Experimental analysis of the Ni-Al phase diagram

Journal of Crystal Growth 87 (1988) 185—192 North-Holland, Amsterdam

185

EXPERIMENTAL ANALYSIS OF THE Ni-Al PHASE DIAGRAM F.J. BREMER, M. BEYSS, E. KARTHAUS, A. HELLWIG and H. WENZL

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T. SCHOBER, J.-M. WELTER

**

Institut fdr Festkorperforschung der Kernforschungsanlage JWich, Postfach 1913, D-5170 JWich, Fed. Rep. of Germany

Received 12 July 1987~manuscript received in final form 11 September 1987

The phase diagram of the binary system Ni—Al has been determined near the Ni~Alcomposition as the basis for the growth of Ni

3Al crystals by the Czochralski and Bridgman methods. The phase morphology of crystals themselves also provides important information about the shape of the phase diagram. The temperature dependence of the long range order has been measured at a Ni1A1 single crystal by means of y-ray diffraction.

1. Introduction A series of fcc Ll2-type ordered binary alloys show brittleness at room temperature combined with increasing yield stress with rising temperature, e.g. Ni3A1, Ni3Ga and Ni3Si. In the case of Ni3A1 ductility and strength are increased by adding small quantities of boron and hafnium. The experimental and theoretical knowledge is summarized in ref. [1]. An experimental analysis of the elementary processes of doping and deformation requires single crystal samples with well defined stoichiometry and defect structure. In planning crystal growth experiments, we have been faced with two conflicting phase diagrams based on publications by Alexander and Vaughan [2] (abbreviated by AV) and Schramm [3] (S); they show astonishingly different behaviour around the relevant Ni3AI concentration range as can be seen from fig. 2b. Also, other investigations [4—8]could not clarify the reaction mechanism by which Ni3 Al is formed, Therefore, we have analyzed again liquidus and solidus lines and solid—solid coexistence lines by *

* *

Now at Max-Planck-Institut. D-4000 DUsseldorf. Fed. Rep. of Germany Now at Tréfimetaux, Centre de Recherche, F-60590 Sérifontaine. France.

rapid solidification of droplets, annealing of samples with well defined composition at various ternperatures and by crystal growth experiments, using transmission electron microscopy (TEM), optical microscopy. differential thermo-analysis (DTA), X-ray and y-ray diffraction and precision density measurements as methods for chemical, structural and morphological analysis.

2. Sample preparation High-purity ntckel (Société Métallurgique Le Nickel, cathodes, vacuum cast for degassing, carbon content 10 at.ppm, purity 5N) and high-purity aluminium (Vereintgte Aluminiumwerke: Kryal 0Z1, refined by the producer, purity 5N) has been used for sample preparation by rf melting in cold crucibles, for splat-cooling of liquid droplets and for crystal growth in Al 203 crucibles by the Czochralski and vertical Bridgman methods [9] in high-purity inert or reducing atmospheres at low pressures. The samples used by AV and S have not been assessed for purity. Nevertheless the quality of .

..

nickel and aluminium available in those days suggests that their samples contained at least 1000 at.pprn carbon as well as other impurities at a

0022-0248/88/$03.50 © Elsevier Science Publishers B.V. (North-Holland Physics Publishing Division)

Fl BrL’mer i’t a!., L’xjerinuoital ana/ysa of X~—.1/ pha.’e diagram

I 56

level which is orders of magnitude larger than in

our samples.

3. Phase identification methods In the examined Ni—Al concentration range. three different phases occur: The /3-phase. a NiAl solid solution, has a cubic body centred structure of the CsCl-tvpe with a lattice constant of ~j/~ 2.881 A. The well known y’-phase (Ni Al) has a cubic face centred structure of the Cu 5Au-tvpe =

with a~ = 3.565 A. The Ni-rich y-phase shows a face centred cubic lattice, the spacing of which increases with the increase in the Al content and is similar to that of Ni

1A1 [161. From all annealed samples. indicated in fig. I. powder X—ray diffraction patterns were taken on a Huber diffractometer unit equipped with a copper source and a monochromator set to diffract only copper Kn radiation. Splat-cooled samples were examined with transmission electron microscopy (Philips EM 400). The splat foils (thickness 10 jam) were jet electropolished in 95~ ethanol and 5~perchioric acid at 15~C. In addition to the phase identi fi— —

cation h diffraction patterns, an X-ray energy dispersive spectrometer (XEDS) was used for element analysis. Single crystals were cut by spark erosion. polished and etched in 30°fnitric acid. I 0~~f sulfuric acid. 1 0~ phosphoric acid and 50~cglacial acetic

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acid at 900 C to detect thc different phases. After etching both /3- and ‘y-phasc appeared dark in a bright y’-rnatrix. If the phases could not he identilied free from doubt by their morphological appearance. the samples were analysed after slightly etching using an energy dispersive X—ray spectrorneter (EDAX) on a scanning electron micro-

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scope (Philips SEM 525M).

4. Results Fig. 1 presents the results of the phase analysis

of splat—cooled samples by transmission electron microscopy [10]. The cooling rate of about 10 K/s was low enough to allow establishment of liquid—solid equilibrium. hut high enough to pre6

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Phase analysis by TEM after splat-cooling (01) and after -

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annealino of rapidly cast samples at 1050 C (120 h) and 7000 C (14 days) (0) by optical microscopy and transmission eleciron microscopy,

vent equilibration by diffusion during the cooling process after crystallization. Because TEM anal’ssis does not provide exact phase fractions of the samples. the symbols in fig. I are only ot qualita— tive importance. The phase lines themselves in fig.

analysis of a series of Ni—-Al samples after anneal. mo at 1050 ( by metallooraphv [II]. .‘ Fig. 2a presents results of DTA runs. The DTA samples were pre—alloved in a cold levitation cruci—

F.J. Bremer et al.

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Fig. 2. (a) Solidus and liquids coexistence lines in Ni—Al as measured by differential thermal analysis. The points indicate temperatures of transition during heating with 2 K/mm for samples with defined composition. (b), (c) The conflicting phase diagrams: (b) Alexander and Vaughan 12]; (c) Schramm [3].

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ble under a reduced argon pressure (Ar 5N7. 100 nihar). The measurements were done in a resistance heated furnace under the same conditions as mentioned above. The samples with a typical i’nass of 0.7 g were placed in specially designed Al ~0’i crucibles which themselves were located on W/ Re( 3 ~ )— W/Re(25 ~f) thermocouples. These were calibrated to the melting point of gold and

nickel. Several heating and cooling runs were taken, so that the final result of each sample is an average of 3 to 5 measurements. A comparison with AV and S is shown in figs. 2h and 2c. There is a remarkable agreement between S and our

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chemical composition and phase morphology show macroscopic variations which are characteristic of the shape of the phase diagram [12]. An analysis of the crystals after growth allows the determination of, e.g., the sign of the slope of the liquidus at the melt composition and whether the liquidus hits a peritectic line before reaching a eutectic point which would result in composition changes during growth due to macroscopic segregation of the components between crystal and melt at the growth interface. Fig.s 3—5 present metallographic pictures of cross sections through crystals grown by the Czochralski and Bridgman methods using different melt compositions. Fig. 6 shows the results of precision density measurements on sections of a single phase y’ crystal grown by the Bridgman method from a melt with an atomic composition of Ni(76)Al(24) [13]. A Sartorius semi-microbalance was used, with which the densities have been determined by hy.

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growth experiments. During crystal growth at noncongruent points on the phase diagram, the ~tr~i~

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\1etallograph~ofcross section through Bridgman crystal grown from Ni(77)Al(23) melt with 10 mm/h pulling rate. Bright: ‘1-phase; dark: y-phase.

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\“~ results, expect for the liquidus temperatures at XNj <75 at~. The phase diagram of AV differs from S and our results: peritectic and eutectic lines in AV are located at least at 20 K higher temperatures: —the peritectic concentration of the Ni 3Al phase and the eutectic concentration are inversed on the concentration coordinate, Due to the limitations in precision of the measurements shown in fig.s 1 and 2, the position of the peritectic point cannot be fixed unequivocally with respect to the eutetic point. We have, therefore, performed an additional series of crystal

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cision up to flp/p © 4 X 10 (where p is the density), depending slightly on the volume of the specimen. The typical sample volume amounted to I cm~. The temperature of the immersion liquid was controlled to within ~ 0.010 C of 25°C.

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solidus lines (dT(liquidus)/d.v~, > 0) in the crystal until essentially only the y’-phase crystallizes finally. This behaviour is only possibly if the pentectic line in the phase diagram is defined by the coexistence of y ‘-phase. liquid and y-phase. In

addition, the slope of’ the liquidus (dT(Iiquidus)// dxN ) has to he positive to explain the segregation behaviour in fig. 4 and fig. 6. which leads to the 5. Phase diagram

crystal becoming enriched in aluminium during

the course of crystallization. The phase diagram in the Ni 3A1 region of the Ni—Al system shown in fig. 7 can be derived from our results with the following arguments. especially as to the question of the sequence of pentectic and eutectic points:— Crystals grown from Ni(75)Al(25) melts (fig. 3) show coexistence of /3 and ‘y’-phases over the whole length indicating that the crystallization operates near the eutectic

of (8)

+ (y’).

Crystals grown from Ni(76)Al(24) melts (fig. 4) show a completely different behaviour: initially y’- and y-phase regions grow together. During growth we observe an increase of the amount of

y’-phase which corresponds to a segregation hehaviour with positive slope of the liquidus and

Crystals grown from Ni(77)Al(23) show coexistence of y’- and y-phase over extended lengths. —

Therefore, this melt composition must he located

between the peritectic concentration of the

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phase and the boundary line of the y-phase at low concentrations of nickel. Cooperative growth of these two phases is possible with the same crystal-

lographic orientation because the edge length of the cubic unit cells of both phases differ only by ~ Obviously, the phase diagram in fig. 7 would have to he extended to the metastable regions + liquid) and (‘y + liquid) to explain cooperative growth near the peritectic point properly in a similar way as cooperative growth of two phases ~

near an eutectic point. --

The concentrations of nickel in the three phases

/3. y’ and y involved in the crystal growth experiments differ only slightly. Therefore, it is very difficult to analyse the question of position of peritectic and eutectic points by conventional

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methods of n1etallographic and chemical analysis. By utilizing the segregation behaviour during crystal growth, subtle topological differences in

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at % Ni Fig 7 Phase diagram of part of the t~i—Alsystem as derived from this work.

The unusual mechanical strength of the inter— metallic compound Ni Al is caused ho the atomic long range order. The extra lines which appear in the diffraction pattern of an ordered alloy are called superlattice lines. The very low attenuation of y—rays in matter allows one to study the teill— perature dependence of ordering by measuring the

Bragg diffraction intensity of a sample located in a furnace. The mntensit\ of a supenlattice reflection

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Experimental analysis of Ni—Alphase diagram

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between the disordered y’-phase and the y-phase seems to be only the lattice constant. 7. Conclusions

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The growth of crystals from a peritectic phase

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possibilities for difficult. successfulIngrowth e,dst: is usually quite the case of Ni3A1, two Growth from melts with compositions between

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the peritectic and eutectic points in the phase

where rAI is the probability that an Al atom site is filled by an Al atom and FAI is the atomic fraction of Al in the alloy. S varies between 0 (complete disorder) and I (complete order). In the case of

diagram (fig. 7). This requires exact control of melt composition, which is difficult over extended crystal lengths due to the incongruent crystallization which leads to macrosegregation. Growth from melts with a composition at the maximum of the metastable (y’ + liquid) phase lines, i.e. near the peritectic composition. Since this composition of congruent growth cannot be derived directly from the phase diagram, this method requires a series of test runs. Two-phase single crystals containing the y- and ‘y’-phases can be grown from melt composition Ni(77)Al(23). An additional annealing at 1250°Cfor 50 h turns the single crystal into pure y’-phase. Recent results of Hilpert et al. [15] in determining the phase diagram of Ni—Al by means of high precision differential thermo-analysis and Knud-

y-ray diffraction, the value for S can be obtained

sen effusion mass spectrometry confirm the char-

in the kinematical approach from the relationship: 1 exp(—2M12001) 3fN1 2 2 m R°°° fAm R~2001exp(~2M°001) 9(toO) 9(200) fAi fN, 2

acteristic topology near the Ni7AI composition reported in this work.

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Fig. 8. Superlattice (100) Bragg diffraction integrated intensity and fundamental (200) Bragg diffraction integrated intensity as a function of temperature measured by y-rav diffraction using a radiation of 316 keV.

is correlated with the long range order parameter S. defined as [14] S

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References where ~ is the measured integrated reflecting power of a hkl reflection, 9~ the Bragg angle, exp( M) the Debye—Waller factor and .t~ the atomic scattering factor. The variation of the integrated intensity with temperature for a Ni(77)Al(23) single crystal is shown in fig. 8. The decrease of the (200) reflection intensity is due to the Debye—Waller factor. The (100) Bragg diffraction intensity is nearly temperature independent below 1300°C, but then declines at 1330°C during heating, indicating the existence of an order—disorder transition of the y ‘-phase slightly below the melting point. The structural difference —

[11 CC. Koch, CT. Liu and N.S. Stoloff, in: High-Temperature Ordered Intermetallic Alloys, MRS Symp. Proc., Vol. 39 (Mater. Res. Soc.. Pittsburgh. PA, 1985). [21W.O. Alexander and NB. Vaughan, J. Inst. Metals 61 (1937) 247. [3] J. Schramm, Z. Metallk. 33(1941) 374. [4] A.J. Bradley and A. Taylor, Proc. Roy. Soc. (London) A159 (1937) 56. [51A. Taylor and R.W. Floyd, J. Inst. Metals 81(1952/53) 25. [61W.R. Floyd, J. Inst. Metals 80 (1951/52) 551. [71H. Grober and V. Hauk, Z. Metallk 41(1950) 283. [81 I. Baker, J.A. Horton and F.M. Schulson, Metallography 19 (1986) 63.

192

F]. Bremer ci a!,

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Experimental ana/rsn.s oJ Ni—Al phase dmagr’mmrn

Brice, Crystal Growth Processes (Black Halsiead.

[13) 1’. Karasinski ci al.. Kernforschungsanlagc .lfihnch. Report

Glasgow, 1986). 10] A. Hellwig. Diploma Work. RWTH Aachen and KFA

No. JlJL-2136 (1987), [14] RD. (‘ulliiy. Elenicnts of X-Ray I)ifiractionr. 2nd ed. )Addison-Weslcs’. Reading. MA, 1978). [15] K. I-lilpert. t). KohcrtL. V. \‘cnugopal. N-I Miller, H. Gerads, F..!. Brcmer and N. Nickel. 7. N:miurforsch. 42a (1987). 16) NV. Ageev. Handbook of Binary Metallic Sssienis, Viii. I (I 966).

J’/mlich (1986). [11] F. Karthaus, Diploma Work. RWTH Aachen and KFA J’iilich (1986). [12] W. Kurz and D.J. Fischer. Fundamenials of Solidification

(Trans. Tech.. Aedermannsdorf, 1984).