when the sample has been partially recrystallized(Figs. 3c, 3d, see the lower inset). In the alloy withoutscandium, the grain�boundary peak is more clearlypronounced, since the heating of a SPD sample to

570°С leads to a more complete occurrence of therecrystallization processes (the temperature of theonset of recrystallization in this alloy is ~470°C,whereas in the alloy with scandium, the maximumtemperature upon the measurements of the TDIF(570°С) is comparable with the temperature of theonset of recrystallization upon continuous heating.The energy of activation of the grain�boundary peak inthe alloy without scandium is Н ≈ 1.9 eV and the fre�quency factor in the Arrhenius equation is τ0 ≈ 2 × 1020 s.

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These parameters are close to the parameters of thegrain�boundary peak in this alloy before ECAP.

A similar effect on the TDIF curves is exerted bythe process of the dissolution of the β phase in thealloys of this system with a greater content of magne�sium, e.g., Al–4.6Mg–Mn–Zr–Sc, after ECAP withdifferent numbers of passages and also after heating ofthe ECAPped samples to 400°С (Fig. 6). The insets ofthese figures show TDIF curves taken upon coolingare given; here, just as in the case of binary Al–Mgalloys, a well pronounced grain�boundary peak of IFcan be seen. Upon cooling in the process of measure�ments of the TDIF at a rate of 1 K/min, the β phasehas no time to precipitate or to block grain�boundarysliding.

It has been shown in [1, 30] that the isothermalholding of deformed binary Al–Mg alloys at a temper�ature of about 260°С leads to the noticeable evolutionof the ADIF curves due to the processes of recrystalli�zation. Upon recrystallization, a sharp decreaseoccurs in the density of dislocations and an increase in

their mobility. As a consequence, the level of the linearstage in the ADIF curves decreases in the course of theisothermal holding and a well�pronounced inflectionis formed that is caused by a decrease in the stressrequired for the onset of microplastic deformation andthe irreversible displacement of dislocations.

In multicomponent Al–4Mg–Mn–Zr alloys, afterECAP, the character of the ADIF curves upon the iso�thermal holding differs substantially from the ADIFcurves in the binary alloy with the same content ofmagnesium (4%). At a temperature of 260°С, no sig�nificant changes in the shape of the ADIF curves wererevealed; the shape of the ADIF curves upon isother�mal holding at a temperature of 390°С (190 min) isshown in Fig. 7a. After ECAP with N = 6, no decreasein the background of IF is observed in the Al–4Mg–Mn–Zr alloys and no visible displacement of the crit�ical deformations to the region of smaller valuesoccurs; i.e., no characteristic signs of the processes ofrecrystallization and related decrease in the disloca�tion density can be seen. The ADIF curves in the

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Fig. 5. Structure of alloys (a,d) Al–Mg–Mn–Zr–Sc and (b,c) Al–Mg–Mn–Zr: (a, b) after ECAP with N = 6 and subsequentannealing (c) at 390°С for 190 min and (d) at 400°С for 110 min.



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entire range of measurements (Fig. 7a) have an almostlinear character. The low�amplitude region (to theamplitude ε ≈ 2 × 10–4 or σ ≈ 15 MPa) remains almostunaltered during holding at 390°С. The high�ampli�tude portion (ε from 2 × 10–4 to 6.5 × 10–4) is alsoapproximated well by a linear dependence, the slope ofwhich decreases with increasing duration of annealingat 390°С. The change in the slope of this portion of theADIF curve (ln(tanα)) is shown in the inset to Fig. 7a.

Similar isothermal measurements of the ADIF at400°С for the SPD multicomponent samples withscandium and with different contents of magnesium(Al–4Mg–Mn–Zr–Sc (Fig. 7b) and Al–4.6Mg–Mn–Zr–Sc (Fig. 7c)) also differ radically from theADIF curves for the binary Al–4Mg and Al–5Mgalloys [1, Figs. 7b, 7c]; in particular, they do not showcharacteristic changes caused by recrystallization. Thebehavior of the ADIF in SPD alloys with scandiumsomewhat differ also from that of the sample withoutscandium. In the samples of the Al–Mg–Mn–Zr–Scalloys, the high�amplitude region of the ADIF curvesremains almost unaltered in the course of the isother�mal holding, and the low�amplitude portion with aweak amplitude dependence of the internal frictionincreases somewhat with increasing duration ofannealing.

Some increase in IF appears to be due to theincreased mobility of dislocation segments upon thecoarsening of fine scandium–zirconium phases(Figs. 5c, 5d), whereas the total density of dislocationsin the material does not change substantially at thistemperature.

According to TEM data, the size of structural ele�ments (grains/subgrains) in the Al–4Mg–Mn–Zralloy after ECAP and annealing at 390°C for 190 minincreases to 1–3 μm (some grains of up to 7 μm areencountered). In the bulk of the grains/subgrains,there have been revealed fine particles of roundedshape somewhat grown upon heating (to 40–110 nm)and particles of platelet shape with an average size ofabout 40 × 150 nm, as well as relatively large particles250–400 nm in size. It is difficult to say from whichparticles obtained by ECAP these coarse particle wereformed, i.e., from fine (less than 65 nm) or coarse (upto 180 nm in size) particles. At the grain boundaries,particles with sizes of up to 400–650 nm wereobserved, which appear to be a result of acceleratedgrain�boundary diffusion. In the Al–Mg–Mn–Zr–Scalloy, analogous structures were observed after ECAPand heating at 400°C for 107 min, but the dimensionsof the structural elements and particles were somewhatsmaller because of the alloying by scandium; the sizeof grains/subgrains was 0.9–2.0 μm (some grainsreached 5 μm), and the size of fine particles in the bulkof grains was 110–330 nm, whereas the size of coarseparticles at grain boundaries was 180–430 nm.

Thus, the temperature and time effects in theamplitude dependences of IF in multicomponentAl⎯Mg–Mn–Zr–Sc alloys subjected to ECAP are

substantially smaller even at 390–400°C than theeffects at 260°С in binary Al–Mg alloys after rolling,and they are due to the coarsening of particles con�taining zirconium and scandium.

4. CONCLUSIONS

Amplitude and temperature dependences of inter�nal friction have been studied in multicomponentAl⎯(4–5)Mg–Mn–Zr–Sc and Al–5Mg–Mn–Cralloys. It has been shown that the main regularities ofthe temperature and amplitude dependences of inter�nal friction of binary Al–(4–5)Mg alloys caused by theformation and dissolution of the β phase are inheritedby the multicomponent alloys of these systems. Anadditional alloying of the Al–Mg alloys by Mn, Cr, Zr,and Sc affects the quantitative parameters of the grain�boundary relaxation and dislocation mobility, as wellas the temperature intervals of the manifestation ofanelasticity. The severe plastic deformation of Al–Mg–Mn–Zr–(Sc) alloys substantially decreases thedislocation mobility, almost completely suppresses thegrain�boundary internal friction, and significantlyincreases the temperature stability of the structureupon heating.

The formation of particles of the Al6(Mn,Cr) typewith an average size of about 50 nm does not affect thedislocation mobility in the elastic region of loading ofthe Al–Mg–Mn–Zr–Sc alloys; a more efficient pro�cess is the pinning of dislocations by magnesium atomspresent in the solid solution. The change in the shapeof the amplitude dependence of internal friction inAl–Mg–Mn–Zr–(Sc) alloys at 390–400°С isrelated to a coarsening of particles of the Al3(Zr,Sc)type and to a weakening of pinning of dislocations bythese particles.

ACKNOWLEDGMENTS

We are grateful to R.O. Kaibyshev, who kindly sup�plied us by samples of the Al–5Mg–Mn–Zr–Sc alloy.

This study was supported by the Ministry of Educa�tion and Science, NUST “MISIS” and President ofthe Russian Federation Grant no. 14.125.13.232�MK.

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Translated by S. Gorin

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Contributions of phase and structural transformations in multicomponent Al-Mg alloys to the linear and nonlinear mechanisms of anelasticity