- Email: [email protected]
Al-rich region of Al–Cu–Mn
Journal of Alloys and Compounds 688 (2016) 957e963
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Al-rich region of AleCueMn B. Grushko a, b, *, S.B. Mi c a
MaTecK, D-52428 Jülich, Germany PGI-5, Forschungszentrum Jülich, D-52425 Jülich, Germany c State Key Laboratory for Mechanical Behavior of Materials, Xi’an Jiaotong University, Xi’an 710049, PR China b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 25 April 2016 Received in revised form 1 July 2016 Accepted 6 July 2016 Available online 7 July 2016
The AleCueMn alloy system was studied above 45 at.% Al between 550 and 910 C by scanning electron microscopy, powder X-ray diffraction and differential thermal analysis. A wide ternary region of the socalled b-phase extending along ~50 at.% Al was confirmed. It has a CsCl-type structure with a z 0.292 e0.298 nm. The total g1/g2-region extends from AleMn up to ~17 at.% Cu and the high-temperature TAl3Mn phase (Pnma, a z 1.48, b z 1.24, c z 1.25 nm) extends up to ~15 at.% Cu. The so-called R-phase (Bbmm, a z 2.41, b z 1.25, c z 0.76 nm) was found to exist in a compositional region of Al74-80Cu512.5Mn12.5-18. The ternary phase earlier reported at Al57.9Cu26.3Mn15.8 was confirmed. It exists below 697 C in a compositional region of Al55-58Cu29-37Mn7.5-14. The decagonal D3-phase was concluded to be stable in a compositional region of Al61.5-68.5Cu19-29.5Mn9-16. Below 631 C an fcc phase (a ¼ 0.5814 nm) was revealed around ~ Al60Cu36.5Mn3.5. Partial isothermal sections at 550, 600, 650, 750, 850 and 910 C were constructed. © 2016 Elsevier B.V. All rights reserved.
Keywords: Transition metal alloys and compounds Phase diagrams
1. Introduction The ternary AleCueMn alloy system has been extensively studied for several decades. The corresponding literature up to 2003 was reviewed in Ref. [1]. Since the first publication on the ternary constitution of this system in 1927 [2], the boundary binary alloy systems has been updated and modified several times, which also required subsequent revisions of the ternary phase diagram in order to bring them into agreement. For the following description the relevant parts of the updated phase diagrams AleCu [3] and AleMn [4] are shown in Fig. 1. The crystallographic data of the phases are provided in Table 1. Of the earlier reports, those in Refs. [5e11] are most relevant to the present work. In contrast to Ref. [2], where no ternary compounds were revealed, two complex orthorhombic structures were reported in Ref. [5] on the basis of the X-ray diffraction study of needle-shaped single crystals. The phase designated T in Ref. [5] (TCu in the following1) with a ¼ 0.769, b ¼ 2.406 and
* Corresponding author. PGI-5, Forschungszentrum Jülich, D-52425 Jülich, Germany. E-mail address: [email protected] (B. Grushko). 1 This is in order to discriminate it from the binary T-Al3Mn phase (or just Tphase). http://dx.doi.org/10.1016/j.jallcom.2016.07.075 0925-8388/© 2016 Elsevier B.V. All rights reserved.
c ¼ 1.248 nm was found to contain 13e15 wt% Cu and 20e25 wt% Mn (i.e. between ~ Al80.1Cu7.8Mn12.1 and ~Al77.7Cu6.9Mn15.4). The other phase designated Y was found to have a ¼ 1.479, b ¼ 1.260, c ¼ 1.243 nm. It was observed in a range of compositions adjacent to AleMn. Particularly, a composition of 5.9 wt% Cu and 32.5 wt% Mn (~Al76.9Cu3.1Mn20.0) was reported. The c-axes of both TCu and Y were selected along the needle axes. In Ref. [6] the existence of the former was confirmed, but no information on the latter was supplied. Instead, a new phase was reported in Ref. [6] at Al60Cu25Mn15 (designated Al11Cu5Mn3, such a different composition did not allow it to be associated with the Y-phase). Indeed, in Ref. [10] this Al11Cu5Mn3 was reported to be orthorhombic with a ¼ 1.210, b ¼ 2.408, c ¼ 1.921 nm. In Ref. [7] the ternary phases of Ref. [5] were only mentioned as being under investigation. The most detailed experimental investigation of the whole ternary AleCueMn phase diagram was reported in Ref. [10]. It was basically accepted in the assessment [1] and only slightly modified in order to be more compatible with the later updates of the boundary binary AleCu and AleMn phase diagrams mentioned above. Particularly, in the boundary AleCu phase diagram used in Ref. [10], the g0, g1 and d phases accepted to date were not distinguished inside a region designated G. Similarly, b0, b, ε1 and ε2 were not distinguished inside the b-region, and h1 and h2 inside the h-region, while the phases a2, z1 and z2 were ignored. As to the
958
B. Grushko, S.B. Mi / Journal of Alloys and Compounds 688 (2016) 957e963
Fig. 1. Relevant parts of the binary phase diagrams of: a) AleCu [3] and b) AleMn [4].
Table 1 Crystallographic data of the relevant phases in AleCueMn. TW means this work. Phase
Space group
Lattice parameters a, nm
q-Al2Cu h1 h2 ε2 Al6Mn l-Al4Mn m-Al4Mn n-Al11Mn4
I4/mcm Pban or Cmmm C2/m P63/mmc Cmcm P63/m P63/mmc
T- Al3Mn (HT- Al11Mn4)
P1 Pnma
g1-Al8Mn5 g2-Al8Mn5
I43m R3m
d-AlMn b
Pm3m
R
U L
Im3m
Bbmm Orth. FCC
0.6063 0.4087 1.2066 0.4146 0.75551 2.8382 2.0015 0.5095
Ref. Comment
a,
b, nm
89.35
1.200 0.4105 e 0.64994 e e 0.8879
b,
c, nm
e
0.4872 0.8635 0.6913 0.5063 0.88724 1.2389 2.4699 0.5051
124.96
100.47
g,
105.08
[3] [3] [3], standardized [3] [4] [20] [14], Al79.8Mn20.2 [21]
1.4873 1.46256 0.9012
1.2420 1.24557 e
1.2547 1.25469 e
[14], Al72.8Mn277.2 [13], Al70.0Cu11.5Mn18.5 [22]
1.2611 1.26403 (23) 0.3063
e e e
0.7927 0.78272 (9) e
[23], Al55Mn45a TW, Al56Cu15Mn29 [23], Al55Mn45, at 957 C
e
e
1.24787 2.408 e
0.76102 1.921 e
TW, Al49.5Cu22Mn28.5 TW, Al48Cu50Mn2 [13] Al80Cu8Mn12 [10] TW
0.29804 (6) 0.29210 (4) 2.41023 1.210 0.5814 (1)
a
The g2estructure can also be described as a distorted g-brass structure with the basic (body-centered) lattice. For Al57Mn43 the corresponding lattice parameters are a z 0.896 nm and a z 89.02 [52]. In this notation, the g2estructure contains the same 52 atoms in the unit cell as the cubic g1estructure.
boundary AleMn phase diagram, updated somewhat later by the same authors in Ref. [11], MnAl(h) and Mn5Al8(h) were still not distinguished in Ref. [10], as well as Mn4Al11(r) and Mn4Al11(h) (the latter is also known as T-Al3Mn). Much later the l-Al4Mn phase was added to the AleMn phase diagram besides m-Al4Mn [12]. The Al-rich part of the ternary diagram assessed in Ref. [1] contains two complex intermetallic compounds t1 and t2 forming in narrow compositional regions and a wide region of the so-called
b solid solution of the W-type structure related to a corresponding cubic modification of Mn, to AleCu b and to high-temperature AlMn. The ternary t1-phase was associated with the abovementioned TCu-phase of Ref. [5], but placed at considerably lower Al concentrations. It was concluded that t1 was formed by the reaction L þ Al8Mn5 4 t1 at 1020 C. The ternary t2-phase forming by the reaction b þ t1 4 t2 at ~700 C was associated with Al11Cu5Mn3 of Ref. [6]. The Y-phase of Ref. [5] was just mentioned as not
B. Grushko, S.B. Mi / Journal of Alloys and Compounds 688 (2016) 957e963
confirmed. On the other hand, as was shown in Ref. [13], the ternary t1phase of Ref. [1] exhibited a structure of the above-mentioned Yphase of Ref. [5], which is actually the ternary extension of binary TAl3Mn. In other words, the t1-phase was associated with the wrong one of the two structures revealed in Refs. [5], which also explains such a significant difference in the compositions reported in Refs. [5] and [10]. The other orthorhombic phase of Ref. [5] was found to be formed at somewhat higher Al concentrations than this ternary extension and quite close to the compositions determined in Ref. [5] (see Fig. 5 of Ref. [13]). Similarly to that in AlePdeMn ([14] and references therein), the later phase was designated R. After the discovery of icosahedral (I) and decagonal (D) quasicrystals in AleMn, many efforts were made to stabilize these metastable structures by the addition of a third element, particularly also by Cu. The I-phase, forming in rapidly solidified Al65Cu20Mn15 and several other alloys, was found to transform rapidly by moderate annealing [15,16]. The final constituencies in three samples studied in Ref. [16] (black scuares in Fig. 2) were concluded to be either the above-mentioned t1, or t2, or t1 þ t2 of Refs. [1,10] without specifying the structures of these phases. A D-phase was observed in the intermediate stages of this transformation and together with the I-phase was concluded to be metastable in AleCueMn. However, this conclusion was premature, as was revealed in Ref. [13] and will be described in more detail below. In the present contribution the Al-rich part of the AleCueMn phase diagram was revised above 45 at.% Al at 550e910 C. The above-mentioned updated boundary binary AleMn and AleCu phase diagrams were accepted. 2. Experimental The alloys were produced by levitation induction melting in a water-cooled copper crucible under a pure Ar atmosphere. The purity of Al was 99.999%, of Cu and Mn 99.95%. The ingot weights
Fig. 2. Overall compositions of the phases in AleCueMn revealed in the present study. The total T-phase region above 650 C is shown by a brighter fill, that at 850 C by a darker fill and at 910 C surrounded by solid line, the upper-Cu limit of this region at 750 C is shown by a broken line (see isothermal sections below). The compositions of the samples studied in Ref. [16] are marked by squares. The broken lines continuing b (B2) towards d (A2) and the lower-Al boundary of b are drawn provisionally.
959
were typically ~5 g. Parts of the ingots were thermally annealed at 910 C for 24 h, at 850 C for 67 h, at 750 C for 148e332 h, at 650 C for 163e310 h, at 600 C for 330e404 h, and at 550 C for 206e1460 h. The alloys were examined by powder X-ray diffraction (XRD), scanning electron microscopy (SEM), and differential thermal analysis (DTA). The local phase compositions were determined in SEM by energy-dispersive X-ray analysis (EDX) on polished unetched cross-sections. Powder XRD was carried out in the transmission mode using Cu Ka1 radiation and an image plate detector. DTA was carried out in alumina crucibles at heating rates of 25 K/min and cooling rates of 50 K/min.
3. Results and discussion 3.1. Binary phases and their ternary extensions In general, the AleCu intermediate phases are formed at lower Al concentrations and lower temperatures than the AleMn phases. In the studied compositional and temperature region the relevant AleCu phases are q, high-temperature h1, low-temperature h2, and high-temperature ε2 (see Fig. 1a and Table 1). All of them were found to dissolve very little Mn. The Al-rich part of the AleMn phase diagram in Fig. 1b contains the equilibrium phases Al6Mn, l-Al4Mn, m-Al4Mn, the T-Al3Mn phase (also known as high-temperature Al11Mn4) and the Al11Mn4 (n-phase). The lowest-Mn Al12Mn phase is only stable below 512 C [17], i.e. below the studied temperatures and was not observed in the present study. The central part of the AleMn phase diagram contains high-temperature Al8Mn5 (g1), low-temperature Al8Mn5 (g2) and a high-temperature AlMn phase designated at present d due to its structural similarity to d-Mn (g-phase in Ref. [4]). The accepted AleMn phase diagram was additionally argued in recent Ref. [18]. Of the AleMn binaries, Al6Mn and l-Al4Mn did not exhibit visible solubility of Cu, while m-Al4Mn and n-Al11Mn4 were found to dissolve up to ~2 at.% Cu. On the other hand, the phases T, g1 and g2 were observed at essentially ternary compositions (Fig. 2). Similarly to the case at binary compositions, the Cu-containing high-temperature g1-phase could not be conserved by water quenching and transformed to g2. The existence of g1 at elevated temperatures was detected by DTA. For example, the thermal effect at 734 C in Fig. 3a indicates the transformation of g2 to g1. According to SEM/EDX examination, a sample of this composition annealed at and quenched from 850 C, i.e. above the transition temperature, consisted of essentially a single phase of Al56Cu15Mn29, whose powder XRD pattern indicated the g2 structure (Fig. 4a, Table 1). With increasing Cu concentration the upper existence temperatures of g1 and g2 were found to decrease and the temperature of the corresponding transition decreased accordingly (compare to the lowest transition temperature 957 C in binary AleMn, Fig. 1b). The total g1/g2 region extends up to at least 17 at.% Cu (see Fig. 2). The ternary region of T-Al3Mn was found to propagate up to 15 at.% Cu, extending more with decreasing temperatures (see Fig. 2). The addition of Cu also decreases the lower existence limit of the T-phase, so below ~895 C it is only stable at ternary compositions (see the corresponding isothermal sections). At 5 at.% Cu it was observed down to 650 C and at the rates applied in our DTA experiments its decomposition was not detected. DTA plots in Fig. 3b and c demonstrate the decrease of the upper temperature limits of the T-phase containing 5 and 11.5 at.% Cu, respectively. In addition to the major T-phase the former sample also contained some n-phase, whose transformation is visible in Fig. 3b starting slightly above 750 C (shallow effect marked by an arrow).
960
B. Grushko, S.B. Mi / Journal of Alloys and Compounds 688 (2016) 957e963
a)
734
b)
Heat flow, a.u.
962
c) 697
815
d) 544
572
660
e) 763
f) 631
Exo 400
500
600
700
800
900
1000
1100
T, °C Fig. 3. DTA plots (heating at 25 K/min) of: a) Al56Cu15Mn29 annealed for 67 h at and quenched from 850 C, b) Al70Cu5Mn25 annealed for 310 h at 650 C, c) Al70Cu11.5Mn18.5 annealed for 332 h at 750 C, d) Al57.5Cu29Mn13.5 annealed for 1460 h at 550 C, e) Al75.3Cu10Mn14.7 annealed for 163 h at 650 C, f) Al65.1Cu24.6Mn10.3 annealed for 404 h at 600 C.
3.2. Ternary phases The formation of a phase reported in Ref. [10] in a wide ternary region below ~50 at.% Al was confirmed in the present study. In a slightly modified version of Ref. [1] this so-called b-phase (Im3m, a z 0.29e0.30 nm) is actually a ternary extension of AlMn (d-phase in Fig 1b) and the isostructural phases in AleCu (outside the region in Fig. 1a) ranging in total between the AleCu and AleMn binary terminals. In both AleCu and AleMn the bcc phases are only stable at elevated temperatures (for example, d-phase above 840 C, see Fig. 1b). In the present work only alloys revealing the Al-rich boundary of the b-region were studied. At these compositions the b-phase was
observed down to 600 C, i.e. below the stability temperatures of the d-phase, and a CsCl-type ordering was revealed in the studied alloys. In a typical powder XRD pattern in Fig. 4b this is indicated by the presence of the (100) reflection, which is forbidden in a bcc structure. The ordering was observed in the whole compositional region depicted in Fig. 2 by the grey b-field and at all studied temperatures. In Ref. [1] the CsCl-type ordering inside the b-region was only mentioned around the AlCu2Mn composition as an alternative to the ordering typical of the Heusler phases (i.e. cF16 vs. cP2). On the other hand, similar to AleFeeMn [18], a link of this region with the region of the d-phase is indeed plausible. With decreasing temperature the Al-rich boundary of the b-phase was found to propagate towards higher Cu and at 650 C it achieved the Al48Cu50Mn2 composition. A phase, first reported in Ref. [6] under the name Al11Cu5Mn3, specified and included in the phase diagram of Ref. [10] (t2 in Ref. [1]), was revealed in the present work in a compositional region of Al55-58Cu29-37Mn7.5-14. Its powder XRD pattern (Fig. 4c) was qualitatively similar to that simulated from Table 1 of Ref. [10] where the intensities of the reflections were provided qualitatively. No additional structural investigation of this phase, designated in the following U, was carried out. It was found in samples annealed at 550e650 C, but not at 750 C. In the DTA plot of Al57.5Cu29Mn13.5 in Fig. 3d a transformation of the essentially single U ephase started at ~645 C with a pronounced effect at 697 C, i.e. at a temperature close to that determined in Ref. [10]. The observation of a temperature range of this transformation is in agreement with the existence of a compositional region of U revealed by SEM/EDX. Similarly to Ref. [10], the U -phase was found to be transformed to a mixture of two phases, one of which was identified with the b-phase. The second transformation product of the Uphase had a composition of ~Al64Cu20Mn16, which is quite far away from that of the T-phase claimed in Ref. [10], but very close to a composition where the metastable decagonal phase was observed in Ref. [16]. SEM/EDX examination of a sample with such a nominal composition indicated that the material was solid after annealing at 750 C for 238 h and subsequent water quenching. Apart from the major phase of ~Al64.2Cu20.3Mn15.5 it contained a little of a second phase of ~Al66.5Cu15.4Mn18.1. From these compositions the latter was rather typical of the T-phase,
Fig. 4. Powder XRD patterns of: a) g2 in Al56Cu15Mn29 annealed for 67 h at and quenched from 850 C, b) b in Al49.5Cu22Mn28.5 annealed for 145 h at 650 C (the (100) reflection is marked by an arrow), c) U in Al57.5Cu29Mn13.5 annealed for 1460 h at 550 C, d) D3 of Al64Cu20Mn16 annealed for 238 h at 750 C, e) T in Al70Cu11.5Mn18.5 annealed for 1460 h at 550 C, f) R in Al80Cu8Mn12 annealed for 1460 h at 550 C, g) m, powdered single crystal (Al79.8Mn20.2), h) L in Al61Cu35.5Mn3.5 annealed for 1460 h at 550 C.
Heat flow, a.u.
B. Grushko, S.B. Mi / Journal of Alloys and Compounds 688 (2016) 957e963
774 717
763
Exo 600
700
800
900
1000
T, °C Fig. 5. DTA plot of Al65Cu18.5Mn16.5 pre-annealed at 750 C, held in the DTA device for one hour at 600 C, cooled down to 500 C (not shown) and afterwards heated to 950 C at 10 K/min. A similar result was obtained by a direct heating at 25 K/min.
961
indicating the formation of an additional phase. The powder XRD pattern (Fig. 4d) of this sample was indeed different from that of the T-phase (Fig. 4e), but rather resembled the XRD pattern of the stable AleFeeMn decagonal D3-phase [18], as was already reported in Ref. [13]. Electron diffraction examinations in Ref. [13] confirmed the D3-structure. Additional examinations revealed that the compositional region of this phase shifts with decreasing temperature towards lower Mn, so at 650 C the abovementioned Al64Cu20Mn16 composition corresponds to the twophase region TþU, but at this temperature the D3-phase is formed around ~ Al65Cu23Mn12. As follows from our annealing and DTA experiments, the D3structure is stable in AleCueMn. The experimental results of Ref. [16], leading its authors to erroneous conclusions of the total
Fig. 6. Partial isothermal sections of AleCueMn at: a) 910 C, b) 850 C, c) 750 C, d) 650 C. The measured compositions of the alloys are marked by solid squares, and those of phases by open squares. Suggested three-phase equilibria are marked by broken lines. The broken lines separating the regions of g1 and g2 are drawn provisionally based on DTA examinations. L is liquid.
962
B. Grushko, S.B. Mi / Journal of Alloys and Compounds 688 (2016) 957e963
instability of the decagonal phase in AleCueMn, were incorrectly interpreted. The D3-phase produced in Ref. [16] indeed transformed by moderate heating but could recombine by heating at somewhat higher temperature. This can be demonstrated by an experiment with a sample of Al65Cu18.5Mn16.5 annealed at and water quenched from 750 C. In the subsequent DTA examination it was first held in the DTA device for one hour at 600 C, cooled down to 500 C and afterwards heated to 950 C at 10 K/min. After quenching from 750 C, the sample consisted of the D3-phase, which decomposed by the first heating to and annealing at 600 C. The plot in Fig. 5 was recorded during the second heating from 500 to 950 C. The first weak minimum at ~717 C corresponds to the recombination of the D3-phase. The transformation of the D3phase starts at ~763 C and above ~774 C this phase does not exist anymore, i.e. at this composition, close to that examined in Refs. [16], it is stable in a quite narrow temperature range. Such a possibility was not tested in Ref. [16], and this study only was on single compositions. The ternary R-phase was found to be formed in a region of Al74-80Cu5-12.5Mn12.5-18 close to the upper-Al limits of the ternary extension of the T-phase (see Fig. 2). With decreasing temperature its compositions exhibited a shift towards higher Al, so that the total regions of the R-phase and T-phase partially overlap. The coexistence of these two phases was not identified metallographically due to probably small compositional differences and also due to their formation in the form of fine weakly connected needles. Their coexistence was also not revealed by powder XRD. Both these phases have complex diffraction patterns with numerous overlaps (see Fig. 4e and f). In Fig. 2 and the isothermal sections (see below) the separation between the T-region and R-region is shown conditionally. On the other hand, the coexistence of R and m was observed by SEM/EDX [13]. The powder XRD pattern of the m-phase (see Fig. 4g) is also complex and contains lines overlapping with those of the T-phase and Rphase. The R-phase was already observed at 750 C but not at 850 C and its solidus temperature decreases with increasing Cu
concentration. In the DTA plot in Fig. 3e an endothermal effect at ~660 C could be associated with the start of the incongruent melting of the R-phase resulting in the formation of the T-phase, which continues to melt at higher temperatures. At lower heating rates the DTA plots were smeared even more, which did not allow the reaction effects to be separated. Apart from the above-mentioned phases, a new compound was revealed in a small region around ~ Al60Cu36.5Mn3.5. This phase exhibited a quite simple powder XRD pattern (Fig. 4h), which was indexed for an fcc structure with a ¼ 0.5814 nm (see Appendix A). Such a structure with a very close lattice parameter was earlier observed in AleCueCr [19] (l-phase). In the following it is designated L in order to prevent confusion with l-Al4Mn. The Lephase was found in samples annealed at 550 and 600 C but not at 650 C. In the DTA plot of a sample, where the L-phase coexisted with D3 (Fig. 3f), the transformation temperatures 631 C and 763 C were associated with the incongruent melting of the former and the latter, respectively. 3.3. Isothermal sections Partial isothermal sections were constructed for 910, 850, 750, 650, 600 and 550 C. A temperature of 910 C was selected in order to include in this section both binary AleMn T and m. A wide stripe adjacent to AleCu is occupied by the liquid region (Fig. 6a). The T-phase extends up to ~6 at.% Cu and is in equilibrium with the liquid and the g-region. The equilibrium of the T-phase with the n-phase is shown provisionally. The g-region is provisionally divided into the subregions g1 and g2, since the transition temperatures revealed by DTA between these phases in the studied samples are below the annealing temperature. The g-region extends up to ~17 at.% Cu and is in equilibrium with the b-phase, extending up to at least 25 at.% Cu. At 850 C (Fig. 6b) the constitution of the studied compositional region is similar to that at 910 C apart from the further propagation of the T-phase region towards higher Cu concentrations and its
Fig. 7. Partial isothermal sections of AleCueMn at: a) 600 C, b) 550 C. The measured compositions of the alloys are marked by solid squares, and those of phases by open squares. Suggested three-phase equilibria are marked by broken lines. L is liquid.
B. Grushko, S.B. Mi / Journal of Alloys and Compounds 688 (2016) 957e963
complete disconnection from the AleMn terminal. The region very close to AleMn was not studied. The three-phase equilibria T þ n þ m and T þ n þ g are plausible considering the observation of the T þ n equilibrium revealed at 750 C. At 750 C (Fig. 6c) the regions of the b-phase and the T-phase were found to propagate further towards higher Cu. The ternary Rphase and D3-phase were already revealed. In Ref. [13] the partial 700 C isothermal section was constructed in order to be comparable with that in Ref. [10]. It was not completed and not repeated in the present contribution as qualitatively similar to the 750 C section. Instead, the 650 C section was studied in more detail (Fig. 6d). At this temperature the regions of the R-phase and D3-phase are somewhat wider than at 750 C. The region of the D3-phase was found to visibly shift towards higher-Cu compositions. The ternary U-phase is formed. At 600 and 550 C only a small stripe close to the AleCu terminal was studied (see Fig. 7) due to difficulty in the equilibration of alloys with higher-Mn concentrations. In addition to ternary R, D3 and U, also the ternary L-phase is formed. At 550 C the liquid region is very small around the AleCu eutectic point. 4. Conclusions The AleCueMn alloy system was studied above 45 at.% Al. The so-called b-phase of the CsCl-type structure (a z 0.292e0.298 nm) exists in a wide ternary region extending along ~ 50 at.% Al. The total g1/g2-region extends from AleMn up to ~17 at.% Cu. The high-temperature T-Al3Mn phase (Pnma, a z 1.48, b z 1.24, c z 1.25 nm) extends up to ~15 at.% Cu. The so-called R-phase (Bbmm, a z 2.41, b z 1.25, c z 0.76 nm) exists in a compositional region of Al74-80Cu5-12.5Mn12.5-18. The ternary so-called U-phase exists below 697 C in a compositional region of Al55-58Cu29-37Mn7.5-14. The decagonal D3-phase is stable in a compositional region of Al61.5-68.5Cu19-29.5Mn9-16. Below 631 C an fcc L-phase (a ¼ 0.5814 nm) exists around ~ Al60Cu36.5Mn3.5. Acknowledgements The authors thank C. Thomas for technical contributions and D. Pavlyuchkov for helpful discussions.
963
Appendix A Diffraction data of the L-phase (Al60Cu36.5Mn3.5, suggested Fm3m, a ¼ 0.5814(1) nm). The indexing is for the best fit (aver. D(2q) ¼ 0.011, max. D(2q) ¼ 0.019 , FOM(9) ¼ 84.0).
No.
h
k
l
d, nm
I
1 2 3 4 5 6 7 8 9
1 2 2 3 2 4 3 4 4
1 0 2 1 2 0 3 2 2
1 0 0 1 2 0 1 0 2
0.33546 0.29064 0.20558 0.17536 0.16787 0.14538 0.13339 0.13000 0.11866
45 7 100 20 2 19 10 2 40
References [1] H.L. Lukas, Aluminium - copper - manganese, in MSIT Workplace (online version), G. Effenberg (Ed.), MSI, Stuttgart. [2] W. Krings, W. Ostmann, Z. Anorg, Allg. Chem. 163 (1927) 145 (in German). [3] N. Ponweiser, C.L. Lengauer, K.W. Richter, Intermetallics 19 (2011) 1737. [4] A.J. McAlister, J.L. Murray, Bull. Alloy Phase Diagr. 8 (1987) 438. [5] H.-G. Petri, Alum. Arch. 14 (1938) 5 (in German). [6] W. Guertler, G. Rassmann, Metallwirtschaft 22 (1943) 65 (in German). [7] G.V. Raynor, J. Inst. Met. 70 (1944) 531. [8] K. Robinson, Philos. Mag. 43 (1952) 775. [9] a) M.A. Taylor, Acta Met. 8 (1960) 256; b) M.A. Taylor, Acta Cryst. 14 (1961) 84. € ster, T. Go €decke, Z. Met. 57 (1966) 889 (in German). [10] W. Ko € decke, W. Ko €ster, Z. Met. 62 (1971) 727 (in German). [11] T. Go [12] J.L. Murray, A.J. McAlister, R.J. Schaefer, L.A. Bendersky, F.S. Biancaniello, D.L. Moffatt, Metall. Trans. A 18 (1987) 385. [13] B. Grushko, D. Pavlyuchkov, S.B. Mi, S. Balanetskyy, J. Alloys Comp. 677 (2016) 148. [14] S. Balanetskyy, G. Meisterernst, M. Heggen, M. Feuerbacher, Intermetallics 16 (2008) 71. [15] A.P. Tsai, A. Inoue, T. Matsumoto, J. Mater. Sci. Lett. 7 (1988) 322. ^mar, M. Harmelin, J. Non-Cryst. Solids 153&154 (1993) 494. [16] S. Maa [17] R.J. Schaefer, F.S. Biancaniello, J.W. Cahn, Scr. Metall. 20 (1986) 1439. [18] S. Balanetskyy, D. Pavlyuchkov, T. Velikanova, B. Grushko, J. Alloys Comp. 619 (2015) 211. rzyn ski, D. Pavlyuchkov, S. Mi, E. Kowalska-Strze˛ ciwilk, [19] B. Grushko, B. Przepio M. Surowiec, J. Alloys Comp. 442 (2007) 114. [20] G. Kreiner, H.F. Franzen, J. Alloys Comp. 261 (1997) 83. [21] A. Kontio, E.D. Stevens, P. Coppens, Acta Cryst. B36 (1980) 435. [22] B. Grushko, G.R. Stafford, Scr. Met. 23 (1989) 557. [23] M. Ellner, Metall. Trans. A 21 (1990) 1669.
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Al-rich region of AleCueMn B. Grushko a, b, *, S.B. Mi c a
MaTecK, D-52428 Jülich, Germany PGI-5, Forschungszentrum Jülich, D-52425 Jülich, Germany c State Key Laboratory for Mechanical Behavior of Materials, Xi’an Jiaotong University, Xi’an 710049, PR China b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 25 April 2016 Received in revised form 1 July 2016 Accepted 6 July 2016 Available online 7 July 2016
The AleCueMn alloy system was studied above 45 at.% Al between 550 and 910 C by scanning electron microscopy, powder X-ray diffraction and differential thermal analysis. A wide ternary region of the socalled b-phase extending along ~50 at.% Al was confirmed. It has a CsCl-type structure with a z 0.292 e0.298 nm. The total g1/g2-region extends from AleMn up to ~17 at.% Cu and the high-temperature TAl3Mn phase (Pnma, a z 1.48, b z 1.24, c z 1.25 nm) extends up to ~15 at.% Cu. The so-called R-phase (Bbmm, a z 2.41, b z 1.25, c z 0.76 nm) was found to exist in a compositional region of Al74-80Cu512.5Mn12.5-18. The ternary phase earlier reported at Al57.9Cu26.3Mn15.8 was confirmed. It exists below 697 C in a compositional region of Al55-58Cu29-37Mn7.5-14. The decagonal D3-phase was concluded to be stable in a compositional region of Al61.5-68.5Cu19-29.5Mn9-16. Below 631 C an fcc phase (a ¼ 0.5814 nm) was revealed around ~ Al60Cu36.5Mn3.5. Partial isothermal sections at 550, 600, 650, 750, 850 and 910 C were constructed. © 2016 Elsevier B.V. All rights reserved.
Keywords: Transition metal alloys and compounds Phase diagrams
1. Introduction The ternary AleCueMn alloy system has been extensively studied for several decades. The corresponding literature up to 2003 was reviewed in Ref. [1]. Since the first publication on the ternary constitution of this system in 1927 [2], the boundary binary alloy systems has been updated and modified several times, which also required subsequent revisions of the ternary phase diagram in order to bring them into agreement. For the following description the relevant parts of the updated phase diagrams AleCu [3] and AleMn [4] are shown in Fig. 1. The crystallographic data of the phases are provided in Table 1. Of the earlier reports, those in Refs. [5e11] are most relevant to the present work. In contrast to Ref. [2], where no ternary compounds were revealed, two complex orthorhombic structures were reported in Ref. [5] on the basis of the X-ray diffraction study of needle-shaped single crystals. The phase designated T in Ref. [5] (TCu in the following1) with a ¼ 0.769, b ¼ 2.406 and
* Corresponding author. PGI-5, Forschungszentrum Jülich, D-52425 Jülich, Germany. E-mail address: [email protected] (B. Grushko). 1 This is in order to discriminate it from the binary T-Al3Mn phase (or just Tphase). http://dx.doi.org/10.1016/j.jallcom.2016.07.075 0925-8388/© 2016 Elsevier B.V. All rights reserved.
c ¼ 1.248 nm was found to contain 13e15 wt% Cu and 20e25 wt% Mn (i.e. between ~ Al80.1Cu7.8Mn12.1 and ~Al77.7Cu6.9Mn15.4). The other phase designated Y was found to have a ¼ 1.479, b ¼ 1.260, c ¼ 1.243 nm. It was observed in a range of compositions adjacent to AleMn. Particularly, a composition of 5.9 wt% Cu and 32.5 wt% Mn (~Al76.9Cu3.1Mn20.0) was reported. The c-axes of both TCu and Y were selected along the needle axes. In Ref. [6] the existence of the former was confirmed, but no information on the latter was supplied. Instead, a new phase was reported in Ref. [6] at Al60Cu25Mn15 (designated Al11Cu5Mn3, such a different composition did not allow it to be associated with the Y-phase). Indeed, in Ref. [10] this Al11Cu5Mn3 was reported to be orthorhombic with a ¼ 1.210, b ¼ 2.408, c ¼ 1.921 nm. In Ref. [7] the ternary phases of Ref. [5] were only mentioned as being under investigation. The most detailed experimental investigation of the whole ternary AleCueMn phase diagram was reported in Ref. [10]. It was basically accepted in the assessment [1] and only slightly modified in order to be more compatible with the later updates of the boundary binary AleCu and AleMn phase diagrams mentioned above. Particularly, in the boundary AleCu phase diagram used in Ref. [10], the g0, g1 and d phases accepted to date were not distinguished inside a region designated G. Similarly, b0, b, ε1 and ε2 were not distinguished inside the b-region, and h1 and h2 inside the h-region, while the phases a2, z1 and z2 were ignored. As to the
958
B. Grushko, S.B. Mi / Journal of Alloys and Compounds 688 (2016) 957e963
Fig. 1. Relevant parts of the binary phase diagrams of: a) AleCu [3] and b) AleMn [4].
Table 1 Crystallographic data of the relevant phases in AleCueMn. TW means this work. Phase
Space group
Lattice parameters a, nm
q-Al2Cu h1 h2 ε2 Al6Mn l-Al4Mn m-Al4Mn n-Al11Mn4
I4/mcm Pban or Cmmm C2/m P63/mmc Cmcm P63/m P63/mmc
T- Al3Mn (HT- Al11Mn4)
P1 Pnma
g1-Al8Mn5 g2-Al8Mn5
I43m R3m
d-AlMn b
Pm3m
R
U L
Im3m
Bbmm Orth. FCC
0.6063 0.4087 1.2066 0.4146 0.75551 2.8382 2.0015 0.5095
Ref. Comment
a,
b, nm
89.35
1.200 0.4105 e 0.64994 e e 0.8879
b,
c, nm
e
0.4872 0.8635 0.6913 0.5063 0.88724 1.2389 2.4699 0.5051
124.96
100.47
g,
105.08
[3] [3] [3], standardized [3] [4] [20] [14], Al79.8Mn20.2 [21]
1.4873 1.46256 0.9012
1.2420 1.24557 e
1.2547 1.25469 e
[14], Al72.8Mn277.2 [13], Al70.0Cu11.5Mn18.5 [22]
1.2611 1.26403 (23) 0.3063
e e e
0.7927 0.78272 (9) e
[23], Al55Mn45a TW, Al56Cu15Mn29 [23], Al55Mn45, at 957 C
e
e
1.24787 2.408 e
0.76102 1.921 e
TW, Al49.5Cu22Mn28.5 TW, Al48Cu50Mn2 [13] Al80Cu8Mn12 [10] TW
0.29804 (6) 0.29210 (4) 2.41023 1.210 0.5814 (1)
a
The g2estructure can also be described as a distorted g-brass structure with the basic (body-centered) lattice. For Al57Mn43 the corresponding lattice parameters are a z 0.896 nm and a z 89.02 [52]. In this notation, the g2estructure contains the same 52 atoms in the unit cell as the cubic g1estructure.
boundary AleMn phase diagram, updated somewhat later by the same authors in Ref. [11], MnAl(h) and Mn5Al8(h) were still not distinguished in Ref. [10], as well as Mn4Al11(r) and Mn4Al11(h) (the latter is also known as T-Al3Mn). Much later the l-Al4Mn phase was added to the AleMn phase diagram besides m-Al4Mn [12]. The Al-rich part of the ternary diagram assessed in Ref. [1] contains two complex intermetallic compounds t1 and t2 forming in narrow compositional regions and a wide region of the so-called
b solid solution of the W-type structure related to a corresponding cubic modification of Mn, to AleCu b and to high-temperature AlMn. The ternary t1-phase was associated with the abovementioned TCu-phase of Ref. [5], but placed at considerably lower Al concentrations. It was concluded that t1 was formed by the reaction L þ Al8Mn5 4 t1 at 1020 C. The ternary t2-phase forming by the reaction b þ t1 4 t2 at ~700 C was associated with Al11Cu5Mn3 of Ref. [6]. The Y-phase of Ref. [5] was just mentioned as not
B. Grushko, S.B. Mi / Journal of Alloys and Compounds 688 (2016) 957e963
confirmed. On the other hand, as was shown in Ref. [13], the ternary t1phase of Ref. [1] exhibited a structure of the above-mentioned Yphase of Ref. [5], which is actually the ternary extension of binary TAl3Mn. In other words, the t1-phase was associated with the wrong one of the two structures revealed in Refs. [5], which also explains such a significant difference in the compositions reported in Refs. [5] and [10]. The other orthorhombic phase of Ref. [5] was found to be formed at somewhat higher Al concentrations than this ternary extension and quite close to the compositions determined in Ref. [5] (see Fig. 5 of Ref. [13]). Similarly to that in AlePdeMn ([14] and references therein), the later phase was designated R. After the discovery of icosahedral (I) and decagonal (D) quasicrystals in AleMn, many efforts were made to stabilize these metastable structures by the addition of a third element, particularly also by Cu. The I-phase, forming in rapidly solidified Al65Cu20Mn15 and several other alloys, was found to transform rapidly by moderate annealing [15,16]. The final constituencies in three samples studied in Ref. [16] (black scuares in Fig. 2) were concluded to be either the above-mentioned t1, or t2, or t1 þ t2 of Refs. [1,10] without specifying the structures of these phases. A D-phase was observed in the intermediate stages of this transformation and together with the I-phase was concluded to be metastable in AleCueMn. However, this conclusion was premature, as was revealed in Ref. [13] and will be described in more detail below. In the present contribution the Al-rich part of the AleCueMn phase diagram was revised above 45 at.% Al at 550e910 C. The above-mentioned updated boundary binary AleMn and AleCu phase diagrams were accepted. 2. Experimental The alloys were produced by levitation induction melting in a water-cooled copper crucible under a pure Ar atmosphere. The purity of Al was 99.999%, of Cu and Mn 99.95%. The ingot weights
Fig. 2. Overall compositions of the phases in AleCueMn revealed in the present study. The total T-phase region above 650 C is shown by a brighter fill, that at 850 C by a darker fill and at 910 C surrounded by solid line, the upper-Cu limit of this region at 750 C is shown by a broken line (see isothermal sections below). The compositions of the samples studied in Ref. [16] are marked by squares. The broken lines continuing b (B2) towards d (A2) and the lower-Al boundary of b are drawn provisionally.
959
were typically ~5 g. Parts of the ingots were thermally annealed at 910 C for 24 h, at 850 C for 67 h, at 750 C for 148e332 h, at 650 C for 163e310 h, at 600 C for 330e404 h, and at 550 C for 206e1460 h. The alloys were examined by powder X-ray diffraction (XRD), scanning electron microscopy (SEM), and differential thermal analysis (DTA). The local phase compositions were determined in SEM by energy-dispersive X-ray analysis (EDX) on polished unetched cross-sections. Powder XRD was carried out in the transmission mode using Cu Ka1 radiation and an image plate detector. DTA was carried out in alumina crucibles at heating rates of 25 K/min and cooling rates of 50 K/min.
3. Results and discussion 3.1. Binary phases and their ternary extensions In general, the AleCu intermediate phases are formed at lower Al concentrations and lower temperatures than the AleMn phases. In the studied compositional and temperature region the relevant AleCu phases are q, high-temperature h1, low-temperature h2, and high-temperature ε2 (see Fig. 1a and Table 1). All of them were found to dissolve very little Mn. The Al-rich part of the AleMn phase diagram in Fig. 1b contains the equilibrium phases Al6Mn, l-Al4Mn, m-Al4Mn, the T-Al3Mn phase (also known as high-temperature Al11Mn4) and the Al11Mn4 (n-phase). The lowest-Mn Al12Mn phase is only stable below 512 C [17], i.e. below the studied temperatures and was not observed in the present study. The central part of the AleMn phase diagram contains high-temperature Al8Mn5 (g1), low-temperature Al8Mn5 (g2) and a high-temperature AlMn phase designated at present d due to its structural similarity to d-Mn (g-phase in Ref. [4]). The accepted AleMn phase diagram was additionally argued in recent Ref. [18]. Of the AleMn binaries, Al6Mn and l-Al4Mn did not exhibit visible solubility of Cu, while m-Al4Mn and n-Al11Mn4 were found to dissolve up to ~2 at.% Cu. On the other hand, the phases T, g1 and g2 were observed at essentially ternary compositions (Fig. 2). Similarly to the case at binary compositions, the Cu-containing high-temperature g1-phase could not be conserved by water quenching and transformed to g2. The existence of g1 at elevated temperatures was detected by DTA. For example, the thermal effect at 734 C in Fig. 3a indicates the transformation of g2 to g1. According to SEM/EDX examination, a sample of this composition annealed at and quenched from 850 C, i.e. above the transition temperature, consisted of essentially a single phase of Al56Cu15Mn29, whose powder XRD pattern indicated the g2 structure (Fig. 4a, Table 1). With increasing Cu concentration the upper existence temperatures of g1 and g2 were found to decrease and the temperature of the corresponding transition decreased accordingly (compare to the lowest transition temperature 957 C in binary AleMn, Fig. 1b). The total g1/g2 region extends up to at least 17 at.% Cu (see Fig. 2). The ternary region of T-Al3Mn was found to propagate up to 15 at.% Cu, extending more with decreasing temperatures (see Fig. 2). The addition of Cu also decreases the lower existence limit of the T-phase, so below ~895 C it is only stable at ternary compositions (see the corresponding isothermal sections). At 5 at.% Cu it was observed down to 650 C and at the rates applied in our DTA experiments its decomposition was not detected. DTA plots in Fig. 3b and c demonstrate the decrease of the upper temperature limits of the T-phase containing 5 and 11.5 at.% Cu, respectively. In addition to the major T-phase the former sample also contained some n-phase, whose transformation is visible in Fig. 3b starting slightly above 750 C (shallow effect marked by an arrow).
960
B. Grushko, S.B. Mi / Journal of Alloys and Compounds 688 (2016) 957e963
a)
734
b)
Heat flow, a.u.
962
c) 697
815
d) 544
572
660
e) 763
f) 631
Exo 400
500
600
700
800
900
1000
1100
T, °C Fig. 3. DTA plots (heating at 25 K/min) of: a) Al56Cu15Mn29 annealed for 67 h at and quenched from 850 C, b) Al70Cu5Mn25 annealed for 310 h at 650 C, c) Al70Cu11.5Mn18.5 annealed for 332 h at 750 C, d) Al57.5Cu29Mn13.5 annealed for 1460 h at 550 C, e) Al75.3Cu10Mn14.7 annealed for 163 h at 650 C, f) Al65.1Cu24.6Mn10.3 annealed for 404 h at 600 C.
3.2. Ternary phases The formation of a phase reported in Ref. [10] in a wide ternary region below ~50 at.% Al was confirmed in the present study. In a slightly modified version of Ref. [1] this so-called b-phase (Im3m, a z 0.29e0.30 nm) is actually a ternary extension of AlMn (d-phase in Fig 1b) and the isostructural phases in AleCu (outside the region in Fig. 1a) ranging in total between the AleCu and AleMn binary terminals. In both AleCu and AleMn the bcc phases are only stable at elevated temperatures (for example, d-phase above 840 C, see Fig. 1b). In the present work only alloys revealing the Al-rich boundary of the b-region were studied. At these compositions the b-phase was
observed down to 600 C, i.e. below the stability temperatures of the d-phase, and a CsCl-type ordering was revealed in the studied alloys. In a typical powder XRD pattern in Fig. 4b this is indicated by the presence of the (100) reflection, which is forbidden in a bcc structure. The ordering was observed in the whole compositional region depicted in Fig. 2 by the grey b-field and at all studied temperatures. In Ref. [1] the CsCl-type ordering inside the b-region was only mentioned around the AlCu2Mn composition as an alternative to the ordering typical of the Heusler phases (i.e. cF16 vs. cP2). On the other hand, similar to AleFeeMn [18], a link of this region with the region of the d-phase is indeed plausible. With decreasing temperature the Al-rich boundary of the b-phase was found to propagate towards higher Cu and at 650 C it achieved the Al48Cu50Mn2 composition. A phase, first reported in Ref. [6] under the name Al11Cu5Mn3, specified and included in the phase diagram of Ref. [10] (t2 in Ref. [1]), was revealed in the present work in a compositional region of Al55-58Cu29-37Mn7.5-14. Its powder XRD pattern (Fig. 4c) was qualitatively similar to that simulated from Table 1 of Ref. [10] where the intensities of the reflections were provided qualitatively. No additional structural investigation of this phase, designated in the following U, was carried out. It was found in samples annealed at 550e650 C, but not at 750 C. In the DTA plot of Al57.5Cu29Mn13.5 in Fig. 3d a transformation of the essentially single U ephase started at ~645 C with a pronounced effect at 697 C, i.e. at a temperature close to that determined in Ref. [10]. The observation of a temperature range of this transformation is in agreement with the existence of a compositional region of U revealed by SEM/EDX. Similarly to Ref. [10], the U -phase was found to be transformed to a mixture of two phases, one of which was identified with the b-phase. The second transformation product of the Uphase had a composition of ~Al64Cu20Mn16, which is quite far away from that of the T-phase claimed in Ref. [10], but very close to a composition where the metastable decagonal phase was observed in Ref. [16]. SEM/EDX examination of a sample with such a nominal composition indicated that the material was solid after annealing at 750 C for 238 h and subsequent water quenching. Apart from the major phase of ~Al64.2Cu20.3Mn15.5 it contained a little of a second phase of ~Al66.5Cu15.4Mn18.1. From these compositions the latter was rather typical of the T-phase,
Fig. 4. Powder XRD patterns of: a) g2 in Al56Cu15Mn29 annealed for 67 h at and quenched from 850 C, b) b in Al49.5Cu22Mn28.5 annealed for 145 h at 650 C (the (100) reflection is marked by an arrow), c) U in Al57.5Cu29Mn13.5 annealed for 1460 h at 550 C, d) D3 of Al64Cu20Mn16 annealed for 238 h at 750 C, e) T in Al70Cu11.5Mn18.5 annealed for 1460 h at 550 C, f) R in Al80Cu8Mn12 annealed for 1460 h at 550 C, g) m, powdered single crystal (Al79.8Mn20.2), h) L in Al61Cu35.5Mn3.5 annealed for 1460 h at 550 C.
Heat flow, a.u.
B. Grushko, S.B. Mi / Journal of Alloys and Compounds 688 (2016) 957e963
774 717
763
Exo 600
700
800
900
1000
T, °C Fig. 5. DTA plot of Al65Cu18.5Mn16.5 pre-annealed at 750 C, held in the DTA device for one hour at 600 C, cooled down to 500 C (not shown) and afterwards heated to 950 C at 10 K/min. A similar result was obtained by a direct heating at 25 K/min.
961
indicating the formation of an additional phase. The powder XRD pattern (Fig. 4d) of this sample was indeed different from that of the T-phase (Fig. 4e), but rather resembled the XRD pattern of the stable AleFeeMn decagonal D3-phase [18], as was already reported in Ref. [13]. Electron diffraction examinations in Ref. [13] confirmed the D3-structure. Additional examinations revealed that the compositional region of this phase shifts with decreasing temperature towards lower Mn, so at 650 C the abovementioned Al64Cu20Mn16 composition corresponds to the twophase region TþU, but at this temperature the D3-phase is formed around ~ Al65Cu23Mn12. As follows from our annealing and DTA experiments, the D3structure is stable in AleCueMn. The experimental results of Ref. [16], leading its authors to erroneous conclusions of the total
Fig. 6. Partial isothermal sections of AleCueMn at: a) 910 C, b) 850 C, c) 750 C, d) 650 C. The measured compositions of the alloys are marked by solid squares, and those of phases by open squares. Suggested three-phase equilibria are marked by broken lines. The broken lines separating the regions of g1 and g2 are drawn provisionally based on DTA examinations. L is liquid.
962
B. Grushko, S.B. Mi / Journal of Alloys and Compounds 688 (2016) 957e963
instability of the decagonal phase in AleCueMn, were incorrectly interpreted. The D3-phase produced in Ref. [16] indeed transformed by moderate heating but could recombine by heating at somewhat higher temperature. This can be demonstrated by an experiment with a sample of Al65Cu18.5Mn16.5 annealed at and water quenched from 750 C. In the subsequent DTA examination it was first held in the DTA device for one hour at 600 C, cooled down to 500 C and afterwards heated to 950 C at 10 K/min. After quenching from 750 C, the sample consisted of the D3-phase, which decomposed by the first heating to and annealing at 600 C. The plot in Fig. 5 was recorded during the second heating from 500 to 950 C. The first weak minimum at ~717 C corresponds to the recombination of the D3-phase. The transformation of the D3phase starts at ~763 C and above ~774 C this phase does not exist anymore, i.e. at this composition, close to that examined in Refs. [16], it is stable in a quite narrow temperature range. Such a possibility was not tested in Ref. [16], and this study only was on single compositions. The ternary R-phase was found to be formed in a region of Al74-80Cu5-12.5Mn12.5-18 close to the upper-Al limits of the ternary extension of the T-phase (see Fig. 2). With decreasing temperature its compositions exhibited a shift towards higher Al, so that the total regions of the R-phase and T-phase partially overlap. The coexistence of these two phases was not identified metallographically due to probably small compositional differences and also due to their formation in the form of fine weakly connected needles. Their coexistence was also not revealed by powder XRD. Both these phases have complex diffraction patterns with numerous overlaps (see Fig. 4e and f). In Fig. 2 and the isothermal sections (see below) the separation between the T-region and R-region is shown conditionally. On the other hand, the coexistence of R and m was observed by SEM/EDX [13]. The powder XRD pattern of the m-phase (see Fig. 4g) is also complex and contains lines overlapping with those of the T-phase and Rphase. The R-phase was already observed at 750 C but not at 850 C and its solidus temperature decreases with increasing Cu
concentration. In the DTA plot in Fig. 3e an endothermal effect at ~660 C could be associated with the start of the incongruent melting of the R-phase resulting in the formation of the T-phase, which continues to melt at higher temperatures. At lower heating rates the DTA plots were smeared even more, which did not allow the reaction effects to be separated. Apart from the above-mentioned phases, a new compound was revealed in a small region around ~ Al60Cu36.5Mn3.5. This phase exhibited a quite simple powder XRD pattern (Fig. 4h), which was indexed for an fcc structure with a ¼ 0.5814 nm (see Appendix A). Such a structure with a very close lattice parameter was earlier observed in AleCueCr [19] (l-phase). In the following it is designated L in order to prevent confusion with l-Al4Mn. The Lephase was found in samples annealed at 550 and 600 C but not at 650 C. In the DTA plot of a sample, where the L-phase coexisted with D3 (Fig. 3f), the transformation temperatures 631 C and 763 C were associated with the incongruent melting of the former and the latter, respectively. 3.3. Isothermal sections Partial isothermal sections were constructed for 910, 850, 750, 650, 600 and 550 C. A temperature of 910 C was selected in order to include in this section both binary AleMn T and m. A wide stripe adjacent to AleCu is occupied by the liquid region (Fig. 6a). The T-phase extends up to ~6 at.% Cu and is in equilibrium with the liquid and the g-region. The equilibrium of the T-phase with the n-phase is shown provisionally. The g-region is provisionally divided into the subregions g1 and g2, since the transition temperatures revealed by DTA between these phases in the studied samples are below the annealing temperature. The g-region extends up to ~17 at.% Cu and is in equilibrium with the b-phase, extending up to at least 25 at.% Cu. At 850 C (Fig. 6b) the constitution of the studied compositional region is similar to that at 910 C apart from the further propagation of the T-phase region towards higher Cu concentrations and its
Fig. 7. Partial isothermal sections of AleCueMn at: a) 600 C, b) 550 C. The measured compositions of the alloys are marked by solid squares, and those of phases by open squares. Suggested three-phase equilibria are marked by broken lines. L is liquid.
B. Grushko, S.B. Mi / Journal of Alloys and Compounds 688 (2016) 957e963
complete disconnection from the AleMn terminal. The region very close to AleMn was not studied. The three-phase equilibria T þ n þ m and T þ n þ g are plausible considering the observation of the T þ n equilibrium revealed at 750 C. At 750 C (Fig. 6c) the regions of the b-phase and the T-phase were found to propagate further towards higher Cu. The ternary Rphase and D3-phase were already revealed. In Ref. [13] the partial 700 C isothermal section was constructed in order to be comparable with that in Ref. [10]. It was not completed and not repeated in the present contribution as qualitatively similar to the 750 C section. Instead, the 650 C section was studied in more detail (Fig. 6d). At this temperature the regions of the R-phase and D3-phase are somewhat wider than at 750 C. The region of the D3-phase was found to visibly shift towards higher-Cu compositions. The ternary U-phase is formed. At 600 and 550 C only a small stripe close to the AleCu terminal was studied (see Fig. 7) due to difficulty in the equilibration of alloys with higher-Mn concentrations. In addition to ternary R, D3 and U, also the ternary L-phase is formed. At 550 C the liquid region is very small around the AleCu eutectic point. 4. Conclusions The AleCueMn alloy system was studied above 45 at.% Al. The so-called b-phase of the CsCl-type structure (a z 0.292e0.298 nm) exists in a wide ternary region extending along ~ 50 at.% Al. The total g1/g2-region extends from AleMn up to ~17 at.% Cu. The high-temperature T-Al3Mn phase (Pnma, a z 1.48, b z 1.24, c z 1.25 nm) extends up to ~15 at.% Cu. The so-called R-phase (Bbmm, a z 2.41, b z 1.25, c z 0.76 nm) exists in a compositional region of Al74-80Cu5-12.5Mn12.5-18. The ternary so-called U-phase exists below 697 C in a compositional region of Al55-58Cu29-37Mn7.5-14. The decagonal D3-phase is stable in a compositional region of Al61.5-68.5Cu19-29.5Mn9-16. Below 631 C an fcc L-phase (a ¼ 0.5814 nm) exists around ~ Al60Cu36.5Mn3.5. Acknowledgements The authors thank C. Thomas for technical contributions and D. Pavlyuchkov for helpful discussions.
963
Appendix A Diffraction data of the L-phase (Al60Cu36.5Mn3.5, suggested Fm3m, a ¼ 0.5814(1) nm). The indexing is for the best fit (aver. D(2q) ¼ 0.011, max. D(2q) ¼ 0.019 , FOM(9) ¼ 84.0).
No.
h
k
l
d, nm
I
1 2 3 4 5 6 7 8 9
1 2 2 3 2 4 3 4 4
1 0 2 1 2 0 3 2 2
1 0 0 1 2 0 1 0 2
0.33546 0.29064 0.20558 0.17536 0.16787 0.14538 0.13339 0.13000 0.11866
45 7 100 20 2 19 10 2 40
References [1] H.L. Lukas, Aluminium - copper - manganese, in MSIT Workplace (online version), G. Effenberg (Ed.), MSI, Stuttgart. [2] W. Krings, W. Ostmann, Z. Anorg, Allg. Chem. 163 (1927) 145 (in German). [3] N. Ponweiser, C.L. Lengauer, K.W. Richter, Intermetallics 19 (2011) 1737. [4] A.J. McAlister, J.L. Murray, Bull. Alloy Phase Diagr. 8 (1987) 438. [5] H.-G. Petri, Alum. Arch. 14 (1938) 5 (in German). [6] W. Guertler, G. Rassmann, Metallwirtschaft 22 (1943) 65 (in German). [7] G.V. Raynor, J. Inst. Met. 70 (1944) 531. [8] K. Robinson, Philos. Mag. 43 (1952) 775. [9] a) M.A. Taylor, Acta Met. 8 (1960) 256; b) M.A. Taylor, Acta Cryst. 14 (1961) 84. € ster, T. Go €decke, Z. Met. 57 (1966) 889 (in German). [10] W. Ko € decke, W. Ko €ster, Z. Met. 62 (1971) 727 (in German). [11] T. Go [12] J.L. Murray, A.J. McAlister, R.J. Schaefer, L.A. Bendersky, F.S. Biancaniello, D.L. Moffatt, Metall. Trans. A 18 (1987) 385. [13] B. Grushko, D. Pavlyuchkov, S.B. Mi, S. Balanetskyy, J. Alloys Comp. 677 (2016) 148. [14] S. Balanetskyy, G. Meisterernst, M. Heggen, M. Feuerbacher, Intermetallics 16 (2008) 71. [15] A.P. Tsai, A. Inoue, T. Matsumoto, J. Mater. Sci. Lett. 7 (1988) 322. ^mar, M. Harmelin, J. Non-Cryst. Solids 153&154 (1993) 494. [16] S. Maa [17] R.J. Schaefer, F.S. Biancaniello, J.W. Cahn, Scr. Metall. 20 (1986) 1439. [18] S. Balanetskyy, D. Pavlyuchkov, T. Velikanova, B. Grushko, J. Alloys Comp. 619 (2015) 211. rzyn ski, D. Pavlyuchkov, S. Mi, E. Kowalska-Strze˛ ciwilk, [19] B. Grushko, B. Przepio M. Surowiec, J. Alloys Comp. 442 (2007) 114. [20] G. Kreiner, H.F. Franzen, J. Alloys Comp. 261 (1997) 83. [21] A. Kontio, E.D. Stevens, P. Coppens, Acta Cryst. B36 (1980) 435. [22] B. Grushko, G.R. Stafford, Scr. Met. 23 (1989) 557. [23] M. Ellner, Metall. Trans. A 21 (1990) 1669.