Effect of alloying addition on structure and thermal stability of Ti40Al60 prepared by mechanical alloying

Effect of alloying addition on structure and thermal stability of Ti40Al60 prepared by mechanical alloying

Vol. 42, No. 2, pp. 411-417, 1994 Printed in Great Britain. All rights reserved 0956-7151/94 $6.00 + 0.00 Copyright © 1994 Pergamon Press Ltd Aeta m...

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Vol. 42, No. 2, pp. 411-417, 1994 Printed in Great Britain. All rights reserved

0956-7151/94 $6.00 + 0.00 Copyright © 1994 Pergamon Press Ltd

Aeta metall, mater.

EFFECT OF ALLOYING ADDITION ON STRUCTURE A N D THERMAL STABILITY OF Ti40A160 PREPARED BY MECHANICAL ALLOYING W. GUO, S. MARTELLI and M. MAG1NIt Amorphous Materials Project, E.N.E.A. Casaccia Dep. INN-PCM, C.P. 2400 A.D., 1-00100 Rome, Italy (Received 23 March 1993; in revisedform 28 June 1993)

Abstract--Ternary (Ti4oA160)100_xMx (X= 5-15at.%, M =Mo, Nb) alloys have been prepared by Mechanical Alloying, and the process was monitored by X-ray diffraction technique. The effects of a third additional element have been examined concerning the alloying process, structure and phase thermal stability. As already observed in Ti40A160matrix prepared with the same conditions, an amorphous alloy was obtained at the end of the process. A different solubility of Mo (very limited) or Nb (total dissolved) into the matrix was detected. Upon thermal treatment the third element addition caused, in both cases, an increased crystallization temperature with respect to the matrix. The more detailed investigation on niobium addition (5-10at.%) evidenced the following path for powder crystallization: amorphousdisordered Al-ordered LI 0 phase. The evolution of the long-range order parameter revealed that the disorder ~ order transition is favoured due to the presence of Nb, but independent of its concentration.

1. INTRODUCTION

By rapid quenching it has been found that the LRO (long range order) could be partially reduced [10] with a consequent improvement of the low temperature ductility [12]. By mechanical alloying of pure elemental powders, followed by an adequate thermal treatment, a completely A1 disordered TiAI intermetallic has been successfully prepared [13]. The A1 structure possesses a high symmetry with respect to the L10 one, and it can be expected that the random atom distribution would restore the slip systems of the f.c.c, structure. However, this face-centred disordered structure holds only up to 740°C, beyond which a thermal induced disorder-order transition takes place [13]. Such a temperature value lies far below the one reached in extruding or compacting techniques (1000-1200°C). It was of primary interest in this work to check whether the thermal stability of the A1 disordered TiA1 phase could be improved by adding a third component during MA. The high melting point metal elements niobium and molybdenum have been selected. Their effects on structure and phase stability are discussed here in light of the present experimental results. The degree of disorder of the Ti-AI ), phase may be followed by measuring the tetragonality of the unit cell. In general, the Ll0 structure made from two kinds of atoms has an axial ratio c/a less than unity when the atomic size is the factor governing the structure. The ordered TiA1 phase has an axial ratio of 1.025-1.030 [14, 15] showing thus that the covalent directional nature of the bonds is prevailing in this case. Any disordering of the structure may be directly followed by measuring the axial ratio and the

Intermetallic titanium aluminides are currently investigated for high temperature structural applications. Indeed, the ordered TiA1 and Ti3A1 alloys increase the yield strength by increasing temperature [l~i] while the restricted mobility, due to the localised covalent bonding in the ordered lattice, inhibits the diffusion processes and enhances the creep resistance. Between them, the TiAl alloys are especially attractive for their high melting point, favourable strengthto-weight ratio and good oxidation resistance. The structure o f TiAl is face-centred tetragonal, L10 (also known as ~ phase) having alternated layers of Ti and A1 atoms along the (002) plane. The rigid covalent bonds between atoms are claimed as responsible of the good mentioned mechanical properties. On the other hand, they are also responsible of a limited number of slip systems thus making the intermetallic brittle and with a low ductility at room temperature [7-10]. If the L10 structure is disordered, it approaches the face-centred cubic structure, and, hence, it can be expected to increase its ductility. However, conventional processing is not suitable to that purpose since the alloy remains ordered up to its melting point and forms without stepping through a high temperature A l f.c.c, disordered phase [11]. In the last years, then, attention has been paid, in order to improve ductility, both to the alloy design (micro and macro addition of elements, aimed thermal treatments in order to refine microstructures) and to their processing (rapid solidification, mechanical alloying). tTo whom correspondence should be addressed. AMM 42/2--F

41

412

GUO et al.: STRUCTURE AND THERMAL STABILITY OF Ti40A160

consequent decrease of the LRO. This criteria has been considered in the present work.

2. EXPERIMENTAL DETAILS

2.1. Milling procedure Based on our previous results [13], ternary alloys were prepared with nominal compositions of: (Ti0.4A10.6)100_xMx, where M = Mo, Nb; x = 5, 10, 15 at.%. The M A process was performed in a conventional planetary ball mill (Fritsch "Pulverisette 5"). Pure Ti (99.0%, - 3 2 5 mesh), AI (99.3%), Mo (99.9%, - 2 5 0 mesh) and Nb (99.8%, - 3 2 5 mesh) were all Alfa products. 10 g charges of appropriate powder mixtures were loaded into cylindrical hardened steel vials, together with corresponding stainless-steel balls of 8 mm in diameter. The ball to powder ratio was approximately 10:1. In order to minimize the oxygen and nitrogen contaminations, all the operations, i.e. loading and sampling, have been done inside a controlled argon atmosphere glove-box (02 < 3 ppm, H 2 0 < 5 ppm). The vials were sealed under overpressure argon ( > 1.5 atm) during milling procedure. The ball milling was carried out at the velocity of 250 rpm with interval of 20 min each followed by an equal rest time. To minimize the average temperature increase of the whole system the vials were cooled by compressed air.

2.2. Thermal treatments The usual commercial titanium powders do contain a given amount of hydrogen due to the preparation method of this element. In view of this fact, the hydrogen evolution, normally registered in the thermograms at about 500°C as an endothermic process, completely masks any feature obtained from the differential scanning calorimeter (DSC) measurement, so that DSC measurements could not be used in the present work in order to determine both crystallization and disorder-order transition temperatures (see next section). A Perkin-Elmer DSC 7, however, has been used to perform thermal treatments up to 700°C. The heating rate normally used was of 40°C/min. After the heating ramp up to the desired temperature value, cooling down at room temperature was governed by the

instrument itself (200°C/min). Pure argon was used in all cases as a covering gas. By the above procedure a large number of DSC have been performed by varying the maximum temperature reached and doing soon after X-ray examination in order to detect the products obtained by thermal treatment. We have been able in the way described to determine the transition temperatures discussed in the next section and reported on Table 1. Due to the pro' cedure adopted, the temperature values reported on the table have a confidence range of about 10°. Thermal treatments over 700°C were performed using a tubular oven with an argon atmosphere. The heating and cooling rates were, in these cases, the same adopted with the Perkin-Elmer instrument. The oxygen and nitrogen contents of the milled and thermal treated samples have been measured with Leco TC 136 gas analyser.

2.3. X-ray measurements The mechanical alloying process was monitored by X-ray diffraction patterns using automatic SEIFERT (PAD VI) and SIEMENS (D500) diffractometers, equipped with Mok~ and Cuk~, radiation respectively. Owing to the milling action, the X-ray patterns are usually composed by broadening and overlapping peaks. In order to more precisely estimate the positions and shapes of single component, the recorded patterns were analysed by a non linear least square fitting procedure, using an ensemble of generalized Lorentzian curves [16]. 3. RESULTS AND DISCUSSION

3.1. Behavior of the Ti4oAl6o composition milled with a third element Milling the Ti40A160 mixture alone, a single amorphous phase has been obtained as fully described in a previous paper [13]. Figure l(a) shows the effect of 36 h of milling on the Ti40A160 powder mixture with different additions of molybdenum powder. It can be seen that the Mo peaks are still clearly recognizable in all the three cases together with the presence of an amorphous halo quite similar to the one obtained with Ti4oAl60 mixture alone. The addition of Nb, on the contrary, promotes the formation of a single amorphous phase and no evident traces of the third

Table 1. Crystallizationtemperaturefor mechanicalalloyed(Ti0.4Al0.~)~0o_xMx (M = Mo or Nb) powdersas a function of addingconcentration.To be noticedthat the percentageonly representsthe stoichiometricstartingcomposition(see text) Composition Crystallizationtemperature(Tx) Orderingtemperature(TD/o) (at.%) (°C) (°C) (Tio.4Al0.6)100 xM~. X= 0 440 700 M ~ Nb X=5 480 575 X = 10 550 575 M ~ Mo X= 5 460 -X = 10 480 -X = 15 510 --

GUO et al.:

STRUCTURE AND THERMAL STABILITY OF Ti4oA16o

(a)

(b) .........

@-Mo

(Ti~Al~)mo-xU°'

• -Ti •

413

l,,,j,,,,,l

.........

• - disordered TiAI

S6h

-AI

{• -

~

MO

I

.........

(Tio~Aloz) 1oo-'M°, •

x=15

_

5" /

I0

15

20

25

30

35

10

15

20 20

2O

25

30

Fig. 1. (a) X-ray diffraction patterns of the (Ti0.4Alo.6)z0o_xMox ( X = 5 , 10) after 36h of milling, The components of the first broad peak, obtained by the fitting procedure, are shown for the three cases examined. (b) X-ray diffraction patterns of the samples in (a) after thermal treatment at 600°C.

component can be detected for both compositions examined after 30 h of milling (see Figs 2 and 3 curves b). The position of the amorphous halos with the addition of both Mo and Nb keeps substantially unaltered, within the experimental uncertainty, with respect to the Ti40A160 alloy. The experimental findings indicate therefore a quite different solution behaviour of molybdenum and niobium. Niobium completely takes part in the alloying process, molybdenum does not. From the steric point of view the atomic radii of

the two third elements are very close (Nb = 0.146; M o = 0 . 1 3 9 n m ) to the average atomic radius of Ti-A1 (0.145nm) so that both elements could be equally well accommodated in the matrix structure. Concerning the affinity, the phase diagrams of niobium and molybdenum with aluminium and titanium are very similar. In the examined composition range, both elements form solid solutions with titanium and intermetallics with alluminium. The diffusivities of niobium and molybdenum in A1 are similar (T = 753 K, DMo.AI = 1.64 x 10 13 cm2/s,

_************************************************** i"-- ordered T i A I disordered TiAI Ti Nb

• ) ._,

. ,_~

(Tio.4AIo.e)~Nb n

_%_."

% _J~

Ooot~II-

ordered T i A I disordered TIAI Ti

(Tio.,Alo~).oNbx °

Nb

AI

o

aooc

7 ~'J

b) ~

_

b) ~

30h

•)

~:

~.Oh

' " " .... i ' " ...... i ....... " i ......... i " ' " ....

I0

15

575Cx 10'

20

25

30

35

20 Fig. 2, X-ray patterns of the (Ti0.4Alo.6)asNb~ composition at: (a) starting element powders; (b) after 30 h of milling; (c) thermal treatment up to 575°C; (d) thermal treatment

up to 575°C and 10 min of rest at the same temperature (see text); (e) thermal treatment up to 900°C. In (c) the fitting procedure for the separation of the single peaks is shown.

.

30h

•) ~

~'''''IZ)'H'ITII'J'['III'IIJ'I~''I

I0

15

20

25

oh Jll~rll'lllll

30

35

20 Fig. 3. X-ray diffraction patterns of the (Ti04Al06)90Nb,0 composition at: (a) starting elemental powders; (b) after 30 h of milling; (c) thermal treatment at 575°C; (d) thermal treatment up to 575°C and l0 rain of rest at the same temperature (see text); (e) thermal treatment up to 900°C. In (c) the fitting procedure for separation of the single peaks is shown.

414

GUO et al.: STRUCTURE AND THERMAL STABILITY OF Ti40Alr0

3.29 x 1 0 -13 cn'12/s). The diffusion coefficients for both elements in cz-Ti are not known, but those in fl-Ti are again very similar (T = 1273 K, DMo~rri = 3.66 x 10-1°cmE/s, DNb~fri=8.95 × 10 -1° cme/s). The major great differences between the two 3rd elements appear when comparing their mechanical properties. Molybdenum possesses unusually high melting point, and is also one of the hardest metals (98 B). Its Young's modulus (329.57 GNm -2) and its shear modulus (119.56 GNm -2) are about two times those of niobium (hardness 71 B, Young's modulus 156.8 GNm -~ and shear modulus 58.80 GNm-2). It is well accepted the mechanism of amorphization by ball milling proceeds, whatever be the driving force (free energy gain, defect enhanced diffusion, etc.), through the repeated formation of clean element/ element surfaces at which the solid state reactions start up [17]. The unusual molybdenum hardness combined with the high strength probably make ineffective the shear stress applied during milling to the particles powder. So that, while Ti and AI proceed in the alloying, molybdenum atoms are practically inhibited.

0,40S

DNb~A I =

3.2. Thermal treatments

Thermal structural stability of the final mechanical alloyed powders has been investigated by heattreatment in the 500-900°C temperature range. The X-ray diffraction patterns of the milled samples after thermal treatments are shown in Fig. l(b) for molybdenum added to the matrix composition and Figs 2 and 3 (patterns from c to e) for added niobium. Table 1, column 2 gives the crystallisation temperature (T~) for the five samples investigated and column 3 the disordering--ordering temperature (TD/o) for the niobium samples. It must be outlined the different meaning of the percentage of the third element added reported in Table 1. Niobium atoms are completely dissolved into the amorphous matrix and its percentage represents the actual alloyed niobium content. The molybdenum percentage only represents the stoichiometric starting composition since the amount taking part in the alloying process is certainly very little and, in any case, unknown. Amorphous-crystalline

(disordered)

transition.

Figures l(b), 2 and 3 (patterns c) indicate that the thermal treatment promotes the transformation from amorphous to crystalline phase. The diffraction lines can be indexed by reflexions with either all odd or all even indices corresponding to the A1 cubic face centred structure. The crystalline phase formed, TiAI(M) (M = Nb, Mo), is therefore disordered with Ti, AI and the third element randomly distributed over all atomic sites. The third element addition stabilizes the amorphous phase with respect to the matrix composition as indicated by the shift of higher values of the transition temperature T~ (see Table 1). The diffusion coefficient of A1 in Ti (T = 900 K, DAI~Ti= 9.3 x

I

I

I em-

Nb Mo

0.404

~.~ 0 . 4 0 3

e, ~

0.408 i

0.401

-

0.400

I 5

0

M

I 10

Composition

I 1,5

20

(at.7,)

Fig. 4. Lattice parameter of disordered TiAI(M) as a function of adding concentration.

10 -9 cm2/s is much higher than those of Nb and Mo in Ti. Thus the addition of a low mobility element to the Ti40A160 matrix is responsible for the increased crystallization temperature. It is then reasonable to conclude that the nucleation of the disordered TiAI(M) phase requires the short-distance rearrangement of all three Ti, AI and M atoms. The lattice parameters of the disordered phase is shown in Fig. 4 as a function of the third element addition. Considering the atomic radii of Egami and Waseda [18] (i.e. the radii determined in metals or amorphous alloys) we would have ( r ) T i A I = 0.145 nm, rNb ~ - 0 . 1 4 6 nm. The niobium addition should be therefore ineffective on the lattice parameters. Also molybdenum, with an unknown but surely limited content in the disordered phase and an atomic radii of 0.139 nm, should not perturb the lattice parameters. The experimental evidence, on the contrary, indicates a not negligible effect of the third element on the lattice parameters. Therefore, very likely, the kind of bonds occurring in the phase, i.e. the electronic structure, plays a major role than the atomic size as already observed by Vujic et al. [10]. Disorder-order transition. In the following we will focus our attention on the samples with added niobium since the unknown quantity of molybdenum dissolved in the matrix alloy does not allow any further consideration with this 3rd element. Figures 2 and 3 (patterns d) show the onset of disorder-order transition. The nucleation of the tetragonal ordered L10 TiAI(Nb) structure is indicated by the appearance of the superlattice peaks in the low angle part of the diffraction patterns ({110} reflex) as well as, at higher temperatures (patterns e) by the splitting of the {200} reflection into the two components {200} and {002}. The overall process from the amorphous to the crystalline ordered state is governed by two kinetic processes as clearly indicated by X-ray patterns c and d (Figs 2 and 3). A

GUO et al.: STRUCTURE AND THERMAL STABILITY OF Ti4oAlr0

higher defect density is originated with Nb addition than without. Hence the decrease in the ordering temperature could be explained by an enhanced Ti and AI interdiffusion due to lattice defects. Of course more structural investigations are needed in order to fully clarify this point. The D/O transition, however, seems to be ineffective with respect to the amount of added niobium (see Table 1). Additional information with respect to this point can be drawn by the observations given in the next section.

normal thermal treatment up to 575°C promotes the amorphous-,A1 transition. Standing for 10min at the same temperature starts up the A I - , L 1 0 transformation. It is worth noting that, at 575°C, the amorphous--,Al transition is more advanced for the 5% Nb (lower Tx) than for the 10% niobium composition (patterns c in Figs 2 and 3). The disorder-order transition temperature (TD/o) registered with added niobium is somewhat lower than that of the matrix alone (see Table 1). The Nb addition, thus, while enhancing the stability of the amorphous phase due to its low mobility, favours the ordering kinetic. To this respect the niobium addition is not at all an improvement to extend at higher temperatures the stability of the disordered state. Vujic et al. in studying the (Ti45Alss),0o_xV x alloys obtained by rapid solidification [10], find a dislocation density higher than that of the matrix. Considering the close similarity of niobium and vanadium (electronic structure, mechanical properties), one can suppose that during the milling process a much (a) 0.405

I

I

I

I

3.3. Tetragonal distortion and long range order (LRO) The stoichiometric ordered-TiA1 phase has an experimental axial ratio of 1.025-1.030 [14, 15]. From the experimental axial ratio of the ordered phase, it is argued that the Ti-A1 bond is longer than Ti-Ti and A1-A1 ones. By raising the temperature of the thermal treatment, the ordered fraction increases leading to an increased number of Ti-A1 bonds along (011) and Ti-Ti and AI-AI bonds along (110) (b)

I

( T i o.4A 1o.e) too-x N b

I

I

I

X=5 X=lO

O-

~

I

(TI o.4Ai o.e)too-xNbx

0.410

0.404 0.403

415

-

0.408

0.402 E~ o.4ol

£

•",, 0.406

C v

0.400

0.399

o 0.404

!

0.398 0.402 0.397

X=5 X=IO

l

0.398 400

I 500

I 000

I 700

Temperature

I 800

I 900

0.400

] 0oo

500

1000

(°C)

Temperature

(c)

1.04

I 700

I

I

I

I

I 800

I 900

I coo

I 0oo

1ooo

(°C)

(Tio.4Alo,e)too_xNbx 1.03 -

O- X=5 II=

¢~ 1.02 o

1.01 -

1.00

500

i 600

I 700

Temperature

1000

(°C)

Fig. 5. Lattice parameters and their axial ratio for the (Ti0.4Al0.6)~_xNbx compositions as a function of thermal treatment temperature. (a) a; (b) c; (c) c/a.

416

GUO et al.: STRUCTURE AND THERMAL STABILITY OF Ti4oAl~o I

O.gO

I

where rri and (1 - rri) are the fractions of Ti and A1 atoms occupying the Ti sites whereas rm and (1 - rm) are the fractions of AI and Ti atoms in the A1 sites. Since according to Warren [19] the LRO parameter is defined as

I

(TI o.4A1o.e) lao_xN b x

o')

~2

0.80

S=

rTi -- XTi

r,~--X~

(4)

- - -

where XTi and X~j are the atomic fractions in the alloy and Y~u, YT~ the fractions of Ti and AI sites in the superlattice (for L10 Y~j = YTi = 1/2), it follows that the F(hkl) factor takes the form

~ 0.70

F(hk0 = [(S/2 + X¢i)fT~ + (1 - X¢~ - S/2)f~] 0.fl0 °

x {exp[ni(h + k)] + 1}

5BO

I 050

I 750

I Ca0

+ [(S/2 + XAI)fAa + (1 -- Xju - S/2)fxi ] 950

x {exp[~ti(h + 1)] + exp[ni(k + 1)]}.

Temperature (°C) Fig. 6. Plot of long range order parameter S vs thermal treatment temperature for ( T i 0 . 4 A l o . 6 ) t 0 o _ x N b x (X = 5, 10 at.%) compositions.

(5)

Remembering that for fundamental peaks h, k and l are unmixed and for superlattiee ones, h + k = even, h + l = odd, equation (5) takes the form Ff = 4(XTifn + X,J'~u)

directions. It is then expected a contraction of a and an expansion of c structural parameters by raising the temperature. This has been indeed found and the results from least squares fitting of the unit cell of the ordered L10 phase are reported on Fig. 5(a--e). The ordering structural evolution with temperature may be followed by measuring the long range order (LRO) parameter. The LRO may be determined from the ratio of the intensities of superlattice and fundamental reflexions. The integrated intensity of one reflection is given by I(hkl) = Ir(hkl)[ 2 * P(hkl) * L(hkl) * D(hkl)

(1)

where L(hkl) is the Lorentz-polarization factor. L(hkl) = (I + cos 2 20)/sin 2 0 cos 0

0 being the Bragg angle, P(hkl) is the multiplicity factor and

Fs = 2S0Cxi -fla).

(6)

We have previously seen that, after thermal treatment, a unique phase, with Ll0 structure, is formed with Nb addition. The Warren method, therefore, can be utilized also for TiA1-Nb ternary alloys. However, before using the previous equations, the atomic site occupation of Nb atoms has to be determined. Kawabata [20] and Konitzer et aL [21] in their works observed that niobium substitutes only for Ti sites. We have the same evidence with our data: e.g. the sample thermal treated at 900°C, shows the ratio (Idl{m}) higher than the theoretical one for ordered Ti40Al60. It is reasonable, therefore, to assume that, as soon as Nb atoms get sufficient mobility, they would replace preferentially Ti atom at the Ti sublattice sites. With this assumption, the atomic scattering factor of titanium in equation (6) must be replaced by

D (hkl) = e x p [ - B(sin 0 {22)]

fTi~

XTLfTi "~ XNhfN b

(7)

XTi " + X N b

where B is the Debye-Waller factor. The structure factor F(hkl) can be generally expressed by F(hkl) = ~ (,fn ) exp[2ni(hu, + kv, + lw,)]

and Xri by XTi + Xt~b, where XNb is the atomic fraction of Nb. Equation (6) then becomes

(2)

r~ = 4(X~d~i + x ~

+ XNJNb)

n

in which ( f n ) is the atomic scattering factor, u,, v,, and w, are the position coordinates of atom n in the unit cell. The Llo TiA1 structure may be subdivided into four simple cubic sublattices whose origins can be located at (0 0 0); (ii0) 11 for Ti and (~10 9, 1. (0 ~11i) for A1. Therefore equation (2) becomes

Fs = 2S(XX~lii + ~ N b

fA,)"

(8)

Considering that D(hkl) is close to unity at room temperature for small angle [22], the expression of S 2 is given by S:=4,~, /r

F(hkl) -- [rr~fri + (1 - rT~)f~]{exp[Iti(h + k)] + 1}

(PL)f * (eL)s"

+ [r~JAl + (1 -- rAa)fTi]{exp[gi(h + 1)] + exp[ni(k + 1)1}

(3)

(9)

Equation (9) has been used to obtain the long

GUO et al.: STRUCTURE AND THERMAL STABILITY OF Ti40Alr0 1.035 /

I.o3o

1

/-~

I

I

I

[

I

I

I

the amorphous matrix Ti40Alroas indicated by the changing of the crystallization temperature of the matrix itself. • Niobium completely takes part in the alloying process with the results of a unique amorphous phase. • By thermal treatment the amorphous phase evolves first towards the disordered intermetallic with the cubic A1 structure and then to the ordered Llo structure. • Niobium addition stabilizes the amorphous, whereas reduces the thermal stability of disordered TiA1 phase.

(Tio.,AIo.e)1oo_xNbx e-

X=5

/ /

i 1.025 ~

I

l-- X = l O

x.oz0 ~

/

,.o,o_

//

417

//

REFERENCES 0.0

0.2

0.4

0.6

0.8

1.0

LROP, S

Fig. 7. Plot of c/a vs S for the two compositions with Nb addition. range order parameter for both niobium ternary alloys. The results are shown in Fig. 6 as a function of the heat treatment temperature. The more relevant result is that the LRO parameter is practically the same for both 5 and 10% of niobium. The ordering kinetic is favoured by niobium but it is not sensitive to the niobium quantity at least in the examined composition range. The result was expected since equal To/o (see Table 1) substantially indicates equal ordering kinetic. In binary systems the correlation between tetragonal distortion and LRO is usually linear like in AuCu [23], Cu3Pd [24] and rapid-quenching Ti45Al~5 [10] alloys. However, in present work with third addition, the relationship between c/a and S is no longer linear (Fig. 7) and may be written as follows c/a = 1 + m S "

where m and n are 0.0427, 0.0471, and 2.990, 2.966 for 5 and 10% Nb, respectively. The addition of niobium substantially does not alter the tetragonal distortion of the matrix (1.025-1.030) confirming thus that the electronic structure plays a role more important than the atomic size.

4. CONCLUSIONS The Ti40A160has been alloyed with the addition of molybdenum or niobium. The experimental finding indicate that: • Molybdenum does not take part in the alloying process even though a fraction of it dissolves into

1. P. H. Thornton, R. G. Davies and T. L. Johnston, Metal. Trans. 1, 207 (1970). 2. S. M. Koply and B. H. Kear, Trans. Anal. Inst. Min. Engrs 239, 977 (1967). 3. N. S. Stoloff and R. G. Davies, Acta metall. 12, 473 (1964). 4. S. J. Liang and D. P. Pope, Acta metall. 25, 485 (1977). 5. C. Lall, S. Chiu and D. P. Pope, Metall. Trans. 10A, 1323 (1979). 6. T. Kawabata, T. Kanoi and O. Izumi, Acta metall. 33, 1355 (1985). 7. D. Shechtman, M. J. Blackburn and H. A. Lipsitt, Metall. Trans. 5, 1373 (1974). 8. H. A. Lipsitt, D. Schechtman and R. E. Schafrik, Metall. Trans. llA, 1369 (1980). 9. S. M. L. Sastry and H. A. Lipsitt, Acta metall. 25, 1279 (1977). 10. D. Vujic, Z. X. Li and S. H. Whang, Metall. Trans. 19A, 2445 (1988). 11. J. L. Murray, Binary Alloy Phase Diagrams (edited by T. B. Massalski), p. 173. ASM, Metals Park, Ohio (1986). 12. D. K. Chaterjee and M. G. Mendiratta, J. Metals 5, 33 (1981). 13. S. Martelli, W. Guo, A. Iasonna, M. Magini, F. Padella, E. Paradiso, N. Burgio and M. della Porta, Advanced Structural Materials (edited by T. W. Clyne and P. J. Withers), Vol. 2, p. 379 (1991). 14. P. Duwez and J. L. Taylor, J. Metals 4, 70 (1952). 15. E. S. Bumps, H. D. Kessler and M. Hansen, J. Metals 4, 609 (1952). 16. D. W. Marquard, J. Soc. Ind. Appl. Math. 91, 431 (1963). 17. R. B. Schwarz and R. R. Petrich, J. less-common Metals 140, 171 (1988). 18. T. Egami and J. Waseda, J. Non-Cryst. Solids 64, 113 (1984). 19. B. E. Warren, X-ray Diffraction. Addison-Wesley, Reading, Mass. (1969). 20. T. Kawabata, High Temperature Order lntermetallic Alloys III (edited by C. T. Lin et al.), M R S Proc. 133, 329. 21. D. G. Konitzer, I. P. Jones and H. L. Fraser, Scripta metall. 20, 265 (1986). 22. H. Lin and D. P. Pope, J. Mater. Res. 763, 5 (1990). 23. A. H. Wilson, Proc. Camb. Phil. Soc. 34, 81 (1938). 24. D. Madoc Jones and E. A. Owen, Proc. Phys. Soc. 1367, 297 (1954).