Crystallization behavior in inter-transition metal alloys

Journal of Non-Crystalline Solids 61 & 62 (1984) 943-948 North-Holland, Amsterdam

943

CRYSTALLIZATION BEHAVIORIN INTER-TRANSITIONMETALALLOYS A.F. Marshall, R.G. Walmsley, Y.S. Lee and D.A. Stevenson* Center for Materials Research and the Department of Materials Science, Stanford University, Stanford, CA 94305, USA The crystallization behavior of several amorphousCu-Zr and Cu-Ti alloys was characterized using differential scanning calorimetry, x-ray diffraction and transmission electron microscopy. Comparisonsbetween the two alloy systems and between two Cu6oTi40 alloys prepared by different quenching techniques were made. Transmission electron microscopy (TEM) analysis of the partially crystallized microstructures gave information on the modes of crystal nucleation and growth. 1.

INTRODUCTION Crystallization studies of metallic glasses can give insight into the

nature and relative s t a b i l i t y of the amorphous state.

The overall transforma-

tion sequence and kinetics, often involving the formation of intermediate metastable phases, may be obtained from calorimetric data and structural analysis by x-ray and electron diffraction.

In addition, details of nuclea-

tion and growth modes may be obtained by characterizing the partially crystallized microstructure in TEM.

I t is useful to compare the Cu-Zr and Cu-Ti

systems as both Zr and Ti are in the same column of the periodic table and both systems are similar in their phase diagrams and glass forming a b i l i t i e s . Cu6oTi40 was also prepared by both a liquid-quenching (LQ) and a vapor-quenching (VQ) technique in order to relate the crystallization behavior to the mode of synthesis. 2.

MATERIALSAND METHODS

The vapor-quenchedalloys, Cu6oZr40, Cu6oTi40, and Cu8oZr20, were prepared by planar magnetron sputter deposition onto rotating glass substrates and were on the order of 5-10 pm thick.

These were removed from the substrate for

subsequent experiments. A Cu60Ti40 alloy was also prepared by melt-spinning and was approximately 40 ~m thick. The compositions were confirmed by electron microprobe analysis,

The as-synthesized specimens were examined by

x-ray diffraction (XRD) and transmission electron microscopy (TEM) to insure that the i n i t i a l structure was amorphous. *This work was supported by the NSF-MRL Program through the Center for Materials Research at Stanford University. 0022-3093/84/$03.00 © Elsevier Science Publishers B.V. (North-Holland Physics Publishing Division)

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A.F. Marshall et al. / Crystallization behavior in inter-transition metal alloys

Calorimetric measurements were carried out in a Dupont 910 Differential Scanning Calorimeter (DSC) with the Model 1090 Programmer/Recorder. Partially crystallized specimens were obtained by continuous heating in the DSC or by isothermal annealing in a vacuum furnace. After annealing treatments specimens were thinned

for TEM by ion

milling

and were examined in a Philips 400

STEM/TEM with energy dispersive spectrometry capabilities. 3.

RESULTSAND DISCUSSION During continuous heating, both VQ1 and LQ Cu6oTi40 form an intermediate

microcrystal]Ine bcc phase; however during isothermal

annealing at

lower

temperatures transformation to the final phase appears to occur d i r e c t l y . is

difficult

to

explain why the microcrystalline transformation is

It not

observed at lower temperatures since i t would be expected to be more favored k i n e t i c a l l y as the temperature decreases.

I t is possible that the phase does

occur but with an even finer grain size than that obtained by continuous heating so that i t is d i f f i c u l t to distinguish i t from the amorphous structure by electron or x-ray d i f f r a c t i o n .

However, in the case of the VQ alloy compar-

ison of isothermal incubation times obtained for the overall transformation by DSC measurements, and for the formation of the ordered phase Cu3Ti2 by TEM measurements (Table 1), are in good agreement suggesting that the microcrystalline phase does not form at these temperatures. the mlcrocrystalline phase is

I t may simply be that

thermodynamically unstable

relative to the

amorphous state at the lower temperatures. Table 1:

Crystallization incubation times in minutes measured by DSC and TEM

T(°C) 378 379 384 390 393

DSC LQ - -

37.2 16.7 9.8 6.6

TEM

VQ 21.7 10.6 5.4 3.5

23.2 -9.8 7.4 --

Such an intermediate phase is not observed for the isocomposittonal Cu6oZr402, 3 during e i t h e r continuous heating or isothemal annealing, presum-

ably because of the larger size difference between the Cu and Zr atoms relat i v e to that for Cu and Tt.

Cu6oZr40 transfoms d i r e c t l y to the equilibrium

phase CulOZr 7. Whereas the Cu3Ti2 phase is ordered in one dimension and is s t r u c t u r a l l y related to the intermediate bcc phase, CulOZr 7 has a large u n i t c e l l ordered in three dimensions and is not based on a simple subunit. At another composition however, the Cu-Zr system does show transfomatton to a

A.F. Marshall et al. / Crystallization behavior in inter-transition metal alloys

a

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Cu8oZr20 S!

?

326 . . . . .

~6o. . . . . . . 486 . . . . . TEMP(°C)

~6o

Figure 1: (a) DSC scan of Cu8oZr20 (b) Characteristic diffraction pattern of microcrystalline CusoZr20 obtained by continuous heating to 400°C. simple intermediate structure.

Cu8oZr20,which can be made amorphous only by

vapor quenching, gives a complex DSC scan involving i n i t i a l transformation to a microcrystalline phase,4 as shown in Figure i .

This structure appears to be

close-packed and may have a preferred orientation in the plane of the film. Since this composition lies between those of the final equilibrium phases, i t is also possible that the microcrystalline phase represents the beginning of phase separation prior to the next crystallization step which is nucleation and growth of the ordered Cu3.6Zr phase (21.7 a/o Zr). The final microstructure is largely Cu3.6Zr with ~-Cu at the grain boundaries. Partially crystallized microstructures of both Cu6oZr40 and Cu6oTi40 (VQ) show a change from homogeneousto heterogeneousnucleation as the temperature is lowered.

Homogeneousnucleation during continuous heating and at the

higher isothermal annealing temperatures is indicated by isolated crystall i t e s with a continuous size distribution (Figure 2, a&b).

In the case of

Cu6oZr40 the few crystallites observed at the lowest annealing temperature are of the same size suggesting a one-time heterogeneousnucleation event with few nucleation sites

(Figure

2c).

"pre-nuclei" configurations3.

This

is

presumed due to

quenched-in

Cu6oTi40 shows an increases in heterogeneous

interface nucleation resulting in clusters of crystallites as the temperature is lowered (Figure 2d).

Howeverhomogeneousnucleation s t i l l predominates in

the temperature range studied, as indicated by isolated crystallites of varying size. Presumablyas the temperature is lowered s t i l l further, the nucleation modew i l l becomet o t a l l y heterogeneousand dependenton quenched-in configurations as in the case of Cu6oZr40. Although LQ and VQ Cu6oTi40 show similar transformation steps, the LQ alloy is more stable as indicated by a higher crystallization temperature when continuously heated, and a longer incubation time when isothermally annealed

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A.F. Marshall et al. / Crystallization behavior in inter-transition metal alloys

Figure 2: (a,b) Homogeneous nucleation in Cu60Zr40 and Cu6oTi40 during continuous heating; (c) isolated crystal due to heterogeneous nucleation in Cu6oZr40 and (d) heterogeneous interface nucleation in Cu6oTi40 during isothermal annealing at lower temperatures. (Table I).

In studying the partially crystallized material in TEM, s i g n i f i -

cant differences are found in crystallite morphology, orientation and structure.

VQ Cu6oTi40 transforms polymorphicaIly into the expected equilibrium

phase, Cu3Ti2, whereas the LQ Cu6oTi40 shows a preponderance of the nonstoichiometric Cu4Ti3 (43 a/o Ti). These phases are closely related in structure and composition and can be distinguished by their superlattice reflections in TEM microdiffraction patterns, as shown in Figure 3.

Lattice images

of f u l l y annealed LQ and VQ Cu6oTi40 show that the two phases can grow coherently in a single grain and that there is always a small amount of the alternate phase present.

B

b

Figure 3: <110> microdiffraction patterns of (a) Cu3Ti2 with ten superlattice reflections t o the f i r s t fundamental reflection, (0,0,10), and (b) Cu~Ti3 with seven superlattice reflections to the f i r s t fundamental reflectlon, (0,0,14).

A.F. Marshall et al. / Crystallization behavior in inter-transition metal alloys

947

The Cu6oTi40 VQ crystals grow in a disk shape, with the tetragonal axis normal to the disk.

Most of the crystals in a given VQ specimen actually

appeared torpedo-shaped which is a cross-section of the disk and indicates preferred orientation. axis in

When the crystals are oriented with the tetragonal

the viewing plane, a boundary is observed dividing the torpedo

lengthwise (Figure 4a).

This boundary is not sharp.

site of crystal nucleation and preferred growth.

I t appears to be the

I t may represent a twin or

stacking fault boundary and may be related to the non-equiaxed disk-shaped rowth.

b Figure 4: Isothermally formed crystallites of (a) VQ, and (b) LQ Cu6oTi40 showing a perpendicular cross-section of Cu3Ti2 disk-shaped crystal with bisecting boundary in (a) and spherulitic growth of Cu4Ti3 in (b) The LQ crystals, on the other hand, appear equiaxed, show no indication of preferred orientation, and appear to have a spherulitic growth structure (Figure 4b).

Such growth is probably related to the formation of the non-

stoichiometric Cu4Ti3 phase. The differences in crystallization temperature and incubation time between the VQ and LQ alloys indicate that the LQ alloy is more structurally relaxed.

This may take the form of increased short range

order (SRO) and, in the context of chemical short range order (CSRO) may represent short range compositional inhomogeneities which are the beginning of amorphous phase separation.

This may well lead to primary nucleation of a

non-stoichiometric phase and a different growth morphology. Recently the occurrence of CSROin the inter-transition metal systems has received increasing attention. Results obtained from neutron scattering data on amorphous Ni-Ti 5 and Cu-Ti6 show a pre-peak in the radial distribution function indicating CSRO. EXAFSstudies on Cu-Zr and Cu-Ti glasses also show CSRO7'8. Several authors 5,9 have pointed out that such ordering may be related to ease and s t a b i l i t y of glass formation and must be taken into account in developing metallic glass theories that include inter-transition metal systems.

In addition, Sakata et al. 10 have demonstrated that CSRO

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A.F. Marshall et al. / Crystallization behavior in inter-transition metal alloys

occurs in Cu66Ti34 in the liquid state and appears to be enhanced in the glassy state.

Sadoc et al. 7 find similar CSROin comparing EXAFSdata from a

liquid-quenched and a sputtered Cu6oZr40 alloy, as well as in comparing the amorphous and crystalline data for this composition.

These similarities may

correlate with the crystallization sequence, and apparent relative s t a b i l i t y of VQ vs. LQ Cu6oZr40 as discussed by Walmsley et al. 3 The fact that Cu6oTi40 shows differences between the LQ and VQ alloy in CSRO and in the crystal phases formed may correlate with the opposite pattern of s t a b i l i t y observed in this

system as discussed above.

One may also find changes in relative

s t a b i l i t y with temperature due to corresponding changes in the mode of transformation (e.g., homogeneousvs. heterogeneousnucleation).

I t is clear

that more detailed correlation of crystallization studies with structural studies, and particularly with analysis of CSROin the liquid, glass (prepared by different quenching techniques), and metastable and stable crystalline states, may further our understanding of metallic glass formation. REFERENCES 1) A.F. Marshall, Y.S. Lee and D.A. Stevenson, Acta Met. (1983) in press. 2) J.M. Vitek, J.B. Vander Sande, and N.J. Grant, Acta Met. 23 (1975) 165. 3) R.G. Walmsley, A.F. Marshall, D. Bouchet and D.H. Stevenson, J. NonCrystalline Solids 54 (1983) 277. 4) A.F. Marshall, R.G. Walmsley and D.A. Stevenson, accepted for publication in Mat. Sci. and Eng., 1983. 5) H. Ruppersberg, Dokyol Lee and C.N.J. Wagner, J. Phys. F: (1980) 1645. 6) M. Sakata, N. Cowlam and H.A. Davies, J. Phys. F: L235. -ibid, J. de Physique C8 (1980) 190.

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7) A. Sadoc, D. Raoux, P. Lagarde and A. Fontaine, J. Non-Crystalline Solids 50 (1982) 331. 8) D. Raoux, J.F. Sadoc, P. Lagarde, A. Sadoc, and A. Fontaine, J. Physique C8 (1980) 207. 9) M. Sakata, N. Cowlam and H.A. Davies, Proc. 4th Int}. Quenched Metals, Sendal, Japan 1981. 10) M. Sakata, N. Cowlamand H.A. Davies, J. Phys. F: L157.

de

Conf. Rapidly

Metal Phys 11 (1981)