Short-range order of Zr62−xTixAl10Cu20Ni8 bulk metallic glasses

Acta Materialia 50 (2002) 305–314 www.elsevier.com/locate/actamat

Short-range order of Zr62⫺xTixAl10Cu20Ni8 bulk metallic glasses N. Mattern a

a,*

, U. Ku¨hn a, H. Hermann a, H. Ehrenberg b, J. Neuefeind c, J. Eckert a

Institut fu¨r Festko¨rper- und Werkstofforschung Dresden, Postfach 270016, D-01171 Dresden, Germany b TU Darmstadt, Material- und Geowissenschaften, Petersenstr.23, D-64287 Darmstadt, Germany c HASYLAB am DESY, Notkestr.85, D-22603 Hamburg, Germany Received 9 July 2001; received in revised form 14 September 2001; accepted 17 September 2001

Abstract The short-range order and crystallization behavior of slowly cooled Zr62⫺xTixAl10Cu20Ni8 bulk metallic glasses have been investigated in terms of the atomic pair correlation function as a function of Ti content x (2ⱕxⱕ7.5). The structural parameters point to the presence of chemical short-range order in these bulk glasses. An enhanced local excess free volume around the Ti atoms is concluded from density measurements. The first stage of crystallization in Zr62⫺xTixAl10Cu20Ni8 bulk glasses is related to changes in the medium-range order while the first neighborhood is retained. The atomic pair correlation functions of the first crystallization products are similar for all titanium contents. There is no indication of any special atomic arrangement for the particular alloy forming quasicrystals upon heating (x=3). In case of Zr54.5Ti7.5Al10Cu20Ni8 an ultrafine microstructure consisting of clusters of 2 nm in size is formed as the first step of crystallization.  2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Metallic glasses; X-ray diffraction (XRD); Phase transformations; Order–disorder phenomena

1. Introduction Amorphous Zr-based alloys like Zr65Al7.5Cu17.5Ni10 belong to the new class of bulk metallic glasses showing a wide supercooled liquid region [1–4]. Changes in chemical composition or addition of further elements such as, for example, Pd, Fe or Ti lead to altered thermal stability, different crystallization modes as well as different

* Corresponding author. Fax: +49-351-4959-452. E-mail address: [email protected] (N. Mattern).

phases upon annealing/crystallization of the glass [5–8]. In particular, it has been recently reported that an icosahedral phase forms as a metastable phase in the first crystallization stage of different Zr-based amorphous alloys [9–16]. Because icosahedral clusters are stable in liquids and glasses [17] it has been postulated that metallic glasses may consist of randomly oriented icosahedra [18]. However, there is only little information until now concerning the short-range order and its relation to the crystallization behavior for multi-component bulk metallic glass-forming liquids. Also no direct experimental evidence for

1359-6454/02/$22.00  2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. PII: S 1 3 5 9 - 6 4 5 4 ( 0 1 ) 0 0 3 4 1 - X

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N. Mattern et al. / Acta Materialia 50 (2002) 305–314

the existence of icosahedral-type short-range order or clusters in bulk metallic glasses/supercooled liquids has not been given yet. In this work we analyze the dependence of the short-range order of Zr62⫺xTixAl10Cu20Ni8 bulk metallic glasses as a function of the Ti content x (2ⱕxⱕ7.5) in the as-cast state as well as after annealing. This series of alloys was chosen because an icosahedral quasicrystalline phase forms for the alloy with x=3 in the first stage of crystallization [10]. In this study the atomic pair correlation functions of the glassy as-cast samples are compared with the data obtained after the first stage of crystallization. This allows us to gain insight concerning the degree of structural pre-formation of distinct structural units (polyhedra) in the supercooled liquid state. It will be shown that there is no hint for a special atomic arrangement indicating distinct icosahedral short-range order in the supercooled liquid even for the alloy forming quasicrystals upon crystallization. However, crystallization is characterized by rearrangements of chemically defined structural units in the bulk glass.

sive electron scanning microprobe (ARL quantometer) and by hot carrier gas extraction (C436 Leco analyzer), revealing that their stoichiometry matched the desired composition within the error limits of the analysis methods. The DSC experiments were performed using discs of 1 mm in height employing a Netzsch DSC 404 calorimeter at a heating rate of 40 K/min. Besides continuous-rate heating scans up to elevated temperatures in order to fully crystallize the material, also heating runs with 40 K/min up to the characteristic temperature above the first exothermic crystallization event were carried out to induce primary crystallization. After reaching the respective temperature, these samples were immediately cooled down to room temperature at ⫺100 K/min. The mass density was determined by the Archimedes principle by weighing samples in air and in dodecan (C12H26). To analyze the short- and mediumrange order X-ray diffraction experiments were conducted by synchrotron radiation at the beamlines B2 and BW5 at the storage ring DORIS (HASYLAB, Hamburg) using a wavelength of l=0.03749 nm and 0.01078 nm, respectively.

2. Experimental

3. Experimental results and discussion

Prealloyed ingots were prepared by arc-melting elemental Zr, Ti, Al, Cu and Ni with a purity of 99.9% or better in a Ti-guttered argon atmosphere. For reaching a high homogeneity, the samples were remelted several times. From these prealloys, cylindrical Zr62⫺xTixAl10Cu20Ni8 (x=2, 3, 4, 5, and 7.5) bulk glassy samples with 3 mm diameter and 50 mm length were prepared by copper mold casting using a modified Bu¨ hler melt-spinning device. The amorphous structure of the samples was checked by standard X-ray diffraction (Philips PW 1050 diffractometer using Co-Kα radiation) and transmission electron microscopy (Jeol 2000 FX microscope operated at 200 kV). The cylinders with x=2, 3, 4, 5, and 7.5 proved to be fully amorphous in the as-cast state. In contrast, for x=0, 1, and 10 all attempts resulted in partially crystalline. Therefore, these samples were not considered further for investigation in this study. The chemical composition was checked by wavelength-disper-

3.1. Short-range order Fig. 1 compares the interference functions I(q) of amorphous Zr62⫺xTixAl10Cu20Ni8 alloys in the as-prepared state (q = 4p sin q/l is the absolute value of the scattering vector). The interference curve I(q) is obtained by extracting the elastic scattering intensity in absolute electron units Ie.u. from the measured diffraction curve using the corrections for absorption, multiple scattering and Compton scattering as described in Ref. 19. The I(q) curves of the as-cast samples show a behavior typical of amorphous metallic alloys with a shoulder in the second diffuse maximum. The first maximum at q1 shifts to higher q-values with decreasing zirconium content as shown in Fig. 2. A lower slope dq1/dx is observed for the Zr62⫺xTixAl10Cu20Ni8 alloys compared with the behavior of q1 versus x of binary Zr100⫺xMx amorphous alloys with M=Cu or Ni [20], as well as for

N. Mattern et al. / Acta Materialia 50 (2002) 305–314

307

From the interference functions I(q) the atomic pair correlation functions g(r) = r(r)/r0 were calculated by the Fourier transform of I(q) between 0ⱕqⱕ150 nm⫺1 according to Ref. 19: 4prr0(g(r)⫺1) ⫽

Fig. 1. Interference functions (I(q)⫺1)q of Zr62⫺xTixAl10Cu20Ni8 bulk metallic glasses.

冕

2 I(q)q sin(qr) dq p

(1)

where r(r) is the atomic pair density distribution function and r0 is the mean atomic density. Fig. 3 shows the calculated g(r) curves. The atomic pair correlation functions of the Zr62⫺xTixAl10Cu20Ni8 bulk glasses are very similar for 2ⱕxⱕ7.5. For all g(r) curves, two components of the first maximum are visible. The second maximum consists of at least three components. The measured interference function I(s) and the estimated g(r) curves represent in the n-component alloy the weighted sum of the partial functions Iij(s) and gij [19]:

冘冘 冘冘 冘

I(s) ⫽

wijIij and g(r)

i

⫽

j

wijgij

i

with wij

(2)

j

⫽ cicjfifj/(

cifi)2

i

Fig. 2. Position of the first maximum in I(q) of bulk metallic glasses versus zirconium content.

(Zr65Al7.5Cu17.5Ni10)100⫺xFex [21] (Fig. 2). The continuous shift of the first diffuse maximum with decreasing zirconium content points to a solid solution-like exchange between zirconium and the metal atoms in the glassy state. In contrast to crystalline materials the position of q1 is not directly correlated to an interatomic distance and this linear shift does not necessarily indicate altered nearest neighbor distances.

Fig. 3. Pair correlation functions g(r) of Zr62⫺xTixAl10Cu20Ni8 bulk metallic glasses.

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N. Mattern et al. / Acta Materialia 50 (2002) 305–314

The weights wij given in Table 1 depend on the concentration ci and the atomic form amplitude fi of the i-atom. The total X-ray pair correlation function is determined mainly by the Zr–Zr and Zr– metal contributions. The metal–metal pairs contribute by ⬇10% to the total atomic pair correlation function g(r) and can be neglected in a first approximation. The identification of the maxima in the total pair correlation function can by derived from comparison with crystalline counterparts. However, this is not unambiguous. The partial distances and coordination numbers vary in the crystalline Zr2M and ZrM phases [22], as given in Table 2. In the case of the tetragonal NiZr2, tetragonal CuZr2, cubic CuZr, and orthorhombic NiZr phases the distances rZr–M and rZr–Zr are separated in the total pair correlation function. For the cubic NiZr2 phase, 2/3 of the metal–zirconium distances are rather close to the zirconium–zirconium distance. The position of maxima in the first neighborhood (Table 1) can be, therefore, attributed to the zirconium–metal (r1Zr⫺M) and to the zirconium–zirconium (r1Zr⫺Zr) distances only in a first approximation. The integration over the total atomic pair density function gives a weighted, averaged nearest neighbor number N1 in the interval between r1 and r2:

冕 r2

N1 ⫽ 4pr2r(r)

(3)

r1

The estimated values of N1 of about 13 (Table 1) are characteristic for densely-packed intermetallic compounds. The positions and the integrals of the two components of the first maximum in g(r) were estimated by a fit of two gaussian curves to the radial distribution function 4pr2r(r). The partial coordination number N1ij can be estimated from the weighted nearest neighbor numbers N1 under the presumption that individual atomic pairs are resolved in the total g(r) curve, i.e. N1ij ⫽ cjN1/wij.

(4)

This is probably not the case for the present alloys because Al and Ti have a large atomic diam-

eter compared to Cu and Ni. If we assume that the first component originates from zirconium–metal pairs we obtain about two metal neighbors around a zirconium atom at r1=0.268 nm. Further contributions of other zirconium–metal distances are probably superimposed on the zirconium–zirconium distribution. Therefore, the estimate of the partial nearest neighbor number N1Zr⫺Zr from the second component gives only approximate values. Nevertheless, the values of the nearest neighborhood are different from that reported for binary amorphous Zr100⫺xMx alloys with similar zirconium content [23–25] as given in Table 2. Fig. 3 also shows the g(r) curve of another alloy derived from the bulk glass-forming Zr65Al7.5Cu17.5Ni10 alloy which was obtained by adding 20 at.% iron. This Zr52Al6Cu14Ni8Fe20 alloy has a rather similar Zr content than the Ti-containing sample with x=7.5. The g(r) function is clearly different especially in the first neighborhood. On the other hand, the parameters of the nearest neighborhood are comparable to those of binary amorphous Zr55M45 alloys [25]. This indicates the existence of chemical short-range order in the Ticontaining bulk glasses. In particular, there is no sign of any essential change in the general behavior of the g(r) curves of the Zr62⫺xTixAl10Cu20Ni8 bulk glass for different titanium contents. This can be understood by an isomorphous replacement of Zr by Ti. The reduction of the fraction of gZr–Zr with decreasing zirconium content is compensated by the enhancement of the weight of the gZr–Ti contributions. The Zr59Ti3Al10Cu20Ni8 glass crystallizes into a primary metastable quasicrystalline phase, as has been shown previously [10] and will be discussed in Section 3.3. However, there is no indication of any difference in the g(r) curve of this particular alloy in the as-cast state compared to the other bulk glasses not forming quasicrystals during crystallization. The formation of quasicrystals is, therefore, not directly correlated to the existence of a special short-range order arrangement of icosahedral clusters in the supercooled liquid. On the other hand, the existence of icosahedral units at all can not be excluded from the present experimental results. However, in the case of their presence it would not be limited to x=3.

x=7.5

x=5

x=4

x=3

x=2

26.05

26.02

25.90

25.97

25.95 ±0.01

0.266

0.268

0.269

0.269

0.268 ±0.02

r1Zr⫺M

(nm⫺1)

2wZr–M wM–M

0.505 0.410 0.085 0.493 0.418 0.089 0.482 0.424 0.094 0.471 0.432 0.097 0.442 0.446 0.112

r1 (nm)

q1

wZr–Zr

0.312

0.313

0.313

0.313

0.313 ±0.02

r1Zr⫺Zr

2.3

2.2

2.1

1.9

2.3 ±0.2

N11

N1

13.4 11.1

13.3 11.1

13.5 11.4

13.4 11.1 ±1.0 13.6 11.7

N21

679

662K

660

655

654 ±1

(K)

Tg

709

714

716

719

758 ±1

(K)

Tx1

6.47 (as-cast) 6.59 (740 K) 6.63 (873 K)

6.57 (as-cast) 6.66 (743 K) 6.54 (as-cast) 6.65 (737 K)

6.62 (as-cast) 6.66 (808 K) ±0.01 6.59 (as-cast) 6.67 (750 K)

(g cm⫺3)

s

Table 1 Weights of partial functions wij, position of the first maximum q1 of I(q), nearest neighbor distances rij and numbers Nij, onset temperatures of glass transition Tg and crystallization Tx, and mass density s of Zr62⫺xTixAl10Cu20Ni8 alloys (typical error bars given for x=2)

N. Mattern et al. / Acta Materialia 50 (2002) 305–314 309

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N. Mattern et al. / Acta Materialia 50 (2002) 305–314

Table 2 Nearest neighbor distances rij and numbers Nij of amorphous and crystalline binary M–Zr compounds rij (nm)/(Nij) M–M

Zr–M

Zr–Zr 0.321/(9.3) 0.326/(7.4) 0.316/(8.6) 0.332/(7.8)

Zr60Co40 [25] Zr55Co45 [25] Zr63.5Ni36.5 [23] Zr50Ni50 [24] t-NiZr2 S.G. 140

0.263/(3.3)

0.267/(7.8) 0.271/(7.4) 0.270/(5.6) 0.273/(6.7)

0.262/(2)

0.275/(4)

0.298/(1), 0.306/(2), 0.336/(4), 0.342/(4)

t-CuZr2 S.G. 139

0.321/(4)

0.282/(4)

0.318/(4), 0.321/(4), 0.334/(1)

c-NiZr2 S.G. 227

0.273/(3)

0.261/(1.5), 0.296/(1.5), 0.305/(1.5)

0.306/(3), 0.318/(3), 0.339/(3)

o-NiZr S.G. 63

0.261/(2), 0.327/(2)

2.68/(4), 2.74/(2), 0.277/(1)

0.327/(2), 0.342/(4), 0.344/(2)

0.325/(6)

0.282/(8)

0.325/(6)

0.255/(5.1)

0.285/(6.7)

0.315/(8.5)

c-CuZr S.G. 221 DRPHS Zr50Ni50

The comparison of the parameters of the first neighborhood with the values for the corresponding crystalline MZr2 and MZr phases (Table 2) shows similar values only for the nearest neighborhood in the case of the cubic NiZr2 phase. Zirconium has 1.5 nearest metal neighbors at 0.261 nm in this structure. The other nearest neighbors at 0.296 nm and 0.306 nm are superposed by the distances of the Zr–Zr shell. To compare the short-range order of the glasses with that of the corresponding crystalline phases a model for a para-crystalline approximation of the amorphous structure was applied [26]. The atomic pair correlation functions of the crystals are smeared out by Gaussian functions. Fig. 4 compares the calculated g(r) functions of crystalline-like structure models based on the corresponding MZr2 phases. The short-range order of tetragonal CuZr2 and tetragonal NiZr2 is quite different from the experimental results. The short-range order of the cubic NiZr2 phase (S. G. No. 227) shows similarity

Fig. 4. Pair correlation functions g(r) of para-crystalline structure models.

N. Mattern et al. / Acta Materialia 50 (2002) 305–314

to the experimental curves only for the first neighborhood. The second maximum in g(r) clearly differs in shape. Therefore, the atomic structure of the amorphous alloys is also different from that of cubic NiZr2 but could contain structure elements. In many intermetallic alloys the atoms are arranged in a complex packing of polyhedra [27,28]. Due to tendency of chemical ordering the formation of such polyhedra may be a common feature in the undercooled metallic melt. The counterpart of chemically defined structural units is the model based on dense random packing of hard spheres (DRPHS) [29]. Fig. 4 shows the atomic pair correlation function of a binary Zr50Ni50 model system. For this model system, the behavior of g(r) is clearly different from the experimental observations for the Ti-containing bulk glasses, which strengthens the picture of chemical short-range order in the amorphous Zr– Ti–Al–Ci–Ni alloys. 3.2. Average atomic volume and excess free volume The mass density s of an amorphous alloy and its behavior upon variation of the chemical composition can give information about the distribution of excess free volume in the material [30]. The mass density of the amorphous Zr62⫺xTixAl10Cu20Ni8 alloys is given in Table 2. We will discuss these data in terms of the average atomic volume V¯ i which can be attributed to the atoms of type i with i=Zr, Ti, Al, Cu, Ni. If the coordinates of all atoms are given, e.g. for a computer generated structure model, the value of V¯ i can be calculated using the radical plane tessellation method introduced by Fisher and Koch [31] and applied to a model for amorphous Pd75Si25 by Gellatly and Finney [32]. Here, we assume that the short-range order of each type of atom does not essentially change for the present alloy if the Ti content x is varied in the range 2ⱕxⱕ7.5. This assumption is justified by the X-ray scattering data given in the previous section. Therefore, the values of V¯ i are characteristic quantities and the volume of a sample can be written as [33,34]: NV¯ ⫽

冘 i

ciV¯ i

(5)

311

where N is the total number of atoms, V¯ is atomic volume averaged over all atoms and ci is the fraction of i-type atoms. Analogously, the average mass m ¯ of the sample is given by: m ¯ ⫽

冘

cimi

(6)

i

where mi is the mass of an i-type atom. Inserting definitions (4) and (5) into V¯ = m ¯ /s one obtains the average atomic volume V¯ (x) 76.66⫺(mZr⫺mTi)x/100 s(x) (1.66×10⫺27 kg)

V¯ (x) ⫽

(7)

Fig. 5 shows the estimated values of the average atomic volume as a function of the titanium content. The behavior of V¯ (x) can be described by linear dependence: V¯ (x) ⫽ V¯ 0 ⫹ (V¯ Zr⫺V¯ Ti)x/100

(8)

Here, V¯ 0 is the average atomic volume of the alloy without Ti (x=0). A least-square fit of expression Eq. (8) to the experimental data (solid ˚ 3, and line in Fig. 5) gives V¯ Zr⫺V¯ Ti = 3.15 A 3 ˚ which is the mean atomic volume V¯ 0 = 19.05 A of the alloy without Ti (x=0). The difference in atomic volume between Zr and Ti as calculated

Fig. 5. Mean atomic volume V¯ (x) of Zr62⫺xTixAl10Cu20Ni8 alloys versus Zr-content x.

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N. Mattern et al. / Acta Materialia 50 (2002) 305–314

˚ ; Ti: 1.448 from the atomic diameter (Zr: 1.60 A 3 ˚ ˚ A) is 4.45 A . The 50% larger value for Ti means that the titanium atoms have a larger atomic volume in the bulk glass. For binary amorphous Zr100⫺xMx alloys with M=Cu or Ni an almost linear V¯ (x) dependence has been reported [25] which is also shown in Fig. 5. The difference between the average atomic volume ˚ 3 for M=Cu and Ni, is about V¯ Zr⫺V¯ M = 11.5 A respectively. The difference in atomic volume cal˚ ; Ni: culated from the atomic diameter (Cu: 1.278 A ˚ 3 for ˚ ) gives 8.42 A ˚ 3 for Cu, and 9.06 A 1.246 A Ni, respectively. The 20% smaller V¯ value means a smaller atomic volume which is due to the dense packing of atoms with different diameter. Using the definition of the local packing density [34]: hi ⫽

4p/3r3i V¯ i

(9)

˚3 introduced in Ref. 31, the result V¯ Zr⫺V¯ Ti = 3.15 A can be written as hTi ⫽

hZr 1.25⫺0.24hZr

(10)

Considering that reasonable values of the packing density in metallic alloys are in the range from 0.65 to 0.8 one obtains that the packing density of Ti is about 5% to 10% smaller than that of Zr. This means that the excess free volume of the present alloy is located mostly around the Ti atoms. The results agree well with recent data from positron annihilation spectroscopy of a bulk metallic glass where open-volume Zr52Ti5Al10Cu17.9Ni14.6, regions localized around Ti were reported [35].

Fig. 6. DSC scans of Zr62⫺xTixAl10Cu20Ni8 bulk glasses (heating rate 40 K/min).

crystallization and stepwise transformation into the equilibrium compounds. Fig. 7 shows the structural features for samples heated to temperatures above the first DSC peak by displaying the corresponding interference functions. The formation of different crystalline phases at small variation of the Ti (Zr)

3.3. Medium-range order and crystallization Fig. 6 shows the DSC scans for the as-cast amorphous Zr62⫺xTixAl10Cu20Ni8 bulk samples. Zr62Ti2Al10Cu20Ni8 (x=2) crystallizes eutectically via one sharp exothermic peak by the simultaneous formation of tetragonal CuZr2 and a quasicrystalline phase, as proved by X-ray diffraction and TEM (not shown here in detail). Upon Ti addition, the crystallization mode changes towards primary

Fig. 7. Interference functions I(q) of Zr62⫺xTixAl10Cu20Ni8 bulk metallic glasses after heating through the first DSC peak.

N. Mattern et al. / Acta Materialia 50 (2002) 305–314

content is manifested especially in the change of the first maximum of the I(q) curves. Additional reflections compared to the as-cast states point to structural transformations as expected from the exothermic DSC peak. The crystallization is also related to an increase of the mass density by 1–2% as given in Table 1. Annealing of Zr59Ti3Al10Cu20Ni8 (x=3) leads to primary precipitation of icosahedral quasicrystals with a size of about 50–100 nm as reported previously [10]. For xⱖ4, the diffraction peaks become weaker in intensity and broader because the size of the precipitates decreases [10]. The shape of the first non-split maximum of the annealed Zr57Ti5Al10Cu20Ni8 ˚ ⫺1 (x=5) alloy with a second component at q=3.1 A indicates the formation of another phase different from nano-quasicrystals. This is similar to what is known for other Zr-based alloys [11,12]. In the case of the alloy with x=7.5 the change of the first maximum in I(q) is only a small enhancement of the height of the maxima compared to the as-cast state. More pronounced differences in I(q) become visible at the second diffuse maximum and for higher q-values. The interference functions of the annealed alloys are quite different in the first maximum in I(q) for the different compositions. On the other hand, the I(q) curves are rather similar in their features for higher q-vales (q⬎40 nm⫺1). Due to the properties of the Fourier transform this means differences in medium-range order, but similarities in short-range order. This is seen in the calculated pair correlation functions shown in Fig. 8. The g(r) curves of the annealed samples are similar for the different Ti-contents. Compared with the as-cast state the first maximum in g(r) becomes somewhat sharper for the annealed samples. However, the position and area remain unchanged within the error limits. The second coordination shell is clearly changed. The positions and the relative heights of the three submaxima are different from that of the amorphous samples. The same observations are found for the other maxima. These features strongly indicate that the medium-range order changes during crystallization with consequences to the second neighborhood. The similarity of the g(r) curves confirms the presence of similar coordination polyhedra in the titanium-containing bulk glasses also after the first

313

Fig. 8. Pair correlation functions g(r) of Zr62⫺xTixAl10Cu20Ni8 bulk metallic glasses after heating through the first DSC peak.

stage of crystallization. The comparison with the atomic pair correlation function of the cubic NiZr2phase is also shown in Fig. 8. There is a good agreement for the first and second shell with the experimental data for the Ti-containing alloys. On the other hand, there are differences in the maxima in g(r) for higher r-values which may be correlated to different crystal structures of the phases formed. In the case of the alloys with xⱖ4 the broadening of the interference function and the behavior of the pair correlation function at higher r-values indicates the formation of a nanostructure after the first DSC peak. The correlation length rc (value of a distance r where for all values r⬎rc the deviation of the atomic density from the average density is less than 1%) has a value rc = 1.4 nm for the ascast state. This correlation length increases only slightly to rc = 1.6 nm after heating the Zr54.5Ti7.5Al10Cu20Ni8 sample over the first DSC peak (740 K). An ultra-fine microstructure consisting of regions of less than 2 nm in size embedded in the remaining amorphous phase is formed as the first step of crystallization. The excess free volume is reduced by the crystallization as indicated by an increase of the mass density (Table 1). The estimated values of the average atomic volume are also shown in Fig. 5 for samples annealed through the first DSC peak. The

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N. Mattern et al. / Acta Materialia 50 (2002) 305–314

behavior means that the excess free volume around the Ti atoms becomes annihilated during crystallization.

4. Conclusions The short-range order of Zr62⫺xTixAl10Cu20Ni8 bulk metallic glasses (2ⱕxⱕ7.5) is characterized by an extended chemical short-range order in the as-cast state. Small variations in the titanium content lead to different metastable phases with similar short- and medium-range order upon annealing. Differences between the structure of the amorphous and the crystalline or quasicrystalline states are due to altered connectivity of the polyhedra being the building blocks of the different phases. On the other hand, we observe the formation of nanoclusters without remarkable growth in the first step of crystallization for Zr62⫺xTxAl10Cu20Ni8 bulk metallic glass. An ultrafine microstructure consisting of clusters with cubic NiZr2 crystal-like regions 2 nm in size is formed by the crystallization of glassy Zr54.5Ti7.5Al10Cu20Ni8. The differences in the arrangement of the coordination polyhedra are correlated to the Ti-content which, on the other hand, is responsible for the amount of excess free volume in the alloy. Therefore, the result that different crystalline phases appear upon annealing of alloys with different Ti content, can be understood as a consequence of the different amount of excess volume.

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