Liquid-solid equilibria in the aluminium-rich corner of the AlCrNi system

ELSEVIER

Liquid-solid

Journal of Alloys and Compounds

233 (1996) 246-263

equilibria in the aluminium-rich system E. Rosell-Laclau,

Institut National Polytechnique

M. Durand-Charre,

corner of the Al-Cr-Ni M. Audier

de Grenoble, Laboratoire de Thermodynamique et Physico-Chimie ENSEEG BP7.5, 38402 Saint Martin d’H&es, France Received

9 February

1995: in final form 20 June

Mttallurgiques,

UA CNRS 29,

1995

Abstract The aspect of the phase diagram related to liquid-solid equilibria in the subsection AlCr-AlNi-Al was investigated. It is shown firstly that close to the aluminium corner, a ternary eutectic transformation occurs at 634 “C. Crystallization paths were studied from different liquid alloys with compositions close to those of the binary Al-Ni or Al-Cr systems. The different liquidus phase fields appear to be related to binary Al-N1 and Al-Cr phases but also to the formation of ternary phases. Transmission electron microscopy studies show that each of these ternary phases has several structures related by polytypism or superstructure ordering. The main feature of this phase diagram is that eutectic valley connects a quasi-binary eutectic cKr/AlNi to the ternary eutectic Al,Ni/f?Al,Cr/Al; several pseudo-peritectic transformations interrupt this monovariant line on either side. Keywords: Al-Cr-Ni

system;

Al-rich

corner:

Ternary

phase

diagram

-

1. Introduction

The Al-Cr-Ni phase diagram has been the subject of a number of studies, mainly of the nickel-rich corner. Most investigations were limited to the AlNiCr-Ni subsection because this is the model system for superalloys. Two documented review articles on these results were published by Merchant and Notis in 1984 [l] and by Rogl in 1991 [2]. The aluminium-rich corner has been little investigated; in both these review articles the same diagram of the primary aluminiumrich phase fields defined on high temperature isothermal sections was proposed on the basis of microstructural analyses of as-cast alloys carried out by Taylor and Floyd [3]. For the quasi-binary AlNi-Cr system interest was focused on the pCr/pAlNi eutectic formed at 1450 “C with an atomic proportion of about 66% AlNi [2]. In addition to this, the formation of a ternary orthorhombic phase of composition close to Al,,Ni,Cr,, was mentioned recently by Colin-Urtado

PI. On account of the binary Al-Ni and Al-Cr phase diagrams the extent of the primary aluminium phase field in Al-Cr-Ni is limited by the eutectic composi092%8388/96/$15.00 0 1996 Elsevier SSDZ 0925-8388(95)01988-X

Science

S.A. All rights

reserved

tion Al/Al,Ni [5] on one side and by the peritectic composition Al/BAl,Cr [6] on the other side. Thus it may be assumed that a eutectic or peritectic transformation takes place for a composition close to the binary eutectic Al/Al,Ni [6], as those occurring in the primary aluminium phase field of the Al-Ni-Fe [7-91 and Al-Ni-Ti [lo] systems. Our first aim was to clarify this point. We were also interested in an identification of several intermetallic phases existing in this part of the Al-Ni-Cr phase diagram, in particular, those related to the binary Al-Cr compounds which are formed by peritectic reactions from ,.$Al,Cr, to BAl,Cr [6]. However, because of the difficulties we have had in interpreting the results of the present work where several structures of ternary compounds found in the Al-AlNi-Cr subsection and never mentioned previously were not consistent with the binary Al-Cr structures, it has been necessary to re-examine carefully the crystallographic data available on the binary Al-Cr system. As a result of this recently published study [ll], a new structure was proposed for Al,Cr similar to that of the hexagonal pAl,Mn phase [12,13], and A&r (or qAl,,Cr,) was found to be orthorhombic instead of monoclinic as had been previously

E. Resell-Laclau

et al. I Journal of Alloys and Compounds

proposed [14]. The structures of other compounds, BAl,Cr and ‘/Al,Cr,, were confirmed. The results reported in this article are related to the three following points: (i) the determination of a ternary transformation close to the binary eutectic Al/Al,Ni; (ii) the determination of Al-Cr-Ni primary phase fields; (iii) the structural characterization of new ternary Al-Cr-Ni phases.

2. Experimental procedure Three: sets of alloys, noted A, B and C were prepared (Table 1). Taking account of a previous work [15], the compositions of the A alloys were chosen close to a ternary invariant transformation: four alloys were prepared by melting high purity constituent elements (99.9%) in an inductive cold crucible under argon atmosphere. Seven other alloys (B and C sets in Table 1) were prepared by the addition of chromium and aluminium to prealloyed Al,Ni or Al,Ni, compounds. The two distinctive alloy sets B and C are respectively related to compositions close to those of the binary Al-Ni and Al-Cr systems. The c:ompositions of the phases were first determined approximately by X-ray energy dispersive spectroscopy in a scanning electron microscope (SEM) and then accurately by X-ray wavelength spectroscopy on a Cameca SX50 microprobe (EPMA). The atomic compositions were established using the ZAF modified PAP correction program [16]. Different phase structures were identified by electron diffraction using a transmission electron microscope (TEM). Specimens were prepared as thinned foils by ion milling or as fine powders by crushing small pieces of alloy ingots, which were deposited onto copper grids coated with a carbon film. The transformation temperatures were measured by differential thermal analysis (DTA) of samples of about 1 g. The cooling and heating rates were 5 “C mini’ or less. Some experiments were stopped by Table 1 Alloy sample

compositions

(X = Fe or Ti or Cr)

Alloy

Al (at.%)

Ni (at.%)

Cr(at.%)

X (at.%)

A ,(I) A ,FCi A ,r1,

3 3 3.5

_

A,,.,,

97 96.9 96.3 96.6

2.9

0.3

Bl B2 B3 Cl c2 c3 c4

75 60 55 82 92 87 77.3

23 38 25 9 4 3 4.3

0.1 0.2

2 2 20 9 4 10 18.4

2.73 (1996) 246-263

247

quenching at the appearance of the first solid phase in order to determine the first compound to be formed. As all the ternary systems investigated using the A alloy samples displayed two invariant reactions in a very narrow range of composition and temperature, the accuracy of our recorded DTA thermograms was not sufficient to determine unambiguously the nature of the reaction based on such small differences in temperature (i.e. a few degrees). For this reason a particular experimental procedure was used: the inert reference for normal DTA experiments was replaced by an Al-Ni sample of composition corresponding to that of the binary eutectic so that the difference in temperature between the binary and ternary alloy samples could be recorded directly on the thermograms. In order to correct for the gradient effect in the furnace, experiments were repeated exchanging places between the sample and the Al-Ni eutectic reference. The heating process was monitored at a slow rate of 2”Cmin~‘. This resulted in a complicated thermogram in which the only significant information was what alloy melted first and how large was the temperature gap between both melting points. Directed solidification experiments were performed in a vertical Bridgman furnace with a moving rate of 4.5 cm h-‘. The ingot sizes were about 1 cm diameter by 13 cm length. Several samples prepared with different solidification and cooling rates were examined. For the clarity of microstructure descriptions, the origin of each sample is given completely, for instance S,,,, S quenched liquidus and ’ annealed at .a-“C during y days’

3. Ternary eutectic transformations The temperatures of the ternary invariant transformations were determined for four systems: Al-Ni, Al-Ni-Cr, Al-Ni-Fe and Al-Ni-Ti (Table 2). The temperature gaps between these invariant ternary points and the close binary Al-Ni eutectic point are small for the three ternary systems. In the case of the A (crj alloy (Al,,,,Ni,.,Cr,.,) it was determined that the nature of the transformation is eutectic with a melting point 5 “C lower than that of the binary Al-Ni Table 2 Phase mixing and temperatures corresponding to ternary invariant transformations in the Al-rich corner of the binary AI-Ni and ternary Al-Ni-Cr, Al-Ni-Ti systems System

Transformation

Constituent

phases

AI-Ni Al-Ni-Cr Al-Ni-Fe Al-Ni-Ti

Eutectic Eutectic Eutectic Peritectic

Al,NilAI AI,NiIBAl,Cr/AI Al,Ni/Al,FeNi/Al AI,NiIAI,Ti/AI

Transformation temperature (“C) 640 635 634 645

248

E. Resell-Laclau

et al. I Journal of Alloys and Compounds

eutectic. The microstructure of the ternary eutectic in the AC,,, alloy is rod like, as for the binary eutectic. Longitudinal and transversal sections of these rods are shown respectively in the SEM images (a) and (b) of Fig. 1. The rods correspond to the Al,Ni phase embedded in the aluminium matrix. No third phase could be observed in a sample solidified in a Bridgman furnace. However, after annealing treatment for 1 week at 600 “C, the three phases Al, A1,Ni and BAl,Cr were observed (Fig. l(c)). Considering the location of

233 (1996) 246-263

the global composition of the initial eutectic liquid in the triangle of the tie-lines joining Al, A1,Ni and expected for the BAl,Cr phases, the proportion BAl,Cr phase is small. In this case, it is commonly observed that the minor phase does form but only after long annealing treatment. For the Act+) alloy (Al,,,Ni,Fe,,,), the eutectic transformation was found to occur at 634 “C, i.e. 4°C lower than the value previously measured [7]. As in the preceding case, the observed microstructure was rod like and the formation of a third phase Al,FeNi was only observed after a long annealing treatment. For the last A(.,,, alloy (Al,,,,Ni,,Ti,,,), the measured peritectic temperature was 645 “C, i.e. corresponding exactly to the value already given in the literature [17]. In this case the rod-like microstructure was faceted and, again, the formation of a third phase Al,Ti was observed after a long annealing treatment.

4. Ternary Al-Ni-Cr

Fig. 1. Rod-like morphology of the ternary AI,Ni/BAl,Cr/Al eutectic: (a) and (b) SEM images showing respectively longitudinal and transversal sections of A1,Ni rods embedded in an Al matrix (observed in a Bridgman sample of the A(,.,, alloy (Al,, hNi2 ,Cr,, 3)); (c) SEM image showing the BAl,Cr precipitation in an A,, ,, alloy sample quenched after annealing for 1 week at 600°C.

phase fields

The binary phases which were considered are summarized in Table 3. In addition to these binary phases, several ternary phases 4, h and p were distinguished. Their structures are presented below. The transformation temperatures of B and C alloys reported in Table 4 result from DTA experiments carried out at a heating and cooling rate of 5 “C mini ’. For this cooling rate the corresponding microstructures are coarser than those observed for as cast or quenched liquidus samples and thus their constituent phases can be easily distinguished and analysed. However, on account of some solid state peritectic transformations occurring during the cooling process, it was not always possible to establish a correspondence between the measured temperatures from DTA thermograms and solidification paths. Then, the sequence of crystallization was determined from the microstructure observed in specimens quenched at different stages of transformation (Table 4). Observations of as-cast microstructures were also considered but as some phases can escape observation when nucleation is difficult (as does qA1, ,Cr, for the binary system [ll]), we made complementary observations on samples annealed for 30 min before quenching. Since the proposed Al-Ni-Cr phase diagram limited to the subsection AlCr-AlNi-Al was derived from the results presented in Table 4, various comments are made in the footnotes to this table in order to remain consistent with the proposed interpretation and to respect the Gibbs phase rule. Fig. 2 shows microstructures observed in the quenched liquidus and slowly solidified Bl (Al,,Ni,,Cr,) and B3 (Al,,Ni,,Cr,,) alloys. The primary phase to solidify was either Al,Ni, or AlNi for

E. Rose%Laclau Table

et al.

I Journal of Alloys and Compounds 233 (1996) 246-X_?

249

3

Crystallographic

data

on binary

AI-Cr

and AI-Ni

Phase

Bravais

Al HAI,Cr or AI,,Cr, nAl,,Cr, or Al,& pAI,Cr

Cubic Monoclinic (type AI,,V,) Orthorhombic Hexagonal Cubic Rhombohedral Orthorhombic Hexagonal Cubic Cubic

YAW, lAl,Cri AI,Ni A&N& AlNi rvCr

Table 4 Transformation

temperatures

phases

lattice

and crystallization

considered

Cell parameters

Fm3m C2lm Cmcm or Cmc2, P6, lmmc

a a a a a a a a a a

R3m Pnma

Pi.r?Jl Pm3m Im3m

sequences

Transformation

Bl(AI,,Ni,,Cr,)

1117,872,825,635 ‘, 1100,1010,976: 1411,1277,1151, 884d 969,907,796,686,635 876,683,635 976,893,686,650,635 1018,1004,898,723,635

R3(Al,,Ni2,Cr2,,) Cl(AI,,Ni,,Cr,) C2(AI,,Ni,Cr,) C3(AI,,Ni,Cr,,,) C4(At,,

,Ni, ,Cr,,,)

observed

temperatures

“ The fcsrmation of the eutectic t9AI,Cr/AI,Ni was not detected ” The temperature was not determined because at a temperature was not ’ The ’ Step ’ The ‘The g The

study

Space group

Alloy

B2(AI,,,Ni,,CrZ)

in the present

in the alloys

= = = = = = = = = =

(A)

Reference

4.049 25.256, b = 7.582, c = 10.955, p = 128.68” 12.4, b = 34.6, c = 20 19.98, c = 24.67 9.123 7.811, LY= 109.13” 6.598, b = 7.352, c = 4.802

4.036, c = 4.897 2.8864

2.884 (a = 2.963 for 32% of Al)

belong

1171 [HI

1111 [18,191 1181 [201

WI P4

[231

to sets B and C

Phases

in the order

of appearance

Al,Ni?, AI,Ni. A/AI,Ni. BAI,Cr/AI,Ni“, ternary eutectic AINi, Al,Ni,, 4,’ AINi, cwCr. AI,Ni?‘, [AI,Cr, 4. A, A/AI,Ni, ternary eutectic nAl, ,Cr,. HAl,Cr/AI,Ni’, ternary eutectic p, nAl, ,Crz, BAI,Cr, BAI,Cr/AI,Ni’, ternary eutectic AI,Cr,, p, qAl,,Cr,, BAI,Cr, ternary eutectic by a significant DTA peak. higher than the maximal temperature

of the DTA furnace.

completely remelted. crystallization path could not be determined beyond this step because too little of the liquid phase carresponding to a solid-state transformation. amount of AI,Ni? phase is very small. amount of eutectic fIAl,Cr/AI,Ni is very small. eutectic constituent (SAI,Cr/AI,Ni (see Fig. 3(d)) does not display a classic eutectic structure.

the B alloys whose compositions are close to those of the Al-Ni binary system (Table 1). For instance, primary dendrites of Al,Ni, phase, identified in the quenched liquidus Bl alloy, appear with a white contrast in the SEM image (a), and primary AlNi dendrites, identified in the quenched liquidus B3 alloy, appear with a grey contrast in the SEM image (c). From a comparison of the SEM images (a) and (b) (Fig. 2) the A13Ni, dendrites appear more faceted and more acrcular in the quenched Bl sample than in the slowly solidified Bl sample. However, whatever the solidification rate for B3 alloy samples (Fig. 2(c) and (d)) the AlNi dendrites remain non-faceted. As observed in Fig. 2(a), the Al,Ni, dendrites are surrounded by Al,Ni in grey contrast. The layered configuration of the A1,Ni phase is typical of a compound produced by peritectic reaction. However, in the slowly solidified Bl sample (Fig. 2(b)), Al,Ni, is not only surrounded by A1,Ni but also by the ternary A phase in darker grey contrast, which exhibits some embedded Al,Ni precipitates. Such an Al,Ni,-h interface results probably from a solid state peritectic reaction between A and Al,Ni. The A1,Ni precipitates in the A phase can be interpreted as coming from

the DTA sample

remained.

eutectic solidification. The last stages of the crystallization path for the quenched liquids Bl sample cannot be easily interpreted because of the small proportion of liquid in the interdendritic groove at this stage of the crystallization sequence and the fineness of the microstructure. Meanwhile, from the phases identified in the slowly cooled Bl (Al,,Ni,,Cr,) sample, i.e. A, Al,Ni, BAl,Cr and Al (Fig. 2(b)), it can be assumed that the solidification of a hlAl,Ni eutectic was followed by the solidification of a binary eutectic BAl,CrI A1,Ni and then the ternary Al,NiIBAl,CrIAl eutectic. For the B3 alloy (A,,Ni,,Cr,,) solidified slowly at 5 “C min-’ in the DTA furnace, the solidification of an c&r solution, and subsequently AlNi was identified (Fig. 2(d)). A very small amount of Al,Ni, phase was also found in this sample. This aCr phase, appearing in clear grey contrast, contains a very fine structure of black contrast which is characteristic of a later solid phase precipitation (inset of Fig. 2(d)). It may be described as the solid state transformation &Zr-+ AlCr, + [Al,Cr, if we consider the existence of a smooth exothermic peak observed at T = 884 “C in DTA thermograms and the transformations proposed for the binary Al-Cr system [ll].

250

E. Rosril-Laclau

et al.

I Journal

of Alloys

Fig. 2. Microstructures observed by SEM in the quenched liquidus ((a) and (c)) and slowly solidified ((b) and (d)) Bl (Al,,Ni,,Cr,) and B3 (Al,,Ni,,Cr,,) alloys (t.e. means ternary eutectic and the insert of (d) shows that a solid-state transformation of the cvCr phase has occurred, probably cvCr 4 AICrz + [Al,Cri).

For C alloys whose compositions are close to those of the binary AI-Cr system, the primary phases were either directly related to binary Al-Cr compounds or

and

Compounds

233 (19%)

246-263

to different ternary intermetallic compounds. The latter phases were identified from TEM studies carried out conjointly with the SEM characterizations of microstructures (cf. Section 5). The Cl (Al,,Ni,Cr,) alloy solidifies primarily as a ternary compound called 4 (Fig. 3(a) and (b)). Its morphology depends on the solidification rate: in ascast and quenched liquidus samples the 4 phase forms long needles bordered by a thin layer of h phase (Fig. 3(a)), but in slowly solidified DTA samples the A phase bordering the 4 dendrites is coarser (Fig. 3 (b)). Such a change in morphology can be attributed to a solid state peritectic transformation 4 -+ A occurring during the slow solidification process. As observed in the SEM images (a) and (b) (Fig. 3), the solidification has ended with the formation of a hlAl,Ni eutectic followed by that of the ternary eutectic Al,Ni/OAl,Cr/ Al. Note that the structure of the eutectic constituents is coarse, as is commonly observed for eutectics involving intermetallic compounds. For samples of the C3 (Al,,Ni,Cr,,) alloy, the primary phase to solidify is another ternary compound called p with a stoichiometry very close to that of the binary A1,Cr compound [ll] but with a different crystalline structure (cf. Section 5). As in the preceding case, its morphology depends on the solidification rate: in quenched liquidus samples it forms faceted and acicular dendrites (Fig. 3(c)) while in slowly solidified samples it is observed surrounded by phases in concentric layers produced by peritectic transformations (Fig. 3(d)). These transformations have given rise to the formation of the nA1, ,Cr2 and 8A1,Cr phases and probably of a eutectic BAl,Cr/Al,Ni (which was not formed by classic cooperative growth of two phases), the solidification ending with the ternary eutectic Al,Ni/BAl,Cr/Al. We have found that the crystallization path of the C4 (Al,,.,Ni,,Cr,,,,) alloy depends on the solidification rate. When solidified slowly, the sequence yAl,Cr,, p, nAl,, Cr,, BAl,Cr and ternary eutectic was observed. A small proportion of A, resulting probably from a solid state reaction, was also observed in an irregular layer located between p and nA1, , Cr,. At high solidification rates, from as-cast or quenched liquidus samples, the sequence was yAl,Cr,, 4, A, BAl,Cr and ternary eutectic. Both results were considered in order to establish afterwards a first schema for the liquidus phase field boundaries, though at high solidification rates the results are more significant for the liquidus-solidus equilibria only. We note that from the SEM observations of Urtado [4] of an Al,,Ni,,)Cr, alloy, it appears that the beginning of the solidification sequence was in fact 4 followed by /\. The phase compositions of the different samples were analysed (Table 5). Because of the small solid

E. Resell-Lnclau

et al. I Journal of Alloys und Compounds

23.3 (1996) 246-26.7

251

Depending on the initial liquid alloy composition some phases exhibit a variation in composition which is actually related to their range of solid solubility. Thus the solid-solubility ranges of the 4 and AlNi phases appear to be quite large. Note that a large solid-solubility range for AlNi has also been observed in many other ternary systems. In the case of the A phase, the solid solubility range must be quite narrow since the composition of h phases resulting from the solidification of different alloys remain concentrated in a narrow range of composition around Al,,Cr ,?Ni,,,. Therefore, taking account of all these crystallization paths, determined from the solidification of different alloys and the measured liquidus temperatures, it is possible to propose a schema for the projection of liquidus phase fields limited to the subsection AlCrAlNi-Al (Fig. 4(a)). The compositions of the initial B and C alloys are given in this diagram. In another anisothermal projection (Fig. 4(b)), the measured compositions of the phases are reported in order to illustrate their solid-solubility ranges over the whole range of temperature. The solid lines in Fig. 4(a) indicate only approximately which liquidus phase field boundaries can be considered on the basis of the present experimental results. The direction of the slope of each monovariant line is indicated by an arrow. As a main feature this diagram exhibits a eutectic valley connecting a quasi-binary eutectic cuCr/ AlNi to the ternary eutectic Al,Ni/BAl,Cr/Al; this eutectic monovariant line is interrupted by several pseudo-peritectic transformations on either side.

5. Structure of Al-Ni-Cr

Fig. 3. MLxostructurcs observed ((a) and (c)) and slowly solidified (AI,,Ni,Cr,,,) alloys.

by SEM in the quenched ((b) and (d) Cl(A1,2NiCr,)

liquidus and C3

solubility of chromium in the Al, Al,Ni, Al,Ni, phases and of nvckel in the aCr, yA&Cr,, [AI,Cr, phases, the composition of such phases remains nearly identical, whatever the solidification rate.

phases

The structures of different phases, found previously using SEM in samples Cl (Al,,Ni,Cr,), C2 (Al,,Ni,Cr,), C3 and C4 (Al,,Ni,Cr,J (A177.3Ni4.3CrIX.4),were identified by electron diffraction. Because many different structures have been found to be related, we propose for the sake of clarity, to follow a description according to a summary of the results presented in Table 6. In the first column the names of several phases mentioned previously are given. In columns 2, 3 and 4 several structures are indicated for each of these intermetallic compounds (excepted for the 9 phase) because modulated structures, polytypes or superstructures have been identified. The samples in which these structures have been found are indicated in the fifth column. Finally, some comments on the main structural characteristics and crystallographic relations are given in the last column. An agreement of the TEM and SEM results has been deduced from several comparisons. Thus in

252 Table 5 Chemical

E. Resell-Laclau

composition

of phases

analysed

et al. I Journal of Alloys and Compounds

by EPMA

in different

alloy samples.

Sample

Phase Al Al,Ni AI,Ni, AlNi BAI,Cr vAl,,Cr>

4 (from[41) A

relation to this summary, a detailed several TEM results is reported. 5.1. Modulated

Al,(Ni,

description

233 (1996) 246-263

of

Cr) structure

For the eutectic structure within the Cl (Al,,Ni,Cr,) and C2 (Al,,Ni,Cr,) samples, we have verified that an orthorhombic structure with cell parameters a = 6.6 A, b = 7.35 w and c = 4.8 A and space group Pnma, corresponds to the orthorhombic A1,Ni compound [20]. This was observed for fine rod-like A1,Ni precipitates of about 2 Km diameter, within an Al matrix. However, for large round particles of 10 pm or more (see Fig. 3(b)), the A1,Ni structure appears to be modulated (Fig. 5(a) and (b)). This modulation gives rise to satellite reflections lying either along both equivalent [Oil]* and [Oil]* directions or along only one of them, the period of modulation being variable (between 1.44d,, , and 2.7d,, , ). In addition, reflections OkO, 001 and Okl, with, with respectively k, 1 and k + 1 odd, which are normally forbidden for the Pnma space group, are observed. As such reflections can be related to a variation in the 011 fringe contrast over a period of 2d,,, (Fig. 5(a)), it might correspond to a superstructure where the b and c parameters are twice those of Al,Ni. Whether these modulated structures and superstructures result from a substitution of Ni atoms by Cr atoms, it could be the effect of a difference between the atomic radius of the Ni and Cr atoms, their Goldschmidt radii corresponding respectively to 1.28 A and 1.25 A. In order to understand how the

Cr (at.%)

Ni (at.% )

Al (at.%)

0.27 0.3 0.4 8.2 12.9 16.2 16.8 20.1 28.3 41.2 60 18.7 15.7 17.9 17 18.4 18.7 16 12.4 12.8 11.7 12.5

0.18 24.9 39.8 36.9 1.7 3.1 2.8 3.2 0.25 1.3

99.55 74.8 59.8 54.9 85.4 80.7 80.4 16.8

7 A.

1.7 5.8 5.6 9.7 9.5 10.1 8 9.5 10.4 9.3 8.5

71.35 57.5 38 79.6 78.5 76.5 73.3 72.1 71.2 76 78.1 76.8 79 79

A1,Ni structure can be modulated, we have considered a structural representation previously proposed for a very similar structure (Fe,(Z) by Hyde et al. [23]. In that case, it can be pointed out that the Al,Ni structure is formed of a network of Al octahedra linked by their vertices and where each Ni atom occupies the central interstitial position between eight octahedra (Fig. 5(c)). These octahedra are neither regular nor parallel with each other. Thus, when substituting just a few Ni atoms by larger Cr atoms, it can be assumed that the distortion and tilted positions of Al octahedra are modified slightly such that the Al,Ni structure becomes modulated; a superstructure might exist for an ordering between Ni and Cr atoms. The composition of such an Al,(Ni, Cr) modulated phase might correspond to Al,,,,Ni,,,,Cr,,, (cf. Table 5).

5.2. 19structure and related polytypes

The structure of the ternary 0 phase is directly related to that existing in the binary Al-Cr system: BAl,Cr of structure type Al,,V, whose atomic structure is known [17]. From the analysis of its chemical composition (Al,,,,Ni,,,Cr,,,,), compared with that of 0A1,Cr (or Al,,Cr,), it appears that both Al and Cr atoms could be substituted by Ni atoms. The 8 structure has been identified in the C2 (Al,,Ni,Cr,), C3 (Al,,Ni,Cr,,) and C4 (Al,,., Ni4.3Cr,,,4) samples. However, in the C2 sample, 13was found to be strongly faulted by a (iii) twinning; a twinning sometimes

E. Resell-Laclau

et al. I Journal of Alloys und Compounds

2.73 (1996) 246-26.1

25.7

twin planes are (111) i.e. planes which are perpendicular to the diffracting vector g,,, . A polytype structure of 0 can be identified in the left part of this image since the twinning is repeated periodically every two (ii 1) planes. For a determination of the cell parameters of such a polytype structure we have considered a primitive triclinic cell of 8, which can be obtained with the following matrix of transformation:

Ni at%

b

Al Ni at% Fig. 4. Proposed Al-Ni-Cr phase diagram limited to the subsection AlCr-AINLAI: (a) projection of liquidus phase fields where the compositions of the B and C alloy sets are reported; (b) anisothermal projesztion of the measured composition of the phases giving a rough idel of the solid solubility of different phases.

repeated periodically so that it gives rise to the formation of polytype structures. Two cases of polytype structures have been identified (Fig. 6 (a) and (b)). In the first case (Fig. 6(a)), the 0 structure is observed along its [i 0 i] zone axis. The indexing of the diffraction pattern shows that two orientations of the monoclinic cell are related either by a mirror plane or by a twofold axis along the khh row or the direction perpendicular to this row. From the corresponding HREM image, it can then be deduced that both orientations result from a twinning where the

where both origins in 000 and 0 l/2 112 of the C2/m space group of the 0 phase have been considered. As the (i 11) planes of the monoclinic structures corresponds in that case to the (010) planes of the primitive cell, the polytype cell can be deduced simply from a stacking of four triclinic cells related by mirror operation every two triclinic cells through its (010) plane (see the drawing Fig. 6(a)). Thus the Bravais lattice we have found is monoclinic and its cell parameters are a = 11 A, b = 24.7 A, c = 13.2A, j3 = 113.85“. Taking into account that the centre of symmetry of the C2/m space group of 19 remains unchanged and that two mirror planes are generated by twinning, the space group of this polytype structure can be again C2/m. In the second case (Fig. 6(b)) the 19 structure is observed along its [112] zones axis where again a (iii) twinning can be deduced from the indexing of the diffraction pattern and analysis of the corresponding HREM image. In the right part of this image, it appears that the periodicity exhibited by a sequence of twin planes is characteristic of another polytype structure: both orientations of the primitive cell are repeated according to a sequence -l-2-1-2-. Then, as before, the Bravais lattice and cell parameters of this second polytype can be deduced directly from the stacking of the primitive triclinic cells (see the drawing Fig. 6(b)). It is triclinic with cell parameters equal to a = 11 A, b = 19.05 A, c = 13.2 A, ff= 79.2”, /? = 113.85”, y = 78.3”. Its space group should be Pl since the centre of symmetry of the C2/m space group of 0 also remains in this case. Note that two equivalent twinning modes (ill) and (111) exist in the 0A1,Cr structure. This has been observed by Zhang et al. [24] in melt spun Al-Cr alloys where 0 coexists with an icosahedral A1,Cr phase. Both these twinning modes may generate a pseudo-fivefold symmetry because the atomic structure of 13 is closely related to the symmetry of the icosahedral point group [ 111. 5.3. 7 structure The 7 structure observed in the C3 (Al,,Ni,Cr,,,) (or sample is identical with that of the nAl,,Cr,

E. Resell-Laclau

254

Table 6 Summary

of the results

obtained

from structure

Phase

Bravais lattice

Modulated Al,(Ni, Cr)

Orthorhombic

0 and polytypes

Monoclinic

et al.

I Journal

identification

Triclinic

Orthorhombic type

nA1,,Cr,

and Compounds

of different

Al-Ni-Cr

Cell parameters

Space

233 (1996)

246-263

phases

(A)

Samples

Comments

Cl” (‘2:’

Wavevectors of the modulation along the [01 I] and (01 I] directions (Fig. 5) Polytype structures resulting in a periodic ( 11i) twinning of H (Fig. 6) Without twinning in C.3 and C4 samples

group N = 6.6. 11= 7.4. (’= 4.8

C2lm

L, = 25.256, h = 7.582. c = 10.955. p = 128.68”

C2lm probably Pl probably

up = 0 = c = /3 = (I =

type Al,,V, Monoclinic

of Alloys

C1HCfil or

Il,h=24.7.c= 13.2, 113.XS” Il. h = 19.05. 13.2, (Y= 79.2” 1 I?&“, y = 78.3” 12.4, h = 34.6, c = 20

Crnc2,

CJand polytype”

Monoclinic

Orthorhombic A and superstructureh

Orthorhombic

a,,, =: 13.4, h,,, = 12.55,

c ]

P2, ltr1

(’I// = 12.5s. p,,, = 1(IO

CT: (‘4”

Pn2,n

a,, = 12.55. h,, = 12.55.

(‘2,:

Imnm

cl> = 26.4 u,, = 1255, I,,, = 12.55. (‘,, L 30.75

(‘2:‘ C4”

P2,

or

probably

p, rhombohedral

R3 or R.3

pz hexagonal

P6,

0, = 12.55, h, = 25.10. c, = 17.75. U L p = 110.7”. y = 90” i,,< = 2X.7. (Y= 36”. or in the hexagonal system ‘I, = 17.7. c, = 80.4 a> = 17.7. cI = 12.4

p3 hexagonal

P6,

a, = 30.7. c, = 12.4

Triclinic

P

:‘:

Structure related to that 01 the pAl,Mn phases. 0 = c-12. h = 12&, V’i. c = Uil ‘The orthorhombic structure results in a periodic (I(K)) twinning of J, (Fig. 7) (I,, = h,r,. h,, -- co,. c,, ~-:2~,,, sin p,,, Epitaxy hetween rb and A monoclinic (Fig. X) (I,, = h,),. h,, = UC/,. co 7rr,!, sin p,!,16 ‘1, = 2m,,. b, = h,,.

c, = (a;l,+ t,, 1 cf,)’2 0 = CY.5 -7u / L’:i E’pitaiy bctwecn p, and p? ,r3 is formed by a reaction between p, and /jl (Fig. 9) Structures related to that of the pAl,Mn phase, ‘12 = a,< ~512, c, 2 c,, 12 (‘3””

‘I Column 1, the orthorhombic polytypc coexists with a minor proportion of hexagonal phase which has not been completely identified: both is (I,,~<= h,?,,,,,, th;se phases exhibit an orientational relationship such that a,,,, ]]c,,,~,~~~ and [I lo] ]]b,,,,,,,, ; the relation between the cell parameters \3. ’ Column 1, in the C2(Al,2Ni,Cr,)* sample, the A, sublattice was sometimes observed to be slightly distorted into a monoclinic or triclinic system. Column 5, symbols * and ** refer to SqucncllcdIlqoldu~and S,,,, respectively.

vAl,Cr) phase that we have identified in the binary system [ll]. From a comparison of the stoichiometries between the binary and ternary phases, it seems that Al atoms rather than Cr atoms are substituted by Ni atoms. From some new results obtained on the binary Al-Cr system [25], it appears that the orthorhombic structure of the 7 phase is in fact closely related to the hexagonal structure of pAl,Cr, isomorphous to pAl,Mn [12,13]. The atomic structure of ,uAl,Mn, which also has pseudo-icosahedral symmetry, exhibits icosahedral atomic clusters which are identical with those found in BAl,Cr [11,24]. From the similar diffraction characteristics observed between n and ,q it was possible to deduce that their cell arameters are related as follows: a,, = ~~12, b,, -L a, J 3. co = a@, with

a,, = 12.4 A, b, = 34.6 A, c, = 20 A and a@ = 19.98 A,

cP = 24.67 A (the accuracy of the cell parameters of 7, determined by electron diffraction is kO.l A while the accuracy of the cell parameters of p, determined by X-ray diffraction, is about 20.01 A).

5.4. + structure and its polytype A monoclinic structure attributed to the 4 phase has been identified in the Cl (Al,,Ni,Cr,), C2 (Al,,Ni,Cr,) and C4 (Al,,,,Ni,,,Cr,,,) samples. The electron diffraction patterns related to the main zone axes of this structure (i.e. [lOOI, [OlO] and [OOl] are shown in Fig. 7(a); from the symmetries observed between the intensities of the reflections, it can be

E. Resell-Laclau

et al. I Journal of Alloys and Conlpourd~

33

(1996) .?46&26.1

25.5

Fig. 5. AI,(Ni. Cr) modulated structure: (a) and (b) high resolution images where the modulation of AI,Ni gives rise to a variation in the axis. satellite reflections are contrast parallel to the (0111 and (01 I) fringes. In the corresponding electron diffraction patterns of the [ IOO] lone ._ observed lying along the Ohh and Olzh rows: (c) representation of the Al,Ni structure viewed along its main zone axes (OIO]. [OOI] and [loo] (see text).

deduced that they are characteristic of the Laue class 12/m 1 and a symmetry point group 2 or 2/m. In addition, since only one condition of reflections for Ok0 with k = 2n was found, there are two solutions for the space group of the 4 structure, either P2, (noncentrosymmetric) or P2, lm (centrosymmetric). Note in that case that the weak Ok0 reflections with k odd, which are observed on the patterns [OOl] and [loo],

come from double diffraction effects. Such effects were identified by carrying out small rotations about the hO0 and 001 rows; on rotation the double diffraction reflections disappeared. An orthorhombic polytype of 4 forms because of a periodic multiple (100) twinning of 4 (Fig. 7(a) and (b)). The diffraction patterns of this polytype exhibit a symmetry belonging to the Laue class mmm (Fig.

256

E. Resell-Laclau

et al. I Journal of Alloys and Compounds

233 (1996) 246-263

Fig. 6. 0 structure and related polytypes resulting in different periodic sequences of the (Til) twinning: (a) the twinned structure is observed along the [loll zone axis of 0, in the left part of the HREM image, the twinning, repeated in a periodic sequence every two (iii) planes which can be symbolized as -2-2-2-Z-), corresponds to the formation of a monoclinic polytype of which the cell parameters have been determined by considering a primitive triclinic cell of 0 (see drawing and text); (b) the twinned structure is observed along the [ii21 zone axis of 0, the periodic sequence of the (iii) twinning, observed in the HREM image, is -l-2-1-2and corresponds to a polytype of triclinic structure.

7(b)). As the reflection conditions are such that h = 2n for hO0 and k01, k = 2n for OkO, 1= 2n for 001 and h + k = 2n for hk0, it results that the space group of this polytype is either Pnma (centrosymmetric) or Pn2, a (non-centrosymmetric and corresponding to another abc setting of the standard space group Pna2, [26]). From the HREM image of this polytype (Fig.

7(c)), it appears that there is no mirror plane perpendicular to the b axis but a twofold axis going through the centre of a monoclinic unit. The symmetry point group is then mm2 and such a structure is therefore non-centrosymmetric with space group Pn2,a. The HREM image in Fig. 7(c) shows an interface between the monoclinic structure of 6 (in the bottom part of

E. Resell-Laclau

et al. I Journal of Alloys and Compounds

..*r.*rrr.-

233 (1996) 246-263

2.57

i

.*-a,*.* ..‘.....“’ .

..**m. *..

‘

---,1_;_y

^.^,

--.-_,

.._--___I___

. *

.

.

*

..,.“..

f-T

“-----I

;

;...****.‘

.

.I

.

.

.

.

.

.

j.::;::;::

* . . .

. I

ic

. .

.

.

..

Lt.*.

-.

.;

^ .

i:

:i

.

..‘ . . .. .

.

f”“;:”

2; ; : ; :

.-. i ; f “p ; . *&-J:

Fig. 7. 4 structure and related polytype: (a) main electron diffraction patterns of the 4 monoclinic structure, the [ IOO]and [OOI] patterns arc related to the [OIO] pattern by 90°C rotations about the 001 and hO0 rows respectively: (b) main electron diffraction patterns of the orthorhombic polytype, as before the patterns [IOO], [llO] and [OOl] are related to the [OOl] pattern by 90°C rotations about the OkO, hh0 and hO0 rows respectively: (c) HREM image showing an interface between the 4 monoclinic structure (bottom part of the image) and its polytype (upper part of the image). The corresponding pattern is composed of the superposition of the [OlO] monoclinic pattern and the [OOl] polytype pattern. Ir between these structures, the monoclinic structure exhibits (100) twin planes spaced irregularly. These twin planes become arranged in a period manner for the polytype (the b and c edge lengths of the projected orthorhombic cell correspond rcspectivcly to 12.55 and 26.3 A). (d) Both structures exhibit a similar projection in HREM along the [OOl] zone axis of the 4 monoclinic structure and the [OIO] zone axis of the polytype.

the image) and its polytype (in the upper part of the image). In between the two structures, the 4 phase exhibits several (100) twin planes which are not in a

periodic sequence. The electron diffraction pattern corresponding to a selected area of this interface is constituted of the superimposition of the [OlO] pattern

E. Resell-Laclau

258

et al. I Journal of Alloys and Compounds

of 4 and [OOl] pattern of the polytype. The orientational relationships between both structures can then be deduced as follows:

t0101,IIWlpoly WI, IIPqo,y ww, IlwJ>,o,, Moreover,

the relation between the cell parameters

is

a+ = (c PO,Ysin-‘/?)/2, b, = apOlyrc+ = bpOly with a+ = 13.4 A, b, =, 12.55 A, c+ = !2.55 A, p = JOO’= and a pO,Y= 12.55 A, bpoly = 12.55 A, cPolY= 26.4 A.

Note that both structures, as well as the (100) twinning of 4, cannot be distinguished when viewed along the [OOl] zone axis of 4, parallel to the [OlO] zone axis of the polytype, because their diffraction patterns are very similar as well as their corresponding HREM images (Fig. 7(d)). 5.5. A structure, superstructures

and polytype

An orthorhombic structure, attributed to the A phase has been identified in C2 (Al,,Ni,Cr,) and C, samples. The diffraction patterns (Al,,.,Ni,.,Cr,,.,) related to the main zone axes [lOO], [OlO], [OOl] and [llO] of this phase are shown in Fig. 8(a). Their symmetry is characteristic of the Laue class mmm and from the reflections we have found (h = 2n for hO0, k = 2n for OkO, I= 2n for 001, k + 1= 2n for Okl, h + 1 = 2n for h01, h + k = 2n for hk0 and h + k + 1 = 2n for hkl), three different point groups of symmetry can be considered, 222, mm2 and mmm, which give rise to the following possibilities for the space group: 1222, Zmm2 and Zmmm. However, as no evidence for a twofold axis has been observed by high resolution imaging along the three basis vectors of A, the solution of the space group Zmmm seems more probable. The cell parameters of A are a, = 12.55 A, b, = 12.55 A, c, = 30.75 A, where it appears that the a and b parameters correspond to those of the orthorhombic polytype of 4. Such a relation can be confirmed from an epitaxy relation existing between A and 4 (Fig. 8 (b)); in this figure, the two diffraction patterns corresponding to a selected area of the A-4 interface are related by a rotation of 45” about the [OOl] axis of A parallel to the hO0 row of 4. Thus the orientational relationships between both these phases are

II(001)* WI, IIWOIA Pw dIIWI A

(lOO),

The HREM image corresponding to the zone axes [OOl], I([loo], shows that between A and 4 there is a coherent interface plane parallel to (lOO), and (OOl),,

233 (1996) 246-263

involving therefore an epitaxy relation between A and 4. In addition, as in the pattern of zone axis [OOl], )I[loo] A, it appears that the vectors g700d, and g,,,, are equal, so the cell parameters of A can be deduced from those of 4: cu, = b,, b, = a+, cA = (7~2, sin &)/6. Note that before the present TEM characterization, an epitaxy relation between A and 4 could be expected from the faceted interface observed between these phases in the SEM image shown in Fig. 3(a). A triclinic superstructure of A with probably a related polytype has been identified in Cl (Al,,Ni,Cr,), C2 (Al,,Ni,Cr,) and C3 (Al,,Ni,Cr ,0) samples. The triclinic superstructure is such that supplementary reflections are observed on the [OlO] pattern of the A structure (see the diffraction pattern of Fig. 8(c). These reflections form rows identical and parallel with those of A along its [OOl] direction but they are shifted by a vector g,,, /2 with respect to those of A. From a 90” rotation about the reflection row parallel to the [OOl] axis of A, we have observed that the diffraction pattern remains identical with the [loo] pattern of A. Thus it results that a triclinic cell has to be considered for this superstructure. The cell parameters can be deduced from those of A: in a plane parallel to the (001) plane of A, one has a, = 2a, b, = b,

the vector c, being parallel to a vector of coordinates (l/2,1/2,1/2), its modulus is c, = (a; + b2, + c2,)“*/2 and the angles of the cell are Ly= p = 110.7” y = 90 This triclinic superstructure is seen oriented along its b axis on the top and bottom parts of the HREM image of Fig. 8(c). In between, there is apparently a (100) twinning of the triclinic structure which if it were repeated periodically every (001) plane could be at the origin of the formation of an orthorhombic polytype. Its cell parameters deduced from those of the A structure should be (yPolY = 2a,, bpoly = b,, cPolY= 2c,. However, owing to an intimate intergrowth of several superstructures and/or polytypes where the indexing of corresponding diffraction patterns becomes confused, we cannot conclude from our present results whether such a polytype exists or not. Though further investigations are required, we can only point out that at least two other structures are directly related to this structure of the A phase: the first exhibits a characteristic diffuse scattering in parallel planes perpendicular to the [OlO], axis which are situated at ng,,,/2 (with n odd integer) (Fig. 8(d)). The other comes from the fact that the A structure can exhibit superlattice reflections which, a priori, can be interpreted as characteristic of

d

Fig. 8. A structure and related superstructure: (a) man electron diffraction patterns of the A orthorhombic structure. the [lOOi, [I 101 and (OlO] patterns are related to the (WI 1pattern hy 90” rotations about the OkO.hh0 and hO0 mws respecttvely: (b) upltaxy relation between 4 and A, diffraction pattern in the left left part of the figure corresponds to an Interface between 4 oriented along its [OIlI zone axis and A oriented along its [IIO] zone axts. The next pattern was obtained from a 45’ rotation about the vertical 001 row of A parallel to the hOO row of 4. It is composed of the superposition of the [Ool] pattern of 4 and the ~IOO]pattern of A. The corresponding HREM magr shows that between both structures there is a coherent mterface parallel to (IGO)+ and (001 )A (the h and c edge lengths of the propxted orthorhombic cell correspond respectively to 12.55 and 30.75 A). (c) The A structure exhibits a triclinic superstructure. the dlffraction pattern corresponds to the [OIO] pattern of A but with superlattice reflections which form an obhque two-dimensional lattice (there is also diffuse scattering in planes perpendicular to [lOO\,). The triclinic structure was deduced from systematic rotations, in particular the 90°C rotation ahour the [Ool], direction which yields a pattern identical with the [IOOJ, pattern. In the corresponding HREM image, the [OlO) projection of the triclinic structure. seen in the top and bottom parts of the image, is character&c of a slight variatwn m the fringe contrast. In between, there are apparently a few (100) twn planes III the triclinic structure which might he considered as giving rise to the formation of a thin sandwich of polytype structure (we text). (d) another phase related to A corresponds to a structure cxhibitmg a diffux scattering situated I” parallel planes perpendicular II) the [OIO], axis and ngo,,,i2 (with n an odd rntegcr).

-

260

E. Resell-Laclau

et al. I Journal of Alloys and Compounds

an orthorhombic structure of Fddd space group with cell parameters two times larger than those of A. 5.6. p phase related structures Two related structures, rhombohedral and hexagonal, have been identified in C3 (Al,,Ni,Cr,,) and C4 (Al,,,,Ni,,Cr,,,,) samples. They are attributed to the p phase. Their compositions can be assumed to be identical or nearly identical since they were not distinguished by SEM. The main diffraction patterns, [ill], [112], [ll_O] for the rhombohedral structure (p,) and [OOOl], [lOlO], [?iiO] for the hexagonal structure (pz), are shown in Fig. 9(a) and (b). For the rhombohedral phase, the symmetry between intense reflections observed on the [ill] pattern is characteristic of the Laue class 3 (point group 3 or 3). As we have found that all reflections are permitted, the solutions for the space group are either R3 (centrosymmetric) or R3 (non-centrosymmetric). For the hexagonal phase, the symmetry observed between intense reflections on the [OOOl] pattern is characteristic of the Laue class 6/m (point group 6 or 6 or 6/m). In that case, the reflection conditions are such that 1 = 2n for 0001. Thus there are two solutions for the space group, either P6,lm (centrosymmetric) or P63 (non-centrosymmetric). Both these phases exhibit an epitaxy relation, identified by HREM (Fig. 9(c)); on this image, the rhombohedral structure, viewed along its [liO] zone axis presents an interface plane (111) parallel to the (0001) plane of the hexagonal structure which is viewed along its [1210] zone axis. In addition, the rhombohedral phase presents a (111) twinning where the twin plane appears to have a thickness corresponding to one hexagonal cell. The orientational relationships between the two structures are (lll)ll(OOOl) -[110] 1)[1210] [112] 11 [lOiO] The habit plane between both structures being coherent, there is also a simple relation between their cell parameters a

233 (1996) 246-263

two preceding phases were not observed in this sample, it can be concluded that the hexagonal phase pi has been formed by a reaction between the two preceding phases. Therefore, p, and p2 are metastable at room temperature. The [OOOl]pattern of this third phase appears to be directly related to the pattern of the previous hexagonal structure; it exhibits additional weak superlattice reflections so that its a parameter becomes & times larger than the a parameter of the previous hexagonal phase; its c parameter remains the same. As before, we have determined that its space group can be either P6,lm or P&.-The peak superlattice reflections observed in the [1210] pattern give rise in HREM to a weak variation in the fringe contrast (Fig. 9(e), see arrows) which seems to be quite typical of chemical ordering. This weak variation in the fringe contrast was not observed for the previous hexagonal phase. In addition, from the same HREM image (Fig. 9(e)), it can be noticed that the symmetry of the projected structure is m (i.e. a mirror plane perpendicular to the c a@). It is therefore characteristic of a point group 6 or 6 among the three previous possibilities deduced from analysis of the diffraction patterns (i.e. 6 or 6 or 6/m). For the point group 6/m, the symmetry of the projected structure should be 2mm. Thus, on account of the reflection conditions (E= 2n for OOOf)only one solution remains for the space group: P6, for both hexagonal structures (pz and p3), Finally, the range of chemical composition for the p phases is very close to that of Al,Cr. However, none of the structures presently identified correspond to the pAl,Cr phase [ll]. This means that a substitution of Al or Cr atoms by Ni atoms in the p structure does not occur or is at least very small. As such a small proportion of Ni cannot give rise, a priori, to an important modification of the local atomic arrangement of the k phase, we have examined what could be the relations between the cell parameters of the present phases with those of the p phase. Thus, it appears that a P2 =awJ3/2ci,z=cP12

with aP = 19.98 A, cK = 24.67 A and ap2 = 17.7 A, cp2 = 12.4 A (see Table 6).

ahrx = 2a, sin(cy/2)

However, there is no particular relation between the parameter chex and the length of the threefold diagonal of the rhombohedral structure. The values of the ccl! parameters are aR = 28.7 A, cy = 36”, and ahex = 17.7 A, chex = 12.4 A. Another hexagonal structure (p3) has been identified in a C3 (Al,,Ni,Cr,,) sample, but quenched just after that the second peak was reached, by DTA (Fig. 9 (d)). As the

6. Discussion Looking for similarities with other intermetallic compounds identified in other Al-based systems, we have not found any structure type related to the ternary structures 4, A and p presently identified. Meanwhile, many Al-based systems which may be thought of as somewhat similar to the Al-Ni-Cr

E. Resell-Laclau

et al. I Journal of Alloys and Compounds

233 (1996) 246-263

. . ,..*.

-. .I.‘.

. .. . .. .

;.

:

.

*

***.

??

.

*

261

*

‘.:.‘~:

.

.

. I

.

.

/

.**.*.*.

‘..

.*

.

‘....... .I’...

.

s..

+-I .

.

.

1 . L

.*,...

1

..,

.

.

..,...r... ,...,,.*.. -.,..****.r .<.....**.. ~*.,-*

*.*.4

..r..,.**..~ .

.

...‘.

*

*

.

1

I

.

1

f

*

.

*

I

, * . . . . . * . . *

.

?? .

s

. 1

*

.

??

? ? ? ?

.

4,

*..

.

??

??

.*_. .

.

.

t

.

Fig. 9. p phase (related structures): (a) main electroi diffraction patterns of the rhombohedral structure, the [ll?] and [liO] patterns are related to the [ill] pattern by 90TC_rotations about the hh0 and hh2h rows respectively; (b) main electron diffraction patterns of the hexagonal structure, tile [lOlO] and [1210] patterns are related to the [OOOl] pattern by 90” rotations about the h2hhO and hOh0 rows respectively; (c) epitaxy relation between both these structures identified by HREM, where it appears that a coherent interface plane is parallel to the (111) plane of the rhombohedral structure and parallel to the (0001) plane of the hexagonal structure. In the left part of the image, the rhombohedral phase exhib’its a (111) twinning through a layer of hexagonal structure. (d) Diffraction patterns of the third hexagonal structure resulting from a reaction between both previous phases. With respect to the above diffraction patterns of the hexagonal structure, the present patterns exhibit additional superlattice reflections in planes parallel to (0001) and situated at ng,,,,,, (n an integer). (e) HREM image corresponding to the [1210] pattern, where it appears that only a slight variation in the fringe contrast results in the superlattice reflections (see text).

262

E. Resell-Laclau

et al. I Journal of Alloys and Compounds

systems have not been extensively studied. This is the case of Al-Mn-Ni [27] where the binary boundary Al-Mn, constituted of a succession of peritectic transformations [28], resembles that of Al-Cr. Moreover, the same hexagonal structure p is formed in both these binary systems. However, as far as determination of the Al-Mn-Ni system has been completed, the only ternary Al,,Mn,,Ni, phase [29], mentioned in the Al-rich part of this system, does not have an equivalent in the Al-Ni-Cr system. In the present study, we were also interested to know whether the Al,Cr quasicrystalline phase [30] obtained by rapid quenching (i.e. melt spinning) could be stabilized by an addition of Ni. In effect, from the literature on this subject an eventual formation of a stable quasicrystalline phase could be expected since many icosahedral and decagonal phases have already been obtained as ternary Al-base alloys, including the transition metals Cr, Mn, Fe, Co, Ni and Cu which follow in the periodic table of the elements, for instance the icosahedral phases Al,,Cu,,Cr,, [31], WI, W&W% [331, Al,,Mn,,Fe,, Al&uZOFe,, [34], Al,,Mn,,Ni, [34] and decagonal phases Al,sFe,5Ni10 1351, Al65Cu2,,Co,5 [361, Al,,Ni,,Co,, [37]. From the present results, it does not seem that a quasicrystalline phase could be obtained by normal casting of an Al-Ni-Cr alloy.

7. Conclusion A first schema of the liquid-solid equilibrium occurring in the aluminium-rich corner of the Al-Cr-Ni system has been established. We have shown firstly that a ternary eutectic transformation liquid -+ Al,NiIBAl,Cr/Al takes place at 634°C as could be assumed from the limits of the primary aluminium phase field, i.e. the Al/Al,Ni eutectic and the AlIBAl,Cr peritectic. This transformation is similar to those observed in the Al-NiFe and Al-Ni-Ti systems. The ternary phase fields in the AlCr-AlNi-Al subsection of the phase diagram are not only related to the different binary Al-Ni and Al-Cr liquidlsolid equilibria but also to the formation of three ternary compounds, namely 4, h and p, whose structures are different from those of the known Al-Ni and Al-Cr compounds. The monoclinic 4 and orthorhombic A structures are quite complicated since each of them exhibits very large cell parameters and different polytypes or superordering. The p phase, consisting of an intimate intergrowth of rhombohedral and hexagonal structures at high temperature, transforms into a unique hexagonal structure at lower temperature. Probably as an effect of nickel substitution, monoclinic and triclinic polytypes of the BAl,Cr compound are

233 (1996) 246-263

formed. A chromium-to-nickel substitution within the Al,Ni compound seems to be the origin of the formation of a modulated structure. The proposed phase diagram shows that two collections of liquidus phase fields can be distinguished on either side of a eutectic valley connecting a quasibinary eutectic aCr/AlNi to the ternary eutectic Al,NiIBAl,CrIAl. On the Al-Ni side, different ternary phase fields AlNi, Al,Ni, and A1,Ni related by pseudo-peritectic transformations, are simply prolongated from the binary Al-Ni liquid-solid equilibria. On the Al-Cr side, the binary liquid-solid equilibria cuCr, [Al,Cr,, yAl,Cr,, pAl,Cr, qAl,,Cr? and BAl,Cr, related by peritectic transformations, extend also in the ternary system; the formation of the ternary compounds p, 4 and h is related by pseudo-peritectic reactions of liquid phases with the compounds yAl,Cr,, pAl,Cr and vAl,,Cr,.

References [l]

S.M. Merchant 47-60.

and

M.R.

Notis,

Mater. Sci. Eng.,

[2] P. Rogl, in G. Petzow and G. Effenberg (eds.), Vol. 4, VCH, Germany, 1991, pp. 400-416. [3] A. Taylor 464.

and R.W. Floyd,

[4] P. Colin-Urtado, Grenoble, 1990.

66 (1984)

Ternary Alloys,

J. Inst. Met., 81 (1952-1953)

Thesis, Institut

National

Polytechnique

451de

[5] M.F. Singleton, J.L. Murray and P. Nash, in T.B. Massalsky (ed.), Binary Alloy Phase Diagrams, ASM, Metals Park, OH. 1990, pp. 181-184, 2nd edn. [6] J.L. Murray, in T.B. Massalsky (ed.), Binary Alloy Phase Diagrams ASM, Metals Park, OH, 1990, pp. 138-140,2nd edn. [7] H.W. Phillips, /. Inst. Met., LXVIIZ (1942) 27-46. [8] G.V. Raynor and P.C.L. Pfeil, /. Inst. Met., 73 (1947) 397-404. [9] P. Buldberg and A. Prince, in G. Petzow and G. Effenberg (eds.), Ternary alloys, Vol. 5, VCH, Germany, 1992, pp. 309323. [lo] K.J. Lee and P. Nash, J. Phase Equilibria, 12 (1991) 551-562. [ll] M. Audier, M. Durand-Charre, E. Rosell-Laclau and H. Klein, J. Alloys Comp., 220 (1995) 225-230. J. Microsc., 146 (1987) 303-312. [12] L. Bendersky, [13] C.B. Shoemaker, D.A. Kesler and D.P. Shoemaker, Acta Crystallogr. B, 45 (1989) 13-20. [14] T. Ohnishi, Y. Nakatani and K. Okabayashi, Bull. Univ. Osaka Pref, 24 (1975) 183-191. [15] E. Laclau, M. Durand-Charre and J.C. Joud, Proc. ECASIA, 3-8 October 1993, Catane, J. Mater. Res., IO (1995) 1154-1158. [16] J.L. Pouchou and F. Pichoir, Microbeam Anal., (1990) 120-126. [17] M. Cooper, Acta Crystallogr., 13 (1960) 257-263. [18] M. Ellner, J. Braun and B. Predel, Z. Metallkd., 80 (1989) 374-383. [19] P. Villars and B. Calvert, (eds.), Pearson’s Handbook for Intermetallic Phases, Vol. 2, 2nd edn., 1991, p. 936; Knapp, Z. Metallkd., 33 (1941) 153. [20] A.J. Bradley et al., Philos. Mag., 23 (1937) 1049-1067. [21] P. Villars and B. Calvert, (eds.), Pearson’s Handbook for Intermetallic Phases, Vol. 2, 2nd edn., 1991, p. 1038; Hsueh, Chinese J. Phys. Engl. Trans. 22 (1966) 479. [22] M. Cooper, Philos. Mag., 8 (1963) 811-821.

E. Resell-Laclau

et al. I Journal of Alloys and Compounds

S. Anderson, M. Bakker, C.M. Plug and M. [23] B.G. Hyde. O’Keeffe, frog. Solid State Chem., 22 (1979) 273-312. [24] H. Zhang, D.H. Wang and K.H. Kuo, Phys. Rev. B, 37 (1988) 6220-6225. [25] M. Audier et al., in preparation. [26] T. Hahn (ed.), International Tables for Crystallography, Vol. A, Reidel. Dordrecht, 1983. [27] T. Suzuki, in G. Petzow and G. Effenberg (eds.). Ternary Allo:n, Vol. 7. 1993, pp. 84-99. [ZS] M. IHarmelin, S. MaPmar, S.G. Fries and H.L. Lukas, Z. Metallkd.. 8.~ (1994) 814-818. [29] K. Robinson, Acta Crystallogr., 44 (1988) 707-710. [30] D. Srechtman, I. Blech, D. Gratias and J.W. Cahn, Whys. Rev. Lett.. 53 (1984) 1951-1953. [31] F. Spaepen, L.C. Chen, S. Ebalard and W. Ohashi, in M.V Jaric

[32] [33] [34]

[35] [36] [37]

233 (1996) 246-263

263

and S. Lundqvist (eds.), Quasicrystals (Proc. Anniversary Adriatic0 Research Conf.), World Scientific, Singapore, 1989, pp. I-18. A.P. Tsai, A. Inoue and T. Masumoto, Jpn. J. Appl. Phys., 26 (1987) L1505-1507; 27 (1988) 1587-1590. X.2. Li and K.H. Kuo, Philos. Mag. B. 66 (1982) 117-124. A. Singh, PhD Thesis, Bangalore, 1994. A. Singh and S. Ranganathan, J. Non-Cryst. Solids, 153-1.~4 (1993) 86-91. Y. Zhuang, Z. Zhang and D.B. Williams, 1. Non-Cryst. Solids, 253-154 (1993) 119-122. L.X. He, Y.K. Wu, X.M. Meng and K.H. Kuo, Philos. Mag. Lett., 61 (1990) 15-19. Mater. Trans. JIM, 30 A.P. Tsai, A. Inoue and T. Masumoto. (1987) 463-467.