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Investigation of the Al–Cu–Rh phase diagram in the vicinity of the decagonal phase
Journal of Alloys and Compounds 305 (2000) 219–224
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Investigation of the Al–Cu–Rh phase diagram in the vicinity of the decagonal phase ´ ´ ´ 1 , M. Yurechko B. Grushko*, J. Gwozdz ¨ Festkorperforschung ¨ ¨ ¨ , Forschungszentrum Julich GmbH, D-52425 Julich , Germany Institut f ur Received 2 January 2000; accepted 1 February 2000
Abstract Phase equilibria were studied in Al-rich Al–Rh and Al–Cu–Rh alloys. We present a partial Al–Rh phase diagram. In addition to previously reported Al–Rh phases, two complicated orthorhombic compounds were found close to the Al 3 Rh composition. Partial 900 and 8008C isothermal sections of the Al–Cu–Rh phase diagram were studied in the vicinity of the decagonal phase. At these temperatures the stable decagonal phase was found to be the only ternary compound. It is in equilibrium with the ternary extensions of three binary Al–Rh phases and at 9008C also with the liquid. 2000 Elsevier Science S.A. All rights reserved. Keywords: Alloy phase diagrams; Intermetallics; Quasicrystals; Al–Rh; Al–Cu–Rh
1. Introduction Quasicrystals have frequently been observed in binary and ternary alloy systems of aluminum with transition metals. In several systems they are thermodynamically stable while in others they can only be produced in metastable states [1,2]. A stable decagonal phase similar to the extensively studied Al–Cu–Co decagonal quasicrystal was reported in Al–Cu–Rh [3]. Both Rh and Co belong to the same column in the periodic table. To date little work has been carried out on the Al–Cu– Rh alloy system. In the past only the extension of cubic AlRh into ternary compositions was investigated [4]. Recently the coexistence of the decagonal phase with a cubic phase and an icosahedral phase was reported in this alloy system in Ref. [5], where only two compositions were studied. Even the binary Al–Rh system is not well known. Several Al–Rh compounds were revealed [6–10] but the corresponding phase diagram has not yet been published. Our knowledge of the Al–Cu–Rh alloy system is important for the study of quasicrystals. In the present
*Corresponding author. 1 On leave from the Institute of Physics and Chemistry of Metals, University of Silesia, Bankowa 12, 40007 Katowice, Poland.
paper we report the investigation of the phase equilibrium involving the Al–Cu–Rh decagonal phase. This required the additional study of binary Al–Rh compounds found in equilibrium with the decagonal phase. Therefore information concerning the binary Al–Rh system is also included.
2. Experimental Alloys of 3–5 g were prepared by inductive melting in a water-cooled copper crucible under an Ar atmosphere. The alloys were thermally annealed at 8008C for 90–1400 h, at 9008C and higher temperatures for 70–120 h and subsequently water quenched. Short heat treatments were carried out in alumina crucibles in a vacuum furnace. For longer annealing times evacuated silica ampoules were used, inside which the alloys were contained in alumina crucibles in order to avoid a reaction of the partially molten samples with the silica. The alloys were studied by powder X-ray diffraction (XRD, CoKa 1 radiation was used), scanning electron microscopy (SEM) and transmission electron microscopy (TEM). The local phase compositions were determined in SEM by energy-dispersive X-ray analysis (EDX) on polished unetched cross-sections. The compositions of
0925-8388 / 00 / $ – see front matter 2000 Elsevier Science S.A. All rights reserved. PII: S0925-8388( 00 )00754-4
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B. Grushko et al. / Journal of Alloys and Compounds 305 (2000) 219 – 224
selected single-phase samples were examined by inductively coupled plasma optical emission spectroscopy (ICPOES). These compositions were used for the corrections of the EDX data. TEM examinations were carried out on powdered materials dispersed on grids with carbon film. Differential thermal analysis (DTA) was carried out for selected samples. Heating and cooling rates of 208C / min were used.
3. Results
3.1. Binary Al–Rh compounds The Al–Rh phase diagram was found to contain a number of phases isostructural to those in Al–Co (for the recent Al–Co phase diagram see Ref. [11]). They are AlRh
[6] of the CsCl-type structure corresponding to AlCo, hexagonal Al 5 Rh 2 [6] corresponding to Al 5 Co 2 and monoclinic Al 9 Rh 2 [7] corresponding to Al 9 Co 2 . In Ref. [6] a high-temperature cubic phase existing around the Al 5 Rh 2 composition was also reported. Its structure was recently determined [9]. In the following, this high-temperature cubic phase is designated C while H is used for the low-temperature hexagonal Al 5 Rh 2 phase. An additional high-temperature monoclinic phase (designated V in the following) was found in [10] at about 70 at.% Al. The typical powder XRD patterns of these phases are shown in Fig. 1; for detailed powder diffraction data of H–Al 5 Rh 2 , C–Al 5 Rh 2 and Al 7 Rh 3 see Ref. [10]. Our powder XRD pattern of Al 9 Rh 2 (Fig. 1a) also fits well with that of Ref. [7]. The lattice parameters of the Al–Rh phases are presented in Table 1. In Ref. [8] an Al 13 Rh 4 phase isostructural to monoclinic Al 13 Co 4 (M–Al 13 Co 4 in [11]) was reported to precipitate
Fig. 1. Powder XRD patterns of: (a) Al 9 Rh 2 , (b) the O 1 -phase, (c) the O 2 -phase, (d) the H–Al 5 Rh 2 phase, (e) the C-phase, (f) the C 1 -phase, (g) the monoclinic V-phase (Al 7 Rh 3 ) and (h) the Al–Cu–Rh decagonal phase; CoKa 1 radiation.
B. Grushko et al. / Journal of Alloys and Compounds 305 (2000) 219 – 224
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Table 1 Al–Rh phases Formula
O1 O2 H C V b a b
Al 9 Rh 2 Al 3 Rh Al 3 Rh Al 5 Rh 2 Al 5 Rh 2 Al 7 Rh 3 AlRh
Space group P21 /a
a a
P63 /mmc Pm3¯ b
Pm3¯ m
Lattice parameters
Ref.
a (nm)
b (nm)
c (nm)
b (8)
1.0149 2.38 2.34 0.7905 0.7680 1.0309 0.2980
0.6290 1.64 1.65 – – 0.3808 –
0.8557 3.28 1.24 0.7861 – 0.6595 –
142.4 – – – – 102.4 –
[7] This work This work [6,10] [9,10] [10] [4]
Orthorhombic. Monoclinic.
by a solid-state reaction from a supersaturated Al–Rh alloy. However, this structure was not observed in our experiments. Instead, two new orthorhombic phases of close compositions were found at about Al 3 Rh. The powder XRD patterns of these phases, designated O 1 and O 2 , are shown in Fig. 1b and c.
The O 2 -phase of a slightly lower Al concentration exhibited a powder XRD pattern similar to that of orthorhombic Al 3 Pd [12], TEM examination also revealed electron diffractograms very similar to those observed in Al–Pd (Fig. 2a–c, see Fig. 9 in [12]). The patterns in Fig. 2a–c can be indexed assuming an orthorhombic unit cell
Fig. 2. Typical electron diffraction patterns of the orthorhombic O 2 -phase: (a) Z.A.5[100], (b) Z.A.5[010], (c) Z.A.5[001]. The latter pattern is also typical of the [001] of the O 1 -phase. The pattern in (d) corresponds to Z.A.5[100] of the O 1 -phase, in (e) to Z.A.5[010] and in (f) to Z.A.5[101] of the O 1 -phase.
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Fig. 3. Partial Al–Rh phase diagram. For the phase designation see Table 1, the liquid is designated L, a is fcc Al.
with a52.34, b51.65 and c51.24 nm, which is close to the corresponding Al–Pd phase 2 . As is mentioned in the literature [13], the Al 3 Pd phase has a structural relation to the decagonal phases. The Al–Rh O 2 -phase was found at 75 at.% Al while according to Ref. [12], Al 3 Pd is formed at the Al concentration which is about 2.5 at.% lower. The O 1 -phase was found at 76 at.% Al. Its electron diffraction pattern (see Fig. 2c–f) can be indexed assuming an orthorhombic unit cell with a52.38, b51.64 and c5 3.28 nm and systematic extinctions of l 5 2n 1 1 for 0kl, h 5 2n 1 1 for hk0, h 1 l 5 2n 1 1 for h0l. It is worth mentioning that c /a¯tg368, the electron diffraction pattern indeed resembles that of the decagonal phase (see below). The lattice parameters of the O 1 -phase and the O 2 -phase are related as a 1 ¯a 2 , b 1 ¯b 2 and c 1 ¯ t 2 c 2 , where t is the golden mean. The available data concerning the phase compositions and transition temperatures are summarized in Fig. 3. All the phases apart from AlRh (also designated b in the following) and the high-temperature C-phase did not show pronounced compositional ranges. The lower temperature limits of the V-phase and C-phase regions have not yet been determined. We have also been unable to clarify the melting temperature of the O 1 -phase as yet. Since in the relevant ternary regions we only observed the ternary extensions of the b-phase, C-phase and O 2 -phase the lack of this information was not crucial for the following.
2
From the more recent work on Al–Pd we suggest the existence of more than one binary phase close to the Al 3 Pd composition. In the current paper we refer to Al 3 Pd as is described in Refs. [12–14] (in [14] this structure was studied in the Al 74 Pd 22 Mn 4 composition and is designated j9-phase). It seems that this basic Al 3 Pd structure can exhibit superstructure ordering. Indeed, inspection of the pattern in Fig. 2a reveals weak additional reflections corresponding to the doubled lattice parameter c of the basic Al 3 Pd structure. In this report we define the lattice parameters for the basic structure. A more detailed study of this phase is in progress.
Fig. 4. Phase equilibrium at 9008C in Al–Cu–Rh. For the phase designation see Table 1, the liquid is designated L. The compositions of the studied alloys are marked by spots.
3.2. Partial isothermal 900 and 8008 C sections of Al– Cu–Rh The partial 9008C section of Al–Cu–Rh is shown in Fig. 4. Three Al–Rh phases exhibited high Cu solubility. The b-phase extends to more than 35 at.% Cu and its Al concentration achieves 55 at.% at high Cu. The C-phase can contain up to about 13 at.% Cu and the O 2 -phase up to about 11 at.% Cu. The Al concentrations of the C-phase and O 2 -phase decrease sharply with the increase of the Cu concentration. The C-phase exhibited a superstructure ordering at its low-Al composition limits. This is in agreement with the observation in Ref. [5] where a cubic structure with a doubled lattice parameter (a51.5380(2) nm, space group of Fm3¯ ) was reported. The corresponding powder diffraction pattern of this phase designated C 1 contains a number of additional reflections (Fig. 1f,e), the strongest of which is at 0.2956 nm (2u 535.238 in Fig. 1f). The powder diffraction pattern calculated using the atomic positions from Ref. [5] was in qualitative agreement with our experimentally observed diffractogram. The existence of subregions of C and C 1 was not studied in detail. Based on the examination of the compositions marked in Fig. 4 the range of the cubic phase is conditionally divided between C and C 1 by a broken line. The Al 9 Rh 2 and H–Al 5 Rh 2 phases were found to contain ,0.5 at.% Cu. The Rh solubility in Al 2 Cu formed in the solidified liquid was ,1 at.%. At 9008C the Al–Cu–Rh decagonal phase (D) is stable in the range of 1–2 at.%, whose center is at about Al 64.5 Cu 16.8 Rh 18.7 . It coexists with the liquid and the ternary extensions of the b-phase, the C-phase (including C 1 ) and the O 2 -phase. The decagonal range in Al–Cu–Rh is smaller than that in Al–Cu–Co at the same 9008C [15] and coincides with the higher-Cu part of the decagonal
B. Grushko et al. / Journal of Alloys and Compounds 305 (2000) 219 – 224
223
Fig. 5. Typical electron diffraction patterns of the Al–Cu–Rh decagonal phase.
range in Al–Cu–Co. The powder XRD pattern of the Al–Cu–Rh decagonal phase (Fig. 1h) and its electron diffraction (Fig. 5) are similar to those of the Cu-rich Al–Cu–Co decagonal phase (Fig. 2b and Fig. 3a–c in [16], respectively). In a wide range of Cu concentrations the C-phase (including C 1 ) is in equilibrium with b, and the O 2 -phase is in equilibrium with the liquid. The regions marked by question marks in Fig. 4 were not studied in detail. The investigated part of the 8008C isothermal section (Fig. 6) is quite similar to that at 9008C. The ranges of the b-phase and O 2 -phase extend to still higher Cu than at 9008C. Zero solubility of Cu in H–Al 5 Rh 2 was found at this temperature. The decagonal phase is in equilibrium with the ternary extensions of the b-phase, the C-phase (including C 1 ) and the O 2 -phase but not with the liquid. Instead, the O 2 -phase was found in equilibrium with b at this temperature. The decagonal region lies around about Al 64.0 Cu 18.0 Rh 18.0 . It is slightly shifted from the decagonal
region at 9008C towards AlCu. This shift is smaller than that observed in Al–Cu–Co [15]. In contrast to Ref. [5] we did not observe any icosahedral phase in Al–Cu–Rh.
4. Conclusions The high-Al range of the Al–Rh alloy system and an adjacent range of Al–Cu–Rh were studied. We report the existence of seven Al–Rh phases: cubic AlRh, monoclinic Al 7 Rh 3 , high-temperature cubic Al 5 Rh 2 and low-temperature hexagonal Al 5 Rh 2 , two orthorhombic phases at about the Al 3 Rh composition and monoclinic Al 9 Rh 2 . At ternary compositions the high-temperature cubic Al 5 Rh 2 phase exhibits a superstructure ordering. We confirm the stability of the Al–Cu–Rh decagonal phase at 800–9008C. This ternary phase is in equilibrium with the ternary extensions of AlRh, cubic Al 5 Rh 2 and one of orthorhombic Al 3 Rh, at 9008C also with the liquid.
Acknowledgements We thank W. Reichert, K. Bickmann and C. Thomas for technical contributions. Financial support from the DFG (project Ur51 / 4-1) is gratefully acknowledged.
References
Fig. 6. Phase equilibrium at 8008C in Al–Cu–Rh. For the phase designation see Table 1, the liquid is designated L. The compositions of the studied alloys are marked by spots.
[1] A.P. Tsai, in: Z. Stadnik (Ed.), Physical Properties of Quasicrystals, Springer, Heidelberg, 1999. [2] B. Grushko, Heter. Chem. Rev. 1 (1994) 255. [3] A.P. Tsai, A. Inoue, T. Masumoto, Mater. Trans. JIM 30 (1989) 666. [4] T. Shishido, H. Takei, J. Less-Common Met. 119 (1986) 75. [5] X.Z. Li, K.-T. Park, K. Sugiyama, K. Hiraga, Metall. Mater. Trans. 29A (1998) 1559.
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[6] L.E. Edshammar, Acta Chem. Scand. 21 (1967) 647. [7] L.E. Edshammar, Acta Chem. Scand. 22 (1968) 2822. [8] Z.A. Chaudhury, C. Suryanarayana, J. Less-Common Met. 91 (1983) 181. [9] Yu. Grin, K. Peters, U. Burkhardt, K. Gotzmann, M. Ellner, Z. Kristallogr. 212 (1997) 439. [10] B. Grushko, M. Yurechko, Z. Kristallogr. 214 (1999) 313. [11] B. Grushko, R. Wittenberg, K. Bickmann, C. Freiburg, J. Alloys Compd. 233 (1996) 279.
[12] B. Grushko, M. Yurechko, N. Tamura, J. Alloys Compd. 290 (1999) 164. [13] Y. Matsuo, K. Hiraga, Phil. Mag. Lett. 70 (1994) 155. [14] M. Boudard, H. Klein, M. de Boissieu, M. Audier, H. Vincent, Phil. Mag. A74 (1996) 939. [15] B. Grushko, Phil. Mag. Lett. 66 (1992) 151. [16] B. Grushko, J. Mater. Res. 8 (1993) 1473.
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www.elsevier.com / locate / jallcom
Investigation of the Al–Cu–Rh phase diagram in the vicinity of the decagonal phase ´ ´ ´ 1 , M. Yurechko B. Grushko*, J. Gwozdz ¨ Festkorperforschung ¨ ¨ ¨ , Forschungszentrum Julich GmbH, D-52425 Julich , Germany Institut f ur Received 2 January 2000; accepted 1 February 2000
Abstract Phase equilibria were studied in Al-rich Al–Rh and Al–Cu–Rh alloys. We present a partial Al–Rh phase diagram. In addition to previously reported Al–Rh phases, two complicated orthorhombic compounds were found close to the Al 3 Rh composition. Partial 900 and 8008C isothermal sections of the Al–Cu–Rh phase diagram were studied in the vicinity of the decagonal phase. At these temperatures the stable decagonal phase was found to be the only ternary compound. It is in equilibrium with the ternary extensions of three binary Al–Rh phases and at 9008C also with the liquid. 2000 Elsevier Science S.A. All rights reserved. Keywords: Alloy phase diagrams; Intermetallics; Quasicrystals; Al–Rh; Al–Cu–Rh
1. Introduction Quasicrystals have frequently been observed in binary and ternary alloy systems of aluminum with transition metals. In several systems they are thermodynamically stable while in others they can only be produced in metastable states [1,2]. A stable decagonal phase similar to the extensively studied Al–Cu–Co decagonal quasicrystal was reported in Al–Cu–Rh [3]. Both Rh and Co belong to the same column in the periodic table. To date little work has been carried out on the Al–Cu– Rh alloy system. In the past only the extension of cubic AlRh into ternary compositions was investigated [4]. Recently the coexistence of the decagonal phase with a cubic phase and an icosahedral phase was reported in this alloy system in Ref. [5], where only two compositions were studied. Even the binary Al–Rh system is not well known. Several Al–Rh compounds were revealed [6–10] but the corresponding phase diagram has not yet been published. Our knowledge of the Al–Cu–Rh alloy system is important for the study of quasicrystals. In the present
*Corresponding author. 1 On leave from the Institute of Physics and Chemistry of Metals, University of Silesia, Bankowa 12, 40007 Katowice, Poland.
paper we report the investigation of the phase equilibrium involving the Al–Cu–Rh decagonal phase. This required the additional study of binary Al–Rh compounds found in equilibrium with the decagonal phase. Therefore information concerning the binary Al–Rh system is also included.
2. Experimental Alloys of 3–5 g were prepared by inductive melting in a water-cooled copper crucible under an Ar atmosphere. The alloys were thermally annealed at 8008C for 90–1400 h, at 9008C and higher temperatures for 70–120 h and subsequently water quenched. Short heat treatments were carried out in alumina crucibles in a vacuum furnace. For longer annealing times evacuated silica ampoules were used, inside which the alloys were contained in alumina crucibles in order to avoid a reaction of the partially molten samples with the silica. The alloys were studied by powder X-ray diffraction (XRD, CoKa 1 radiation was used), scanning electron microscopy (SEM) and transmission electron microscopy (TEM). The local phase compositions were determined in SEM by energy-dispersive X-ray analysis (EDX) on polished unetched cross-sections. The compositions of
0925-8388 / 00 / $ – see front matter 2000 Elsevier Science S.A. All rights reserved. PII: S0925-8388( 00 )00754-4
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B. Grushko et al. / Journal of Alloys and Compounds 305 (2000) 219 – 224
selected single-phase samples were examined by inductively coupled plasma optical emission spectroscopy (ICPOES). These compositions were used for the corrections of the EDX data. TEM examinations were carried out on powdered materials dispersed on grids with carbon film. Differential thermal analysis (DTA) was carried out for selected samples. Heating and cooling rates of 208C / min were used.
3. Results
3.1. Binary Al–Rh compounds The Al–Rh phase diagram was found to contain a number of phases isostructural to those in Al–Co (for the recent Al–Co phase diagram see Ref. [11]). They are AlRh
[6] of the CsCl-type structure corresponding to AlCo, hexagonal Al 5 Rh 2 [6] corresponding to Al 5 Co 2 and monoclinic Al 9 Rh 2 [7] corresponding to Al 9 Co 2 . In Ref. [6] a high-temperature cubic phase existing around the Al 5 Rh 2 composition was also reported. Its structure was recently determined [9]. In the following, this high-temperature cubic phase is designated C while H is used for the low-temperature hexagonal Al 5 Rh 2 phase. An additional high-temperature monoclinic phase (designated V in the following) was found in [10] at about 70 at.% Al. The typical powder XRD patterns of these phases are shown in Fig. 1; for detailed powder diffraction data of H–Al 5 Rh 2 , C–Al 5 Rh 2 and Al 7 Rh 3 see Ref. [10]. Our powder XRD pattern of Al 9 Rh 2 (Fig. 1a) also fits well with that of Ref. [7]. The lattice parameters of the Al–Rh phases are presented in Table 1. In Ref. [8] an Al 13 Rh 4 phase isostructural to monoclinic Al 13 Co 4 (M–Al 13 Co 4 in [11]) was reported to precipitate
Fig. 1. Powder XRD patterns of: (a) Al 9 Rh 2 , (b) the O 1 -phase, (c) the O 2 -phase, (d) the H–Al 5 Rh 2 phase, (e) the C-phase, (f) the C 1 -phase, (g) the monoclinic V-phase (Al 7 Rh 3 ) and (h) the Al–Cu–Rh decagonal phase; CoKa 1 radiation.
B. Grushko et al. / Journal of Alloys and Compounds 305 (2000) 219 – 224
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Table 1 Al–Rh phases Formula
O1 O2 H C V b a b
Al 9 Rh 2 Al 3 Rh Al 3 Rh Al 5 Rh 2 Al 5 Rh 2 Al 7 Rh 3 AlRh
Space group P21 /a
a a
P63 /mmc Pm3¯ b
Pm3¯ m
Lattice parameters
Ref.
a (nm)
b (nm)
c (nm)
b (8)
1.0149 2.38 2.34 0.7905 0.7680 1.0309 0.2980
0.6290 1.64 1.65 – – 0.3808 –
0.8557 3.28 1.24 0.7861 – 0.6595 –
142.4 – – – – 102.4 –
[7] This work This work [6,10] [9,10] [10] [4]
Orthorhombic. Monoclinic.
by a solid-state reaction from a supersaturated Al–Rh alloy. However, this structure was not observed in our experiments. Instead, two new orthorhombic phases of close compositions were found at about Al 3 Rh. The powder XRD patterns of these phases, designated O 1 and O 2 , are shown in Fig. 1b and c.
The O 2 -phase of a slightly lower Al concentration exhibited a powder XRD pattern similar to that of orthorhombic Al 3 Pd [12], TEM examination also revealed electron diffractograms very similar to those observed in Al–Pd (Fig. 2a–c, see Fig. 9 in [12]). The patterns in Fig. 2a–c can be indexed assuming an orthorhombic unit cell
Fig. 2. Typical electron diffraction patterns of the orthorhombic O 2 -phase: (a) Z.A.5[100], (b) Z.A.5[010], (c) Z.A.5[001]. The latter pattern is also typical of the [001] of the O 1 -phase. The pattern in (d) corresponds to Z.A.5[100] of the O 1 -phase, in (e) to Z.A.5[010] and in (f) to Z.A.5[101] of the O 1 -phase.
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Fig. 3. Partial Al–Rh phase diagram. For the phase designation see Table 1, the liquid is designated L, a is fcc Al.
with a52.34, b51.65 and c51.24 nm, which is close to the corresponding Al–Pd phase 2 . As is mentioned in the literature [13], the Al 3 Pd phase has a structural relation to the decagonal phases. The Al–Rh O 2 -phase was found at 75 at.% Al while according to Ref. [12], Al 3 Pd is formed at the Al concentration which is about 2.5 at.% lower. The O 1 -phase was found at 76 at.% Al. Its electron diffraction pattern (see Fig. 2c–f) can be indexed assuming an orthorhombic unit cell with a52.38, b51.64 and c5 3.28 nm and systematic extinctions of l 5 2n 1 1 for 0kl, h 5 2n 1 1 for hk0, h 1 l 5 2n 1 1 for h0l. It is worth mentioning that c /a¯tg368, the electron diffraction pattern indeed resembles that of the decagonal phase (see below). The lattice parameters of the O 1 -phase and the O 2 -phase are related as a 1 ¯a 2 , b 1 ¯b 2 and c 1 ¯ t 2 c 2 , where t is the golden mean. The available data concerning the phase compositions and transition temperatures are summarized in Fig. 3. All the phases apart from AlRh (also designated b in the following) and the high-temperature C-phase did not show pronounced compositional ranges. The lower temperature limits of the V-phase and C-phase regions have not yet been determined. We have also been unable to clarify the melting temperature of the O 1 -phase as yet. Since in the relevant ternary regions we only observed the ternary extensions of the b-phase, C-phase and O 2 -phase the lack of this information was not crucial for the following.
2
From the more recent work on Al–Pd we suggest the existence of more than one binary phase close to the Al 3 Pd composition. In the current paper we refer to Al 3 Pd as is described in Refs. [12–14] (in [14] this structure was studied in the Al 74 Pd 22 Mn 4 composition and is designated j9-phase). It seems that this basic Al 3 Pd structure can exhibit superstructure ordering. Indeed, inspection of the pattern in Fig. 2a reveals weak additional reflections corresponding to the doubled lattice parameter c of the basic Al 3 Pd structure. In this report we define the lattice parameters for the basic structure. A more detailed study of this phase is in progress.
Fig. 4. Phase equilibrium at 9008C in Al–Cu–Rh. For the phase designation see Table 1, the liquid is designated L. The compositions of the studied alloys are marked by spots.
3.2. Partial isothermal 900 and 8008 C sections of Al– Cu–Rh The partial 9008C section of Al–Cu–Rh is shown in Fig. 4. Three Al–Rh phases exhibited high Cu solubility. The b-phase extends to more than 35 at.% Cu and its Al concentration achieves 55 at.% at high Cu. The C-phase can contain up to about 13 at.% Cu and the O 2 -phase up to about 11 at.% Cu. The Al concentrations of the C-phase and O 2 -phase decrease sharply with the increase of the Cu concentration. The C-phase exhibited a superstructure ordering at its low-Al composition limits. This is in agreement with the observation in Ref. [5] where a cubic structure with a doubled lattice parameter (a51.5380(2) nm, space group of Fm3¯ ) was reported. The corresponding powder diffraction pattern of this phase designated C 1 contains a number of additional reflections (Fig. 1f,e), the strongest of which is at 0.2956 nm (2u 535.238 in Fig. 1f). The powder diffraction pattern calculated using the atomic positions from Ref. [5] was in qualitative agreement with our experimentally observed diffractogram. The existence of subregions of C and C 1 was not studied in detail. Based on the examination of the compositions marked in Fig. 4 the range of the cubic phase is conditionally divided between C and C 1 by a broken line. The Al 9 Rh 2 and H–Al 5 Rh 2 phases were found to contain ,0.5 at.% Cu. The Rh solubility in Al 2 Cu formed in the solidified liquid was ,1 at.%. At 9008C the Al–Cu–Rh decagonal phase (D) is stable in the range of 1–2 at.%, whose center is at about Al 64.5 Cu 16.8 Rh 18.7 . It coexists with the liquid and the ternary extensions of the b-phase, the C-phase (including C 1 ) and the O 2 -phase. The decagonal range in Al–Cu–Rh is smaller than that in Al–Cu–Co at the same 9008C [15] and coincides with the higher-Cu part of the decagonal
B. Grushko et al. / Journal of Alloys and Compounds 305 (2000) 219 – 224
223
Fig. 5. Typical electron diffraction patterns of the Al–Cu–Rh decagonal phase.
range in Al–Cu–Co. The powder XRD pattern of the Al–Cu–Rh decagonal phase (Fig. 1h) and its electron diffraction (Fig. 5) are similar to those of the Cu-rich Al–Cu–Co decagonal phase (Fig. 2b and Fig. 3a–c in [16], respectively). In a wide range of Cu concentrations the C-phase (including C 1 ) is in equilibrium with b, and the O 2 -phase is in equilibrium with the liquid. The regions marked by question marks in Fig. 4 were not studied in detail. The investigated part of the 8008C isothermal section (Fig. 6) is quite similar to that at 9008C. The ranges of the b-phase and O 2 -phase extend to still higher Cu than at 9008C. Zero solubility of Cu in H–Al 5 Rh 2 was found at this temperature. The decagonal phase is in equilibrium with the ternary extensions of the b-phase, the C-phase (including C 1 ) and the O 2 -phase but not with the liquid. Instead, the O 2 -phase was found in equilibrium with b at this temperature. The decagonal region lies around about Al 64.0 Cu 18.0 Rh 18.0 . It is slightly shifted from the decagonal
region at 9008C towards AlCu. This shift is smaller than that observed in Al–Cu–Co [15]. In contrast to Ref. [5] we did not observe any icosahedral phase in Al–Cu–Rh.
4. Conclusions The high-Al range of the Al–Rh alloy system and an adjacent range of Al–Cu–Rh were studied. We report the existence of seven Al–Rh phases: cubic AlRh, monoclinic Al 7 Rh 3 , high-temperature cubic Al 5 Rh 2 and low-temperature hexagonal Al 5 Rh 2 , two orthorhombic phases at about the Al 3 Rh composition and monoclinic Al 9 Rh 2 . At ternary compositions the high-temperature cubic Al 5 Rh 2 phase exhibits a superstructure ordering. We confirm the stability of the Al–Cu–Rh decagonal phase at 800–9008C. This ternary phase is in equilibrium with the ternary extensions of AlRh, cubic Al 5 Rh 2 and one of orthorhombic Al 3 Rh, at 9008C also with the liquid.
Acknowledgements We thank W. Reichert, K. Bickmann and C. Thomas for technical contributions. Financial support from the DFG (project Ur51 / 4-1) is gratefully acknowledged.
References
Fig. 6. Phase equilibrium at 8008C in Al–Cu–Rh. For the phase designation see Table 1, the liquid is designated L. The compositions of the studied alloys are marked by spots.
[1] A.P. Tsai, in: Z. Stadnik (Ed.), Physical Properties of Quasicrystals, Springer, Heidelberg, 1999. [2] B. Grushko, Heter. Chem. Rev. 1 (1994) 255. [3] A.P. Tsai, A. Inoue, T. Masumoto, Mater. Trans. JIM 30 (1989) 666. [4] T. Shishido, H. Takei, J. Less-Common Met. 119 (1986) 75. [5] X.Z. Li, K.-T. Park, K. Sugiyama, K. Hiraga, Metall. Mater. Trans. 29A (1998) 1559.
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[6] L.E. Edshammar, Acta Chem. Scand. 21 (1967) 647. [7] L.E. Edshammar, Acta Chem. Scand. 22 (1968) 2822. [8] Z.A. Chaudhury, C. Suryanarayana, J. Less-Common Met. 91 (1983) 181. [9] Yu. Grin, K. Peters, U. Burkhardt, K. Gotzmann, M. Ellner, Z. Kristallogr. 212 (1997) 439. [10] B. Grushko, M. Yurechko, Z. Kristallogr. 214 (1999) 313. [11] B. Grushko, R. Wittenberg, K. Bickmann, C. Freiburg, J. Alloys Compd. 233 (1996) 279.
[12] B. Grushko, M. Yurechko, N. Tamura, J. Alloys Compd. 290 (1999) 164. [13] Y. Matsuo, K. Hiraga, Phil. Mag. Lett. 70 (1994) 155. [14] M. Boudard, H. Klein, M. de Boissieu, M. Audier, H. Vincent, Phil. Mag. A74 (1996) 939. [15] B. Grushko, Phil. Mag. Lett. 66 (1992) 151. [16] B. Grushko, J. Mater. Res. 8 (1993) 1473.