ON THE AMORPHISATION REACTIONS OF CuTi BY MECHANICAL ALLOYING

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ON THE AMORPHISATION REACTIONS OF CuTi BY MECHANICAL ALLOYING

G. Cocco, I. Soletta, S. Enzo, M. Magini, N. Cowlam

To cite this version:

G. Cocco, I. Soletta, S. Enzo, M. Magini, N. Cowlam. ON THE AMORPHISATION REACTIONS

OF CuTi BY MECHANICAL ALLOYING. Journal de Physique Colloques, 1990, 51 (C4), pp.C4-

181-C4-187. �10.1051/jphyscol:1990422�. �jpa-00230782�

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COLLOQUE DE PHYSIQUE

Colloque C4, suppl6ment au n014, Tome 51, 15 juillet 1990

ON THE AMORPHISATION REACTIONS OF CuTi BY MECHANICAL ALLOYING

G. COCCO, I. SOLETTA, S. ENZO*, M. MAGINI"* and N. COWLAM*'*

Fipartimento di Chimica, Via Vienna 2 . I-07100 Sassari, Italy Dipartimento di Chimica Fisica, 0.0.2137, I-30123 Venezia, Italy

"'ENEA Tib, CRE-Casaccia, Roma, Italy

* * t

Department of Physics, University of Scheffield, GB-Scheffield S1 350, Great-Britain

Resume

Les alliages des poudres de Cu et de Ti ont bte realise au moyen d'un moulin a bille.

Dans un intervalle de concentration 45< Cu at% 5 7 5 la phase amorphe se dkveloppe directement B partir des Blements d'origine, alors qu'a basse concentration, cette meme phase se verifie avec l'apparition d'un composk cristallin de CuTiz et CuTi. Si le processus mecanique se prolonge, on aboutit a une transformation presque complete des composes cristallins en phase amorphe. On observe un mkcanisme d'amorphisation different lorsqufon arrive a un refroidissement a 253 K.

Abstract

Crystalline powders of Cu and Ti were mechanically alloyed by a ball-mill apparatus.

For 45< Cu at% 5 7 5 amorphous phases form directly from starting elements, whereas

CuTiz and CuTi interrnetallics develop concurrent with an amorphous fraction in the lower Cu content range. Further milling causes the transformation of the intermediate compounds into an almost complete amorphous condition. A different amorphisation mechanism is operative when the milling tool vial is cooled to 253 K.

1

-

INTRODUCTION

While neither a great deal of experimental results nor an increasing theoretical effort have yielded a definitive explanation of the Mechanical Alloying (MA) process, there is a firm basis for the belief that its mechanism can be successfully described in terms of thermodynamic quantities and kinetics constraints /l-2/. It was argued that To curves, which define the thermodynamic limits for polymorphous transformation, can also be used as a useful guide to identify possible glass forming regions when the amorphous phase is treated as an undercooled liquid, provided that single phase compound formations are supressed. Thus, To and glass transition temperature curves define a temperature composition range where the formation of an amorphous phase is accessible via solid state reaction.

A wealth of essentially consistent experimental data supports the val- idity of To criterion. The CuTi system, however, constitutes a noticeable exception. Actually, To curves calculated by Schwarz et a1 /2/ by making use of different models, intersect above the experimental glass transition temperature thus predicting no glass formation for this system. Similar behaviours have been reported by Murray /3/ and Massalski and Woychik / 4 / .

In contrast, Politis and Johnson / 5 / , found by MA an exceptional wide glass forming region extending from 0.10 to 0.87 Cu at%. This discrepancy was also underlined by Hellstern and Schultz /6/ who discussed the problem in terms of excess free entalpy partially stored in the highly defective crystalline lattice of the reactant elements. Weeber and Bakker /7/, by an extension of Miedema8s semiempirical model / 8 / , presented recently an entalphy diagram for metastable CuTi predicting an amorphisable region extending from 0.28 < xcu < 0.75, smaller than the experimental one.

Article published online by EDP Sciences and available at http://dx.doi.org/10.1051/jphyscol:1990422

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C4-182 COLLOQUE DE PHYSIQUE

It appears that the MA process in- CuTi is in need of more study. The purpose of this and a related paper /9/ is to give a more general framework concerning the CuTi MA process.

2

-

EXPERIMENTAL METHODS

Appropriate amounts of Cu (Alfa Product, purity 99.99 %, 200 mesh) and Ti (Alfa Product, purity 99.9 %, 325 mesh) were milled under argon atmo- sphere in a hardened steel vial with a Spex Mixer-Mill, model 8000. The ball (1/2Iq) to powder weight ratio was 8.3. The following compositions were

studied: C U ~

,

~^C T^U ~^~

,

~^~^C~^T^U ^~^~

,

^~^~^C ^~^T^U ^~^~

,

^~^~^C ^~^T^U ^~^~

,

^~^~^C ^~^T^U ^~^~

,

^~^~^C ^~^T^U ^~^~

,

^~^~^C ^~^T^U ^~^~

,

^~^~ ^~^T ^~ ^~ ^~

here after coded as Cu,, etc..

The powders were 'processed under two experimental conditions: in the former, reported as "room temperaturew, heat was removed by a fan working in front of the milling tool; in the latter the vial was inserted into a self- locking metal casing and cooled by means of a refrigerating liquid flowing through the resulting jacket at 253 K. Milling was carried out till selected times without interruption and structural characterization /10/ was performed at the end of each milling run. CuKo: radiation was employed.

Samples were also submitted to differential scanning calorimetry (DSc) analysis: a Perkin Elmer DSc-7 was used with purified argon as purging gas.

A heating rate of 0.5 K S-' was generally employed.

3

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RESULTS AND DISCUSSION

The structural changes effected on the starting mixtures after 16 hrs of MA can be evaluated from Figure 1 in which X-ray patterns are reported for the quoted concentrations. The amorphisation path relevant to 75, 70, 60, 55 and 50 Cu at % is accomplished after 16 hrs (Figure 1, left side). Moreover the behaviour observed at intermediate reaction times (not shown) is in agreement with that reported in /5/ for the same system: because of fragmentation phenomena and disordering effects, the line profiles of separate elements become broader, the crystalline phases disappear progressively and amorphous alloys form with X-ray maxima at an angular position intermediate between the most intense Ti 101 and Cu 111 peaks.

30 35 40 45 50 55 2 THETA

30 35 40 45 50 55 I

2

THETA

Figure 1

-

CuKa X-ray scans for several mechanically alloyed CuxTi samples after 16 h of milling at room temperature. Cux is quoted. Left side: for Cu at

kl-9

⁵⁰an X-ray amorphous state is achieved; right side: at lower Cu content the absence of Cu and Ti lines and the development of modulated features show the achievement of crystalline conditions.

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From this behaviour an amorphisation process can be inferred /7/ which occurs by interdiffusion similar to the mechanism operative in multilayered structure during the heating /11/. However, at higher Ti content, different levels of structural organization are achieved (Figure 1, right side):

crystalline Ti and Cu lines disappear but now, instead of an amorphous phase, new X-ray signals are observable for CU,.~, and further diffraction peaks grow in C U ~ ~ and CuJO samples. Thus, at the low Cu content side, the reported trend differs from the Politis and Johnson results /5/. In their paper Weeber and Bakker /7/ properly stress the role of milling equipment and consequent local conditions in determining different reaction paths.

However, an immediate comparison between our results with those by Politis and Johnson is not possible since they did not quote the milling apparatus employed.

2 THETA

33 44 55 I

2

THETA

Figure 2 - C u b X-Ray patterns of Cu TieO sample milled at room temperature after increasing milling times, quoted as hours. A :CA0; v :Ti.

Figure 3

-

Comparison between the CuKa X-Ray patterns of the C U ~ sample milled 16 h at ~ T ~ ~ ~ room temperature and after heating up to 623 K and subsequent quench. The X-ray pattern after an isothermal annealing at 823 K for l h is also reported (upper curve). Inset: DSC trace; the selected quenching temperature corresponds to the end of the crystallization signal :CuTi2; : gamma-CuTi.

The attempt to amorphise the Cu,,Ti,, is presented in Figure 2 where the X-ray patterns are now depicted as a function of milling time. Crystalline Cu and Ti broadened lines are clearly recognizable up to three hours concurrent with a diffuse halo which arises between Cu and Ti peaks. Thus in the early stages, the alloying process conforms to the reactions already described for higher Cu content samples. The formation of the amorphous phase after relatively short times of milling is particularly worthy of notice. After 4 h, Ti and Cu crystalline peaks are very weak, but now the main feature is the occurrence of two new signals which develop progressively. After 8 h the angular position of these new maxima can be already assessed, but their assignement to a definite intermetallic is hindered by occurence in this angular range of several diffraction lines relevant to various CuTi compounds. However, as it can be observed in

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C4-184 COLLOQUE DE PHYSIQUE

Figure 3, the X-ray pattern of the crystalline CuTi2 and gamma-CuTi compounds /3/, found in the sample after a DSc run up to the end temperature of the crystallization (see the inset of Figure 3), fits quite well that of the same sample before annealing, particulary as far as the CuTi2 lines are concerned. This analysis strongly suggests that CuTi, and CuTi intermetallics are already present in the untreated specimen and form directly as a product of the milling process. Thus, after an initial interdiffusion reaction common to Cu rich samples, after 4 h, the alloying path of C U ~ ~ enters a different regime in which the formation rate of the crystalline compounds overcome that of the competing amorphous phase.

Coming back to Figure 2, further milling brings to light new details.

After 32 h one can notice a smearing out of the X-ray signal of CuTi, phase and, as shown in Figure 4, the main contribution to the intensity dis- tribution now arises from the amorphous fraction. The maximum is located at 2-theta of about 40.90 and agrees with an amorphous alloy of Cu3*Ti6, composition /5,12/.

2 THETA

Figure 4

-

A possible decomposition of the main contributions to Cuka X-ray pattern from the sample milled 32 h. The crystalline trace still observable is relative to CuTiz phase.

The almost total transformation of intermediate CuTi, and CuTi into an amorphous state is not unexpected. In fact, it has been reported that milling of pure components can initially lead to intermetallic compounds before transforming into an amorphous phase /13-15/. Moreover Lee et a1 /16/

obtained an amorphous alloy by MA of CuTiz

+

CuTi intermetallics with the average composition C U ~ ~ T ~ ~ ~ . TO this regard, it is also worth noting that Askenazy et a1 /17/, after cold rolling of a melt-spun CuTi ribbon, observed the amorphisation of CuTi, and Cu3Ti2 crystallites formed in the glassy matrix during the quench. It is evident that a quite different mechanism explains the crystal-to-amorphous transformation occurring at this stage.

The formation of CuTi, and CuTi intermetallics implies an almost total reduction of the thermodynamic driving force which sustains interdiffusion phenomena. Going on with the milling the amorphisation reaction is driven by the destabilization of the crystalline lattice owing to defects introduced by shear deformations concurrent with an enhanced level of chemical disorder /18,19/. A further comment concerns with the annealing behaviour of this sample: it is remarkable and to some extent puzzling that CuTi, phase is not any longer the main crystallization product; this point will be touched later.

The sequence presented in Figure 2 shows that the amorphous state is not the final product of the milling process. In fact, after 42 h of treatment, new humps develop in the patterns which become more evident after 52 h. The identification of this new developing phase has not been possible.

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To gain further information, alloys milled 16 h were submitted to DSc analysis. The crystallization temperature, Tx, was evaluated at the onset of the exothermal signals which appear well defined for all examined samples (see the inset of Figure 3 as an example). Our results are summarized in Figure 5 as open circles and compared with the crystallization data obtained by Politis and Johnson as well as with those relevant to CuTi amorphous ribbons /20,3,12/. The essential feature is the Tx drop relevant to the samples with lower Cu content, which reflects their different transformation path observed in the milling process.

Although some of the features so far observed find precise references in a variety of experimental results pertaining to MA process, the emerging framework gives rise to several questions. The substantial disagreement concerning the glass forming range accessible by MA is a central issue in this study. Our results also differ from the amorphisable compositions, extending from 30 to 75 at % Ti, relevant to the rapid quenching technique /20/, whereas it is generally maintained that the MA process is able to extend glass forming range. Moreover, a further question pertains to the lowering of the crystallization temperature observed for C U ~ ~ , C ~ ~ ~ a n d CuSo.

Figure 5

-

Crystallization temperature,T , versus composition of amorphous CuTi alloys. Open and solid circles refer to samples mifled 16 h at room temperature and 13 h at 250 K respectively (see text for explanation). Dotted line: milled sample from reference /5/.

Broken line: melt spinning samples /20/.

800.

750

C5 y 700

Y

650 600

Figure 6

-

C u K a X-ray patterns relevant to CU,~T~,,, Gu,,Ti,,, Cu,,Ti,, samples milled 13 h at 250K.

0

.."'....

...'.o "%.

S . . . . .

. ,----%?-

..."'

^,^l^--H@ \ 0

,

^"S.

... ".

..'

^/

...'

^\

.S..-. ' ⁰ t

2-

0

Focusing on this last item, it can be initially noted that, in the case of amorphous ribbons, the presence of crystallized nuclei was invoked to explain the lowering of activation energy needed for the accomplishment of the crystallization process /21/. Further results /22/ support this view, showing that the energy barrier for nucleation does overcome that for crystal growth. Schwarz and Petrich /19/ found these results to be rather general also in amorphous alloys synthesized by grinding powders of crystalline intermetallics. Therefore the formation of CuTiz and CuTi crystallites acting as heterogeneous nucleation centres during the first hours of milling can explain the corresponding lowering of crystallization temperatures. To this regard thermal analysis and X-ray results are in agreement with and sustain each other. As for their formation, we would advance a possible explanation involving the role of the morphology of the processed powders. The compressive nature of ball milling lead to the formation of a multilayered system of alternating bands of metal components:

a lamellar microstructure develops and continuously refines as milling proceeds / 2 3 , 2 4 / . It is now interesting to note that both crystalline phases

550;)

20 40 60 80 100 30 35 40 45 50 55

ATOMIC PERCENT COPPER 2 THETA

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C4-186 COLLOQUE DE PHYSIQUE

here involved, i.e. CuTiz and gamma-CuTi are constituted of alternating layers of Cu and Ti atoms along c. direction /25/. To some extent, their formation can be favoured by the natural evolution of the structural texture of refined powders.

On the other hand temperature rises during milling might be of paramount relevance. There is a spread of estimated values concerning this parameter starting from a T lower limit of 45 K /26/- 100 K /19/ up to the melting point of the powders involved. However T rises in typical alloy processed in a Spex apparatus were evaluated to be 1300 K /27/. Schultz /28/

has shown that the temperature can rise to the crystallization temperature of some NiZr glasses. This point is strictly connected with the monotonically descending trend of crystallization temperature in the

compositional range of interest here. It occurs below 650 K for both C U ~ ~ T ~ ~ ~ alloy prepared by MA and melt spinning (as quoted in Figure 5), and this

value is already competitive with the thermal spikes originating in the milling tool. Therefore, the drop of crystallization temperature relevant to low Cu samples can be of great concern since the hypothesis can be advanced that, in the present experimental conditions, the local temperature during milling process overcomes the crystallization temperature.

To pursue this consideration somewhat further, the starting pure powders were processed at about 250 K and an amorphous state for Cu4,., C U ~ ~ , was achieved after 13 h as it can be seen in Figure 6. A fraction of CuTi, compound is still evident in the pattern of C U ~ ~ . In any case the CuTi amorphisation range becomes wider and it seems likely that further cooling may enhance this effect.

The different structural states achieved under the two settled thermal conditions are testified by the new crystallization temperatures obtained in the same DSc conditions. It is noticeable that now the new trend, quoted as solid circles in Figure 5, adapts better to the data reported in the cited literature.

4

-

CONCLUDING REMARKS

The main results of this study are exemplified by the alloying path evidenced on C U ~ ~composition: T ~ ~ ~ after an initial stage where an amorphisation reactlon actually occurs, later the formation rate of CuTi and CuTiz intermetallics prevails. By cooling the milling vial at 250 K it is possible to inhibit the intermetallic nucleation in favour of the amorphous phase as it can be inferred from Tx values relevant to these samples. Thus, selection of a working temperature for the MA process effects the product formation pointing out the importance of kinetic constraints. Temperature increases during the process should sustain diffusive phenomena reducing the characteristic time for diffusion compared with that for nucleation and growth of crystalline compounds and therefore favouring the formation of the amorphous phase. However, an opposite behaviour is observed in the low Cu content range for sample milled at room temperature where the characteristic time for nucleation of CUT& and CuTi compounds appears favoured. On the other hand, it is interesting to remind the limited presence of CuTi, compound in the crystallization products of the 32 h milled C U ~ ~ T ~ ~ , , sample which looks almost completely X-ray amorphous (Figure 4). Thus we advance the hypothesis that the competing mechanism leading to the amorphous phase concerns the nature of the intermetallics which initially form.

This subject will be discussed in the forthcoming paper /29/.

ACKNOWLEDGEMENT

This work was financially supported by ENEA under Contract 3965 and by CNR-Rome, Contract 89.00623.68.

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