Experimental determination of the ternary Co–Al–Nb phase diagram

Intermetallics 18 (2010) 2191e2207

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Experimental determination of the ternary CoeAleNb phase diagram O. Dovbenko 1, F. Stein*, M. Palm, O. Prymak MPI für Eisenforschung GmbH, Max-Planck-Str. 1., D-40237 Düsseldorf, Germany

a r t i c l e i n f o

a b s t r a c t

Article history: Received 21 May 2010 Received in revised form 7 July 2010 Accepted 7 July 2010 Available online 4 August 2010

Isothermal sections of the CoeAleNb system at 800, 1000, 1150, 1200, and 1250
C were experimentally determined by combining results of light-optical as well as scanning electron microscopy, electron-probe microanalysis, and X-ray diffraction measurements with Rietveld analysis of heat-treated alloys and adding some results obtained from solid/liquid diffusion couples. Three polytypes of Laves phases, which are the cubic C15 and the hexagonal C14 and C36 variants, occur as stable phases in the central part of the system. The phase equilibria involving these phases were in the focus of the present study. The composition-dependence of the crystallographic properties of the Laves phases as well as the crystal structure parameters of all other occurring intermetallic phases was investigated. As in case of the binary CoeNb system, the hexagonal Laves phases C14 and C36 are only stable at off-stoichiometric Nb contents. The preferential site occupations of the Co-, Al-, and Nb-atoms in the hexagonal lattices of the C14 and C36 Laves phase polytypes were determined by Rietveld analyses. For all binary phases except the Nbrich phase Nb3Al, the solubility for the third element was investigated and the extensions of all phase fields were established. Besides the three Laves phases, the most striking features of the CoeAleNb phase diagram are the occurrence of the slowly forming L21-type Heusler phase Co2AlNb and the extended phase field of the m phase. The existence of a m0 phase, which was reported in the literature, is not confirmed. Ó 2010 Elsevier Ltd. All rights reserved.

Keywords: A. Ternary alloy systems A. Laves phases B. Phase diagrams A. Intermetallics, miscellaneous D. Site occupancy

1. Introduction The binary CoeNb system belongs to the very few examples where three different polytypes of the Laves phase occur as stable phases [1e3]. As part of an inter-institutional research project on “The Nature of Laves Phases” [4], a study on the structure and stability of these Laves phases in case of the addition of Al as ternary element was performed in dependence on temperature and composition. Here we report on the resulting ternary CoeAleNb phase diagram focusing on phase equilibria which include the Laves phases. The three binary subsystems forming the boundaries of the ternary phase diagram are well-investigated. The most recent versions of the binary phase diagrams, which were used for the present investigations, are given in Refs. [2,3] for the CoeNb system, in Refs. [5,6] for AleCo, where [6] contains an up-dated version of the Al-rich part according to Refs. [7e9], and in Ref. [10] for AleNb.

* Corresponding author. Tel.: þ49 211 6792 557; fax: þ49 211 6792 299. E-mail address: [email protected] (F. Stein). 1 Present address: MSI Materials Science International Services GmbH, Industriestrasse 25, D-70565 Stuttgart, Germany. 0966-9795/$ e see front matter Ó 2010 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2010.07.004

The previous, sparse knowledge of the ternary CoeAleNb system is summarized in Refs. [11e13]. The only complete isothermal section of the ternary phase diagram which is available in the literature was established by Burnashova et al. [14] for a temperature of 800
C, see Fig. 1. They report the existence of the three ternary phases H, l1, and m0 . The so-called H phase is the Heusler phase Co2AlNb (L21 structure type), which was already found earlier by Markiv et al. [15] and confirmed later on by Buschow et al. [16] after heat-treating alloys of respective composition at 800
C for 480 h [15] and 900
C for 240 h [16]. Ziebeck and Webster [17] studied an alloy of the same composition in the ascast state and found it to be multiphase (without giving further details). This observation indicates that the formation of Co2AlNb from cast material is a slow process. The ternary phase l1 reported in Ref. [14] is the hexagonal, C14-type Laves phase Nb(Co,Al)2. It occurs at an off-stoichiometric Nb content of 36.5 at.% and Al contents ranging from about 5e55 at.% Al. In the binary CoeNb system, this phase is not stable at 800
C but occurs at higher temperatures between about 1250 and 1424
C [2,3]. The cubic, C15-type Laves phase NbCo2 is indicated by Burnashova et al. [14] to have a low solubility of less than about 2 at.% for Al at 800
C. As they did not analyse phase compositions, no exact values for the solubilities are given in their paper. A different result was reported by

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(>75 at.% Al) were not investigated. Results on the liquidus surface and reaction scheme of this system will be published elsewhere in Ref. [22]. 2. Experimental methods

Fig. 1. Isothermal section of the CoeAleNb phase diagram at 800
C redrawn from [14].

Blazina and Trojko [18], who investigated two alloys with compositions Coe3.3Ale30Nb and Coe6.7Ale26.7Nb (all compositions given in at.% throughout this paper) by X-ray diffraction (XRD) and found them to be single-phase C15 Laves phase after heat treatments in the range 800e1200
C. Due to the contradiction of these results to the data of Ref. [14], they were explicitly not accepted for the assessment of the ternary system in Ref. [12]. More recently, von Keitz et al. [19] prepared three ternary alloys with a nominal Nb content of 33.3 at.% and Al contents of 2, 4, and 6 at.%, which were annealed for 6 h at 1200
C and slowly cooled to room temperature. XRD measurements revealed that the first two alloys consisted of the C15 Laves phase whereas the Al-rich alloy only showed the C14 structure. The third ternary phase shown in Fig. 1 [14] is denoted as m0 and occurs near to the extended phase field of the m phase at an approximate composition of CoAlNb2. This phase had been already reported before by Hunt and Raman [20] who performed XRD investigations on heat-treated (1000
C/240 h) ternary alloys with compositions in the range between the m phase (Co, Al)7Nb6 and the AleNb phase Nb2Al. As they could not identify the crystal structure but found it to be very similar to that of the m phase, they denoted this phase as m0 . In an investigation of the effect of different alloying additions on the properties of the Al-rich compound NbAl3, Subramanian and Simmons [21] studied two Al-rich CoeAleNb alloys with compositions Coe63.6Ale24.3Nb and Coe73.1Ale20.8Nb after annealing at 1200
C. They present a partial isothermal section of the Al corner of the ternary system consisting of the two tie-triangles NbAl3 þ C14 þ CoAl and NbAl3 þ CoAl þ L (L is the liquid phase coming from the Al corner). The solubility of Co in NbAl3 was determined to be less than 1 at.%. In the present investigation, five isothermal sections of the CoeAleNb system at 800, 1000, 1150, 1200, and 1250
C, excluding the Nb- and the Al-corner, were determined by combining results of light-optical as well as scanning electron microscopy (SEM), electron-probe microanalyses (EPMA), and XRD measurements of heat-treated alloys and diffusion couples. The focus is on the phase equilibria involving the different polytypes of Laves phases. Alloys in the Nb corner above about 70 at.% Nb and in the Al corner

From the pure metals Co (99.95 wt.%), Al (99.999 wt.%), and Nb (99.9 wt.%), 29 alloys were prepared by crucible-free levitation melting in an argon atmosphere and drop-casting into cold copper moulds. Another five alloys were produced by arc-melting under argon. The nominal compositions of all investigated alloys are given in Table 1. For further examinations, the alloys were sectioned by electro-discharge machining (EDM). Wet chemical analyses of selected as-cast samples and samples annealed at 800
C yielded impurity contents in the range 40e70 wt.ppm oxygen, 20e30 wt.ppm nitrogen and 15e30 wt.ppm carbon. Nominal and analysed sample compositions (Tables 2e6) usually agree within 1 at.%. For heat treatments at 800, 1000 and 1150
C, pieces of the alloys were wrapped in niobium foil and encapsulated in quartz ampoules which were evacuated and back-filled with argon several times before sealing them. After heat treatments for up to 5000 h at 800
C, 1000 h at 1000
C, and 500 h at 1150
C, the samples were quenched in iced brine (10% NaCl solution). The individual times of the heat treatments of all samples are listed in Tables 2e4. In some cases two different heat treatment times were chosen in order to check if the times were sufficient to achieve equilibrium. Heat treatments at 1200 and 1250
C for up to 240 h were at first performed in an argon atmosphere with samples wrapped in niobium foil. The experimental examination of these samples showed that a lot of the Nb-rich samples contained significant amounts of an additional, impurity-stabilized phase (see also Section 3.5). Therefore, the heat treatments were repeated with new samples which, after wrapping in niobium foil, were placed into closed alumina crucibles which were filled up with titanium filings acting as oxygen getter. Rapid cooling was performed in a jet of argon gas. Heat treatment times and sample compositions are given in Tables 5 and 6. Solid/liquid diffusion couples consisting of single-phase, binary NbCo2 (exact composition Coe33.0Nb) and Al were prepared by machining a cylindrical hole of 5 mm in diameter and 5 mm in height into a massive piece of NbCo2 by EDM, pressing an Al cylinder of the same diameter into the hole and covering it with a slice of NbCo2. The whole assembly was wrapped in niobium foil, heat-treated at 1000
C in an argon atmosphere, and cooled to room temperature by turning off the furnace. Two diffusion couples of this type were prepared one of them heat-treated for 2 h and the other one for 24 h. For metallographic observation of the microstructure of the heat-treated alloys by light-optical microscopy, the samples were etched with “Ti2”, an etchant consisting of glycerine (68 vol.%), 70%-HNO3 (16 vol.%), and 40%-HF (16 vol.%). EPMA measurements were carried out with a Cameca SX 50 and a Jeol JXA-8100 instrument using pure Co, Al, and Nb as standards. Phase compositions were analysed with a beam probing sample volumes of about 1 mm3. The relative error of the resulting compositions is 1%. The actual bulk composition of the heat-treated alloys was determined by averaging the EPMA results of at least 1000 spot analyses per sample with spacings of 5 or 10 mm in four to five rectangular fields of about 100e300 mm in width and length. Comparison with chemical analysis by inductively coupled plasma optical emission spectroscopy (ICP-OES) gave good agreement of the analysed bulk compositions within 1 at.%. The compositions of the different phases were determined by measuring at least fifteen points per

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Table 1 Nominal compositions of investigated alloys and phases identified by XRD after heat treatments (alloys are sorted according to their Al content). Alloy no.

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 a b c d e

Nominal composition

Phases identified by XRD

at.% Co

at.% Al

at.% Nb

800
C

1000
C

1150
C

1200
C

1250
C

75.5 62 64 77 65 70.5 67 59 63 66 61 59 80 70 65 64 62 60 48 75 55 30 20 60 50 33 26 40 17 47 32 15 2 22

1 1 1.5 2 2 2.5 3 3 4 5 6 8 10 10 10 10 10 10 10 13 13 20 20 24.5 25 25 25 35 35 38 48 60 65 68

23.5 37 34.5 21 33 27 30 38 33 29 33 33 10 20 25 26 28 30 42 12 32 50 60 15.5 25 42 49 25 48 15 20 25 33 10

Nb2Co7 þ C15 C15 þ C14dþCo7Nb6 C15 þ C14d þ Co7Nb6 Co þ Nb2Co7 C15 þ C14d þ Co7Nb6 C15 þ C36 C15 þ Co2AlNb C14 þ Co7Nb6 C15 þ C14 C15 þ C36 þ Co2AlNb C14 C14 Co þ C36 þ CoAl C36 þ CoAl C36 þ CoAl þ Co2AlNb C36 þ Co2AlNb C15 þ C36 þ Co2AlNb C15 þ C14 þ Co2AlNb C14 þ Co7Nb6 e C14 þ Co2AlNb e e C36 þ CoAl þ Co2AlNb C14 þ Co2AlNb C14 þ Co7Nb6 e C14 þ CoAl þ Co2AlNb C14 þ Co7Nb6 þ Nb2Al C14 þ CoAl þ Co2AlNb e e C14 þ NbAl3 þ Nb2Al CoAl þ Co2Al5 þ NbAl3

Co þ Nb2Co7aþ C36 C15 þ C14d þ Co7Nb6 C15 þ C14d þ Co7Nb6 Co þ C36 C15 þ C14 þ Co7Nb6 C15 þ C36 C15 C14 þ Co7Nb6 C15 þ C14 C15 þ Co2AlNb C14 C14 Co þ C36 þ CoAl C36 þ CoAl C36 þ Co2AlNb C15 þ C36 þ Co2AlNb C15 þ C14 þ Co2AlNb C14 þ Co2AlNb C14 þ Co7Nb6 Co þ C36 þ CoAl C14 þ Co2AlNb e e C36 þ CoAl þ Co2AlNb C14 þ Co2AlNb C14 þ Co7Nb6 e C14 þ CoAl þ Co2AlNb C14 þ Co7Nb6 þ Nb2Al C14 þ CoAl þ Co2AlNb e e C14 þ NbAl3 þ Nb2Al CoAl þ Co2Al5 þ NbAl3

Co þ C36 C15 þ C14 þ Co7Nb6 C15 þ C14 Co þ C36 C15 C15 C15 C14 þ Co7Nb6 C15 þ C14 C15 C14 C14 Co þ C36 C36 þ CoAl C36 þ Co2AlNb C15 þ C36 þ Co2AlNb C15 þ C14 þ Co2AlNb C14 þ Co2AlNb C14 þ Co7Nb6 Co þ C36 þ CoAl C14 þ Co2AlNb Co7Nb6 Co7Nb6 þ Nb2Al C36 þ Co2AlNb C14 þ Co2AlNb C14 þ Co7Nb6 Co7Nb6 þ Nb2Al C14 þ CoAl þ Co2AlNb C14 þ Co7Nb6 þ Nb2Al C14 þ CoAl þ Co2AlNb C14 þ CoAl C14 þ CoAl þ NbAl3b C14 þ NbAl3 þ Nb2Al CoAl þ Co2Al5 þ NbAl3

Co þ C36b C14 þ Co7Nb6b ee e ee C15 ee C14 þ Co7Nb6 ee C15 ee ee Co þ C36 C36 þ CoAl C36 þ Co2AlNb C15 þ C36 þ Co2AlNb C14 þ Co2AlNb b C14 þ Co2AlNb C14 þ Co7Nb6 e C14 þ Co2AlNb e e C36 þ Co2AlNb C14 þ Co2AlNb e e C14 þ CoAl þ Co2AlNb C14 þ Co7Nb6 þ Nb2Al e e e C14 þ NbAl3 þ Nb2Al e

Co þ C36c ee C15 þ C14 e ee e ee ee ee C15 ee ee e e C36 þ Co2AlNb C15 þ C36 þ Co2AlNb C14 þ Co2AlNb C14 þ Co2AlNbb ee e C14 þ Co2AlNb e e C36 þ Co2AlNb C14 þ Co2AlNb C14 þ Co7Nb6 e C14 þ CoAl þ Co2AlNb C14 þ Co7Nb6 þ Nb2Al e e e C14 þ NbAl3 þ Nb2Al e

Phase only present in small amounts, XRD peaks strongly affected by C36 peaks. From EPMA (no XRD). Has been C36 þ liquid at 1250
C. Very probably metastable phase. Sample contained additional TiNi2-type impurity phase after heat treatment.

phase. Concentration profiles of the diffusion couples were obtained by recording line scans along the concentration gradient with a step width of 1 mm. In some cases, phase compositions of heat-treated alloys were measured by SEM-EDX instead of EPMA using a Hitachi S-530 instrument equipped with an EDAX detector. Microstructures of heat-treated alloys were investigated with both the EPMA instrument in SEM mode and the Hitachi S-530 SEM. XRD investigations on crushed and powdered samples with particle sizes <90 mm were performed using monochromatic Co-Ka1 radiation on a Huber and an Inel diffractometer and Cu-Ka radiation on a PHILIPS PW-1827 diffractometer. The lattice parameters were refined by the least-squares fitting software LCLSQ 8.5 [23] and by the Rietveld refinement program package FullProf [24], which was also used for the determination of site occupation factors in Laves phases of selected alloys. The quality of the XRD measurements was not sufficient to give precise values of the site occupation factors. Therefore, the results were only used to establish the preferred lattice sites of the different atom species. 3. Results Table 1 gives a list of the nominal compositions of all investigated alloys and summarizes all phases identified by XRD in the alloys after the heat treatments at 800 to 1250
C. Tables 2e6 give the

analysed overall compositions of the heat-treated samples along with the compositions and lattice parameters of all phases as obtained by EPMA and XRD. The crystallographic structure types and composition ranges of all phases (excluding the hexagonalclose-packed Co variant and the Al solid solution, which both are not stable in the investigated temperature range) and the solubilities for the third component in the binary phases are listed in Table 7. The resulting isothermal sections of the CoeAleNb phase diagram are shown in Figs. 2e6 and representative microstructures of threephase alloys are presented in Fig. 7. In the following sections, we will go through the ternary system phase by phase. The central part of the phase diagram is dominated by the extended phase field of the Laves phase which will be discussed first and in most detail. After a brief section on the phase Nb2Co7, which is structurally related to the Laves phases, the results on the Heusler phase Co2AlNb as the only truly ternary phase of the system are presented. The subsequent two sections are devoted to the m and m0 phases. In case of the latter phase, the results indicate that it does not exist in the ternary system. In the remaining sections, the AleNb phases Nb2Al and NbAl3, the Al-rich CoeAl phases, the extended B2-type phase CoAl, and the aCo solid solution are briefly discussed. The final section deals with the liquid phase which penetrates into the ternary system from two points one of which is the Al corner with its low melting temperatures and the other is the Co-rich eutectic in the CoeNb system.

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Table 2 Alloy and phase compositions as analysed by EPMA/SEM-EDX and lattice parameters determined by XRD after heat treatments at 800
C. Alloy no.

1

Analysed alloy composition at.% Co

at.% Al

at.% Nb

74.4

1.2

24.4

Heat treatment time [h]

Phases

5000

C15 Nb2Co7

2

62a

1a

37a

5000

3

64a

1.5a

34.5a

2000

4

76.3

1.8

21.9

2000

5

62.8

2.0

35.2

1000

6

70.5a

2.5a

27a

2000

7

67a

3a

30a

5000

8

58.7

3.1

38.2

1000

9

62.3

3.8

33.9

1000

10

64.2

5.9

29.9

2000

11 12 13

60.6 58.6 79.4

6.0 7.8 10.2

33.4 33.6 10.4

1000 1000 1000

14

69.7

9.8

20.5

1000

15

64.8

10.1

25.1

1000

16

62.7

10.6

26.7

2000

17

62a

10a

28a

2000

18

59.1

10.6

30.3

5000

19

47.7

9.6

42.7

1000

21

55.0

12.6

32.4

1000

24

58.0

26.3

15.7

2000

26

32.1

25.4

42.5

2000

28

40.2

34.2

25.6

1000

29

17.5

33.6

48.9

1000

30

47a

38a

15a

2000

33

1.4

63.1

35.5

1000

34

22.0

67.0

11.0

1000

a b c d

Nominal composition. Phase only present in small amounts, probably metastable. C14-C15 only distinguishable in XRD. Phase too small for EPMA analysis.

Phase composition at.% Co

C15 C14b Co7Nb6 C15 C14b Co7Nb6 Co Nb2Co7 C15 C14b Co7Nb6 C15 C36 C15 Co2AlNb C14 Co7Nb6 C15 C14 C15 C36 Co2AlNb C14 C14 Co C36 CoAl C36 CoAl C36 CoAl Co2AlNb C36 Co2AlNb C15 C36 Co2AlNb C15 C14 Co2AlNb C14 Co7Nb6 C14 Co2AlNb C36 CoAl Co2AlNb C14 Co7Nb6 C14 CoAl Co2AlNb C14 Co7Nb6 Nb2Al C14 CoAl Co2AlNb C14 NbAl3 Nb2Al CoAl Co2Al5 NbAl3

Lattice parameters

at.% Al

70.5  0.1 75.2  0.2

1.3  0.1 1.0  0.1

at.% Nb

a [Å]

28.2  0.2 23.8  0.2

6.730 a ¼ 4.582 b ¼ 8.161 6.794 4.806 4.904 6.792 4.816 4.903

88.6  0.5 75.0  0.3

7.1  0.1 1.7  0.3

4.3  0.6 23.3  0.1

63.8  0.4 62.3  0.2 51.3  0.3

2.1  0.1 2.2  0.2 0.9  0.1

34.1  0.4 35.5  0.2 47.9  0.3

66.3 52.8 60.7 51.0 62.5 61.9

     

0.2 0.1 0.3 0.4 0.2 0.2

60.6  0.2 58.6  0.3 88.7  0.5 72.6  0.6 62  1 72.0  0.3 61  1 69.6  0.3 59  1 53.3  0.8 68.0  0.3 51.8  0.3

62.4 62.4 49.8 49.5 45.7

31.9 32.7 33 ed d e 23.6 21.0 5.0

10.6 0.22 0.2 45.8 28.7 1.5

    

0.2c 0.2c 0.4 0.7 0.9

 0.2  0.6 1

2.9 22.3 3.7 1.4 3.9 3.7

     

0.2 0.2 0.2 0.2 0.1 0.1

6.0  0.2 7.9  0.3 8.7  0.5 4.0  0.9 32  3 3.2  0.1 33  2 3.4  0.1 32  1 25.8  0.9 3.3  0.2 27.7  0.4

4.3 4.3 25.5 13.2 6.2

    

0.2c 0.2c 0.4 0.6 0.7

 0.2  0.5 1

 0.6  0.1  0.3

31.8 18.7 33 ed d e 40.8 30.6 29.0

     

52.5 72.9 42.2 54 71.3 72.9

0.6 0.05 0.2 0.9 0.2 0.2

30.8 24.9 34.6 47.6 33.6 34.4

     

0.2 0.2 0.5 0.4 0.2 0.1

33.4  0.2 33.5  0.2 2.6  0.9 23.4  0.9 62 24.9  0.3 61 27.0  0.4 91 20.8  0.8 28.7  0.3 20.5  0.4

33.3 33.3 24.7 37.4 48.1

    

0.1c 0.1c 0.2 0.7 0.3

 0.2  0.1 1

 0.5  0.3  0.4

36.3 48.6 34 ed d e 35.6 48.4 66.1

     

36.9 26.9 57.6 0.2 0.04 25.7

     

0.9 0.4 0.8 1 0.2 0.3

 0.6  0.2  0.4

0.9 0.4 0.8 0.4 0.06 0.3

a ¼ 4.572 b ¼ 8.164 6.786 4.823 4.911 6.736 4.760 6.766 5.947 4.834 4.913 6.801 4.817 6.770 4.772 5.943 4.825 4.835 3.569 4.756 2.876 4.750 2.873 4.769 2.882 5.935 4.772 5.935 6.770 4.784 5.952 6.790 4.807 5.960 4.883 4.937 4.852 5.965 4.764 2.880 5.875 4.930 4.962 4.973 2.923 5.933 4.999 5.027 9.884 4.972 2.908 5.916 5.039 3.841 9.910 2.858 7.680 3.840

c [Å] c ¼ 6.344

b ¼ 110.2
7.826 26.21 7.836 26.20

c ¼ 6.338

b ¼ 110.3
7.831 26.20 15.52

7.838 26.22 7.839 15.56

7.835 7.846 15.48

15.48 15.54

15.55 15.59

7.831

7.909 26.26 7.872 15.53

8.003 26.81 8.026

8.094 27.24 5.159 8.028

8.288 8.618 5.186 7.615 8.619

O. Dovbenko et al. / Intermetallics 18 (2010) 2191e2207

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Table 3 Alloy and phase compositions as analysed by EPMA/SEM-EDX and lattice parameters determined by XRD after heat treatments at 1000
C. Alloy no.

1

Analysed alloy composition at.% Co

at.% Al

at.% Nb

75.5a

1a

23.5a

Heat treatment time [h]

Phases

1000

Co C36 Nb2Co7b C15 C14c Co7Nb6 C15 C14c Co7Nb6 Co C36 C15 C14 Co7Nb6 C15 C36 C15 C14 Co7Nb6 C15 C14 C15 Co2AlNb C14 C14 Co C36 CoAl C36 CoAl C36 Co2AlNb C15 C36 Co2AlNb C15 C14 Co2AlNb C14 Co2AlNb C14 Co7Nb6 Co C36 CoAl C14 Co2AlNb C36 CoAl Co2AlNb C14 Co7Nb6 C14 CoAl Co2AlNb C14 Co7Nb6 Nb2Al C14 CoAl Co2AlNb C14 Nb2Al NbAl3 CoAl Co2Al5 NbAl3

2

62a

1a

37a

1000

3

64a

1.5a

34.5a

1000

4

77a

2a

21a

1000

5

64.3

2.0

33.7

260

6

69.9

2.5

27.6

1000

7 8

67a 58.4

3a 3.0

30a 38.6

1000 235

9

62.0

3.7

34.3

260

a

10

66

11 12 13

60.3 58.5 79.8

5

a

29

6.1 7.9 9.9

a

33.6 33.6 10.3

1000 260 260 235

14

69.6

10.3

20.1

235

15

64.8

9.7

25.5

245

16

63.4

9.9

26.7

1000

17

61.3

10.1

28.6

1000

18

59.7

9.7

30.6

1000

19

48.3

9.7

42.0

235

20

75

a

13

a

12

a

1000

21

54.9

12.7

32.4

235

24

59.7

24.8

15.5

1000

26

33a

25a

42a

1000

28

39.7

34.4

25.9

245

29

17.7

33.2

49.1

245

30

47a

38a

15a

1000

33

1.3

63.9

34.8

235

34

20.8

68.0

11.2

245

a b c d e

Phase composition at.% Co

Nominal composition. Phase only present in small amounts, XRD peaks strongly affected by C36 peaks. Phase only present in small amounts, probably metastable. C14/C36/C15 only distinguishable in XRD. Phase too small for EPMA analysis.

    

Lattice parameters

at.% Al

at.% Nb

0.5 0.4 0.2 0.2d 0.2d

2.0  0.1 2.2  0.2 1.2  0.2 2.5  0.1d 2.5  0.1d

61.5  0.2 51.0  0.3 62.3  0.4d 62.3  0.4d 65.9  0.2 51.7  0.2 61.0  0.3 58.7  0.2 86.2  0.1 71.5  0.2 66  1 71.1  0.2 63.8  0.2 68.3  0.2 53.8  0.3 66.9  0.3d 66.9  0.3d 52.5  0.4 64.8  0.3d 64.8  0.3d 51.4  0.2 62.2  0.5 51.4  0.3 50.3  0.5 46.1  0.4

3.5  0.1 1.6  0.1 3.6  0.3d 3.6  0.3d 4.6  0.2 24.7  0.2 5.6  0.2 7.8  0.2 11.4  0.1 4.4  0.1 29  1 4.5  0.2 31.4  0.5 5.0  0.1 27.3  0.7 4.9  0.4d 4.9  0.4d 25.1  0.6 4.8  0.4d 4.8  0.4d 24.5  0.2 5.7  0.6 23.4  0.6 13.0  0.5 6.3  0.3

64.7 62.8 51.5 68.8 68.8

55.7 50.6 68.6 57.3 55.1

    

0.6 0.2 0.1 0.2 0.2

    

0.5 0.6 0.1 0.3 0.3

    

0.3 0.4 0.2 0.3d 0.3d

35.0  0.2 47.4  0.3 34.1  0.3d 34.1  0.3d 29.5  0.2 23.6  0.3 33.4  0.2 33.5  0.2 2.4  0.1 24.1  0.2 51 24.4  0.2 4.8  0.4 26.7  0.2 18.9  0.9 28.2  0.3d 28.2  0.3d 22.4  0.3 30.4  0.3d 30.4  0.3d 24.1  0.2 32.1  0.3 25.2  0.4 36.7  0.4 47.6  0.2

34.0  0.2 25.9  0.5 26.3  0.2 7.3  0.3 16.0  0.3

 0.5

0.5 0.5 0.5 0.3

35.1 ee 30.2 40.1 31.4 29.4

   

0.7 0.5 0.8 0.3

34.1 ee 20.5 35.1 47.8 66.8

11  1 0.1  0.1 0.1  0.1 47.1  0.2 28.3  0.2 1.2  0.2

54 38.5 73.7 52.9 71.6 73.2

     

1 0.8 0.2 0.2 0.2 0.2

35  1 61.4  0.8 26.2  0.1 0.01  0.0 0.1  0.2 25.6  0.2

30.8 ee 49.3 24.8 20.8 3.8

 0.5

10.3 23.5 5.1 35.4 28.9

33.3 35.0 47.3 27.7 27.7

   

 0.3    

0.6 0.3 0.4 0.3

a [Å]

c [Å]

3.554 4.742

15.46

6.793 4.807 4.903 6.797 4.809 4.908 3.566 4.745 6.787 4.817 4.909 6.746 4.772 6.763 4.832 4.912 6.803 4.815 6.775 5.952 4.828 4.835 3.583 4.756 2.901 4.754 2.879 4.771 5.908 6.766 4.787 5.940 6.783 4.798 5.957 4.807 5.953 4.883 4.937 3.58(1) 4.75(1) 2.87(1) 4.850 5.963 4.772 2.898 5.879 4.954 4.978 4.972 2.891 5.920 4.999 5.028 9.897 4.976 2.892 5.916 5.039 9.932 3.842 2.857 7.677 3.839

7.836 26.22 7.840 26.22

15.47 7.845 26.20 15.55 7.844 26.22 7.846

7.838 7.845 15.47 15.49 15.55

15.60

7.820 7.830 7.909 26.26 15.47 7.866 15.55

8.011 26.88 8.037

8.090 27.24 5.161 8.029

8.290 5.177 8.617 7.616 8.615

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O. Dovbenko et al. / Intermetallics 18 (2010) 2191e2207

Table 4 Alloy and phase compositions as analysed by EPMA/SEM-EDX and lattice parameters determined by XRD after heat treatments at 1150
C. Alloy no.

Analysed alloy composition

Heat treatment time [h]

at.% Co

at.% Al

at.% Nb

1

74.5a

1a

23.5a

500

2

62a

1a

37a

500

3

64

a

4

76.9

1.9

21.1

500

5 6 7 8

64.7 70.5 67 58.8

1.9 2.5 3 3.1

33.4 27 30 38.1

168 500 500 113

9

62.6

3.9

33.5

168

10 11 12 13

65.9 60.3 58.6 79.3

4.7 6.0 7.8 10.0

29.5 33.8 33.6 10.7

500 168 168 113

14

69.7

9.9

20.4

113

15

64.8

9.9

25.3

136

16

64

10

26

500

17

61.7

1.5

a

9.8

34.5

a

28.5

500

500

18

60a

10a

30a

500

19

48.1

9.6

42.3

113

20

a

75

13

a

12

a

1000

21

55.3

12.4

32.3

113

22 23

31.2 19.6

20.4 21.0

48.4 59.4

500 500

24

60

24.5

15.5

500

26

32.2

25.4

42.3

500

27

26.2

24.9

48.9

500

a

a

28

40

35

29

17.9

33.4

48.7

136

30

46.5

37.7

15.8

500

31

33.1

46.2

20.7

500

32

15a

60a

25a

500

33

1.6

62.1

36.3

136

34

20.9

67.6

11.5

136

a b c

25

a

136

Phases

Phase composition at.% Co

Co C36 C15 C14 Co7Nb6 C15 C14 Co C36 C15b C15 C15 C14 Co7Nb6 C15 C14 C15 C14 C14 Co C36 C36 CoAl C36 Co2AlNb C15c C36c Co2AlNb C15c C14c Co2AlNb C14 Co2AlNb C14 Co7Nb6 Co C36 CoAl C14 Co2AlNb Co7Nb6 Co7Nb6 Nb2Al C36 Co2AlNb C14 Co7Nb6 Co7Nb6 Nb2Al C14 CoAl Co2AlNb C14 Co7Nb6 Nb2Al C14 CoAl Co2AlNb C14 CoAl C14 CoAl NbAl3 C14 NbAl3 Nb2Al CoAl Co2Al5 NbAl3

at.% Nb

0.3 0.1 0.5 0.1

4.8  0.2 1.5  0.1 1.8  0.2 2.5  0.1

60.8  0.2 50.9  0.2 62.6  0.3 61.7  0.2 65.9  0.2 60.3  0.4 58.6  0.2 84.0  0.1 70.9  0.2 70.4  0.2 67.5  0.2 66.8  0.1 54.8  0.1 66.0  0.2 66.0  0.2 53.5  0.1 50.1  0.2 50.1  0.2 52.1  0.2 61.0  0.2 51.7  1.0 51.0  0.4 45.6  0.3

3.5  0.1 1.8  0.1 3.7  0.2 4.1  0.1 4.7  0.2 5.9  0.1 7.8  0.3 12.4  0.1 5.1  0.1 5.4  0.1 26.9  0.1 6.6  0.1 26.5  0.1 6.5  0.1 6.5  0.1 25.4  0.1 6.7  0.1 6.7  0.1 24.5  0.2 7.8  0.2 23.4  2.0 12.3  0.3 7.1  0.3

91.3 73.4 65.3 70.1

   

Lattice parameters

at.% Al

55.5  0.3 50.8  0.2 31.2 29.4  0.2 11.4  0.4 67.8  0.2 57.7  0.2 35.0  0.1 30.7  0.2 26.9  0.3 9.6  0.3

24.6  0.5 19.6  0.5 4.1  0.6 32.3  0.3 50.2  0.2 50.5  0.3 23.0  0.4 49.7  0.2 15.2  0.2 46.6  0.3 0.9  0.1 10.6  0.4 0.1  0.1 0.2  0.1 44.9  0.2 28.4  0.1 1.3  0.3

11.4 24.2 20.4 21.0 20.7 6.4 29.3 28.4 21.6 24.9 23.4

39.9 32.8 29.0 35.6 43.5 33.0 44.1 49.2 52.3 53.1 74.9 53.5 79.8 35.5 55.1 71.6 73.3

 0.2  0.3        

                

0.1 0.2 0.1 0.2 0.1 0.3 0.4 0.3

0.5 0.5 0.4 0.3 0.4 0.2 0.2 0.2 0.2 0.1 0.2 1.0 0.1 0.2 0.2 0.1 0.2

3.9 25.1 32.9 27.5

   

0.1 0.1 0.4 0.2

35.8  0.2 47.3  0.2 33.7  0.2 34.2  0.1 29.5  0.2 33.7  0.4 33.6  0.2 3.7  0.1 24.0  0.2 24.2  0.2 5.7  0.1 26.6  0.2 18.7  0.2 27.5  0.2 27.5  0.2 21.2  0.2 29.5  0.3 29.5  0.3 23.5  0.2 31.2  0.1 24.8  0.8 36.7  0.4 47.3  0.2

33.2 25.0 48.4 49.7 67.9 25.8 13.0 36.6 47.7 48.2 67.0

 0.2  0.3        

0.2 0.2 0.2 0.2 0.1 0.1 0.3 0.3

35.5  0.5 47.6  0.2 66.9  0.3 32.1  0.3 6.3  0.6 16.5  0.3 32.9  0.5 1.1  0.1 32.5  0.3 0.2  0.4 24.2  0.2 35.8  1.0 26.1  0.1 64.3  0.3 0.02  0.02 0.03  0.01 25.4  0.2

a [Å] 3.571 4.740 6.798 4.825 4.902 6.801 4.824 3.580 4.753 6.789 6.738 6.763 4.841 4.912 6.802 4.821 6.778 4.826 4.836 3.591 4.754 4.756 2.881 4.775 5.895 6.754 4.779 5.916 6.776 4.794 5.939 4.817 5.950 4.883 4.935 3.58(1) 4.75(1) 2.88(1) 4.849 5.959 4.988 5.005 9.848 4.772 5.840 4.963 4.987

c [Å] 15.46 7.841 26.21 7.844 15.49

7.854 26.22 7.837 7.835 7.844 15.48 15.50 15.55

15.57

7.811 7.816 7.919 26.30 15.48 7.863 26.96 27.06 5.145 15.54 8.020 26.95

4.967 2.901 5.901 5.002 5.033 9.892 4.976 2.904 5.911 5.003 2.865

8.031

5.040 3.842 9.936 2.856 7.677 3.840

8.290 8.619 5.176 7.615 8.617

Nominal composition. Sample contains an additional very small amount of C14 Laves phase and traces of a very fine, Nb-rich impurity phase in the vicinity of C14 Laves phase. C14/C36/C15 only distinguishable in XRD.

8.101 27.27 5.160 8.047

8.131

O. Dovbenko et al. / Intermetallics 18 (2010) 2191e2207

2197

Table 5 Alloy and phase compositions as analysed by EPMA/SEM-EDX and lattice parameters determined by XRD after heat treatments at 1200
C. Alloy no.

Analysed alloy composition

Heat treatment time [h]

at.% Co

at.% Al

at.% Nb

1

74.9

1.0

24.1

96

2

62a

1a

37a

96

6 8

70.1 58.5

2.1 3.1

27.8 38.4

240 72

10 13

65.0 79.7

5.1 9.9

29.9 10.4

240 6

14

69.7

9.9

20.4

6

15

65a

10a

25a

96

16

64a

10a

26a

240

Phase composition at.% Co

17

61.1

10.5

28.4

240

18

59.6

9.8

30.6

240

19

48a

10a

42a

72

20

76.3

13.2

10.5

96

21

54.7

12.6

32.7

72

24

60a

24.5a

15.5a

28

40a

35a

25a

96

29

17a

35a

48a

96

33

1.3

62.5

36.2

96

a

Phases

240

Co C36 C14 Co7Nb6 C15 C14 Co7Nb6 C15 Co C36 C36 CoAl C36 Co2AlNb C15 C36 Co2AlNb C14 Co2AlNb C14 Co2AlNb C14 Co7Nb6 Co C36 CoAl C14 Co2AlNb C36 Co2AlNb C14 CoAl Co2AlNb C14 Co7Nb6 Nb2Al C14 NbAl3 Nb2Al

93.4 74.5 63.0 51.8 69.3 60.4 50.7 64.9 83.8 70.6 70.1 67.7 66.4 55.1

63.1 51.5 60.5 51.2 50.7 45.2 83.4 70.5 69.6 55.2 50.6

             

          

0.1 0.2 0.6 0.2 0.1 0.2 0.2 0.4 0.2 0.8 0.2 0.3 0.2 0.1

0.1 0.2 0.2 0.1 0.3 0.2 0.2 0.2 0.2 0.2 0.2

10.9  0.4 0.1  0.1 0.2  0.1

Lattice parameters

at.% Al

at.% Nb

a [Å]

c [Å]

2.2  0.1 0.7  0.1 1.0  0.1 0.6  0.1 2.6  0.1 3.4  0.1 1.8  0.1 4.8  0.2 12.3  0.1 5.1  0.1 5.7  0.1 26.9  0.4 7.1  0.1 26.7  0.1

4.4  0.1 24.9  0.2 36.0  0.5 47.6  0.2 28.1  0.1 36.2  0.2 47.5  0.2 30.3  0.3 3.9  0.3 24.3  0.3 24.2  0.2 5.4  0.2 26.5  0.2 18.2  0.1

3.573 4.746

15.47

7.4  0.2 25.2  0.1 8.4  0.1 24.5  0.2 12.3  0.2 7.3  0.2 12.5  0.1 5.5  0.1 24.5  0.2 11.5  0.2 24.2  0.2

53.7  0.6 73.2  0.9 36.0  0.6

29.5  0.1 23.3  0.1 31.1  0.1 24.3  0.1 37.0  0.3 47.5  0.2 4.1  0.1 24.0  0.2 5.9  0.1 33.3  0.2 25.2  0.1

35.4  0.4 26.7  0.8 63.8  0.7

6.734 4.849 4.921 6.766 3.583 4.754 4.755 2.879 4.777 5.890 6.774 4.800 5.939

4.821 5.945 4.884 4.935

4.850 5.958 4.769 5.836 4.965 2.901 5.895 5.005 5.034 9.894 5.039 3.842 9.930

7.853 26.26

15.47 15.49 15.56

15.62

7.813 7.922 26.30

7.859 15.53 8.020

8.097 27.29 5.162 8.291 8.618 5.178

Nominal composition.

3.1. The Laves phases C14, C15 and C36 3.1.1. Homogeneity ranges In the binary CoeNb system, the C14 and C36 polytype of the Laves phases only exist at temperatures above about 1250 and 1040
C, respectively [2,3]. As Figs. 2e6 show, this is completely different in case of the ternary system, where all three Laves phase polytypes C14, C15 and C36 occur as stable phases in the entire investigated temperature range between 800 and 1250
C, i.e., the addition of aluminium stabilizes the hexagonal polytypes to lower temperatures. C15 Laves phase. The cubic C15 Laves phase exists in a wide composition range in the binary CoeNb system, which is between about 27 and 35 at.% Nb at 800
C and reaches its maximum extension at 1250
C (about 26e35.5 at.% Nb) [2]. Within this temperature range, the solubility of Al in the C15 Laves phase is found to increase continuously from about 4 at.% at 800
C to 8 at.% at 1250
C. Even though the shape of the Al-rich phase boundary of the C15 phase field could not be exactly fixed due to the very narrow two-phase fields with the C14 and the C36 Laves phases, the results clearly show that the maximum solubility is not reached at the stoichiometric Nb content of 33.3 at.% but at about 28e29 at.% Nb, i.e., in the direction of the Heusler phase Co2AlNb with which the C15 phase forms a two-phase field.

C14 Laves phase. The homogeneity range of the C14 Laves phase extends far into the ternary system ending at an only slightly temperature-dependent maximum Al content which increases from 52.5 at.% at 800
C to 55.5 at.% at 1250
C. On the Al-lean side of the homogeneity range of the C14 Laves phase, the minimum Al content needed to stabilize the C14 Laves phase decreases with increasing temperature from <3.7 at.% at 800
C and 2.2 at.% at 1000
C to <1.0 at.% at 1200
C. In line with the binary CoeNb phase diagram [2,3], this value will reach 0 at 1250
C. In the binary system, C14 is only stable at off-stoichiometric compositions in the range between about 36 and 37 at.% Nb. Even though the homogeneity range for Nb in the ternary system is significantly larger, the same trend is observed in the ternary C14 Laves phase and the Al-rich end of its phase field is at an offstoichiometric Nb content of 36  1 at.% for all investigated temperatures. At low Al contents between about 5 and 10 at.%, the C14 phase field also extends in the direction of the Co corner forming a narrow two-phase field with the C15 Laves phase. Due to the narrowness of this two-phase field, the exact width is difficult to determine experimentally. From the few samples containing C15 þ C14 Laves phase, the width of the two-phase field can be concluded to be below 1 at.%. In some cases both phases could not even be distinguished by EPMA but only by XRD; cf. Tables 2e6.

2198

O. Dovbenko et al. / Intermetallics 18 (2010) 2191e2207

Table 6 Alloy and phase compositions as analysed by EPMA/SEM-EDX and lattice parameters determined by XRD after heat treatments at 1250
C. Alloy no.

Analysed alloy composition

1

76.8

1.0

3

64b

1.5b

10 15

66b 64.8

16

63.2

at.% Co

17

60.4 b

Heat treatment time [h]

Phases

22.2

240

34.5b

168

5b 10.2

29b 25.0

168 72

10.1

26.7

168

Coa C36a C15 C14 C15 C36 Co2AlNb C15 C36 Co2AlNb C14 Co2AlNb C14 Co2AlNb C14 Co2AlNb C36 Co2AlNb C14 Co2AlNb C14 Co7Nb6 C14 CoAl Co2AlNb C14 Co7Nb6 Nb2Al C14 NbAl3 Nb2Al

at.% Al

10.4 b

at.% Nb

29.2 b

at.% Co

168

18

60

10

21

54.9

12.5

32. 6

48

24

58.7

25.7

15.6

168

b

b

30

b

240

25

50

25

26

32.0

24.9

43.1

168

28

40b

35b

25b

72

29

17b

35b

48b

72

66.5

32.5

72

33

a b c d

1.0

Phase composition

25

240

92.4  0.2 74.5  0.1

65.4  0.2 54.8  0.2 63.9  0.1c 63.9  0.1c 52.4  0.2 61.4  0.1 51.1  0.3 60.2  0.3 51.9  0.3 54.9  0.1 50.7  0.2 65.3  0.1 56.7  0.1 49.0  0.7 50.5  0.1 34.2  0.2 29.5  0.2 31.7  0.2 ed 49.6  0.4 24.4  0.4 19.8  0.5 4.9  0.9 9.2  0.3 0.3  0.2 0.5  0.2

Lattice parameters

at.% Al 2.4  0.1 0.8  0.1

8.1  0.1 27.9  0.1 8.4  0.1c 8.4  0.1c 27.0  0.1 8.9  0.1 25.7  0.2 9.1  0.1 24.2  0.4 12.1  0.2 24.3  0.3 8.1  0.1 29.8  0.1 16.8  0.6 24.7  0.3 28.4  0.2 21.3  0.2 34.1  0.1 ed 33.2  0.2 39.4  0.9 32.0  0.6 27.6  0.7 55.5  0.3 73.2  0.5 40.2  0.6

at.% Nb

a [Å]

5.1  0.2 24.7  0.1

3.592 4.791 6.809 4.829 6.766 4.778 5.881 6.763 4.784 5.909 4.806 5.933

26.5 17.3 27.7 27.7 20.6 29.7 23.2 30.6 23.9 33.0 25.0 26.6 13.5 34.2 24.8 37.4 49.2 34.2 ed 17.2 36.2 48.2 67.5 35.3 26.5 59.3

                 

0.1 0.1 0.1c 0.1c 0.1 0.1 0.2 0.2 0.2 0.2 0.2 0.1 0.1 0.2 0.2 0.1 0.2 0.3

      

0.4 0.9 0.3 0.5 0.4 0.5 0.7

4.850 5.956 4.771 5.840 4.874 5.956 4.960 4.982 4.962 2.904 5.887 5.005 5.035 9.893 5.039 3.842 9.925

c [Å] 15.64 7.850 15.56

15.57 7.792

7.859 15.54 7.910 8.007 26.92 8.016

8.099 27.29 5.163 8.291 8.617 5.179

Alloy was (liquid + C36) at heat treatment temperature; Co has crystallized during cooling to room temperature. Nominal composition. C15/C36 only distinguishable in XRD. Phase too small for EPMA analysis.

In Tables 2 and 3, it is also indicated that some samples contained small amounts of metastable C14 phase after heat treatments at 800 or 1000
C. This effect is well known from the binary system [2]. The transformation of metastable C14 phase, which forms on cooling when crossing the C14 stability range, to the equilibrium phases C15 þ Co7Nb6 is a very slow process at low temperatures. C36 Laves phase. Like the C14 Laves phase, the C36 Laves phase was found to exist in the ternary system at all investigated temperatures including 800
C what is in contrast to the isothermal section of Burnashova et al. [14] (Fig. 1). The C36 Laves phase could be clearly identified in various alloys forming two- or three-phase equilibria with the Co solid solution, Nb2Co7, CoAl, Co2AlNb, and the C15 Laves phase depending on composition and temperature. Phase equilibria with the C14 Laves phase can not occur because at all temperatures a two-phase field C15 þ Co2AlNb exists as mentioned above. At 800
C, the maximum Al content of the C36 phase is 4 at.%. This value continuously increases with temperature reaching 8.4 at.% at 1250
C. On the Al-lean side, the C36 phase field is separated from the binary CoeNb system by a C15 þ Nb2Co7 twophase field at 800
C. Increasing temperatures result in a continuous shift of the phase boundary of the C36 phase field towards the bounding binary CoeNb system, and above 1000
C, the phase field is connected with the binary system. This is nicely shown by alloy 1 (Coe1Ale23.5Nb), which is two-phase C15 þ Nb2Co7 at 800
C, but contains the C36 Laves phase at all higher temperatures (Table 1 and Figs. 2e6). As in the binary system, the Nb content of the ternary C36 Laves phase is always significantly below the stoichiometric value of 33.3 at.% ranging from about 24 to 28e29 at.% at

all investigated temperatures. As in case of the C14 Laves phase, the C36 phase forms a narrow two-phase field with the C15 phase the width of which is estimated to be below 1 at.%. 3.1.2. Crystallographic parameters C14 Laves phase. The shape of the C14 phase field indicates that Al atoms preferentially substitute for Co atoms leaving the Nb content more or less unchanged. Since the Al atoms are larger than the Co atoms, an increase of the lattice parameters is to be expected on alloying with Al what is confirmed by the results of the XRD experiments shown in Fig. 8. For clarity, only data of samples from the Nb-rich boundary of the C14 phase field with a constant Nb content of 36  1 at.% are plotted in the figure including the lattice parameters of the binary C14 Laves phase reported in Ref. [2]. It is noted that both lattice parameters a and c increase approximately linearly with the Al content with the exception of the values for the Al contents higher than 50 at.% which show a negative deviation in case of a and a positive deviation in case of c from the linear behaviour. Negative and positive deviations are such that they eliminate each other with respect to the volume per atom (Fig. 8c). The reason for this behaviour is not yet understood. The twelve atoms of the hexagonal unit cell of a prototypic C14 Laves phase AB2 are distributed over three crystallographically independent sites with different symmetry which are the 4f sites for the bigger A atoms and the 2a and 6h sites for the smaller B atoms. The B atoms are arranged in a tetrahedral network with infinite chains of double tetrahedra forming trigonal bipyramids along the c axis of the hexagonal lattice, see, e.g., Ref. [25]. The top and bottom atom sites of the bipyramids are the 2a sites and those

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Table 7 Crystallographic data and composition ranges of all phases occurring in the investigated temperature range [2,5e7,9,10], and solid solubility ranges for the third component in the binary phases and the Co solid solution (values for the composition ranges and solubilities are for a temperature of 1150
C). Phase

Structure type

Pearson symbol

Space group

Struktur-bericht designation

Binary composition range [at.%]

Solubility for the third component [at.%]

aCo

Cu

cF4

Fm3m

A1

Nb

W

cI2

Im3m

A2

0e15 Al [5] 0e4.3 Nb [2] 0e9 Al [10] 0e3.5 Co [2]

3.7 Nb 13 Ala n.d. n.d.

CoAl Co2Al5 Z (CoAl3)

CsCl Co2Al5 e

cP2 hP28 e

Pm3m P63/mmc Cecentered monoclinic

B2 D811 e

29e53 Al [5] 71e72 Al [6,7] 74.5 Al [7]

9.2 Nb 0.1 Nb n.d.

MeCo4Al13b OeCo4Al13b Co2Al9b

Co4Al13 e Co2Al9

mC102 oP102 mP22

C2/m Pmn21 P21/c

e e D8d

75.8 Al [7,9] 76e76.5 Al [6,7] 81.2 Al [7]

n.d. n.d. n.d.

Nb2Co7b NbCo2 (C15) NbCo2 (C14) NbCo2 (C36) Co7Nb6 (m)

e MgCu2 MgZn2 MgNi2 Fe7W6

mC18 cF24 hP12 hP24 hR39

C2/m Fd3m P63/mmc P63/mmc R3m

e C15 C14 C36 D85

22.1e22.3 Nb [2] 26.0e34.6 Nb [2] 35.8e37.4 Nb [2] 24.5e25.5 Nb [2] 47.5e54 Nb [2] d

2 Ala 6.5 Al 53.5 Al 6.6 Al 32.8 Al

NbAl3 Nb2Al Nb3Al

Al3Ti CrFe Cr3Si

tI18 tP30 cP8

I4/mmm P42/mnm Pm3n

D022 D8b A15

25.0e25.3 Nb [10] 65e71Nb [10] 78e82Nb [10]

1.3 Co >11.4 Co n.d.

Co2AlNb

Cu2AlMn

cF16

Fm3m

L21

e

e

c

n.d.: not determined. a Estimated from the present data. b Not stable at 1150
C; data are for 800
C. c Not stable in the binary system at 1150
C; binary composition range is for 1350
C. d Estimated from data in [2].

of the basal triangles are the 6h sites. Single-crystal structure refinements of the off-stoichiometric, Nb-rich binary NbCo2 C14 Laves phase have shown that antisite atoms are the only type of structural defect (i.e., there are no constitutional vacancies) and that the excess Nb atoms replace Co only on the 2a sites leaving the 6h sites untouched [25,26]. For the present ternary C14 Laves

phase, Rietveld refinements were performed for Nb-rich samples with different Al contents. As in the binary case the resulting site occupation factors clearly show that the excess Nb atoms preferentially occupy the 2a sites. However, in contrast to the binary case a small amount of the excess Nb atoms seems also to occupy the 6h sites.

Fig. 2. Isothermal section of the CoeAleNb phase diagram at 800
C as obtained from the present EPMA and XRD data (squares mark measured phase compositions and circles give the overall compositions of the samples where black, half-filled, and white circles indicate three-, two-, and single-phase alloys, respectively; triangles mark the overall composition of samples which were only studied by XRD).

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O. Dovbenko et al. / Intermetallics 18 (2010) 2191e2207

Fig. 3. Isothermal section of the CoeAleNb phase diagram at 1000
C as obtained from the present EPMA and XRD data (crosses mark data from the diffusion couple, other symbols as in Fig. 2).

As mentioned above, it is evident from the shape of the C14 phase field that Al atoms will preferentially replace Co atoms. Rietveld analyses show that e similar but less pronounced than in case of the excess Nb atoms e Al atoms preferentially occupy 2a sites. This is in agreement with recent results for the closely related ternary C14 Laves phase Nb(Cr,Al)2 where both Rietveld

refinements of experimental XRD patterns as well as calculations by a statistical mechanics approach reveal that Al prefers to occupy the 2a sites [27]. Summarizing the issue of site occupation in the ternary C14 Laves phase, Nb as well as Al atoms, which both are larger atoms than Co, tend to replace Co not evenly on the two possible sites with

Fig. 4. Isothermal section of the CoeAleNb phase diagram at 1150
C as obtained from the present EPMA and XRD data (black square marks the composition of the liquid obtained from the liquidus surface [22]; other symbols as in Fig. 2).

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2201

Fig. 5. Isothermal section of the CoeAleNb phase diagram at 1200
C as obtained from the present EPMA and XRD data (black square marks the composition of the liquid obtained from the liquidus surface [22]; other symbols as in Fig. 2).

2a and 6h symmetry, but with a more or less pronounced preference for the 2a sites. C36 Laves phase. The lattice parameters a and c as well as the volume per atom are shown as a function of the Al content for a constant Nb content of 24.5  0.5 at.% in Fig. 9. The values for the

binary sample are taken from Ref. [28]. As in case of the C14 Laves phase, all lattice parameters continuously increase on replacing Co by Al. The unit cell of the C36 Laves phase contains five crystallographically independent atom sites. In the ideal case of a stoichiometric AB2 composition, all A atoms are on sites with one of the two

Fig. 6. Isothermal section of the CoeAleNb phase diagram at 1250
C as obtained from the present EPMA and XRD data (black squares mark the compositions of the liquid obtained from the liquidus surface [22]; other symbols as in Fig. 2). Alloy 1 (Coe1.0Ale22.2Nb) contained liquid (þC36) at 1250
C which decomposed during cooling into aCo þ C36, see Table 6.

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O. Dovbenko et al. / Intermetallics 18 (2010) 2191e2207

Fig. 7. Back-scattered electron micrographs of the microstructures of four representative three-phase alloys: a) Coe10.2Ale10.4Nb (alloy 13) after 1000 h at 800
C, Co (grey) þ C36 (light) þ CoAl (dark); b) Coe33.2Ale49.1Nb (alloy 29) after 240 h at 1000
C, C14 (dark) þ Co7Nb6 (grey) þ Nb2Al (light); c) Coe66.5Ale32.5Nb (alloy 33) after 72 h at 1250
C, C14 (grey) þ NbAl3 (dark) þ Nb2Al (light); d) Coe67.6Ale11.5Nb (alloy 34) after 136 h at 1150
C, CoAl (grey) þ Co2Al5 (dark) þ NbAl3 (light).

different symmetries 4e and 4f. In case of the Co-rich binary C36 Laves phase, which has a strongly off-stoichiometric composition of about Coe25Nb, the excess Co atoms preferentially occupy the 4e sites approximately twice as much as the 4f sites [28]. Rietveld refinements of two ternary C36 alloys with the same Nb content and about 1 and 5 at.% Al confirm the preferential occupation of the 4e site by Co for the sample with 1 at.% Al, whereas for the higher Al content the excess Co atoms are distributed about equally on both original Nb sites 4e and 4f. It is likely that this behaviour is a result of the Al addition as the Rietveld analysis shows that all Co sites of the C36 lattice remain completely occupied by Co atoms and Al occupies the remaining Nb sites 4e and 4f in an approximate ratio of 3:1, i.e., the Al atoms e like the excess Co atoms in the binary case e prefer to occupy the 4e sites. C15 Laves phase and relation to the C14 and C36 polytypes. As in the case of the ternary C14 and C36 Laves phases, the addition of Al increases the lattice parameter of the cubic C15 Laves phase. Fig. 10 shows the volumes per atom as a function of the Al content for an approximately stoichiometric Nb content of 33.3 at.% covering the stability ranges of the C15 and C14 Laves phases. The continuous increase of the volume per atom is unaffected by the change of the

structure type from C15 to C14. A similar behaviour is observed in Fig. 11, where instead of the dependence on the Al content the volume per atom as a function of the Nb content is shown for the C36, C15, and C14 Laves phases. For a constant Al content of 5 at.%, the volume per atom continuously increases with increasing Nb content across the phase boundaries from C36 to C15 and from C15 to C14. A comparison of Figs. 10 and 11 indicates that replacing Co by Nb results in a significantly stronger increase of the volume per atom (about 0.060 Å3/at.%Nb) than replacing Co by Al (about 0.037 Å3/at.%Al). The atomic radii of Al and Nb in the pure metals (i.e., in fcc-Al and bcc-Nb) are approximately the same (1.43 Å), but their average covalent radii differ significantly (Al: 1.21(4) Å, Nb: 1.64(6) Å [29]). The observed higher effectiveness of Nb in expanding the unit cell dimensions therefore is in accordance with the fact that chemical bonding effects play an important role in Laves phases, see, e.g., Refs. [30,31]. 3.2. Nb2Co7 The monoclinic phase Nb2Co7 has a very small homogeneity range in the binary CoeNb system of about 0.2 at.% and decomposes

O. Dovbenko et al. / Intermetallics 18 (2010) 2191e2207

Fig. 8. Lattice parameters a and c and volume per atom of the hexagonal C14-type Laves phase as a function of the Al content along the Nb-rich boundary of the phase field (36  1 at.% Nb). The values for 0 at.% Al were taken from Ref. [2].

above 1086
C [2]. In agreement with that, Nb2Co7 was observed in the ternary system only in samples annealed at 800 and 1000
C. The maximum Al content in this phase found at 800
C is 1.7 at.%, and at 1000
C the solubility is below 1 at.% Al as can be concluded from alloy 1 (Coe1.0Ale23.5Nb), which is three-phase Co þ C36 þ Nb2Co7 at 1000
C. In contrast to the isothermal section reported by Burnashova et al. [14] (Fig. 1), the present results clearly show that Nb2Co7 does not form equilibria with the Heusler phase Co2AlNb and CoAl but is separated from these phases by a two-phase field C36 þ Co solid solution. Due to the low solubility for Al, the lattice parameters of Nb2Co7 with 1.0 and 1.7 at.% Al (Table 2) are very similar to those for the binary case [2] and the unit cell volumes differ from that of the binary phase by less than 0.2%.

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Fig. 9. Lattice parameters a and c and volume per atom of the hexagonal C36-type Laves phase as a function of the Al content for a constant Nb content of 24.5  0.5 at.% (the values for 0 at.% Al were taken from Ref. [2]).

exists down to Nb contents below 13 at.% and up to Al contents above 33 at.%. The maximum Nb content at all temperatures is 25 at.%. The fcc crystal structure of the Heusler phase is of the L21 Cu2AlMn structure type. Lattice parameters reported in the literature are 5.946 Å [15] and 5.935 Å [16], in both cases the exact composition of the investigated samples was not given. As Fig. 12 shows, the lattice parameter of the Heusler phase increases

3.3. The Heusler phase Co2AlNb Burnashova et al. [14] reported the occurrence of the Heusler phase Co2AlNb at 800
C and plotted it as line compound in their isothermal section (Fig. 1). They did not perform composition analyses. The present investigations confirm the existence of this phase, which was found in the entire investigated temperature range. However, our results show that the Co2AlNb Heusler phase is not a line compound but has an extended homogeneity range which increases with increasing temperature. From the stoichiometric composition Co2AlNb, the phase field extends in the direction of (Al þ Co)-enriched compositions. At 1250
C, the Heusler phase

Fig. 10. Volume per atom of C15 and C14 Laves phase with a constant, nearly stoichiometric Nb content of 33.5  0.6 at.% as a function of the Al content (the value for 0 at.% Al was taken from Ref. [2]).

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Fig. 11. Volume per atom of C36, C15 and C14 Laves phase with a constant Al content of 5.0  0.4 at.% as a function of the Nb content.

linearly as a function of the Nb content along the Al-lean side of the homogeneity range from a ¼ 5.840 Å for 13.0 at.% Nb to the maximum value of a ¼ 5.959 Å for 25.0 at.% Nb. The Heusler phase does not melt congruently but forms in a slow peritectic reaction. Its slow formation has already been reported in the literature, as has been discussed in Section 1 ‘Introduction’. At the stoichiometric Co2AlNb composition, the C14 Laves phase solidifies as the primary phase. The alloy Coe25Ale25Nb contains significant amounts of C14 phase in the as-cast state and long heat treatment times are needed to dissolve this phase. As an example, Fig. 13 shows the XRD patterns of the alloy Coe25Ale25Nb after heat treatment at 1250
C for 48 h and 240 h. Even after the longer annealing time, some very small peaks of the metastable C14 phase are still visible, though a comparison of the two patterns shows that the remnant amount of C14 phase has further decreased after the longer heat treatment time. Due to the presence of non-negligible amounts of metastable, inhomogeneous C14 phase, the results after heat treatments of the alloy Coe25Ale25Nb at 800 to 1200
C were not used for drawing the respective isothermal sections in Figs. 2e5. 3.4. The m phase Co7Nb6(Al) The binary m phase Co7Nb6 has an extended homogeneity range from about 47 to 56 at.% Nb. The solubility for Al in this phase

Fig. 12. Lattice parameter a of the cubic Heusler phase as a function of the Nb content along the Al-lean side of the homogeneity range. Data are taken from alloys heattreated at 1150, 1200 and 1250
C.

Fig. 13. XRD patterns of the alloy Coe25Ale25Nb after heat treatment at 1250
C for 48 h and 240 h showing the presence of small amounts of metastable C14 phase (lCoKa1 ¼ 1.78897 Å, ‘H’: Heusler phase).

amounts to about 32  1 at.% for all investigated temperatures. As in case of the C14 Laves phase, the phase field of the m phase primarily extends parallel to the CoeAl axis. Data for both the Nb-poor and the Nb-rich boundary of the homogeneity range were only measured at 1150
C (Fig. 4). The results indicate that the width of the phase field strongly reduces with increasing Al content. The lattice parameters a and c and the resulting volume per atom of the close-packed hexagonal lattice of the m phase along a constant Nb content of 48  1 at.%, i.e., along the Nb-lean boundary of the phase field are shown as a function of the Al content in Fig. 14. A linear fit to the volume per atom data in Fig. 14c gives a slope of 0.042 Å3/at.%Al which is approximately 10% larger than in case of the Laves phase. 3.5. The m0 phase As mentioned in Section 1, Hunt and Raman [20] studied the occurrence of m phases in a series of ternary transition-metalbased systems and frequently observed additionally a so-called m0 phase, which is also included in the 800
C section of Burnashova et al. (Fig. 1) [14] in the composition range between the m phase and Nb2Al. However, the existence of the m0 phase remained somehow doubtful as neither the exact composition was measured nor its crystal structure could be determined in both studies. In the present investigation, the existence of this phase could not be confirmed. In all cases, XRD patterns of alloys in the respective composition range could clearly be indexed with the peaks of the m phase and Nb2Al. However, after heat treatments at 1200 and 1250
C several Nb-rich alloys contained a phase of the fcc NiTi2 structure type with Nb contents similar to that of the m phase. The occurrence of such a NiTi2-type phase with lattice parameter of about a ¼ 11.3 Å is well known from a number of different binary and ternary transition-metal-based systems and it was shown that this phase is stabilized by small amounts of impurities, see, e.g., Refs. [32e35]. Samples containing this impurity-stabilized phase were not used for the present phase diagram evaluation. As described in Section 2, more careful heat treatments at 1200 and 1250
C of several new samples were performed resulting in material free or nearly free of the NiTi2type phase.

O. Dovbenko et al. / Intermetallics 18 (2010) 2191e2207

Fig. 14. Lattice parameters a and c and the resulting volume per atom of the hexagonal m phase as a function of the Al content along a constant Nb content of 48  1 at.% (the values for 0 at.% Al were taken from Ref. [2]).

3.6. Nb2Al and NbAl3 The Nb-rich phase Nb2Al was not investigated in detail. The 1150
C section shows that this phase has a high solubility for Co of at least 11.4 at.%. Whereas the 800
C isothermal section of Burnashova et al. [14] indicates a low solubility of approximately 2 at.% Co, the present investigations show that the solubility for Co is at least 5 at.% (alloy 29) at 800
C. The maximum Co content as well as the shape of the Nb-rich boundary of the ternary phase field was not determined. The lattice parameters of the tetragonal lattice decrease by the addition of Co from a ¼ 9.943 Å and c ¼ 5.186 Å, as given for the binary, nearly stoichiometric phase in Ref. [36], to a ¼ 9.848 Å and c ¼ 5.145 Å for 11.4 at.% Co. In contrast to Nb2Al, the solubility for Co in NbAl3 is comparably small, the maximum value in the investigated temperature range is 1.5  0.2 at.% Co. The lattice parameters of the body-centred tetragonal unit cell of all investigated samples have the same values of a ¼ 3.841 Å and c ¼ 8.617 Å within 0.002 Å, and these values are similar to the lattice parameters a ¼ 3.841 Å and c ¼ 8.609 Å given in the literature for the binary compound [37]. 3.7. The Al-rich CoeAl phases The phase Co2Al5, which melts at 1188
C in the binary system [6], was found in ternary samples after heat treatments at 800,

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1000, and 1150
C. The solubility of Nb is 0.1 at.% at all temperatures. The measured Al contents are 71.5  0.2 at.% in good agreement with the reported composition of the binary phase (Table 7). For the lattice parameters of the hexagonal lattice of Co2Al5, values of a ¼ 7.678  0.002 Å and c ¼ 7.615  0.002 Å were obtained, which, probably as an effect of the small amount of dissolved Nb, are slightly higher than the binary literature values [7] of a ¼ 7.658 Å and c ¼ 7.604 Å. Alloys in the Al corner containing phases with Al compositions beyond Co2Al5 were not prepared in the present investigations. On the basis of the information from the binary CoeAl [6] and AleNb [10] systems and supposing that the solubility for Nb in the Al-rich CoeAl intermetallic phases Co2Al9 (occurs only in the 800
C section), O-Co4Al13 and M-Co4Al13 (only 800 and 1000
C section) and Z (CoAl3) (800e1150
C sections) is very low, respective phase equilibria were tentatively drawn with dotted lines in Figs. 2e4. Some additional information supporting the isothermal sections as drawn here comes from results of the NbCo2/Al diffusion couple, which had been annealed at 1000
C. Fig. 15 shows a back-scattered electron image of the diffusion couple after 24 h at 1000
C. In diffusion couples in ternary systems single- as well as two-phase layers can form depending on the diffusion path through the system. If the diffusion path crosses a two-phase field from one to the other adjoining three-phase field, it cuts the tie lines and a locally equilibrated two-phase layer can form in the diffusion couple as has been frequently discussed in the literature; see, e.g., Refs. [38e42]. In the present case the Al-rich corner of the isothermal section at 1000
C shows a succession of several, partially very narrow, two-phase fields separated by three-phase fields connecting the series of binary CoeAl intermetallic phases with the NbeAl phase NbAl3 and the C14 Laves phase (Fig. 3). The micrograph of the NbCo2/Al diffusion couple in Fig. 15 shows that the diffusion path obviously has crossed this series of two- and three-phase fields. Four two-phase zones have formed confirming the existence of the phase equilibria CoAl þ C14-NbCo2, CoAl þ NbAl3, Co2Al5 þ NbAl3, and Z (CoAl3) þ NbAl3. Due to the high annealing temperature compared to the melting point of Al, a certain amount of the liquid Al has evaporated during the experiment, and because of the high diffusivity of the Al atoms at this temperature, the equilibration process obviously has already reached such an advanced state that neither Al nor the Al-richest phases O-Co4Al13 and M-Co4Al13 were found in the diffusion couple. On the other hand, the Al-poorest two-phase zone CoAl þ C14-NbCo2 with measured Al contents of 40 and 50 at.% has a thickness of only about 5 mm and the annealing time obviously was not long enough to form layers of visible thickness of phases with less than 40 at.% Al. In the last two-phase zone, the particles are large enough (>1 mm in diameter) to allow composition analyses of the two phases. For the Z phase, a very low solubility for Nb of <0.1 at.% and an Al content of 74.3 at.% e in good agreement with the value of 74.5 at.% Al reported for the binary Z phase (Table 7) e were obtained. The second phase is NbAl3 with 1.0 at.% Co in solution and a Nb content of 25.6 at.%. 3.8. CoAl The binary, cubic B2-type phase CoAl has a broad homogeneity range which according to Ref. [5] extends from about 40 to 53 at.% Al at 800
C and broadens to 25 to 53 at.% Al at 1250
C. Although the shape of the ternary phase field was not studied in detail in the present investigations, a significant solubility for Nb of up to 9 at.% was observed. CoAl and the L21 Heusler phase Co2AlNb form a two-phase field which becomes smaller with increasing temperature. From other ternary transition-metal-based systems with coexisting B2 and L21 phases like FeeAleTi [43e47] or

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O. Dovbenko et al. / Intermetallics 18 (2010) 2191e2207

Fig. 15. Back-scattered electron image of a NbCo2/Al solid/liquid diffusion couple after 24 h heat treatment at 1000
C. Four successive two-phase zones have formed. A very thin (about 5 mm) diffusion zone of CoAl þ C14 Laves phase is only visible in the magnification on the left side.

CoeAleTi [46,48,49], it is known that there is a continuous transition (i.e., no two-phase field) between the B2 and L21 phase above a certain tricritical point, which for the CoeAleTi system is at temperatures between 1000
C (two-phase field, no continuous transition) and 1100
C (B2 and L21 phase fields in contact, continuous transition). From the presently available experimental data, it can not be concluded whether the two-phase field between the Co2AlNb Heusler phase and CoAl will disappear at 1150
C or higher temperatures. Therefore, the respective phase equilibria are drawn with dotted lines in Figs. 2e6. The lattice parameter of the cubic B2 phase, which in the binary system has a maximum value of a ¼ 2.862 Å at the stoichiometric composition [50], increases with the addition of Nb, e.g., for CoAl containing 43.5 at.% Al and 6.3 at.% Nb (alloy 30, 1150
C) a value of a ¼ 2.904 Å was measured. 3.9. The Co solid solution The solubilities of both Al and Nb in the Co solid solution are similar to those in the binary systems and increase with temperature. The crystal structure was found to be the high-temperature fcc aCo modification in all investigated samples (alloys 1, 4, 13, 20). As expected, Al and Nb additions both increase the lattice parameter. 3.10. The liquid phase With increasing temperature, the first liquid occurs in the Al corner of the system and its phase field continuously grows into the ternary system in the investigated temperature range. The black squares in Figs. 4e6 marking the maximum extension of the liquid phase field are taken from an investigation of the liquidus surface

on the basis of differential thermal analysis and microstructural analysis of the as-cast alloys [22]. The 1250
C isothermal section (Fig. 6) shows another liquid phase field which is located in the Co-rich corner and results from the deep eutectic between the Co solid solution and the Laves phase. As in the case of the phase field of the Al-rich liquid, its shape was estimated from the liquidus surface studied in Ref. [22].

4. Summary Five isothermal sections of the CoeAleNb system in the temperature range 800e1250
C have been established and lattice parameters of all phases have been measured in dependence on composition. In contrast to the 800
C section reported by Burnashova et al. [14] which was the only available isothermal section before, the present investigations showed that the C36 Laves phase exists already at 800
C in the ternary system. The phase fields of both the C14 and C36 Laves phase, which appear in the 800
C section as ternary phases, are connected with the binary CoeNb system at higher temperatures, i.e., the addition of Al stabilizes the C36 and C14 Laves phase polytypes to lower temperatures. The actually only ternary phase of the CoeAleNb system is the slowly forming Heusler phase Co2AlNb. It is not a line compound as indicated by the work of Burnashova et al. [14], but certain amounts of Nb, which increase with temperature, can be replaced by Co and Al. This results in a phase field covering the composition range 50e60 at.% Co, 25e33 at.% Al, and 13e25 at.% Nb. The m0 phase reported in Refs. [14,20] was not observed in the present investigations. However, a number of Nb-rich samples annealed under insufficiently clean conditions were found to

O. Dovbenko et al. / Intermetallics 18 (2010) 2191e2207

contain an impurity-stabilized, NiTi2-type phase with Nb contents similar to those of the m phase. Nb2Co7 does not form equilibria with the Heusler phase Co2AlNb and CoAl as indicated by the work of Burnashova et al. [14], but is separated from these phases by a two-phase field C36 þ Co solid solution. The solubility for the ternary element is below 0.1 at.% in the phases Co2Al5 and Z (CoAl3) and takes low values of 1e2 at.% in Co7Nb2 and NbAl3. In contrast to these phases, CoAl, Nb2Al and the m phase Co7Nb6 have high solubilities for the third element which amount to more than 9 at.%, more than 10 at.%, and about 32 at.%, respectively. Whereas the maximum solubility observed for Al in the C15 and C36 Laves phases is below 10 at.%, the phase field of the C14 Laves phase extends far into the ternary system up to Al contents beyond 50 at.%. Like in the binary system, the Nb contents of the C14 and C36 Laves phases remain off-stoichiometric in the ternary system. No equilibrium exists between the C14 and C36 Laves phase, as the C15 Laves phase forms a two-phase field with the Heusler phase. The results indicate that the widths of the two-phase fields between C15 and C14 and between C15 and C36 are below 1 at.%. As a result of the close relationship between the crystal lattices of the three Laves phase polytypes, the volume per atom of the Laves phase increases continuously as a function of Al or Nb content independent of the structural change from C36 to C15 or C15 to C14. The effect of Nb additions on the lattice parameters is much stronger than that of Al additions in the Laves phases as well as in all other investigated intermetallic phases of the system. On deviating from the ideal AB2 stoichiometry of the Laves phases, both types of defect atoms, i.e., Nb and Al in case of C14 and Co and Al in case of C36, prefer to occupy the same atom site. For the Nb-rich C14 Laves phase, Nb as well as Al replaces Co preferentially on the 2a sites, and in case of the Co-rich C36 Laves phase, both Co and Al tend to occupy the Nb 4e sites. Acknowledgements The authors would like to thank Mr. R. Staegemann and Mr. M. Kulse for the preparation of alloys, Mr. G. Bialkowski for sample cutting, Mrs. A. Bobrowski and Mrs. H. Bögershausen for preparation of metallographic sections, Mrs. I. Wossack and Mr. U. Wellms for EPMA analyses, and Dr. C. He for providing three ternary samples. Financial support by the Max Planck Society within the framework of the inter-institutional research initiative ‘The Nature of Laves Phases’ is gratefully acknowledged. References [1] Hari Kumar KC, Ansara I, Wollants P, Delaey L. J Alloys Compd 1998;267:105. [2] Stein F, Jiang D, Palm M, Sauthoff G, Grüner D, Kreiner G. Intermetallics 2008;16:785. [3] Okamoto H. J Phase Equilib Diffus 2010;31:94. [4] The nature of Laves phases. An inter-institutional research project of the Max Planck Society, http://Laves.mpie.de.

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