Journal of Alloys and Compounds 570 (2013) 125–132

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Journal of Alloys and Compounds

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The experimental investigation of phase equilibria in the Al-rich cornerwithin the ternary Al–Mn–Be system

0925-8388/$ - see front matter � 2013 Elsevier B.V. All rights reserved.http://dx.doi.org/10.1016/j.jallcom.2013.03.048

⇑ Corresponding author. Tel.: +386 2 220 7863; fax: +386 2 220 7990.E-mail addresses: franc.zupanic@um.si, franc.zupanic@uni-mb.si (F. Zupanic).

Franc Zupanic a,⇑, Boštjan Markoli b, Iztok Naglic b, Tonica Boncina a

a University of Maribor, Faculty of Mechanical Engineering, Smetanova ulica 17, SI-2000 Maribor, Sloveniab University of Ljubljana, Faculty of Natural Sciences and Technologies, Askerceva 12, Ljubljana SI-1000, Slovenia

a r t i c l e i n f o a b s t r a c t

Article history:Received 29 December 2012Received in revised form 7 March 2013Accepted 8 March 2013Available online 30 March 2013

Keywords:Phase diagramHorizontal cross-sectionIsothermal annealingAluminiumManganeseBeryllium

This work investigated the constitution of the Al-rich corner within the ternary Al–Mn–Be phase diagramusing SEM + EDS, AES, XRD and DSC. With respect to the results, an isothermal cross-section at 600 �Cwas established, as well as a prediction of the apparent liquidus projection in the Al-corner. Be4AlMnis a thermodynamically stable phase in the Al-rich corner of the ternary phase diagram. The other ternaryT-phase, usually designated as Al15Mn3Be2, formed during solidification in alloys with Be:Mn atomicratios of less than 4:1, and having more than 1.5 at.% Mn. This phase is not a stable phase in the Al-richcorner at 600 �C. In contrast, the k-Al4Mn phase is a stable one. The T-phase is stable over a rather largepart of the phase diagram at least within a temperature range close to 750 �C, where it is in equilibriumwith the Al-rich liquid phase, and Be4AlMn.

� 2013 Elsevier B.V. All rights reserved.

1. Introduction

Over recent years, interest in ternary Al–Mn–Be alloys has at-tracted much attention because Be additions considerably increasethe quasicrystalline forming ability of Al–Mn alloys [1,2]. As a re-sult, several new casting alloys have been developed having veryinteresting properties [3], having the potential for applicationswhen requiring combinations of high strength, low density andductility. Investigations into these alloys in as-cast state havepointed out certain discrepancies with the established Al–Mn–Bephase diagram. For example, Be4AlMn was present in alloys inthe Al-corner [4], raising the question as to whether this phase isstable or not.

A short overview concerning the current knowledge will be gi-ven regarding the Al–Mn–Be phase diagram, and boundary Al–Beand Al–Mn.

The Al–Be phase diagram is a simple eutectic one [5,6]. The eu-tectic reaction

L ð2:4 at:%BeÞ ! a-Al ð0:19 at:% BeÞ þ a-Be ð0 at:% AlÞ ð1Þ

takes place at 644 �C. L denotes the liquid phase, whereas a-Al anda-Be denote Al-rich and Be-rich solid solution, respectively. The sol-ubility of Al in Be is negligible.

The Al–Mn phase diagram is much more complicated [7,8]. Therelevant phases occurring in the Al-corner are: a-Al, Al6Mn, l-Al4.12-

Mn and k-Al4Mn. The more significant invariant reactions are:

L ð14:2 at:%MnÞ þ l-Al4:12Mn! k-Al4Mn; at � 722 �C ð2Þ

L ð2:4 at:%MnÞ þ k-Al4Mn ð19 at:% MnÞ! Al6Mnð14:2 at:%MnÞ; 700 �C ð3Þ

L ð1:0 at:% MnÞ ! a-Al ð0:62 at:% MnÞ þ Al6Mn; 657 �C ð4Þ

There are some discrepancies regarding these equations in theliterature. Grushko and Balanetskyy [9] argued that k-Al4Mn formsvia a peritectiod (l-Al4.12Mn + Al6Mn ? k-Al4Mn) and not via peri-tectic reaction Eq. (2).

The only systematic investigation of the ternary Al–Mn–Be sys-tem was carried out by Raynor et al. [10]. Their results were usedfor compilations of phase diagrams by Mondolfo and Stiltz [11,12].Raynor et al. [10] discovered a ternary phase Al15Mn3Be2, whichStiltz [12] designated as a T-phase, whereas Kim et al. [1] denotedit as phase H1. In this presented work, this phase will be symbol-ized as T-phase. Carrabine [13] reported the existence of a ternarycompound Be4AlMn. Mondolfo [11] expressed doubt aboutBe4AlMn being the equilibrium phase in the Al-rich corner. Re-cently, Boncina et al. [14] discovered this phase in a slowly cooledAl86Mn3Be11-alloy.

The following invariant reactions were predicted by Mondolfo[11] with respect to the metallographic results of Raynor et al. [10]:

Table 1Chemical compositions of the investigated alloys as determined using inductively-coupled plasma - atomic emission spectroscopy.

Sample Unit Al Mn Be

MOB1 Al98.4Mn0.7Be0.9 mas.% 98.2 1.49 0.31at.% 98.35 0.73 0.92

MOB2 Al97.9Mn1.2Be0.9 mas.% 97.18 2.52 0.30at.% 97.85 1.24 0.90

MOB3 Al97.3Mn1.9Be0.8 mas.% 95.99 3.75 0.26at.% 97.35 1.86 0.79

MOB4 Al98.1Mn1.2Be0.7 mas.% 97.25 2.52 0.23at.% 98.06 1.24 0.69

MOB5 Al97.0Mn1.8Be1.2 mas.% 95.99 3.61 0.40at.% 97.00 1.78 1.20

MOB6 Al79Mn17Be4 mas.% 68.71 30.07 1.22at.% 78.89 16.92 4.19

MOB8 Al73Mn14Be13 mas.% 68.47 27.48 4.05at.% 72.81 14.32 12.87

MOB9 Al66Mn4Be30 mas.% 77.93 10.22 11.85at.% 65.85 4.23 29.92

MOB10 Al96.4Mn0.6Be3.0 mas.% 97.71 1.30 0.99at.% 96.45 0.57 2.98

Fig. 1. Backscattered electron micrographs of the alloys: (a) Al98.1Mn1.2Be0.7 and (b)Al97.3Mn1.9Be0.8.

126 F. Zupanic et al. / Journal of Alloys and Compounds 570 (2013) 125–132

Lþ l-Al4:12Mn! Al6Mnþ T ð5Þ

Lþ a-Be! a-Alþ T ð6Þ

Fig. 2. DSC-traces for (a) heating and (b) cooling of alloys Al98.1Mn1.2Be0.7 andAl97.3Mn1.9Be0.8.

L! a-Alþ Al6Mnþ T ð7Þ

Mondolfo [11] also foresaw the liquidus projection in the Al-corner with primary solidification fields for a-Al, a-Be, Al6Mnand T. The partial isothermal section at 600 �C predicted theexistence of two three-phase fields: a-Al + a-Be + T anda-Al + Al6Mn + T.

The aims of the current investigation were

� to verify whether Be4AlMn is a stable phase in the Al-corner ornot,� to check the stability range of T-phase,� to determine apparent liquidus projection and� the horizontal cross-sections at 600 �C and 750 �C in the Al-

corner,� and to predict the more probable invariant and univariant reac-

tions taking place in the Al-corner.

2. Experimental

The alloys were prepared from pure aluminium (99.997% Al), manganese (99.9%Mn) and beryllium (99.97% Be) using an arc-melting furnace (Edmund BühlerGmbH, Compact Arc Melter MAM-1) under an argon atmosphere. Chemical compo-sitions of the investigated alloys were determined using inductively-coupled plas-ma - atomic emission spectroscopy (Table 1).

Arc melted alloys were heat-treated in quartz ampullae filled with argon at600 �C and 750 �C for 24 h, 168 h (1 week) and 720 h (1 month). The samples thatannealed at 750 �C were additionally placed within an alumina crucible becausethey had partly melted. Differential scanning calorimetry (DSC) measurements ofarc-melted alloys were performed using Netzsch STA 449C Jupiter in an argonatmosphere at heating and cooling rates of 10 �C/min.

The samples were investigated using light microscopy (Nikon, Epiphot 300),scanning electron microscopy (FEI, SIRION NC), energy dispersive spectroscopyEDS (Inca 350, Oxford Instruments), Auger electron spectroscopy AES (PHI 680 Au-ger Nanoprobe) and X-ray diffraction (synchrotron radiation, wavelength 0.1 nmcarried out at Elettra, Sincrotrone Trieste, Italy).

3. Results and discussion

3.1. Solidification microstructures

Investigations using XRD, SEM + EDS, AES, and a method fordetermining electron backscattering coefficients [15] revealed that

Table 2Phase composition of selected alloys in the as-cast and heat-treated conditions.

Alloy As-cast 600 �C 750 �C

Al97.3Mn1.9Be0.8 a-Al, Be4AlMn, Al6MnT

k-Al4Mn, a-Al, Be4AlMn (considerable change in microstructure) Above the liquidus temperature

Al79Mn17Be4 a-Al, k-Al4Mn,Al11Mn4

a-Al, Al6Mn, k-Al4Mn, Al11Mn4 (moderate change inmicrostructure)

a-Al, l-Al4Mn (considerable change inmicrostructure)

Al73Mn14Be13 a-Al, T,Be4AlMn,Al11Mn4

a-Al, T,Be4AlMn, Al11Mn4 (no change in microstructure, diffusiontoo low)

a-Al, T,Be4AlMn (considerable change inmicrostructure)

Al66Mn4Be30 a-Al, a-Be, Be4AlMn a-Al, a-Be, Be4AlMn (considerable change in microstructure) a-Al, a-Be, Be4AlMn (considerable change inmicrostructure)

Fig. 3. Backscattered electron micrographs of alloys: (a) Al96.4Mn0.6Be3.0 and (b)Al66Mn4Be30.

Fig. 4. A DSC trace for the alloy Al96.4Mn0.6Be3.0.

Fig. 5. Backscattered electron micrographs of alloys after DSC: (a) Al79Mn17Be4 and(b) Al73Mn14Be13.

F. Zupanic et al. / Journal of Alloys and Compounds 570 (2013) 125–132 127

alloys Al98.4Mn0.7Be0.9, Al97.9Mn1.2Be0.9 and Al98.1Mn1.2Be0.7 con-sisted of three phases a-Al, Al6Mn and Be4AlMn. These alloys con-tained up to 2 at.% Be + Mn, and the atomic ratios Be:Mn were lessthan 4:1. Fig. 1 shows a typical microstructure of these alloys. Theycontained dendritic a-Al and two types of heterogeneous struc-tures. The two-phase (a-Al + Al6Mn) structure was akin to the (a-

Al + Al6Mn) eutectic occurring in the boundary Al–Mn phase dia-gram. The other heterogeneous structure comprised of a-Al andBe4AlMn. DSC (Fig. 2) indicated solidus temperature of around650 �C (650 ± 2 �C). Some undercooling was observed prior to crys-tallization of a-Al with a large exothermic peak appearing at648 �C. Closer examination of this peak (Fig. 2b) revealed two devi-ations, indicating processes leading to the formations of the twoheterogeneous structures (a-Al + Al6Mn) and (a-Al + Be4AlMn).

Metallographic analysis showed that the alloys with a slightlyhigher Be + Mn content (Be + Mn � 3 at.%) and Mn content higherthan 1.5 at.%, contained particles of the T-phase in addition tophases encountered in the alloys with up to 2 at.% Be + Mn(Fig. 1b, Table 2).

The solidus temperatures of these alloys were apparently thesame as those in alloys with Be + Mn � 2 at.% (Fig. 2). Upon cooling,an additional peak appeared before the largest exothermic peak,indicating the formation of T-phase. In the microstructure,(a-Al + Al6Mn) and (a-Al + Be4AlMn) heterogeneous structureswere also present.

The alloy with the approximate Be:Mn ratio 5:1(Al96.4Mn0.6Be3.0) possessed a two-phase microstructure consistingof a-Al and a-Be (Fig. 3a), and did not contain any other Mn-richphase. This Mn-content corresponded approximately to the

Fig. 6. DSC-traces for (a) heating and (b) cooling of alloys Al79Mn17Be4 andAl73Mn14Be13.

Fig. 7. XRD-traces of the alloy Al97.3Mn1.9Be0.8 under different conditions. Bottomcurve: as-cast condition; middle curve: 600 �C, 24 h; upper curve: 600 �C, 720 h.

Fig. 8. Microstructures of the alloy Al97.3Mn1.9Be0.8 after annealing at 600 �C: (a) for24 h (light-optical micrographs) and (b) for 720 h (backscattered electron image).

128 F. Zupanic et al. / Journal of Alloys and Compounds 570 (2013) 125–132

maximum solubility of Mn in Al. Thus the all Mn was dissolved ina-Al. The microstructure resembled that of the binary eutectic inthe Al–Be phase diagram. The same two-phase matrix appearedin the much more concentrated alloy Al66Mn4Be30. XRD-resultsand the microscopic investigations clearly revealed the presenceof Be4AlMn and large a-Be particles (Fig. 3b). This indicated thatMn, undissolved in a-Al, formed Be4AlMn, and that the redundantBe was present as a-Be. DSC revealed that the solidus temperaturewas around 640 �C (Fig. 4), which was within the range of the eu-tectic temperature in the Al–Be system (644 �C) [5,6].

These results indicated that two completely different micro-structures occurred upon solidification in the Al-rich corner ofthe ternary system Al–Mn–Be. The bordering line was approxi-mately the tie-line Al–Be4AlMn; and approximately coincided with

the line having the atomic ratio Be:Mn equals 4:1. In fact, the posi-tion of Be4AlMn slightly deviates from this line because, in Al–Mn–Be alloys, Be4AlMn contains more Al than Mn. In addition, a-Al alsodissolved at around 0.3 at.% Be and 0.6 at.% Mn. Be4AlMn and a-Alappeared on the Be-rich side of the tie-line, whereas a-Al, Be4-

AlMn, Al6Mn and T-phase could be found on the Mn-rich side ofthe tie-line.

Fig. 5 shows the as-cast microstructures of alloys with higherMn- and Be-contents. In both cases the Be: Mn ratio was lowerthan 4:1. The alloy Al79Mn17Be4, which was positioned very closeto the binary Al–Mn system, consisted of Al11Mn4, partly envel-oped by k-Al4Mn, and a-Al (Fig. 5a). Neither Be-rich phases norAl6Mn were present. Al11Mn4 was undetected using XRD, and thesame fact was reported by Song et al. [16]. DSC showed morepeaks, pointing to several processes occurring upon melting andsolidification (Fig. 6). Upon heating an exothermic peak occurredat 686 �C, indicating the transition of the metastableAl11Mn4-phase to more stable phases (e.g. l-Al4.12Mn). The moreprobable reactions on cooling with the reference to the binaryAl–Mn system are: (1) L ? Al11Mn4 (HT); (2) L + Al11Mn4

(HT) ? l-Al4.12Mn, (3) L + l-Al4.12Mn ? k-Al4Mn and (4) metasta-ble L ? a-Al + k-Al4Mn.

In the Al73Mn14Be13 alloy, the T-phase prevailed (Fig. 5b), and aconsiderable amount of Be4AlMn was also present. Additionally,the Al11Mn4-phase occasionally appeared, however its fractionwas very small. DSC-traces showed several peaks occurring uponheating and cooling (Fig. 6). Reactions taking place during solidifi-cation could be: (1) L ? Al11Mn4 (HT), (2) L + Al11Mn4 (HT) ? T; (3)L ? T + Be4AlMn and (4) metastable reaction L ? a-Al + Be4AlMn + T.

3.2. Isothermal annealing

Samples solidified within a water-cooled copper mould in thearc melting furnace were used for isothermal annealing. Their

Fig. 9. Secondary electron image (SEI) and Auger electron maps for Al, Be and Mn in alloy Al97.3Mn1.9Be0.8 heat-treated for 720 h at 600 �C.

F. Zupanic et al. / Journal of Alloys and Compounds 570 (2013) 125–132 129

microstructures were therefore much finer than those of the DSC-samples. This is a prerequisite for faster approaching towards ther-modynamically-stable conditions.

3.2.1. Isothermal annealing at 600 �CFig. 7 shows the diffraction patterns for the alloy

Al97.3Mn1.9Be0.8. The as-cast condition comprised of a-Al, Al6Mnand Be4AlMn, and a small fraction of the T-phase. The peaks of firstthree phases were clearly visible in the corresponding XRD-trace.The annealing at 600 �C caused the appearance of thek-Al4Mn-peaks, and a decrease in the Al6Mn-peaks. After 720 hthe Al6Mn-peaks completely disappeared. Fig. 8a shows the trans-formation of Al6Mn to k-Al4Mn. Phase Al6Mn is on the left, andk-Al4Mn is on the right. The transformation could take place bythe following reaction: Al6Mn ? k-Al4Mn + a-Al. Thus, a homoge-neous Al6Mn-phase transformed into a two-phase structurek-Al4Mn + a-Al. Fig. 8b shows the fully transformed microstructureafter annealing for 1 month. It consisted of a bright k-Al4Mn, a-Almatrix and Be4AlMn particles. These results were unambiguouslyconfirmed by the AES. Fig. 9 shows the Auger electron maps forAl, Mn and Be. The chemical compositions of Be4AlMn and k-Al4Mnwere Be66 ± 2Al20 ± 1Mn14 ± 1 and Al75 ± 3Mn18 ± 2Be7 ± 2, respec-tively. Thus, k-Al4Mn dissolved a considerable fraction of Be thusincreasing the stability of k-Al4Mn.

Fig. 10 shows the Auger electron maps for Al, Mn and Be in thealloy Al66Mn4Be30 after annealing for one month at 600 �C. Thesame three phases as present in the as-cast state (a-Al matrix,slightly brighter Be4AlMn particles, and black a-Be particles) stillpersisted in the microstructure. The difference between a-Be andBe4AlMn was evident. The shapes and sizes of the phases had chan-ged considerably, demonstrating that the diffusion must have ta-ken place to a considerable extent, thus allowing a fast approachtowards the equilibrium condition. It should be stressed that noMn-rich phase, such as Al6Mn, T-phase or k-Al4Mn appeared in thisalloy. Some Mn was dissolved in the matrix (less than 1 at.%), andthe main fraction of Mn was bound in Be4AlMn. The composition of

Be4AlMn was Be64 ± 2Al23 ± 1Mn13 ± 1; therefore it contained moreAl and less Mn than in the alloy Al97.3Mn1.9Be0.8. It could be statedthat the equilibrium phases in this part of the Al–Mn–Be phase dia-gram are a-Al, Be4AlMn and a-Be.

No evident variations in the microstructure and phase compo-sition were observed after the annealing of Al73Mn14Be13

alloy. This showed that this type of microstructure is highlyresistant to any changes. The composition of T-phase wasAl72±2Mn19±2Be9±2. This phase is metastable, but it cannot be eas-ily transformed into other phases. Its thermodynamic stability isprobably only slightly smaller than that of k-Al4Mn, thus the driv-ing force for the transformation is very low. Both the T-phase andk-Al4Mn have very similar chemical compositions, and they prob-ably consisted of clusters arranged in a similar way. In addition,the reaction area was limited only to the surface of the T-phaseparticles, which was also larger in the faster cooled samples. Songet al. [16] observed that the T-phase had not changed uponannealing at 540 �C, and they concluded that it is a stablephase. In fact, it could be a stable phase, but outside thea-Al-k-Al4Mn–Be4AlMn triangle.

Fig. 11 gives the isothermal section of the Al–Mn–Be phase dia-gram at 600 �C within the region limited by Al, Be, Be4AlMn andk -Al4Mn based on the current results. In this part of the diagramtwo three-phase fields dominated: a-Al�a-Be�Be4AlMn anda-Al�Be4AlMn and k-Al4Mn. The one-phase fields for a-Be andAl6Mn are represented by points, whilst the one-phase fields fora-Al, Be4AlMn and k-Al4Mn are small areas, allowing smallvariations in compositions. The two-phase fields a-Al + a-Be,a-Al + Be4AlMn,a-Al + k-Al4Mn,a-Al + Al6Mn and Al6Mn + k-Al4Mnare very narrow strips, and some of them could be represented bylines. The three-phase region a-Al + Al6Mn + k-Al4Mn is predictedto be close to the boundary Al–Mn phase diagram.

This horizontal cross-section differs considerably from the pub-lished compilations of this ternary phase diagram. It is evident thatin this region the T-phase is not an equilibrium phase inside the re-gion limited by Al, Be, Be4AlMn and k-Al4Mn. The composition of

Fig. 10. A secondary electron image (SEI) and Auger electron maps for Al, Mn andBe in alloy Al66Mn4Be30 heat-treated for 720 h at 600 �C.

Fig. 11. Schematic partial isothermal cross-section at 600 �C.

130 F. Zupanic et al. / Journal of Alloys and Compounds 570 (2013) 125–132

the T-phase is shifted towards the Mn-corner relative to theBe4AlMn�k-Al4Mn tie-line. It could be inferred that it is slightlyless thermodynamically stable than k-Al4Mn. It could be formedduring solidification in alloys with Be: Mn < 4:1, and Mn-contenthigher than approximately 1.5 at.%. However, it is afterwardstransformed into k-Al4Mn. Thus the role of the T-phase in theAl–Mn–Be phase diagram is similar to that of the l-Al4.12Mn inthe binary Al–Mn system.

3.2.2. Isothermal annealing at 750 �CFig. 12 shows microstructures of the alloys Al79Mn17Be4,

Al73Mn14Be13 and Al66Mn4Be30 after annealing for 720 h at750 �C and then quenched, and Fig. 13 shows the correspondingdiffraction patterns. The alloy Al79Mn17Be4 contained onlyk-Al4Mn, and a small fraction of a-Al. Thus, this alloy consistedof l-Al4.12Mn and L at 750 �C, which transformed into k-Al4Mnand a-Al upon cooling. The alloy Al73Mn14Be13 consisted ofBe4AlMn, the T-phase and a-Al. The latter phase formed duringsolidification. The sizes of the Be4AlMn and T-phase particles in-creased significantly when being held at 750 �C because of thepresence of the liquid phase. Normally, the diffusion coefficientsin the liquid phase are several orders of magnitude higher than

in the solid, resulting in a faster approach towards theequilibrium state. The chemical compositions of Be4AlMn andthe T-phase were Be62 ± 2Al22 ± 2Mn16 ± 1 and Al71 ± 2Mn19 ± 1Be10 ± 1.

The alloy Al66Mn4Be30 was composed of Be4AlMn, a-Be and a-Al.At the annealing temperature, a large fraction of Be4AlMnwas present in the liquid phase. Its composition wasBe58 ± 2Al26 ± 2Mn16 ± 1. The fraction of a-Be was much smaller. Adendritic morphology of a-Be indicated that it formed predomi-nantly on cooling.

These results clearly indicated that the T-phase is thermody-namically more stable than the competitive l-Al4.12Mn and k-Al4-

Mn at 750 �C. It appeared in the three-phase equilibrium with theliquid phase and Be4AlMn, over a rather wide concentration range.However, only in alloys with the atomic Be:Mn ratio smaller than4. In alloys with the Be:Mn ratios greater than 4:1, the resultsstrongly suggested the three-phase equilibrium between Be4AlMn,the liquid phase, and a-Be.

Also at these temperatures, other one-phase and two-phasefields close to Al–Be and Al–Mn boundary systems are very shal-low. It is unclear how considerable solubility of Be in k-Al4Mn af-fects its thermodynamic stability. Its stability may increase,therefore it could be thermodynamically stable at higher tempera-tures than 722 �C in the binary Al–Mn system [8].

3. 3 Prediction of the liquidus surface in the Al-corner of the Al-Mn-Bephase diagram

The results of the investigation indicate the following liquidussurfaces (primary solidification areas): a-Al, a-Be, Be4AlMn, Al6Mn,k-Al4Mn, l-Al4.12Mn (very close to the Al–Mn system), the T-phase,and Al11Mn4.(Fig. 14).

The exact phase diagram, especially at the Al–Mn side of thediagram, is rather complicated [8]. The DSC-results strongly indi-cate that the temperature decreases from the binary eutectic reac-tion in the Al–Mn system (657 �C) towards the eutectic reaction inthe binary Al–Be system (644 �C), i.e. from point e1 towards e2. Inbetween, two transition reactions take place: L + AlMn6 ?a-Al + Be4AlMn (T1) and L + Be4AlMn ? a-Al + a-Be (T2). A binaryeutectic reaction L ? a-Al + Be4AlMn takes place between thesetransition reactions.

Fig. 12. Backscattered electron micrographs of the alloys Al79Mn17Be4 (a), Al73-

Mn14Be13 (b) and Al66Mn4Be30 (c) after annealing for 720 h at 750 �C, andquenching.

Fig. 13. XRD-traces of the alloys Al79Mn17Be4, Al73Mn14Be13 and Al66Mn4Be30 afterannealing for 720 h at 750 �C, and quenching. The insert shows the trace for thealloy Al66Mn4Be30 between 32� and 34� where two Be-peaks can be clearly seen.

Fig. 14. Schematic presentation of liquidus projections in the Al-corner of the Al–Mn–Be phase diagram.

F. Zupanic et al. / Journal of Alloys and Compounds 570 (2013) 125–132 131

The univariant reactions close to the Al-corner are:

e1—T1 : binary eutectic reaction L! a-Alþ AlMn6 ð8Þ

T1—T2 : binary eutectic reaction L! a-Alþ Be4AlMn ð9Þ

T2—e2 : binary eutectic reaction L! a-Alþ a-Be ð10Þ

The invariant reactions:

T1 : Lþ AlMn6 ! a-Alþ Be4AlMn ð11Þ

T2 : Lþ Be4AlMn! a-Alþ a-Be ð12Þ

There should also be three additional invariant reactions, inwhich the T-phase takes part. There is no clear evidence for allthese reactions because they occur very close to the Al–Mn system,but we can suggest the following reactions:

T3 : Lþ T! Be4AlMnþ Al6Mn ð13Þ

T4 : Lþ k-Al4Mn! Tþ Al6Mn ð14Þ

T5 : Lþ l-Al4:12Mn! Tþ k-Al4Mn ð15Þ

T6 : Lþ Al11Mn4ðHTÞ ! l-Al4:12Mnþ T ð16Þ

3.4. Comparison with the published ternary Al–Mn–Be phase diagram

This investigation did not bring any new information regardingthe boundary Al–Be, Al–Mn and Be–Mn phase diagrams. However,there were established many differences with the existing phasediagram that based on the work of Raynor et al. [10], and was com-piled by Mondolfo [11] and Stiltz [12]. The main difference repre-sented the presence of Be4AlMn as an equilibrium phase in the Al-corner. This phase possessed its region of primary solidification,and was present in the isothermal sections at 600 �C and 750 �C.The T-phase found by Raynor et al. [10] can appear in the as-caststate of alloys close to the Al-corner, but it transformed to the equi-librium k-Al4Mn phase at 600 �C. Nevertheless, T-phase possessedits region of primary solidification as in published phase diagram,and can be an equilibrium phase in the alloy at higher tempera-tures, in the presence of the liquid phase (e.g. at 750 �C).

The described differences caused completely different invariantreactions upon solidification. Instead of reactions (6) and (7), reac-tions (11) in (12) takes place, in which Be4AlMn takes part and theT-phase is absent.

132 F. Zupanic et al. / Journal of Alloys and Compounds 570 (2013) 125–132

4. Conclusions

The results of this work lead us to the following conclusions:

Be4AlMn is a thermodynamically stable phase in the Al-richcorner of the ternary system Al–Mn–Be. The two three-phaseequilibria a-Al + a-Be + Be4AlMn and a-Al + k-Al4Mn + Be4AlMnoccupy the largest part of the partial isothermal cross sectionAl–Be–Be4AlMn-k-Al4Mn at 600 �C. At this same temperature,and in this part of the Al–Mn–Be phase diagram, the ternarycompound T-phase (Al71Mn19Be10) is not an equilibrium phase.This phase formed via a binary peritectic reaction L + Al11Mn4 (-HT) ? T-phase on cooling. It is thermodynamically stable at750 �C, within a very large three-phase region: liquid phase(with small fraction of Be and Mn), Be4AlMn and the T-phase.The T-phase did not appear in alloys with Be:Mn > 4:1 underany investigated conditions.The solidification microstructures and DSC-results strongly indi-cate that the temperature decreases from the binary Al–Mneutectic point towards the binary Al–Be eutectic point. Betweenthese two points there are three univariant binary eutectic reac-tions: L ? AlMn6 + a-Al, L ? a-Al + Be4AlMn, and L ? a-Al + a-Be. Univariant binary eutectic troughs are interconnected withtwo invariant transition reactions: L + AlMn6 ? a-Al + Be4AlMnand L + Be4AlMn ? a-Al + a-Be.This investigation showed that the current results cannot beexplained by the published ternary phase diagram Al–Mn–Be.Further investigations and computer modelling of the alloysclose to the binary Al–Mn phase diagram are required to fullyexplain the constitution of the Al–Mn–Be system in the Al-richcorner.

Acknowledgements

This work was partly financed by the research programme P2—0120 (Slovenian Research Agency – ARRS). Part of the work wascarried out with the support of the European Community. We

appreciate the support of the European Research InfrastructureEUMINAfab (funded under the FP7 specific programme Capacities,Grant Agreement Number 226460), its partner Karlsruhe Instituteof Technology (KIT) and Tobias Weingärtner. The XRD-investiga-tions at Elettra, Sincrotrone Trieste, Italy, were funded by the Euro-pean Community’s Seventh Framework Programme (FP7/2007-2013) under Grant Agreement No. 226716. The authors also wishto thank Mrs. Vesna Krapez for the metallographic preparation ofsamples and Prof. Dr. Joze Medved, University of Ljubljana, Facultyof Natural Sciences and Engineering, for the DSC analyses. Theauthors would also like to acknowledge Dr. Paul McGuiness fromthe Jozef Stefan Institute, Slovenia, for the synthesis of the investi-gated alloys.

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